Polymer Blends Handbook, Volume 1 and Volume 2

POLYMER BLENDS HANDBOOK

POLYMER BLENDS HANDBOOK Volume 1

Edited by

L. A. Utracki

KLUWER ACADEMIC PUBLISHERS DORDRECHT / BOSTON / LONDON

A C.I.P. Catalogue record for this book is available from the Library of Congress.

ISBN 1-4020-1110-5 Volume 1 ISBN 1-4020-1111-3 Volume 2 ISBN 1-4020-1114 -8 Set

Published by Kluwer Academic Publishers, P.O. Box 17, 3300 AA Dordrecht, The Netherlands. Sold and distributed in North, Central and South America by Kluwer Academic Publishers, 101 Philip Drive, Norwell, MA 02061, U.S.A. In all other countries, sold and distributed by Kluwer Academic Publishers, P.O. Box 322, 3300 AH Dordrecht, The Netherlands.

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LIST OF AUTHORS Dr. ABDELLAH AJJI National Research Council Canada, Industrial Materials Institute, 75 de Mortagne, Boucherville, QC, Canada J4B 6Y4, Tel: +1 (450) 641-5244; Fax: +1 (450) 641-5105 E-mail:

Dr. MICHEL M. DUMOULIN National Research Council Canada, Industrial Materials Institute, 75 de Mortagne, Boucherville, QC, Canada J4B 6Y4, Tel: +1 (450) 641-5181; Fax: +1 (450) 641-5105 E-mail: <[email protected]>

Dr. KRIS AKKAPEDDI Honeywell International, Nylon Systems, Specialty Materials Division 101 Columbia Rd. Morristown, NJ 07962,USA Tel: +1 (973) 455-3231; Fax: +1 (973) 455-4892 E-mail:

Dr Eng. VALIJA EVERAERT Centexbel Technologiepark 7 B-9052 Zwijnaarde, Belgium Tel:: +32 (9) 2438235; Fax: +32 (9) 2204955 E-mail:

Dr. KAMARUDIN BAHARI Malaysian Institute for Nuclear Technology Research, Radiation Processing Technology Division, Bangi, 43000 Kajang, Malaysia Tel: +60 (3) 8925-0510; Fax: +60 (3) 8920-2968 E-mail:

Dr. S. BRUCE BROWN General Electric Co. R&D Center, K-1/3A71, One Research Circle, Niskayuna, NY 12309, USA, Tel: (518) 387-6607; Fax: (518) 387-7751, E-mail:

Professor JOHN M. G. COWIE Heriot-Watt University, Department of Chemistry, Edinburgh, SCOTLAND EH14 4AS, Tel: +44 (0131) 451-3106; Fax: +44 (0131) 451-3180 E-mail: <[email protected]>

Professor LLOYD A. GOETTLER The University of Akron Department of Polymer Engineering Akron, OH 44325-0301, USA, Tel: +1 (330) 972-7467, Fax: +1 (330) 258-2339 E-mail:

Professor SUAT HONG, GOH National University of Singapore, Department of Chemistry, 3 Science Drive 3 Singapore 117543 Tel: +(65)-874-2844; Fax: +(65)-779-1691 E-mail:

Professor GABRIËL GROENINCKX Catholic University of Leuven, Department of Chemistry, Laboratory of Macromolecular Structural Chemistry, Celestijnenlaan 200F, 3001 Heverlee, Belgium, Tel: +32 (16) 327-440; Fax: +32 (16) 327-990 E-mail:

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List of Authors

Dr. MARIA RONG HU Atofina Chemicals, Inc. 900 First Ave. King of Prussia, PA 19406, USA Tel: +1 (610) 878-6703; Fax: +1 (610) 878-6260 E-mail: <maria.hu@atofina.com>

Professor TAKASHI INOUE Yamagata University, Department of Polymer Science and Engineering, 992-8510 Yonezawa, Japan, Tel: +81 (0238) 26-3060; Fax: +81 (0238) 26-3060 E-mail:

Professor MUSA R KAMAL McGill University, Department of Chemical Engineering, 3480 University Street, Montreal, QC, Canada H3A 2A7, Tel: +1 (514) 398 4262; Fax: +1 (514) 398 6678 E-mail:

Dr. I. J. McEWEN Department of Chemistry, Heriot-Watt University, Edinburgh EH14 4AS, Scotland

Dr. R. McINTYRE Millenium Inorganic Chemicals PO Box 26 Grimsby, NE Lincolnshire DN37 8DP, UK

Professor SONJA KRAUSE Rensselaer Polytechnic Institute, Department of Chemistry, Troy, NY 12180, USA, Tel: +1 (518) 276- 8445; Fax: +1 (518) 276-4887 E-mail:

Dr. V. N. S. PENDYALA Indian Petrochemicals Corporation Ltd., Research Centre, Distr. Vadodara 391 346, Gujarat, India, Tel: +91 (265) 272011/272031; Fax: +91 (265) 272098 /270040 E-mail:

Dr. LLOYD M. ROBESON Air Products and Chemicals, Inc., Corporate Science and Technology Center, 7201 Hamilton Blvd, Allentown, PA 18195, USA, Tel: +1 (610) 481-5026; Fax: +1 (610) 481-6517 E-mail:

Dr. JAMES J. SCOBBO, Jr. General Electric Plastics, Structured Products Bldg. 4, 1 Lexan Lane, Mt. Vernon IN, 47620, U.S.A Tel: +1 (812) 831-7103, dial comm: 8*393-7103 Fax: +1 (812) 831-4955, dial comm: 8*393-4955 E-mail: <[email protected]>

Dr. GERARD Z.-H. SHI Rehau Inc., P.O. Box 1706, 1501 Edwards Ferry Road, Leesburg, VA-20176, USA, Tel: +1 (703) 777-5255, ext. 2129; Fax: +1 (703) 777-3053 E-mail:

Dr. AJIT SINGH 822 Gulfview Place, Victoria, BC V8Y 2R6, Canada, Tel: +1 (250) 658-5575 E-mail: <[email protected]>

List of Authors

Professor LESLIE H. SPERLING Lehigh University, Materials Research Center, Whitaker Lab., 5 E. Packer Ave., Bethlehem, PA 18015-3194, USA, Tel: +1 (610) 758-3845; Fax: +1 (610) 758-3526 E-mail: Dr. LESZEK A. UTRACKI National Research Council Canada, Industrial Materials Institute, 75 de Mortagne, Boucherville, QC, Canada J4B 6Y4, Tel: +1 (450) 641-5182; Fax: +1 (450) 641-5105 E-mail: Dr. MYRIAM VANNESTE Centexbel Technologiepark 7 B-9052 Zwijnaarde, Belgium Tel:: +32 (9) 2438235; Fax: +32 (9) 2204955 E-mail: <[email protected]>

Dr. S. FRANCIS XAVIER Indian Petrochemicals Corporation Ltd., Distr. Vadodara, 391 346 Gujarat, India, Tel: +91 (265) 272011/272031 Ext. 3113; Fax: +91 (265) 272098 / 270040 E-mail:

vii

PREFACE

Science as a methodical investigation of Nature’s capacities evolved from the humble craft tradition. Its goal is to provide the most general and the simplest possible description of the observable character of Nature. In the past the singular concept of “science” comprised all aspects of intellectual endeavor: the arts, the sciences and the crafts. It was Diderot’s Encyclopédie ou Dictionnaire Raisonné des Sciences, des Arts et des Métiers of 1751-66 that first divided the old “science” into these three parts. The next split — that between the basic and applied sciences — is barely a century old. Basic science has been described as motivated by the desire to discover connections between natural phenomena, while applied science is the application of the discovered laws of nature for the material benefits of mankind. The boundary between the two is not rigid since experimental observation frequently provides a spur to fundamental discoveries. In the golden times of the scientific institutions in Europe and North America, the most prominent scientists, often the Nobel Prize winners directed the work. In the USA, during the years 1945-75, basic scientific research was considered “essential for the national security, economic growth and survival of the basic democratic values” [J. Krige & D. Pestre, Science in the 20th Century, 1997]. In the 1960s several major corporations supported research institutes with total freedom of the research topics. During last two decades of the twentieth century there has been an apparent reversal in the appreciation of science. Except for a few domains (e.g., astrophysics or atomic physics, project genome), intellectual efforts are being directed toward short-term developmental work of a commercially pertinent nature. This tendency is global, evident in the industrial, academic as well as state-supported laboratories. The CEOs hired for a contract to manage an institution are focused on the present. Managing has become a profession divorced from technical knowledge — a research institute, finance company, or pig farm may “benefit” from guidance by the same person. These tendencies are reflected in the evolution of polymer science and technology. The history of synthetic polymers is incredibly short. The term, polymer, was introduced in 1832. The first synthetic polymer (phenol-formaldehyde) was commercialized as Bakelite in 1909, while the first thermoplastic (polystyrene, Trolitul) six years later. The early polymer industry was developed by entrepreneurs that had little if any technical background. The commercial successes (and less known, but more numerous failures) predated even the fundamental idea of what constitutes the polymeric species. As late as 1926, Hermann Staudinger unsuccessfully advocated the concept of a linear, covalently bonded macromolecule. This idea was finally accepted during the Faraday Society meeting in 1935, only after Carothers reported on his polymerization studies, and demonstrated the validity of the polycondensation theory, developed by his younger colleague from du Pont de Nemours, Paul Flory. The theory provided the relationships between the molecular weights and the reaction kinetics, thus making it possible to ascertain validity of the newly formulated polycondensation principles that postulated sequential addition of bi-functional units to form linear macromolecules.

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Preface

In 1900 the world production of plastics was 25 kton, doubling during the following 30 years, then re-doubling in five. The most spectacular growth was recorded in the early 1940s when the demand created by the convulsions of World War II engendered a spectacular growth of 25% per annum. During the first 30 post-war years the global plastics’ industry sustained an average growth rate of 15%/year. By 1992 the world production of plastics had reached 102 million m3/year, while that of steel was 50 million m3/year. Furthermore, from 1980 to 1990 plastics production increased by 62% while that of steel decreased by 21%. Only during the past 20 years or so has plastics consumption shown smaller and more erratic advances. It is expected that by the year 2000 the world production of plastics will be 151 million tons/year. Considering the uneven polymer consumption around the world, polymer production has the potential to increase tenfold by the mid-21st century. Polymers are the fastest growing structural materials. Rubber blending predates that of thermoplastics by nearly a century. In 1846 Parkes introduced the first blends of trans- and cis-1,4-polyisoprene, i.e., natural rubber (NR) with gutta-percha (GP). By varying the composition and/or adding fillers the blends were formed into a variety of flexible or rigid articles. Polymer blends were developed alongside the emerging polymers. Once nitrocellulose (NC) was invented, it was mixed with NR. Blends of NC with NR were patented in 1865 — three years before the commercialization of NC. The first compatibilization of polyvinylchloride (PVC) by blending with polyvinylacetate (PVAc) and their copolymers date from 1928. PVC was commercialized in 1931 while its blends with nitrile rubber (NBR) were patented in 1936 — two years after the NBR patent was issued. The modern era of polymer blending began in 1960, after Alan Hay discovered the oxidative polymerization of 2,4-xylenols that led to polyphenyleneether (PPE). Its blends with styrenics, Noryl, were commercialized in 1965. At present, polymer alloys, blends, and composites consume over 80 wt% of all plastics.

In addition, the polymer blends segment of the plastics’ industry increases at about three times faster than the whole plastics’ industry. Blending has been recognized as the most versatile, economic method to produce materials able to satisfy complex demands for performance. By the year 2000 the world market for polymer blends is expected to reach 51 million tons per annum, worth well over US$ 200 billion. The tendency is to offer blends that can be treated as any other resin on the market; hence their processability must closely match that of single-phase polymer, but offer a much greater range of performance possibilities. In the economically advanced countries, plastics have displaced conventional materials for most applications. Today the market pressure forces the resin manufacturers to provide better, more economic materials with superior combinations of properties, not as a replacement for wood or steel, but rather to replace the more traditional polymers. This has resulted in: • Increased scale of production. • Use of multicomponent and multiphase materials. • New processing methods. For example, twin-screw extruders with 80 tons/hr throughput and injection (100,000 kN) molding presses with shot size of 100 liters of polymer are available. Composites where the matrix is a polymer blend that comprises six different polymers have been introduced. Gas and multiple injection processes, melt-core technology, solid-state forming, microcellular foams all lead to new products with advanced performance. The polymer industry is becoming increasingly sophisticated. To support these new tendencies the research community has been asked to provide better predictive methods for the multicomponent blends as well as improved sensors for the closed-loop process control. In particular, the evolution of morphology during the compounding and processing steps is of paramount importance. Microrheology and coalescence are the keys to describing the structure evolution of polymer blends.

Preface

In the early 1990s, the first mathematical models capable of predicting the evolution of morphology during compounding of polymer blends were developed. The fully predictive model provided good agreement with the experimentally determined variation of morphology inside a twin-screw extruder. However, it must be recognized that the morphology developed inside the compounding or the processing unit is dynamic. Upon removal of stress and in the absence of effective compatibilization the morphology changes with time. Today, very few unmodified resins are being used. Some polymers require less modification then others. For example, the semi-crystalline polymers that already have a two-phase structure may need modification less urgently. By contrast, the amorphous resins, such as PVC, PS, PPE, or polycarbonate of bis-phenol-A (PC), are brittle and require blending more frequently. The advantages of blending fall into two categories. A. Blending may improve resin or product performance by: 1. Producing materials having a full set of the desired properties at lowest cost. 2. Extending the engineering resins’ performance by incorporation of less expensive polymers. 3. Improvement of specific properties. • Toughening brittle polymers, thus eliminating the need to use low molecular weight additives (e.g., plasticizer in the flexible PVC formulations). • Blending with more rigid and more heat resistant resin may lead to improved modulus and dimensional stability. • Incorporation of semicrystalline polymer into an amorphous resin to improve solvent and chemical resistance (e.g., in blends of PC with PEST). • Incorporation of non-flammable resin into a flammable one improves flame resistance (e.g., styrenics or acrylics with PVC). • Blends with polymers having either –OH or –SH functionality leads to permanently anti-static blends (e.g., ethylene oxide-coepichlorohydrin with ABS/PC blend). • Biodegradable materials produced by incorporation of a biodegradable resin.

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• Blending makes it possible to produce integrated multi-layer structures. 4. Providing means for recycling of industrial and/or municipal plastics waste. 5. The blending technology makes it possible to rebuild high molecular weights of partially degraded polymers, thus to produce high performance articles from the plastics waste. B. Blending may lead to improved processability: 1. Incorporation of a miscible resin with a lower glass transition temperature (Tg) makes it possible to process the high Tg resin at temperatures well below the thermal degradation limit (e.g., PS/PPE blends). 2. Incorporation of an immiscible, low viscosity resin makes it possible to reduce pressure drop across dies or runners, thereby increasing productivity (e.g., LCP/PEEK blends). 3. Blending with resin that either by itself shows high strain hardening (SH) (e.g., LDPE in blends with another PO), or when reactively blended forms long chain branches (e.g., PS in blends with PO) results in blends having a controllable degree of SH. These materials show better processability in technologies where the extensional flow field is important, viz., film blowing, blow molding, wire coating, foaming, etc. 4. Incorporation of elastomeric particles improves nucleation of gas bubbles, hence it stabilizes the foaming process, reduces bubble size and the final foam density. 5. Incorporation of a degradable resin into an engineering or specialty one provides the means for generation of a controllable amount of the foaming gas during the ensuing stages of processing, viz., injection molding. 6. Blending different grades of the same resin broadens the molecular weight distribution, which in turn results in easier, more stable processing (as well as better mechanical performance). 7. Blending improves product uniformity (scrap reduction) and plant economy. 8. Blending ascertains quick formulation changes, thus plant flexibility and productivity.

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Preface

9. Blending reduces the number of grades that need to be manufactured and stored. 10. Blending technology offers methods for producing higher esthetic value materials, e.g., films or coatings without gel particles (or “fish eyes”), moldings with streak-free surface finish, etc. The aim of the Polymer Blends Handbook (PBH) is to provide the most comprehensive information on all aspects of polymer blend science and technology. The book will be useful for students entering the field as well as to seasoned professionals. The contributors to PBH are renowned experts from eight countries and four continents, who work in academe, government laboratories and industry. In consequence, the book may be considered comprising two parts: 1. Fundamental principles (nine chapters), and 2. Technology (eight chapters and four Appendices). Each chapter provides an introduction to the pertinent topic, discusses the principal aspects, the typical approaches used by the experts in the area, provides numerical values of pertinent parameters, and gives extensive references that facilitate further topical studies.

PBH comprises seventeen chapters that contain 1. Introduction to polymer blends, 2. Thermodynamics, 3. Crystallization, 4. Interphase and compatibilization by addition of a compatibilizer, 5. Reactive compatibilization, 6. Interpenetrating Polymer Networks, 7. Rheology, 8. Morphology, 9. Compounding, 10. Processing, 11. Use of radiation, 12. Properties and performance, 13. Applications, 14. Degradation and aging, 15. Commercial blends, 16 Role of polymer blends’ technology in polymer recycling, and 17. Perspectives. Furthermore, the Appendices provide information on 1. International abbreviations for polymers and polymer processing, 2. Miscible polymer blends, 3. Examples of commercial polymer blends, and 4. Dictionary of terms used in polymer science & technology. Finally, the Editor wishes to express thanks and personal appreciation to the contributors. They invested much time outside their regular duties, collecting the material and setting it into uniform text. They showed a high spirit of cooperation and great patience. The Polymer Blends Handbook is a testimonial of their efforts.

L. A. Utracki, Editor Montreal, 25.12.1999

CONTENTS

List of Authors . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . v Preface . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . ix

1. 1.1 1.2 1.3 1.4 1.5 1.6 1.7 1.8 1.9 1.10

2. 2.1 2.2 2.3 2.4 2.5 2.6 2.7 2.8 2.9 2.10

3. 3.1 3.2 3.3 3.4 3.5 3.6

Introduction to Polymer Blends . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1 L. A. Utracki Early polymer industry . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2 Polymer structure and nomenclature . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 7 Introduction to polymer blends . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 12 Commodity resins and their blends . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 24 Engineering resins and their blends . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 64 Specialty polymers and their blends . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 80 Biodegradable blends . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 93 Blending and recycling . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 93 Conclusions and outlook . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 95 References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 96

Thermodynamics of Polymer Blends . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . L. A. Utracki Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Thermodynamic principles . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Thermodynamics of a single component systems . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Polymeric liquid mixtures . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Theories of liquid mixtures . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Phase separation . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Methods of measurements . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Summary and conclusions . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . List of abbreviations . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

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Crystallization, Morphological Structure, and Melting of Polymer Blends . . . . . . . . . . . G. Groeninckx, M. Vanneste and V. Everaert Abstract . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . General introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Crystallization, morphological structure, and melting behavior of miscible polymer blends . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Crystallization, morphological structure and melting behavior of immiscible polymer blends . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Nomenclature . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

203

123 124 127 134 138 168 175 193 194 194

203 204 205 241 285 289

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4. 4.1 4.2 4.3 4.4 4.5 4.6 4.7 4.8

5 5.1 5.2 5.3 5.4 5.5 5.6 5.7 5.8 5.9 5.10 5.11 5.12 5.13 5.14 5.15

6. 6.1 6.2 6.3 6.4 6.5 6.6 6.7 6.8

7. 7.1 7.2

Contents

Interphase and Compatibilization by Addition of a Compatibilizer . . . . . . . . . . . . . . . . . A. Ajji Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Theoretical aspects of the interface . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Determination of the interfacial parameters . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Compatibilization by addition of a compatibilizer . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Patented blends with added compatibilizer . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Conclusions and outlook . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . References. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Nomenclature . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

Reactive Compatibilization of Polymer Blends . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . S. B. Brown Introduction and purpose . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Definitions of compatibilization and polymer alloys . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Types of polymer blends . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Characteristics of immiscible polymer blends . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . General strategies for compatibilization of immiscible polymer blends . . . . . . . . . . . . . . . . . Compatibilization by redistribution to produce block and random copolymer . . . . . . . . . . . . Compatibilization by graft copolymer formation . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Compatibilization by block copolymer formation . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Compatibilization by direct covalent crosslinking . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Compatibilization by covalent crosslinking mediated by an activating agent . . . . . . . . . . . . . Compatibilization by covalent crosslinking mediated by a coupling agent . . . . . . . . . . . . . . . Compatibilization by ionic interaction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Summary . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Abbreviations . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

Interpenetrating Polymer Networks . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . L. H. Sperling and R. Hu Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Synthetic methods . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Morphology and glass transitions . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Latex-based materials . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Actual and proposed applications of IPNs . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Summary . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Abbreviations and acronyms . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

295 295 296 308 316 326 329 331 336

339 339 341 342 343 345 347 352 381 388 394 398 401 404 405 414

417 417 419 421 427 438 442 442 443

The Rheology of Polymer Alloys and Blends . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 449 L. A. Utracki and M. R. Kamal Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 449 Rheological models for miscible blends . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 457

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7.3 7.4 7.5 7.6 7.7 7.8 7.9

Model systems for immiscible blends . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Rheology of miscible blends . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Rheology of immiscible blends . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Conclusions . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Notation and abbreviations . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Appendices . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

458 482 489 523 524 535 537

8.

Morphology of Polymer Blends . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Takashi Inoue Abstract . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Microscopic methods . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Quantification of morphology . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Crystallization from a single-phase mixture . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Spinodal decomposition under non-iso-quench depth . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Melt blending of highly immiscible polymers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . "Structure" in single phase mixtures . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Concluding remarks . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Acknowledgement . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Nomenclature . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

547

8.1 8.2 8.3 8.4 8.5 8.6 8.7 8.8 8.9 8.10 8.11 8.12

9. 9.1 9.2 9.3 9.4 9.5 9.6 9.7

10. 10.1 10.2 10.3 10.4 10.5 10.6 10.7 10.8 10.9

Compounding Polymer Blends . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . L. A. Utracki and G. Z.-H. Shi Fundamentals of polymer mixing . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Blending methods and equipment . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Non-mechanical methods of polymer blending . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Reactive processing (compatibilization) . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Summary . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Nomenclature and abbreviations . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

Polymer Blends Forming . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Michel M. Dumoulin Abstract . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Introduction to polymer forming . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Forming polymer blends . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Extrusion . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Injection molding . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Blow molding . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Other forming processes . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Forming of filled and/or reinforced blends . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Quality and process control . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

547 548 548 556 558 562 567 571 572 572 572 574

577 577 601 625 631 640 641 650

653 653 654 660 686 699 727 732 742 745

xvi

Contents

10.10 10.11 10.12 10.13

Reprocessing . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Conclusions . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Notation and abbreviations . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

748 750 752 754

11.

Use of high-energy radiation in polymer blends technology . . . . . . . . . . . . . . . . . . . . . . . . A. Singh and K. Bahari Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Radiation chemistry of organic compounds . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Radiation effects on polymers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Irradiation of polymer blends . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Concluding remarks . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Acknowledgement . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

757

11.1 11.2 11.3 11.4 11.5 11.6 11.7

12.

758 758 764 774 853 854 854

Properties and Performance of Polymer Blends . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . S. F. Xavier 12.1 Abstract . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 12.2 Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 12.3 Low speed mechanical properties . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 12.4 High-speed mechanical properties . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 12.5 Miscibility and solubility (by V. N. S. Pendyala) . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 12.6 Thermal properties . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 12.7 Flammability . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 12.8 Other properties . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 12.9 Conclusions . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 12.10 Acknowledgement . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 12.11 References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 12.12 Appendices . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

861 862 863 872 894 905 915 922 930 931 931 942

13.

951

Applications of Polymer Alloys and Blends . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . J. J. Scobbo, Jr. and L. A. Goettler 13.1 General Principles for the Use of Polymer Blends . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 13.2 Automotive applications . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 13.3 Lawn and garden . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 13.4 Electrical and Electronics Applications . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 13.5 Medical Applications . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 13.6 Building and Construction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 13.7 Business Machines and Communications . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 13.8 Packaging . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 13.9 Miscellaneous . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 13.10 References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 13.11 List of abbreviations . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

861

951 954 964 965 967 969 970 973 974 975 976

Contents

14 14.1 14.2 14.3 14.4 14.5 14.6 14.7

15 15.1 15.2 15.3 15.4 15.5 15.6 15.7 15.8 15.9 15.10 15.11 15.12 15.13 15.14 15.15 15.16

16.

xvii

Aging and Degradation of Polymer Blends . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 977 J. M. G. Cowie I. J. McEwen and R. McIntyre Abstract . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 977 Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 978 Physical aging . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 978 Chemical aging . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1000 Weathering of polymer blends . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1005 Summary, conclusions and outlook . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1014 References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1015

Commercial Polymer Blends . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . M. K. Akkapeddi Abstract . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Polyolefin blends . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Styrenic blends . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Vinyl resin blends . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Acrylic blends . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Elastomeric blends . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Polyamide blends . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Polycarbonate blends . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Polyoxymethylene blends . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Polyphenyleneether (PPE) blends . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Thermoplastic polyester blends . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Specialty polymer blends . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Thermoset blend systems . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Conclusions . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

Role of Polymer Blends’ Technology in Polymer Recycling . . . . . . . . . . . . . . . . . . . . . . . L. A. Utracki 16.1 Abstract . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 16.2 Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 16.3 The principles of polymer blending . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 16.4 Principles of polymer degradation and stabilization . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 16.5 Use of polymer blend technology in plastics recycling . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 16.6 Polymer blends' recycling . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 16.7 Post-consumer waste (PCW) recycling . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 16.8 Biodegradable blends . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 16.9 Conclusions . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 16.10 Nomenclature . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 16.11 References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

1023 1023 1024 1032 1042 1053 1056 1058 1063 1082 1092 1093 1097 1101 1106 1110 1110

1117 1117 1117 1126 1131 1137 1139 1146 1151 1157 1158 1160

xviii

17.

Contents

Perspectives in Polymer Blend Technology . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . L. M. Robeson 17.1 The past and present: relevance to the future . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 17.2 Perspectives in the future developments in commodity polymer blends . . . . . . . . . . . . . . . . 17.3 Engineering polymer blends: future perspective . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 17.4 High temperature polymer blends . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 17.5 Electrically conductive polymer blends . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 17.6 Biodegradable polymer blends . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 17.7 Future trends in the science and technology of polymer blends . . . . . . . . . . . . . . . . . . . . . . . 17.8 Additional areas of emerging and future blend technology . . . . . . . . . . . . . . . . . . . . . . . . . . 17.9 Postscript . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 17.10 References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

1167 1167 1171 1174 1178 1182 1186 1187 1190 1194 1194

APPENDIX I. International Abbreviations For Polymers And Polymer Processing . . . . . . . . 1201 compiled by L. A. Utracki APPENDIX II. Miscible Polymer Blends . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1215 prepared by S. Krause and S. H. Goh 1. Introductory remarks . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2. Table I Chemically dissimilar polymer pairs miscible in the amorphous state at room temperature. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1219 3. Table II Polymer pairs containing one monomer in common, miscible in the amorphous state at room temperature. . . . . . . . . . . . . . . . . . . . . . . . 1269 4. Table III Chemically dissimilar polymer triads (and tetrads) miscible in the amorphous state at room temperature. . . . . . . . . . . . . . . . . . . . . . . . 1276 5. Table IV Polymer pairs miscible in the amorphous state at room temperature. Molecular weight dependence investigated. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1279 6. Table V Polymer pairs that appear to have high temperature miscibility although immiscible at or below room temperature (UCST behavior). . . . . . . . . . . . . . . . . . . 1282 7. Table VI Polymer pairs miscible at room temperature that appear to have a lower critical solution temperature (LCST) above room temperature. . . . . . . . . . . . . . . . . . . . . . . 1284 8. Table VII Polymer pairs appear to have both a lower critical solution temperature and a upper critical solution temperature. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1294 9. Table VIII Polymer pairs that co-crystallize and form mixed crystals. . . . . . . . . . . . . . . . . . . . 1293 10. References 1297 APPENDIX III. Examples of Commercial Blends . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1323 compiled by L. A. Utracki APPENDIX IV. Dictionary of Terms Used in Polymer Science and Technology . . . . . . . . . . . . 1333 compiled by L. A. Utracki

Subject Index . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1387

CHAPTER 1

INTRODUCTION TO POLYMER BLENDS

L. A. Utracki

National Research Council Canada, Industrial Materials Institute, Boucherville, QC, Canada

An Introductory Note In this introductory chapter the basic information on polymer blends (with a special emphasis on the commercial alloys) is presented in the sequence: (i) a historical perspective on the polymer science and technology, (ii) polymeric structures and nomenclature, (iii) fundamental concepts in polymer blend science, and (iv) evolution of polymer blends’ technology. The world production of plastics in 1900 was about 30,000 tons — in the year 2000 it is expected to reach 151 Mt. The projected saturation level on the global scale (an increase by a factor of ten) is expected to be reached in the middle of the 21st century. The rapidity of the plastics expansion can be best judged by comparing it with steel — already in 1992 the annual world production of plastics more than doubled in volume the world production of steel, and nearly tripled its value. Polymers are the fastest growing structural materials. It is noteworthy that the polymer blend segment of the plastics industry increases at a rate about three times higher than the whole. Polymers are classified as either natural that resulted from natural biosynthesis, or synthetic. The natural (polysaccharides, proteins, nucleic acids, natural rubbers, cellulose, lignin, etc.) have been used for tens of thousands of years. In Egypt the musical string instruments, papyrus for writing, and styrene [in a tree balsam] for embalming were used 3,000 BC. For millennia shellac has been used in Indian turnery [Chattopadhyaya, 1986]. The natural rubber was used by Olmecs at least 3000 years ago [Stuart, 1993]. The term synthetic polymer refers equally well to linear, saturated macromolecules (i.e., thermoplastics), to unsaturated polymers (i.e., rubbers), or to any substance based on crosslinkable monomers, macromers, or pre-polymers (i.e., thermosets). The focus of this handbook is on blends of thermoplastics made of predominantly saturated, linear macromolecules. There are many sources of information about polymer history [Martuscelli et al., 1987; Seymour and Cheng, 1987; Vogl and Immergut, 1987; Alper and Nelson, 1989; Morris, 1989; Seymour, 1989; Sperling, 1992; Mark, 1993; Sparke, 1993; Utracki, 1994, 1998a]. The abbreviations used in this text are listed in Appendix 1. International Abbreviations for polymers and polymer processing.

L.A. Utracki (Ed.), Polymer Blends Handbook, 1-122. © 2003 Kluwer Academic Publishers. Printed in the Netherlands.

2

L. A. Utracki

1.1

Early Polymer Industry

1.1.1

The Beginnings

The polymer industry traces its beginning to the early modifications of shellac, natural rubber (NR — an amorphous cis-1,4-polyisoprene), gutta-percha (GP — a semi-crystalline trans-1,4-polyisoprene), and cellulose. In 1846, Parkes patented the first polymer blend: NR with GP partially co-dissolved in CS2. Blending these two isomers resulted in partially crosslinked (co-vulcanized) materials whose rigidity was controllable by composition. The blends had many applications ranging from picture frames, table-ware, ear-trumpets, to sheathing the first submarine cables. 1.1.2

Modified Natural Polymers

The first man-made polymer was nitrocellulose (NC). The main use of the NC resins was a replacement of the natural and expensive materials, viz., ivory, tortoise shell products, amber, ebony, onyx or alabaster. The use of cellulose acetate (CA) as a thermoplastic began in 1926. Cellulose ethers and esters became commercially available in 1927. Casein crosslinked by formaldehyde gave horn-like materials — Galalith™ has been used to manufacture shirt buttons, or as imitation of ivory and porcelain [Pontio, 1919]. 1.1.3

Synthetic Rubbers

The first polymerization of isoprene in sealed bottles was reported in 1884 by Tilden. Methyl rubber was thermally polymerized at 70°C — the reaction required 3 to 6 months, giving poor quality products. In 1926 BASF developed sodium-initiated polymerization of butadiene known as Buna™ (for BUtadiene + Natrium). The first successful, general purpose rubbers were copolymers of butadiene with either styrene, Buna-S, or acrylonitrile, Buna-N [Tschunkur and Bock, 1933; Konrad and Tschunkur, 1934]. Poly(2-chlorobutadiene), chloroprene [Carothers et al., 1931], was introduced in 1931 by DuPont. Elastomeric polysulfides [Patrick, 1932] were

commercialized in 1930 as Thiokol™. In 1937 butyl rubber (copolymer of isobutylene with isoprene) was invented. The synthetic rubber production took a big leap during the second world war (WW-2) [Morton, 1982]. 1.1.4

Synthetic Thermosetting Polymers

The first commercially successful synthetic polymer was phenol-formaldehyde (PF) [Smith, 1899]. The resin was introduced in 1909 by Baekeland as Bakelite™. The urea-formaldehyde resins (UF), were discovered in 1884, but production of Beetle™ moldable resin commenced in 1928. Three years later, Formica™, phenolic paper covered with decorative layer protected by UF, was introduced. The thiourea-formaldehyde molding powders were commercialized in 1920, while in 1935, Ciba introduced Cibanite™, aniline-formaldehyde (AF) molding materials, then two years later, the melamine-formaldehyde (MF). Epoxy compounds were discovered by Prileschaiev in 1909, but its importance was realized only during WW-2. In 1956, glass fiber reinforcements were introduced. The thermoset polyesters (TS) were developed by Ellis in 1933-4. The first use of glass-reinforced TS dates from 1938. 1.1.5

Synthetic Thermoplastic Polymers

The synthetic polymers are divided into three categories: 1. Commodity, 2. Engineering, and 3. Specialty. The five large-volume polymeric families that belong to the Commodity resins are: polyethylenes (PE), polypropylenes (PP), styrenics (PS), acrylics (PMMA), and vinyls (PVC). Their world market share (see Table 1.1) remains relatively stable — they represent 79% of all plastics. The five engineering polymer families are: polyamides (PA), thermoplastic polyesters (PEST), polycarbonates (PC), polyoxymethylenes (POM), and polyphenylene ethers (PPE). They constitute about 11% by volume and 34% by value of

3

Introduction to Polymer Blends

Table 1.1. No. 1. 2. 3. 4. 5. 6. 7.

World Market Share (MS) and Annual Growth Rate (AGR) by Resin Type (1995-2000) Resin Type

MS (%)

Low density polyethylene (LDPE + LLDPE) High density polyethylene (HDPE) Polypropylene (PP) Polystyrene and copolymers (PS + ABS) Polyvinylchloride (PVC) Other thermoplastics (TP) Thermosets (TS)

20 13 17 11 18 11 10

the plastic’s consumption. The engineering and specialty polymers show high mechanical performance, and the continuous use temperature 150 ” CUT(°C) ” 500. The polymer industry increasingly favors high technology and high value-added materials. These are obtained either by means of new polymerization methods, new processing technologies, or by alloying and reinforcing. For example, new syndiotactic PP or PS (sPP or sPS, respectively) surpasses the performance of their predecessors. The gel spun PE fibers have 200 times higher tensile strength than standard PE. New aromatic polyester (EKF from Sumitomo) has tensile strength of 4.1 GPa, to be compared with 70 MPa of a standard polyester resin (see Table 1.2).

Table 1.2. No.

5.1 5.2 6.6 4.7 5.2 4.5 -2.1

New types of polymers are also being introduced, e.g., dendritic-structure polymers [Fréchet et al., 1992], carbosilane dendritic macromolecules [Roovers et al., 1993], the “hairy rod” molecular structures where rigid-rod chain macromolecules are provided with short and flexible side branches [Wegner, 1992], etc. However, the polymer technology invariably moves away from the single phase materials to diverse combinations of polymers, additives, and reinforcements. While synergistic effects are often cited, the main reason is a need for widening the range of properties, for development of materials that would have the desired combination of properties — tailor-made polymeric systems. At present, about 36% of the synthetic resins are used in blends and about 39% in composites.

High performance materials — a comparison Material

Strength (GPa) Theoretical

1. 2. 3. 4. 5. 6. 7. 8. 9. 10.

AGR (%)

Polyethylene (standard) Polyethylene gel-spun Polyester (standard) Polyester oriented Aromatic Polyamide Aromatic Polyester (EFK) Poly(phenylene benzothiazole) Polyazomethin Carbon fiber Steel

21 21 24 24 21 29

Modulus (GPa)

Observed

Theoretical

” 0.03 6.0 0.07 1.2 3.6 4.1 4.2 4.7 3.1 2.1 to 3.5

316 316 124 124 190 371 -

Observed 0.2 220 2.2 21 125 139 365 125 235 210

4

1.1.6

L. A. Utracki

Compounding and Processing

The first mixer was an annular container with a spiked rotor for rubber compounding [Hancock, 1823]. The calender/two roll mill was patented by Chaffe in 1836 and manufactured by Farrel Co. A counter rotating twin shaft internal mixer with elliptical rotating discs or sigma-blades was developed by the end of the last century [Freyburger, 1876; Pfleiderer, 1880]. The first hand-operated extruder was a ram press, used for forming NR or GP, then later NC. The first, belt-driven extruders with Archimedean screw were patented nearly a century later [Gray, 1879]. In 1939, Paul Leistritz Maschinenfabrik built electrically heated, air cooled extruder, with nitrited barrel, having L/D = 10, an automatic temperature control, variable screw speed. The machine is considered a precursor of the modern single-screw extruders, SSE. During the WW-2 breaker plates, screen packs, crosshead dies, coextrusion, monofilament extrusion, film blowing and biaxial sheet orientation were introduced. In the 1950’s, a coextrusion process, venting, and two-stage screws were developed. In 1980’s the microprocessor control evolved into computer integrated manufacturing, the helical grooved feed barrels, high pressure gear pumps, air lubricated die flow, biaxial film orientation were introduced [Utracki, 1991a, c]. Pfleiderer patented the first modular counter-rotating twin-screw extruder (TSE) in 1882. An intermeshing, corotating TSE, the predecessor of the modern machines, was designed for extrusion of CA. The TSE was used by I. G. Farbenindustrie for the production of PA-6 [Colombo, 1939]. In 1959, Werner & Pfleiderer introduced ZSK machines (vented, intermeshing, corotating, with segmented screw and barrel, twin-screw extruders). These provided good balance between the dispersive and distributive mixing at relatively high output rates. In 1979 Japan Steel Works (JSW) developed TEX-series TSE’s for reactive compounding, permitting an easy change of the screw direction from co- to counter-rotation. In collaboration with Sumitomo Chem., barrel elements with sampling ports were designed, providing ready access to the processed material for deter-

mining the reaction progress and morphology [Nishio et al., 1990]. American Leistritz has been active in designing TSE kneading elements that improved mixing capability by maximizing the extensional flow field. More information on the evolution of the extrusion technology can be found in Chapter 9. Compounding Polymer Blends. Injection molding of NC dates from 1872. The early machines were hand operated. They used an axially movable screw or plunger and were equipped for devolatilization. The commercialscale injection molding of PS has begun in 1931. In 1932 Gastrow developed the first automatic unit, Isoma-Automat (30 g capacity per shot), with torpedo-type heating chamber. In 1951, Willert invented an in-line reciprocating screw plasticization that revolutionized the injection molding industry. The first automated injection molding plant was developed by Eastman Kodak in 1950. Hayatt used blow molding in 1880 to produce baby rattles out of CA tubes or sheets. In 1942, Plax Corp. started manufacturing squeezable LDPE bottles. By the end of 1950’s blow molding was the most rapidly developing processing method. In 1965 Wyeth, using the stretch blow molding, produced PET bottles. In 1972, Toyo Seikan started to produce multilayered blow molded bottles from PP and EVAl. In 1976 Ishikawajima-Harima introduced intermittent coextrusion blow-molding system for large parts. 1.1.7

Development of Polymer Science

1.1.7.1 Polymerization Cellulose modification dates from 1833 (Braconnot). In 1838 Regnault photo-polymerized vinylidene chloride. A year later Simon observed that heating styrene in the presence of air generated a tough gelatinous material — a low molecular weight PS. Polyoxymethylene (acetal) was discovered in 1859. In 1872 several new polymers were announced, viz., PVC, polyvinyl bromide (PVB) and phenolformaldehyde (PF). Polymethacrylates were discovered by Kahlbaum in 1880, polymethylene in 1897, one year later polycarbonate (PC) by Einhorn, polyamide-6 (PA-6) in 1907, etc. In the

Introduction to Polymer Blends

1920’s the list of polymers rapidly started to increase, viz. polysulfide (PSF), polyvinyl alcohol (PVAl), poly(styrene-co-maleic anhydride) (SMA), polyvinylformal (PVFO). During the next decade polyacetylene (PACE), styrene-acrylonitrile copolymer (SAN), low density polyethylene (LDPE), polyvinylidenechloride (PVDC), epoxy resins (EP), polyamides (e.g., PA-66; PA-610; PA-106), polysiloxanes (PDMS), polychlorotrifluoroethylene (PCTF), polytetrafluoroethylene (PTFE), and many others, were discovered [Utracki, 1989a]. Most early thermoplastics, e.g., PVC or PS, were obtained in the free radical polymerization, initiated either by heat or by sunlight. The first systematic studies of the free radical chemistry commenced 80 years later [Ostromislensky, 1911, 1915, 1916]. Fikentscher empirically determined which one of the 30-or-so monomers liked or disliked to copolymerize with each other. The advantage of latex-blending was also established. The theory of the free radical copolymerization was finally developed in the 1940’s [Alfrey et al., 1952]. The polycondensation reactions have been known since the mid-1800 [Lourenço, 1859; Wurtz, 1859, 1860]. In 1927, Carothers and his colleagues provided the basis for understanding the nature of these reactions. Good agreement between Flory’s theoretical predictions and the experimental observations of the average molecular weight (MW) provided convincing arguments for the acceptance of the linear macromolecule model. The alkyl-lithium initiated, living anionic polymerization of elastomers was described in 1928 by Ziegler. To polymerize styrene-isoprene block copolymers Szwarc et al., [1956] used sodium naphthalene as an anion-radical di-initiator, while Shell used an organolithium initiator. The polymerization mechanism was described by Bywater [1965]. In the early 1950’s, Ziegler found that in the presence of ZrCl4 + AlR3 ethylene can be polymerized at low temperature and pressure into linear, high density polyethylene (HDPE). The catalysts developed by Ziegler, and later by Natta become known as Ziegler-Natta, or Z-N catalysts. These can be defined as polymerization initiators created from a catalyst (1) and co-catalyst

5

(2), where (1) is halide or oxyhalide of transition metals from groups IV to VII, and (2) is an organometallic compound of metal from groups I to III. The Z-N-catalyst are prepared by mixing ingredients (1) and (2) in a dry, oxygen-free solvent [Natta and Danusso, 1967]. Recently Z-N catalysis is in renaissance, with new MgCl2supported catalysts that have a hundred-fold more active sites per mole of Ti and about ten times higher propagation rate [Rieger et al., 1990]. The new, single-site metallocene catalysts make it possible to control MW, molecular weight distribution (MWD), comonomer placement, stereoregularity, and life-time of the reactive chain-end [Kaminsky et al., 1985, 1992; Kaminsky, 1998; Swogger, 1998]. The use of either (Cp)pR’s(Cp)MeQ3-p, or R”s(Cp)2MeQ’ {where: Cp is cyclopentadienyl (substituted or not) radical, Me is metal from Group 4b, 5b, or 6b, R’, R”, Q and Q’ are radicals (viz. aryl, alkyl, alkenyl, alkylaryl, or arylalkyl), s = 0-1, p = 0-2}, for the polymerization of ethylene copolymers, provides independent control of MW and density. The catalyst is used in combination with a large amount of alumoxanes. In 1975 Mitsui Petrochemicals introduced metallocene-made LLDPE Tafmer™, with controlled comonomer placement, but rather low MW. In 1991, Dow Plastics produced developmental quantities of ethylene copolymers with up to 25 mole % of butene, hexene or octene, Affinity™ resins. The use of a metallocene catalyst with a single cyclopentadiene ring, resulted in a certain degree of randomization of the polymerization process. The catalyst produced PP with narrow molecular weight distribution, and a long chain branching, similar to LDPE. The metallocene catalysts are also used to produce high melting point polymers out of commodity monomers, e.g., sPS, with Tg = 100°C, and Tm = 266°C, or syndiotactic poly(p-phenyl styrene), (sPhPS), with Tg = 196°C, Tm = 352°C, and the decomposition temperature, Tdecomp = 380°C. Since sPhPS is miscible with sPS in the whole range of concentration, blends of these two syndiotactic polymers can be processed at any temperature above 266°C [Watanabe et al., 1992].

6

L. A. Utracki

Polycyclohexylethylene (PCHE) is a new metallocene resin, developed by Dow as a replacement for PC in the production of optical discs. PCHE has low shrinkage (0.02% after 24 hrs), higher light transmission than PC (91.9% vs. 89.8%, respectively) and high flex modulus of 71 GPa. Commercial production is to start by the year 2000. In 1999 Equistar Chem. Introduced high performance, non-metallocene single-site catalyst for PE’s. 1.1.7.2 Polymer Physics Molecular Weight (MW) Osmotic pressure measurements for the determination of MW were used in 1900 to characterize starch. Twenty years later, the solution viscosity measurements were introduced by Staudinger for this purpose. However, it was Mark and his collaborators who developed the concept of the intrinsic viscosity ([η]) and demonstrated that it provides information on the volume of individual colloidal particles, thus on MW. For the freely rotating chains the dependence (today known as Mark-Houwink-Sakurada equation) was obtained [Guth and Mark, 1934]: [η] ≡ lim [(η/ηΟ) - 1] / c = KM va

(1.1)

c→0

where η and ηο are viscosities of the solution and solvent, Mv is the viscosity-average molecular weight, and K and a = 0.5-0.7, are equation parameters. In 1933 the ultra-centrifugation was introduced [Kraemer and Lansing, 1933]. Utility of light scattering for the determination of MW was demonstrated eleven years later [Debye, 1944, 1946]. Free Volume Concept The free volume theory of liquids dates from the beginning of the 20th century. Two expressions for the free volume fraction, f, have been proposed, either f = (V - Vocc)/V or less frequently used fD = (V - Vocc)/Vocc, (Vocc is the occupied volume). The theory was used to interpret the temperature (T) and pressure (P) dependencies of liquid viscosity [Batschinski, 1913]. The Vocc was defined as the specific volume at which the liquid viscosity

is immeasurably high, η → ∞. Good correlation was found between Vocc and either the critical volume or the van der Waals constant b, viz. V occ / (V crit / 3) = Vocc / b = 0.921 ± 0.018. Batschinski wrote: η = a0 + a1 / f = a0 + a1 V / (V - Vocc)

(1.2)

where ai are equation parameters. Forty years later, more accurate data of viscosity (spanning several orders of magnitude) and specific volume for a series of paraffin’s with molecular weight MW = 72 to 1000 g/mol led to the logarithmic dependence [Doolittle, 1951]: ln η = a0 + a1 Vo / (V - Vo)

(1.3)

where Vo is the value of V at a characteristic solidification temperature, To, at which the fluid viscosity increases to infinity. Eq 1.3 provided a basis for the derivation of well-known WLF time-temperature shift factor aT [Williams et al., 1955]. The free volume model has been also incorporated into thermodynamic theories of liquids and solutions [Prigogine et al., 1957] and it is an integral part of theories used for the interpretation of thermodynamic properties of polymer blends [Utracki, 1989a]. In particular, it is a part of the most successful equation of state (EoS) derived for liquids and glasses [Simha and Somcynsky, 1969], critically examined using data for 56 principal polymers [Rodgers, 1993]. Since the mid-1960’s, the lifetime of ortho-positronium has been used to measure the free volume fraction f. Accordingly, f increases linearly with the temperature: f = -0.13556 + 6.2878(T/T*) for 0.0165 ” T/T* ” 0.0703, where T* is the temperature reducing parameter in Simha-Somcynsky theory [Utracki, 1998b]. More detailed analysis indicated that the free volume should be discussed in terms of distribution of the holes. For example, the measurements showed that above Tg the number of holes does not increase, but their volume does [Kobayashi et al., 1989]. In PS/PPE blends, the size of the free volume spaces in PS was found smaller than that in PPE [Li et al., 1999].

Introduction to Polymer Blends

Viscoelasticity In 1874, Boltzmann formulated the theory of viscoelasticity, giving the foundation to the modern rheology. The concept of the relaxation spectrum was introduced by Thompson in 1888. The springand-dashpot analogy of the viscoelastic behavior (Maxwell and Voigt models) appeared in 1906. The statistical approach to polymer problems was introduced by Kuhn [1930]. Busse [1932] observed that “green” rubber under stress shows a dual behavior, suggesting presence of two types of interactions: few widely separated strong ones, acting as physical crosslinks, and many weak ones of the van der Waals type, that make it possible for one macromolecule to slip by the others. This postulate was the first connotation of the chain entanglement. Bueche [1952, 1956, 1962] adopted the entanglements’ concept for the interpretation of polymer flow. He calculated the molecular friction constant per statistical segment as the unit force needed to pull the undeformed macromolecule through the surrounding medium at unit speed, fo = F/N (with N being the number of statistical segments per macromolecule), deriving the relations (see Eq 1.4) between the diffusion constant, D, or zero-shear viscosity, ηο, and such molecular parameters as density, ρ, molecular weight, M, and radius of gyration, Rg:

(

)

(1.4)

Dηo = ( ρN A / 36) Rg2 / M kBT ; and

(

)

ηo = ( ρN A / 36) Rg2 / M N * fo for: M ≤ 2 Me

N * = M / Mo

for: M > 2 Me N * = β ( M / Mo )( ρN A / 48)( M / Me )

2

(

M Rg2 / M

)

3/ 2

where numerical constant β ≅ 0.6. The dependence predicts that for low molecular weight liquids (M below the value of the critical molecular weight for entanglement, Mc = 2Me, where Me is the molecular weight between entanglements) ηo should be proportional to M, while for high molecular weight macromolecules (above Mc) to M3.5. Thus, predictions of the entanglement-based theory

7

are in good agreement with the observations: ηo ∝ M3.4. The correlation between the plateau modulus and entanglement concentration soon followed [Ferry et al., 1955]. The long disputes on the nature of entanglement led to defining it as “a special type of interactions, affecting mainly the largescale motions of the chains, and through them, the long time end of the viscoelastic relaxation time spectrum” [Graessley, 1974].

1.2

Polymer Structure and Nomenclature

1.2.1

Basic Considerations

Polymer is a substance composed of macromolecules, built by covalently joining at least 50 molecular mers, or the Constitutional Repeating Units or CRU. The longest sequence of CRU defines the main chain of a macromolecule. The main chain may be composed of a series of subchains, identified by some chemical of physical characteristic (e.g., tactic placement). The main chain may also contain long or short side chains or branches, attached to it at the branch points. A small region in a macromolecule from which at least four chains emanate constitutes a crosslinking point. A macromolecule that has only one crosslink is the star macromolecule. A macromolecule consisting of several crosslinked chains, but having a finite molecular weight is a micronetwork. A highly ramified macromolecule in which each CRU is connected to every other CRU is a polymer network. When the main chain of a macromolecule has numerous branch points from which linear side chains emanate, it is comb macromolecule. The CRU is defined as a bivalent organic group, not necessarily identical to the source from which the macromolecule was prepared — it is the largest identifiable group in the polymer dictated by the macromolecular structure. To discuss the structure of polymer molecules, one may consider the chemical nature of CRU, type of the linkages, the global macromolecular arrangement, and the topochemical character of the macromolecule, tacticity, etc. These are summarized in Table 1.3.

8

L. A. Utracki

1.2.2

Polymer Nomenclature

4. The source of the compounds (viz., synthetic, natural, and derived products).

Macromolecular compounds can be classified according to: 1. The chemical structure of the repeating unit (viz. polyamides, polyesters, polyolefins). 2. The structure (viz. linear, branched, ladder, or crosslinked). 3. The phenomenological behavior or technological use. Table 1.3.

Macromolecular structures

No. Characteristic 1. 1.1. 1.2. 1.3. 1.4. 2. 3. 4. 4.1. 4.2. 4.3. 4.4. 4.5.

Recurring Constitutional Repeating Units, CRU Structure Joining similar CRU Joining different CRU Joining polymer segments The nature of bond between CRU Macromolecular structure Topochemical characteristics of macromolecule Geometrical isomers Optical isomers Tacticity Helical structures Head-to-tail, head-to-head

Table 1.4.

The Commission on Macromolecular Nomenclature defined 52 terms related to polymer structure, including polymer, constitutional units, monomer, polymerization, regular polymer, tactic polymer, block polymer, graft polymer, monomeric unit, degree of polymerization, addition polymerization, condensation polymerization, homopolymer,

Examples

Aliphatic, aromatic, heterocyclic, metallo-organic, ... Homopolymers (linear, branched, dendritic, crosslinked, etc.) Copolymers, multipolymers, polyadducts, polycondensates, ... Block copolymers, graft copolymers, ladder polymers, ... e.g., ether, ester, amide, urethane, sulfite, ... Linear, branched, cross-linked, dendritic, ... e.g., rubber and gutta-percha are poly(1,4-isoprene), cis- and trans-, respectively Having optically active C*; e.g., polypeptides, polysaccharides, ... Isotactic, syndiotactic, and atactic Polypeptides, tactic polymers Example: PIB or PS

Polymer nomenclature proposed by the IUPAC

No. Title

Reference

1. 2. 3. 4. 5. 6.

Huggins et al., 1966 IUPAC, 1974; 1996 IUPAC, 1976 IUPAC, 1981 IUPAC, 1984 IUPAC, 1985a

7. 8. 9. 10. 11. 12. 13.

Report on Nomenclature Dealing with Steric Regularity in High Polymers Basic Definitions of Terms Relating to Polymers Nomenclature of Regular Single-Strand Organic Polymers Stereochemical Definitions and Notations Relating to Polymers Note on the Terminology for Molar Masses in Polymer Science Nomenclature for Regular Single-Strand and Quasi-Single-Strand Inorganic and Coordination Polymers Source-Based Nomenclature for Copolymers Use of Abbreviations for Names of Polymeric Substances Definitions of Terms Relating to Individual Macromolecules, their Assemblies, and Dilute Polymer Solutions Definitions of, Terms Relating to Crystalline Polymers A Classification of Linear Single-Strand Polymers Compendium of Macromolecular Nomenclature Source-Based Nomenclature for Non-Linear Macromolecules and Macromolecular Assemblies

IUPAC, 1985b IUPAC, 1987 IUPAC, 1989a IUPAC, 1989b IUPAC, 1989c Metanomski, 1991 Jenkins et al., 1993

Introduction to Polymer Blends

copolymer, bipolymer, terpolymer, copolymerization [IUPAC, 1974]. The Commission remains the leading nomenclature body in the polymer field. Table 1.4 lists the pertinent sources for information on the nomenclature of polymeric materials. Since there are difficulties in assigning systematic and unique abbreviations to polymers, only a short list has the IUPAC’s official sanction. An extensive list of internationally used abbreviations is provided in Appendix I. The IUPAC Macromolecular Nomenclature Commission has published three sets of rules for naming polymers: 1. Traditional, trivial names are sanctioned by the historical use and approved by IUPAC as an alternative (examples are listed in Table 1.5), 2. Structure-based nomenclature, and 3. Source based nomenclature proposed by the Commission.

Table 1.5.

9

1.2.2.1 Structure-based Nomenclature For organic, regular, single-strand polymers the structure-based system of naming polymers should be used. This nomenclature describes chemical structures rather than substances. Three steps are to be followed in a sequence: 1. Identify the constitutional repeating unit, CRU. 2. Orient the CRU. 3. Name the CRU. The name of the polymer is poly(CRU). The preferred CRU is one beginning with the subunit of highest seniority. The order of seniority is: heterocyclic rings, chains containing heteroatoms, (in the descending order O, S, Se, Te, N, P), carbocyclic rings, chains containing only carbon. The seniority is expressed by brackets and internal parentheses (see examples in Table 1.5).

Traditional and systematic names of polymers

No.

Traditional name

Systematic name

1. 2. 3. 4. 5. 6. 7. 8. 9. 10. 11. 12. 13. 14. 15. 16. 17. 18. 19. 20. 21.

polyethylene polypropylene polyisobutylene polybutadiene polyisoprene polystyrene polyacrylonitrile polyvinyl alcohol polyvinylacetate polyvinylchloride polyvinylidenefluoride polytetrafluoroethylene polyvinylbutyral polymethylacrylate polymethylmethacrylate polyformaldehyde polyethylene oxide polyphenylene ether Polyethyleneterephthalate poly-ε-caprolactam polyamide-6,6 or polyhexamethyleneadipamide

poly(methylene) poly(propylene) poly(1,1-dimethyl ethylene) poly(1-butenylene) poly(1-methyl- 1-butenylene) poly(1-phenyl ethylene) poly(1-cyano ethylene) poly(1-hydroxy ethylene) poly(1-acetoxy ethylene) poly(1-chloro ethylene) poly(1,1-difluoro ethylene) poly(difluoro methylene) poly[(2-propyl-1,3-dioxane-4,6-diyl) methylene] poly[1-(methoxycarbonyl) ethylene] poly[1-(methoxycarbonyl)-1-methyl ethylene] poly(oxy methylene) poly(oxy ethylene) poly(oxy-1,4-phenylene) poly(oxyethylene-oxyterephthaloyl] poly[imino(1-oxohexamethylene)] poly[imino(1,6-dioxohexa methylene) iminohexa methylene]; or poly(iminoadipoyliminohexa methylene)

10

L. A. Utracki

After the CRU and its orientation, reading left to right, have been established, the CRU or its constituent subunits are named. The name (the largest identifiable unit) includes description of the main chain and the substituents. The subunits are named according to the rules for nomenclature of organic chemistry. The name of the CRU is formed by citing, in order, the names of the largest subunits within the CRU. More complicated, regular single-strand polymers can be represented as multiples of repeating units, such as [ABC]n. The name of the polymer is poly(ABC), where (ABC) stands for the names of A, B, and C, taken in the order of seniority. An extension of the structure-based method to linear inorganic and/or coordination polymers is limited by the general lack of a system for naming bivalent radical. Few polymers with inorganic, covalently-bonded backbones have trivial names [viz., poly(dimethylsiloxane) or poly(dichlorophosphazene)], some can be named (as organic polymers) by using bivalent radicals; e.g., poly[oxy(dimethyl silylene)] or poly[nitrilo (dichlorophosphoranylidyne)]. Structure-based nomenclature is also applicable to copolymers having a regular structure, regardless of the starting materials used [viz. poly(oxyethylene-oxyterephthaloyl)]. In principle, it should be possible to extend the existing structure-based nomenclature beyond regular, single-strand polymers to polymers that have reacted, cross-linked polymers, ladder polymers, and other more complicated systems. 1.2.2.2 Source-based Nomenclature Traditionally, polymers have been named by attaching the prefix poly to the name of the CRU, real or assumed monomer, the source from which it is derived. Thus PS is the polymer made from styrene. When the name of the monomer consists of two or more words, parentheses should be used, but for common polymers such as polyvinylchloride, polyvinylacetate, etc., it is customary to omit them. Different types of polymerization can take place with many monomers, and there are different ways for obtaining a polymer.

For example, name such as polyvinyl alcohol refers to a hypothetical source, since this polymer is obtained by hydrolysis of polyvinylacetate. In spite of deficiencies, the source-based nomenclature is still entrenched in the literature. It is also the basis for naming and classifying copolymers (see Table 1.6). 1.2.3

Copolymers

When mers are not identical the polymerization leads to a copolymer. For divalent mers a linear copolymer is obtained, but when at least some mers are able to join more than two units, the polymerization leads to branched or crosslinked copolymer. When the polymerization starts on a polymer chain of different chemical character that the one that is subsequently forming, the resulting structure is known as grafted copolymer. Thus, the arrangement of the different types of monomeric units must be specified. Several types of arrangements are shown in Table 1.6, where A, B, and C represent different CRU. The systematic source-based nomenclature for copolymers involves identification of the constituent monomers, and description of their arrangement. This is achieved by citing the names of the constituent monomers after the prefix “poly”, and by placing between the names of each pair of monomers an italicized connective to denote the kind of arrangement by which those two types of monomeric units are related in the structure. The structures listed in Table 1.6 are divided into three categories: Short sequences, Long sequences, and Networks. Within the first category a sequence of placement of individual CRU is considered, within the second the placement of long sequences of CRU defines the copolymer type, while to the third belong crosslinked networks, crosslinked polymers, and chemical-type interpenetrating polymer networks. The network is a crosslinked system in which macromolecules of polymer A are crosslinked by macromolecules of polymer B [Sperling, 1992]. The composition can be expressed as, e.g., block-co-poly(butadiene/styrene) (75:25 wt%), or graft-co-poly[isoprene/ (isoprene; acrylonitrile)] (85:15 mole %).

11

Introduction to Polymer Blends

Table 1.6. No.

Nomenclature of copolymers

Type

Connective

Example

1. 1.1. 1.2. 1.3. 1.4. 1.5.

SHORT SEQUENCES unknown or unspecified statistical random alternating periodic with at least three monomeric units

-co-stat-ran-alt-per-

poly(A-co-B) poly(A-stat-B) poly(A-ran-B) poly(A-alt-B); (AB)n poly(A-per-B-per-C); (ABC)n; (ABAC)n; (AABB)n, ...

2. 2.1. 2.2.

LONG SEQUENCES block graft (polymeric side chain different)

-block-, or -b-graft-or -g-

2.3. 2.4.

star star-block

-star-star-...-block-

poly A-block-poly B; poly(A-b-B) poly A-graft-poly B; poly(A-g-B); AAAAA(g-BBB)AAAAAAA star-poly A star-poly A-block-poly B

3. 3.1. 3.2. 3.3.

NETWORKS crosslinked interpenetrating conterminous

-cross-inter-cross-

cross-poly A cross-poly A-inter-cross-poly B poly A-cross-poly B

1.2.4

Macromolecular Assemblies

To describe polymers or polymer blends with greater precision the qualifiers listed in Table 1.7 have been suggested [Jenkins et al., 1993].

Table 1.7.

comb star

Polymer Blend Terminology

The terminology used in polymer blends’ science and technology is summarized in Table 1.8 [Utracki, 1989, 1998a].

Descriptors for non-linear macromolecules and macromolecular assemblies

Polymer Structure cyclic branched

1.2.5

cyclo (unspecified) short-chain long-chain with f (give numerical value) branch points comb (unspecified) with f (give numerical value) arms

network micronetwork polymer blend interpenetrating polymer network semi-interpenetrating polymer network macromolecule-macromolecule complex

Descriptor

branch sh-branch l-branch f-branch star f-star net m-net blend ipn sipn compl

12

Table 1.8.

L. A. Utracki

Terminology of polymer blends

Term

Definition

Polymer

A substance composed of large molecules, the macromolecules, built by covalently joining at least 50 molecular mers, segments, or recurring Constitutional Repeating Units, CRU. Commercial polymers may contain up to 2 wt% of another polymeric modifier. Polymeric material synthesized from more than a single monomer. Processable polymeric material, capable of being formed to precise and stable dimensions, exhibiting high performance at the continuous use temperature CTU > 100°C, and having tensile strength in excess of 40 MPa. Mixture of at least two macromolecular substances, polymers or copolymers, in which the ingredient content is above 2 wt%. Mixture of two homologous polymers (usually a mixture of narrow molecular weight distribution fractions of the same polymer). Polymer blend, homogenous down to the molecular level, associated with the negative value of the free energy of mixing: ∆Gm § ∆Hm ” 0, and a positive value of the second derivative: ∂2¨Gm/∂φ2 > 0. Operationally, it is a blend whose domain size is comparable to the dimension of the macromolecular statistical segment. Polymer blend whose free energy of mixing; ∆Gm § ∆Hm > 0. Term to be avoided! At best a utilitarian, non-specific term indicating a marketable, visibly homogeneous polymer mixture, with enhanced performance over the constituent polymers. Immiscible, compatibilized polymer blend with modified interface and morphology. Process of modification of the interfacial properties in immiscible polymer blend, resulting in reduction of the interfacial tension coefficient and stabilization of the desired morphology, thus leading to the creation of a polymer alloy. Third phase in binary polymer alloys, enhanced by inter-diffusion or compatibilization. Thickness of this layer varies with the blend components and compatibilization method from 2 to 60 nm. Polymer or copolymer that either added to a polymer blend or generated there during reactive processing modifies its interfacial character and stabilizes the morphology. Compatibilization by incorporation of a compatibilizer, usually either a copolymer or multipolymer. Compatibilization by physical means: high stress field, thermal treatment, irradiation, etc. Compatibilization during reactive processing, extrusion or injection molding. Polymer blend or polymer alloy that either contains or has properties of an engineering polymer. Polymer alloy, containing two or more polymers in the network form, each chemically crosslinked. Sequential, simultaneous (SIN) and latex type IPNs are known. Polymer alloy, containing two or more polymers in a co-continuous network form, each physically crosslinked. The crosslinking originates in crystallinity, ion cluster formation, presence of hard blocks in copolymers, etc.

Copolymer Engineering polymer (EP)

Polymer blend (PB) Homologous polymer blend Miscible polymer blend

Immiscible polymer blend, Compatible polymer blend

Polymer alloy Compatibilization

Interphase

Compatibilizer Chemical compatibilization Physical compatibilization Reactive compatibilization Engineering polymer blend Interpenetrating polymer network (IPN) Thermoplastic IPN

1.3

Introduction to Polymer Blends

Polymer blends constitute ca. 36 wt% of the total polymer consumption, and their pertinence

continues to increase. About 65% of polymer alloys and blends are produced by polymer manufacturers, 25% by compounding companies and the remaining 10% by the transformers.

Introduction to Polymer Blends

1.3.1

Benefits and Problems of Blending

The following material-related benefits can be cited: (i) Providing materials with full set of desired properties at the lowest price. (ii) Extending the engineering resins’ performance. (iii) Improving specific properties, viz. impact strength or solvent resistance. (iv) Offering the means for industrial and/or municipal plastics waste recycling. Blending also benefits the manufacturer by offering: (i) Improved processability, product uniformity, and scrap reduction. (ii) Quick formulation changes. (iii) Plant flexibility and high productivity. (iv) Reduction of the number of grades that need to be manufactured and stored. (v) Inherent recyclability, etc. 1.3.2

Compatibilization

The topic is extensively treated in Chapters 4 and 5, thus only the key features are mentioned below. Recently published books and reviews also provide extensive information on the topic [Kotliar, 1981; Porter et al., 1989; Porter and Wang, 1992; Brown, 1992; Ajji and Utracki, 1996; Datta and Lohse, 1996; Utracki, 1998a]. It is noteworthy that in the absence of the configurational entropy effects (see Chapter 2), the miscibility of polymer blends depends on the balance of small enthalpic and non-configurational entropic effects. Sensitivity of this balance to small variation of the macromolecular structure is illustrated in the series of papers on miscibility of model polyolefins — e.g., see [Rabeony et al., 1998]. Another example is provided by the photoisomerization initiated, reversible phase separation of PVME blends with stilbene-substituted PS [Ohta et al., 1998]. While miscibility is limited to a specific set of conditions (viz. Appendix 2) the immiscibility dominates — most polymers form immiscible blends that require compatibilization. Alloys’ performance depends on the ingredients, their concentration, and morphology. The alloying process must result in stable and reproducible properties of the polymer blend. Thus, the morphology must either be stable, unchanged during the forming steps, or the changes

13

must be well predicted. The alloying makes use of an appropriate dispersing method (viz. mechanical mixing, solution or latex blending) and compatibilization. The latter process must accomplish three tasks: 1. Reduce the interfacial tension, thus finer dispersion, 2. Stabilize the morphology against thermal or shear effects during the processing steps. 3. Provide interphasial adhesion in the solid state. The compatibilization strategies comprise: (i) addition of a small quantity of co-solvent — a third component, miscible with both phases, (ii) addition of a copolymer whose one part is miscible with one phase and another with another phase, (iii) addition of a large amount of a coreshell copolymer — a compatibilizer-cum-impact modifier, (iv) reactive compounding that leads to modification of at least one macromolecular species that result in development of local miscibility regions, (v) mechano-chemical blending, etc. Commercial alloys may comprise six or more polymeric ingredients. The increased number of components, n, increases the number of interfaces between them: N = n(n-1)/2. For such complex systems it may be necessary to use an ingredient with highly reactive groups, capable of interacting with several polymeric components, e.g., such multicomponent copolymer as ethylene-glycidyl methacrylate, triglycidylisocyanurate, etc. Alternatively, one may carry a sequential blending, incorporating one polymer within another, then combining the pre-blends into the final alloy, hence reducing the number of interfaces that must be controlled at any given time. While the reduction of the interfacial tension coefficient, ν, is relatively easy by introduction of a macromolecular “surfactant”, the stabilization of morphology and improvement of the interphasial adhesion in the solid state, may not be so. One may use either a single compatibilizer that can perform all three compatibilization tasks, or a combination of agents, each playing one or two different roles. For example, stabilization of the desired dispersion (accomplished by addition of “surfactant” to mechanically mixed compound), may be accom-

14

L. A. Utracki

Figure 1.1 Density profile across the interface, defining thickness of the interphase.

plished by partial crosslinking of one of the three phases: matrix, dispersed and the interphase. In the latter case the interphasial adhesion in the solid state is also improved. The density profile across the interface follows an exponential decay (see Figure 1.1). The intercepts of the steepest tangential line with the horizontal lines defining the volume fraction of either one of the two polymeric ingredients, φ = 0 and 1, define the thickness of the interphase, ∆l [Helfand and Tagami, 1971, 1972]. Experimentally ∆l varies from 2 to 60 nm [Kressler et al., 1993; Yukioka and Inoue, 1993, 1994]. Measurements of ∆l have been recently used to map the miscibility region of PC/SAN blends when varying the AN-content and temperature [Li et al., 1999]. For high molecular weight (M → ∞) binary blends, the Helfand and Tagami theory predicts that in binary blends: (i) the interfacial thickness, ¨l’ is inversely proportional to the interfacial tension coefficient,v’, the product, ¨l’v’ being independent of the thermodynamic interaction parameter, χ, (ii) the surface free energy is proportional to χ1/2; (iii) the chain-ends of both polymers concentrate at the interface; (iv) any low molecular

weight third component is repulsed to the interface; and (v) the interfacial tension coefficient increases with molecular weight to an asymptotic value: v = v’ - aoM-2/3. The value of ν can be measured or calculated from the molecular structure of two polymers, e.g., by means of the Hoy’s group contribution method. The computed and experimental values of ν for 46 polymer pair showed good correlation with an average error of ±36% [Luciani et al., 1996, 1997]. The concentration dependence of ν and drop diameter, d, may be expressed as [Tang and Huang, 1994; Ajji and Utracki, 1996]: ν = νCMC + (νο − νCMC) exp{−aχZcφ} ; d = dCMC + (dο − dCMC) exp{−aχZcφ}

(1.5)

where a and a1 are adjustable parameters, and ZC is the copolymer’s degree of polymerization, and subscript CMC indicates the “critical micelles concentration”. It is important to note that ν and the diameter of the dispersed phase follow the same mathematical dependence. The amount of compatibilizer required to saturate the interface, wcr, can be expressed by the two limiting equations:

Introduction to Polymer Blends

wcr = 3φM / aRNA ; wcr = 27φM / [〈r2〉RNA] (1.6) where φ is the volume fraction of the dispersed phase, R is the radius of the dispersed drop, NAv is the Avogadro number, M is the copolymer molecular weight, a • 5 nm2 is the specific crosssectional area of the copolymer macromolecule, 〈r2〉 = KM is square end-to-end distance of the copolymer, and K is the characteristic parameter of the polymeric chain. The first equation in Eq 1.6 was derived assuming that all compatibilizer’s molecules cross the interface once, while the second assuming that di-block copolymer macromolecular coils are randomly deposited on the interface [Mathos, 1993; Ajji and Utracki, 1996]. The reality is somewhere in between these two ideal cases. Both relations predict that the amount of copolymer required to saturate the interface is proportional to the total interfacial area expressed as φ/R. The measurements of ν for the commercial resins’ blends depended on the contact time of the two polymers. Helfand et al. theory predicts that, owing to diffusion of low molecular weight ingredients to the interphase, ν should decrease with time. This indeed was observed for most blends, but an opposite effect was also seen for some PA/PO blends. The effect depended on the nature and amount of antioxidants and stabilizers in each resin. PO’s frequently contain acidic stabilizers, viz. sterically hindered phenols, (hydro)-peroxide decomposers [e.g., tris(2,4-di-tert-butyl phenyl)-phosphite], radical scavengers such as thio-derivatives, etc. [Herbst et al., 1995, 1998]. When blending thus stabilized PO with PA, chemical reactions between the acidic stabilizers and -NH2 of the PA chain ends result in formation of a rigid membrane. Measuring the interfacial tension coefficient as a function of the contact time shows increasing values of the interfacial tension coefficient. These time-effects should be incorporated when predicting the blends’ morphology [Luciani et al., 1996, 1997]. Initially, the most common method of compatibilization was an addition of a third polymeric

15

component, either a block or a graft copolymer. It was assumed that the compatibilizer would migrate to the interface, broadening of the segmental concentration profile, ∆l. There are several reports indicating that addition of a block or graft copolymer reduces ν and alters the molecular structure at the interface, but it rarely increases the interphase thickness. Another disadvantage of the addition method is the tendency for a copolymer to migrate to at least five different locations, forming saturated solutions and micelles in both phases, as well as the interphasial layer (block copolymers may also form mesophases). Hence, the copolymeric compatibilizer that is to be added to a blend should have: (i) maximum miscibility with the respective polymeric components, (ii) molecular weight of each block only slightly higher than the entanglement Me, and (iii) concentration just above CMC. In industry the time effects are important — higher is the viscosity of the blend’s components longer is the diffusion time, thus slower the processing. The morphology of commercial blends usually is far from equilibrium. Preparation of the alloys must take thermodynamic and kinetic parameters into account if the desired effects are to be achieved. The effects of copolymer addition of the dispersion size and blend performance have been studied [Hobbs et al., 1983; Fayt et al., 1986; Armat and Moet, 1993]. From the economic as well as the performance points of view, the reactive compatibilization is most interesting (see Chapter 5). The process involves: (i) sufficient dispersive and distributive mixing to ascertain required renewal of the interface, (ii) presence of a reactive functionality, suitable to react across the interphase, (iii) sufficient reaction rate making it possible to produce sufficient quantity of the compatibilizing copolymer within the residence time of the processing unit. The method leads to particularly thick interphase, thus good stability of morphology. The reactive blending has been used since the beginning of the plastics industry. For example,

16

L. A. Utracki

two polyisoprene isomers, NR and GP, were softened by addition of SCl2 + CS2 and milled in a hot rubber mill. During milling, the solvent partially decomposed co-vulcanizing NR with GP [Parkes, 1846]. A patent from 1939 describes reactive blending of PVAl with multicomponent acrylic copolymers containing maleic anhydride [I. G. Farbenindustrie, 1939]. In the early 1940’s BASF used a corotating TSE for reactive extrusion of PA-6. Since the mid-1960, the reactive extrusion has been used for toughening and general modification of the engineering resins, viz. PA, PET, PC, or PBT. In 1971, Exxon patented styrene-grafting of PE in reactive extrusion followed by blending with PPE for improved processability and performance. Oxazoline-grafted polymers were used as compatibilizers in reactive blending of PC with PA. In 1975 DuPont started to manufacture the super-tough PA, Zytel-ST™, by reactive blending of PA-66 with maleated ethylene-propylene-diene elastomer (EPDM-MA). The reactive compatibilization of the PPE/PA was developed in 1977 [Ueno and Maruyama, 1979]. More details on reactive compatibilization can be found in Chapter 5 as well as in the monograph [Utracki, 1998a]. 1.3.3

Morphology

The morphology depends on the blend concentration. At low concentration of either component the dispersed phase forms nearly spherical drops, then, at higher loading, cylinders, fibers, and sheets are formed. Thus, one may classify the morphology into dispersed at both ends of the concentration scale, and co-continuous in the middle range. The maximum co-continuity occurs at the phase inversion concentration, φI, where the distinction between the dispersed and matrix phase vanishes. The phase inversion concentration and stability of the co-continuous phase structure, depend on the strain and thermal history. For a three-dimensional, 3D, totally immiscible case the percolation theory predicts that φperc = 0.156. In accord with the theory, the transition from dispersed to co-continuous structure occurs at an average volume fraction, φonset = 0.19 ± 0.09

[Lyngaae-Jørgensen and Utracki, 1991b; LyngaaeJørgensen et al., 1999]. The co-continuity contributes to synergism of properties, e.g., advantageous combination of high modulus and high impact strength in commercial blends. Detailed discussion of the phase co-continuity and its effect on morphology and rheology is given in Chapter 7. When discussing the morphology it is useful to use the microrheology as a guide. At low stresses in a steady uniform shear flow, the deformation can be expressed by means of three dimensionless parameters: the viscosity ratio, the capillarity number, and the reduced time, respectively: λ ≡ ηd / ηm ; κ = σd / ν ; t* = t y / κ = γ / κ (1.7) where σ is the local stress, ηd and ηm is the dispersed phase and matrix viscosity, respectively, y is the deformation rate, and d is the droplet diameter. The capillarity number may be used in its reduced form κ∗ ≡ κ / κcr , where the critical capillary number κcr is defined as the minimum capillarity number sufficient to cause breakup of the deformed drop. The drop can break when 1 < κ* < 2. For κ* > 2 the drops deform into stable filaments, which only upon reduction of κ* disintegrate by the capillarity forces into minidroplets. The deformation and breakup processes require time — in shear flows the reduced time to break is tb* • 100. When values of the capillarity number and the reduced time are within the region of drop breakup, the mechanism of breakup depends on the viscosity ratio, λ — in shear flow when λ > 3.8 the drops may only deform but they can not break. Dispersing in extensional flow field is not subjected to this limitation. Furthermore, for this deformation mode κcr (being proportional to drop diameter) is significantly smaller than that in shear [Grace, 1982]. The use of microrheology for the description of drop deformation and break was found to provide a surprisingly good agreement with experimental observations for the morphology evolution during compounding in a TSE [Utracki and Shi, 1992; Shi and Utracki, 1992, 1993]. The predictive model (without adjustable parameters) was further improved by incorporation of the coalescence

Introduction to Polymer Blends

[Huneault et al., 1995]. A similar model was more recently proposed [Moon and Park, 1998]. The flow affects the blend morphology, but the structure variations also engender changes to the rheological response. The flow affects morphology in two ways — it changes the degree and type of dispersion on a local level, and imposes global changes of morphology in formed parts. The latter effects originate from the flow-imposed migration of the dispersed phase that, for example, may cause formation of skin-core structures, weld lines, etc. The flow-imposed morphologies can be classified as: (i) Dispersion (mechanical compatibilization), (ii) Fibrillation, (iii) Flow coalescence, (iv) Interlayer slip, (v) Encapsulation, and others [Utracki, 1995]. Flow may also cause mechano-chemical degradation that generates reactive components, viz., radicals, peroxides, acids, etc. Trans-esterification, trans-amidation and ester-amide exchange reactions during processing are well documented (their rate depends on the total interphasial area that in turn depends on flow). These reactions may be responsible for the formation of compatibilizers that increase the interphasial area, affect the phase equilibria, and the regularity of the main chain, thus modify the degree of dispersion, blend’s crystallinity, hence performance. Miscibility of the blend components has an obvious effect on morphology (for more detailed discussions see Chapter 8). During processing, the hydrostatic and shear stresses can change the lower critical solubility temperature (LCST) by at least 60°C. This may result in formation (inside the processing unit) of a miscible blend. The blend emerging from the extruder may phase separates by the spinodal decomposition mechanism into a co-continuous structure, whose degree of dispersion can be controlled; example: PET/PC blends. Stress-induced fibrillation occurs in a steady state shearing or extension, when the capillarity ratio κ > 2. Under these conditions, flow is co-deformational. Since, κ ∝ d, it is easier to fibrillate coarser b dispersions where φ > aλ (the numerical value of the a and b parameters depends on the composition of the blend) [Krasnikova et al., 1984].

17

Flow through a capillary of POM dispersed in a copolyamide (CPA) at T = Tm(POM) + 6°C resulted in fibrils with diameters of about 20 µm and length 3.2 mm. Fibrillation of POM in EVAc strongly varied with λ. For λ = 1 the finest morphology was found [Tsebrenko et al., 1976, 1982]. At temperatures slightly above the melting point, T > Tm, coalescence combined with stress-induced crystallization resulted in the formation of long fibers. The effect has been explored for performance improvement of blends comprising liquid crystal polymers (LCP) [La Mantia, 1993; Champagne et al., 1996]. The shear-induced interlayer slip was theoretically predicted — it creates a tree-ring structure in the extrudates [Utracki et al., 1986; Utracki, 1991b; Bousmina et al., 1999]. The relation may be used to describe the steady-state viscosity of antagonistically immiscible polymer blends, such as PP/LCP [Ye et al., 1991; Utracki, 1991b]. The shear-induced segregation takes place in any system comprising flow elements with different friction coefficient, either miscible or immiscible [Doi and Onuki, 1992]. Migration of the low viscosity component toward the high stress regions may result in a flow-induced encapsulation. The effect has been well documented and successfully explored in polymer processing [Utracki, 1987, 1988, 1989a, 1991a, 1995]. For example, the high viscosity engineering resins with poor resistance to solvents, e.g., PC, PEST, or PEEK, can be blended with a low melt viscosity LCP. Extrusion through a die with sufficiently long land causes LCP to migrate toward the high stress zone near the die land, thus lubricating the die flow, improving the throughput and enveloping the resin in a protective layer of LCP [Cogswell et al., 1981, 1983, 1984]. 1.3.4

Rheology

The rheology of polymer blends is discussed in Chapter 7. Here only an outline will be given. Since the flow of blends is complex, it is useful to refer to simpler system, e.g., for miscible blends to solutions or a mixture of polymer fractions; for immiscible blends to suspensions or emulsions,

18

L. A. Utracki

and to compatibilized blends to block copolymers [Utracki, 1995]. It is important to remember that the flow behavior of a multiphase system should be determined at a constant stress, not at a constant deformation rate. For miscible blends, the free volume theory predicts a positive deviation from the log-additivity rule, PDB. However, depending on the system and method of preparation, these blends can show either a positive deviation, negative deviation, or additivity [Utracki, 1989a]. Upon mixing, the presence of specific interactions may change the free volume and degree of entanglement, which in turn affect the flow behavior [Steller and Ŧuchowska, 1990; Couchman, 1996]. For immiscible blends the flow is similarly affected, but in addition there are at least three contributing phases: those of polymeric components and the interphase in-between. Flow of suspensions provides good model for blends with high viscosity ratio, λ > 4, while for blends with λ = 1 the emulsion model is preferred. The block copolymer is a good model for well-compatibilized polymer alloys. The fundamental assumption of the classical rheological theories is that the liquid structure is either stable (Newtonian behavior) or its changes are well defined (non-Newtonian behavior). This is rarely the case for flow of multiphase systems. For example, orientation of sheared layers may be responsible for either dilatant or pseudoplastic behavior, while strong interparticle interactions may lead to yield stress or transient behaviors. Liquids with yield stress show a plug flow. As a result, these liquids have drastically reduced extrudate swell, B ≡ d/do (d is diameter of the extrudate, do that of the die) [Utracki et al., 1984]. Since there is no deformation within the plug volume, the molecular theories of elasticity and the relations they provide to correlate, for example either the entrance pressure drop or the extrudate swell, are not applicable. The concentration dependence of the constant-stress viscosity, provides information on the inherent flow mechanism. The experimental data should be evaluated considering the log-additivity rule, lnηb = ∑φilnηi. There are five possible types of behavior, described as: 1. positively

deviating blend (PDB), 2. negatively deviating blends (NDB), 3. log-additive blends, 4. PNDB, and 5. NPDB. These can be described combining the emulsion model of polymer blends with the interlayer slip [Utracki, 1991b; Bousmina et al., 1999]. Owing to variability of the blend structure with flow, the rheological responses are sensitive to the way they are measured. Since the structure depends on strain, the responses measured at high and low values of strain are different. For this reason, the selected test procedure should reflect the final use of the data. When simulation of flow through a die is attempted, the large strain capillary flow is useful. However, when the material characterization is important, the dynamic tests are recommended. The dynamic measurements of polymer blends at small strains are simple and reliable. The storage and loss shear moduli (G’ and G”, respectively) should be first corrected for the yield stress then analyzed for the relaxation spectrum [Utracki and Schlund, 1987; Riemann et al., 1995; Friedrich et al., 1995]. Two types of rheological phenomena can be used for the detection of blend’s miscibility: (1) influence of polydispersity on the rheological functions, and (2) the inherent nature of the two-phase flow. The first type draws conclusions about miscibility from, e.g., coordinates of the relaxation spectrum maximum; cross-point coordinates (Gx, ωx) [Zeichner and Patel, 1981]; free volume gradient of viscosity: α = d(1nη) / df; the initial slope of the stress growth function: S = d(lnη + E )/dlnt; the power-law exponent n = d(lnσ12)/dlny ≅ S, etc. The second type involves evaluation of the: extrudate swell parameter, B ≡ D/Do, strain (or form) recovery, apparent yield stress, etc. Compatibilization enhances dispersion, increases the total apparent volume of the dispersed phase, rigidifies the interface, and increases interactions not only between the two phases, but also between the dispersed drops. These changes usually increase the blends’ viscosity, elasticity and the yield stress. The compatibilizer effects are especially evident at low frequencies. There are two mechanisms that may further affect these behaviors: (i) the

Introduction to Polymer Blends

copolymer may form micelles inside one or both polymeric phases instead of migrating to the interphase, and (ii) an addition of compatibilizer may increase the free volume resulting in decreased viscosity. The time-temperature, t-T, superposition principle is not valid even in miscible blends well above the glass transition temperature, Tg [Cavaille et al., 1987; Ngai and Plazek, 1990; Chung et al., 1994]. In miscible blends, as either the concentration or temperature changes, the chain mobility changes and relaxation spectra of polymeric components in the blends show different temperature dependence, thus the t-T principle cannot be obeyed. Furthermore, at the test temperatures, the polymeric components are at different distance from their respective glass transition temperatures, T - Tg1  T - Tg2, what affects not only the t-T superposition, but also the physical aging time — [Maurer et al., 1985]. In immiscible PO blends, such as PE/PP, at best, the superposition is limited to the melt within narrow temperature ranges [Dumoulin, 1988]. For most blends, the morphology changes with the imposed strain. Thus, it is expected that the dynamic low strain data will not follow the pattern observed for the steady-state flow. One may formulate it more strongly: in polymer blends the material morphology and the flow behavior depend on the deformation field, thus under different flow conditions, different materials are being tested. Even if low strain dynamic data could be generalized using the t-T principle, those determined in the steady state will not follow the pattern. Chuang and Han [1984] reported that for blends at constant composition the plots of N1 vs. σ12 and G’ vs. G” are independent of T. However, for immiscible blends, the steady state relation may be quite different from the dynamic one. The agreement can be improved by means of the Sprigg’s theory [Utracki, 1989a]. Four measures of melt elasticity have been used: the first normal stress difference, N1, the storage modulus, G’, and the two indirect ones, the entrance-exit pressure drop, Pe (or Bagley correction), and the extrudate swell, B. In homogeneous melts, the four measures are in a

19

qualitative agreement. In polymer blends where the dispersed phase is rigid, B and is small. By contrast, for the readily deformable dispersed phase, the deformation-and-recovery mechanism provides a potent mechanism for energy storage leading to a large elastic response. In short, neither Bagley’s entrance-exit pressure drop correction, Pe, nor the extrudate swell, B, should be used as a measure of blends’ elasticity. In both cases not the molecular deformation, but the form recovery dominates the observed dependence. Two contributions to the tensile stress growth function, η1+, should be distinguished: one due to the linear viscoelastic response, η+EL, and the other originating in the structural change of the specimen during deformation, η+ES. The first can be calculated from any linear viscoelastic function, while the second, depends on the intermolecular interactions or entanglements, and its value depends on the total strain, , and either strain rate or straining time, t [Utracki, 1988, 1989, 1995; Takahashi, 1996]. Owing to the industrial importance of strain hardening, SH ≡ log(η+ES/η+EL), a large body of literature focuses on the optimization of blend composition to maximize SH. Since SH depends on the entanglement, either inter-chain reactions that lead to branched macromolecules, blending linear polymers with branched ones, synthesizing bimodal resins, or widening the molecular weight distribution may result in improved SH. Extensive work on SH has been done for PE blends, especially the ones comprising LDPE [Utracki and Schlund, 1987]. More recently, several other resins with long chain branching (viz. bPC, bPP, or a biodegradable polybutylenesuccinate, etc.) have been introduced as special grades for, e.g., film blowing, blow molding, wire coating, or foaming [Imaizumi et al., 1998]. The convergent flow at the die entrance provides strong elongational flow. In 1989 Laun and Schuch, derived for Newtonian liquids that Pe § 1.64σ12, The relation is satisfactory for homopolymers, but for the blend, the prediction is about one decade too low. On the other hand, this type of flow provides excellent means for mixing highly viscous dispersed phase. An extensional

20

L. A. Utracki

flow mixer (EFM) was developed. The device provides good mixing for multicomponent polymer systems, e.g., for blends with components having widely different viscosities, viz. PE with UHMWPE, PP with high elasticity EPR, PC with PTFE [Nguyen and Utracki, 1995; Utracki and Luciani, 1996; Luciani and Utracki, 1996]. 1.3.5

Developing Commercial Blends

There are several methods of blending, viz., mechanical (dominant), solution, latex, fine powder, as well as several techniques adopted from the IPN technology. Not always the finest dispersion is desirable — the size and shape of the dispersed phase must be optimized considering the final performance of the blend. Polymer blends’ performance depends on the properties of ingredients’, their content, and morphology. Since the cost is virtually fixed by the material and the compounding method, the economy depends on blend’s morphology, tailored for a specific application. Blends have been developed for the economic reasons, viz., improvement of either a specific property (e.g., impact strength) or engendering a full set of required properties, extending engineering resin performance, improving processability, recycling, etc. There are several approaches to evaluation of the blend economy. For example, the cost of a blend equals the weight-average of material cost, plus the compounding cost per unit mass, e.g., Cb = ΣwiCi + K. Another approach is to calculate the cost-to-performance ratios for diverse materials and/or compositions. For example, one may ask how much a unit of the tensile modulus or the strength at yield will cost, and optimize the composition accordingly. However, with growing frequency the blend economy is based on the replacement calculations, comprising the total cost, that of material, compounding, forming, assembling, customer satisfaction, esthetics, service life-spans, then the ease of disposal or recycling. For the resin manufacturers, blending provides means to improve and broaden the resin performance and therefore it enhances the demands

and sale. By contrast, the resin users start with a set of performance parameters that the material must possess. In both cases the basic preposition is the same: to have a blend with desired characteristics, one must use a component that already shows this characteristic, or simply, one cannot create something out of nothing. While extension of the engineering resin performance constitutes the largest part of the high performance blends’ production, the most difficult and interesting task is the development of blends with a full set of desired properties. To achieve this goal a systematic approach has been developed [Utracki, 1994]. The procedure starts with the selection of blend components, each possessing at least one of the desired properties. For example, to improve impact strength an elastomer should be used, to induce flame retardancy a nonflammable polymer, to improve modulus a stiffer resin should be incorporated, etc. Since for each property there are several candidates to select from, the selection is guided by the principle of the compensation of properties — advantages of one component should compensate for deficiencies of the other, e.g., the disadvantages of PPE (processability and impact strength) can be compensated for by those of HIPS. Next, the method of compatibilization, compounding, and processing must be selected. Since polymer blends’ performance depends on morphology, the goal is to ascertain the desired structure by selecting an appropriate resin grade (rheology), as well as the methods of compatibilization, compounding, and processing. Interesting studies on the morphology development during dispersive mixing were published by Kozlowski [1994, 1995]. In this fundamental work, a rotating disc mixer was used. The disk had a milled grove in which stationary spreader was inserted. The gap clearance, speed of rotation, temperature, shape of the spreader, and pressure were controlled. The device simulated the dispersive processes that take place in internal mixers and extruders. A model of step-wise generation of morphology was proposed, where the original pellet (of the dispersed phase) undergoes deformation into elongated plates, that under stress

Introduction to Polymer Blends

break into fibers and finally into drops. The final morphology is a result of dispersion and coalescence processes that depend on the viscoelastic character of the component, the interfacial tension properties, and the stress history (see Chapter 9 on compounding). The ideal compounding unit should have: (i) uniform elongational and shear stress field, (ii) flexible control of temperature, pressure and residence time, (iii) capability for homogenization of liquids having widely different rheological properties, (iv) efficient homogenization before onset of degradation, and (v) flexibility for the controllable change of mixing parameters. Most blends described in the patent literature have been prepared using either an internal mixer or a single screw extruder (SSE). In standard configuration SSE is inadequate for the preparation of blends with controlled morphology. Furthermore, due to presence of “dead spaces,” the run-to-run reproducibility of the SSE-extruded blends may be poor. SSE should not be used for reactive blending. However, there are several designs of mixing screws, profiled barrel elements, and add-on mixing devices that ameliorate SSE mixing capability. From between the latter devices, RAPRA’s cavity transfer mixer (CTM), or the newly patented extensional flow mixer (EFM) should be mentioned. The first of these is a sort of “dynamic motionless mixer”, where material is transferred from the cavities in the barrel to those on the screw, enhancing the distributive mixing [Gale, 1980]. The EFM is a motionless device in which the extensional forces provide dispersive mixing for blends with components having widely different viscosities, viz. PE with UHMWPE, PP with high elasticity EPR, PC with PTFE, gel particles in reactor powders, etc. [Utracki and Luciani, 1996]. More expensive but easier to control is a twin screw extruder, TSE. Owing to the modular design with many types of elements fulfilling different functions, TSE can be optimized for specific tasks. The ratio of the dispersive to distributive mixing can be adjusted, and the width of the residence time can be controlled. TSE is excellent chemical reactor for polymerization, modification

21

of polymers, and reactive compatibilization. As a result, the blend quality and run-to-run reproducibility are improved. Computer models have been developed to predict variation of blend’s morphology along the screw length in these machines [Shi and Utracki, 1992, 1993; Huneault et al., 1995]. 1.3.6

Blends’ Performance

The quality of compounded blend affects the processing and performance. Layering, poor weld lines in injection molded parts, skin-core extrudate structure with low notched Izod impact strength, all indicate poor blend quality — either not adequate dispersion or poor stabilization morphology. Compounding demands precise control of process variables. At the present most alloys are prepared by reactive processing. It has been reported that pellet-blending of two blend lots may lead to apparent immiscibility and bad weld line strength. Evidently, even a small variation in the extend of reaction may make them immiscible. The mixed lots may pass standard tests, but still yield unacceptable products. Historically, blending was used to improve impact strength of the early resins, i.e., toughening of PS, PVC, PMMA, PET, PA, etc. With time, blends evolved into multi-polymer systems, that not only have to be impact-modified, but also compatibilized. Many blends have been formulated with a multicomponent modifier that simultaneously compatibilizes and impact-modifies the mixtures. The following observations can be made: (1) The maximum toughening of brittle polymers has been obtained dispersing ca. 10 vol% of a ductile resin with domain diameter of d ≈ 0.1-1.0 µm — the lower limit is for resins that fracture by the shear-banding, whereas the higher for those that fracture by the crazing-and-cracking [Bucknall, 1977; Bucknall et al., 1984]. (2) The phase co-continuity provides the best balance of properties, e.g., high rigidity in the presence of large deformability (or elongation). The properties depend on the thickness of the interpenetrating strands, thus also on compatibilization. (3) For the

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L. A. Utracki

best barrier properties the lamellar structure is desirable. To create it, the blend should comprise large but stable drops with diameter d § 5 - 50 µm. During bi-axial stretching (e.g., in blow molding or film blowing processes) the large drops easily deform into lamellae. Two types of mechanical tests are recognized: low speed (tensile, compressive, or bending), and high speed (impact). Poor compatibilization affects both. For example, in tensile tests the maximum strain at break, and the yield stress can be dramatically reduced by poor inter-domain adhesion. Similarly, the lack of adhesion is responsible for low impact strength — the specimens are brittle. Several toughening mechanisms have been proposed, viz. crazing, shear-banding, cavitation, particle debonding, elastic deformation of the toughening particles, etc. [Arends, 1996]. Polymeric systems are roughly classified as either brittle or pseudo-ductile. The first type has low crack initiation as well as propagation energy and it fails by the crazing-&-cracking mechanism. Typical examples are: PS, PMMA and SAN. The second type has high crack initiation energy, but low crack propagation energy and it tends to fail by yielding and shear-banding. Typical examples are: PA, PEST and PC. As usual,

there are some polymers, e.g., POM and PVC, that show intermediate behaviors — in many systems the fracture takes place by a mixed mode. The transition from brittle to ductile mode of fracture depends on the intrinsic properties of the material as well as on the external variables such as geometry, temperature, loading mode, test rate, etc. To detect the mechanism of fracture the stress-strain, and the volume-strain dependencies should be known. “Toughness” is defined as the total area under the stress-strain curve, thus abruptly ending curves without the yield point are characteristic of brittle materials. The volumestrain dependence provides means for quantitative identification of the fracture mode — pure shear banding shows no volume expansion, whereas pure crazing and cracking show the maximum volume expansion. Wu [1985, 1990] postulated that the brittle/ ductile behavior of a neat amorphous polymer is controlled by two intrinsic molecular parameters: the entanglement density, νe, and the chain stiffness (given by the characteristic chain constant C’). Assuming that crazing involves chain scission, the stress, σz, should be proportional to νe½, and the yield stress, σy, proportional to C’. In consequence: σ z / σ y ∝ ν e 3Mv / ρa , where Mv

Figure 1.2. Izod impact strength at room temperature as a function of diameter of elastomeric particles in methylmethacrylatebutadiene-styrene copolymer used for toughening polyvinylchloride resin [after Bertelo and Mori, 1994].

Introduction to Polymer Blends

is the average molecular weight of a statistical segment, and ρa is the polymer density. For brittle polymers: νe < 7.5. Effectiveness of the toughening process depends on the diameter of the elastomeric particles and their concentration. At constant concentration the toughness vs. particle diameter dependence goes through a maximum — abscissa of its location, dopt, does not depend on concentration, whereas the ordinate does (see Figure 1.2). The optimum particle size, dopt, depends on the entanglement density of the matrix resin, as well as on the fracturing and toughening mechanisms. However, the determination of the dopt may be ambiguous, owing to the polydispersity of sizes as well as to inclusion of the matrix polymer inside the elastomeric particle. For example, it has been accepted that to toughen PS into HIPS, the optimum diameter is defined as a diameter of the elastomeric particles expanded by inclusion of the PS mini-drops. In PVC, the diameter of the elastomer was defined as the diameter of the original butadiene latex particle before grafting it with styrene and methylmethacrylate. In PC, the “optimum” diameter was defined by availability of the core-shell toughening agent — it is difficult to find toughners with elastomeric particles having diameter d < 100 nm. The strategy for the preparation of polymer blends with stable morphology demands that blends have thick interphase, ∆l ” 60 nm. Frequently it is impossible to decide how far the toughening by rubber core extends into the interphase. Many impact-resistant engineering resin blends have been formulated using a core-shell multicomponent copolymer with a rigid core and elastomeric shell whose thickness and affinity with the matrix resin was adjusted. In many cases, toughening of a brittle polymer can be achieved by introduction of stiffness heterogeneity, viz. incorporation of an elastomer, immiscible polymer, solid particles, gas bubbles (i.e., foaming or micro-foaming), etc. However, the size and concentration of these heterogeneities should be optimized. For most thermoplastic’s the optimum diameter of the dispersed elastomeric particle is d < 3 µm and its volume fraction: 0.05 < φ < 0.10. The accepted mechanism of

23

toughening considers the heterogeneity to be a stress concentrator, generating excessive crazing and/or shear banding of the matrix, thus requiring higher amount of energy to cause fracture. The stress concentration factor was defined as: γ ∝ 1/(1-aφ2/3) where a depends on the matrix [Bucknall, 1977; Partridge, 1992]. For blends with pseudo-ductile matrix dopt depends on concentration, as it is important to keep the distance between the elastomeric particles approximately constant. During the early works on compatibilization of PE/PS blends in Prof. Heikens laboratories it was noted that addition of a small amount of one polymer to another improved impact properties. Since these two polymers are antagonistically immiscible and upon solidification void formed around the dispersed particles, it was concluded that it is presence of the voids that accounts for the toughening effect [Heikens, 1982]. About 15 odd years later, the microcellular blends have been introduced. For example, micro-foamed blends of HDPE with PP (using CO2 in an autoclave) showed significantly improved impact strength [Dorudiani et al., 1998]. Similar enhancement of mechanical performance was reported earlier for N2-microfoamed PS, SAN or PC [Collias and Baird, 1995]. During the last few years, microfoaming is being used not only to reduce the material consumption, part weight (by 30 to 50%), but also to improve the mechanical performance, especially of the injection molded parts. Under the triaxial stresses in the region ahead of the sharp crack, a particle may cavitate at a certain strain, changing the stress field of the matrix from the dilatation to the distortion dominated. Thus, the matrix may deform plastically, what consumes energy. The mechanism depends on the size of the dispersed toughening particles and the inherent plastic deformation capability of the matrix [Borggreve and Gaymans, 1989; Lazzeri and Bucknall, 1993; Yee and Shi, 1995; Groeninckx et al., 1995]. In PC cavitation occurred well before shear yielding [Parker et al., 1992]. Blends comprising relatively high concentration of two engineering resins may require toughening of both phases by sequential reactive blending.

24

L. A. Utracki

Formation of co-continuous structures in blends of either a brittle or pseudo-ductile resin with an elastomer may result in a quantum jump of toughness, without greatly affecting the key engineering properties of the high performance resin. Commercial blends of this type of, e.g., POM, PA, PC, or PET with an elastomer, are available (viz. Triax™ series). 1.3.5

Evolution of Polymer Alloys and Blends

The historical evolution of the polymer blend technology is presented in the following order: 1. Commodity resins (Styrenics, PVC, Acrylics, PE’s, PP). 2. Engineering resins (PA, PEST, PC, POM, PPE), and 3. Specialty resins (PSF, PAE, PARA, PAr, PPS, LCP, PEI, PEA, etc.). Blends of polymer-A with polymer-B will be discussed following the adopted rules: (i) symbol A/B is used to identify any mixture of polymer A with B, independently of the concentration range or morphology, and (ii) the A/B blends are discussed under the name of the lower category polymer, i.e., blends of engineering or specialty polymers with a commodity resin are discussed in the category of commodity resin blends, blends of specialty polymers with engineering resins are discussed in the category of engineering resins, hence “Specialty resin blends” consider only mixtures of two (or more) specialty resins.

1.4

Commodity Resins and their Blends

Five large-volume polymeric groups belong to this category: polyethylenes, polypropylenes, styrenics, acrylics and vinyls. Their world market share remains relatively stable — the commodity resins represent 71% of all consumed plastics.

1.4.1

Polystyrene (PS)

Simon in 1839 named the distillate of Storax officinalis a styrol. By 1845, the thermal polymerization of styrene as well as the thermal depolymerization of PS were known. In 1915 I. G. Farbenindustrie started commercial production of PS, Trolitul™. Until the 1950’s, PS was produced in small quantities — the resin was brittle, thermally unstable, with poor solvent and scratch resistance. The main use of styrene was in the manufacture of styrenics, viz. Buna-S, SBR, or ABS. Common PS is atactic and amorphous. It has good optical clarity, low dielectric loss factor, modulus E = 3.2 GPa, strength σ = 45-65 MPa, 3 density ρ = 1050 kg/m , and CUT = 50-70°C. Because of brittleness and low chemical resistance the demand for neat PS has decreased, and except for foaming, PS is rarely used. PS can also be polymerized into crystalline forms: isotactic (iPS) or syndiotactic (sPS) with Tm = 230 or 272°C, respectively. The former was polymerized using Ziegler-Natta catalyst [Ishihara et al., 1986], while the latter using a single-site metallocene titanium-based catalyst [Imabayashi et al., 1994]. The high-impact PS, HIPS, has been known since 1911 [Matthews, 1911, 1913]. In the USA, Ostromislensky [1924, 1926-28] patented copolymerization of styrene with “rubber, balata or other elastic and plastic gum”. Production of HIPS, Victron™, by the Naugatuck Chemical started in 1925, but soon it was discontinued. PS is miscible with several polymers, viz. polyphenyleneether (PPE), polyvinylmethylether (PVME), poly-2-chlorostyrene (PCS), polymethylstyrene (PMS), polycarbonate of tetramethyl bisphenol-A (TMPC), co-polycarbonate of bisphenol-A and tetramethyl bisphenol-A, polycyclohexyl acrylate (PCHA), polyethylmethacrylate (PEMA), poly-n-propyl methacrylate (PPMA), polycyclohexyl methacrylate (PCHMA), copolymers of cyclohexyl methacrylate and methyl methacrylate, bromobenzylated- or sulfonatedPPE, etc. Other miscible blends are listed in Appendix 2.

Introduction to Polymer Blends

Similarly, poly-α-methylstyrene is miscible with PMMA, PEMA, PBMA, and PCHMA. Poly-p-methylstyrene, and poly-p-t-butylstyrene show miscibility with polyalkyl(meth)acrylates. However, PS is immiscible with PMMA, PMA, polyethylacrylate (PEA), polybutylacrylate (PBA), or PBMA [Somani and Shaw, 1981]. In miscible blends it is important that both components are in the entangled state. In particular, during processing in the extensional flow field (e.g., blow molding, film blowing, wire coating, calendering or foaming) an enhancement of strain hardening (SH) can only be obtained when the concentration of the high molecular weight component is at least comparable to the critical concentration of entanglement, c • c*. Under these circumstances large increases of SH were observed, e.g., for PS blended with ultra high molecular weight PS (UHMW-PS), or SAN blended with ultra high molecular weight PMMA (UHMW-PMMA). By contrast, addition of immiscible UHMW-PS to SAN did not show any improvement of SH [Takashi, 1996; Takahashi et al., 1996; Koyama et al., 1997; Minegishi

Table 1.9.

25

et al., 1997, 1998]. Examples of blends that were evaluated for SH are listed in Table 1.9. 1.4.1.1 PS/Commodity Resin Blends The most common immiscible PS blends are those prepared to improve the impact strength of PS or its copolymers, viz. HIPS or SBR, viz. Table 1.10. It was reported that incorporation of 0.1-18 vol% of either acrylic or olefinic elastomer particles (e.g., in HIPS) into a thermoplastic (viz. PE, PP, PS, SAN, PEST, PPE/HIPS, PC, PEI, PA, fluoropolymers, etc.) resulted in excellent control of the foaming process [Campbell and Rasmussen, 1994]. The bubble diameter could be calculated from the concentration of rubber particles. When these were lightly crosslinked, the stretched membrane provided an excellent barrier against coalescence of gas bubbles. Thus, reliable nucleation and absence of coalescence lead to foaming stability. For example, in autoclave foaming of PS with N2, the cell size was less than 40 µm, independently of the saturation pressure, and only slightly increasing with the foaming temperature.

Strain hardening in molten polymer blends

Polymers

Type

SH behavior [Reference]

PMMA/UHMW-PMMA

miscible

PS/UHMW-PS

miscible

SAN/UHMW-PS SAN/UHMW-PMMA PE/UHMW-PE Bimodal PO’s LLDPE/LLDPE

immiscible miscible miscible miscible miscible

LLDPE/LDPE

immiscible

LDPE/PS

with 0 - 5 wt% of SEBS compatibilized miscible

General rule: large enhancement of SH only for: c(UHMW) > c* = 1.2 wt %; SH ” 10 [Minegishi et al., 1997] large enhancement of SH increasing with T; SH ” 10 [Minegishi et al., 1998] at 145°C no effect on SH; SH ” 2 [Koyama et al., 1997] large enhancement of SH; SH ” 10 [Koyama et al., 1997] large enhancement of SH; SH ” 10 [Okamoto et al., 1998; Kotaka, 1998] large enhancement of SH; SH ” 10 [Münstedt and Kurzbeck, 1998] average SH for the narrow & broad MWD LLDPE’s [Schlund and Utracki, 1987] Presence of LDPE increased linearly SH of LLDPE [Schlund and Utracki, 1987] enhancement of SH for c • 10 wt% LDPE SEBS (micelles in LDPE phase) reduced the strain at break; additive SH; yield stress [Utracki and Sammut, 1990] better SH; the blends are suitable for foaming [Horiuchi, 1998] 30 wt% of PVDF reduced SH of PMMA at 160°C [Kotaka, 1998]

LDPE/PS PMMA/PVDF

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L. A. Utracki

Table 1.10. Toughening of PS Composition

Reason

Reference

PS with SBR PS with PIB and PSIB PS with PB and SBR PS with SBR and a crosslinked SBR PS with SBR and PEG

impact performance toughness toughness toughness toughness, adhesion, electrostatic dissipation to improve weatherability enhanced properties low density PO foams for marine or submarine applications

Seymour, 1951 Sparks and Turner, 1952 Hayes, 1956, 1967 Conrad and Reid, 1963 Briggs and Price, 1963

toughness enhanced properties transparent HIPS

Grancio et al., 1981, 1983 Gausepohl et al., 1982 Asahi Chemical Industry, 1982

enhanced properties for adhesiveness

Murray, 1982 Shiraki et al., 1986

HIPS (PS toughened by styrene-grafted EPR) HIPS with SBS PB reacted with styrene, ethyl- or methyl-styrene, t-butyl styrene and/or vinyl silanes PS with SBR, (SB)n and PP HIPS with a star block copolymer, (SB)n HIPS with a SB-block copolymer having small size of the rubber particles HIPS with HDPE and SEBS HIPS with BS(B’S’)n terminated with 2,4-di tert-butyl-p-cresol and tris(nonyl-phenyl)phosphite

Hostyren™ Durst, 1970 Dawans and Binet, 1981

Postulating that the rubber particles are stretched to membranes all having the same thickness, the foam cell size can be expressed as: 3 Dcell = Do + ( drubber / nt )1/ 2 ; n = 3 to 6

(1.8)

where Dcell is the cell size, drubber is the initial diameter of rubber particle, Do is the diameter of foam cell in the absence of rubber particles and t is thickness of the rubber shell after foaming. Depending on the initial assumption of the cell geometry, the geometrical factor can be calculated as: n = 3 to 6. Adequacy of Eq 1.8 is shown in Figure 1.3. The second large group of styrenic blends comprises these with polyolefins — they are summarized in Tables 1.11 and 1.12. These blends are mainly used in packaging. Formulated for extrusion, injection and blow molding, they show excellent processability, improved impact strength, low moisture absorption and shrinkage.

Figure 1.3. Cell size vs. the initial rubber particle diameter. Data points — [Campbell and Rasmussen, 1994], solid line — Eq 1.8 with n = 6 and the cell wall thickness: t = 85 nm.

27

Introduction to Polymer Blends

The performance characteristics (e.g., modulus, toughness, ductility, transparency, or gloss) can be controlled by the composition and morphology. Other patents described similar blends prepared either by different methods or comprising different compatibilizer. For example, PO was mixed with styrene in the presence of an initiator that caused polymerization at temperatures below melting point of PO [Vestberg and Lehtiniemi, 1994]. Interesting blends, having a broad range of properties, were prepared in two steps: 1. BR was grafted and crosslinked with either styrene or methylmethacrylate to produce a core-shell copolymer. 2. Next, it was blended with PO for improved processability, impact resistance, rigidity, etc. [Aoyama et al., 1993, 1994]. Structural blends of styrene-grafted PP with either SBR, SBS, or an acrylic elastomer were developed [DeNicola and Conboy, 1994]. Since the early 1990’s the constrained geometry metallocene catalysts have been used by Dow to produce either alternating or pseudo-random ethylene-co-styrene interpolymers (ESI) [Stevens et al., 1991]. ESI with up to 50 wt% styrene is semicrystalline, it is known to compatibilize PE/PS

blends since it forms domain structures into which the homopolymers can dissolve. ESI also has good melt strength, mechanical, impact and damping characteristics [Ellebrach and Chum, 1998]. Flow and processing information on ESI were published [Karjala et al., 1998]. Himont (now Montell Polyolefins) introduced a family of PP-based blends under the trade name of Hivalloy™. Some grades seem to be mechanical alloys of PP/PS, compatibilized and impact modified by incorporation of SEBS and EPR. Others are reactor-made — here porous grains of PP serve as reaction beds for the polymerization and grafting of PS, SMA, acrylics, etc. PS is one of the most frequently foamable thermoplastic resin. Blends that belong to this category are presented in Table 1.13. Recently blends containing • 50 wt% PS (MW = 200 kg/mol) and acrylic copolymers were described as particularly useful for the manufacture of low density foams. The acrylic copolymer contained methylmethacrylate and, e.g., 5 wt% of ethylacrylate. The presence of copolymer facilitated foaming, but it reduced the foam compressive strength. The best balance was obtained using about 22 wt% of the copolymer.

Table 1.11. PS or HIPS with polyolefins Additive

Reference

PE and CSR PE and styrene-ethylene bulk copolymer PO and EVAc Either PP or PE 5-95 wt% PS with 95-5 wt% PO and 0.5 to 10 wt% SEBS (foaming) PP and SEBS 1 to 99 wt% of either LDPE, LLDPE, HDPE, EVAc, PP or EPR (chemical foaming to open-cell structures) LLDPE

Herbing and Salyer, 1963 Gorham and Farnham, 1964 Yamamoto et al., 1971 Ogawa et al., 1973 Zeitler and Mueller-Tamm, 1977 Holden and Gouw, 1979 Tashiro et al., 1983

Reprocessed PE/PS, compatibilized with 0.5 - 40 wt% of EVAc • 60 wt% of partially neutralized ionic PS (e.g., a copolymer of styrene and acrylic acid) with • 1 wt% PE ionomers, foamed with 3 - 20 wt% of NaHCO3 PE or PP compatibilized with a non symmetrical 3-block copolymer, S1-D-S2 HDPE and either SBS or SIS

Canterino & Freudemann, 1985; Canterino et al., 1986 McCullough and Stevens, 1985 Park, 1986a, 1986b, 1988 Hoenl et al., 1993 Swartzmiller et al., 1993, 1994

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Table 1.12. Addition of styrenics to PO Additive to PE

Reason

1. Addition of Styrenics to PE SAN 5-20 wt% SBR SB, SBS, or a p-methylstyrene-b-isoprene copolymer Styrene-grafted HDPE ABS with CPE or low molecular weight PS SEBS SBS or SEBS as compatibilizers for PS/PE HIPS with HDPE and SEBS PS/LDPE; MI-ratio R = MIPS/MIPE = 7 to 90, weight ratio

improved crack resistance impact strength improved crack resistance higher modulus impact strength mechanical properties recyclability impact strength

Jankens, 1963 Dow Chem., 1963 Minekawa et al., 1971 Yui et al., 1978 Kamosaki et al., 1978 Davison & Gergen, 1980 Lindsey et al., 1981 Castelein, 1982

flexibility, dumping small vibrations, stiffness, heat insulation, low water permeability

Hoki and Miura, 1987

dimensionally stable, ρ = 20-30 kg/m3, cells 1.5 mm studies of the rheology during closed cell foaming

Park, 1986c, 1987

Y = 10 to Y = [394.7 + 1.18R - 295.1 log D]; density: D = ρ = 15 to 30 kg/m3. R & Y control the co-continuous blends morphology 5-50 wt% PS with LDPE a chemical blowing agent and a peroxide (0.05 - 0.1 pph DCP) LDPE/PS = 80/20 blends compatibilized in a twin-screw extruder with supercritical CO2 2. Addition of Styrenics to PP PS or HIPS ABS with either CPE or low molecular weight PS PS with HIPS and SEBS PS was compatibilized by adding either SBS or (SB)n PS or HIPS and a non-symmetrical, linear 3-block copolymer of styrene and butadiene, S1-D-S2, where the polystyrene blocks S1 • S2 PS with recycled PP — co-continuous morphology sPP-co-sPS from single-site metallocene catalyst PP, EPR, EVAc, and PS blended with 1 to 50 wt% of the silane-modified based resin, crosslinking catalyst and 1 to 20 wt% of a foaming agent

The foam had closed cells with cell diameter varying from 0.1 to 1 mm [Smith and Cross, 1996]. 1.4.1.2 PS/Engineering Resin Blends The majority of PS blends that belong to this category are mixtures with PPE. Discovery of PPE miscibility with PS led to a family of Noryl™ blends, commercialized in 1965. Since that time,

for nacreous soda-straw tubes impact strength mechanical performance higher modulus processability, impact & stress cracking resistance, impermeability to H2O performance, recyclability compatibilization of sPP/sPS resilient foams with superior, compression strength and heat-insulating properties

Reference

Lee et al., 1998

Ogawa et al., 1973 Kamosaki et al., 1978 Holden and Gouw, 1979 Grancio et al., 1981, 1983 Hoenl et al., 1993

Morrow et al., 1994 Razavi, 1994 Kobayashi et al., 1997

the PPE/PS blends were modified by the incorporation of a variety of additives. Worldwide, there are six manufacturers of PPE. The PPE/PS blends show the glass transition temperature, Tg = 100 to 210°C, continuously increasing with PPE content. The most often used compositions contain less than 30 wt% of PPE (PPE is about three times more expensive than PS).

29

Introduction to Polymer Blends

Table 1.13. Foamable PS blends Blend

Comment

Reference

PS with low concentration of siloxane-oxyalkylene copolymers

reduced interface tension, easier bubble nucleation and growth, smaller cells molded articles had thin skin and uniformly foamed interior

Granda et al., 1977

Siedenstrang and Thorsrud, 1984

chemical foaming (NaHCO3)

Sprenkle, 1980

PS with radial teleblock SBS, PMS, naphthenic extender oil, formed at P = 1.4 MPa then foamed using microwaves at 10 to 100 MHz PS with SAN, ABS and 15 - 50 wt% of SMA

during injection molding 75 to 98 wt% of PS, 2 to 25 wt% of a BR easy to foam due to fine dispersion (MW = 200 to 300 kg/mol) of BR spheres Mixtures of emulsion copolymers comprising: acrylonitrile, low density foams for non-wovens, butadiene, styrene, and acrylic or methacrylic acid carpets, fleece or cardboard SBR blends with 10 to 50 wt% of plastisol PVC for foamed, flame-resistant carpet backing SBR, ABS, MABS, and/or SBS, with either PS, PO, PVC, for high impact strength moldings PPE, PA, POM, PC, PSF, or PEST Polymer alloys of SMA and cellulose esters reaction between anhydride and at a ratio varying from 1:100 to 100:1 cellulosic -OH facilitated foaming and gave good product performance 5 to 35 wt% of SEBS, 65 to 90 wt% of PB-1, foams had excellent bending capability, 5-30 wt% of EPR or EPDM, 2-15 wt% of a LDPE tear strength, stiffness, and HDT Two SBR copolymers, (1) with of 53-75 wt% of cured foams for shoe soles with high styrene and (2) with 42-75 wt% of styrene shock absorption ABS compositions foamed with supercritical CO2 at P ” 5 MPa 45-90 wt% of PS or styrene-acrylic acid copolymer physically foamed products had with 10-55 wt% of PVDC or improved O2 and H2O permeability, vinylidenechloride-methyl acrylate copolymer toughness, and flame resistance PS, SMA, SAN, PMS or HIPS blended with SBS and A general method for the production an extending oil, then incorporated into: PA, PEST, of a variety of foamable injection PPS, SAN, ABS, ASA, PC, PPE, PO, moldings their copolymers or blends 75-97 wt% of either PS or HIPS and 3-25 wt% of foamable materials with good an elastomeric (co)polymer having a Tg < -20°C performance characteristics Latex copolymers were blended, crosslinked and foamed: open-cell foams with great absorbency (1) 20 wt% styrene, 20 wt% divinylbenzene, 60 wt% were prepared for baby dippers, 2-ethylhexyl acrylate with (2) 80 wt% styrene, 20 wt% for paint rollers, filters, etc. divinylbenzene ABS with ASA and two SAN copolymers were foamed easy formability, excellent physical with a physical foaming agent properties and Freon resistance Blends of NR, SBR, BR and SB copolymer that had foamable rubber blend, suitable for 0-30 wt% of styrene, and MW • 30 kg/mol tires or belts

Henn et al., 1996 Matner et al., 1977 Morgan and Ribaudo, 1980 Aoki et al., 1981 Heslinga and Greidanus, 1982 Hwo, 1996 Hashimoto and Ohashi, 1985 Kumar et al., 1995 Romesberg, 1991

Burnell, 1993

Blumenstein et al., 1994 Brownscombe et al., 1997

Kim and Choi, 1998 Kawauzra et al., 1997

30

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Table 1.14. Examples of PS/PPE blends Modifier of PPE/PS Blend

Reference

Elastomers such as PB, SBR, or NBR PB Poly(methylmethacrylate-co-styrene) and PO Either SBR or ABS Vinyl-terminated ethylene-propylene-styrene terpolymer (SEP) PA-66 Foaming with dry gases generated by the thermal decomposition of a dihydro-oxadiazinone + azodicarboxylic acid amide or ester PPE-polyolefin graft copolymer and NBR Epoxy-terminated liquid PB, with either PP-MA or SEBS SBR and SBS copolymer ABS and SAN Hydroxynaphthoic acid Ethylene glycol-propylene glycol copolymer (PEG-PPG) SBR and radial-SB copolymer PPE/PS closed-cell insulating foams, with high compressive strength PPE/SAN with crosslinking and C3 - C6 hydrocarbon blowing agents PPE with, e.g., PS, PMS, PES, PEI, PC, PA, PEST, PP or PE), and the blowing and nucleating agents

Lauchlan and Shaw, 1970 Huels, 1971 Izawa et al., 1973 Nishioka et al., 1973 Haaf, 1979 Mitsubishi Petrochemical, 1982 Kochanowski, 1982

PPE is the most “natural” additive that upgrades performance of PS to the required level. PS/PPE blends have been used as a replacement for PS in applications where higher HDT and/or impact strength is required. These alloys are easy to foam for the manufacture of, e.g., hot water piping insulation, in automotive applications, etc. Examples of PPE/PS blends are listed in Table 1.14. Five Caril™ grades of expandable PPE/PS beads (diameter 0.3 to 0.5 mm) offer HDT up to 120°C, thus are suitable for the production of microwaveble and steam-cleanable packaging with the wall thickness • 1 mm. The recommended density of molded product is ρ = 60 kg/m3. Other foamable, flame retardant PPE/PS blends, with good acoustic and thermal insulation properties, have been produced in suspension polymerization of a PPE solution in styrene and pentane. Resulting beads had diameter d = 0.5 to 1.0 mm and could

Mitsubishi Petrochemical, 1983 Mitsubishi Petrochemical, 1983 Mitsubishi Gas, 1985 Japan Synthetic Rubber, 1985 Tamura, 1985 Vaughan, 1985 Sugio et al., 1987 Allen et al., 1989; Weber et al., 1990 Hahn et al., 1992 Bland and Conte, 1993

be steam pre-foamed and compression molded in a standard equipment. The cited advantages are: high HDT, non-flammability, dimensional stability, strength, stiffness, low molding cost, low density, easy lamination with decorative and weatherresistant ASA, and recyclability [Koetzing and Diebold, 1995]. PS is antagonistically immiscible with all other engineering resins, viz. PA, PC, POM, and PEST. PS has been added to these polymers to improve processability and reduce cost without unduly affecting the performance (the so-called, “extension of the engineering performance”). Non-compatibilized blends of PS with either PEST or PEST and PMMA have been used for decorative applications or as the so-called plastic paper [Kamata et al., 1980]. Similarly, PAr blends with either SAN [Brandstetter et al., 1983], or high performance blends of LCP with thermoplastic polymers (e.g., PP, PS, PC, PI) [Haghighat

31

Introduction to Polymer Blends

et al., 1992], showed adequate performance for the envisaged applications. However, most PS blends with engineering resins require compatibilization. Thus, for example, PS with PA-6 was compatibilized by addition of either methylmethacrylate-styrene copolymer (SMM) [Fayt et al., 1986], or SMA (e.g., used in PARA/PS blends) [Lee and Char, 1994]. POM was blended with a small amount of either PS, poly(α-methyl styrene) (MPS), or SAN, and with particulate fillers [Tajima et al., 1991]. PAr/PS blends were compatibilized with PAr-PS segmented copolymer [Unitika Ltd., 1983]. Several blends comprising PC and diverse styrenics, viz., ABS, SAN, SB, SBS, MBS, etc., are known (see Table 1.15). Similarly as for PVC blends (see Table 1.16) the strong interactions between AN and carbonyl groups of PC (in PVC it is the tertiary carbon) are responsible for the good performance. An interesting variation of the compatibilization procedure involved dispersing PC in water with vinyl monomer(s) that subsequently were polymerized. The in situ formed graft copolymer acted as a compatibilizer [Kanai et al., 1978; Kakizaki et al., 1979]. In 1974, polyphenylenesulfide, PPS, was blended with either PS or a styrene-copolymer [Miyanishi, 1976]. Recently, to provide a complete set of the required performance characteristics, multicomponent blends are being promoted. For example, PC, PPE, ASA, SAN, PS, phosphate esters, PTFE, and SEBS [Niessner et al., 1993], or PC, PEST, ABS modified by incorporation of alkyl (meth)acrylates and glycidyl methacrylate, and PPE with either

PS, HIPS, or SEBS, and a polyalkyl(meth)acrylate [Laughner, 1993]. 1.4.2

Acrylonitrile-Butadiene-Styrene (ABS)

The first mechanical blends of NBR with SAN, known as “type-A ABS,” date from 1936. In the mid-1940’s, Dow started emulsion polymerization of “ABS type-G”. By the late 1950’s, the high heat ABS were invented, viz. interpolymers of α-methylstyrene and acrylonitrile [Irving, 1961]; a mixture of methylmethacrylate-α−methylstyrene either with styrene-grafted polybutadiene (SBR) or with an ABS [Kanegafuchi, 1967]; a mixture of SMA and ABS [Stafford and Adams, 1972]; a mixture of SMA with ABS and MBS [Tatuhiko and Akira, 1982]; a mixture of SMA-MMA with ABS, etc. Acrylonitrile-styrene-acrylate terpolymers, known as either ASA or AAS, constitute another class of ABS resins, viz. Centrex™, Luran™ S, Richform™, etc. These materials may also contain reactive groups, viz. maleic anhydride or glycidyl methacrylate. Weather-resistant ABS can be obtained either by the incorporation of EVAc [Fukushima and Mitarai, 1971], or by replacing PB with EPDM, to obtain AES [Wefer, 1984, 1985, 1988]. Alternatively, blends of SAN with maleated EPDM and CPE may be used [Kim et al., 1994]. However, the non-weatherable styrenics are frequently prepared by dissolving an elastomer in methyl methacrylate, and either styrene or α-methylstyrene, then polymerizing them into methyl methacrylate-butadiene-styrene graft copolymers

Table 1.15. Compatibilization of PS/PC blends by SAN Additive to PC

Reason

Reference

Either SAN & styrene-allyl methacrylate-butyl acrylate-methyl methacrylate copolymer, or with a multilayered copolymer from styrene, allyl methacrylate, benzyl acrylate, divinylbenzene SAN and a styrene-grafted acrylic rubber PS and MBS

toughening, high mechanical performance, solvent resistance

Kishida et al., 1978

improved mechanical properties higher modulus

Kamata et al., 1979 Lee, 1980

32

L. A. Utracki

(MBS) [Ruffing et al., 1964; Schmitt et al., 1967]. There is a great diversity of the MBS copolymers, viz. graft, core-shell, or multilayer type — lately also with acidic or epoxy groups [Lee and Trementozzi, 1979, 1980-82; McKee et al., 1982; Keskkula et al., 1984]. 1.4.2.1 ABS/SMA Blends The ABS/SMA blends show excellent processability, high heat deflection temperature (HDT) low warpage, stiffness at high temperature, good impact strength, as well as solvent and chemical resistance. They successfully compete with PPE or PC alloys for the automotive applications (trim, instrument panels, roof linings, hubcaps, headlight housings), electronics, and electrical industry, houseware, appliances, power tools,

industrial machinery, plumbing products, parts for washing machines and vacuum cleaners, etc. An example of commercial blends is Cadon™. 1.4.2.2 ABS/PVC Blends There are several reasons for blending PVC with ABS-type copolymers, viz. to improve processability, mechanical properties, and low-temperature toughness. Good properties of these blends originate from the miscibility between PVC and SAN-part of ABS. In some commercial blends, viz. Geloy™, ABS may be replaced by ASA to obtain improved miscibility and weatherability. For enhancement of HDT, SMA may also be added. The blends with more than 30 wt% PVC are selfextinguishing, but are more difficult to process.

Table 1.16. PVC/ABS-type blends Additive to PVC

Reason

Reference

ABS

for either phonographic records, or artificial leathers high impact strength, mechanical properties improved processability, impact strength, and thermal stability improved toughness improved impact strength impact strength processability, and high impact resistance impact strength processability, high impact strength, mechanical properties processability, high impact strength, mechanical properties abrasion resistance processability, high notched impact strength economy, high HDT and impact strength

Parker, 1951; Schule, 1952 Himei et al., 1967

5-30 wt% of either acrylonitrile — butadiene — styrene-methylmethacrylate (MABS) 1-50 wt% ABS and post-chlorinated PVC ABS grafted with acrylonitrile-ethyl acrylate-styrene ABS and SBS MBS or MABS Multilayer butadiene-styrene-divinylbenzene-butylacrylate-methyl methacrylate Poly(2-cyano-5-norbornene) and ABS PB grafted with styrene, methylmethacrylate, and maleic anhydride (ABSM-MA) or a mixture of ABS and SMM-MA Methylstyrene-styrene-acrylonitrile-grafted polybutadiene or with maleated styrene-methyl-methacrylate-butadiene (ABS-MA) ABS and vinylchloride-ethylhexyl acrylate Core-shell copolymer: EPDM grafted with styrene-butadiene methacrylate or allyl cyanurate CPVC and PMMA, methylstyrene-acrylonitrile-methyl methacrylate, methylstyrene-acrylonitrile-styrene, imidized-PMMA, imidized-SMA, and SAN

Kojima et al., 1970 Tanaka et al., 1971 Minekawa et al., 1971 Kumabe et al., 1973 Usami and Ochiai, 1976 Matsuura et al., 1978 Dufour, 1982 Dufour, 1988 Greenlee et al., 1992 Siol et al., 1993 Soby et al., 1994

33

Introduction to Polymer Blends

1.4.2.3 ABS/PC Blends Blends of PC with 5-70 wt% ABS were developed in the early 1960’s. The basic technology has been used to produce such alloys as Bayblend™, Cycoloy™, Idemitsu™ PC/ABS, Iupilon™, or Triax™ 2000. The consumption of ABS/PC blends is increasing as the cost-to-performance ratio is low, and properties are predictable. The alloys combine good processability of ABS with excellent mechanical properties, impact and heat resistance of PC. The opaque blends show dimensional stability, low shrinkage and moisture absorption, high stiffness and hardness, good impact resistance at temperatures T • -50°C, excellent UV stability, processability, mechanical properties, heat resistance, flame retardancy, good chemical resistance, but poor to gasoline, aromatic hydrocarbons, esters, ketones and some chlorinated hydrocarbon. The ABS/PC blends are being manufactured with either a dispersed or co-continuous morphology. There are many similarities between ABS/PVC and ABS/PC blends. Both are immiscible, having three distinct phases of PVC or PC, SAN, and an elastomer [Suarez and Barlow, 1984]. The blends are compatibilized by the dipole-dipole interactions between PC and SAN, particularly evident

in SAN with • 25 wt% AN [Kim and Burns, 1988]. ABS/PC blends can also be compatibilized by incorporation of either acrylic, acidic or epoxy groups (see Table 1.17). In the late 1970’s, the reactive blending of PC/ABS has begun to dominate the technology. Initially, the PC blends with ABS modified by incorporation of the maleic anhydride moieties (ABS-MA), later ABS with acrylic acid groups (ABS-AA) were developed. The third generation blends comprise ABS modified by copolymerization with glycidylmethacrylate (ABS-GMA). Examples are listed in Table 1.18. In 1983, Monsanto developed blends with cocontinuous morphology, Triax™ 2000. These alloys comprised PC, ABS, and styrene-methylmethacrylate-maleic anhydride (SMMA-MA) [Jones and Mendelson, 1985]. One year later, PC was reactively blended with either ABS, SAN-GMA, and NBR, or with graft copolymers of acrylonitrile-butadiene-αmethyl styrene-methyl-methacrylate (MeABS) and acrylonitrile-α-methyl styrene-methyl methacrylate copolymer (MeSAN) [Kress et al., 1986]. The blends were commercialized by Bayer as Bayblend™. In 1992, low-gloss, moldable blends, with electrostatic discharge properties were developed. They comprised PC, ABS, and either a graft copolymer of styrene, hydroxyethyl methacrylate, and

Table 1.17. PC/ABS-type blends Reason

Additive to PC/ABS

Reference

Toughness HDT and stiffness HDT an impact resistance Processability, HDT, impact resistance Flow, weatherability, thermal, and mechanical performance Pearl-like iridescence, dyeability Heat resistance, dimensional stability Mechanical performance Processability, impact strength

10-70 wt% ABS styrene-α-methylstyrene-acrylonitrile, PSF MBS and PAES polyethersulfone, PES EVAc

Grabowski, 1964a Grabowski, 1970-72 Yamauchi et al., 1974 Weaver, 1972 Hasegawa et al., 1974

PMMA PVC CPE grafted with SAN skin-core graft copolymers of styrene and acrylonitrile on elastomeric latex particles MBS and acrylic elastomer

Ikura et al., 1974 Hardt et al., 1975; Kabuki et al., 1973 Sakano et al., 1978

Solvent and impact resistance

Kitamura, 1986

34

L. A. Utracki

acrylonitrile bonded onto a 1,3-butadiene rubber (ABS-HEMA), styrene-acrylonitrile-methacrylic acid copolymer (SAN-MAc), styrene-acrylonitrilehydroxyethyl methacrylate (SAN-HEMA) or an acrylonitrile polymer containing gels [Vilasagar and Rawlings, 1994]. Cycoloy™ is the PC/ABS blend from General Electric Co. 1.4.2.4 ABS/PA Blends ABS/PA mixtures are immiscible, hence the standard three strategies are applicable: (i) Addition of a small amount of ABS to improve PA toughness without a compatibilizer. (ii) Generation of non-compatibilized blends with co-continuous morphology. (iii) Compatibilized blends in the full range of composition. ABS is an amorphous resin, while PA’s are semi-crystalline, hence it is advantageous to incorporate ABS as either a dispersed or a co-continuous phase — the latter being preferred. However, addition of semicrystalline PA to ABS increases mold shrinkage, thus addition of filler is advised. Owing to high processing temperatures of PA, it is essential to use high heat ABS. For the adequate impact performance at least 10 wt% of ABS should be

added, but at this level, the compatibilization is required. The reactive compatibilization involves the use of ABS that has been modified by incorporation of either acrylic acid, maleic anhydride, or polyvinylphenol (PVPh). The reason for blending ABS with PA is to reduce moisture sensitivity, improve toughness, and reduce shrinkage and warpage of the latter resin. The alloys show good processability, surface finish, high heat stability, a chemical, oil, wear, and abrasion resistance, dimensional stability, low temperature impact strength, reduced moisture sensitivity, and economy. Synergistic properties have been reported. Examples of commercial alloys are: Stapron™ N, Novalloy™-A, Techniace™ TA, Triax™ 1000, Ultramid™, Macslloy™ [Utracki, 1994]. The first ABS/PA blends were announced in 1961 [Grabowski, 1964b, 1966], and much later introduced as Elemid™. Triax™-1000 is an alloy of PA-66 with ABS-MA, having the phase co-continuity [Lavengood et al., 1986, 1988]. PA-6 was also blended with BR grafted with styrene and MA (SBMA) [Asahi-Dow Ltd., 1981]. Later, transparent blends of copolyamide(s), PA, and ABS were developed [Fox et al., 1989].

Table 1.18. PC/ABS reactive blends Composition

Reason

Reference

PC with ABS and rubber-modified SMA PC with SAA and EAMM PC with ACM and SAA PC/ABS with EAA acidic compatibilizer PC/ABS with SMA-AA PC/ABS with SMM-GMA PC/MBS with SAN and PEST PBT, PC, ABS and PB grafted with acrylate esters and AN, ACM PC, PEST, polyester carbonate, etc., with 30-90 wt% of SMA + ABS and 2.5-20 wt% of a chemical blowing agent PC, PEST or PEI with 1-50 wt% of ABS and a chemical foaming agent

processability, impact strength, heat resistance impact strength, mechanical properties impact strength and HDT processability and impact strength high HDT and impact strength processability, impact strength, and heat resistance impact strength and thermal stability rigidity, flowability, solvent resistance, impact strength, dimensional stability foamable engineering blends having excellent physical performance

Henton, 1980, 1982 Thomas, 1982 Henton, 1984 Grigo et al., 1984 Brandstetter et al., 1982-83 Daicel, 1982-84 Teijin Chem., 1980 Bier and Indner, 1982; Neuray et al., 1982 White and Krishnan, 1989

moldable blends for chemical foaming

Allen and Avakian, 1987

35

Introduction to Polymer Blends

Blending either ABS-MA or EPR-MA, with amine-terminated PA or PEST resulted in alloys with excellent performance [Akkapeddi et al., 1990, 1992]. Similarly, either ABS-MA or ABSGMA copolymer was used to compatibilize and to toughen PA blends with other resins, viz. PC, PEST, or PAr [Yuichi and Suehiro, 1989]. Recently the role of elastomer, its type and location in the PA-66/SAN/Elastomer system has been studied [Nair et al., 1998]. 1.4.2.5 ABS/PEST Blends The thermoplastic polyesters (PEST) are dominated by two resins: polyethyleneterephthalate (PET) and polybutyleneterephthalate (PBT). There are similarities between ABS/PA and ABS/PEST blends. In blends with ABS a part of PEST may be replaced by PC, and 10-20 wt% of an impact modifier may also be added, e.g., MBS, poly(methylmethacrylate-g-butadiene-co-styrene), poly(MMA-g-n-BuA), high rubber ABS (•50 wt% PB), or ASA with •50 wt% acrylate rubber, etc. Examples of commercial blends are: Alphaloy™ MPB, Cycolin™, Diaaloy™ B, Malecca™ B, Maxloy™, Lumax™, Triax™ 4000, Ultrablend™ S. The alloys show excellent moldability, low post-molding shrinkage and warpage, stress-crack resistance, high gloss, high temperature stiffness, toughness and mechanical strength, high heat resistance at temperatures T • 140°C,

low shrinkage, good dimensional stability, impact strength, good wear and abrasion resistance, good thermal and weathering resistance, as well as solvent (e.g., to gasoline and motor oils) and chemical resistance. An abbreviated evolution of the PEST/ ABS technology is summarized in Table 1.19. 1.4.2.6 ABS/TPU Blends Developed in the early 1960’s ABS/TPU blends combine TPU’s toughness and paintability with ABS’s low temperature impact strength and adequate HDT. The main advantage is the excellent impact behavior at T ” -40°C. Furthermore, TPU improves anti-friction properties, abrasion and chemical resistance. Stiffness is also increased and the flowability of injection molding compounds is good. ABS is usually dispersed in the TPU matrix. TPU was also successfully blended with SBR grafted with acrylonitrile, acrylate or methacrylate esters (MABS) [Abe et al., 1977], with SMM-MA copolymer, [Gomez, 1992], and with bulk-polymerized ABS [Henton et al., 1992]. Depending on the type of TPU, compatibilization may be necessary. Examples of commercial alloys are Prevail™ and Techniace™ TU. 1.4.2.7 ABS/PSF Blends In these blends, ABS role is to improve flowability and reduce cost, while that of PSF is to

Table 1.19. PEST/ABS blends Composition

Reason

Reference

PET with MBS PAr with ABS PBT/ABS/SEBS PBT with carboxyl-modified ABS PBT + PET or PC, with ABS or ACM (rubber particle diameters d ≅ 0.4 µm) PBT with either ABS-MA or ABS-GMA PET or PBT with an AES-GMA (Techniace™)

notched impact strength, & embrittlement resistance processability and impact strength stable morphology chemical, solvent, and impact resistance impact strength and balance of other properties

Sauers and Barth, 1970 Koshimo, 1973 Gergen and Davison, 1978 Tanaka et al., 1979 Bier and Indner, 1982; Binsack et al., 1982 Orikasa et al., 1989 Hirai et al., 1988, 1989, 1992

heat, chemical & impact resistance flowability and good balance of properties

36

L. A. Utracki

improve the shape retention at high temperatures. ABS/PSF blends are compatibilized either by Phenoxy, EVAc-GMA, or SMA copolymers. They have good processability, high notched Izod impact strength, plateability, hydrolytic stability, and economy. However, they may show poor surface, and weld-line strength. Arylon™, and Mindel™ A are examples of the commercial ABS/PSF alloys, while Ucardel™ is an example of PSF blends with SAN. Evolution of ABS/PSF blends’ technology is summarized in Table 1.20. 1.4.3

SBS Block Copolymers

In 1961, using lithium catalyst, a series of styrene-isoprene (SI) and styrene-butadiene (SB) block copolymers were synthesized [Bull and Holden, 1977]. The resins had Tg = -90 to +90°C. Full-scale production started in 1965. Since then, numerous two- and three- block copolymers have been developed. More recently, hydrogenated and during the last few years maleated block copolymers are being offered. With the world consumption of 330 kton/y, the block copolymers constitute the largest part of the commercial TPE market. Large quantity of SBS resin is used

in blends. Commercial resins include: Elexar™, Collimate™, Finaclear™, Kraton™, Thermolastic™, Tufprene™. 1.4.3.1 SBS/SEBS Blends with Commodity Resins SBS copolymers are used in blends as compatibilizers, impact modifiers, or stabilizers of morphology and performance. As shown in Table 1.21, they have been frequently incorporated along other styrenics. There are many patents for mixtures of PO with SBS-type copolymers. Their history is outlined in Table 1.22. 1.4.3.2 SBS Blends with Engineering Resins Since the anionically polymerized block copolymers are relatively expensive they have been more frequently used in blends with engineering than commodity resins. Owing to miscibility of styrene blocks with PPE, the SBS and SEBS are “natural” tougheners for this polymer. However, for blending with PEST, PC, POM or PA, the copolymer should be modified by incorporation of acidic, acid anhydride, or epoxy moieties.

Table 1.20. PSF/ABS blends Composition

Reason

Reference

PSF with 40-52 wt% ABS, and poly(α-methyl styrene-co-AN) Polyarylethersulfone (PAES) with AES ABS with equal amount of PSF and PC PAES with EVAc and/or MBS PAES with PC and either MBS or ABS PSF with 10 wt% anhydride-terminated PSF, and MABS PSF with AES PSF with crosslinked acrylate copolymer, crosslinked SAN, and uncrosslinked SAN PSF/ABS with EVAc-GMA PSF with 25-45 wt% semi-crystalline PPS, and 0-10 wt% MBS

excellent flow, good impact resistance, non-flammability toughness, and impact strength processability, rigidity, and impact strength tensile, flexural, and impact strength high HDT, and impact resistance excellent HDT, and impact resistance

Ingulli & Alter, 1969, 1970

thermal stability, and impact strength good tensile modulus, yield strength, impact resistance, and respectable HDT = 106°C processability, HDT and impact strength resistance to impact, high temperatures, and adverse environmental conditions

Sumitomo, 1982 Robeson, 1985

Barth, 1970 Grabowski, 1971-72 Lauchlan, 1971 Yamauchi et al., 1974 Aya et al., 1979

Orikasa & Sakazume, 1990 Golovoy & Cheung, 1994

37

Introduction to Polymer Blends

Table 1.21. SBS blends with other styrenics Composition

Reason

Reference

PVC with ABS and SBS SBS with PS, EVAc and other ingredients HIPS with 12.5 wt% PB & SBS SIS with PS and/or IR PS and/or HIPS with PP and SEBS Poly-p-methylstyrene (PpMS) with SBS AN-grafted SEBS with SAN PS with (SB)n and SBR SEBS dissolved in styrene, methacrylic acid, and isoprene, then polymerized SEBS-type IPN with carbon black, CB PS with AXBXA or (AXB)n (A = styrene, B = butadiene, X = AB tapered block) SBS (acidic, amino, imido-terminated) and PA, PEST, TPU, POM, PVAl, PC, PSF, PPE, PPS, or PVC SMMA, a tapered SB and ductile SBS, Zylar™

toughness & performance for elastic films excellent impact strength optical and mechanical properties impact & tensile strength, solvent resistance impact strength, and clarity weatherability, impact strength impact strength, transparency thermoplastic IPN, with superior mechanical properties electrically conductive blends impact strength, and transparency, superior to that observed for SBS/PS blends water-swelling materials for civil engineering, construction, etc. transparent (extremely low haze), impact resistant chemical, solvent, and stress cracking resistance transparent, low haze, high impact, craze, and γ-radiation resistance

Minekawa et al., 1971 Hinselmann et al., 1973 Durst, 1970, 1975 Kawai et al., 1978 Holden and Gouw, 1979 Sherman, 1981, 1983 Paddock, 1981 Asahi Chemical, 1982 Siegfried et al., 1984

HIPS, PE, and either SBS or SIS; a co-continuous morphology SMMA and either a mixture of SBS and a tapered BSB triblock copolymer, or SBR

Sorensen, 1984 Toyama et al., 1985 Shiraki and Hattori, 1986 Blasius, 1992, 1994 Hoenl et al., 1993; Seelert et al., 1993 Hauser et al., 1993

Table 1.22. SBS blends with polyolefins Composition

Reason

Reference

PP with 20 wt% of either SBS or SIS

high impact strength, without adverse effect on other properties transparency and impact strength processability, mechanical properties performance, co-continuous morphology dispersed drops of d = 200 nm moldability and mechanical properties superior performance over HDPE/PS impact strength transparent, impact resistant impact strength, lack of yellowing rigid blends with good impact strength and adhesion to solids moldable, good impact strength recycled commingled scrap Elastomeric films moldable resin with good impact strength, scratch and abrasion resistance stand alone structural materials

Japan Synthetic Chemical Ind., 1971

PP blends with 6-8 wt% SEBS SBS with 20-30 wt% PO PS with HIPS, PP and SEBS SEBS with diverse polymers, including PE PP/PS with either SBS or (SB)n HDPE with PS and SEBS HIPS with HDPE SEBS LLDPE with SEBS PS, LLDPE and SEBS PP and PET reactively blended with maleated SEBS PA/PO/SEBS compatibilized by SMA PO with PA, PET and styrenics SBS, EVAc, PS and LLDPE or ULDPE PO with SEBS, SEPS, SEB grafted with maleic anhydride acrylic or sulfonic acid PP with either SBR, SBS, or an acrylic elastomer, and PP grafted with styrene

Porter, 1972 Tabana and Maki, 1976 Holden & Gouw, 1979 Davison & Gergen, 1980 Grancio et al., 1981, 1983 Lindsey et al., 1981 Castelein, 1982 Holden & Hansen, 1989 Seelert et al., 1993 Tekkanat et al., 1993, 1994 Chundury, 1993, 1994 Weber et al., 1994 Djiauw, 1994 O’Leary and Musgrave, 1993 DeNicola and Conboy, 1994

38

L. A. Utracki

1.4.3.2.1 SBS with Polyphenylene ether (PPE) Evolution of PPE blends with SBS-type block copolymers is summarized in Table 1.23. SBS or its derivatives have been frequently used to stabilize the morphology in the newer, more complex blends. In Table 1.24 examples of this type of system are presented. 1.4.3.2.2 SBS with Polyamides (PA) SBS or SEBS has been used as an impact modifier in PPE/PA blends, with PA usually being the matrix, and PPE an organic, low density filler. The blends were developed in the early 1970’s by the Asahi Chemical. By the end of the decade the first reactive blends were announced by the Sumitomo Chemical [Ueno and Maruyama, 1981] and General Electric [Van der Meer et al., 1989]. The simple, SBS/PA blends were in parallel development with the PPE/SBS/PA ones. Addition of SBS to PA improved the tensile and impact strength of the latter resin. The blends

comprise either 1-25 wt% SBS as a dispersed or at higher concentration as co-continuous phase (see Table 1.25). 1.4.3.2.3

SBS with Thermoplastic Polyesters (PEST) The development of PEST/SBS blends parallels that of PA/SBS ones. First, blends of PBT, ABS, and SEBS were disclosed [Gergen and Davison, 1978]. Four years later the reactive compatibilization was discovered — PBT was blended with SEBS and SMA [Durbin et al., 1983]. By the end of 1970’s, multicomponent blends comprising, PBT, PET, PC, and either SEBS, (SB)n, butadiene-caprolactone-styrene, or butadiene-caprolactone block copolymer, were developed [Wambach and Dieck, 1980]. Reactive compatibilization of PEST/SEBS by addition of MA was disclosed in 1984. The method was general, applicable to polyamides as well as to polyesters [Shiraishi and Goto, 1986].

Table 1.23. SBS blends with polyphenyleneether PPE Modifiers

Reason

Reference

PPE with 10-90 wt% SBS PPE with SEBS SBS or EPDM-modified PS and SEBS HIPS and SB di-block copolymer Styrene-phenyl-maleimide copolymer and either SB, SBS, or SEBS HIPS, styrene-grafted EPDM and/or SB grafted with EGMA SBR, SBS and EPR HIPS, SEBS and PE Styrene-grafted PPE, PPE-S with SBR and SB

processability and toughness processability flow, impact and thermal properties impact strength & solvent resistance HDT, chemical, solvent & impact resistance

Kambour, 1970 Haaf, 1979 Lee, 1979-80, 1982-83, 1985 Brandstetter et al., 1982 Fukuda and Kasahara, 1982

impact strength and processability

Ueno et al., 1982

processability & impact strength processability and impact strength processability, gloss, toughness and tensile strength impact strength, processability, and flame retardancy moldability, toughness, strength plateability & mechanical properties high impact strength cracking and impact resistance

Mitsubishi Gas, 1982 Haaf, 1983 Izawa et al., 1983

HIPS, SBS, SBR, EPR and hydrogenated poly(bisphenol-A-phosphite) ABS and SEBS PS, SBR and SBS HIPS, SEBS and LLDPE HIPS and either SB, SBS, or (SB)n

Sugio et al., 1984 Ueda and Sasame, 1986 Mitsubishi Gas Co., 1985 Hambrecht et al., 1986 DeMunck & Lohmeijer, 1986

39

Introduction to Polymer Blends

Table 1.24. SBS in PPE multicomponent blends Additives

Reference

PPE with PA-66, PS, maleated-PP and SEBS PPE reactively blended with SBS & MA, then mixed with PA or PEST Reactive blends of carboxylated PPE with PA-66 and SEBS PPE, PA, SBS and a reactive mixture of styrene-glycidylmethacrylate with styrene and a peroxide PPE, PBT, SEBS, and PC [PBT — matrix; PPE + SEBS — dispersed phase; PC at the interface] PPE, HIPS, PEST, PS with reactive (2-oxazoline) groups, PC and SBS PPE with PBT (or PET), SEBS, PC and mica PPE-grafted with fumaric acid, reactively blended with PC and SEBS Grafted PPE, blended with dimethylsiloxanes, PC, PBT, and SEBS PPE, with PBT, PC, SEBS and/or acrylate copolymer PPE, HIPS, an ethylene-methacrylic acid copolymer, EMAA, SEBS and SGMA PPE with PVDF, SEBS, and poly(styrene-co-methyl-methacrylate) PPE with HIPS, PE and SEBS PPE with PP and SEBS PPE/PET or PPE/PA reactively compatibilized with SEBS-GMA PPE/PET or PPE/PA reactively compatibilized with SEBS-MA PPE/PBT, toughened by addition of urea-butylated resin and SEBS PPE with PC, PBT, and either SBS, SEBS or a core-shell copolymer Epoxy- or phosphate- functionalized PPE, with PBT or PET, palmitamide, SEBS and PC PPE with PA-66, SEBS, SB, a styrene-butadiene radial copolymer, citric acid and either citric acid or chloro-epoxy triazine PPE with PA-6 or PA-66, MA and toughened with SB PPE-MA or PPE-GMA, with sPS, SEBS and fillers

Fujii et al., 1987 Van der Meer and Yates, 1987 Grant and Jalbert, 1987, 1989 Mawatari et al., 1987 Brown et al., 1987 Avakian et al., 1988 Yates, 1987, 1989 Ishihara, 1989 Brown, 1992 Yates and Lee, 1990 Fuji and Ting, 1987 Van der Meer et al., 1989 Ting, 1990 Lee, 1990 Mayumi and Omori, 1988 Modic and Gelles, 1988 Mizuno and Maruyama, 1990 Brown and Fewkes, 1992, 1994 Yates, 1993 Gianchandai et al., 1993 Lee, 1994 Okada and Masuyama, 1994

Table 1.25. Evolution of the SBS/PA blends Composition

Reason

Reference

PA-66, SEBS, Phenoxy and bisphenol-A-epichlorohydrin PA-6 with SEBS PA-6 or PA-66 with SBS PA-6 with SEBS-MA and/or LLDPE PA-6 and SEBS-MA

tensile and impact strength of PA

Freed, 1975

toughness, balance of properties improved toughness and reduced modulus tensile and impact strength processability, mechanical properties, impact strength high notched Izod impact strength moldable alloys with high impact strength good balance of strength and toughness processability, mechanical properties, and low temperature toughness processability and toughness

Bull and Holden, 1977 Cerny & Troncy, 1981 Mitsubishi Chem., 1982 Asahi Chem., 1983, 1984, 1987

PA, or PEST with SEBS-MA PA-66 with SEBS, SEBS-MA, PO COPO, with PA-6 and SEBS-MA PA + acidified SEBS, EPR or EPDM; the adduct incorporated into PA PA-66 + 1:1 blend SEBS and SEBS-MA

Gelles et al., 1987 Gelles et al., 1988 Machado, 1992 Ohmae et al., 1991, 1992 Gelles et al., 1994

40

L. A. Utracki

1.4.3.2.4 SBS with Polycarbonates (PC) Similarly to blends of SBS with PA or PEST, these with PC were first described in 1976. However, owing to weak interactions between SBS and PC, these systems require compatibilization. Thus, either SBS must be acidified (e.g., with SEBS-MA), or acidified acrylate added, viz. MABS, MBS, SMA, etc. Selected examples are listed in Table 1.26. 1.4.3.3 SBS Blends with Specialty Resins SEBS must be processed below 280°C, thus its use with specialty resins has been limited to polyphenylenesulfide (PPS) [Garcia and Martinovich, 1984]. Sometimes SBS and a specialty resin are parts of a multicomponent blend, viz. PPS, PPE, either PA-6 or PA-12, SEBS (Kraton™ G), an acidified polyolefin (e.g., EPR-MA, PE-GMA, or EVAc-GMA), and reactive compatibilizer (citric, maleic, or agaric acid). Here, PPS was a matrix, PA was the dispersed phase that contained PPE/ SEBS and a filler [Ishida and Kabaya, 1994]. 1.4.4

Polyvinylchloride (PVC)

PVC was first synthesized by Regnault in 1835. The first patent on vinylchloride monomer (VCM) polymerization was granted in 1912

to Ostromislensky. However, to make commercially viable articles, PVC must be stabilized and either plasticized or blended. In 1927, B. F. Goodrich started production of plasticized PVC, Vinylite™ [Semon, 1933]. The first patented PVC alloys were prepared by latex blending with PVAc, and poly(vinylchloride-co-vinylacetate) (PVCAc) [Voss, and Dickhäuser, 1930, 1933, 1935, 1936]. I. G. Farbenindustrie commercialized PVC extruder-blended with polyacrylic ester — the so called “rigid formulation” [Fikentscher and Heuce, 1930; Fikentscher, and Wolff, 1931]. Troluloid™ and Astralon™ were the first commercial thermoplastic polymer blends. PVC blended with Buna-N produced excellent thermoplastic materials [Badum, 1942]. These blends were prepared either in a rubber-mill, by latex blending, or powder blending then extruding. The rigid PVC not only had higher heat HDT than the flexible one, but it was permanently plasticized. In 1940 also B. F. Goodrich patented the NBR/PVC blends. Many forms of PVC and its copolymers have been developed over the years to fit specific uses, viz. latex, plastisol, organosol, flexible, and mostly the rigid formulations. In 1991 world production of PVC was 22.0 Mton, or 21.6 wt% of the thermoplastic resin market.

Table 1.26. SBS/PC blends Composition

Reference

PC with 30 wt% SEBS for co-continuous morphology PC with PE and hydrogenated and chlorosulfonated SBS PEST with PC, SEBS and mineral filers PC/ABS/PP compatibilized and toughened by SEBS-MA PC with SB-teleblock and SEBS PC with either SBS and MBA, or SEBS, EEA and LLDPE PC with PE and SEBS PC with either SBS, EGMA or MBS PC with SMA and SBS PC, COPO, PEST, SEBS + butylacrylate-methylmethacrylate grafted rubber PC, PPE, ASA, SAN, PS, phosphate esters, PTFE and SEBS PC, modified SEBS and hydroxyethyl acrylate terminated ε-caprolactone

Gergen and Davison, 1977 Bussink et al., 1978 Dieck and Wambach, 1980 Gallucci and Bookbinder, 1989 Lee, 1983 Liu, 1982, 1984 Idemitsu Kosan, 1983 Sumitomo Chemical, 1982, 1983 Daicel Chem., 1984 Laughner et al., 1992 Niessner et al., 1993 Wilkey, 1994

Introduction to Polymer Blends

1.4.4.1 PVC/NBR Blends The PVC/NBR blends were commercialized in 1936 by Bergisch-Gladbach. Nearly identical alloys, Geon™ Polyblends, were introduced by B.F. Goodrich 1947. To ascertain adequate interaction between PVC and NBR the AN content in NBR should be at least 25 wt% [Matsuo et al., 1969]. Most commercial blends contain 50-90 wt% NBR that acts as a solid plasticizer and processing aid. PVC blends with crosslinked NBR have been foamed since the 1940’s, initially for the production of buoyancy vests, shock absorption and insulation [McCracken, 1984], later for shoe soles. More recently, acidification of NBR made it possible to incorporate the NBR/PVC blends into PA, PC, or PEST [Iwanaga et al., 1990]. It was also found that NBR provides good compatibilization and toughening in blends of PVC with carbon monoxide-ethylene-vinyl-acetate copolymer (COPO-VAc) [Lund and Agren, 1993]. There are several commercial PVC/NBR blends, viz., Geon™, JSR NV, Krynac™ NV, Nipol™, Paracril™ OZO, or Vynite™. 1.4.4.2 PVC/Acrylics Blends The most common acrylic, PMMA, shows limited miscibility with chlorinated hydrocarbons (e.g., PVDC, PVC, CPVC, or CPE). The miscibility depends on the type of chlorinated polymer, tacticity of PMMA, and molecular weights of the two polymers. The origin of miscibility is the interaction between the -CHCl- group of PVC and the carboxyl group of the acrylic [Jager et al., 1983]. Acrylics have been incorporated into PVC blends quite early [Small and Small, 1951]. In spite of PVC miscibility with PMMA, blends of these two polymers are not commercially important. To be useful the blends must be toughened, preferably by acrylic elastomers. Some of the toughening agents are listed in Table 1.27. Examples of commercial PVC/acrylics blends are: Acrylivin™, Decoloy™, Enplex™, Hostalit™, Kane-ace™, Kydene ™, Metabulen™ or Vinidur™. The blends have been used for industrial, commercial, and consumer goods, in medical, electrical, and chemical engineering equipment, for food

41

or beverage, as aircraft or mass transit interior components, for power tool housings, etc. 1.4.4.3 PVC/Elastomer Blends These blends, usually with 30-60 wt% PVC, are represented by: Carloy™, OxyBlend™, or Vynaprene™. They have been formulated for extrusion, calendering, injection, or blow molding, e.g., into bottles, sheets for exterior signs, window accessories, cables and hoses, printing plates and rollers, shoe soles, profiles, military coax jacketing, etc. PVC blends with ABS and modified-ABS were already discussed. In Table 1.28 an abbreviated list of PVC blends comprising non-acrylic elastomer(s) is provided. 1.4.4.4 PVC/Polyolefin Blends PVC is antagonistically immiscible with PO. Thus, the standard strategies are applicable: (i) addition of a small amount of PO to improve processing and impact strength, (ii) co-continuous morphology, (iii) incorporation of PO as part of a copolymer comprising miscible with PVC segments, and (iv) compatibilized blends. Owing to difficulties in compatibilization, the PVC/PO blends are not yet commercial [Liang et al., 1999]. Evolution of these systems is outlined in Table 1.29. 1.4.4.5 PVC/CPE and PVC/CSR Blends PVC blends with CPE were patented and commercialized in 1956 as Hostalit™. Blends with CSR soon followed. By the mid-1970 the emphasis shifted toward blends with acrylic elastomers. Ternary alloys were developed, viz. of PVC with CPE, and poly(methyl methacrylate-co-butyl acrylate) (MMBA) [Maruyama et al., 1977]; or PVC, CPVC, and either MABS or a mixture of PMMA with imidized-PMMA or imidized-SMA [Soby et al., 1994]. These blends have been used for outdoor applications, flame retardant wall coverings and automobile interiors. Injection molded components include gullies in sewage systems, caps for road reflector posts and bench slats, etc. Evolution of these blends is traced in Table 1.30.

42

L. A. Utracki

Table 1.27. Toughened PVC/Acrylic blends Modifier

Reference

Butadiene-MMA-styrene copolymer (MBS, Acryloid™) PMMA and poly(butadiene-g-MMA) Ethylene-ethylacrylate copolymer MBS with controlled size of the elastomeric particles; transparent Copolymer of vinylchloride, alkyl acrylate, and vinylidene chloride Butadiene-styrene-methylacrylate-ethylacrylate Core-shell: crosslinked ABS with grafted onto it PMMA shell PB-grafted with MMA, styrene and vinyl acetate Poly(butadiene-co-butyl acrylate-co-styrene) Core-shell: poly(AN-co-MMA), or poly(AN-co-ethylhexyl acrylate-co-MMA) Poly(styrene-co-AN-co-MMA-g-butyl acrylate-g-MMA) 5-20 phr of MMA-AN-butadiene-styrene (MABS) with 10-40 wt% AN and/or styrene, 50-80 wt% 1,3-butadiene, and 25-75 wt% MMA for foamed profiles, bottles, pipes, boards, moldings, etc. Core-shell: poly(butadiene-co-styrene-divinylbenzene-co-butylacrylate-co-MMA) ” 20 wt% of either poly(vinylchloride-co-vinyl acetate) or EVAc — the blends were though and easy to foam. CPE and poly(MMA-co-butyl acrylate) MMA and styrene grafted onto an acrylic elastomer AN-b-MMA block copolymer PMMA with dehydrochlorinated PVC were found miscible and easy to foam for the cryogenic insulation in space vehicles Copolymer of ethylene, 1 - 60 % acrylic ester and 1-30 % CO2 or SO2

Fujii and Ohtsuka, 1954 Jarrett and Williams, 1960 Van Cleve and Mullins, 1962 Saito, 1975 Hoshi and Kaneko, 1962-63, 1965 Ichinoe, 1967 Michel, 1969 Kakefuda and Ito, 1971 Ide and Deguchi, 1971 Tanaka et al., 1971, 1972 Ide et al., 1972 Parks, 1976

Core: crosslinked silicone rubber; inner shell: crosslinked acrylate elastomer; outer shell: and styrene-AN copolymers

Usami and Ochiai, 1976 Barth et al., 1976; Goswami, 1977 Maruyama et al., 1977 Kishida et al., 1977 Iwata et al., 1979 Jayabalan, 1982; Jayabalan and Balakrishnan, 1985 Rys-Sikora, 1983, 1984 Lindner et al., 1990

Table 1.28. PVC toughening by non-acrylic elastomers Modifier

Reference

PIB, NR, IR, or CR Thio-rubbers (TM) CR and NBR Chlorinated polybutadiene (CPB) Polyisobutylene (PIB) Di-butyl fumarate and butadiene copolymer TM and CSR BR and/or poly(ethylene-co- vinyl or acrylic monomer), e.g., EVAc EPDM & polynorbornene having carboxylic, and carboxylic ester groups PVC blends with crosslinked NBR for foamed floating devices DOIP plasticized PVC blended with TPU and EVAc

B. F. Goodrich, 1941 Rittershausen, 1949 Signer & Beal, 1953 Esso, 1960 Lonza Ltd., 1964 Koenig et al., 1964 Allied Chemicals, 1965, 1966 Kasuya et al., 1969 Mitsubishi Chem., 1983 McCracken, 1984 Shin et al., 1998

43

Introduction to Polymer Blends

Table 1.29. PVC blends with PO Composition

Reason for Blending

Reference

PO with PVC PVC with either PP, PE, PS, or SBR and ethylene-vinylchloride PVC with either PE or PP and MBS

for extrusion or milling HDT, flame resistance, impact and tensile strength

Rosenfelder & Rosen, 1962 Montgomery, 1966

impact strength formulations for pipes or electrical insulation impact performance layered, wood-like materials low-temperature impact resistance flame retardancy compatibilized blends mechanical and impact performance

Baer and Hankey, 1967

PVC with either PE or PP and ABSM PP with PVC and either PMMA or PC PP and either EVAc-VC, EVAc or HDPE PVC was copolymerized with PP PVC, HDPE and CPE PVC/PO + poly(ethylene-co- alkyl, aryl, alkaryl or methyl methacrylate ester)

Himei et al., 1967 Yahata et al., 1971 Kojima & Tanahashi, 1972 Unitika Ltd., 1984 Nippon Zeon, 1984 Williams and Ilenda, 1993

Table 1.30. PVC blends with either CPE or CSR Composition

Reference

PVC with CPE PVC with CSR PVC with CPB Latex blending: PVC with CR PVC with PO, compatibilized by either CPE or MBS Solution-blended PVC with CPE or CSR PVC with CPE and a diamine (an interfacial agent) PVC with equal amount of CSR and SAN PVC with equal amount of CPE or CSR PVC with EVAc and CPE PVC with chlorinated-SBR (C-SBR) PVC with chlorinated EPDM (C-EPDM) powder blending: PVC and PE, then chlorinating and milling PVC with two heterogeneously chlorinated LDPE’s PVC with CPVC and ABS PVC with PS and C-SBR PVC/CPE with SMA PVC with HDPE and C-SBR PVC/CPE with glycidyl p-tert-butyl-benzoate PVC/CPE with either SMMA or PS-VAc

Frey, 1958 Matsuda, 1960 Esso, 1960 Nyori et al., 1962 Baer and Hankey, 1963 Beer, 1963 Hankey and Kianpour, 1964 Salyer and Holladay, 1964 Hedberg and Magner, 1965 Dow, 1965 Takkosha Co., 1967 Watanabe et al., 1967 Kato et al., 1967 Willott, 1968 Kojima et al., 1970 Falk, 1981 Bourland, 1983 Nippon Zeon, 1984 Sugiara and Takayama, 1988 Liou and Sun, 1993

44

L. A. Utracki

1.4.4.6 PVC/TPU (mainly polyester-type) Blends Initially, two-component, PVC/TPU blends were proposed [B. F. Goodrich Co., 1960], but soon, PVC/TPU blends with a modifier, e.g., ABS [Waugaman et al., 1963]; NBR or PA [Képes, 1959] were disclosed. Blending was also carried out by mixing PVC with polyols and isocyanates then polymerizing these two [Dainichiseika Color & Chemicals, 1983]. Commercial PVC/TPU (with NBR) are represented by Duralex™. The materials are usually formulated for extrusion, e.g., for wire and cable insulation, hoses and packaging. More recently, foamable, recyclable PVC blends were disclosed. They comprise PVC plasticized with DOP and/or epoxidized soybean oil, blended with either TPU and/or EVAc. The formulation could be foamed either during extrusion or injection molding. The material is used to produce anti-slip shoe soles with good abrasion resistance [Shin et al., 1998]. 1.4.4.7 PVC/EVAc and PVC/EVAc-VC Blends The miscibility of PVC with EVAc depends on the VAc content. Blends of PVC with PVAc were patented in 1938. PVC, or poly(vinylchlorideco-acetate) (PVCAc) was also blended with polyvinylacetal (PVA) [Lonza El. & Chem., 1948]. In later patents PVC instead of being mechanically blended with PVAc was copolymerized with vinyl acetate. The copolymer still required toughening, thus it was emulsion polymerized in the presence of styrene-butadiene-vinylacetate latex [Farbwerke Hoechst, 1970]. Latex blending (followed by spray drying) was a simple and efficient mixing method [Hammer, 1971]. Similarly, PVC and/or PVCAc was blended with a variety of butadiene-butyl acrylate-styrene copolymers [Ide and Deguchi, 1971]. PVC blends with ethylene-vinyl acetate-carbon monoxide copolymer (EVAc-CO) and a methylmethacrylate graft copolymer, Kane-Ace™, are also commercially interesting [Mitsui Petrochemicals, 1983].

Commercial PVC blends with either EVAc, or PVC-VAc have been offered for outdoor applications since the 1970’s as high impact strength, rigid formulations (e.g., Denkovinyl™, Hostalit™, Vinidur™, Solvic™, or Trosiplast™). The resins show good hardness, rigidity, adequate heat, chemical, and flame resistance. 1.4.4.8 PVC Blends with COPO The first PVC/COPO blends were developed in 1960 [Mullins, 1964]. It was reported that PVC melt viscosity decreased by addition of COPO [Hammer, 1973]. Later, the compositions were modified — PVC was blended with ethylene-carbon monoxide-vinyl acetate copolymer (COPO-VAc) and BMMM [Reardon, 1982]. 1.4.4.9 PVC Blends with Engineering Resins Owing to poor thermal stability of PVC, the high temperature blending must be avoided. Thus only few PVC/engineering resin blends are known. These are summarized in Table 1.31. The first commercial blend of this type is Cylon™. Here PVC is the matrix, and PA (that melts below 215°C!) the dispersed phase. The two resins were compatibilized using the well known PVC plasticizer — Elvaloy™ (a terpolymer of ethylene, carbon monoxide and acrylics). These soft to semi-rigid alloys were commercialized for wire coating, automotive applications and blow molding [Grande, 1997; Hofmann, 1998]. They are flame, abrasion and chemicals resistant, easy to process, and tough. 1.4.5

Polyvinylidenechloride (PVDC)

Polyvinylidenechloride, PVDC, was discovered in 1838 by Regnault but commercialized 100 years later as Saran™. The commercial PVDC is modified by the incorporation of either 15-20 wt% vinyl chloride, or 13 wt% vinyl chloride and 2 wt% acrylonitrile. PVDC blends are summarized in Table 1.32.

45

Introduction to Polymer Blends

Table 1.31. PVC blends with engineering resins Composition

Reference

PVC with 5-25 wt% polypropylene-terephthalate (PPT) PVC + copolymer of PET with polybutyleneglycol, PBG and 1,4-butanediol PVC with PC and ABS Low friction coefficient blends of PVC with POM Styrene-grafted PC with neat PC, PS and other styrenics, acrylics or PVC PVC with polyethylene carbonate High HDT blends: PVC, MBS and polyimide (PI) PVC with poly(butanediol-terephthalate-adipate) and 30 wt% GF Vinylchloride polymerized in the presence of PI, blended with PVC and MBA PVC with poly(methylmethacrylate-co-maleimide-co-vinyl cyanide) and styrene-cyclohexyl-maleimide-grafted butadiene PVC with imidated polymethacrylate (polyglutarimide, PGI) Plasticized PVC with an aliphatic polyester-b-aromatic polyester Miscible blend: PVC/PC and a bishydroxyphenyl-hexafluoropropane (6F-PC)

Hurwitz and DeWitt, 1970 Crawford and Witsiepe, 1972 Hardt et al., 1975 Doerffurt and Waeteraere, 1977 Kakizaki et al., 1979 Dixon and Ford, 1979 Kopchik, 1981 Yang, 1987 Clikeman et al., 1987 Ito et al., 1990 Fromuth et al., 1992 Jean and Devauchelle, 1993 Drzewinski, 1993, 1994

Table 1.32. Examples of PVDC blends Composition: PVDC with

Reference

NBR and CR Polyurethanes Ethylene-carbon monoxide-vinyl chloride copolymer; Alcryn™ blends PO and ethylene-methylacrylate compatibilizing ionomer PA-6, PA-1212 or PARA, and poly(ethylene-co-alkyl (meth)acrylate-co-vinyl acetate-co-CO-co-maleic anhydride) PVDC-VC with vinylidene chloride-methyl acrylate copolymer (PVDC-MeA)

Signer and Beal, 1951 McCready, 1976 Loomis and Statz, 1984, 1986 Burgert, 1987 Hofmann, 1994

1.4.6

Paleari and Fornasiero, 1994

Polyvinylidenefluoride (PVDF)

PVDF was patented in 1948 and commercialized by Pennsalt in 1962. It is a semi-crystalline polymer with Tg = -56 to -35°C, 50 % crystallinity, and Tm = 160-180°C. PVDF has been blended mainly with PMMA [Lin and Burks, 1993]. The blends are suitable for the use as stable electrets or weather-resistant architectural coatings. Nearly 25 wt% of PVDF consumption is in

weather resistant architectural spray finishing or coating to metals, roofing, curtain walls, wall panels, window frames, doors, hand rails, fascias, awings, louvers, and canopies. PMMA/PVDF blends are commercially available, e.g., Polycast™ from Royalite. PVDF blends are summarized in Table 1.33.

46

1.4.7

L. A. Utracki

Acrylic Blends

Polymethylacrylate (PMA) and polymethylmethacrylate (PMMA) were discovered, respectively, in 1880 and 1930. The resins have been used for the production of transparent plastic sheets, viz. Plexiglas™ or Perspex™, used for the military aircraft cockpit canopies, gunner’s turrets and the like [Riddle, 1954]. Acrylic elastomers (ACM or ANM) were developed by Röhm in 1901, and commercialized in 1948 as Hycar™ vulcanizable copolymers of ethyl acrylate, allyl maleated lactones, chloroethyl vinyl ether, butadiene, isoprene, acrylonitrile, etc. [Mast et al., 1944]. Since the 1950’s, a wide variety of acrylic compatibilizers and impact modifiers have been developed.

1.4.7.1 Co-poly(meth)acrylates (MBA & MBS) In the 1950’s, the core-shell, emulsion type methylmethacrylate-butadiene-styrene terpolymer (MBS) was developed to toughen PVC or PC. These blends could also contain other polymers, viz. SAA [Murdock et al., 1960]; SMM and PS [Murdock et al., 1962]; SMM-AN [Schmitt et al., 1967]; high heat ABS [Kanegafuchi Chemical Industry, 1967]; HIPS [Ward, 1970]; MMVAc-AA [Holland et al., 1970]; SMMA [Blasius, 1992], etc. Table 1.34 traces evolution of these systems. Later, these multipolymers were modified by incorporation MA, AA, or GMA units to serve as reactive compatibilizers and toughening agents for PA, PEST or PC blends.

Table 1.33. Examples of PVDF blends Composition: PVDF with

Reference

PCTFE for wire coating 30 wt% PMMA for outdoor films with good clarity, chemical & UV stability PMMA and polyethylacrylate Solution-blended with PA-610 to lower PA’s water absorption PC & acrylic copolymer for clear, yellowish films with single Tg = 120°C PPE/SEBS and MMAS for weatherability, chemical and solvent resistance Compatibilized PA for impact resistance and gas barrier properties COPO, PVP, PSF, polyester rubbers or poly-2-oxazoline POM for resistance to frictional wear, heat and UV stability

Kaufman, 1963 Koblitz, 1966 Schmitt and Miller, 1970 Saito, 1975 Leibler and Ringenberg, 1986 Van der Meer et al., 1989 Hizasumi et al., 1989 Gergen and Lutz, 1989 Shibata et al., 1992

Table 1.34. Acrylic compatibilizers-cum-impact modifiers, MBA & MBS Composition

Reference

Methylmethacrylate-butadiene-styrene terpolymer (MBS) Methyl methacrylate-butadiene-styrene-α-methylstyrene, for weather resistance Butadiene-styrene-methylacrylate-ethylacrylate (ASA) Partially crosslinked ABS core and PMMA shell (a MABS) Copolymers of styrenics (e.g., PS, SAN, SMMA, etc.) with, e.g., 0.1 wt% of hydroxyethylacrylate

Fujii and Ohtsuka, 1954 Ruffing et al., 1964 Ichinoe, 1967 Michel, 1969 Rubens, 1986

47

Introduction to Polymer Blends

1.4.7.2 Impact Modification of PMMA PMMA, like PS, is brittle and requires toughening. These efforts are summarized in Table 1.35. 1.4.7.3 PO Blends with Acrylic Polymers PMMA is antagonistically immiscible with a polyolefins — blends of this type have been used in non-critical applications, viz. PP/PMMA blends with EVAc were used as plastic-paper [Yamamoto et al., 1971], while those with PVC (or CA) as wood-like materials [Yahata et al., 1971]. For more demanding applications, either PO should be blended with an acrylic copolymer comprising a PO block, or PO should be grafted with acrylic moieties. Examples of the PO/Acrylics blends are listed in Table 1.36. Blends of a PO (PE, PP, PB, P4MP, their blends and copolymers, e.g., with 1-alkenes, vinyl esters, vinyl chloride, methacrylic esters, and methacrylic acid) with 0.2-50 wt% of a graft

copolymer showed high tensile modulus and high sag resistance without increased melt viscosity. The blends could be shaped into foamed profiles at T = 200 to 230°C. To prepare the graft copolymer, a PO (MW = 50 to 1,000 kg/mol) was either dissolved or swollen in an inert hydrocarbon, monomers (•80 wt% of a methacrylic ester, CH2=C(CH3)COOR) and an initiator was added to the heated mixture while stirring. As a result, acrylic branches of a relatively high molecular weight (MW = 20 to 200 kg/mol) were grafted onto the PO macromolecules. The graft copolymer could be used as a compatibilizer-cum-impact modifier in a variety of polymers selected from between: PO, acrylic polymers, SAN, EVAc, PA, PEST, PC, POM, PAr, PVC, ABS, PVDC, cellulosics, polyesterpolyether block copolymers, PEA, PEEK, PEI, PES, CPVC, PVDF, PPE, PPS, PSF, TPU, PAI, PCL, polyglutarimide, blends of PEST with PC or PVC [Ilendra et al., 1992, 1993].

Table 1.35. Impact modification of PMMA PMMA Impact Modifier

Reference

PVAc Copolymers of methacrylonitrile, ethylacrylate and/or α-methylstyrene Ethylene-vinyl acetate-vinyl chloride copolymer (EVAc-VC) SMA and methylmethacrylate-methylacrylate copolymer (MMMA) MBA: copolymer of butadiene, butylacrylate and methylmethacrylate Styrene-acrylonitrile copolymer (SAN) SAN, PS, and poly(methyl norbornene-2-carboxylate) IPN: crosslinked PBA, crosslinked & uncrosslinked SAN Poly(p-hydroxy styrene), PVPh and EVAl Acrylic core-shell copolymer and either PBT or PET Poly(allyl methacrylate-co-butyl acrylate-co-butanediol dimethacrylate-co-styrene-co-methyl methacrylate) or poly(acrylonitrile-co-butyl acrylate-co-tricyclodecenyl acrylate-co-styrene) Poly(acrylate-N-cyclohexyl maleimide), PMI, and a copolymer: PMMA — core, crosslinked butyl acrylate-styrene copolymer — middle layer, and PMMA shell; d = 200-300 nm PEG/atactic PMMA blends were characterized by PVT at T = 20-200°C and P = 0-200 MPa. Free volume fraction was calculated from an equation of state

E. I. DuPont de Nemours, 1942 Coover and Wooten, 1962 Kishikawa et al., 1971 Bronstert et al., 1971 Kotama, 1972 Sugimura et al., 1972 Ikeda et al., 1976 Silberberg, 1982 La Fleur et al., 1992, 1994 Bright et al., 1993 Farwerck et al., 1993

Shen, 1994

Schmidt and Maurer, 1998

48

L. A. Utracki

Table 1.36. Polyolefin/acrylic blends PMMA Impact Modifier

Reference

40 - 90 wt% of ethylene-co-acrylic or methacrylic acid with ethylene-co-vinyl acetate or ethyl acrylate for foam production Ionomer with (1) a terpolymer of ethylene, vinylacetate and CO or SO2,

Park, 1978, 1980

and (2) an elastomer (e.g., NR, IR, PU) PE with EVAc, CPE, BR, etc., have been chemically foamed at T = 150-210°C 5 to 95 wt% of LDPE or LLDPE with EAA PE with either poly(ethylene-co-vinylcarboxylate) or an acrylate PP with an ionomer and EBA-GMA PP with an ionomer, EBA-GMA and EPDM PO with a core-shell graft copolymer MBS-type PP with acidified-PP, or a carboxylic acid-modified EPR, SMM-MA, and either EMMA-GMA or EVAc-GMA LLDPE, PMMA and SEBS, EPR, or ethylene-styrene block copolymer (ES) At least two elastomers and an ethylene-methacrylate-acrylic acid ionomer PE with alkyl acrylate or alkyl methacrylate

Enderle, 1984 Kuhnel and Spielau, 1981 Park, 1985, 1986d Broadhed, 1987 Saltman, 1988, 1989, 1992 Dawson, 1993 Aoyama et al., 1993, 1994 Abe et al., 1994 Dobreski and Donaldson, 1994 Arjunan, 1994, 1995 Godfrey, 1995

1.4.7.4 PC Blends with Acrylic Polymers

1.4.7.5 PEST Blends with Acrylic Polymers

PMMA has been blended with PC since 1971. Two types of PMMA/PC systems are of interest: (i) impact modified alloys, and (ii) miscible blends. To the first category belong Meta-marble™ blends of PMMA/PC with ABS [Ikura et al., 1974] or with ASA [Giles and Sasserath, 1986]. Blends of PC with two acrylic copolymers showed good processability, notched impact strength and HDT [Eckel et al., 1993]. Acrylic polymers are recognized for their miscibility with a variety of polymers, viz. miscibility of PMA with PVAc [Kern, 1957]. PMMA is miscible with standard PC at T < LCST = 140°C. The miscibility range can be greatly increased by modifying the PC chain ends (LCST ” 300°C) [Kambour, 1988]. Demixing PMMA/PC blends by the spinodal decomposition mechanism generated alloys with excellent mechanical properties [Kyu, 1990]. PMMA is also miscible with fluorinated-PC [Drzewinski, 1993, 1994].

Blends of PEST with acrylic polymers are limited to systems with acrylic elastomers. Examples are listed in Table 1.37. PBT and PET were reported to form miscible blends with either a poly-p-methoxyphenyl methacrylate or poly(phenyl methacrylate) [Siol et al., 1993, 1994]. 1.4.7.6 PPE Blends with Acrylic Polymers Poly(2,6-dimethyl-p-phenylene ether) (PPE) was rarely blended with acrylics, viz. with styrene-methylmethacrylate-co-cis-polyisoprene [Abolins and Reinhardt, 1976]; PMMA [Izawa and Nakanishi, 1973; Matsunaga et al., 1974]. 1.4.7.7 PA Blends with Acrylic Polymers Polyamides, PA, can be impact-modified by addition of acrylic multipolymers, e.g., methyl methacrylate-co-methacrylic acid-co-ethylacrylate [Halliwell, 1965, 1966]; ethylene-2-hydroxy ethyl methacrylate-methyl methacrylate [Hepworth et

49

Introduction to Polymer Blends

al., 1970]; or ethylene-ethylacrylate-acrylic acid ionomer [Meyer and Tacke, 1978].

with EGMA [Takahashi and Kobayashi, 1992], EGMA/AS, EGMA/PMMA, or their mixture [Kobayashi and Shinohara, 1993].

1.4.7.8 POM Blends with Acrylic Polymers 1.4.8 These systems are not of industrial importance. However, addition of an acrylic was reported to improve processability, abrasion resistance, and weatherability of POM. For example, to improve weatherability POM was blended with polythioisocyanate, TPU, PMMA, and benzotriazole [Endo et al., 1990]. POM/TPU with EMMA and benzotriazole shown enhanced performance [Okuda, 1990]. POM was also blended

Polyethylenes (PE)

1.4.8.1 Homopolymers Properties of PE depend on: molecular weight (MW), molecular weight distribution (MWD), as well as on the degree and type of branching. The density and modulus of PE’s increase with crystallinity. As shown in Table 1.38, seven principal categories of PE are recognized.

Table 1.37. Polyester/acrylic blends 1. Acrylic Impact Modifiers for Pest Ethylene-methylmethacrylate copolymer (EMMA) Graft copolymer: acrylonitrile-butadiene-styrene-methylmethacrylate, ABSM & PDMS Ethylene-hydroxyethyl methacrylate (EHEMA) Ethylene-vinyl acetate (EVAc) Methyl methacrylate-methyl acrylate copolymer (MMMA) Ethylene-vinyl acetate-methacrylic acid copolymer

Reference Dijkstra and Jones, 1969 Sauers and Barth, 1970 Jones et al., 1971 Jones et al., 1971 Kamata et al., 1974 Gander et al., 1977

2. Blends of PMMA with: 1,4-butanediol terephthalate-co-polybutylene glycol (PBT-PBG) PET and PS PBT and CH3NH2 (to convert PMMA into polyglutarimide)

Charles & Gasman, 1979 Kamata et al., 1980 Toray Ind., 1984

Table 1.38. Polyethylenes No Type

Code

1 2 3 4 5 6 7

UHMWPE HDPE MDPE LDPE LLDPE VLDPE ULDPE

Ultra-High-MW High density Medium density Low density Linear low density Very low density Ultra low density

Density (kg/m3) ρ = 969 941-969 926-940 910-925 910-925 900-910 855-900

Characteristics MW > 3,000 kg/mol High MW and crystallinity Long chain branching, Tm = 115°C Ziegler Natta type with short branching, Tm = 120-135°C Tm = 40-85°C

50

L. A. Utracki

The first polymethylene was obtained in 1897 by the thermal decomposition of diazomethane. In 1931, about half a gram of PE was obtained in a free radically polymerization at high T and P. In 1937, Telcothene™, a blend of PE and polyisobutylene (PIB) was produced for submarine cables, and in 1939, the first LDPE, Alketh™, plant with 100 tons/year capacity went into operation [Kennedy, 1986]. In 1951, HDPE was polymerized using the Z-N catalyst [Zletz, 1954]. In 1957, DuPont Canada developed LLDPE, Sclairtech™ [Lank and Williams, 1982]. In the 1980’s, new catalysts made it possible to polymerize VLDPE and ULDPE, commercialized in 1986. The newest PE’s (Tafmer™ was introduced by Mitsui Petrochemicals in 1975) are prepared using the single-site metallocene catalysts. These new resins have controlled MW, MWD, comonomer placement, and density. Theirs melting point, Tm = 70-120°C, increases linearly with density, ρ = 880-930 kg/m3. In 1991, the world production of LDPE, HDPE, and LLDPE was respectively 16.3, 15.4, and 7.8 Mton/year, amounting to 39 wt% of all thermoplastics.

1.4.8.2 PE Blends PE blends can be divided into three categories: (1) PE lots blended to meet standard specifications for density and melt flow, (2) PE modified with ”15 wt% of other polymer(s), and (3) PE bends with other thermoplastics or thermoplastic elastomers. PE’s are immiscible with nearly all polymers, thus the standard strategies are applicable: (i) noncompatibilized blends with low concentration of the dispersed phase, e.g., blends of either PP or PE with 2 wt% PVAl; (ii) non-compatibilized blends for the use in non-critical applications; (iii) non-compatibilized blends having co-continuous morphology, e.g., PE, blended with neoprene rubber at a ratio 1:1, then irradiated by electron beam; and (iv) compatibilized blends. 1.4.8.2.1 PE/Elastomer Blends Polyolefins have been modified by the incorporation of elastomers to improve low temperature impact strength and elongation. Table 1.39 provides examples of these systems.

Table 1.39. PE/Elastomer blends Modifier

Reason

Reference

Rubber Cyclo-rubber PIB CSR BR Poly-1-butene BR and EVAc An ionomer, with or without EVAc 0.1-5 wt% aPP EVAc and EVAl Polytransoctanamer (PTO) Poly(ethylene-co-vinylcarboxylate) ULDPE/CSR or CPE, dynamically vulcanized, then dispersed into fresh CSR or CPE Starch and at least one ionic compound

impact modification adhesion to metal transparent, impermeable, shrink-wrap films for films or coatings improve elongation processability & ESCR improved extensional behavior films with good tear, and yield strength blown or stretched packaging films transparency & impact strength impact modification high impact strength processability, hot-weld strength, adhesion and crack resistance for single-ply roofing membranes high frequency sealable packagings

Standard Oil, 1937 Child et al., 1942 Briggs et al., 1958 Boger and Thomas, 1958 Cole, 1959 Rudin & Schreiber, 1964 Ceresa et al., 1968 Willott, 1968 Nakamura et al., 1973 Pritchett, 1980, 1981 Kita & Hashimoto, 1987 Broadhed, 1987 Ainsworth, 1990, 1994 Dehennau et al., 1994

Introduction to Polymer Blends

1.4.8.2.2 PE/EPR or EPDM Blends The first patent on PE/EPR blends was deposited before commercialization of EPR [Corbellini, 1962]. Several similar inventions were disclosed, viz. HDPE blends with EPR [Crawford and Oakes, 1963]; PE with EPDM [Prillieux et al., 1962]; PE/EPDM blends with either PP or PB [Schreiber, 1966]; PE with EPR and ethylene-acetoxybicycloheptene copolymer [Shirayama and Iketa, 1971]; or VLDPE with EPR, EPDM or their mixtures [Nishio et al., 1992]. To improve PE/EPDM adhesion to polar materials, PE was first grafted with MA and then blended with EPDM [Honkanen et al., 1983]. The first reactor-type thermoplastic polyolefin (R-TPO) was LLDPE/PP [Yamazaki and Fujimaki, 1970, 1972]. The three-component R-TPO’s (PE with PP and EPR) soon followed [Strametz et al., 1975]. PE was also polymerized in the presence of active catalyst and an olefinic copolymer [Morita and Kashiwa, 1981]. Blending amorphous co-polyolefins with crystalline PO’s (HDPE, LLDPE, PP), and a filler resulted in moldable blends, characterized by excellent sets of properties [Davis and Valaitis, 1993, 1994]. Blends of polycycloolefin (PCO) with a block copolymer (both polymerized in metallocene catalyzed process) and PE, were reported to show outstanding properties, viz. strength, modulus, heat resistance and toughness [Epple and Brekner, 1994]. More recently, blends of a partially crosslinked thermoplastic elastomer with 5-40 parts of a PO (viz. LLDPE, PP, EPR, or PB-1) were developed for low density, foamable alloys [Okada et al., 1998a]. The density was reduced at least by a factor of two. In the following patent 1-17 wt% of a long-chain branched PP was also added [Okada et al., 1998b]. The extruded foam was free of surface roughness caused by defoaming, was soft to the touch and showed excellent heat and weathering resistance. 1.4.8.2.3 PE/PE Blends Molten polyethylenes of different type chain structure usually are immiscible (see Chapter 2). Upon crystallization the spherulites of one PE

51

(having higher Tm) are encapsulated by those of the other PE. Co-crystallization of two PE’s into a single type isomorphic cell is rare [Utracki, 1989a]. However, due to low interfacial tension coefficient, the phase coarsening is slow. Alloys of different PE’s constitute a large and important part of the PO technology. For example, in some countries 70 wt% of PE is sold after blending (e.g., LLDPE with LDPE). As the technology evolves, these blends are prepared from resins of widely different rheological character, giving the viscosity ratios λ ≥ 10,000. Usually, they do not require compatibilization, but owing to such a large value of the viscosity ratio, blending in shear flow is inefficient. Mixing in the extensional flow field is the potential solution [Luciani and Utracki, 1996; Utracki and Luciani, 1996]. Once cooled below the crystallization point of one component, the blend’s morphology is fixed by crystalline crosslinks. Blending of different grades and types of PE improves processability and mechanical performance. Blending, as it will be evident from the examples in Table 1.40, also may lead to transparency, improved abrasion resistance, stress crack resistance, etc. 1.4.8.2.4 PE/PP Blends PE has been used to improve the low temperature impact strength of PP (see Table 1.41). The blends are mostly immiscible, compatibilized either by addition of EPR, EPDM, by reactive blending, or by post-blending co-crosslinking, e.g., by electron beam or γ-radiation [Utracki and Dumoulin, 1995]. 1.4.8.2.5 PE/Other Commodity Polymer Blends PE’s are frequently used as impact modifiers for a variety of other thermoplastics. For example, addition of either PE, CPE, or CSR to PVC improved its moldability, stability, impact strength and chemical resistance [Matsuda et al., 1960]. Blends of PO/PVAl were developed to improve the antistatic properties [Minekawa et al., 1969]. LDPE was blended with poly(2-ethyl-2-oxazoline) (PEOX) for improved adhesion, e.g., to PET [Hoenig et al., 1984]. Blends of PE, PP, PS or their copolymers with ethylene-fluorinated vinyl ether copolymer

52

L. A. Utracki

Table 1.40. PE/PE blends Blend

Advantage

Reference

LDPE with LLDPE

processability, stiffness, abrasion resistance, H2O vapor permeability soft, thin films processability, impact strength, mechanical performance improved processability

Wissbrun et al., 1962; Golike, 1962 Sakane et al., 1979 Larsen, 1982

LDPE, HDPE and EPDM or aPP Two types of LLDPE LLDPE with: LDPE, PP, TPO’s, rubbers, EVAc, PP-MA, EPR, ... HDPE with LLDPE LDPE with HDPE, PP and EP-block LLDPE, LDPE and PP or EPR HDPE with either LLDPE or LDPE VLDPE and LLDPE Metallocene LLDPE and ionomer High and low molecular weight PE Reactor blends of LLDPE 100 parts of LDPE (ρ = 890 to 925 kg/m3), 1 to 110 parts of a HDPE LLDPE with EVAc (10-20 % VAc) LDPE with LLDPE 70-98 wt % of LDPE (with ≤ 60 wt% LLDPE) and 2-30 wt% HDPE Bimodal PE’s (LCB = 0.01- 3; Mn1/Mn2 > 7) was lightly crosslinked. The PE-1 was prepared in the first reactor, and PE-2 was prepared in the presence of 15 - 65 wt% of PE-1 in the second reactor. The reaction could be carried out in the slurry, solution, or gas phase.

improved strength and transparency modulus, strength, no sagging high stiffness, and film clarity high stress crack resistance processability heat shrinkable films processability and physical properties improved MD/TD tear balance expandable compositions for a small diameter electric wire insulation processability, tear strength, transparency improved tear strength and haze for physical foaming of recycled HDPE blends (ρ ≤ 885 kg/m3) were used for: wire/cable coating, weather-stripping, seals, foamed articles with closed, open or mixed cells, containers, medical appliances, drapes and coverings, fibers, tapes, tubes, pipes and hoses, bellows, boots, gaiters, footwear, etc.

were used for the electrical insulation of highvoltage, submarine cables [Barraud et al., 1993]. Blends of LLDPE with EVAc or EEA have comparable physical properties and cost to plasticized PVC [Rifi, 1994]. 1.4.8.2.6 PE/PA Blends The reasons for blending PE with PA are (1) a desire to improve the impact strength and moisture absorption of PA, and (2) to improve rigidity and barrier properties (to oxygen and solvents) of PE. Films and containers manufactured from the latter blends show overlapping lamellar structures, that causes high tortuosity for molecular diffusion, and significant reduction of oxygen or solvent (e.g., gasoline) permeability. The technology became particularly attractive after the reactive grafting

Haas & Raviola, 1982; Cowan, 1983; Fukui et al., 1983 Showa Denko, 1983 Shin-Kobe El. Mach., 1984 Bahl et al., 1985 Boehm et al., 1992 Godbey and Martin, 1993 Babrowicz et al., 1994 Coutant, 1994 Ali et al., 1994 Sakamoto et al., 1994 Benham and McDaniel, 1994 Benham et al., 1995 Lee, 1995, 1996 Cree et al., 1998

of PO with either maleic anhydride, acrylic acid or glycidyl methacrylate was invented [Steinkamp and Grail, 1976]. These modified PO’s could be directly used in blends with either PA or PEST [Davis, 1975]. In Table 1.42 examples of PE/PA blends are given. 1.4.8.2.7 PE/PC Blends To increase rigidity of PE, the resin has been blended with about 5 wt% of a high-modulus polymer, e.g., PC [Peters and Schuelde, 1963]. PC also stabilized PO against the thermal degradation [Schutze et al., 1972]. Addition of 3-5 wt% PO toughened PC [Yamada, 1963]. For good dispersion, the blending should be carried out at T > 290°C, using PE grades with the viscosity ratio: η(PE)/η(PC) ≤ 0.3-0.9. Several impact-modified

53

Introduction to Polymer Blends

Table 1.41. PE/PP blends Composition

Reason

Reference

PP/LLDPE PP/LLDPE PP/LDPE PE/PP compatibilized with EPR PE/PP compatibilized with PIB Isotactic PP with sPP PE/PP compatibilized with EPDM PE grafted with methacrylic acid & PP with dimethylaminoethylmethacrylate PE/PP compatibilized with EPR Isotactic PP with aPP PE/PP compatibilized with EP-block copolymers Reactor-blends: PE with PP and EPR PP/EPR and 5-30 wt% of hexene- or octene-type LLDPE PP/HDPE = 1:1

mechanical properties at low temperature impact resistance, and low-T brittleness impact strength and low brittle temperature low-T brittleness and Izod impact strength low temperature impact strength low temperature impact strength improved impact properties of PE’s blended at a ratio 1:1 showed excellent mechanical properties high impact strength impact strength at low temperature mechanical, low temperature impact, and optical properties reactor- thermoplastic polyolefin, R-TPO improvement of mechanical properties

Holzer and Mehnert, 1963 Martinovitch & March, 1963 Sun Oil, 1964 Rayner et al., 1964 Lehane, 1964 Emrick, 1966 Schreiber, 1966 Langworth, 1967

PP with 5-20 wt% LDPE and EPR PP with EPR, then with PE Blending PP with EPR, then with PE PP, LDPE, HDPE and an EP-block copolymer PP, LLDPE, LDPE and/or EPR PP, LLDPE and a Plastomer™ (a metallocene ethylene-co-butene) EPR with Plastomer™ HDPE with PP autoclave-foamed with CO2

processability, weld-line strength, low-T impact strength transparency and mechanical performance co-continuous morphology — impact and mechanical properties co-continuous morphology, high performance films with good modulus, tear strength and sagging properties high modulus and clarity for melt-spun or melt-blown fibers or fabrics packaging films, tubes, and trays foam with • 109 cells/mL and cell diameter, d ” 10 µm. High impact strength.

Sumitomo Chem., 1968 Tanahashi and Kojima, 1970 Leugering & Schaum, 1970 Yamazaki & Fujimaki, 1970 Shirayama et al., 1972 Moorwessel et al., 1974 Oita et al., 1978 Huff, 1978 Huff, 1980 Shin-Kobe Electric Machinery, 1984 Bahl et al., 1985 Bartz et al., 1993a Mehta and Chen, 1994 Dorudiani et al., 1998

Table 1.42. PE/PA blends Composition

Reason

Reference

LDPE or PP with either PA-6 or PA-66 PE with a PE-PA-6 copolymer PE mixed with a lactam, then polymerized PA-6/LDPE or PIB/N-stearyl steramide PA-66, PE, PBT and PC with SEBS Aromatic polyamide (PARA) compounded with PE

for sheets, films, fibers or bottles transparency and low permeability low water absorption, strength low temperature impact; tenacity co-continuous morphology processability, elongation at break, tensile and impact strength for improved impact resistance processability, mechanical properties even after water immersion rigidity and low-T impact strength resistance to thermal degradation adhesion to fillers, excellent performance of filled compositions

Mesrobian & Ammondson, 1962 Craubner et al., 1962 Hill et al., 1970 Gilch and Michael, 1970 Gergen and Davison, 1978 Paschke et al., 1983

PA with carboxylated EPDM and PE PA blended with a mixture of PP, HDPE and EPR PA with HDPE, EPR and maleated-PP PARA with PE or PP-MA and hydrazine PO-g-GMA, acrylamide, vinylpyrrolidone, acrylic and/or methacrylic acid ester, then blended with either PA or PEST

Unitika Ltd., 1983 Hasuo et al., 1985 Kondo and Tominari, 1987 Yoshihara, 1990 Teraya et al., 1994

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grades of PC (viz., Lexan™ EM) comprise PE [Freitag et al., 1991]. However, as Table 1.43 illustrates, most PE/PC blends also contain a polymeric compatibilizer-cum-impact modifier. 1.4.8.2.8 PE/PPE Blends The PPE/PE blends are not commercial, but a small amount of PE is frequently added to PPE/ HIPS or PPE/SBR blends to improve process-

ability and solvent resistance (see Table 1.44). To enhanced rigidity of PO, 5-35 wt% of modified PPE was added. In this application PPE can be considered a low density filler. Similar effects can be obtained adding a small quantity of other polymers, viz. PC, POM, PPS, etc. 1.4.8.2.9 PE/PEST Blends Addition of PE to PEST is known to improve

Table 1.43. Compatibilized PE/PC blends Compatibilizer/Impact Modifier

Reference

0.01-2 wt% ABS Hydrogenated chlorosulfonated butadiene-styrene block copolymer Maleated-LLDPE Ethylene-glycidyl methacrylate copolymer (EGMA) Methyl-phenyl siloxane Acrylic and hydrocarbon elastomers (viz. BR, EPR, EPDM, IR, IIR) EPR or EPDM 1-5 wt% SEBS Acrylic rubber or maleated-PO 4 wt% poly(butylacrylate-co-methylmethacrylate) copolymer ABS with polysiloxanes containing Si-H bonds Ethylene-ethyl acrylate copolymer (EEA) and SEBS 0.2-15 wt% acrylic impact modifier (MBA, Acryloid™ KM) TPE (acrylic rubber, butyl rubber, EPDM, or SBS)

O’Connell, 1974 Bussink et al., 1978 Mitsubishi Chemical Ind., 1980 Sumitomo Chem., 1982, 1983 Rosenquist, 1982 Teijin Chem., 1982, 1983 D’Orazio et al., 1982 Idemitsu Kosan, 1983 Idemitsu Petrochem., 1983, 1984 O’Connell, 1983 Liu, 1983, 1984 Overton and Liu, 1984, 1985 Endo and Ishii, 1984 Kozakura et al., 1992, 1994

Table 1.44. PPE/styrenics blends with PE Composition

Reference

PPE with LDPE PPE with either SBR or HIPS and LDPE PPE was blended with SEBS and PE PPE with HIPS, SEBS and PE PPE with PS, SEBS and PE PE with 0-35 wt% PP, PDMS, and 5-35 wt% PPE, PC, PET or PA PPE with hydrogenated SB block copolymer and LDPE PPE with HIPS, SEBS and LLDPE PPE with styrenics and high molecular weight HDPE

General Electric, 1966 Summers et al., 1972 Yonemitsu et al., 1976 Haff and Lee 1978 Haaf, 1979 Plochocki et al., 1979 General Electric, 1984 Hambrecht et al., 1986 Bopp and Balfour, 1993

55

Introduction to Polymer Blends

impact strength, processability, solvent resistance and weatherability. When more than 5 wt% of PE is required, compatibilization is advisable. Examples of these systems are listed in Table 1.45. 1.4.8.2.10 PE/POM Blends POM is difficult to compatibilize, and without compatibilization only ≤10 wt% of POM in PE, or vice versa PE in POM, can be used. For example, addition of a small amount of PE to POM improves its processability, impact and abrasion resistance, hardness, surface finish, and rigidity, while addition of POM to PE improved its modulus and abrasion resistance — see Table 1.46. 1.4.8.2.11 PE/Specialty Resin Blends Most specialty resins are processed at temperatures

that limits the possibility of blending them with PE. The PE/specialty resin blends usually contain low concentration, ≤ 5 wt%, of either component. Addition of PE improves the processability, surface finish, chemical, solvent and impact resistance. Addition of specialty polymer to PE may improve rigidity and processability (viz. PE/LCP). Examples are given in Table 1.47. 1.4.9

Polypropylene (PP)

1.4.9.1 Homopolymers There are three types of polypropylene: amorphous (aPP), isotactic (PP), and syndiotactic (sPP). Performance of these resins depends on the tacticity content. PP was commercialized in 1957 by Hoechst.

Table 1.45. PEST/PE blends Composition

Reference

PET with 0.5-50 wt% PE; for impact strength PET with poly(ethylene-co-acrylic acid) ionomer PET with oxidized and carboxylated PE and glass fibers Poly(ethylene-2,6-naphthalene dicarboxylate) (PEN) with MDPE and/or PP PBT (or PET) with PC and LLDPE PBT (or PET) with PC and PB PET with LLDPE PET with either PE, PP, PO-GMA, vinyl- or acrylic-grafted PO, and GF PBT with PC, and PE, PP and/or EPR grafted with GMA or MA HDPE and copolymer of n-butylterephthalate with ethylene- and propylene-glycol PET with 30-70 wt% LDPE, HDPE or LDPE and EEA-GMA

Glanzstoff, 1967 Cope, 1969 Segal, 1973 Tokai and Sakai, 1973 Boutni and Liu, 1983 Dieck and Kostelnik, 1983 Smith and Wilson, 1984 Mukohyama, 1993 Fujie, 1993 Abu-Isa and Graham, 1993 Natarajan et al., 1993, 1994

Table 1.46. POM/PE blends Composition

Reference

POM was blended with 1-5 wt% of either PE, EVAc, PEG POM was blended with 90-99 wt% of LDPE PE or PP with EVAc and either POM, PMMA, PS or SMM POM was blended with EVAc and HDPE POM was first blended with TPU, then with either PA, LLDPE, PP, PBT or PET

Burg et al., 1972 Rudin and Schreiber, 1964b Yamamoto et al., 1971 Ishida and Masamoto, 1974 Flexman, 1992

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Table 1.47. Specialty polymer blends with PE or PP Composition

Reference

Polyarylene polyether sulfone (PAES) with ” 5 wt% of PE or PP Polysulfone (PSF) with either PE, PP, BR, EPR or EPDM Polyarylamideimide (PAI) with a small amount of PE Polysulfide (PPS) blends with PDMS and either PE or PP PPS/PE compatibilized by addition of an aromatic nitro-compound Polyoxycyanoarylene (POCA) with PO; compatibilized by EGMA PE blends with liquid crystal polyester (LCP)

Gowan, 1969 Hart, 1971 Toray Ind., 1981 Liang, 1987 Köhler et al., 1992 Hashimoto et al., 1990 Alder et al., 1993

The slurry process in hexane used the Ziegler-Natta (AlEt2Cl + TiAlCl6) catalyst [Sailors and Hogan, 1982]. The new metallocene catalysis leads to isomer purity in excess of 96%. It is also possible to produce branched, high melt strength PP, with extensional stress hardening, similar to that of LDPE [Phillips et al., 1992]. The new PP’s show the melting point, Tm = 120-164°C. To enhance the performance, PP is usually blended (in the reactor or outside the reactor) with much more viscous PP-copolymers. As a consequence, one of the most serious industrial problem is homogenization of these materials [Luciani and Utracki, 1996; Utracki and Luciani, 1996]. The advantage of sPP over PP is that impact strength and tensile modulus of sPP are significantly higher. While PP has a planar zigzag helical structure, the sPP has a three-dimensional one that leads to lower crystallinity and melting point: Tm(PP) § 165 vs. Tm(sPP) § 133°C. 1.4.9.2 PP Blends PP is brittle, especially at T ” Tg = 0°C. The resin fractures by the crazing-cracking mechanism [Friedrich, 1983]. The discovery of PP immediately followed by search for methods of improvement the low-T impact behavior. PP was blended with: EPR or EPDM [Hogan and Banks, 1953, 1955]; PE [Holzer and Mehnert, 1963]; sPP [Emrick, 1966]; aPP [Tanahashi and Kojima, 1970], etc.

1.4.9.2.1 PP/Other Polyolefin Blends PP blends with PE have already been discussed. PP blends with elastomers will be discussed in the following parts. In Table 1.48 few examples of PP blends with other PO’s are given. 1.4.9.2.2 PP/Elastomer Blends These blends constitute a large, commercially important group. Usually 5-20 wt% of elastomer has been used. Alloying improves processability (e.g., in blow molding) and impact strength at low temperature. Diverse elastomers have been used, e.g.: EPR, BP, PIB, BR, uncured PB and SBR; dynamically co-vulcanized CBR; BR, CSM, and EPDM [Reid and Conrad, 1962; Dow Chem., 1963; Gessler and Haslett, 1962; Esso R&E, 1962; Coran and Patel, 1978]. Blends with amorphous CSR showed good mechanical properties [Shikata et al., 1973]. Partially vulcanized blends of CSR with PP and/or ULDPE had good processability, hot-weld strength, interplay adhesion and crack resistance [Ainsworth, 1990, 1994]. Addition of CPE improved PP’s processability and properties [Newe et al., 1984]. Many EPR and EPDM elastomers show a block copolymer behavior. When blended with PP, they form emulsion-like dispersions. For the ease of compounding a small amount of PE may also be added. Furthermore, if the elastomeric phase is lightly crosslinked the morphology is more stable. The PP/EPR blends can be processed by all methods used for PP. They are characterized by good processability, dimensional stability, low

57

Introduction to Polymer Blends

Table 1.48. PP/PO blends Elastomer Added To PP

Comment

PP/PIB miscible (?) blends PP miscible with aPP or polybutene-1

low temperature impact properties reduction of Tg by up to 20°C

10-60 wt% PP or LLDPE with EPDM and ”95 wt% PIB or butyl rubber PP with polytransoctanamer (PTO) PP with polyoctadecene (POD) PP/PB and Plastomer™ PP/PB and a poly(1-butene-co-ethylene)

Soft, easy foamable blends (due to strain hardening) five-fold increased impact strength temperature sensitive transparency impact & mechanical properties processability, impact strength, and optical properties aPP was immiscible with PP and partially miscible with sPP to strengthen spherulites boundary

High MW aPP blended with high MW of either PP or sPP Addition of EPR to PP

shrinkage, high stiffness, tear strength and softening temperature, good mechanical properties (at T = -40 to 150°C), ozone resistance, fatigue and abrasion resistance (see Table 1.49). These materials have been used in more than 200 applications, in automotive industry, appliances, hardware and plumbing, medical, shoe industry, sports equipment, toys, etc. Examples of commercial PP/EPR blends are: Buna™, Dutral™, Epcar™, Epichlomer™, Epsin™ and Santoprene™, Esprene™, Ferrocline™, Gafply™, Intolan™, Kelburon™, Larflex™, Milastomer™, Nordel™, Royaltherm™, Trilene ™, Vistalon™. PP/EVAc blends are immiscible, thus in twocomponent systems only a small amount of EVAc can be used, e.g., to improve dyeability, flexibility, electrostatic dissipation, or barrier properties. The hydrolyzed EVAc (EVAl) was also used [Minekawa et al., 1969]. In most cases the PP/EVAc blends are part of more a complex, multicomponent system comprising a reactive compatibilizer (see Table 1.50). 1.4.10

Thermoplastic olefin elastomers (TPO)

Ziegler-Natta catalyst makes it possible to polymerize α-olefins into elastomers with controlled degree of crystallinity and crosslikability. The first EPR’s

Reference Ranalli, 1958 Romankevich & Frenkel, 1980 Matsuda et al., 1980, 1981 Kita and Hashimoto, 1987 Tanaka et al., 1988 Bartz et al., 1993b Hwo, 1994 Silvestri and Sgarzi, 1998 Lustiger et al., 1998

were manufactured in 1960; three years later the first EPDM. It is advantageous to produce block copolymers with PP being the rigid and PE the soft block. A direct sequential polymerization of propylene and ethylene-propylene mixture leads to the reactor-blends (R-TPO) [Cecchin and Guglielmi, 1990]. EPR may be crosslinked by peroxides, while EPDM by the standard methods of the rubber industry. By varying the composition and process variables a wide range of properties can be obtained. The resin with vulcanized, dispersed phase has CUT = 125°C, higher than standard TPO, and they are known as the thermoplastic vulcanizates, TPV [Fritz and Anderlik, 1993]. Diverse TPO’s with properties that range from flexible to rigid (but tough) are manufactured by the large resin producers, as well as by the compounding houses [Utracki and Dumoulin, 1995]. Polytransoctanamer (PTO) has been used as a high performance elastomer and in blends with commodity and engineering resins. Polyoctadecene (POD) blends with PP are thermochromic. The most interesting are the metallocene-type polycyclic polyolefins, e.g., polycyclopentene or polynorbornene, either syndiotactic or isotactic with Tm = 400-600°C.

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Table 1.49. PP/EPR blends Elastomer(s) Added To PP

Comment

Reference

0.1-60 wt% EPR, containing 2-25% ethylene — the earliest patents EPDM (EPR- with -dicyclopentadiene, or -ethylidene-norbornene), partially crosslinked with peroxides EPDM Dynamic vulcanization of PP with either EPR or EPDM; Santoprene™ Dynamically blended PP/EPR, and a peroxide-containing co-polyolefin Sequential compounding of PP, first with EPR then with PE Amorphous EPR + crystalline EPR Bimodal EPR PP/PE, EPR, EPDM, SBS, ionomers, EVAc, EEA, or ESI. Styrene-grafted PP added and “visbreaking” Reactive preblend of PP with either EPR or EPDM (in a ratio from 1:0.01 to 1:0.5) added to PP Dynamically vulcanized BR or CBR and EPDM PP, PE, PS, PMMA or PVC was blended with either EPR or EPDM EPR blends with sPP EPR, EPDM or their mixtures, VLDPE, PP and talc Ethylene-α-olefin copolymers, stereo-block

increased impact and tensile strength, superior mechanical properties alloys could be shaped into articles with good properties without further vulcanization unexpectedly high tensile strength a range of Shore hardness, toughness, elongation, impact strength a master-batch that subsequently was blended with EPDM into TPO co-continuous morphology; good impact and mechanical properties balance of properties, impact strength impact strength & mechanical properties blends foamed with isopentane; good dimensional stability

Schilling, 1966; Short, 1967; Shirayama et al., 1971 Fischer, 1972

polypropylene or EPR Isotactic EPR, Tg < -20°C Random crystalline terpolymers — EPR EPDM with ethylene-methacrylate-zinc, glycidyl methacrylate-acrylate, or epoxy Dynamically vulcanized blends of PP with EPDM, mineral oil and PDMS Poly(4-methyl-1-pentene) with PIB Olefinic, partially crosslinked elastomer: EPDM, EPR, BR, NR, IR, CBR, etc. Metallocene PP was alloyed with EPR PP blended with EPR and EPDM, then irradiated in the presence of O2 Long branch containing PP with partially crosslinked EPR and a foaming agent

improved homogeneity, heat-resistance, impact-resistance and greater flowability high tensile strength sequentially cured and foamed blends transparent, low-T impact resistance moldability, surface appearance, hardness, and impact resistance

Stricharczuk, 1977 Coran and Patel, 1978 Yamamoto and Shimizu, 1979 Huff, 1980 Galli and Spataro, 1983 Makino et al., 1986 Fudge, 1987

Yeo et al., 1989 Puydak et al., 1990, 1992 Cakmak and Dutta, 1992; Dutta and Cakmak, 1992 Asanuma et al., 1992 Nishio et al., 1992

enhanced inter-spherulitic and interlamellar strength modulus, low-T impact strength for fibers with high resiliency and shrinkage, for pile fabrics to improve the scuff resistance

Lustiger, 1993 Winter et al., 1993 Clementini et al., 1993

for slush molding large plastic parts

Hikasa et al., 1994

for food overwrap films for manufacturing automotive components

Nagase et al., 1994 Ellul, 1994

for low-temperature heat sealability easily foamed blends with high strain hardening reduction of density by •2; smooth surface; heat & weathering resistance

Shichijo, 1994 DeNicola et al., 1997

Dawson, 1993

Okada et al., 1998

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Introduction to Polymer Blends

Table 1.50. PP/EVAc blends Elastomer Added To PP

Comment

Reference

PP with 18-32 wt% EVAc and/or ethylene-ethylacrylate copolymer (EEA) PP with EVAc

impact strength, elongation, and low-T brittleness dyeability, flexibility, barrier properties, and toughness for paper-like films low-temperature impact resistance melt strength and rigidity for films, moldings and extrusions excellent solvent resistance miscible (?) alloy for intermediate layer in recyclable barrier films low permeability by gases or liquids tough, radiation-resistant, heat-sealable

Miller & Reddeman, 1962

PP/EVAc and POM, PMMA, PS or SMM 70 wt% PP and EVAc, PVCAc, HDPE PP/EVAc with maleated-LDPE 35 wt% PP, 50% PIB and 15% EVAc PP/PC with 2-35 wt% EVAc 10-90 wt% PP, 5-60 wt% EVAc and 5-50 wt% PEOX PP/EVAc with PP-MA PP/EVAc with EAA and polybutene

1.4.11

PP/Engineering Resin Blends

1.4.11.1 PP/PA Blends There are three types of PP/PA blends: (1) with a small percentage of PO, either acidified or not, (2) alloys with high component’s ratio where PA is a matrix, (3) and blends with a small amount of dispersed PA to increase rigidity. Table 1.51 gives some examples of these systems. To toughen PA, 2-5 wt% of either PO, elastomer, ionomer, acidified or epoxidized copolymer may be added. PA/PO blends of type (2) were developed to improve dimensional stability and to reduce water absorbency of PA. Alloying PA with PO reduces the rate of water migration to and from the blend, but not the inherent water absorption of PA [Utracki and Sammut, 1991, 1992]. The alloying is either a two- or three-step reactive process: (1°) acidification of PO, (2°) preparation of a compatibilizer, and (3°) compounding PP, PA, and the compatibilizer. Usually, the reactive blending is carried out in a twin screw extruder [Nishio et al., 1990; Hu and Cartier, 1998]. Since it may cause reduction of the blend crystallinity (thus performance), the extend must be optimized. The rigid PA/PP blends usually comprise PA:PP =

Sakata et al., 1968 Yamamoto et al., 1971 Kojima & Tanahashi, 1972 Idemitsu Petrochem., 1983 Shulman, 1984 Giles & Hirt, 1986 Sanchez et al., 1991 Kamal et al., 1992 Wilfong & Rolando, 1993

3:2 with 12 wt% of a compatibilizing copolymer. Finally, in type (3), incorporation of PA improves processability, solvent resistance, CUT, HDT, and surface finish. For enhanced performance the blends should be compatibilized. 1.4.11.2 PP/PC Blends PP and PC are immiscible, thus excepting the exploratory use as a “plastic paper,” only the two ends of the concentration range have been explored, viz. addition of 5 wt% PP to PC (to improve processability of PC) [Dobkowski, 1980], or addition of ”10 wt% of PC (to improve PP processability, enhance crystallinity and crystallization temperature, the appearance, modulus, and impact strength) [Liang and Williams, 1991]. For concentrations •10 wt% compatibilization is necessary. This is accomplished using: ethyleneacrylic copolymer, cellulosics, PA, PVAc, or TPU [Goldblum, 1963, 1964]; an acrylic elastomer, acrylic elastomer with PP-MA and either butyl rubber, or isobutene-isoprene rubber [Teijin Chem., 1982, 1983]; SBR and EEA [Liu, 1984]; MBS [Overton and Liu, 1984]; or EVAc [Giles and Hirt, 1986].

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Table 1.51. PP/PA blends Composition 1. Toughened PA PA, PP and 0.5 wt% of PP-MA PA/PO and ionomer PA/PO and PO grafted with MA or a compound containing two epoxy groups PO grafted with glycidylmethacrylate, acrylamide, vinylpyrrolidone, acrylic acid esters and/or methacrylic acid esters, then blended with PA or PEST 2. Rigid PA/PP Blends Non-compatibilized PP/PA blends PP with PA-6 or PA-66 and GF PP/PA-6 compatibilized by EPR-MA PP was maleated then reactively blended with PA to obtain 12 wt% of PP-co-PA PA was blended with acidified PP and GF Reactive blending of PP/PA-6 with either EPR-MA or SEBS-MA 3. PA/Elastomer blends PP, PA-6-co-PA-66, PP-MA with NBR and EPDM PP, EPR, PA, and SMA, SEBS-MA, EPR-MA, MBA, ASA, etc. PP, EPR, PA, and 2 compatibilizers from between: SEBS, SEBS-MA, or poly(ethylene-co-acrylic ester-co-MA) PP-MA was blended with PB

Comment

Reference

compatibilized blends Surlyn™ as a compatibilizer high impact strength

Davis, 1975 Toyobo, 1981 Subramanian, 1983, 1984

Two-stage, reactive impact modification: preparation of reactive compatibilizer, then blending

Teraya et al., 1994

for films, sheets or bottles with good printability and low liquid permeability rigid, non-compatibilized blends impact resistant blends two-step blending: maleation of PP, incorporation of PA moldability, water resistance, HDT, low-T impact strength SEBS-MA gave better impact strength, yield stress, toughness and modulus

Mesrobian and Ammondson, 1962 Asahi Fiber Glass, 1981 Katsura, 1986 Glotin et al., 1989

impact strength, high stress at break, high elongation, good ductility and high gloss for automotive, electrical, electronics, building, furniture, small appliances, etc. recyclable high impact strength blends

Tokas, 1981

adhesive to metals and polar polymers: PA, PET or PC

Lee et al., 1994

Iwanami et al., 1989 Rösch and Mülhaupt, 1994

Chundury and Scheibelhoffer, 1994 Chundury, 1994

1.4.11.3 PP/PEST Blends Blends of PP/PEST are at an early stage of development. PP is antagonistically immiscible with PEST, and when the concentration of the dispersed phase exceeds 5-10 wt% compatibilization is necessary. Initially, the uncompatibilized

blends were formulated within the low concentration region of the dispersed phase. By the end of 1980’s, reactive compatibilization started to dominate the technology. Examples of PP/PEST blends are listed in Table 1.52.

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Introduction to Polymer Blends

Table 1.52. PP/PEST blends Composition

Comment

Reference

PAr with low concentration of PO, elastomer or EVAc PBT with PP PET, PO and either glycidyl methacrylate-modified PO or vinyl-grafted PO PBT/PC with PE, PP or EPR modified with 0.05-15 wt% glycidyl methacrylate and/or 0.1-2 wt% MA PPE, PA, or PEST with EPDM, SBS, SEBS, or EPR, and PP-grafted with styrene and a glycidyl moiety Recycled PET/PP = 2:3-3:2 and SEBS-MA PP with PET or PBT and a polyolefin-polyester graft copolymer PEST, PA, PP-MA and either PET or PBT with sodium dimethyl 5-sulfo-isophthalate

uncompatibilized blends

Koshimo et al., 1973

improved moldability of PBT high heat and moisture absorption resistance processability, toughness, rigidity, strength, dimensional stability, and flexural modulus flexural and tensile strength

Seydl and Strickle, 1976 Mukohyama, 1993 Fujie, 1993

for molding battery containers compatibilized

Tekkanat et al., 1993, 1994 Fujita et al., 1994

high tensile and impact strength

Tajima et al., 1994

Okamoto et al., 1993, 1994

1.4.11.4 PP/PPE Blends

1.4.12

Two-component blends of PP with PPE are unknown. To get a reasonable performance out of a PP/PPE mixture, first PPE must be toughened using a styrenic resin, then the blend compatibilized. Thus (see Table 1.53), PP is only one ingredient in multicomponent PPE blends.

Most specialty resins require high processing temperatures, while PP usually has Tproc < 250°C. Thus, only few blends of this type are known, none commercial. For example, addition of PP enhanced throughput of PAES [Gowan, 1969]; PEI; PC/PEST blends, etc. Linear, aromatic polyamides (PARA) are either liquid crystalline, semicrystalline, or amorphous. Mainly the latter resins are used in blends with PP, viz. PARA with PP-MA [Iwanami et al., 1990]; PARA with PE-MA, or PP and hydrazine [Yoshihara, 1990]. Blends of copolyphthalamide (PPA) with PP were compatibilized using either PP-MA [Paschke et al., 1993, 1994]; or PP-grafted with acrylic acid [Brooks et al., 1993, 1994]. PP blends with a small amount of LCP are of industrial interest for two reasons: (i) to improve processability, or (ii) to improve the mechanical performance. The second effect depends on the blend’s morphology, i.e., on the orientation of LCP domains. The latter, depends on the concen-

1.4.11.5 PP/POM Blends PP/POM blends are antagonistically immiscible, not available on the market. The early blends contained a small amount of one resin in another, viz. either PP or LDPE, with 1-10 wt% POM showed improved processability, good appearance and improved performance [Rudin and Schreiber, 1964]. However, blends containing higher concentration of ingredients must be compatibilized, viz. by addition of muconic acid-grafted PP [Chen et al., 1991], or by TPU and EBA-GMA [Subramanian, 1992].

PP/Specialty Polymer Blends

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L. A. Utracki

Table 1.53. PP/PPE blends Composition

Comment

Reference

PPE-MA with PP-MA, SBR, glycidyl methacrylate copolymer and/or phenylenediamine binder PPE dispersed in: PP, PE, PA, PEST, POM, PPS, or PEEK, compatibilized by: EPR-MA, EVAc-GMA, and either maleic anhydride or bis(4-phenyl isocyanate) PPE, PP and SEBS PPE with PP-grafted with styrene and MA and EVAc-GMA, styrene-grafted EPDM PPE, PP and PPS with hydrogenated styrene-isoprene block copolymers PPE modified by glycidol or epichlorohydrin with PP-MA or PP-GMA and PA or PEST PPE with PP, compatibilized with styrene-grafted propylene-methyloctadiene (or hexadiene) PPE/PS, styrene-grafted polypropylene (PP-PS), SEBS and PP PPE/PP compatibilized with PO grafted with ethylenically unsaturated t-alkylcarbamate

solvent resistance, moldability, impact and mechanical properties

Togo et al., 1988

processability, heat resistance, and mechanical properties

Nishio et al., 1988, 1994

modulus, rigidity, tensile strength and HDT rigidity, as well as high heat, chemical, and impact resistance heat, impact and solvent resistance

Lee, 1990 Furuta & Maruyama, 1990

moldability and mechanical strength

Nakano et al., 1992

stiffness and impact strength

Tanaka et al., 1992

processability, HDT, impact and tensile strength, stiffness solvent resistance, gloss copolymers used as decomposable compatibilizing agents

DeNicola and Giroux, 1994

tration, viscosity and elasticity ratios, interfacial tension coefficient, flow type and intensity, total strain, draw-down ratio, etc. Three stages of orientation are: (1) drop deformation, (2) fibrillation of the domains, and (3) stretching of the LCP chains [Champagne et al., 1996]. Only the latter provides a reasonable cost-to-performance ratio. Examples of PP/LCP systems are listed in Table 1.54.

1.5

Engineering Resins and their Blends

The term engineering polymer is applied to a processable resin, capable of being formed to precise and stable dimensions, having high performance

Maruyama & Mizuno, 1990

Campbell & Presley, 1995

at CUT • 100°C, and the tensile strength of σ • 40 MPa [Utracki, 1989a]. Five polymer families belong to this category: PA, PEST, PC, POM, and PPE. While the relative size of the engineering resin market varies from country to country, these polymers constitute about 13% by volume and 34% by value of the total plastic’s consumption. Since the blends of engineering/commodity resins have already been discussed, here only engineering/engineering and engineering/specialty resins will be considered. Engineering resins and their blends have been foamed using mainly chemical foaming agents, e.g., hydrazodicarboxylates, benzazimides or 5-phenyltetrazole. However, products of decom-

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Introduction to Polymer Blends

Table 1.54. PP/LCP blends Composition

Comment

Reference

Low concentration of LCP in a resin 10 wt% of LCP was blended with, e.g., PP, PS, PC or PI LCP &: PP, PET, PA, PC, PE, PVC, PVDC, PPS, PVDF, PVF or PMMA PP was blended with LCP

to improve processability molecular orientation imposed by extrusion through a rotating die to produce oriented films with protrusions that resulted in low friction 2 extruders with a static mixer, to stretch LCP into microfibrils (processing as described above) (1) prepregs from an extruder, mixer, and rolls; (2) prepregs consolidated at T < Tm.

Cogswell et al., 1981 Haghighat et al., 1992

PP/LCP compatibilized with PP-MA LCP with either: PP, PPE, PPE/PS, PC, PEI, etc.

position of these agents (e.g., alcohol, ammonia, water, etc.) were found to hydrolyze PC, PA PC/ABS, or PEST, reducing the performance of foamed products. For this reason, dihydrooxadiazinones were proposed as the preferred chemical foaming agents for the engineering resins, their compositions and blends (including PPE/ HIPS) [Nizik, 1978; 1979]. Another method for foaming high temperature polymers profits from thermal instability of polypropylene carbonate (PPCO, MW = 1 to 1000 kg/mol) [Kuphal et al., 1990]. The process involves blending PPCO with a resin or its blend that is to be foamed. The blend is then heated to T > 300°C, sufficient to melt the principal resin and to decompose PPCO. Talc, wollastonite, clay, CaCO3 or citric acid can be used as a nucleating agent. The densities of the molded articles were reduced by at least 50%. Most recently, foaming of the engeneering resims by supercritical CO2 was introduced. 1.5.1

Polyamides (PA)

Polyamides are abbreviated as PA-, followed by the number of carbons in a di-amine and dicarboxylic acid, viz. PA-66 has six carbons in each component, etc. Poly-ε-caprolactam or polyamide-6 (PA-6) was first studied by Von Brown in 1905-10, and 30 years later commercialized as Perlon™. In the year 1936, PA-66 and several other

Wong, 1990, 1994 Sukhadia et al., 1991, 1992 Baird and Datta, 1992 Isayev, 1991, 1993

aliphatic and semi-aromatic polyamides were disclosed [Carothers, 1937]. Two years later, DuPont introduced Nylon™, an aggregate of New York and LONdon. PA-11 was commercialized in 1955, PA-12 in 1966, PA-612 in 1970, and PA-46 in 1987. In 1976 DuPont started production of the super tough PA-blends [Damm and Matthies, 1990]. In 1991 consumption of the thermoplastic PA’s was estimated at 1.2 Mton. Comparative properties of the better known PA’s are listed in Table 1.55. Liquid crystal aromatic polyamides (PARA), poly(meta-phenylene isophthalamide), Nomex™, and poly(para-phenylene terephthalamide), Kevlar™-49, were commercialized in 1961 and 1965, respectively. Amorphous aromatic polyamide, Trogamid™, was introduced in 1969, and polyphthalamide Amodel™ in 1991. PA started to be blended with other engineering resins at the end of the 1960’s, viz. PA with POM [Moncure, 1969; Asahi Chem., 1969]; PA with POM and PET [Fujiwara, 1971]; PA-6 reactively compatibilized with PET [Reimschussel and Dege, 1969]; PA-6-co-diisocyanate copolymer with PET [Illing, 1970], etc. By the year 1970 the number of PA blends rapidly started to increase. The main efforts were directed toward improvement of toughness and processability of PA. Reactive compatibilization and impact modification became an integral part of the PA-blends’ technology.

64

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Table 1.55. Comparative properties of PA’s Property

PA-6

PA-66

PA-11

PA-12

PA-46

PA-610

PA-63T

Tensile Modulus (MPa) Tensile Stress (MPa) Maximum Elongation (%) Density (kg/m3)

1400 40 200

2000 65 150

1000 50 500

1600 45 300

1700 59 60

1500 40 500

• 2000 • 75 50-150

1130 57 220 80

1140 57 255 105

1040 46 185 55

1020 37 180 140

1100

1080 50 215 95

1060-1120 145 240 150

Tg (°C), Tm (°C) HDT (°C)

295 90

1.5.1.1 PA(1)/PA(2) Blends

1.5.1.3 PA Blends with POM, PEST or PC

These blends are commercially available, e.g., Zytel 3100™, Grilon™ BT. They show improved processability, solvent resistance, elongation, lowtemperature impact and tensile strength, as well as enhanced barrier properties (see Table 1.56). They have been also incorporated into more complex, multicomponent systems, e.g., PA/PARA = 1:1, PPE, PCL, ionomer, EPR, a monomeric mixture of oxide and/or carbonate [e.g., ethylene carbonate, ethylene oxide, etc.] and a polyhydric alcohol [e.g., ethylene glycol, or trimethylene glycol]. The alloys were used to mold parts for the automotive, electrical or electronic industries [Hamada et al., 1994].

The immiscible blends with ”5 wt% of either component were introduced first (e.g., POM with PA or PARA) before adequate methods of compatibilization were developed. Owing to the crystalline nature of these resins, the blends should also be impact modified. It is noteworthy that in blends of semicrystalline resins the total crystallinity tends to increase [Nadkarni and Jog, 1991]. The compatibilization and impact modification are often accomplished using a multi-polymer. For example, POM/PA-66 blends have been modified by adding either an ethylene-methylacrylate copolymer (EMAC), PEG [Kohan, 1982], or a melamine-derivative “dispersant” [Tsukahara and Niino, 1992, 1994]. In PA/PEST blends, PA improves the processability, mechanical properties and solvent resistance of PEST. The blends are being commercially used, but they are not available on the open market. Examples of the developed PA/PEST blends are listed in Tables 1.58-1.59. PA blends with PC are similar to those with PEST. Since addition of PA to PC may lead to crystallization of the latter resin, also these blends should be compatibilized and impact modified (see Table 1.60).

1.5.1.2 PA/PPE Blends PA’s are excellent candidates for blending with PPE — each ingredient compensates for deficiency of the other. Since the resins are immiscible and brittle, they must be compatibilized and toughened. In consequence, PA/PPE blends comprise minimum four polymeric components: PA, PPE, a styrenic modifier and an acidic compatibilizer. Usually PA is the matrix in which PPE/styrenic resin domains are dispersed. As time progresses, these blends are getting more complex — examples are shown in Table 1.57. Commercial blends are: Artley, Dimension™, Luranyl™, Lynex™ A and Xyron™, Noryl™ GT, Remarry™, Ultranyl™, Vestoblend™, etc.

1.5.1.4 PA/Specialty Polymer Blends Addition of a small amount of PA improves processability of the specialty resin and is beneficial

65

Introduction to Polymer Blends

Table 1.56. PA(1)/PA(2) blends PA(1)

PA(2)

Comment

Reference

PA-66

PA-6 or PA-610

Stott and Hervey, 1958

X-linked PA-66 PARA PA(1) PA-6I6T PA-6 PA-66

PA-6 or PA-11

reduced residual stresses in moldings excellent impact resistance to improve the impact strength miscibility/immiscibility studies tough and strong fibers or films toughened by sulfonated-EPDM printability, clarity, barrier, & dimensional stability resistant to cracking when exposed to metal halides SMM-MA & mineral filler with aminosilane-treated GF tough, high-T, films; good barrier against O2 permeation

Dynamit Nobel, 1969 Zimmerman et al., 1973 Unitika, 1982 Weaver, 1983; 1985 Mollison, 1984

PA-6

PA-6, PA-66, PA-610, PA-11, or PA-12 PA(2) another PA at 1:1 PA-11 either PA-6, PA-612, PA-11, or PA-12 poly(trimethyl hexamethylene terephthalamide)

PA-66 PA-66 PA-mXD6

PA-6 PA-6 PA-66 and PBT; or PA-6

PA-12 PA-6I6T

PARA PA-612 or PA-666, with: PA-6, PA-11, or PA-12 semi-aromatic copolyamide

PA-6 or PA-66 PA-66

PARA

to performance of the GF-reinforced systems. Addition of specialty resin to PA enhances the thermal behavior and rigidity of the latter resin. For higher concentration of these ingredients, compatibilization and impact modification are required. Since PA is sensitive to heat, oxygen and moisture, the compounding requires a special care. Examples are listed in Table 1.61. 1.5.2

Thermoplastic Polyesters (PEST)

Aromatic polyesters show good performance and have high Tm and HDT (see Table 1.62). Polyethyleneterephthalate (PET) was invented in 1941 and commercialized as Terylene™ fibers. The catalytic transesterification also lead to PPT, PBT, PHMT, PEN, etc. Polybutylene-terephthalate

transparent blends at least 5% shrinkage at 90°C in at least one direction compatibilized with either SMA, maleated EPDM, or MBA maleated elastomer, and filler; low-T impact strength and HDT

Uniroyal, 1968

Ube Ind., 1984 Asahi Chem. Ind., 1985 Toray Ind., 1985 Mitsubishi Gas Chem., 1985 Maj and Blondel, 1993 Vicik, 1994 Schmid and Thullen, 1994 Heger and Oeller, 1994

(PBT) has better processability than PET and lower Tm. Poly(ethylene-2,6-naphthalene dicarboxylate) (PEN) has high modulus, strength, HDT, and excellent barrier properties. Polypropyleneterephthalate (PPT) is a newer resin from Shell. It combines the high crystallization rate of PBT with performance of PET. Polyarylates, [-O-φ-C(CH3)2-φ-CO2-φ-CO-]n (PAr) are linear amorphous polyesters, usually of bisphenol-A with isophthalic and terephthalic acids — they show intermediate performance between that of PC and PEST — high strength, stiffness, excellent impact strength, and HDT. Polyethyleneterephthalateglycol (PETG) is an amorphous copolymer of ethyleneglycol and 1,4-cyclohexanedimethanol, with terephthalic and isophthalic acids. The block copolymers, having

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Table 1.57. PA/PPE blends Composition

Comment

Reference

PPE/PA-12 PPE/PA-66 = 1:1 with SBR-MA PPE/PA and a compatibilizer PPE/PA and a compatibilizer PPE/PA and a compatibilizer PPE/PA and a compatibilizer PPE/PA and a compatibilizer

processability and performance first reactive compatibilization compatibilized by SMA polycarboxylic acid derivatives with citric acid with oxalic dihydrazide acid chloride of trimellitic, terephthalic, or 1-acetoxyacetyl-3,4-dibenzoic acid processability, impact strength, solvent C resistance, HDT = 156° core butylacrylate-MMA; shell SMA mechanical properties, low-T ductility addition of SEBS-MA Dimension™ commercial blends with high elongation, impact strength, HDT, and dimensional stability either EPR-MA, EVAc-GMA, MA, or bis(4-phenyl isocyanate) PDMDPhS with carboxylic acid, amine, epoxy, anhydride or ester groups

Komoto, 1972 Ueno and Maruyama, 1981 Kasahara et al., 1982 Abolins et al., 1985 Gallucci et al., 1985 Lohmeijer et al., 1986 Aycock and Ting, 1986, 1987, 1994

PPE blended with PTO and MA then compounded with PA-12 PPE/PA and a compatibilizer Amino-terminated PPE with PA-MA PPE/PA and a compatibilizer PPE-MA with bis-2-hydroxy ethyl fumaramide, PA, and functionalized ethylene-α-olefin elastomer PPE/PA and a compatibilizer PPE functionalized with trimellitic anhydride acid chloride and dimethyl-n-butyl amine, PA & PDMDPhS PPE/PA and a compatibilizer PPE/PA and a compatibilizer The above blends were modified with styrene-butadiene radial copolymer The above blends were modified with tapered block copolymer PPE capped with trimellitic anhydride acid chloride and blended with PA-66 PPE/PA and a compatibilizer Modified PPE with PA and SEBS PPE modified with glycidol or epichlorohydrin, with PA or PEST PPE-MA with PA and SEB PPE capped with salicylic ester & SBS dispersed in: PS, PA, PEST or PEI PPE, PA, POM and a Lewis acid (e.g., trialkylboran or borate, boric acid or halogenated boron) PPE dissolved in lactam(s), then compatibilized & polymerized

compatibilized by organic diisocyanates by aromatic nitro-derivative 3-fold increase of the Izod impact strength

Droescher et al., 1986 Van der Meer and Yates, 1987 Fujii et al., 1987 Modic and Gelles, 1988 Akkapeddi et al., 1988, 1992

Nishio et al., 1988; 1994 Smith et al., 1990, 1994

Pernice et al., 1992 Bencini and Ghidoni, 1993 Gianchandai et al., 1993

further improvement of the impact strength Yates, 1993 toughened by addition of SEBS

Aycock and Ting, 1994

either EPR-MA, E/GMA, or E/VAc/ GMA, citric, malic, or agaric acid moldability and mechanical properties either PP-MA, PP-GMA, ionomers, or EVAl was added selectively hydrogenated S-B di-block resistance to loss of impact strength after recycling for automotive applications

Ishida and Kabaya, 1994

the compatibilizer was either PPE-MA, SAA, or SMA

Kodaira et al., 1994 Arashiro et al., 1994 Lee, 1994 Richards and White, 1994 Takayanagi et al., 1994

Samuels, 1994

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Introduction to Polymer Blends

Table 1.58. PA/PET blends Composition

Compatibilizer

Reference

PA/PET Blends PET/PA-6 PET with PA-6-co-diisocyanate PET with PA-66 Reactive PA/PEST blends Reactive PA/PEST blends Reactive PA/PEST blends PA/PET blends Reactive PA/PET blends Reactive PA/PET blends Reactive PA/PET blends

with α,α-dimethylol-propionic acid compatibilized and toughened blends crystallization rates and mechanical properties direct coupling catalyst, p-toluenosulfonic acid with Phenoxy polyamide-polyester block copolymer butylacrylate-methyl-acrylate block copolymer phosphoryl azide reactive coupler PET containing toluenosulfonic acid groups

Reimschussel & Dege, 1969 Illing, 1970 Nakamura & Neki, 1981 Mitsubishi Petrochem., 1984 Pillon & Utracki, 1984, 1986 Robeson, 1988 Maresca & Shafer, 1988 Tsumiyama et al., 1988 Bhattacharje et al., 1990 Van Sluijs et al., 1992

to enhance PBT crystallinity acidified ethylene copolymer maleated-PS epoxy compounds either SGMA, or SMA low molecular weight PBT transparent, impact resistant processability, high gloss, chemical, mechanical, electrical properties and HDT PA-co-PAr added; single Tg

Toray Ind., 1983 Sheer, 1982 Mitsubishi Petrochem., 1985 Urabe & Ikuhara, 1989 Watanabe & Inozuka, 1991 Goetz et al., 1993 Asahara et al., 1977 Unitika, 1982- 1984

PA Blends with PBT or PAr Non-compatibilized PA/PBT Reactive PA/PBT blends Reactive PA/PBT blends Reactive PA/PBT blends Reactive PA/PBT blends Compatibilized PA/PBT 1:1 PAr with PA, PET or PBT PAr with PA-6 and Si-compounds or PET PAr with PA or PARA

Dean, 1990, 1992

Table 1.59. Multicomponent PA/PEST blends Composition

Compatibilizer

Reference

Amorphous PA or PARA, with either PEST, PC, PEC or PAr PA blends with either PAr or PC PAr/PA PET, PA-6 and PO PEST, EVAl, PA, and PEST (with Na- dimethyl 5-sulfoisophthalate groups) PEST/PA

a polyamide-polyester block copolymer, PA-b-PEST EGMA either ABS-MA or ABS-GMA EEA-GMA an ionomer, and PP-MA

Maresca and Shafer, 1988 Yuichi and Suehiro, 1989 Yasue et al., 1989 Natarajan et al., 1994 Tajima et al., 1994

reactively blended in solid state

Al Ghatta, 1994

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Table 1.60. PA/PC blends Composition

Comment

PA-12/PC PA-12/PC with PSF, PPE, or PET PA/PC PARA/PC

Reference

for electrical insulation moldability and mechanical performance toughened with SEBS pearly-looking, resistance to oils & water, good mechanical properties Polyestercarbonate (PEC) with PA compatibilized and toughened by MABS PC/PA-6 compatibilized by addition of SMA PA/PC commercial alloys Dexcarb™ polyesteramide, an elastomer, and either PP-MA or EPR-MA PC end-capped with trimellitic moldability, excellent Izod impact strength, anhydride acid chloride, then reactively and elongation blended with PA-6 and MBS PA-6I with PC high impact- and tensile strength PC/PA-6; compatibilized polyethyloxazoline, PEOX, added PA blends with, either PAr or PC compatibilized and modified by EGMA PA blends with, either PAr or PC with ABS-MA or ABS-GMA PA with, PAr or PC and PA-co-PC with glycidyl isocyanurate Branched, bPC, and •1 polymer from: processability, solvent and impact resistance, mechanical properties mechanical properties PC and/or PEC with PA-6 with acrylic elastomer PC/PA-6/ABS with imidized polyacrylates PC/PA-6 with PEI and/or TPU with: butyl glycidyl ether, EPR-MA or EPDM-MA PA-6/PC/SEBS/SEBS-MA 20 fold increase of impact strength, 50-fold increase of elongation

hard PBT segments and soft polyetherglycol ones, are versatile elastomers (e.g., Hytrel™ from DuPont). PEST’s and PC’s are known to have low melt strength and are difficult to process in operations involving elongational flows, viz. blow molding, stretching or foaming. During foaming the cell size distribution is broad, the wall thickness variable and the mechanical properties are poor. To improve the melt strength a diacid anhydride and a metal compound may be added during the extrusion [Hirai and Amano, 1993]. These additives induce branching, increase MW and strain hardening. Alternatively, linear and branched resins may be blended.

Okuzono and Kifune, 1975 Okuzono and Kifune, 1976 Gergen and Davison, 1978 Mitsubishi Chem., 1980 Sakano et al., 1981, 1982 Dainippon Ink., 1983 Perron, 1984, 1988 Hathaway and Pyles, 1988, 1989

Gambale et al., 1988; 1994 Thill, 1989 Yuichi and Suehiro, 1989 Yasue et al., 1989 Derudder, 1990 Kozakura et al., 1992, 1994 Heger et al., 1992 Leitz et al., 1992 Perron et al., 1993 Ind. Technol. Inst., Japan, 1996

The branched macromolecules can be produced in reactions with polyols having 3 to 6 hydroxyl groups. This approach has been successful even for recycling post-consumer PET. For example, • 25 wt% of a bPET with recycled PET were extruder-blended with ” 20 wt% of a chain extender (e.g., partially neutralized ethylene-methacrylic acid ionomers, copolymers of maleic anhydride or glycidyl methacrylate, etc.). Downstream, up to 5 wt% of a chemical or physical blowing agent was incorporated. The blends were used to produce rigid insulation, trays, food packaging, microwave cookware, oxygen and moisture barrier films, etc. [Muschiatti and Smillie, 1995].

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Introduction to Polymer Blends

Table 1.61. PA specialty polymer blends Composition

Comment

Reference

1. PA/PSF Blends PSF with PA-11 PA/PSF with poly(sulfone-g-lactam) PA-6/PSF PARA/PES PA-46/PES PARA/PES PAES or PES blended with PA-6T6

for improved impact properties processability and mechanical properties processability, thermal & mechanical properties HDT = 172°C & mechanical properties mechanical and thermal properties mechanical and thermal performance high heat resistance and stiffness

Nield, 1971 McGrath and Matzner, 1972 Kyo and Asai, 1978 Hub et al., 1986 Koning and Vroomans, 1992 Bapat et al., 1992 Weber and Muehlbach, 1993

moldability, HDT, & impact strength impact and mechanical properties, HDT, heat and chemical resistance heat and chemical resistance, HDT, mechanical properties mechanical properties compatibilized-cum-impact modified mechanical properties, HDT, dimensional and chemical stability tensile strength, arc tracking, and heat resistance

Shue and Scoggins, 1981 Chiba et al., 1978

2. PA/PPS Blends PARA with a small amount of PPS PA-46 with PPS PPS with polyphthalamide (PPA) PPS with PPA and GF PPS/PA with EPDM and MA PPSS with either PA or PEST PPS, either PA-66 or PA-MXD6 and Mg(OH)2 3. PA/Phenoxy Blends PA-6 with Phenoxy PA-66 with Phenoxy and SEBS PA with PEST, Phenoxy and MBS 4. PA/PEA Blends PA-12 with PEA PEBA/PA and (ABS, MBS, NBR, SBR or EPDM) 5. PA/PAI Blends PAI with PA-66 or PARA and inorganic filler PI, PA, PAI or PI+PA and aromatic PI 6. PA/PEI Blends PEI with PA or PEST PEI with PA-6 and PEI-b-PA PEI with PA-6 or PA-66 PEI with PA-12 Copolyesteretherimide, PEEI Lomod™, blended with either PA or PARA

Chen and Sinclair, 1990 Davies, 1990 Yu and Beever, 1992 Ishio and Kobayashi, 1992 Dubois et al., 1993

excellent ESCR tensile and impact strength tensile and impact properties

Schober, 1973 Freed, 1975 Robeson, 1988

tough, flexible, heat and hot oil resistant impact strength at low-T, Rilsan™

Sumitomo Chem., 1984 Arraou, 1986

moldings with high mechanical performance

Toray Ind., 1979

miscible blends for selective permeation

Ekiner and Simmons, 1993

processability impact strength compatibilized by nonyl-phenolic reduced shrinkage and water absorption thermal aging behavior

Giles and White, 1983 Robeson and Matzner, 1984 Gallucci, 1988 Giles, 1987 Angeli, 1992

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Table 1.62. Comparative properties of PEST’s Property Tensile Modulus (MPa) Tensile Strength (MPa) Maximum Elongation (%) Density (kg/m3) Tg (°C), Tm (°C) HDT (°C)

PC

PET

PPT

PBT

PEN

PETG

PAr

2300 62 120

2800 81 70

2500 68 —

2600 52 200

2400 82 100

6700 34 110

2200 60-70 7-100

1200 149 220 280

1375 98 255 167

1350 80 225 149

1300 60 223 136

117 337 109

1250 88 — 70

1190-1210 > 180 — 120-175

1.5.2.1 Polyester Blends The largest group comprises the impact modified PEST’s — these were summarized while discussing commodity resin blends. The commercial blends with, e.g., SMA or acrylic rubber, show good processability, rigidity, impact and tensile strength, as well as excellent weatherability, viz. Arloy™ 2000, Bexloy™, Celanex™, Pibiter™ HI, Rynite™, Ultradur™ KR. The second large group comprise different combination of PEST, for example, PET/PBT, PBT/PAr, or PET/PEN. These have been mainly developed for improved processability, good surface properties, HDT, impact strength and dimensional stability, viz. Celanex™, EMC™, Valox™, etc. Examples of their formulations and performances are given in Table 1.63. 1.5.2.2 PEST Blends with PC The PEST/PC blends are immiscible (Tg of PC is depressed by ca. 20°C) and brittle, requiring toughening. Usually, PC blends with PEST contain 10-20 wt% of ABS, ASA, or MBS. In most commercial blends PC is the matrix, but blends formulated for low-T impact strength (down to -40°C) have co-continuous morphology, and are reinforced by addition of ”30 wt% of GF. It is vital to avoid reduction of PET crystallinity during blending or processing. The main advantage of the PC/PEST alloys is the increased stiffness, reduced susceptibility to stress cracking on contact with

fuels, and an improved resistance to chemicals and fuels. The blends show good processability, heat resistance, ductility, HDT, high modulus at high temperature, good electrical properties, thermal stability, impact, tensile and flexural strength over a wide temperature range, low shrinkage, good dimensional stability, but they may have poor weatherability, and their solvent resistance (while superior to that of PC) is moderate. The commercial blends include: Alphaloy™ MPB, Baitaloy™ VL, Cycolin™, Dialoy™, Ektar™, Hyperlite™, Eastalloy™, Idemitsu™ S, Impact™, Lumax™, Malecca™ B, Makroblend™, Maxloy™, Novadol™, Novaloy™ B, Petsar™, Pocan™, Sabre™, Stapron™, Techniace™ TB, Triax™ 400, Ultrablend™, Valox™, Xenoy™. The blends are mainly used for automotive body panels, in outdoor power or recreational equipment, appliance housings, telecommunications, etc. The PC/PEST blends were first described in 1966. The first, three-component blends were disclosed in 1972 (see Table 1.64). Many multicomponent alloys comprise PC and PEST. From between them, these with PPE are particularly important — see the following Part. 1.5.2.3 PEST Blends with PPE Blending PPE with either PEST or PC poses similar problems — the polymers are immiscible and brittle, hence require compatibilization and toughening. The PEST/PPE blends are multicomponent systems, with • 5 components: PEST, PPE,

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Introduction to Polymer Blends

Table 1.63. Examples of PEST blends Composition

Comment

Reference

First PET/PBT blends PBT with 40 wt% PAr First modified PET/PBT blends PET with PA PAr with either PC or PEST, PBT with PET or PHT PBT with PC PBT with PPE PET with PC PAr with PEST and PA-6 or PPS PET/PBT blends PAr with acetate-capped POM PAr with either PET or PBT PET/PBT/SEBS PBT/PET/PC PBT/PET/Phenoxy and a toughener PAr with PCT PAr/PEST with TPU PAr, PET, and EEA, TPU, etc. PET/PBT or PEST/PC; toughened PBT with either POM or PA PET/PBT and, e.g., EGMA PBT reactively blended with PAr PET with PAr and PA-6 PAr with PEST and PC, ABS, PA PAr/PEST with EGMA or EPR grafted with glycidyl oxybenzyl acrylamide PET/PC/PCT, EPR and/or MBA POM, POM-copolymer, aliphatic polyester and polyether POM/TPU/EBA-GMA/PA or PEST PEST/POM and isocyanate or isothiocyanate coupling agent Solution-prepared PAr/PET Blend of linear and branched PET PET with PCT PET with PEN

enhanced crystallizability, miscibility for electrical insulation, films, moldings toughening with butyl rubber for good frictional resistance moldability and impact strength high elongation and impact strength impact resistance & elongation processability & mechanical properties high stress corrosion resistance further enhancement properties modified by of acrylic elastomers processability of POM transparency, HDT, & impact strength impact, heat, and tensile strength impact strength, rigidity, strain at break butylacrylate-glycidylmethacrylate-MMA processability, weatherability, impact high modulus, while retaining strength to improve impact resistance and HDT toughener: MBS, ACM, or ABS, etc. impact-modified by an ionomer processability & performance high impact strength processability, gloss, HDT, transparency processability blends, set of properties excellent processability and performance

Heywang, 1966 Wiener, 1969 Hiri & Kotama, 1971 Fujiwara, 1971 Koshimo et al., 1973 Tsunawaki et al., 1973 Matsukane & Azo, 1973 Nakamura & Toyomoto, 1974 Horiuchi & Kamiya, 1974 Asahara et al., 1977 Kamata et al., 1978 Gale, 1978 Fukushima et al., 1979 Wambach & Dieck, 1980 Goedde et al., 1980 Charles & Coleman, 1981 Robeson, 1981 Robeson et al., 1981 Robeson, 1981, 1982 Bier & Indner, 1982 Sheer, 1982 Memon & Myers, 1983 Toga & Okamoto, 1983 Unitika Ltd., 1983 Robeson, 1985 Toyoda et al., 1986

high impact resistance processability, crystallization, low shrinkage, HDT, & mechanical prop. processability and mechanical properties processability, mechanical properties and low gloss immiscible: two Tg’s and χ > 0 for high-quality mineral water bottles high impact strength at low-T transesterification reducing crystallinity

Romance, 1990 Makabe et al., 1991 Subramanian, 1992 Katsumata & Matsunaga, 1993, 1995 Chung & Akkapeddi, 1993 Greaves et al., 1993 Stewart & Bell, 1995 Wu and Cuculo, 1998

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Table 1.64. PEST/PC blends Composition

Comment

Reference

PBT with PC and MBS PET/PC with a compatibilizer-impact modifier PCT with PC and branched PC

processability, chemical resistance, impact butylacrylate-styrene-triallyl isocyanurate, grafted with MMA & styrene good clarity, tensile strength and notched Izod impact strength moldability and high impact strength high impact strength & strain at break

Nakamura and Toyomoto, 1974 Kamata et al., 1978

Motz, 1980 Dieck & Cohen, 1980

Xenoy™ alloys, with mineral filers or not

Dieck & Wambach, 1980

with (meth)acrylic esters (and acrylonitrile, acrylate rubber, ACM, EVAc, and/or ABS) good overall performance high impact strength alloys good sets of mechanical properties

Cohen, 1980, 1982; Bier & Indner, 1982 Teijin Chem., 1983 Boutni & Liu, 1983 Teijin Chem., 1983

tensile strength, impact resistance & electrical conductivity; Tafmer™ good moldability, mechanical properties and heat resistance improved HDT, elongation, and impact strength greatly improved impact strength for uniaxially drawn shrink films high impact resistance high impact strength moldable compositions

Mitsubishi Chem., 1983

PBT/ PC with Acryloid™ PBT/ PC & PEC PEST/PC, SEBS, PVAc or (SB)n & butadiene-caprolactone-styrene PEST/PC, BR and an impact modifier; Makroblend™ PET/PC, butyl rubber (BR) PEST, PC, PO, and MBA PEST, PC, BR, an acrylic elastomer, HIA-15, and PMP Reactively blended PEST, PC and an acidified elastomer PEST/PC/styrene-α-methylstyreneacrylonitrile-butadiene PEST/PC/2-methyl-1,3-propanediol PBT/PC with AES PC/PEST with EVAc PET/PC and 10-15 wt% Acryloid™ PEST, PC, and an ethylene-acrylic acid-butylacrylate PC/PEST/elastomer/Phenoxy PC/PET, or PETG, EEA, r and an impact modifie PEST, PC, & MBS-AA or GMA PC/PEST and siloxane copolymer PC/PET with •90 wt% PES PC/PEST with acrylic elastomer PC/PEST and PE-MA wax PET/ PC/EPDM/PCL/EEA-MA PBT/PC, PO-GMA or PO-MA PBT/PC/ABS/PE-GMA or -MA, vinyl, styrene or (meth)acrylate

impact modified blends extrusion-blended, foamed, then injection molded; ρ = 900 kg/m3, good performance high impact strength; Ultrablend™ chemical, weather, impact resistance; Dialoy™ PES foamed by degradation of PET/PC improved impact strength improved impact strength chemically resistant, impact properties moldability, toughness, strength and stability impact strength conditions

Mohn et al., 1979

Biglione & Fasulo, 1983 Yoga et al., 1983 Sumitomo Naug., 1983 Weyer et al., 1984 Hepp, 1984 Portugall et al., 1984 Liu & Giles, 1986 Avakian and Jodice, 1986 Lausberg et al., 1987 Hongo et al., 1987 Haggard, 1987 Sakai et al., 1989 Liang, 1989 Natarajan & Mininni, 1991 Fujie, 1993 Steendam et al., 1993

73

Introduction to Polymer Blends

styrenic copolymer, compatibilizer, and impact modifier. Examples of commercial blends are: Dialoy™ X, Gemax™, or Iupi-ace™. For improved modulus and dimensional stability, they are usually reinforced with GF. These alloys are known for excellent processability, high solvent resistance, and dimensional stability. Evolution of these systems is outline in Table 1.65.

1.5.2.4 PEST Blends with Specialty Resins PEST’s are immiscible with polyphenylsulfides (PPS) or polyphenylenesulphidesulfone (PPSS). The customary three types of uncompatibilized blends are known, with 5-10 wt% of the dispersed (either PPS or PEST) phase and the phase co-continuity. The compatibilized blends (with a copolymer containing either epoxy or acid anhydride groups) show high tensile and impact

Table 1.65. PEST blends with PPE Composition Blends with unmodified PPE PPE/PBT/SBS PPE, PC, SEBS and either PBT or PET PPE/PET/PC/EGMA PPE/PET with SEBS-GMA PPE/PET, EPR-MA, EVAc-GMA, MA or bis(4-phenyl isocyanate) PPE, PEST, PS, PC, SEBS and/or acrylic modifier PPE/PBT with PC-b-PS-b-PC, PC and either SEBS or MBS PPE, PEST or TMPC, PS-VPh, HIPS PPE, PEST, PC-PBT and SEBS Blends with modified PPE PPE modified with unsaturated dicarboxylate, e.g., fumaric acid PPE modified with dimethylsiloxanes, PPE modified with an epoxy-compound PPE modified with end-capped with salicylic acid ester PPE modified with either glycidol or epichlorohydrin

Comment

Reference

processability, impact and tensile strength stable morphology, processability, high impact strength, and solvent resistance for automotive applications excellent solvent resistance, moldability, impact and mechanical properties processability, heat resistance, and mechanical properties High impact strength PPE alloys; Gemax™

Lee, 1978 Brown et al., 1987 Yonetani et al., 1987 Mayumi & Omori, 1988 Nishio et al., 1988, 1994 Yates & Lee, 1990

processability, tensile elongation and strength, as well as chemical and impact resistance PEST is miscible with PVPh, thus PS-VPh is an efficient compatibilizer excellent rigidity and impact strength

Brown & Fewkes, 1992, 1994 Colby et al., 1993, 1994

reactively blended with PBT, PC, and SEBS for improved mechanical performance blends with PPE, PEST, and SEBS for enhanced solvent and impact resistance blended with PPE, PEST and hydrogenated poly-α-olefin for processability and impact resistance blended with SEBS, then dispersed in either PA, PEST, PEI, or PS; for processability, resistance to loss of impact strength after recycling blended with PP-MA, PP-GMA, ionomers, EVAl, PA, and PEST; for moldability, solvent and heat resistance, mechanical strength

Ishihara, 1989

Chambers et al., 1995

Brown, 1989 Hasson & Yates, 1993 Richards & White, 1994

Arashiro et al., 1994

74

L. A. Utracki

strength [Nishiyama et al., 1990, 1991] that can be further improved by addition of a TPE [Nishiyama and Nakakita, 1991]. PPS/PEST blends were also compatibilized by addition of a PPS-PEST copolymer [Suenaga and Ishikuwa, 1991]. The alloys could be reinforced with GF, talc, mica, wollastonite or clay [Gary, 1993]. Polyarylethersulfone: (-[O-φ-O-φ-SO2-φ-]0.25n[-O-φ-SO2-φ-]0.75n-) (PAES) blended with either PAr, PEST, PC, or their mixtures showed well-balanced mechanical properties and good environmental stress cracking resistance [Robeson and Harris, 1985, 1986]. For improvement HDT, mechanical properties and flame retardancy, PET was blended with a PAES: (-[- φ-SO2-φ-O-φ-C(CH3)-φ-O]n-) [El-Hibiri et al., 1992]. Blends of this type PAES with PC, PET, ZnBO3, and PTFE had high flame retardancy [Jack et al., 1993]. Polyetherimide (PEI) was blended with PAr for improved moldability and mechanical properties [Holub and Mellinger, 1981]. Similarly, polyestercarbonate (PEC) was blended with either PEI, PA, PI, PAI, or their mixtures, to give alloys characterized by high HDT and tensile strength [Quinn, 1984]. Blends of PEI, PAr, and PC showed good moldability, flexural strength and modulus [Holub, 1990]. Also PEI, PET, and PEC blends had good processability and impact strength similar to PEI [Quinn and Holub, 1986]. Adding ABS to PEI/PEST blends enhanced the impact resistance [Gaafar, 1990]. PEI was also blended with PBT and a crosslinking triallyl cyanurate and triallyl isocyanurate [Hosoi, 1991]. To produce soft, flexible, dimensional stable, and water resistant materials, polyesteramide (PEA) was blended with PET (and/or PA-6) [Toray Industries, Inc., 1982]. The low temperature phosphate glasses (LTG: -[-Zn-O-P(O2Na)-O-P(O2Li)-O-Zn-O-P(O2Na)-O]-) have been blended with: PET, PBT, PEK, PEEK, PPS, PEI, LCP, PC, etc. [Bahn et al., 1991]. The blends containing up to 65 vol% of LTG were reported to have good mechanical properties [Frayer, 1993, 1994]. The technology makes it possible to generate controlled morphology of the dispersed LTG phase, as well as to ascertain its good bonding to the organic matrix.

1.5.3

Polyurethanes (TPU)

The polyester-type thermoset polyurethanes were commercialized in 1942, and the linear thermoplastic ones (TPU) ten years later. Polyester-type TPU, Texin™ resins for extrusion and injection molding, were introduced in 1961, whereas polyether-type, Roylar™, in 1971. Owing to great diversity of the ingredients, the TPU performance can be readily modified. For this reason, as well as because of the cost, TPU’s are seldom blended. Their use can be divided into three groups: (i) blends with POM, (ii) blends in which TPU is used as a compatibilizer and impact modifier, and (iii) others. 1.5.3.1 POM/TPU Blends The oldest and most common method of POM toughening is by incorporation of TPU. Delrin™ POM/TPU blends were commercialized in 1960. The others followed, e.g., Celcon™, Duraloy™, Formaldafil™, Fulton™ KL, Hostaform™, RTP 800, TC, or Ultraform™. These alloys have high impact strength and elastic recovery that depend on composition, morphology, as well as on the compounding and processing methods. Examples of these blends are listed in Table 1.66. 1.5.3.2 PC/TPU Blends TPU has been used to toughen PC, to enhance its mechanical behavior and ESCR. The blends have been used in industrial and medical applications [De Boer and Heuschen, 1988; Pinchuck, 1991]. Blends of PC/PET/TPU with EVAc-GMA and optionally MBS or ABS, have good flexural modulus, strength, weld line strength, solvent resistance, and impact behavior [Laughner, 1994]. PC blends with a polycaprolactone-polyurethane resin, TPU Pellethane™, and either MBS or MBA showed similar behavior [Henton et al., 1993]. Texin™ 3000 and Texin™ 4000 are the commercial blends. In the first TPU is the matrix, while in the second PC plays this role. The resins are used for the production of gears, tubings, housings, top-lifts, extruded profiles, for the automotive industry and consumer goods.

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Introduction to Polymer Blends

Table 1.66. POM/TPU blends Composition

Comment

Reference

POM with OH- NCO- or NCS-terminated TPU POM/TPU with mineral fillers POM/TPU, & polycarbodiimide and ethylene-bis-stearamide POM/TPU POM, PC and TPU POM with polythioisocyanate-TPU POM/TPU and acrylics POM/TPU and ABS

to improve POM elasticity for enhanced stiffness moldability and impact strength (POM was acetate end-capped) high impact strength at low-T high impact strength impact and flexural strength abrasion resistance and weatherability processability, thermal & dimensional stability, chemical & creep resistance for lower mold shrinkage, good stiffness, elongation, toughness, etc.

McAndrew, 1971 Reske and Wolters, 1984 Richardson, 1984

POM/TPU, with SAN, ABS, AES, PC, PA, PAr, PPE, HIPS, acrylics, imidized acrylics or SMA POM/TPU and EBA-GMA, with PA-612, PA-6, PP or PET POM/TPU and di-glycerin, pentaerythritol,

1.5.4

good modulus, impact strength and processability tensile strength & elongation at break,

Polycarbonate (PC)

Polycarbonates are polyesters of polyhydric phenols and carboxylic acid. Except for the lack of crystallinity, their properties resemble those of PEST. The most common polycarbonate, that of bisphenol-A (PC), was commercialized in 1956. PC is tough, transparent, self-extinguished, dimensionally stable, resistant to salts and oxidation, but susceptible to abrasion, stress cracking and attacks by solvent, acids and alkali. Its Tg = 149°C, but the ductile-brittle transition is at 0-10°C. The resin ought to be toughened, for example by addition of ABS, MBA or MBS. The annual consumption and its growth rate of PC are, respectively, 700 kton and 5%. About 40 % of PC is used in blends. Originally, the commercial PC resins were linear polymers with high shear viscosity and low melt strength, thus difficult to process in operation involving extensional flows, viz. blow molding stretching, foaming. Several years ago branched PC (bPC) became available. The resin is usually blended with linear PC at the ratio that on the

Drzal et al., 1986 Silvis et al., 1990 Sugiyama and Mochizuki, 1990 Endo et al., 1990 Guest et al., 1991 Flexman, 1992

Subramanian, 1992 Nagai et al., 1993, 1994

one hand is economically viable and on another that provides sufficient melt strength for the required process. For example, 60-95 wt% bPC (MW = 32-45 kg/mol) was blended with 5-40 wt% PC (MW = 15 to 27 kg/mol). The blends with MFR = 2 - 8 g/10 min were suitable for extrusion, injection molding, blow molding, and/or foaming at Tprocess = 250 - 310°C [Van Nuffel et al., 1998]. 1.5.4.1 PC Blends with PPE Blends of PPE with PC are immiscible and brittle, thus they must to be compatibilized and toughened (see Table 1.67). 1.5.4.2 PC/POM Blends These blends are immiscible, thus should be compatibilized and toughened. Addition of POM to PC improves the solvent and chemical resistance [Miller, 1972]. PC blends with POM and TPU were easy to mold into articles having high impact

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Table 1.67. Examples of PPE/PC blends PPE/PC and ...

Comment

Reference

PS or SMMA and PEST

compatibilizer: acrylic elastomer, Phenoxy; or Cl-SBR processability, impact strength, solvent resistance advantageous set of performance characteristics

Izawa & Nakanishi, 1973

mechanical properties and low gloss high impact strength excellent rigidity and impact strength

PBT and SEBS HIPS/PEST/SBS and PS grafted with 2-oxazoline (PSOX) Fumaric acid-grafted PPE and SEBS PS, PEST, SEBS and/or MBA PBT, PC-PS-PC block copolymer and impact modifier SAN/PS — compatibilized PEST/ABS with PS, HIPS, SEBS,

poly(butylacrylate-co-styrene-co-acrylonitrile) multi-component blends prepared by

Ishihara, 1989 Yates & Lee, 1990 Brown & Fewkes, 1992, 1994 Niessner et al., 1993 Laughner, 1993; 1994

PA, PC and/or PEST PEST, PC-PBT copolymer and SEBS

reactive processing good rigidity and impact strength

Chambers et al., 1995

strength [Silvis et al., 1990]. POM-b-PC was used either as a compatibilizer, or as a modifier of performance for such polymers as PES, PEEK, PA, PAN [Dhein et al., 1993]. 1.5.4.3 PC Blends with Specialty Resins There is a great diversity of polyimides (PI) having Tg = 180-420°C. Several were blended with PC to improve its stiffness, HDT and strength. PEI/PC blends were commercialized in 1992 as Ultem™ LTX, for injection molding or extrusion. They show higher impact resistance than PEI and higher heat resistance than PC, as well as they retain the stain, chemical resistance, and the hydrolytic stability of PEI. Fluoropolymers are notoriously immiscible with any other polymer. Usually, they are dispersed in blends of engineering and specialty polymers either to improve processability or to induce lubricity and abrasion resistance. Examples of the PC/Specialty resin blends are listed in Table 1.68.

1.5.5

Brown et al., 1987 Avakian et al., 1988

Polyoxymethylene (POM)

Polycondensation of formaldehyde was reported by Butlerov in 1859, but only in 1950 DuPont developed end-capping that prevented unzipping. POM is crystalline, thus rigid, brittle, and chemically non reactive. Production of Delrin™, and Celcon™ started in 1959 and 1962, respectively. The world consumption of POM and its annual growth rate are: 500 kton and 5%. The most common POM blends are homologous mixtures of POM’s having different molecular structure (linear, branched, crosslinked) [Matsuzaki, 1991], different molecular weight [Ishida and Sato, 1970], or with different end-groups [Nagasaki et al., 1991; Hanezawa and Ono, 1991]. On the second place are blends of POM with TPU, preferably polyester-type. POM’s are also blended with core-shell acrylic elastomers, MBS or MBA. Commercial blends of POM with PEST are available. To improve weatherability of POM, the resin was blended with PMMA and a fluoropolymer (viz. PTFE, PVF, PVDF) [Katsumata, 1991].

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Introduction to Polymer Blends

Table 1.68. PC blends with specialty polymers Composition 1. PC/Siloxane resin Blends PC with PDMS PC and/or PEST with siloxane-based vinyl copolymer PC with poly(dimethyl siloxy biphenylene oxide) PC with siloxane/vinyl-based copolymer PC, or polyestercarbonate with PC-b-PDMS PC/PArSi and SBS 2. PC/PSF Blends PAES: [-φ-C(CH3)2-φ-O-φ-SO2-φ-O-]n, with PC PC with PAES and MBS or AES PC with PAES and CHR PSF with linear and branched PC PSF with PEC PC and polycarbonate-sulfone grafted with ethylhexyl acrylate PAES with PC, PAr or PEST PSF/PC with PET or PBT and GF PC/PAES with MBA

Comments

Reference

solution cast films with good properties for chemical, weather, and low temperature impact resistance transparent, flame and impact resistant alloys thermal stability, ductility at low-T, and impact resistance low-flammability, impact strength over a wide T excellent mechanical properties

Caird, 1961 Hongo et al., 1987 Jordan & Webb, 1992 Derudder & Wang, 1993 Hoover, 1993 Jordan & Webb, 1994

impact, tensile strength, elongation at break impact and tensile strength, HDT good impact and fire resistance improved solvent resistance impact strength and flame retardancy high impact strength

Union Carbide Corp., 1966 Grabowski, 1971 Lauchlan & Snodgrass, 1973 Binsack et al., 1979 Quinn & Rosenquist, 1982 Tyrell et al., 1984

well-balanced mechanical properties performance, solvent and chemical resistance BR grafted with MMA, styrene, and/or AN; then (shell) MMA, styrene, AN

Robeson & Harris, 1985-86 Militskova et al., 1993 Weber & Muehlbach, 1993

3. PC/Fluoropolymer Blends PC with ETFE PC with PP and with PTFE PC/PVDF with co-polyacrylics PC with PTFE PC with PTFE, ABS and cresol novolak phosphate oligomers

processability, lubricity & abrasion improved performance miscible, yellowish films, Tg = 120°C processability, lubricity and abrasion improved performance, flame resistance

Kawai & Miyauchi, 1974 Kishimoto, 1976 Leibler & Ringenberg, 1986 Akega, 1991 Fuhr et al., 1992

4. PC/Polyimide Blends PC with SMI PC with PEI PEI/PAr, PC and phenyl phosphate PEI with PAr, PC and HIPS PEI with PAr and PC PC with PAI PC with a carbodiimide: (X)m- (-N=C=N-Y-)p-(-N=C=N-X)m

improved stiffness, HDT and strength. processability, flexural & impact strength improved melt stability flexibility and impact resistance HDT, flame and impact resistance mechanical and antistatic properties for sheath optical fibers

Fava, 1979 Giles, 1983, 1984 Peters & Rock, 1989 Holub & Rock, 1989 Holub, 1990 Shimamura & Suzuki, 1991 Kamps et al., 1994

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For the manufacture of sliding parts, POM blends were developed with: PTFE [Ishioka, 1991]; with PVDF [Shibata et al., 1992]; with either wax, PTFE, silicone oil, or PEG, and EBA-GMA [Takahashi and Kobayashi, 1993]. Recent blending technology of POM involves introduction of the reactive end-groups — it makes compatibilization with other polymers relatively simple. Addition of PPE/PS to POM was used to improve processability, HDT, and mechanical properties [Ishida and Masamoto, 1974]. More recently, POM with a Lewis acid was incorporated into PPE/PA blends to improve compatibilization and induce high heat and impact resistance [Takayanagi et al., 1994]. POM blends with specialty polymers are formulated either to take advantage of POM (the resistance to abrasion), or of the specialty resin (e.g., to improve stiffness and wear resistance by incorporation of PEEK or PEI) [Suzuki and Nagahama, 1987]. POM is miscible with polyvinylphenol, PVPh, thus addition of PVPh to blends of POM with COPO compatibilizes the system [Machado, 1993]. 1.5.6

Polyphenylene ether (PPE)

In 1956, by oxidative coupling of 2,6-dimethyl phenol, poly(2,6-dimethylphenyl ether) was obtained (PPE) [Hay, 1959, 1964]. The resin was commercialized in 1964. PPE is amorphous (Tg = 210°C), but it can crystallize (Tm = 257°C). It is thermally stable only to T ” 150°C (CUT = 125°C). It has good rigidity, creep resistance, dimensional stability, high electrical, chemical, moisture and flame resistance. The main disadvantages are: processability, oxidative degradation, low impact strength, and weatherability. The resin is usually “plasticized” by blending with styrenics. The annual production and growth rate of PPE are respectively, 300 kton and 9%.

modified by blending (e.g., with HIPS, ASA, MBA, SBS or SEBS, etc.), by grafting [Brown, 1989], or by reacting its end-groups [Richards and White, 1994]. Modified PPE is mainly blended with other engineering resins — most of these were already discussed. The principal types of PPE blends are with styrenics, with PA, and with PEST. Owing to miscibility of PPE with PS, the compatibilization is relatively simple. However, blending PPE with either PA or PEST is more challenging, since these systems require reactive compatibilization. Development of PPE grafted with acidic functionality was motivated by this need. PPE has been blended with most specialty resins — the latter usually as a minor component. Exceptions are the PPE/PPS alloys. Their performance depends on the level of PPS crystallinity. The commercial blends, DIC PPS (introduced in 1982) were developed for the electrical, electronic, and mechanical industry. They show good processability, reduced flash in injection molding, toughness, high heat and chemical resistance. Several blends of PPE with specialty resins are parts of multicomponent systems. For example, PI was blended with PPE, then cured. The alloys were used as rigid and stable matrices for the manufacturing fiber reinforced composites [Camargo et al., 1986]. Similarly, end-capped PPE was blended with SEBS, then dispersed in either PEI, PA, PEST, or PS, to give improved resistance to loss of impact strength after thermal recycling [Richards and White, 1994]. Blends of PAI with PPS and at least one of the following resins: PA, PEST, PC, PPE, PSF, PES, PEI, PEK, PEEK, PPS, PEST, PA, PEA, or siloxanes, were compatibilized with aromatic polyisocyanates. The systems showed excellent flowability, high heat resistance, and mechanical strength [Kawaki et al., 1995]. Examples of PPE/Specialty polymer blends are listed in Table 1.69.

1.5.6.1 PPE Blends 1.5.6.2 Miscible PPE Blends The first PPE blends with PS and polydiphenylsiloxane (PDPS) were reported to be transparent, colorless solids that turn into liquid at 85°C [Boldebuck, 1962]. Since then PPE has been

Miscibility of PPE with PS has been known since 1962. Later, two other PPE blends were announced miscible. PPE blends with polytrans-

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Introduction to Polymer Blends

Table 1.69. PPE specialty polymer blends Composition 1. PPS/PPE Blends PPS with PPE/PSF and/or PC, GF PPS with a styrene-grafted PPE and polyetherester rubber PPS/PPE with EPR-MA, EBA-MA, SGMA, EVAc-GMA, etc.

Comments

Reference

the first, reactively compatibilized blends good toughness, impact and tensile strength

Bailey, 1977 Dainippon Ink & Chem., Inc., 1982 Nishio et al., 1988, 1994

reactive blending, yielded reinforcing spherical PPE particles d = 0.01-10 µm

PPS/PPE/core-shell graft copolymer PPS/PPE with SEBS and PP

toughened by silicone elastomer particles tough blends

PPS + ABS, PPE, PC, PA, PEST PPSS with PPE, PC, PA or POM PPE/PPS & polymethylene-phenylene-isocyanate with GF PPE/PPS/core-shell MBS or SEBS

presence of macromers with epoxy group impact strength, mechanical properties mechanical & welding properties, solvent resistance PPE reacted with trimellitic anhydride acid chloride and dimethyl-n-butylamine impact resistance, stiffness, heat resistance, moldability, appearance processability, good mechanical properties

PPS/PPE/EBA-GMA blends with particle size d = 0.001-10 µm PPE with co-poly(arylene sulfide), [(-φ-S-)1-x(-φ-S-S-)x]n

Sasaki et al., 1989 Maruyama & Mizuno, 1990 Tsuda & Azuma, 1991 Ono et al., 1991 Gotoh & Nagaoka, 1993 Dekkers, 1994 Orikasa & Sakazume, 1994 Bagrodia et al., 1994

2. PSF/PPE Blends PSF, POM with PPE PC/PA-12 and PPE/PSF

high HDT moldability & mechanical performance

Ikeguchi & Nagata, 1974 Okuzono & Kifune, 1975-76

3. PI/PPE Blends PPE/SBR-grafted (SBMI) PPE/SBS, styrene-phenyl-maleimide PI with PPE, PPS, PEI or PSF PEI with PPE-MA

improved performance HDT, impact, and solvent resistance moldability, stability, mechanical strength mechanical performance

Two PPE’s with polyetherimide-silane copolymer

processability, flame retardant properties and impact strength

Fava & Doak, 1980 Fukuda & Kasahara, 1982 Ohta et al., 1988 White & van der Meer, 1989 Haaf, 1992

4. PEBA/PPE Blends PEBA with PPE and triglycidyl isocyanurate

5. PPE/Fluoro-polymer Blends PPE with poly(hexafluoro propyleneco-vinylidene fluoride) PPE with PVDF, SMMA and SEBS

non-delaminating behavior, good rigidity and strength

Brown et al., 1992

thermoformable, high HDT, and flame resistance improved impact strength and elongation

Snodgrass & Lauchlan, 1972 Van der Meer et al., 1989

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octanylene (PTO) can be processed at 260°C and have HDT • 194°C. The patents suggested that PTO is miscible with PPE [Jadamus et al., 1986, 1987]. The PPE blends with polyphenyl methacrylate, poly(p-methoxy phenyl methacrylate), poly(benzyl methacrylate), or poly(3-phenyl propyl methacrylate) were reported to have the lower critical solution temperature, LCST = 105-150°C [Fischer and Siol, 1993, 1995]. Immiscible, but transparent, are blends of PPE with a copolymer of cyclohexanedimethanol, ethylene glycol, and terephthalic acid (PCTG) — the clarity was achieved by closely matching the refractive indices at the use temperature. The alloys also show good processability, thermal dimensional stability, and economy [Stewart and Massa, 1993].

1.6

Specialty Polymers and their Blends

The specialty resins are expensive, produced in relatively small volumes either for a specific application or looking for a market niche. Their Tg > 200°C and modulus > 3 GPa. In 1991 the total world consumption of polysulfones (PSF) and polyethersulfones (PES) was 8.5 kton. Blends of the following polymers are known: polyfluorocarbons, polysiloxanes, sulfur-containing polymers (PPS, PPSS, PES, and PSF), polyetherketones (PEK, PEEK, PEKK), polyimides (PI, PEI, and PAI), PAr, COPO, polyphosphazene (PHZ) and LCP. 1.6.1

Fluorocarbon Polymers

Known in Germany since 1933, polytetrafluoroethylene (PTFE) is a semicrystalline resin (92-98% crystallinity), with Tm = 342°C and melt viscosity of η = 10 GPas. Others, more common fluoropolymers are: polychlorotrifluoroethylene (PCTFE), Hostaflon™ commercialized in 1934, fluorinated ethylene-propylene (FEP), Teflon™-FEP introduced in 1972, and numerous copolymers with Tm = 260 to 304°C, processable at Tprocess = 315 to 425°C, and having the degradation temperature, Tdeg = 425 to 440°C. The fluoropolymers are

characterized by stability at high temperatures, toughness and flexibility at low temperatures, low friction, insolubility and inertness to chemicals, low dielectric losses and high dielectric strength. The world consumption of fluoropolymers in 1991 was 72 kton. In blends, fluoropolymers are used in small quantities to enhance throughput, reduce the frictional properties and increase the wear resistance. Blends comprising 0.3-50 wt% of a low molecular weight PTFE (Tm ” 350°C) with engineering resin showed improved anti-friction properties [Asai et al., 1991]. PPS and PEEK blended with a fluoro(co)polymers and reinforced with either CF or GF were wear resistant with a short break-in period for forming a self lubricating film [Davies and Hatton, 1994]. Many commercial blends contain fluoropolymers (primarily PTFE) for the improved weatherability, wear and solvent resistance: SUPEC™ — “self-lubricating” blend of crystalline PPS with PTFE and 30 wt% GF, Lubricomp™ blends from LNP and similar RTP™ blends from RTP Co. (e.g., 15 wt% PTFE, 30 wt% GF and any of the following resins: ABS, PA, PEST, PC, PE, PEI, POM, PP, PPE, PPS, PS, PSF, PVDF, SAN, TPU, PEEK, PES, etc.), Sumiploy™ from Sumitomo Chem. Co., etc. [Utracki, 1994]. 1.6.2

Siloxane Polymers

Polysiloxanes, [-O-Si(RR’)-], are linear resins that can be branched or crosslinked into elastomers. They have high compressibility, permeability to gases, low Tg and viscosity, exceptional weatherability, low surface tension coefficient and are relatively expensive. Siloxane polymers or copolymers have been incorporated into engineering or specialty resins to improve processability, toughness, HDT, solvent and weather resistance. The main polymers of this type are: polydimethylsiloxane (PDMS), and polymethylphenylsiloxane (PMPhS). Their Tg = -127 and -86°C, respectively. They start oxidizing at 290 and 375°C and undergo structural rearrangement at 435 and 410°C. Polysiloxanes have been used as high temperature impact modifiers that improve

Introduction to Polymer Blends

the flame resistance, processability and optical properties. Several commercial blends are on the market, viz. Rimplast™ (high tensile, flexural, and Izod impact strength PA’s), Dialoy™ (PC/PET blends with good chemical, weather, and low-T impact resistance), etc. Evolution of polysiloxane blends with engineering and specialty resins is summarized in Table 1.70. 1.6.3

Polyarylene Sulfide (PPS)

Polyarylenesulfides (PPS), (-φ-S-)n, was commercialized in 1971 as Rayton™ R. The resin is semicrystalline with Tg = 194°C and Tm = 288°C, thus Tprocess • 290°C. PPS is difficult to mold — it tends to adhere to the mold surface and to flow into mold crevices. It has also relatively poor impact resistance. Blends have been developed to alleviate these problems, e.g., with 25 wt% of either PSF, PPE, or PC [Bailey, 1977]. Commercial PPS blends are available with PPE (e.g., DIC PPS commercialized in 1982, Noryl™ APS), PARA (RTP 1300), or PTFE (Lubricomp™ PPS). They show good processability with reduced flash, are tough, have excellent wear, as well as high heat, solvent, chemical, and oxidation resistance. PPS has been frequently blended with PSF’s. The latter resins are mainly amorphous, frequently transparent, with Tg = 196-288°C, able to maintain high performance characteristics over a wide temperature range, but poor weatherability, notched impact strength and ESCR. The PPS/PSF blends have been developed to improve PPS processability and/or the mechanical performance over a wide range of temperatures, to improve PSF weatherability, impact and ESCR characteristics. Evolution of these blends is outlined in Table 1.71. PAr’s are aromatic amorphous polyesters, viz. U-polymer™, Ardel™ D-100, Durel™, Arylon™, etc. Their Tg = 188°C and HDT = 120-175°C. Blends with PPS have been developed to improve the performance of PAr — processability, rigidity and hydrolytic stability. To the category of amorphous aromatic polyamides (PARA) belong polyphthalamides (PPhA), e.g., Amodel™ (Tg = 127°C, Tm = 310°C, HDT = 285°C, CUT = 180°C). PPS/PARA

81

blends were formulated to increase the reinforcing effects of GF on PPS. They show good processability, mechanical performance, and resistance to thermal degradation. Polyimides (PI) were introduced in 1962 as thermally non-processable Kapton™. To improve processability, the main-chain flexibility was enhanced by incorporating segments with higher mobility, viz. polyamide-imide (PAI), polyetherimide (PEI), polyimide-sulfone (PISO), etc. These polymers are characterized by high Tg = 150-420°C and thermal resistance. They are blended with PPS to enhance its moldability, thermal stability and mechanical performance. Polyaryletherketone, [-φ-CO-φ-O-φ-O]n, was commercialized in 1980 as Victrex™. It is a tough, resin with Tg = 143°C and Tm = 334°C. Blends of PEEK with PPS show synergistic effects toward tensile and flexural strength, as well as the impact resistance. The inorganic low temperature glasses (LTG) with Tg ” 300°C, are durable, and water resistant. LTG was blended with either PPS, PET, PBT, PEK, PEEK, PEI, LCP, PC, or fluorinated polymers [Frayer, 1993, 1994]. Oxidation of PPS by addition of N2O4 in a sulfuric acid solution, results in the incorporation of surfoxide groups, leading to either polyphenylenesulfidesulfoxide or polyphenylenesulfoxide. Their blends with high temperature resins (viz. PSF, PES, PPS, PEI, PAr, PEEK, PC, PI, PAI, LCP, fluoropolymers, cycloolefins, and their alloys or composites) produced high temperature resistant foams by heating for 5-60 min at T = 300-470°C [Scheckenbach et al., 1998]. The process reduced the moldings’ density by at least 50%. 1.6.4

Polysulfone (PSF)

Polyarylsulfones (PSF or PSU), or polyarylethersulfones (PAES) have the chain structure: (-φ-SO2-)n. The commercial resins include: Udel™ PSF, [-φ-C(CH3)2-φ-O-φ-SO2-φ-O-]n (Tg = 196°C and CUT = 160°C), Astrel™ [-φ-φ-SO2-φ-O-φ-SO2-]n (Tg = 288°C), Victrex™ PES [-φ-SO2-φ-O-]n (Tg = 228-232°C), Radel™ R PPSF [-φ-φ-O-φ-O-φ-SO 2 -φ-O-] n (T g = 220°C),

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Table 1.70. Polysiloxane blends Composition

Comments

Reference

1. PA Blends PA, vinyl-terminated PDMS, siloxane with Si-H groups & Pt-catalyst Acidified PPE, PA, PDMDPhS

tensile, flexural, and notched Izod impact strength flame-resistance

Arkles, 1983

2. PEST Blends PET with PDMS and MABS PEST and/or PC with siloxane-based vinyl-grafted copolymer PEST/PC/PPE/star-block copolymer PEST and siloxane-acrylic elastomer

impact and embrittlement resistance chemical, weather, and low temperature impact resistance impact modified engineering resins impact strength at low temperature

Sauers & Barth, 1970 Hongo et al., 1987

3. PC Blends PC with PDMS PC with siloxane & elastomer PC with PArSi PC or PEC with PC-b-PDMS PC with elastomeric polysiloxane/ polyvinyl-based graft copolymer

solution cast films impact resistance transparent, flame & impact resistant low-flammability and good impact strength thermally stable, low-T ductility, impact- and heat- resistance

Caird, 1961 De Boer & Heuschen, 1988 Jordan & Webb, 1992, 1994 Hoover, 1993 Derudder & Wang, 1993

4. POM Blends POM/PDMS adsorbed on silicone

for sliding parts with high wear resistance

Takayama et al., 1991

clear, transparent, colorless solids at 65°C enhanced solvent and impact resistance HDT, reduced melt viscosity flame retardancy

Boldebuck, 1962 Brown, 1989 Herrmann-Schoenherr & 1993, 1994 Jordan & Webb, 1994

5. PPE Blends PPE, PS, and PDPS PPE-g-siloxane, and SEBS Poly(bisphenol-A dimethyl-siloxane) with PPE, PAr, PI, PEST or PC PArSi with PPE and SBS

Smith et al., 1990, 1994

Hoxmeier, 1994 Yamamoto et al., 1992, 1994

6. PEI Blends PEI with poly(carbonate-b-siloxane), and EPDM, ABS, MBS or MMBA PEI with PArSi PArSi with PPE and SBS

processability, impact strength

Giles and White, 1983

processability, flame retardancy flame retardancy

Jordan and Webb, 1994 Jordan & Webb, 1994

7. PPS Blends PPS/PDMS, trialcoxysilane and PO PPS/PBT, silane and GF PPS with silicone and acrylate elastomer lattices PPS-g-amine with PDMS-g-epoxy

processability and impact strength chemical resistance and toughness improved heat- and impact- resistance improved tensile elongation and strength

Liang, 1987 Serizawa et al., 1992 Koshirai et al., 1992, 1994 Han, 1994

8. PHZ Blends PHZ or its copolymer with a siloxane polymer, and/or elastomer

bisazoformamide (foaming agent) gave semi-rigid, highly flame retardant foams

Dieck and Quinn, 1977

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Introduction to Polymer Blends

Table 1.71. PPS blends Composition

Comments

Reference

1. PPS/PSF PSF/ PPS with 45 parts of a styrene-butadiene rubber (SBR) PPS with PSF and 5-40 wt% PTFE PSF/PPS/ PSF-b-PPS copolymer PPS/PPSS { PPE, PC, PA, POM PPS with either PSF or PPSS 45-60 wt% PSF, 25-45 wt% PPS, and 0-10 wt% MBS

the first PPS/PSF blends — to improve the impact strength processability, and resistance to corrosives impact strength, uniformity, and cohesion impact strength and mechanical properties improved interfacial adhesion and moldability resistance to impact, high-T performance, weatherability — for car body panels

Asahi Chem. Ind. Co., Ltd., 1981

2. PPS/PAr PAr with 40 wt% PET and PPS PAr with 1-99 wt% PPS PAr/PPS, & chloro-hydro-dimethanodi-benzocyclo octene PAr/PPS, ABS or acrylic elastomer

enhancement of properties processability, impact, fire & abrasion resistance good hydrolytic stability, moldability, and flame retardancy excellent hydrolytic stability

Kyo et al., 1978 Matsunaga et al., 1978 Salee, 1980, 1981

Shue & Scoggins, 1981 Davies, 1990

PPS with 25-95 parts of either PA-66 or PA-MXD6 and Mg(OH)2

moldability, HDT and impact strength mechanical properties, adhesion: NH2 with GF, and aromatic parts of PPhA with PPS high degradation temperature, chemical resistance, HDT, mechanical properties high resistance to heat and thermal aging; improved melt flow excellent tensile strength, as well as arc, tracking, and heat resistance

4. PPS/PI PPS with 60 wt% PI PI with PPS, PPE, PEI or PSF PPS, 20-65 wt% PAI, and 4,4’-diphenyl methane diisocyanate

to improve moldability at 310°C to improve the moldability of PI processability, heat, chemical, and solvent resistance, mechanical strength

Alvarez, 1977 Ohta et al., 1988 Kawaki et al., 1992, 1994, 1995

processability, strength and impact resistance mold release and reduced molding flash rigidity, dimensional stability, strength

Robeson, 1982 Hindi et al., 1994 Frayer, 1993, 1994

3. PPS/PARA PARA with a small amount of PPS PPS with a small amount of PPhA and GF Reinforced blends of PPS with PPhA and POCA 95-5 wt% PARA with PPS

5. PPS with PEEK and LTG PPS with 10-90 wt% PEEK PPS/PEEK/PMP, GF, organosilane Low temperature glasses (LTG) with 35 wt% PPS

Ultrason™ E, Talpa™ 1000, Sumilite™, polyimidesulfone, PISO, Amoron™ polythioethersulfone, PTES, etc. The latest is the sinterable polyphenylenesulfone, Ceramer™ [-φ-SO2-φ-]n

Bailleux et al., 1984 Hashimoto, 1986 Ono et al., 1991 Bagrodia et al., 1993, 1994 Golovoy and Cheung, 1994

Salee, 1982

Chen and Sinclair, 1990 Yamamoto & Toyota, 1992 Dubois et al., 1993

(Tg = 360°C, Tdecomp. > 450°C), used as an “additive” to high performance polymers used in harsh environment [Ceramer, 1996].

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L. A. Utracki

PSF’s are transparent, flame resistant, have high strength, modulus and hardness, HDT > 200°C. They show excellent resistance to thermal and irradiation degradation, but are difficult to process (high melt viscosity), have low ESCR and poor weatherability. The latter properties can be improved by blending and/or reinforcing. PSF blends comprise high performance resins, viz. PPE, PPS, PTFE, etc., with such compatibilizers/impact modifiers as Phenoxy, EVAc-GMA [Gaafar, 1990], SMA-copolymers [Golovoy and Cheung, 1994], siloxane-polyarylene polyether copolymers, or high temperature MBS. Mindel™ A and Arylon™ are examples of the commercial ASA/PSF blends. They show good processability, toughens, plateability, heat and water resistance. Addition of a small amount of PSF to a variety of resins improves hardness, the notched Izod impact strength, plateability, hydrolytic stability, and shape retention at high temperatures. Many PSF blends of with engineering resins have been developed, viz. with PA, PEST, PC, PPE, or POM. They have high HDT, heat resistance, strength, stiffness, mechanical properties and ESCR. Polysulfone blends have been foamed using water and either N2 or CO2 [Bland and Conte, 1991]. The blend comprised at least two sulfone polymers, e.g., PES and PSF, and at least one non-sulfone polymer (e.g., PS, PPE, PEI, PC, PA, PEST, PP or PE). The nucleating agent was either talc, mica, silica, Zn-stearate, Al-stearate, TiO2, or ZnO. The foams were used as insulation for high temperature structural applications. Since in the preceding part PPS blends with PSF were described, in Table 1.72 examples of PSF blends with other specialty resins are listed. 1.6.5

1.6.6

Polyimides (PI, PEI or PAI)

Polyimides (PI) have imide group, -R-N=(CO)2=R’-, in the main chain. Owing to a variety of possible R and R’ groups, their Tg = 180-420°C. To improve processability, flexible groups were incorporated into the main chain. Examples of blends of these resins are listed in Table 1.74. Polyamideimides (PAI) were obtained by polycondensation of imides with aromatic di-amines, [-N=(CO)2=φ-CO-NH-R-NH-CO-φ=(CO)2=N-]n (Tg = 275°C, HDT = 265-280°C). The resin has high tensile and impact strength from T = -190 to T = 260°C, dimensional stability, good dielectric properties, solvent and chemical resistance, flame retardancy, good UV stability and low outgassing in high vacuum. To improve processability, PAI was blended with PA, PSF, or PEST [Toray Industries, Inc., 1979]. Polyetherimide (PEI), [-N(CO)2-φ-O-φ-C(CH3)2φ-O-φ(CΟ)2CN-φ−]n (Tg = 215-220°C, HDT = 217°C, CUT = 170°C, no weight loss at T ” 400°C) was commercialized as Ultem™. The resin has high tensile modulus (even at elevated temperatures), approaching that of many glass-reinforced resins. Commercial PEI blends include these with PC (Ultem™ LTX introduced in 1990) or with PPS (Supec™ CTX530) [Utracki, 1994]. Polyimidesulfone (PISO) was introduced in 1986 Celazole™ (Tg = 249-349°C). This transparent resin with flexural modulus of 4.8 GPa and tensile strength of 63 MPa, shown excellent solvent and creep resistance. 1.6.7

Aromatic Amorphous Polyamides (PARA)

Polyetheretherketone (PEEK)

Polyaryletherketones (PAEK) are aromatic polymers with ether and ketone linkages in the chain, viz. PEK, PEEK, PEEKK, etc. Polyetheretherketone (Victrex™ PEEK), [-φ-CO-φ-O-φ-O-]n, was commercialized in 1980 (Tg = 143°C, Tm = 334°C). Commercial blends of PEEK include, Sumiploy™ PEEK/PES/PTFE, PEEK/LCP, Cortem™ PEEK/ LTG, etc. Evolution of PEEK blends’ technology is outlined in Table 1.73.

There is a great diversity of amorphous aromatic or semi-aromatic polyamides (PARA). The commercial resins include Trogamid™, Quiana™, Amodel™, etc. They have been blended to improve the mechanical properties and impact strength, as well as to enhance the barrier properties of the matrix resin to permeation by gases or liquids. Examples of blends with PARA are listed in Table 1.75. Blends of PARA were also discussed along other blends of polyamides.

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Introduction to Polymer Blends

Table 1.72. PSF/PI blends Composition

Comments

Reference

improved melt flow, and good mechanical properties

Toray Industries, Inc., 1979, 1980, 1981

PES/PPBA (Tg = 200-300°C)

for lacquers & homogenous, clear films

Patton & LaMarre 1983

PEI with: PSF, PP, PEC or PAr; with PC and PEST, PAr or PA, etc.

for good processability, improved flexural and impact strength

Giles, 1983, 1984

PI with 0.01-10 phr PSF solution cast

for films with good blocking resistance

Mitsubishi Chem., 1984

PI with: PEI, PES, PAr, PC, PEEK, or PPE, e.g., PI:PEEK:PEI = 1:1:1

Tg = 175°C, used as crack-resistant coatings with good adhesion to metal

Camargo et al., 1986

1. PSF/PI PAI with 0.1-50 wt% of either PSF, PA or PARA

PI with: PSF, PPS, PPE, or PEI

moldability, heat & chemical resistance

Ohta et al., 1988

PES with PEI

improved HDT

Melquist, 1993

PES/sulfonyl bis(phthalic anhydride)-cobis(p-amino cumyl) benzene

processability, solubility, mechanical and thermal properties

El-Hibri & Melquist, 1993

LCP- type PI with PES, PI, PEI, PAI, PEK or PEEK

remarkably good processability and excellent thermal stability

Okawa et al., 1994

PES/ PI (e.g., XU-218, or PI-2080) miscible blends (single Tg, UCST)

high moduli, tensile strengths, and impact strengths

Karasz & MacKnight, 1994

PSF/PPS/5-40 wt% PTFE fibrils

processability, lubricity, anti-corrosive

Bailleux et al., 1984

PES/0.3-50 wt% low-MW PTFE

self-lubricity

Asai et al., 1991

PSF/PC or PET, ZnBO3 and PTFE

flame-retardant, synergistic properties

Jack et al., 1993

Fluorine-containing polycyanurates with PSF, PP or PEEK

flame retardant, low thermal expansion, Tg = 180-320°C, stable to 430-500°C

Ardakani et al., 1994

PES, with poly(p-phenyleneether-cop-phenylenesulfonyl)

miscible, transparent solvent-cast films, with good water and chemical resistance

Newton, 1981

PSF with 70 wt% polyether-amide (PEA)

moldability, high HDT and mechanical properties

Hitachi Chemical Co., Ltd., 1983

2. PSF Blends with Fluoropolymers

3. PSF Blends with Other Specialty Resins

PSF/acrylic elastomer/ polyphosphates

thermal stability, flame retardancy, toughness

Schmidt, 1983

PEEK/PAES with HDT = 157°C

low warpage and shrinkage, rigidity, stress cracking, solvent, & impact resistance

Harris & Robeson, 1986, 1987

PES, PEEK and 20 wt% PEI

mechanical properties, and heat resistance

Rostami, 1987

PES/95-75 wt% PEEK

chemical and hot-water resistance

Tsumato et al., 1987

PAEK, PAE, or PPE, blended with LCP, PI,

processability (warp-free moldings),

Harris & Michno, 1988

PAES or PEST

mechanical properties, and high HDT

PES with Phenoxy

chemical resistance and tensile strength

Kraus et al., 1991

PSF with PVP or PEG, radiation crosslinked

for selectively permeable membranes or hollow fibers

Kobayashi & Tanaka, 1992

Biodegradable PLA with either PSF, PC, PI, PPE, etc.

to improve the thermal properties

Nemphos & Kharas, 1993

PES dissolved in oligooxybenzoyl acid, then polymerized to POBA

molecular composites, polymerizing while shearing

Tochioka, 1993

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L. A. Utracki

Table 1.73. PEEK blends with specialty resins Composition

Comments

Reference

PEEK/PAI and optionally with PPS

solvent resistance, hydrolytic stability

Harris & Gavula, 1986

95 - 75 wt% PEEK with PES

chemical and hot-water resistance

Tsumato et al., 1987

POM/10 phr of PEEK and/or PEI

wear resistance without loss of slipperiness

Suzuki & Nagahama, 1987

PEK/PAI and zinc sulfate hydrate

good moldability and high impact strength

Smyser & Brooks, 1990

LTG with either PEK, PEEK, PPS, PEI, LCP, PC, PET, PBT or fluorinated polymers

processability, mechanical properties, stiffness —Cortem™ Alloys with either LCP or with PEEK

Bahn et al., 1991

Crystalline & amorphous PEK with 35-2 wt% of a PAr

good flowability and processability

Falk & HerrmannSchoenherr, 1992

POM-b-PC with PES, PEEK, PA or PAN

film-forming thermoplastic polymeric alloys

Dhein et al., 1993

ASA, PC, PEST, PEC, PPE, PPS, PEEK, PES, PSF and/or PPE

toughened by 30-80 wt% elastomer, e.g., SEBS and core-shell graft copolymers

Niessner et al., 1994

PPS blends with PEEK

improved mold release and reduced flash

Hindi et al., 1994

PI, PAI, PSF, PEI, PES, PEEK, PPS or PPE and a polyether-b-polyamide or polyether-b-polyester

easy to mold blends, flexible and elastic, with excellent chemical and thermal resistance

Movak et al., 1994

PEKK/PEI = 70/30 w/w co-reacted through the terminal amine group of PEI and ketone one of PEKK

resulted compound with strain hardening was water foamable at T = 335 - 350°C, whereas neither PEKK nor PEI can be foamed

Brandom et al., 1997

1.6.8

Polyarylates (PAr)

These polyesters, [-O-φ-C(CH3)2-φ-CO2-φ-CO-]n (Tg = 188°C, and HDT = 120-175°C), were introduced in 1974. The commercial resins include U-polymer™, Ardel™, Durel™, and Arylon™. Their advantages include transparency, good weatherability and high HDT. PAr has been blended with nearly all resins, including: ABS, EPDM, Ionomers, LCP, PA, PB, PBI, PBT, PC, PEI, PEK, PET, Phenoxy, PMB, PS, PPE, PPS, etc. Three types of PAr blends are of particular importance — those with polyesters, PEST, polyamides, PA, and with polyphenylenesulfide, PPS. A summary of PAr blends is provided in Table 1.76.

1.6.9

Aliphatic Polyketone (COPO)

This copolymer of carbon monoxide with ethylene and propylene is semicrystalline, with Tg = 15-20°C, Tm = 110-242°C [Ballauf et al., 1941]. Carilon™ resin (introduced in 1995) is a strictly alternating copolymer, [-CO-C2H4-]n, obtained using metallocene catalyst. It has Tm = 220°C, tensile strength σ = 80 MPa, and elongation at break ε = 25%. The moldings have outstanding wear and friction resistance, high resilience over a wide temperature range, low sensitivity to water and organic solvents, good barrier properties, but they are sensitive to UV. Several blends of COPO have been patented, e.g., with SAN (miscible blends), PA-6 and SEBS-MA [Machado, 1992], with TPU [George, 1992], with POM and PVPh [Machado, 1993].

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Introduction to Polymer Blends

Table 1.74. PI blends Composition

Comments

Reference

PEI with PAr PEI/PAI = 1:1 Polyarylethers with PEI PEI, poly(carbonate-b-siloxane), and EPDM, ABS, MBS or MMBA PEI, PA-6 and PEI-b-PA PEI with polyestercarbonate (PEC) PEI/0.5-20 wt% of a fluoropolymer PI with PAI in the full range of composition

moldability and mechanical properties mechanical properties and ESCR good ESCR processability and impact strength

Holub & Mellinger, 1981 Maresca et al., 1981 Robeson et al., 1981 Giles & White, 1983

moldability and impact strength high HDT & tensile strength mold release, heat resistance and shrinkage foamed during the final stage of the condensation reaction at T = 120 - 320°C moldability, heat stability, chemical resistance, and mechanical strength mechanical performance processability, mechanical strength, as well as heat and flame resistance high thermal deformation resistance, and HDT

Robeson & Matzner, 1984 Quinn, 1984 Sumitomo Chem., 1985 Long and Gagliani, 1986

high modulus, mechanical performance impact-resistant materials with excellent processability and HDT high flow and no flash moldings: flexible, elastic, self-lubricating, having excellent chemical and thermal resistance remarkably good processability and excellent thermal stability thermo-oxidative stability

Bahn et al., 1991 Durfee & Rock, 1993

PI with PPS, PPE, PEI, and PSF PEI with PPE-MA PEK with either PES, PEI, PEEK, PEST, PAr or PPS and filler PEI, PBT and triallyl cyanurate and triallyl isocyanurate Low temperature glasses with PEI Polyether-b-polyimide-b-siloxane copolymer with low MW PEI PPS/PEI with 30 wt% GF Fluoro-elastomers dispersed in a resin, e.g., PI, PAI, PSF, PEI, PES, PEEK, PPS, PPE, etc. LCP-PI with 99.9-50 wt% of either PI, PEI, PAI, PES, PEK or PEEK PBI with 0-95 wt% of PEI

1.6.10

Blends with Rigid-rod Polymers

Three types of blends belong to this group: (1) molecular composites, i.e., the molecular LCP solutions, (2) immiscible blends of LCP, and (3) blends of electro-conductive polymers. 1.6.10.1 Molecular Composites In fiber-reinforced composites, the absolute size of the reinforcing fibers is not important, but good adhesion to matrix and the length-to-diameter

Ohta et al., 1988 White & van der Meer, 1989 Murakami et al., 1991 Hosoi, 1991

Supec™ CTX530 Movak et al., 1994

Okawa et al., 1994 Haider & Chenevey, 1994

ratio of the fiber, L/D • 500 are [Piggott, 1986]. Accordingly, reduction of the reinforcing particle size from, e.g., GF or CF, to rigid-rod molecules seems desirable. If the reinforcement is to be provided by individual macromolecules, the rigid-rod polymer must form molecular solution in selected thermoplastic resin. Such systems are known as molecular composites, MC, first generated in the late 1970’s (see Table 1.77). MC can be prepared by dissolution of either the rigid-rod polymer in a monomer that subsequently can be polymerized, or by dissolving

88

L. A. Utracki

Table 1.75. Blends with aromatic amorphous polyamides (PARA) Composition

Comments

Reference

PARA with POM PARA with semicrystalline PA PAI with PA-66 or PARA PARA with PC PARA with 5-95 wt% PPS PARA with PA-6I6T PARA with PES PPS with PPhA and GF PARA with maleated PP or PE

toughness and impact strength improved oxygen barrier properties processability & mechanical properties nacreous, resistant to oils and boiling water improved moldability, HDT and impact strength for strong fibers or films increased HDT & mechanical performance stiff, high performance reinforced alloys sliding electrical parts, resistance to thermal degradation in contact with Cu transparent, having near-zero birefringence processability, mechanical properties, solvent, weather, HDT, impact and stress-crack resistance

Asahi Chem. Ind., 1969 Dynamit Nobel, 1969 Toray Ind., 1979, 1981 Mitsubishi Chem., 1980 Shue & Scoggins, 1981 Unitika Co., Ltd., 1982 Hub et al., 1986 Davies, 1990 Iwanami et al., 1990

PARA with rubber-modified PS PARA with PAr are miscible

Angeli & Maresca, 1990 Bapat et al., 1992

Table 1.76. Blends with linear, aromatic polyesters, PAr Composition

Comments

Reference

PAr with PC and PET PAr blends with PET

improved impact resistance transparent and impact resistant commercial blends

Koshimo, 1973 U 8000 from Unitika or Ardel™ D-240 from Amoco Asahara et al., 1977 Matsunaga et al., 1978

PAr/PET = 1:1 with PA-6 or PPS PAr with 1-99 wt% PPS

enhanced mechanical properties processability, impact strength, fire and abrasion resistance PAr/PPS and dodecachloro-dodecagood hydrolytic stability, moldability, hydro-dimethano-di-benzocyclooctene and flame retardancy PAr/PET/PPS with ABS or MBS improved hydrolytic stability PAr with PA-6, U-polymer™ X-9 processability, mechanical properties PAr with polybenzimidazole (PBI) miscible blends

monomer of the rigid-rod polymer in a thermoplastic resin, then polymerizing it. The selections of the soluble monomer/polymer pair, as well as control of the polymerization and phase separation rates are critical. The IPN approach may “lock” the dispersed structure into a metastable system with sufficient stability for processing [Utracki, 1994a].

Salee, 1980, 1981 Salee, 1982 Unitika, Ltd., 1982; 1983 Chen et al., 1990

1.6.10.2 Liquid Crystal Polymers (LCP) There are several commercial LCP’s, viz. Ekkcel™ (degrades at Tprocess = 400°C), Xydar™, Vectra™, E-konol™, X-7G, Ultrax™, KU 1-90, Granlar™, Novoaccurate™, Rodron™, Victrex™, etc. LCP’s are mainly used for injection molding of parts that require exact dimensions and high performance.

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Introduction to Polymer Blends

Table 1.77. Molecular composites Composition

Comments

Reference

Poly(p-phenylenebenzobisthiazole) with poly(2,5,(6’)-benzimidazole) Poly(p-phenylene terephthalamide) (PPTA) with PA-6 or PA-66 in methanesulfonic acid Poly(2,5,(6’)-benzimidazole) with PAr Rigid-rod [-CO-φ(CF 3)-φ(CF3)-CONH-φ(CF3)-φ(CF3)-NH-]

MC of oriented macromolecules; films and fibers had high modulus and strength coagulation resulted in MC that upon thermal treatment phase separated MC miscible system

Hwang et al., 1983

polymerization of the monomeric solvent in vinylpyridine or pyrrolidone resulted in MC Polybenzimidazoles, 5-95 wt% with PARA blended miscible by: film transparency, single Tg, in DMF FTIR, and X-ray diffraction PES dissolved in polyoxybenzoyl or p-aceto-aminobenzoic solvent-free MC; high modulus and strength; acid, polymerizing the latter at the shear rate of 2.0-13 1/s for the manufacture of fibers, or rod-shaped extrudates 70-99 wt% PGI, and either PBI or LCP, with PET or PC good balance of toughness, tensile modulus and HDT Dissolution of PA or PO in lactams, then polymerizing it MC: N-(p-amino benzoyl)caprolactam into a rigid-rod, e.g., poly(p-aminoethyl benzoate) (PAEB) in molten maleated-PP, PA-6, PA-66, or PARA, then polymerized; PA modulus x2 Poly[2,2’-(m-phenylene)-5,5’-bis benzimidazole] mechanical, thermal & chemical properties; (PBI) with 0-65 wt% PEEK for gaskets, seals, valve seats and O-rings

Large quantity of LCP is used in blends. These are immiscible, highly oriented systems, where LCP domains provide reinforcement. Since LCP’s are expensive, either the desired performance must be achieved using a small amount of melt processable LCP, or the other component of the blend is similarly priced. In blends LCP can: (i) improve processability of engineering and specialty polymer [Froix, 1981; Cogswell et al., 1981, 1983, 1984], (ii) enhance crystallization of semicrystalline polymers [Hong et al., 1992], (iii) improve stiffness and other mechanical properties in applications where fatigue strength is important [Yamauchi et al., 1991], provide external protective layer for solvent and/or abrasion sensitive resins, etc. Excepting those with PP, the LCP blends with commodity resins are scarce (see Table 1.78).

Chuah et al., 1989 Chen et al., 1990 Stein et al., 1992 Calundann et al., 1992; 1994 Tochioka, 1993

Hallden-Abberton et al., 1994 Mülhaupt et al., 1994 Andres et al., 1995

1.6.10.3 Electro-Dissipative and -Conductive Blends Most organic polymers are insulators. However, there are applications requiring dissipation of the electrostatic charge (ESD) or even electrical conductivity (ECP) that would be comparable to that of metals. The ESD materials should have the surface resistivity 1012 • R • 105 Ω cm. The resistivity of ECP should be: 105 • R • 10-2 Ω cm. The ESD behavior can be provided by blending in a flexible-chain polymer with an active -OH or –SH group, viz. polyvinyl alcohol (PVAl), ethylene-vinylacetate (EVAc), polyvinylphenol (PVPh), a copolymer of ethylene oxide and epichlorohydrin (EO-CHR), maleated copolymer, aliphatic polysulfides, etc. These low performance resins have been incorporated into a variety of alloys and blends (see Table 1.79).

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Table 1.78. Liquid crystal polymer blends Composition 1. LCP blends with Commodity resins PP/LCP with PP-MW as a compatibilizer PP/LCP PP/LCP LCP/PE LCP/cycloolefins (COP) LCP & radiation crosslinkable resins 2. LCP blends with Engineering resins Poly(1,4-benzamide) or terephthalamide dispersed in PA PET with 2 phr of poly[bis(4-methoxy phenyl)terephthalate] 30 wt% PET with LCP

Comments

Reference

LCP macromolecules stretched by simultaneous flow and crystallization in a static mixer LCP macromolecules stretched in a counter-rotating pipe die LCP macromolecules stretched under high injection molding stresses viscosity reduction processability for articles that strain-recover upon heating

Baird & Datta, 1992

rigid microfibrils enhanced modulus and improved the thermal behavior processability, and excellent mechanical properties

processability, mechanical properties, HDT = 167°C PC with a wholly aromatic polyester, LCP mechanical, tensile and flexural, properties PPE or PAEK with LCP, PI, PAES or PEST processability, mechanical properties, and HDT PS/PPE/2-98 wt% of a LCP, stretched tensile strength, high modulus, satisfactory into sub-microscopic fibers elongation, good impact strength, and high HDT PBT with p-hydroxybenzoic co-reaction to increase η, thus orientability and acid-ethyleneterephthalate (Tm ” 300°C) PC with PAr and LCP LCP with PP, PS, PC, PI, etc. LCP with PC and PET or PBT PBT with segmented block copolymer, LCP-b-PBT PEI-LCP with PC, PBT or PA LCP block copolymer of the type [rod]x-[coil]y- with PET, PBT, PA LCP dispersed in either PEST, PC, PA or modified PPE LCP with: PA, ABS, PC, PBT, PPE, PP, PC, or their blends LCP with 3-15 wt% PAr

processability, and mechanical behavior for spinning fibers with high mechanical properties and low shrinkage replacements for fiber reinforced plastics — recyclable blends compatibilized blends, used as replacement for glass fiber composites matrices for conventional composites

LCP with PEST, PC, PA’s, PI, etc. Poly(p-phenylene terephthalamide)/ PA or PARA, and PEKK or PAN

polymerization of LCP in polymeric matrix biphasic solution in sulfuric acid, spun, coagulated, stretched into PPD-T fibrils

mechanical performance low anisotropy, high HDT, heat and impact resistance multi-axial molecular orientation of LCP ductility, toughness, strength, HDT modulus for fibers having high modulus and strength

Haghighat et al., 1992 Heino et al., 1993 Alder et al., 1993 Epple et al., 1992 Toy et al., 1994

Takayanagi et al., 1980 Toray Industries, Inc., 1980 Celanese Corp., 1981, 1984 Froix, 1981 Harris & Michno, 1988 Isayev, 1991, 1993, 1994 Dekkers et al., 1992 Izumi et al., 1992 Haghighat et al., 1992 Cottis et al., 1993 Farris & Jo, 1993 Bonfanti et al., 1993 Dashevsky et al., 1993, 1994 Tomita et al., 1993, 1994 Tomita et al., 1993, 1994 Roemer & Schleicher, 1993, 1994 Gupta et al., 1994 Coburn & Yang, 1994

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Introduction to Polymer Blends

Table 1.78. Continued Composition

Comments

Reference

Hydroxyalkylated PPE, hydroxyl-containing PO, PA and LCP Compatibilized blends of PET with 10-15 wt% LCP

moldability, solvent and heat resistance, mechanical strength processing, non-transparent material with good mechanical properties

Arashiro et al., 1994

improved flowability and processability processability, and physical properties LCP fibers, tensile strength, modulus, elongation, processability, good mechanical properties, and high HDT oriented films with small protrusions that resulted in low friction processability, thermal stability, and mechanical properties Cortem™ Alloys: matrix LCP & ” 80 wt% of dispersed LTG; E = 14 GPa high tensile strength

Cogswell et al., 1981 Froix, 1981 Isayev & Swaminathan, 1989 Harris et al., 1988

3. LCP blends with Specialty resins PES with a small amount of LCP PPS, LC-polyesters & LC-poly(esteramides) PEI with 35-95 wt% LCP, LCP with PI, PAES or PEST and either PAEK, PAE or PPE • 0.01 wt% LCP with: PET, PA, PC, PE, PP, PVC, PVDC, PPS, PVDF, PVF LCP with either a Phenoxy or an esteramide-based LC LTG/LCP Amino-terminated PEI with polyester-type LCP LCP with 2-98 wt% PEK PPS with polyesteramide-type LCP LCP from p-hydroxybenzoic and 2,6-hydroxynaphthoic acids, with a non-thermotropic polymer, silane PEI with 5-95 wt% LCP and p-amino benzoic acid or pyromellitic anhydride 99.5-50 wt% PI with polyimide-type LCP, Tm • 300°C PI with PEK and/or LCP and other additives; Aurum™ PI/LCP LC-type polyimide with either: PI, PEI, PAI, PES, or PEK Polyglutarimide with PBI, or LCP, and PET or PC PAI, with 3-30 wt% LCP PPS with a polymer of: 6-hydroxy-2-naphthoic acid and 4-hydroxybenzoic acid, or terephthalic acid and 4-amino phenol LCP blended with another LCP

toughness, excellent elasticity, and impact strength accelerated crystallization rate of PPS good phase morphology, interfacial adhesion, good thermal and mechanical behavior compatibilized moldable blends, useful as matrix for composites. processability, chemical resistance, flame retardancy

Bonis & Adur, 1995

Wong, C. P., 1990, 1994 Koning et al., 1990 Bahn et al., 1991 Bookbinder & Sybert, 1992 Falk & Hermann-Schoenherr, 1992 Minkova et al., 1992 Haider et al., 1993

Roemer & Schleicher, 1993 Asanuma et al., 1993

and mechanical strength processability, HDT • 230°C, strength, thermal and Tsutsumi et al., 1994 chemical resistance remarkably good processability and excellent Okawa et al., 1994 thermal stability good balance of toughness, tensile modulus and HDT Hallden-Abberton et al., 1994 lower viscosity, unaffected mechanical properties processability and properties; used to mold parts for the electronic industry, particularly connectors

Lai et al., 1994 Yung & Linstid, 1995

processability, HDT • 200°C, impact strength

Charbonneau et al., 1995

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Table 1.79. Electro dissipative and conductive blends Composition

Comments

1. Electro Dissipative Blends — ESD systems Aliphatic polysulfides (TM) to mold static charge-free rolls and guides with polybutadienes (PB) for textile industry PO with 2 wt% PVAl mechanical, hygroscopic, antistatic properties high tensile strength, good impact resistance, PC, with PET and • 1 wt% of an elastomer containing acidic groups and electrical conductivity ABS and ” 20 wt% EO-CHR antistatic thermoplastic compositions EO-CHR with ABS, HIPS, MBS, SMA rapid dissipation of static charge, reduced or PS/PPE and an acrylic (co)polymer delamination and improved ductility PVC, CPVC, PC, PEST, EP, PF, antistatic properties or styrenics with EO-CHR PC/PAI and a C2-10 diamine processability, impact strength, appearance, mechanical, and antistatic properties PS with EO-CHR and PCL static dissipative and tensile elongation 2. Electro Conductive Blends — ECP systems PAc was polymerized into PE PE with catalyst exposed to acetylene Polypyrrole electrochemically electrically conducting material with improved polymerized within a matrix resin mechanical properties over those of PPy PVC with “doped” PANI and an additive intrinsically electrically conductive films or fibers Amine-terminated PANI materials were suitable for the use melt-blended with SMA as electric conductors Polyaniline tosylate (PANI), and PETG for films, inks, fibers, and coatings, in shielding, antistatic and adhesives Poly-3-octyl thiophene with PP, PVC, PS, blends were formed into desired shapes, and PE, EVAc, PVC/ABS, etc., and dopant, used either as EMI or ESD materials e.g., I2 PANI with dodecylbenzene sulfonic acid soluble thermoplastic ECPB’s could be heat treated, then mixed with either modified by mixing with protonic PS, PE, PS, ABS or PP acid and metallic salts Fluorine-containing polycyanurates and heat- or electrically conducting materials, a thermoplastic polymer, for electronic packaging, adhesives, e.g., PSF, PPE, PEEK in the fabrication of electronic parts PANI or PPy with polymeric dopant: electrically conductive polymeric systems sulfonated:-PE, -SEBS, -PS, etc. with good mechanical properties Matrix polymer and an electrically matrix: PO, EPR, CPE, CSR, PS; dispersed: conducting thermotropic PANI, PAc, PPy, poly(3-undecylthiophene), liquid-crystal polymer, LCP poly(3-dodecylthiophene), or polyparaphenylene

Reference

Patric, 1942 Minekawa et al., 1969 Mitsubishi Chemical Industries Co. Ltd., 1983 Federl & Kipouras, 1986 Gaggar et al., 1988; 1989 Barnhouse & Yu, 1988; Yu, 1988 Shimamura & Suzuki, 1991 Giles & Vilasagar, 1994

Galvin & Wnek, 1982 Lindsey & Street, 1985 Kulkarni & Wessling, 1992-1994 Jongeling, 1993 Shacklette et al., 1993 Kokkonen et al., 1994

Karna et al., 1994

Ardakani et al., 1994

Cross & Lines, 1995 Ho & Levon, 1995

Introduction to Polymer Blends

By contrast, the ECP must have conjugated rigid-rod macromolecules. Several such polymers show high electrical conductivity (usually after doping), viz. polyacetylene (PAc), polyaniline (PANI), polypyrrole (PPy), polyparaphenylenes (PPP), or poly-3-octyl thiophene (POT). The resins are expensive, difficult to process, brittle and affected by ambient moisture, thus blending is desirable. For uniaxially stretched fibers the percolation threshold is 1.8 vol%, hence low concentration of ECP (usually 5-6 vol%) provides sufficient phase co-continuity to ascertain conductivity similar to that of copper wires (see Table 1.79).

1.7

Biodegradable Blends

Biodegradability has been explored in agriculture to prevent excessive moisture loss and weeds growth, and to alleviate the recyclability problems — an agricultural film should last as long as it is needed, then disintegrate under the influence of either microorganisms and/or UV irradiation. Most biopolymers are biodegradable, e.g., a large family of polysaccharides. They have been used in biodegradable blends with synthetic polymers. Some synthetic polymers, viz. PET, is susceptible to biodegradation when copolymerized with polylactones. Polymers with controlled, reversed miscibility, viz. polyglycoles, are also biodegradable. Polymers with carbon backbones, viz. PE or PP may be susceptible to biodegradation after incorporation of ketone side groups, -C(R)(COR’)- [Guillet, 1973]. Biodegradable polybutylenesuccinate or adipate, Bionolle™, has been commercially introduced in 1996 by Showa Denko. Similarly, Novamont introduced fully biodegradable Mater-Bi™. The latter materials are blends of starch and other polymers, viz. poly-ε-caprolactone, ethylene-vinyl alcohol, etc. The biodegradable polymer blends are often prepared by blending a thermoplastic resin with a biodegradable one (see Table 1.80). Blending must produce dispersion that after disintegration of the biodegradable part the thermoplastic powder will not contaminate the environment.

1.8

93

Blending and Recycling

Recycling is becoming increasingly important. Its methods depend on the polymer type and source. Within the resin manufacturers’ plants recycling is the easiest. In processing plants, where commingled polymeric scrap is generated, it is more difficult. The most difficult is recycling of the post-consumer polymers. Three basic methods of recycling have been used: (i) direct, where cleaned resins are incorporated into virgin material, (ii) reprocessing the commingled plastics either by blending, or transforming into plastic wood or plastic concrete, and (iii) feedstock type that may involve depolymerization or pyrolysis. To the following text, only the method (ii) is important. It can be subdivided into: (1) Compatibilization and upgrading of resins in direct recycling, (2) Compatibilization and upgrading of commingled plastics for reprocessing, and (3) Recycling of polymer blends [Akovali et al., 1998]. To compatibilize a multicomponent blend one may: (i) add at least one ingredient with highly reactive groups that can interact with several polymeric components, e.g., ethylene-acrylate-maleic anhydride, glycidylmethacrylate-ethylene-vinylacetate, ethylene-glycidylmethacrylate-methyl methacrylate (the copolymer may compatibilize and toughen), (ii) add a low molecular weight additive that at different stages of the reactive blending binds to different components, viz., ethylene-glycidylmethacrylate, triglycidylisocyanurate, etc. (iii) add a co-solvent, for example, Phenoxy. A significant penetration of properly designed copolymer into the homopolymer phases has been reported [Brown, 1989]. The morphology can be stabilized by: (i) thick interphase, (ii) partial crosslinking, or (iii) addition of an immiscible polymer with a suitable spreading coefficient [Yeung et al., 1994]. The adhesion between the phases in the solid state is improved by: (i) addition of a copolymer that covalently bonds the phases, (ii) reduction of size of the crystalline domains, (iii) adequate adhesion, e.g., by the use of polyetherimine, PEIm [Bjoerkengren and Joensson, 1980], and (iv) dispersing at high stresses, either in the melt

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Table 1.80. Biodegradable polymer blends Composition

Comments

Reference

PVAl/ vinylacetate grafted starch PHBA, with 10-40 wt% CPE PLA/PEO, EVAc, EVAl, EPDM, SBR, etc. EVAl/poly(hydroxybutyrate-valerate) PS, PO, NR, SBR, PI, PB or CA, a polysaccharide and bioagent LDPE, starch and a copolyacrylate Maleated starch, PE or PP and 1-35 wt% acrylic copolymer Amylose/PA/PEST/POM/gelatins Starch, latex of either polymer or elastomer and 0-20 phr fillers Starch/poly [unsaturated fatty acids + diamines + diol-based glycols] PO/PHB/A-B block copolymers, of poly(meth)acrylic esters Plasticized polylactic or a lactic acid-hydroxycarboxylic acid copolymer PLA/PC or PSF, PI, PPE, siloxanes, silicones, PMMA, etc. Hydroxypropyl-starch or urea-starch, and either PA, PEST, etc. PS, PE, PP, TPU, PEST, PA, etc., with 5-99 wt% of either carbohydrates, proteins or lipids A polar polymer, polysaccharide and fatty acid (hydroxy) peroxide Synthetic polymer, peracid & starch LLDPE with starch and •1 ionic compound Non-consumable agricultural products with an adhesive PEG with: PA, PE-co-acrylic or methacrylic acid, EVAc, EVAl TPU and/or Phenoxy, EVAl, COPO, cellulose and/or polyalkylene oxide

biodegradable, better properties that PVAl biodegradability, impact properties and HDT biodegradable, flexible alloys biodegradability, and good impact properties cellulose with 1 wt% of bacteria, fungi and/or enzymes biodegradable blends biodegradable films with good mechanical properties for manufacturing transparent packaging films the mixture could be molded or extruded to form parts useful for food packaging the blends were used to manufacture packaging films or moldings for disposable napkins, ostomy bags, and ordinary wrapping flexible and hydrolyzable materials, useful for absorbing oils and body fluids improved HDT of biodegradable polymers

Yoshitake et al., 1978 Holmes et al., 1982 Kharas & Nemphos, 1992 Webb et al., 1992 Guttag, 1992, 1994

Starch with biologically degradable aliphatic polyesters, hydrophobic protein, PVAl, or cellulose esters and a hygroscopic material

for the manufacture of printable moldings or films reactively blended biodegradable inter-polymers, with good mechanical properties, and limited water absorption good performance until exposed to suitable environment for degradation superior mechanical properties, biodegradability for high frequency sealable multilayer packaging films, biodegradation biodegradable tableware from impact-molded, coated particles degradable and/or recyclable plastic articles with inverse solubility characteristic attractive physical, optical and barrier properties, and were melt-processable without PVAl degradation absorbed water was released during either extrusion or injection molding at T = 200°C, causing the mixture to foam to density ρ = 160 kg/m3

Willett, 1992 Tomka, 1992; Tomka et al., 1993 Meier, 1993 Munk, 1993 Ritter et al., 1993

Ballard & Buckmann, 1993 Morita et al., 1993 Nemphos & Kharas, 1993, 1994 Buehler et al., 1993, 1994 Vaidya & Bhattacharya, 1994 Chapman & Downie, 1994 Hsu et al., 1994 Dehennau et al., 1994 Liebermann, 1994 Petcavich, 1994 La Fleur et al., 1994

Tomka, 1998

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Introduction to Polymer Blends

[Patfoort, 1976] or in solid state [Shaw, 1993; Khait, 1994, 1995]. Chemical re-stabilization of recycled material against the thermal and light-induced degradation is essential. Addition of 0.1-0.5 wt% of a sterically hindered phenol and a phosphite at a ratio varying from 10:1 to 1:10 is recommended [Pauquet et al., 1994]. For outdoor applications, hindered amine light stabilizers with UV-absorbers of the benzotriazole type are to be used [Herbst et al., 1995, 1998]. Examples of blends used for polymer recycling are listed in Table 1.81. For more details see Chapter 16 in this book.

1.9

Conclusions and Outlook

Blends like composites are integral parts of the plastics industry. Their sales are estimated at more than US$ 100 billion per annum. The blends

provide widening selection of performance characteristics, tunable for specific applications at a reasonable cost. In effect, this technology is a shortcut to development of complex polymeric species. However, blending requires a broad knowledge of many disciplines, from the thermodynamic principles of miscibility and compatibilization, to the surface and interphasial characteristics, morphology, rheology, processing, and performance. The inter-relation between these elements is of crucial importance, viz. flow affects the interfacial energy, miscibility, morphology, thus performance. Furthermore, most commercial blends are in a non-equilibrium state. The product performance depends on the selected processing method and process variables. As the fundamental knowledge improves, and the demand for higher material’s performance grows, the role of blending is becoming increasingly important.

Table 1.81. Polymer blends for recycling Composition

Comments

Reference

PS with 1-10 wt% PE and CSR

improved impact, elongation, and strength [CPE or CSR is a compatibilizer for scrap] blends with good mechanical properties SEBS is expensive but useful in recycling recycling requires higher concentration of stabilizers than virgin resins recyclable, impact resistant, good elongation and retention of physical properties high properties due to stable co-continuous morphology without compatibilization recyclable blends, good performance blending at T in between the melting points of the components cryogenically comminuting rubber, drying it, blending with plasticizer/binding agent, heating and blending with polymers compounding in a TSE, devolatilization, filtering of paint flakes; closed-loop control system for properties foaming and extruding the mixture

Herbing & Salyer, 1963

LLDPE/PS shear-compatibilized PS or HIPS, with PP and SEBS PS with PO, S-b-B, stabilizers sPS and sPS-copolymer with MA or PO with 30-40 wt% PS recycled PA with PA/LDPE, EGMA • 2 PO, PS, polydienes — either virgin, recycled or both Rubber scrap with rosin and fatty acids, esters and unsaponifiables), and PE, PET, TPU, PU, PVC, etc. Automotive scrap plastics parts comprising PC, PEST, ABS, PA, etc., with 5-15 wt% MBS Branched PET with recycled PET and a chain extender or a crosslinking agent

Patfoort, 1976 Holden & Gouw, 1979 Sadrmohaghegh et al., 1985 Okada, 1994 Morrow et al., 1994 Timmermann et al., 1994 Lai & Edmondson, 1995 Segrest, 1995

Lieberman, 1995

Muschiatti & Smillie, 1995

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L. A. Utracki

Computational methods are being developed in the science and technology of polymer blends. Molecular dynamics provide means of computation of the specific interactions and miscibility [Coleman et al., 1991] or the interfacial energy in polymer blend [Yao and Kamei, 1995]. Commercial computer programs are available for designing blends with specific sets of properties. A computer program capable to predict variation of blends’ morphology inside a twin-screw extruder has been developed [Utracki and Shi, 1992; Shi and Utracki, 1992; Huneault et al., 1995]. Possibilities of modeling of the reactive compatibilization are being investigated. The ultimate program should combine three subroutines: flow mechanics, kinetics, and description of the morphology evolution.

1.10

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1979, to Teijin. Yamamoto, S., and Toyota, A., Europ. Pat. Appl., 489,437, 10 June 1992, to Mitsui Petrochem. Yamamoto, S., Honda, S., and Shimizu, H., Jap. Pat., 043,470, 23 Dec 1971, to Sekisui Chem. Yamauchi, M., Magome, T., Takahara, A., and Kajiyama, T., Reports Prog. Polym. Phys. Japan, 34, 229 (1991). Yamauchi, S., Yasuno, K., and Kitamura, S., Jap. Pat., 013,855, 03 Apr 1974, to Sumitomo Chem. Yamazaki, I., and Fujimaki, T., Jap. Pat., 007,627, 03 Mar 1970; 007,141, 29 Feb 1972, to Showa Denko. Yang, W. Y., Europ. Pat. Appl., 212,449, 04 Mar 1987, to B. F. Goodrich. Yao, S., and Kamei, E., Nihon Reor. Gakk., 23, 103 (1995). Yasue, K., Marutani, T., Fukushima, Y., and Ida, T., Europ. Pat. Appl., 301,663, 01 Feb 1989, to Stamicarbon. Yates, J. B., III, and Lee, G. F., Jr., Europ. Pat. Appl., 349,747, 10 Jan 1990, to General Electric Co. Yates, J. B., III, PCT Int. Appl., 005,311, 11 Sep 1987; Europ. Pat. Appl., 303,077, 15 Feb 1989; Europ. Pat. Appl., 550,208, 07 Jul 1993, to General Electric Co. Ye, Y.-C., La Mantia, F. P., Valenza, A., Citta, V., Pedretti, U., and Roggero, A., Europ. Polym. J., 27, 723 (1991). Yee, A. F., and Shi, Y., SPE Techn. Pap., 41, 2642 (1995). Yeo, J. K., Oh, J. M., Chang, S. K., and Lee, E. S., H., US Pat., 4,829.125, 09 May 1989, to Lucky. Yeung, C., Desai, R. C., and Noolandi, J., Macromolecules, 27, 55 (1994). Yoga, Y., Okamoto, I., and Watanabe, K., Ger. Offen., 3,313,442, 27 Oct 1983, to Daicel Chem. Yonemitsu, E., Sugio, A., Masu, M., Kawaki, T., and Sasaki, Y., Jap. Pat., 057,754, 20 May 1976, to Mitsubishi GasChem. Yonetani, K., Nakamura, S., and Inoue, S., Jap. Pat., 121,757; 121,760, 03 June 1987, to Toray. Yoshihara, Y., Europ. Pat. Appl., 397,531, 14 Nov 1990, to Mitsui Petrochem. Yoshitake, T., Tasaka, T., and Sato, R., Jap. Pat., 050,254; 050,264, 08 May 1978, to Kuraray. Yu, M. C. C., and Beever, W. H., Europ. Pat. Appl., 473,038, 04 Mar 1992, to Phillips. Yu, S. H. P., Europ. Pat. Appl., 282,985, 21 Sep 1988; Jap. Pat., 63 314,261, 22 Dec 1988, to B. F. Goodrich. Yui, H., Kakizaki, T., and Sano, H., Jap. Pat., 014,752, 09 Feb 1978, to Mitsubishi Petrochem.

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Yuichi, O., and Suehiro, S., Europ. Pat. Appl., 304,041, 22 Feb 1989, to Nippon Petrochem. Yukioka, S., and Inoue, T., Polymer, 34, 1256 (1993); ibid., 35, 1182 (1994). Yung, P C., and Linstid, H. C., III, US Pat., 5,418,281, 23 May 1995, to Hoechst Celanese Corp. Zeichner, G. R., and Patel, P. D., Second World Congress Chem. Eng., Montreal, Que., Canada, 6, 333 (1981). Zeitler, G., Mueller-Tamm, H., US Pat., 4,020.025, 26 Apr 1977, to BASF. Zimmerman, J., Pearce, E. M., Miller, I. K., Muzzio, J. A., Epstein, I. G., and Hosegood, E. A., J. Appl. Polym. Sci., 17, 849 (1973). Zletz, A., US Pat., 2,692,257, 19 Oct 1954, to Standard Oil of Indiana.

CHAPTER 2

THERMODYNAMICS OF POLYMER BLENDS

L. A. Utracki

National Research Council Canada, Industrial Materials Institute, Boucherville, QC, Canada

2.1

Introduction

Performance of polymer blends depends on the properties of polymeric components, as well as how they are arranged in space. The spatial arrangement is controlled by the thermodynamics and flow-imposed morphology. The word “thermodynamics” invariably brings to mind “miscibility.” However, thermodynamics has a broader use for the practitioners of polymer science and technology than predicting miscibility. The aim of this chapter is to describe how to measure, interpret, and predict the thermodynamic properties of polymer blends, as well as where to find the required information and/or the numerical values. Determination of such thermodynamic properties as the phase diagram or the Huggins-Flory binary interaction parameter, χ12, is difficult. The difficulties originate in high viscosity of macromolecular species, thus slow diffusion toward the equilibrium, heat generation when mixing and dangers of degradation. For these reasons, there is a tendency to use low molecular homologues or solutions. Furthermore, it is an accepted practice to purify the polymers before measuring their thermodynamic properties. However, the industrial polymers have high molecular weights, and are modified by incorporating low molecular weight additives. Furthermore they are processed under high flow rates and stresses that preclude the possibility of thermodynamic equilibrium. For these and other reasons, a direct application of the laboratory data to industrial systems may not always be advisable. Another difficulty originates in the lack of theories able to predict variation of thermodynamic properties for commercially attractive systems with modifiers. Different additive compositions are used by different manufacturers of the same polymer. These are being “used-up” during processing and products’ life time, their content and chemical structure change. They may significantly affect the thermodynamic properties of a polymeric mixture, by the physical, viz. that of a co-solvent, and the chemical effects. For example, additives of one polymeric component of a blend may chemically react with additives of another polymeric component, mutually neutralizing each other. In particular, these effects may be large as far as the surface and interface energies are concerned. L.A. Utracki (Ed.), Polymer Blends Handbook, 123-201. © 2003 Kluwer Academic Publishers. Printed in the Netherlands.

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2.2

Thermodynamic Principles

2.2.1

Definitions

(or intrinsic) that do not depend on the mass, and extensive that do. For example, volume, entropy and total energy of a system, are extensive variables, but the specific volume (or its reciprocity — the density), molar volume or molar free energy of mixing are intensive. It is advisable to use whenever possible the intensive variables.

For convenience, the thermodynamic systems are assumed closed, isolated from the surroundings. The laws that govern such systems are written in terms of two types of variables: the intensive

Table 2.1.

Definition of principal terms

Term

Symbol

Total energy Entropy

U S

Helmholtz free energy Enthalpy Gibbs free energy Extensive properties of a mixture Change of an extensive property due to mixing Ideal solutions

F H G Ζm ∆Ζ m ∆GI

Regular solutions

∆GR

Excess properties Athermal solutions Combinatorial entropy

∆ZE ∆GA ∆SComb

Equation of state (EoS) Corresponding states

PVT

Reducing variables are characteristic materials’ parameters

P*, V*, T*

Free volume

Vf

Solubility parameter δ (can also be calculated from the molecular structure).

δ

Definition It comprises heat, Q, and work, W. S = k ln M; where M is a measure of the multiplicity of states F = U - TS H = U + PV G = U - TS + PV = F + PV = H - TS Ζ m = Ε m or Sm or Fm or Hm, etc. ∆Ζ m = Ζ m - ΣxiZi = Σxi (Zi - Zi0) = Σxi ∆Zi = ∆Z ∆GI = -T∆SI = kTNΣxilnxi; ∆ΗΙ = 0; ∆VΙ = 0, thus total randomness of molecular placement. ∆GR = ∆ΗΡ - T∆SI; ∆H R ∝ ( ε11 + ε12 ) / 2 - ε12 . The molecular interactions are non-specific, without associations, hydrogen or dipole-dipole bonding, distribution, orientation, etc. ∆ZE = ∆Z - ∆ZI ∆GA = -T∆S; ∆ΗΑ = 0 That part that originates from the number of possible placements of molecules in the lattice of an athermal solution (the latter assumption usually is abandoned in the following derivations). The relation between P, V, and T for a material A relation describing behavior of material in form of the reduced variables, e.g., P ≡ P / P*; V ≡ V / V*; T ≡ T / T*. Used to reduce corresponding independent variables, making them to follow respective corresponding states equation(s). Vf = V - Vo, where Vo ≡ lim V T→0

δ = √(∆ΗV − RT) / V , where the heat of vaporization: ∆HV = RT2 (∂ln P / ∂T) sat.

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The main independent variables are: the temperature (T), pressure (P), and composition (expressed as either molar, xi, or volume fraction, φi). The principal thermodynamics’ terms are listed in Table 2.1. Table 2.2 provides values of constants used in the thermodynamic calculations. 2.2.2

Three Laws of Thermodynamics

Thermodynamics focuses on the state of material, usually in a closed system. Historically, the laws were empirically formulated using the wellaccepted process of hypothesis, observation, and analysis. However, since the energetic state of a closed system is a sum of all the molecular and atomic motions, the statistical physics re-derived these laws from the first principles [Waldram, 1985; Gupta, 1990]. The First Law (Conservation of Energy) In a closed system the total energy remains constant (also known as sine perpetuum mobile). Thus, the change of internal energy: dU = dW + dQ = 0

(2.1)

where W represents the work done on or by the system, and Q represents the change of the thermal energy content. Note that there are several forms of energy that can be classified as “work”, for example compression, friction, electromagnetic interaction, etc.

Table 2.2.

For a perfect gas at constant pressure, P = const., Eq 2.1 gives the relationship between the heat capacity at constant pressure and constant volume, viz. Cp - Cv = R (R is the gas constant). Similarly, at T = const., Eq 2.1 predicts that the external work can only be performed at a cost of the internal energy: PdV = -dU. The Second Law (The Principle of Entropy Increase) The energy always flows from the higher to a lower level, hence the system tends to move toward a more uniform distribution of the local energetic states. Since the entropy provides a measure of randomness, in a closed system the entropy never decreases: dS = dQ/T • 0

(2.2)

The second law of thermodynamics is also known as the Carnot cycle principle that specifies that “heat can never pass from a colder to a warmer body without some other changes, connected therewith, occurring at the same time” (Clausius). Eq 2.2 can also be written as: (∂S / ∂V)T = (l / T)(∂Q / ∂V)T = (∂P / ∂T)V (∂S / ∂V)V = (l / T)(∂Q / ∂P)V = (∂V / ∂T)S

Useful constants

Constant Avogadro’s number Boltzmann constant Molar gas constant Molar volume at STP Origin of Celsius scale Planck’s Standard atmospheric pressure Velocity of light

Symbol NA kB R = NAk B Vo = RTo/P 0 To h Po c

Value 6.02205x1023 [mol-1] 1.38066x1023 [JK-1] 8.31441 [JK-1mol-1] 22.41383x10 -3 [m3mol-1] 273.15 [K] 6.62618x10 -34 [Js] 0.101325x106 [Nm-2] 2.997925x108 [ms-1]

(2.2a)

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The Third Law (Entropy Vanishes at Absolute Zero) In 1906, Nernst showed (experimentally) that the equilibrium entropy tends toward zero as the absolute temperature approaches zero:

(2) the relation between three variables — x, y, z: ˜z ˜z ˜y = ∂xy ∂yy ∂xy

(2.7)

(3) the commutation rule: lim S = lim dS = 0 T→0

(2.3)

T→0

 ˜  ˜z  ˜  ˜z  = ∂xy∂yx ∂yx ∂x y

As a corollary, near zero K the change of entropy in any process is negligibly small. Thus the third law is empirical, but so far no contradictory observation was reported. As a consequence of Eq 2.3, near the absolute zero the coefficient of thermal expansion and the pressure coefficient are expected to vanish:

T→0

Thus, a series of useful identities (known as Maxwell relations) was derived: ˜P ˜T  ˜P  ˜S =− ; = ; ∂SV ∂V S ∂TV ∂VS (2.9) ˜V ˜T ˜S ˜V = ; =∂S P ∂PS ∂PT ∂T P

lim (∂S/∂P)T = lim (∂S/∂V)T = 0

T→0

(2.8)

(2.3a)

lim (∂V/∂T)P = -lim (∂P/∂T)V = 0

The heat capacity at constant pressure and volume are related:

2.2.3

CP ≡ −

T→0

T→0

Interrelations Between Thermodynamic Variables (Tompa, 1956)

The full differentials (or the total changes) of the principal thermodynamic potentials cast in terms of their natural variables are: dF = -SdT - PdV; dH = TdS + VdP; dG = -SdT + VdP

˜H  ; ∂T P

CP ˜P  / ˜P = CV ∂VS ∂VT

From Eq 2.10 one can express the enthalpy and entropy gradients between two temperatures. Similarly, the compressibility and the thermal expansion coefficients are written as, respectively:

(2.4) KT ≡ (˜P / ˜ ln V)T ; α ≡ (˜ ln V / ˜T)P

Note that from Eq 2.4 the thermodynamic definitions of a pressure, P, and temperature T, are: P = (∂S/∂V)E / (∂S/∂U)V = (∂F/∂V)T ; T = (∂H/∂S)P = (∂U/∂S)V

(2.11)

The parameter KT is also known as the isothermal bulk modulus, and α as the volume expansion coefficient. They are interrelated by means of the Grüneisen constant, γ:

(2.5) γ = − (V / CV)[( ∂V / ∂T)P / (∂V / ∂P)T] = Vα / CVKT

These relations are important when deriving an equation of state by statistical methods. There are three rules for the partial derivatives of a function Z = Z(x, y, z): (1) the chain rule: ∂z = ∂z ∂u ∂xy ∂uy ∂xy

(2.10)

(2.6)

(2.12)

2.2.4

Multicomponent Systems

In a multicomponent, closed system comprising Ni moles of components, i = 1, 2, 3, ..., the changes

Thermodynamics of Polymer Blends

in the extensive functions of state, X = U, F, G, S, H, Cp, Cv, etc., caused by variation of composition are given by: dX = ΣXidNi , or X = ΣXiNi , i

i

thus Xi = (∂X / ∂Ni)T,P,N

(2.13)

ji

where the functions Xi are known as partial molar quantities. Thus, the full differentials of the state functions (viz. Eq 2.4) can be expressed as: dU = TdS − PdV + ΣµidNi i

dF = −SdT − PdV + ΣµidNi i

127

Eq 2.16 also introduced the free volume concept — note that as T → 0, V → b. Van der Waals considered that molecules move in “cells” made by the surrounding molecules with a uniform potential. The volume within which the center of a molecule can freely move, is what defines the free volume. Thus, one may distinguish: 1. Total volume, V. 2. Occupied volume, Vo (usually defined as V at T = 0 K). 3. Free volume, Vf = V - Vo. 4. Doolittle’s free volume fraction, fD = Vf / Vo. 5. Free volume fraction, f = Vf / V = 1 / (1 + 1 / fD), to be used in this text.

(2.14)

dH = TdS + VdP + ΣµidNi i

dG = −SdT + VdP + ΣµidNi i

where µi = (∂G / ∂Ni)ji is the chemical potential — in a homogenous system G = ΣµiNi. i

The relationships in Eq 2.14 indicate that in a closed, multicomponent system any change of the independent variables must be reflected in a change of the chemical potential: −SdT + VdP − ΣdNidµi = 0

(2.15)

Eq 2.15 is known as Gibbs-Duhem relationship.

2.3

Thermodynamics of a Single Component Systems

2.3.1

Equation of State (EoS) or PVT Relationships

In his Ph.D. thesis of 1873, van der Waals proposed the first EoS. The relation is frequently written in terms of reduced variables, indicating expected observance of the corresponding states principle: P = RT - a ; or P = 8T - 3 (2.16) V-b V2 3V - 1 V2 where: P* = a / 9b2; V* = 3b; T* = 8a / 9Rb

Detailed methods of computation of the van der Waals excluded volume (for any chemical structure) have been developed [Bondi, 1964; van Krevelen, 1976]. Thermodynamically, the free volume is expressed in terms of the entropy of vaporization: P = (RT / Vf) exp {-∆HV / RT} = (RT / V) exp {-∆SV / R} = RT / V or: ∆SV = R ln(V / Vf)

(2.17)

Over the years, many versions of the EoS theories have been proposed. Several comprehensive reviews of the EoS’s used in the polymer thermodynamics have been published. For example, Curro [1974] discussed applications of EoS within a full range of materials and variables, viz. to crystals, glasses, molten polymers and monoatomic liquids. The review discusses fundamentals of the theories as well as it provides a list of available experimental data. The comparison between different EoS was made on two levels, first by comparing the derived expressions for physical quantities (e.g., the characteristic reducing parameters, cohesive energy density, or internal pressure), then comparing how well the EoS describes the observed PVT dependencies for polymers. The second review focused on the molten state [Zoller, 1989]. The author examined four EoS theories, viz. [Spencer-Gilmore, 1949] (S-G),

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L. A. Utracki

[Flory-Orwoll-Vrij, 1964] (FOV), [SanchezLacombe, 1976-8] (S-L), and [Simha-Somcynsky, 1969] (S-S). Large deviations (” 0.01 mL/g) were observed for S-G, the two following relationships were useful only at low P and over small P-ranges, whereas S-S consistently provided the best representation of data over extended ranges of T and P, with deviations ” 0.003 mL/g, comparable to the experimental uncertainties. The FOV model can be expressed as:

The recent review of EoS theories also focused on the molten state [Rodgers, 1993]. The author collected PVT data for 56 polymers at P ” 200 MPa and T-range ∆T = 50 to 150°C (for details see Table 2.3). In addition to FOV, S-L, and S-S theories, the validity of these proposed by Prigogine [1953; 1957], Dee and Walsh [1988] (D-W), and Hartmann and Haque [1985] (H-H), were also examined. The D-W “modified cell model” resulted in the following EoS:

PV / T = (1 - V-1/3)-1 - 1 / VT

PV / T = (1 - 0.8909qV-1/3)-1 (2 / T)(1.2045V-2 - 1.011V-4) q ≅ 1.07; V = V / V*; T = T / T*; P = P / P* (2.21)

V = V / V*; T = T / T*: P = P/P* (2.18) 3 V* = ρ* ; T* = s*η* / (c*k); P*= c*kBT* / V* where: ρ*, s*, η*, c* are respectively: the “hardsphere” radius, number of contacts per segment, the segment-segment interaction energy, and the coordination number (kB is the Boltzmann constant). The S-L relation can be expressed as: PV / T = -V[ln(1 - p) + (1 - l / r)p] - 1 / TV p = 1 / V = V* / V; T = T / T*; P = P / P* where r = MWP* / RT*ρ*; T* = ε* / kB (2.19) with: Mw - the weight-average molecular weight, while ρ* is the characteristic density parameter. The parameter r represents the number of lattice sites occupied by the r-mer — its presence in the EoS negates the principle of corresponding states. The latter can be recovered only for r→’. More recently, Sanchez et al. [1993] reported that plots of the compressional strain, ln(V / Vo), vs. the reduced pressure, ∆P = (P - Po) / Bo, follow the same general curve independent of T. Here B is the bulk modulus, and Bo is its isothermal value evaluated at a reference pressure, Po. Thus, the general, corresponding states dependence can be expressed as: ln(V / V0) = -∆P + 1 B1 (∆P)2 - 1 (2B21 - B0B2) 2 3 (∆P)2 + … (2.20) An excellent superposition of data was achieved, indicating generality of the relation.

The H-H dependence is noteworthy by the simplicity of its form and respectable agreement with the experimental data: PV5 = T3/2 - lnV V = V / V0; T = T / T0; P = P / B0

(2.22)

In Eq 2.22 the characteristic pressure reducing parameter, B0, has been identified as the isothermal bulk modulus extrapolated to T = 0 and P = 0. Within the low pressure region Rodgers found all six EoS performing adequately — the worst performance was of S-L and the best of S-S theories. However, as the pressure increased, FOV and S-L dependencies started to perform poorly, while S-S and D-W were found to provide good description. Thus, the overall difference between the experimental and computed specific volumes for the 56 polymer melts was calculated for two ranges of pressures — from zero to 50 MPa, and within the full available range of pressure (see Table 2.3). The following differences were cited by Rodgers: for FOV ∆Vx104 = ±22(7), for S-L ∆Vx104 = ±33(10), for Prigogine ∆Vx104 = ±8(6), for H-H ∆Vx104 = ±9(6), for D-W ∆Vx104 = ± 6(5), and for S-S ∆Vx104 = ±7(4). Here, the deviations computed for the full pressure range are given first, then in parentheses, these for the range P ” 50 MPa.

129

Thermodynamics of Polymer Blends

The observations reported by Rodgers were later confirmed by Rudolf et al. [1995]. The authors examined suitability of FOV, S-L, and D-W theories to describe their own sets of data for PS and PI melts. Of the three EoS the one by D-W performed the best. Since this chapter is not intended to be a review of polymer thermodynamics, but to provide information on diverse thermodynamic aspects pertinent to polymer blends, only EoS derived by Simha and Somcynsky [1969], will be discussed in some detail. The S-S theory is based on the lattice model, where a liquid is represented as a mixture of unoccupied, h = 1 - y, and occupied, y, sites. Following the “Significant Liquid Structures” nomenclature [Eyring and Jhon, 1969], the model considers a liquid as being an intermediate between solid and gas. For liquid “i”, S-S used the following nomenclature: yi = si Ni / ( si Ni + Nhi) fraction of occupied lattice sites si number of segments per chain of molar mass Mi Ni and Nhi are the number of occupied and vacant sites, respectively 3ci number of external degrees of freedom per chain Moi segmental molar mass Mi / si v*ii molar intermolecular repulsion volume between segment pair of constituent liquid i ε*ii maximum molar intermolecular attraction energy between segment pair of constituent liquid i A = 1.011, B = 1.2045, constants and z = 12 Qi = 1/ (yiVi), and dimensionless quantities ηι = 2-1/6yiQi1/3 NA and hP Avogadro’s and

qiz = si(z-2) + 2 Pi* = zqiε*ii / siv*ii; Ti* = zqiε*ii / Rci; and Vi* = v*ii / Moi

Planck’s numbers, respectively number of intermolecular contacts the three principal reducing (or scaling) parameters; thus: Pi*Vi* / Ti*Moi = R(ci / si)

To derive the EoS the authors first calculated the partition function, Z, for all possible number of arrangements of occupied sites and empty holes in a lattice with z coordination number. Since the Helmholtz free energy is directly given as: F = -kBT lnZ, its differentiation (see Eq 2.5) gives the pressure, thus the equation of state: PiVi / Ti = (1 - ηi)−1 + 2yiQi2 (AQi2 - B) / Ti (2.23) Minimization of F gave the second relation: 3ci[(ηi - 1/3) / (1 - ηi) − yiQi2 (3AQi2 - 2B) / 6Ti] + (1 - s) - si ln[(1 - yi) / yi] = 0 (2.24) Thus, of all the EoS used for polymeric species, only that derived by S-S explicitly gives the hole fraction, h = 1 – y, which is directly related to f. Eqs 2.23 and 2.24 provide a corresponding states description of PVT behavior of any liquid. Once the four characteristic parameters: P*, V*, T*, and 3c/s are known, the specific volume and all its derivatives are known in the full range of P and T. For linear molecules the external degrees of freedom are proportional to number of segments: 3ci = si + 3. Thus, for linear polymers, where si >> 3, the external degree of freedom: 3ci/si = 1, i.e., for polymers, only three parameters: P*, V*, and T*, are required. Typical values of P*, V*, and T* for selected polymers are listed in Table 2.3. At atmospheric pressure, P = 0, and within the range of the reduced volume, 0.95 < V < 1.40, Eqs 2.23 and 2.24 predict that the volume expansion with T should follow the dependence [Simha and Weil, 1970]: ln Vi = S1(si, ci) + S2(si, ci)Ti3/2

(2.25)

130

Table 2.3.

Polymer

HDPE LDPE PS PoMS PVAc PMMA PcHMA PnBMA BPE LPE HMLPE PIB PDMS PMP PTFE PSF PBD PEG PTHF LDPE-A LDPE-B LDPE-C i-PP i-PB PET PPE PC PAr Phenoxy PEEK PVME PA-6 PA-66 PMA PEA PEMA TMPC HFPC BCPC PECH

L. A. Utracki

List of polymers, the temperature and pressure ranges of data, as well as the S-S characteristic parameters and an average difference between the measured and computed volumes [Rodgers, 1993] T-range

P-range

P*

T*

V*

∆Vx104

(°C)

(bar)

(bar)

(K)

(mL/g)

(mL/g)

0-1960 0-1960 0-2000 0-1800 0- 800 0-2000 0-2000 0-2000 0-2000 0-2000 0-2000 0-1000 0-1000 0-1960 0- 390 0-1960 0-2835 0- 685 0- 785 0-1960 0-1960 0-1960 0-1960 0-1960 0-1960 0-1765 0-1765 0-1765 0-1765 0-2000 0-2000 0-1960 0-1960 0-1960 0-1960 0-1960 0-1600 0-2000 0-2000 0-2000

5063 7936 7159 7461 9474 9264 7722 8560 6923 7864 9213 6866 5014 5453 6581 11160 8150 9145 7255 7162 7036 7188 5730 6037 11940 9294 10200 10030 11390 10860 8481 5499 7069 9691 8308 9870 8192 8510 9878 9131

12290 9595 12840 13080 9348 11940 12030 10310 10390 9793 9207 11360 7864 11030 8126 12770 9225 10150 10280 10580 10860 10660 11060 10920 11800 10580 11830 12390 11730 12580 10360 16870 12640 10460 10040 10190 11540 10550 12190 11370

1.2190 1.1380 0.9634 0.9814 0.8126 0.8369 0.9047 0.9358 1.1674 1.1406 1.1278 1.0940 0.9592 1.2050 0.4339 0.7903 1.0766 0.8812 1.0087 1.1664 1.1734 1.1679 1.1884 1.1666 0.7426 0.8602 0.8156 0.8091 0.8529 0.7705 0.9632 0.8327 0.8195 0.8431 0.8773 0.8710 0.8794 0.6317 0.6975 0.7343

14.3 22.6 3.5 4.6 1.4 1.0 6.6 13.1 9.9 9.0 11.0 2.0 4.8 17.9 13.6 3.6 3.6 4.1 4.0 6.5 5.9 6.1 12.6 7.5 5.7 9.1 3.7 2.6 3.5 8.5 6.8 4.4 5.2 9.3 8.8 8.1 9.6 6.6 4.7 3.1

140- 203 121-175 115-196 139-198 35-100 114-159 123-198 34-200 125-198 142-200 137-200 53-110 25-70 241-319 330-372 202-371 4-55 88-224 62-166 112-225 112-225 112-225 170-297 133-246 274-342 203-320 151-340 177-310 68-300 346-398 30-198 236-296 246-298 37-220 37-217 113-161 218-290 159-280 155-284 60-140

131

Thermodynamics of Polymer Blends

Table 2.3.

Polymer

PCL PVC a-PP EP 50 EVAc 18 EVAc 25 EVAc 28 EVAc 40 SAN 3 SAN 6 SAN 15 SAN 18 SAN 40 SAN 70 SMMA 20 SMMA 60

Continued: List of polymers, the temperature and pressure ranges of data, as well as the S-S characteristic parameters and an average difference between the measured and computed volumes [Rodgers, 1993] T-range

P-range

P*

T*

V*

∆Vx104

(°C)

(bar)

(bar)

(K)

(mL/g)

(mL/g)

0-2000 0-2000 0-1000 0- 625 0-1765 0-1765 0-1765 0-1765 0-2000 0-2000 0-2000 0-2000 0-2000 0-2000 0-2000 0-2000

7845 8495 6277 5720 7056 6978 7472 7539 7642 8238 7792 7853 8118 8747 7640 7911

10870 12350 9494 12220 10630 10440 10310 10360 12070 11490 12360 12380 12900 13790 11800 11780

0.9173 0.7230 1.1274 1.2227 1.1341 1.1040 1.0949 1.0446 0.9416 0.9352 0.9299 0.9255 0.9124 0.8906 0.9186 0.8739

3.0 4.2 7.4 15.8 4.7 5.6 6.7 5.4 8.1 7.0 5.0 4.2 5.6 3.6 4.2 5.5

100-148 100-150 80-120 140-250 112-219 94-233 94-233 75-235 105-266 96-267 132-262 104-255 100-255 100-270 110-270 110-270

(It is noteworthy that for P = 0, the H-H Eq 2.22 also predicts that ln V ∝ Ti3/2.) For linear polymers where 3c/s = 1, S1 = -0.1034 and S2 = 23.835 [Simha et al., 1973]. Within the same range of the reduced variables: yio Vio = K(Tio), ∴ hio = 1 - yio = [Vio(T) - K(T)Vio] / Vio (T)

(2.26)

Thus, the vibrational contribution is expected to be nearly constant, K(T) = 0.95, indicating that an effective excluded volume: K(T)Vi* = Vi* (T = 0) = Vo . Furthermore, there is a correlation between the self-repulsion volume in S-S theory and the van der Waals group volume [Yahsi, 1994; Simha and Yahsi, 1995]: v*ii = 1.59Vw + 1.00. Eqs 2.23-2.26 are valid for any liquid. When upon cooling the temperature reaches the glass transition region, part of the free volume fraction is no longer accessible for the molecular motion.

However, the S-S theory can also be used in the glassy region assuming that a part of the free volume is frozen at the glass transition temperature, Tg. Experimentally, the frozen fraction of the free volume, FF, depends on the absolute value of Tg. This finding was first reported for several polymers at ambient pressure [Simha and Wilson, 1973]. Recently, the generality of this observation was confirmed by analyzing isobaric thermal expansion of PS for a wide range of pressures [Utracki and Simha, 1997]. As shown in Figure 2.1, the frozen fraction of the free volume, FF = FF(Tg), follows the same dependence whether Tg changes are caused by the polymeric chemical structure or imposed pressure. The observed, general dependence follows the relation: FF = 0.997 - 4.75×10-4Tg - 1.52×10-6T2g ; R = 0.975 (2.27)

132

L. A. Utracki

It is gratifying to see that as the glass transition temperature approaches the absolute zero, Tg→0 K, Eq 2.27 predicts that all free volume should freeze, FF(Tg = 0) = 1. On the other hand, within the high temperature range, as Tg exceeds T = 669 K = 396°C all free volume should be accessible to thermal motion in the glassy state, i.e., FF(Tg • 400°C) = 0. Thus, it is to be expected that the high temperature polymers will have the same thermal expansion coefficient across the glass transition temperature, α L = α G, where subscripts L and G indicate liquid and glassy state, respectively. Recently, new EoS was proposed [Jung, 1996]: PV / T = -V ln(1 - 1 / V) - 1 - (4 / T) (1.2045 V-2 - 1.011 V-4) V = V / V*; T = T / T*; P = P / P*

(2.28)

To examine ability of the relation to describe PVT dependencies, the author used experimental data of eight polymers (T-span = 30°C, P-span = 0 - 20 MPa). For comparison, FOV, S-L, S-S, and D-W relations were also used. The evaluation was performed computing errors in describing the volume, thermal expansivity and isothermal compressibility. As in the previous evaluations, S-S dependence performed the best. For the

description of PVT the new EoS performed as well as that of D-W, but for the expansivity and compressibility it outperformed the latter EoS. 2.3.2

Solid-liquid and Vapor-liquid Equilibria

For a single component, the phase diagram provides a map of the solid, liquid, and vapor states, as well as their coexistence regions. It is customary to construct it as a log-log plot of the P vs. V relation at constant T. Figure 2.2 shows a phase diagram of argon. Dividing the three variables by the corresponding reducing parameters, the dependence can be cast in a general plot of reduced pressure vs. reduced volume, P vs. V, at constant reduced temperature, T. For low molecular weight liquids, the van der Waals Eq 2.16 provides a reasonable description. Guggenheim demonstrated that near the critical point the coexistence curve of liquids and vapors follows a simple proportionality: T - 1 ∝ (1 - p)3 / 3

(2.29)

This dependence is shown in Figure 2.3 as a solid line — the broken one represents the van der Waals parabolic prediction.

Figure 2.1. Frozen fraction of free volume vs. glass transition temperature. Full triangles — values for different polymers at ambient pressures [Simha and Wilson, 1973]. Squares — PS data at pressures P = 0 - 400 MPa (data [Rehage, 1980]; calculations [Utracki and Simha, 1997]).

133

Thermodynamics of Polymer Blends

Figure 2.2. The van der Waals-type phase diagram for one mole of argon. C = 150 K is the critical point, TP is the triple point line of the coexistence between solid, liquid and vapor. The upper shaded area represents the liquid-vapor coexistence, while the lower that of solid-vapor. Vapor exists on the right hand side of the shaded areas, while liquid and solids on their left. The supercritical region is that above the critical point, where (at high pressures) the “vapor” density is comparable to that of liquid. Supercooled vapor is also indicated by a series of solid points. After [Flowers and Mendoza, 1970].

T/T

c

0.9

0.7 Guggenheim van der Waals

0.5 0

1

r/r

c

2

Figure 2.3. The liquid-vapor coexistence curves, T vs p, as predicted by Guggenheim (solid line) and van der Waals.

134

2.3.3

L. A. Utracki

Gibbs Phase Rule

The state variables (intensive, i.e., independent of the amount of matter, and extensive that are proportional to it) are these quantities that describe a system, for example, by means of the “equation of state.” For the discussions of phase diagrams it is important to know how many of the state variables one may change without going through a phase transition. The total number of variables required to describe a system by Eq 2.14 is N + 2, where N is number of components and “2” stands, e.g., for V and T. For a closed system the number of intensive variables (also known as a degree of freedom), #f, is given by the “Gibbs the phase rule:” #f = N + 2 - Π

(2.30)

where Π is number of phases. For example, for a phase equilibrium of a two component blend (e.g., PVC with PCL), N = 2, hence Π = 2, thus #f = 2. Thus, Eq 2.30 predicts that to follow the two phase coexistence two variables must be simultaneously changed (e.g., T and composition), but within a single miscible phase three variables are available, for example, composition, T and P.

is a large number of possible arrangements of macromolecules in space containing small solvent molecules [Yamakawa, 1971]. Traditionally, solutions have been used to characterize the polymer — to measure its molecular weight averages, e.g., number, weight and z-average: Mn, Mw, and Mz, or the size of its macromolecular coil. The latter maybe expressed as the unperturbed end-to-end distance (six times larger) or the radius of gyration, viz.: 〈ro2〉 = 6〈so2〉 = σ2 nl2[(1 + 〈cosθ〉) / (1 - 〈cosθ〉] (2.31) In Eq 2.31 σ is the steric hindrance factor of the macromolecular chain, n is number of statistical segments, each of length l and with bond angle between two consecutive segments equal to θ. The solutions are also used to measure of the thermodynamic interactions between the polymer segments and solvent molecules. The latter is best discussed in terms of the virial coefficients, Ai. The change of solvent chemical potential upon dissolution of polymer is given by: ∆µ ≡Ż(∂∆Gm / ∂N2)T,P,N = RT lnao = -V1π 1

2.4

Polymeric Liquid Mixtures

Polymeric liquid mixtures are conveniently divided into solutions (containing a low molecular weight liquid) and blends, containing only macromolecular species. They will be briefly described in the following parts, 2.4.1 and 2.4.2, respectively. More detailed discussion of the theoretical foundations is given in Part 2.5. 2.4.1

Polymer Solutions

By contrast with the mixtures of low molecular weight specimens, where concentration is customarily expressed as a mole fraction, xi, that of polymer solution is given either as volume fraction, φi, or (for dilute solutions) wt/vol concentration, c (g/100 mL). In solutions, the solubility originates mainly from the entropic effects — there

(2.32)

where ao is activity. The osmotic pressure, π, is usually expressed as: π = RTΣAi; A1 = Mn-1; A2 = 4π3/2NAMn-2〈s2〉3/2Ψ; (2.33) The second virial coefficient in Eq 2.33 contains two functionals that depend on the binary interactions — the perturbed radius of gyration and the function ψ = ψ(z). Here, the parameter z and the function ψ are: z = (4 / 33/2) CMΨ(1 - Θ / T)Mn1/2 = (4/33/2)CM(½ - χ12)Mn1/2

(2.34)

Ψ = [1 - exp{-5.73z}] / 5.73 As z increases from zero (the theta conditions) to 2 (good solvent) the function also increases

135

Thermodynamics of Polymer Blends

from zero to about 0.2. Thus, A2 provides a direct measure of the intermolecular interactions in polymer solutions. For simplicity, usually dilute solutions are used to avoid the need for determination of higher virial coefficients. However, since A3 ∝ Α22 Mn, Eq 2.33 can be written as: (πMn / RT)1/2 = 1 + MnA2 / 2

(2.35)

ηr = η / ηo; ηsp = (η − ηo) / ηo; [η] ≡ lim ηsp / c = lim(lnηr) / c c→0

(2.37)

c→0

Many relations have been proposed between the intrinsic viscosity, [η], and measures of the polymer/solvent interaction parameters. One of the better known is that credited to Inagaki et al. [1966]:

The light scattering also makes it possible to determine A2, viz.:

[η]4/5 / Mv2/5 = 0.786K4/5 + 0.454K2/15 Φo2/3 B2/3

Kc / RΘ = ∑iAi; A1 = 1 / Mw 〈P1(Θ)〉;

here K = Φo[〈Ro2〉 / Mv]3/2, the parameter Φo ≅ 2.5×1023 (cgs), B = (2ρs2 / Vo)(½ - χ), with ρs being segmental density and Vo the molar volume of the solvent. It is worth noting that the three methods of evaluation of the solution behavior: osmometry, light scattering, and intrinsic viscosity, provide different molecular weight averages, respectively, number, weight, and the “viscosity” average: Mn, Mw, and Mv. Knowing at least two of them one can estimate the molecular weight polydispersity factor, Mw/Mn, reasonably well.

(2.36)

i=1

where K is an experimental quantity, RΘ is the Rayleigh ratio, and 〈P1(Θ)〉 is the intramolecular interference factor. Similarly as Eq 2.33, also Eq 2.36 makes it possible to calculate virial coefficients. The thermodynamic interactions and the size of polymer coil enter dependencies that describe the transport behavior of polymer solutions, viz. viscosity, diffusion, sedimentation, etc. To complete this short summary, the intrinsic viscosity should be mentioned. Defining η as the solution viscosity and ηo, as that of the solvent, the following: relative, specific and intrinsic viscosities are expressed as respectively,

Table 2.4.

Mv1/3

2.4.2

(2.38)

Polymer Blends — Definitions and the Miscibility of Blends

By contrast with solutions, the polymer blends are mostly immiscible. As shown in Appendix 2 by

Terms related to polymer blend miscibility (see also Chapter 1. Introduction)

Term

Definition

Miscible polymer blend

Polymer blend, homogenous down to the molecular level, in which the domain size is comparable to macromolecular dimension, associated with the negative value of the free energy and heat of mixing, ∆Gm § ∆Hm ” 0, and ∂2∆Gm / ∂φ 2 > 0. Polymer blends whose free energy of mixing; ∆Gm § ∆Hm > 0. Immiscible, compatibilized polymer blend with the modified interface and morphology. Third phase in binary polymer alloys, engendered by interdiffusion or compatibilization. Its thickness ∆l = 2 to 60 nm depends on polymers’ miscibility and compatibilization. Process of modification of the interphase in immiscible polymer blends, resulting in reduction of the interfacial energy, development and stabilization of the desired morphology, leading to the creation of a polymer alloys with enhanced performance.

Immiscible blends Polymer alloy Interphase Compatibilization

L. A. Utracki

Krause and Goh, over 1000 cases of miscibility have been found. However, these “islands” are so infrequent and poorly defined that one may consider them exceptions from the general rule of polymer/polymer immiscibility. While the thermodynamic definition of miscibility is unambiguous, (see Table 2.4) there is a significant amount of discussion as to the methods of miscibility detection and the size of heterogeneity in miscible blends. The methods of miscibility detection (sometimes wrongly labeled as “compatibility”) will be discussed later in this chapter. Opacity can provide only limited information on miscibility — light scattering can be observed only when the size of heterogeneity is larger than 100 nm and the difference in refractive index is greater than about 0.01. The most widely used tool for the “detection of miscibility” is the glass transition temperature, Tg. There is a believe that blends that display a single Tg are miscible. The glass transition temperature, Tg, is relatively simple to measure, but there are inherent uncertainties of the measurements that needs to be examined [Utracki, 1989]. Tg is insensitive when the amount of the second component is less than about 10 wt%. Furthermore, the method should not be used for blends containing polymers whose Tg‘s do not differ at least by 10°C from each other. As it has been shown, first by Schultz and Young [1980] then by many others, Tg is not sensitive to the thermodynamic miscibility of the components, but rather to the degree of dispersion. For example, in solvent-mixed PS/PMMA blends that were not allowed to phase separate, a single Tg has been detected, but when the specimens were annealed, double peaks were observed. More recently, solvent cast blends of PVC with caprolactone-grafted lignin showed a single Tg. At the same time, the measured domain size (from 10 to 30 nm) indicated immiscibility [De Oliveira and Glasser, 1994]. For PAEK/PI blends the Tg’s of neat components are about 90°C apart [Sauer et al., 1996]. However, depending on the chemical nature of the components, blend composition and blending conditions, three types of the glass transition behavior were observed.

250

PEK/N-TPI Tg (°C)

136

200

150 0

40

PEK (wt%)

80

Figure 2.4. The glass transition temperature vs. composition for PEK/N-TPI blends. The two large black circles represent neat polymers, open circles represent data for blends prepared at T = 455°C, and small solid circles represent these prepared at 400°C. The lines are only to guide the eye. Data from [Sauer et al., 1996].

For example, PEK was blended with N-TPI (N-TPI stands for: “New Thermoplastic PolyImide”, obtained from condensation of 4,4-bis(3-aminophenoxy)biphenyl with pyromellitic dianhydride). At “low” temperature of T = 400°C, the mixture showed two Tg’s, when blended at T • 440°C it showed one (see Figure 2.4), but when blended at intermediate temperatures, 400 < T°C < 440, a single, broad, overlapping Tg transition was obtained. Clearly, the glass transition behavior reflected different degree of dispersion of the system (the degradation temperature is 410°C). Frequently, on the basis of Tg measurements, blends of engineering resins with such copolymers as SBS or ABS-type (e.g., PPE/SBS, PC/ABS, or PPE/MBS) are identified as miscible. Obviously, the method focuses on the interactions between the viscoelastic components of the blend, neglecting the elastomeric domains. These remain immiscible and may show (or not, depending on their content) another Tg at lower temperatures. In any liquid, either single or multicomponent one, there is a significant amount of heterogeneity, usually expressed by the density fluctuation function. Thus, it is fair to ask two questions: (1) What is the maximum size, dd, of heterogeneity in polymer blend that fulfills the condi-

Thermodynamics of Polymer Blends

tions of the thermodynamic miscibility, viz. ∆Gm § ∆Hm ” 0, and ∂2∆Gm / ∂φ2 > 0? (2) At what level of heterogeneity, a monotonic function of Tg vs. composition is to be found? Many answers have been given to the first question. On the basis of the size of the cooperative segmental motion required at Tg, it was estimated that dd = 2-3 nm [Boyer, 1966; Warfield and Hartmann, 1980]. The local segmental dynamics of a flexible polymer chain was found to be affected by the neighboring chains lying within 2 to 7 nm [Chung and Kornfield, 1994]. Other authors consider that miscibility is achieved when the heterogeneity diameter becomes comparable to the unperturbed radius of gyration of the macromolecule, thus dd = 〈so2〉1/2 ≅ 3 to 10 nm [Silberberg and Kuhn, 1952; Wolf, 1980, 1984]. Many others’ estimates fall within these limits [e.g., Kaplan, 1976; Bair and Warren, 1980; Cowie, 1989]. Thus, it seems that the true, thermodynamic miscibility is associated with the size of compositional heterogeneity dd < 10 nm. The answer to the second question has been given as well. Depending on the chemical nature of the system and its morphology, double peak of Tg has been reported for the domain sizes of dd • 15 to 20 nm [Frisch et al., 1982; Utracki, 1989; De Oliveira and Glasser, 1994]. From the fundamental point of view, the glass transition reflects a change in the molecular mobility upon cooling, and it is associated with “freezing” of a portion of the free volume. However, the frozen fraction depends on the absolute value of Tg — as its absolute value increases more and more free volume becomes accessible to the thermal motion in the glassy state. Judging by data presented in Figure 2.1, for organic macromolecules the fundamental mechanism of the glass transition is not expected to be valid for Tg > 400°C. In short, even under the most favorable conditions, Tg is unable to discriminate between presence and absence of thermodynamic miscibility — if a system is miscible, a single Tg will indeed be found, but a single Tg will also be observed for immiscible systems having finely dispersed phases. In many cases Tg may be able

137

to detect the “technological miscibility,” i.e., to identify the systems that are capable to be so well homogenized that the phase domains will not be affected by the processing conditions. The method may be used as a pragmatic scan for industrially useful blends, or evaluation of a compatibilization scheme. It is noteworthy that “homogeneity at fairly fine level is necessary for optimum performance, but some degree of microheterogeneity is usually desirable to preserve the individual properties of respective polymer components” [Hess et al., 1993]. Note that nearly all commercial polymer blends (the exception is blend of PVDF with PMMA) are immiscible. One tends to study miscibility not to develop single-phase commercial blends, but mainly to design better compatibilizers and compatibilization strategies. Since the standard test methods for detecting miscibility, viz. Tg, microscopy, small angle X-ray scattering, etc., are limited to dd • 15 nm, other methods for studies of the thermodynamic miscibility should be used. To these belong the advanced scattering methods, such as high resolution nuclear magnetic resonance techniques, the NMR-spin diffusion, non-radiative energy transfer, excimer fluorescence, thermally stimulated depolarization current, small angle neutron scattering, SANS, Fourier transfer infra-red spectroscopy, FTIR, etc. Even simple NMR measurement of the spin-lattice relaxation times, T1, is capable to distinguish the domain size of 2-3 nm. The method can be used for either molten or solidified specimens. For example, these methods produced the following results [McBrierty and Packer, 1993]: a-PVC/PMMA: Homogenous at 20 nm, but heterogeneous at 2 nm. After deuteration of PMMA the miscibility extended to 1-2 nm. PS/PVME, PnBMA/PS (2 mol % -OH): Homogenous at 20 nm. SAN/PMMA: Homogenous at 20 nm, but heterogeneous at 2 to 15 nm [McBrierty et al., 1978]. a-PET/PC, a-PVDF/PMMA, a-PCL/PVC: Homogenous at 2 nm [Douglass and McBrierty, 1978; Douglass, 1979].

138

L. A. Utracki

PS/PPE: Only about 30 wt% blend components participate in intimate interactions on the scale of less than 2 nm. The rest show the same nuclear resonance pattern as that recorded for he two homopolymers [Stejskal et al., 1981; Takahashi et al., 1990].

2.5

Theories of Liquid Mixtures

2.5.1

Lattice, Cell, and Hole Theories

The statistical mechanics methods that use a pseudo-crystalline model of regularly placed elements on a “lattice” are known as lattice theories. Many theories, known under the names of free volume, cell-hole, tunnel, Monte Carlo, or molecular dynamics belong here. Of these, only two will be mentioned. The first, and the best known, was originally developed by Huggins [1941] and by Flory [1941], then extended by many authors [e.g. Utracki, 1962; Koningsveld, 1967]. The second, is the cell-hole Simha-Somcynsky [1969] theory that has been incessantly evolving during the intervening years. The theory makes it possible to interpret and predict different material’s behavior in a wide range of states and independent variables, viz. gases, low molecular weight organic liquids, metals, polymers in a glassy or molten state, to determine miscibility of gases or liquids in polymers, to compute the phase diagrams of polymer blends, etc. 2.5.1.1 Huggins-Flory theory For binary systems that contain an ingredient i = 1 or 2 (traditionally, for polymer solutions the subscript 1 indicates solvent, and 2 polymer) the Huggins-Flory, H-F, relation has been expressed in several equivalent forms: ∆Gm / RTV = (φ1 / V1) lnφ1 + (φ2 / V2) lnφ2 + (χ12 / V1)φ1φ2 ∆Gm = RTV[(φ1 / V1) lnφ1 + (φ2 / V2) lnφ2] + Bφ1φ2 χ12 / V1= χ′12; B ≡ χ12 RT(V / V1)

(2.39)

In Eq 2.39, φi is the volume fraction and Vi is the molar volume of the specimen “i”. The first two logarithmic terms give the combinatorial entropy of mixing, while the third term the enthalpy. For polymer blends Vi is large, thus the combinatorial entropy is vanishingly small — the miscibility or immiscibility of the system mainly depends on the value of the last term, χ′12 φi φ2. Applying to Eq 2.39 the conditions for phase separation: (∂2∆Gm / ∂φ22)P,T = (∂3∆Gm / ∂φ23)P,T = 0, and treating the so call binary interaction parameter, χ12, as a constant, the critical conditions for the phase separation have been expressed as: χ′12,cr ≡ χ12,cr / V1 = (V1-1/2 + V2-1/2)2 / 2 ≅ (N1-1/2 + N2-1/2)2 / 2

(2.40)

Βcr = (RTV / 2) (V1-1/2 + V2-1/2)2 ≅ (RT / 2)(N1-1/2 + N2-1/2)2 where Ni is degree of polymerization of polymer i = 1 or 2. Eq 2.40 gives the miscibility conditions for systems with species of different molecular weight. Three cases can be distinguished: 1. solvent solutions: V1 = V2 or N1 = N2 = 1, are miscible if χ12 < χ12,cr = 2 2. polymer solutions: V1 / V2 → ’ or N1 = 1 and N2 → ’, are miscible if χ12 < χ12,cr = 1/2 3. polymer blends: V2 and V1 → ’ or N1 = N2 → ’, are miscible if χ12 < χ12,cr = 0 Originally, after Hildebrand, the parameter χ12 • 0 was assumed to have a single, characteristic value for a given system. However, it was soon found that even for polymer solutions χ12 is a complex function of many independent variables, viz. concentration, temperature, pressure, molecular weight, molecular weight distribution, etc. For calculating isobaric phase diagram, the influence of the first two variables should be expressed as [Koningsveld, 1967]: n”2

m”2

i=0

j=-1

χ12 = ∑aiφi2 and ai = ∑aijTj

(2.41)

139

Thermodynamics of Polymer Blends

In this notation χ12 is expressed as composed of the enthalpic and entropic parts, χH and χs respectively, each being a function of concentration, molecular weight of both polymers, and other independent variables. In other words, Eq 2.43 takes into account contribution of the noncombinatorial entropy. This idea is particularly important for polymer blends. Here the miscibility mainly originates in strong interactions that are expected to cause changes of intersegmental orientation, hence entropic effects. Table 2.6 gives few examples for recent values of Eq 2.43 parameters for selected, simple polymer blends. Similar, large variability of χ12 has been observed for the concentration dependence in polymer blends. For example, in Figure 2.5 the experimental values of this parameter are shown for blends of a hydrogenated polymer with its deuterated homologue, i.e., PE with d-PE at T = 443K, and PS with d-PS at T = 433K [Londono et al., 1994]. The data can be described in terms of Eqs 2.41 or 2.43. A formal analysis of these data was published by Bidkar and Sanchez [1995]. Another drawback of the H-F theory was the initial assumption that all lattice cells are occupied by either solvent molecules or polymeric segments that are of equal size. As a consequence the free volume contribution was neglected. Maron [1959] pointed out that dissolution of polymer is associated with volume changes — his modification of

Thus, it takes nine parameters to describe variation of χ12 with concentration and temperature, at constant pressure. For less demanding thermodynamic calculations, Eq 2.41 can be simplified. Thus, to express conditions of miscibility in PS blends with poly(styrene-co-4-bromostyrene) the binary interaction parameter per one mer of styrene (within the T = 440-500 K region) was expressed as [Strobl et al., 1986]: χ12 / N = -0.02208 + 20 / T - 0.01482φ2. Similarly, for PS blends the following relation was used [Ahn et al., 1997]: B12(φ2,T) = B0 + B1φ2 + B2 × T / 1000 B12(φ2=0.5,T) = Bs+B2 × T / 1000

(2.42)

where the experimental values of Bi parameters are listed in Table 2.5. More generally, Eq 2.41 can be written as: χ12(φ2,T,M1,M2...) = χH (φ2,M1,M2...) + χs(φ2,M1,M2,...) / T B12(φ2,T,M1,M2...) = T × BH (φ2,M1,M2...) + Bs(φ2,T,M1,M2...) (2.43)

Table 2.5.

Parameters of Eq 2.42 (Ahn et al., 1997)

Polymer-1

Polymer-2

PS, MW = 48 PS, MW = 48 SBR, MW = 150 SBR, MW = 50 SBR, MW = 50 PS SBR (68% styrene) SBR (60% styrene) PBD PBD

PVME, MW = 81 PVME, MW = 180 PVME, MW = 81 PVME, MW = 81 PVME, MW = 180 PVME PVME PVME PVME PS

Bo

B1

B2

- 0.423 - 0.375 - 0.338 - 0.299 - 0.283 -

0.027 0.028 0.007 0.012 0.019 -

1.074 0.970 0.920 0.894 0.852 1.020 0.920 0.870 - 0.670 - 1.970

B3 - 0.385 - 0.335 - 0.285 0.855 1.620

Notes for Table 2.5: N is the degree of polymerization, d-PE - deuterated PE, d-PS - deuterated PS, PpMS - poly(p-methyl styrene).

140

L. A. Utracki

H-F theory was extended to ternary systems comprising polydisperse polymer by Utracki [1962]. Starting in the early 1960’s considerable effort was made to develop what become known as the equation of state theories [Flory et al., 1964; Eichinger Table 2.6.

and Flory, 1968; Simha and Somcynsky, 1969; Patterson, 1969, 1982; Patterson and Robard, 1978; Sanchez and Lacombe, 1976, 1977; Sanchez, 1983, 1984].

Enthalpic and entropic contributions to the H-F binary interaction parameter, χ12 (see Eq 2.43)

Polymer-1

Polymer-2

PE, N = 2538 PE, N = 3308 PE, N = 3308 PE, N = 4598 PE, N = 4598 PE, N = 4598 PE, N = 4598 PE, N = 4598 PE, N = 4598 PS. N = 15,400 PpMS, N = 498 PpMS, N = 498 PpMS, N = 498 PpMS, N = 1108 PpMS, N = 835

d-PE, N = 2464 d-PE, N = 3275 d-PE, N = 3275 d-PE, N = 4148 d-PE, N = 4148 d-PE, N = 4148 d-PE, N = 4148 d-PE, N = 4148 d-PE, N = 4148 d-PS, N = 8,700 d-PS, N = 291 d-PS, N = 291 d-PS, N = 291 d-PS, N = 291 d-PS, N = 3123

φ2

χHx104

0.221 0.087 0.457 0.044 0.087 0.131 0.221 0.457 0.708 0.500 0.250 0.500 0.749 0.249 0.5

- 0.012 2.761 0.242 1.628 0.759 0.808 - 0.907 - 0.843 - 0.859 - 2.900 - 0.011 - 0.0081 - 0.0081 - 0.0119 -2

χ s (K) 0.132 0.069 0.089 0.325 0.138 0.100 0.139 0.127 0.133 0.200 70 57 61 72 2.16

References Londono et al., 1994

Londono & Wignall, 1997

Jung & Fischer, 1988

Figure 2.5. The binary interaction parameter for blends of PS with deuterated-PS at T = 433 K and PE with deuterated-PE at T = 443K [Londono et al., 1994].

141

Thermodynamics of Polymer Blends

2.5.1.2 Equation of State Theories

where, xi is the molar fraction, and the values in angle brackets, < >, are compositional averages:

The equation of state, EoS, theories of mixtures are based on the principles discussed in Part 2.31. Formally, the computation of the partition function for a single component or for a mixture of components is similar, yielding the Helmholtz free energy of mixing. Different EoS use different measures of the binary interactions between components. For example, in Flory’s extension of the H-F approach the binary interaction parameter, χ12, is still present [Flory, 1970]. In S-L EoS (see Eq 2.19) the characteristic pressures for the mixture of species are assumed to be pairwise additive: P* = ∑φiPi* - ∑∑φiφj∆P* i

(2.44)

i<j

For such systems, the difference of the interaction energy density, ¨P*, has been considered a measure of the binary interactions between polymeric segments, proportional to either χ12 or B [Sanchez, 1989]. The most successful statistical thermodynamic theory of liquids is that proposed by Simha and Somcynsky, S-S [Simha and Somcynsky, 1969; Jain and Simha, 1980, 1981, 1982, 1984]. From the partition function, the Helmholtz free energy, then the pressure was computed using the general derivatives, Eq 2.5. For a single component system the S-S yielded PVT relationships, Eqs 2.23 and 2.24. Initially the theory has been used to compute PVT-behavior of homopolymers. Later, it was extended to full thermodynamic description of multicomponent systems, viz. thermodynamics of liquid mixtures, gas-liquid and liquid-liquid phase equilibria, etc. Thus, using the derived expression for free energy of neat components, the molar Helmholtz free energy was derived for a binary mixture [Jain and Simha, 1980; 1984; Stroeks and Nies, 1988]:

<s> = x1s1 + x2s2 ; = x1c1 + x2c2 ; <Mo> = (x1s1Mo1 + x2s2Mo2) / <s> * *k * *k * *k <ε*>k = X12 ε11 v11 + X22 ε22 v22 + 2X1X2 ε12 v12; k = 2,4 (2.46)

where X1 = 1 - X2 = x1[s1(z - 2) + 2] / In Eq 2.46 there are two cross-interaction parameters that characterize the binary system: 1. the interaction energy, ε*12 = Sε (ε*11 ε*22), and 2. the repulsion volume, v*12 = Sv[(v11*1/3 + v22*1/3) / 2]3. For nearly athermal systems, the proportionality factors, Sε and Sv, are taken as equal 1. Thus, for the systems without strong interactions the binary parameters were well approximated by the geometric and algebraic averages. For example, for PS/PVME blends the assumption: Sε = Sv = 1 resulted in 0.1% deviation of the experimental values of the cross-parameters [Xie et al., 1992; Xie and Simha, 1997]. It is to be expected that for systems with strong intermolecular interactions experimental values of these factors may have to be found. However, the least squares fit of Eqs 2.45-2.46 to experimental values of CO2 miscibilities in PS (in a wide range of P and T) showed that the best values of Sε and Sv are close to 1 [Xie et al., 1997]. The S-S equations were used for the description of PVT dependencies for PS, PPE and their blends [Jain et al., 1982]. The data are presented in Figure 2.6, where the best fit values for each composition: Sε(P, T, φ) and Sv(P, T, φ) are plotted vs. PPE volume fraction, φ(PPE). The plot

Fm / RT = x1 ln x2 + x2 ln x2 + (<s> / y)(1 - y) In(1 -y) - (<s> - 1) ln[(z - 1) / e] - {ln[(1 - η)3 / Q] - (yQ2 / 2T)(AQ2 - 2B)} - (3/2){x1c1 ln[2π < Mo1 > RT(NAh)-2] + x2c2 ln[2π < Mo2 > RT(NAh)-2]}

(2.45)

142

L. A. Utracki

show concentration dependence of these two parameters, not observed previously for mixtures of solvents or for polymer solutions. The variation is not large — the mean averages of Sε and Sv are, respectively, 1.15 ± 0.06 and 0.86 ± 0.04. From the theoretical point of view it is highly satisfying that these values are not far from unity. Furthermore, the larger (than 1) value of Sε and smaller (than 1) value of Sv indicate presence of the specific interactions between PS and PPE — larger than average energetic interactions and a contraction of the binary volume contribution. It should be stressed that as shown by the original authors, the Sε and Sv values provided excellent description of the PVT dependencies in the full range of variables. The observed compositional variability may be caused by inadequacy of: the theory, the computational method, or the experimental data. Analysis of these and other blends by means of the S-S EoS is being pursued.

and “2”) with holes (subscript “0”). The derived equations were used to describe the vapor-liquid equilibrium of n-alkanes. They also well predicted the phase separation of n-alkanes/PE systems. The lattice-gas model gives the non-combinatorial free energy expressed in terms of composition and the surface interacting parameters, γi, [Kleintjens et al., 1980; Nies et al., 1983]: ∆Gm / VRTφ1φ2 = {α1 + χL11[(1 − γ1) / (1 − γ1φ1 − γ2φ2)]}(φο / φ2) + {α2 + χL22[(1 − γ2) / (1 − γ1φ1 − γ2φ2)]}(φο / φ2) + {αm + χLm[(1 − γ2) / (1 − γ1φ1 − γ2φ2)]} where

γi = 1 − si / so

(2.47)

In these relations the super-script L is used only to indicate that the relation was derived from the gas-lattice model. The parameter, si, is the interacting surface of species “i”, and the binary interaction parameters χii depend only on temperature. According to the gas-lattice theory, four factors determine the polymer/polymer miscibility [Koningsveld et al., 1982: Koningsveld, 1986]: 1. Interacting surface areas of segments. 2. Coil dimensions (dependent on T, φ, and MW). 3. Molecular weight polydispersity. 4. Free volume fraction.

2.5.1.3 Gas-lattice Model The gas-lattice model considers liquid to be a binary mixture of randomly distributed, occupied and vacant sites. P and T can change the concentration of holes, but not their size. A molecule may occupy m sites. Binary liquid mixtures are treated as ternary systems of 2 liquids (subscripts “1”

PS/PPE blends T = Tg — 330°C P = 0 — 180 MPa

Se, or Sv

1,3

<Se> = 1.15 ± 0.06

1,1

0,9 <Sv> 0.86 ± 0.04

0,7 0

0,4

f(PPE)

0,8

Figure 2.6. The binary interaction parameters of the Simha-Somcynsky EoS, Sε and Sv, for blends of PS with PPE: T = Tg to 330°C and P = 0 to 180 MPa [Jain et al., 1982].

Thermodynamics of Polymer Blends

During the last several years the lattice cluster theory (LCT) was developed [Dudowicz et al., 1991; Dudowicz and Freed, 1991, 1993, 1995, 1996; Freed and Dudowicz, 1995, 1996]. LCT is based on the lattice model. In comparison to H-F theory it incorporates two modifications: (1) it includes local packing and interactions, and (2) it distinguishes different structures of the monomeric units, allowing the monomeric units to occupy different adjacent sites, dictated by the structure. The theory represents polymer chains as strings of freely jointed by flexible bonds beads, called “united atoms.” For example, it considers CHX (x = 0, 1, 2, or 3) as a single bead (or group) that occupies one lattice site. LCT also incorporates vacant sites (free volume) and uses the non-random mixing principle. Thus, the free energy of mixing is given by: ∆Fm / N1k B T =

∑ (φ

/ Mi ) ln φi + φ v ln φ v +

i

i =1

∑ g (φ )φ φ + ∑ g ({φ })φ φ i

i =1

i

i v

ij

k

i j

+ ...

i≠ j

(2.48) The first term of Eq 2.48 is the H-F combinatorial entropy. The second term (where the subscript “v” indicates the free volume fraction) represents the free volume contribution to the entropy of mixing. The third term represents the non-combinatorial contribution [gi(φi) is the non-combinatorial energy on a molten state of polymer “i” having the free volume fraction φV]. The fourth term represents the energetic contribution originating from interaction between unlikely species, i  j. Here gij({φk}) is the interaction term expressed as polynomial with coefficients that depend on the structure of the polymer chains. These coefficients are computed as double expansions in 1/z ( z is the lattice coordination number), and εab/kBT (εab is the van der Waals interaction energies between groups a and b).

143

Through the curve fitting method to experimental data it has been shown that the H-F binary interaction parameter, χ12, is a thermodynamic function of such independent variables as: T, P, φi, molecular weights, and others. LCT shows that χ12 has strong sensitivity to composition, monomeric structure, and local correlation. The theory was quite successful describing observed dependencies for PS/PVME blends using four parameters: εS/S; εVME/VME; εS/VME and the cell volume. Except for the hetero-contact parameter, εS/VME (which must be determined by a fitting procedure), the other three are determined from the PVT dependencies of neat resins. LCT (originally developed for di-block copolymers) was found to be particularly useful to explain miscibility of polyolefin blends where the two resins differ in the type and size of short chain branching. The structural units of a polymer with two carbons in the main chain can be written as: PE = (CH2-CH2)n, PP = [CH2-CH (CH3)]n, poly2-butene (P2B) = [CH (CH3)-CH (CH3)]n, PIB = [CH2-C (CH3)2 ]n, poly(4,4-dimethyl 1-butene) (PDMB) = [CH2-CH (C4 H9)]n, etc. Three structural parameters (ratio of end to interior groups) have been used to distinguish PO structure: r, p, and q. Their values for the model macromolecules discussed above are listed in Table 2.7. The miscibility is expected for blends of PO’s having similar value of the structural parameters. In Table 2.8 examples of the computed binary interaction parameters for 1:1 composition of PO blends at 500 K are shown [Freed and Dudowicz, 1996]. The experimental values of these parameters have also been measured [Balsara et al., 1992; 1994; Graessley et al., 1993; 1994 a, b; 1995; Krishnamoorti et al., 1994 a, b; 1995; 1996; Lin et al., 1996; Schipp et al., 1996; Reichart et al., 1997]. The data were determined using either the small angle neutron scattering (SANS), cloud point curve determination (CPC), or PVTmeasurements. The results will be discussed later, viz. SANS measurements.

144

Table 2.7.

L. A. Utracki

Structural parameters for model PO [Dudowicz and Freed, 1996]

Polyolefin

Code

Polyethylene Polyethylenepropylene Polypropylethylene (polypentene-1) Polyethylethylene (polybutene-1) Polypropylene Poly-2-butene Poly(4,4-dimethyl 1-butene) Polyisobutylene

Table 2.8. Blends PE/PIB PEP/PIB PE/P2B PP/PIB PP/PE PEP/P2B P2B/PIB PE/PEP PP/P2B PEP/PP

2.5.2

PE PEP PPE PEE PP P2B PDMB PIB

r

p

q

1 1.2 1.2 1.25 1.33 1.5 1.67 1.75

0 0.2 0.2 0.25 0.33 0.5 0.5 0.5

0 0.25 0.25 0.33 0.5 1 1 1

Binary interaction parameters: χ total = χS + χH / T , for model PO 1:1 blends at 500 K [Freed and Dudowicz, 1996] Tc (K)

r

1000χS

χH

100χtotal

488 477 432 395 383 365 351 340 328 311

0.75 0.55 0.5 0.417 0.333 0.3 0.25 0.2 0.167 0.133

-0.3248 0.4560 1.877 1.840 1.356 1.327 1.049 0.6486 0.5463 0.3249

5.8188 2.8443 1.6352 0.9062 0.5150 0.3259 0.2119 0.1249 0.0569 0.0424

1.13 0.615 0.510 0.354 0.229 0.187 0.135 0.083 0.061 0.037

Off-lattice Theories

where U1 and U2 are the attractive and repulsive energies, respectively, and q is the degeneracy number. Depending on the relative magnitude of U1 and U2, Eq 2.49 predicts either UCST or LCST.

2.5.2.1 Strong Interactions Model For incompressible systems having strong interactions, e.g., acid-base type, the directional-specific model of segmental interactions may be used [Walker and Vause, 1982; ten Brinke and Karasz, 1984]. The following expression was derived: ¨Gm / VRT φ1φ2 = (U2 / RT) + ln (1-λ) + (1+q-1) 1/λ = 1+q × exp{(U1-U2) / RT)}; 0 ” λ ” 1 (2.49)

2.5.2.2 Heat of Mixing Approach When the configurational entropy of a polymer blend is vanishingly small, the enthalpic effects dominate, hence the adiabatic calorimetry should be able to predict polymer/polymer miscibility [Cruz et al., 1979)].

145

Thermodynamics of Polymer Blends

¨Gm ≈ ¨Hm = Bφ1φ2; Β ≡ χ12RT(V / V1)

(2.50)

B = B13φ′1 + B23φ′2 - B12φ′2φ′1, or

After experimentally confirming validity of this idea, the principal authors attempted to use this approach for explanation of the so call “miscibility windows” [Paul and Barlow, 1984]. The latter term refers to either polymer/copolymer blends that show miscibility only within a limited rage of the copolymer composition [viz. Balazs et al., 1985; Fernandes et al., 1986; Goh and Lee, 1987], or to blends of two copolymers having a common monomer [Shiomi et al., 1986]. As in earlier Koningsveld’s treatment of χ12, here also the parameter B has the enthalpic and the non-combinatorial entropic contribution. For multicomponent systems Eq 2.50 can be generalized [Barlow and Paul, 1987]: ∆Gm / V ≈ ∆Hm / V =

∑∑ B φ φ

ij i j

∑Ψ ∑∑ B φ φ

k k ij i j

k

k

i

−

j ≠i

i

;

j ≠i

∑φ

i

≡ 1;

i

¨Gm / V ≈ ¨Hm / V = BΨ1Ψ2, where:

∑φ

k i

≡1

i

(2.51) where Ψk is the volume fraction of polymer k. For simple systems containing copolymer A (mers #1 and #2) and either polymer B (mers #3) or a copolymer B (mers #1 and #3), Eq 2.51 can be simplified to read, respectively:

B = B12φ2(φ′2 - φ′′3) + B13φ3(φ′′3 - φ′2) + B23φ′2φ′′3 (2.52) where φ′i and φ′′i denote the volume fraction of unit “i” in copolymers A and B, respectively. Over the years, values of the parameter B for many polymer mixtures were published. As shown in Table 2.9, similarly to χ12, also B varies with composition, temperature, and other independent variables. Determination of B for a series of blends made it possible to establish empirical rules for the observed miscibilities. Thus, for example, Prud’homme [1982] reported systematic variation of miscibility in a series of halogenated polymer/ aliphatic polyester blends. The highest miscibility was observed for the chlorinated polymer, e.g., PVC, and the lowest for the fluorinated, PVF, with the brominated, PVB, having an intermediate behavior. Furthermore, when the chlorinated polymer was mixed with a series of polyesters, the highest miscibility was found when the ratio of the -CH2- to -COO- groups reached an optimum value. This optimum value depended on the chemical nature of the halogenated polymer — as shown in Figure 2.7, for PVDC blends with aliphatic polyesters this value is 5 to 6.

PVDC / Aliphatic polyester blends

B (cal/mL)

2

0

-2

2

4

6

8

10

CH2/COO group ratio Figure 2.7. The binary interaction parameter B for PVDC/aliphatic polyester blends plotted as a function of the number of -CH2 groups per one ester -COO- group in the latter polymer. Data from [Woo et al., 1986].

146

L. A. Utracki

c 12

0

-0,004

PIB1/EB85: c = 0.019385 -6.3646/T PIB1/EB114: c = 0.023163 -8.306/T

-0,008

PIB2/EB114: c =0.022809 -8.1355/T PIB1/EB73: c =0.015013 -5.1485/T PIB1/hhPP: c =0.025084 -10.47/T

-0,012 0,0022

0,0026

0,003

0,0034

1/T (K) Figure 2.8. The SANS binary interaction parameters for two PIB resins (MW = 81.6 and 114 kg/mol) with either ethylene-butene copolymers (of different MW and composition) or with an atactic head-to-head PP. After [Krishnamoorti et al., 1995].

0,006

d-PP/hhPP: c = -0.00639 + 3.305/T d-hhPP/PEB: c = -0.00137 + 1.011/T d-hhPP/PEP: c = -0.00036 + 0.517/T d-PP/EB97: c = -0.00883 + 4.200/T d-PP/EB78: c = -0.00320 + 1.685/T

c 12

0,004

0,002

0 0,0022

0,0026

0,003

0,0034

1/T (K) Figure 2.9. Temperature dependence of the binary interaction parameter, χ12, for blends of deuterated amorphous polypropylene (either head-to-tail or head-to-head, hh) with polyethylene copolymers. The values were determined by SANS. Data [Graessley et al., 1995].

Thermodynamics of Polymer Blends

In Table 2.9 along the parameters χ12 and B, the available values of ¨P* are also listed. The latter parameter originates from S-L EoS, and it has been considered a measure if the interaction energy, proportional to χ12 or B [Sanchez, 1989]. All these parameters, being proportional to each other, are expected to show similar variability with the independent variables. For example, in Figures 2.8-2.9 the variability with composition and temperature of χ12, and in Figure 2.10 these of ¨P*, is shown.

147

For example, taking the case illustrated in Figure 2.7, one may consider that PVDC is an alternate copolymer of units -CH2- and -CCl2-, whereas the aliphatic polyesters are composed of units -CH2- and -COO-. Eqs 2.51-2.52, predict that even systems with all positive values of the binary interaction parameter, Bij > 0, (repulsive interactions) may have a window of miscibility, where the overall parameter, B < 0. The magnitude of this effect depends primarily on the value of the repulsive interactions within the copolymer molecule, B12 > 0. Schematic representation of Eq 2.52 is shown in Figure 2.11. The idea of Paul and Barlow [1984] was ingeniously simple and far reaching. If Bij characterizes the enthalpic and non-configurational entropy of interaction between segments “i” and “j”, then one could tabulate these parameters, and in turn use them for predicting miscibility in any arbitrary system composed of polymers and/or copolymers built from these segments. Table 2.10 provides examples of Bij values published over the years. As in the original publications the units of Bij (cal/mL), are retained (to convert these to Joule/m3, multiply the listed values by 4.187x106).

Figure 2.10. (top) Compositional variation of the interaction parameter, ¨P*, in Phenoxy/ polyvinylmethylether blends at the spinodal temperature, Ts. (bottom) Temperature dependence of ¨P*. After [Etxeberria et al., 1997].

The observed regularity in the miscibility behavior for a series of polymers can be better understood considering that two polymers are composed of individual interacting groups.

Figure 2.11. Compositional variation of the interaction parameter, B, in a copolymer (mers #1 & #2)/ homopolymer (mers #3) blends. Even when all three binary interaction parameters are positive (B23 = 1, B13 = 2, and B12 = 0 - 10), for a copolymer with strong repulsive interactions, B12 = 10, a “window of miscibility” is predicted. See text.

148

Table 2.9.

L. A. Utracki

Binary interaction parameters: B, ∆P* or χ12. (see earlier data in [Utracki, 1989]) Note: To convert cal/mL to Joule/m3, multiply the listed values by 4.187x106

Polymer-1

Polymer-2

T(°C)

PS PS PS PS 50 wt% PS PS PS PS PS PS 50 wt% Pα-MS Pα-MS PS PS PS PS PS PS PS PS Pα-MS Pα-MS Pα-MS Pα-MS Pα-MS Pα-MS Pα-MS Pα-MS PMMA PMMA PMMA PMMA PMMA PMMA PMMA PMMA PS PS PS PS PS PS

TMPC TMPC Pα-MS Pα-MS PMMA PMMA PMMA PMMA PMMA PMMA PMMA PMMA PC DMPC TMPC TCPC PCZ HFPC BCPC TMPC-P PC DMPC TMPC TCPC PCZ HFPC BCPC TMPC-P PC DMPC TMPC TCPC PCZ HFPC BCPC TMPC-P PSF DMPSF TMPSF HMBIPSF PES HFPSF

30 300 50 160 245 153 250 195 150 250 50 50 240 50 50 300 50 180 50 50 180 300 200 50 300 180 50 196 214 300 50 160 150 235 248 300 228 300 300 300

B(cal/mL)

∆P* (cal/mL)

- 0.13 - 0.17 ± 0.01 0.11 - 0.17 ± 0.01 0.012-0.025 0.011 - 0.025 - 0.0833 + 0.001034 T 0.542 0.620 0.464 0.532 0.457 0.520 0.392 0.455 0.21 ± 0.02 0.24 ± 0.05 0.542 + 10 -4 T 0.354 0.458 0.43 0.44 0.20 - 0.49 0.036 - 0.17 > 0.72 0.28 > 14 > 1.6 1.5 > 0.33 0.16 0.023 0.39 - 0.44 0.42 - 0.49 > 0.18 0.26 0.068 > 0.31 > 0.24 0.22 - 0.72 0.12 - 1.4 > 0.33 > 0.44 0.21 - 0.29 0.001 - 0.006 0.057 -0.066 0.043 0.16 0.29 0.40 0.45 - 0.77 0.17 - 0.072 - 0.30 0.01 - 0.077 0.22 0.31 0.85 1.18 > 0.67 0.31 0.33 > 1.25 - 1.27 > 136 -

χ12 (-)

Ref. 1 1 2 2 2 2 2 2 2 2 2 2 3 3 3 3 3 3 3 3 3 3 3 3 3 3 3 3 3 3 3 3 3 3 3 3 4 4 4 4 4 4

149

Thermodynamics of Polymer Blends

Table 2.9.

Continued

Polymer-1

Polymer-2

PS PS PS Pα-MS Pα-MS Pα-MS Pα-MS Pα-MS Pα-MS Pα-MS Pα-MS Pα-MS PMMA PMMA PMMA PMMA PMMA PMMA PMMA PMMA PMMA PAN PET PBT PBT PBT PBT PBT PBT PBT PBT PBT PBT PS PS PS PB PB PB PMMA PMMA PiBMA PS

TMHFPSF TMHFPSF TMPSF-P PSF PSF DMPSF TMPSF HMBIPSF PES HFPSF TMHFPSF TMPSF-P PSF DMPSF TMPSF HMBIPSF PES HFPSF TMHFPSF TMHFPSF TMPSF-P PSF PETG 20 % PAr 40 % PAr 60 %PAr 80 %PAr 65 wt % PEE 70 wt % PEE 75 wt % PEE 80 wt % PEE 85 wt % PEE 90 wt % PEE 75% Pp-MS 50% Pp-MS 25% Pp-MS 50% d-PB 50% d-PB 43% d-PB 50% PnBMA 50% PiBMA 50% PnBMA PCSt

T(°C) 50 240 174 50 300 300 50 300 300 50 300 300 50 300 231 300 300 50 50 293 300 50 280 249 246 236 232

140 140 140 130 121 130 25 25 25 30

B(cal/mL) 0.63 1.12 0.34 0.32 > 0.43 > 0.30 > 0.35 > 0.29 > 0.31 > 0.20 > 0.30 > 0.29 0.25 - 0.34 > 0.77 0.39 > 0.76 > 0.78 0.10 - 0.15 0.25 0.77 > 0.76 3.6

∆P* (cal/mL)

χ12 (-)

0.37 > 0.53 > 0.36 0.19 - 0.27 0.44 4.1 - 0.12 - 0.65 - 0.40 - 0.31 - 0.22 0.13 0.10 0.07 0.048 0.032 0.012 - 0.008 - 0.006 - 0.006 0.00053 0.00055 0.00032 0.081 0.068 0.0019 0.07 ± 0.02

Ref. 4 4 4 4 4 4 4 4 4 4 4 4 4 4 4 4 4 4 4 4 4 4 5 6 6 6 6 7 7 7 7 7 7 8 8 8 9 9 9 10 10 10 11

150

Table 2.9.

L. A. Utracki

Continued

Polymer-1

Polymer-2

PVDF PVDF PVDF PVDF PC PC PCEMA PCEMA PS Pα-MS Pα-MS BR BR PVC PVAl PA-6 PA-6 PA-6 PVDF PVDF PEMA PA-6 PA-6 PA-6 PA-6 PS PVC PI PHS PHS PHS PCL PCL PC PVDC PVDC PVDC PVDC PVDF PVME PVME PB PSiαMS

PMMA MMA-GMA8 MMA-GMA14 MMA-GMA28 PNP PMS PCL PHS PAN PAN PS PAN PVC PAN PAA Zn-SPS Li-SPS complex PS PMMA PEMA PMMA Mn-SPS, 10% Mn-SPS, 20% Mn-SPS, 30% Mn-SPS, 10% PPE PCL 50% d-PB 50% PCEMA PCMMA PVC SAN(25AN) PC SAN(25AN) PDPS PDPA PCL PCDS PBA d-PS 50 wt% d-PS 70.6 wt% d-PB PS 50 wt%

T(°C) 170 170 170 170 200 200

B(cal/mL) -2.93 -4.1 -4.5 -3.7

4.59 6.02 0.022±0.001 8.60 0.72 ± 0.07 3.84 ± 0.43 - 19.9 220 240 240 160 160 160 190 180 175 166 150 250 150

175 100-150 100-150 -50-80 100-200

-215 28.7 - 4.43 - 2.66 3.25 - 1.9 - 2.0 - 1.8 - 1.5

- 4.67 - 8.39 - 3.8 - 0.61 - 0.39 0.2 ± 0.3 - 0.2 - 1.1 - 2.0 - 3.1 - 1.0

∆P* (cal/mL)

χ12 (-)

Ref.

12 12 12 12 0.175 13 0.031 13 - 0.99 14 - 0.48 14 15 15 16 17 17 17 - 1.24 18 - 1.3 19 20 20 21 21 21 22 22 22 22 - 0.2 23 - 0.5 23 0.0023 24 - 0.76 25 - 1.2 25 - 0.76 26 27 27 27 viz. Fig. 2.5 28 viz. Fig. 2.5 28 viz. Fig. 2.5 28 viz. Fig. 2.5 28 - 0.19 29 0.0702-30.9/T 30 0.0817-36.8/T 30 0.5 to 1.29, resp. 30 0.0032-5.46/T 31

151

Thermodynamics of Polymer Blends

Table 2.9.

Continued

Polymer-1

Polymer-2

PαMS PMMA P4VP P4VP P4VP PIB-MW=82k PIB-MW=82k PIB-MW=160k PIB-MW=82k PIB-MW=82k d-PP d-PP d-PP d-hhPP d-hhPP PVDF P(VF2-co-VF4) HDPE HDPE HDPE PEP 25 wt% PEP 57.5 wt% PEP 89.1 wt% PEMA 70 wt% PEMA 50 wt% PEMA 30 wt% PCL PVP PVP PVP PVP PMMA PMMA PEG PA-6 PEEK PP PCl (high φ) PCl (low φ) PCl (high φ) PCl (low φ) PS PS

PS 50 wt% PS PS PS PS EB- MW =85k EB-MW =114k EB-MW =114k EB- MW =73k hhPP-MW=27.5k hhPP EB-97 EB-78 PEB PEP PMA PMA LDPE LDPE LDPE PEB PEB PEB CR CR CR P4HS CDA 90 wt% CDA 65 wt% CDA 40 wt% CDA 15 wt% Phenoxy PEG Phenoxy MXD PEI SEBS + oil PVDC PVDC P(VCl2-VAc) P(VCl2-VAc) PPE PPE

T(°C)

B(cal/mL)

∆P* (cal/mL)

180-300 165 180 183 (see Fig. 2.6) (see Fig. 2.6) (see Fig. 2.6) (see Fig. 2.6) (see Fig. 2.6) (see Fig. 2.7) (see Fig. 2.7) (see Fig. 2.7) (see Fig. 2.7) (see Fig. 2.7) 160 160 150 180 190 27-167 27-167 27-167 42 40 39 50 24.5 24.5 24.5 24.5 170

275 180 160

210 210

- 0.89±0.04 - 0.31±0.15

χ12 (-) 0.0044-0.0046 0.006-0.022 7.5 ± 2.5 3.5 ± 1.5 0.4 0.0194-6.36/T 0.0232-8.306/T 0.0228-8.14/T 0.0151-5.149/T 0.0194-6.36/T -0.00639+3.305/T -0.00883+4.200/T -0.00320+1.685/T -0.00137+1.011/T -0.00036+0.517/T -0.221 ± 0.002 -0.005 ± 0.0005 0.000402±4x10 -5 0.000390±4x10 -5 0.000387±4x10 -5 -0.00167+0.954/T -0.00143+0.883/T -0.00219+1.138/T - 0.122 - 0.053 - 0.030 - 1.1 - 4.20 - 1.64 - 0.60 - 0.36 - 0.61 - 0.35 - 1.90 - 0.185/ - 0.194 - 0.3 - 0.043 - 0.02 - 0.21 - 0.01 - 0.28

Ref. 31 32 33 33 33 34 34 34 34 34 35 35 35 35 35 36 36 37 37 37 38 38 38 39 39 39 40 41 41 41 41 42 43 44 45 46 47 48 48 48 48 49 50

152

Table 2.9.

L. A. Utracki

Continued

Polymer-1

Polymer-2

PS PS PS PS PS PS PES-C

PPE PPE P1,4PE PC TMPC PAN PEG

∆P* (cal/mL)

T(°C)

B(cal/mL)

232 210 30 30 30 30 65

- 1.62±0.07 - 1.46 + 0.00238T - 0.005±0.43 0.41 ± 0.13 0.19 ± 0.34 7.63 ± 0.12 4.65

χ12 (-)

0.121 - 77.9/T

Ref. 51 52 53 54 54 54 55

Abbreviations for Table 2.9: BCPC - bisphenol chloral polycarbonate, BR - polybutadiene, CDA - cellulose diacetate, CR - polychloroprene, d-hhPP - deuterated head-to-head amorphous polypropylene, DMPC - dimethyl bisphenol-A polycarbonate, DMPSF - dimethyl bisphenol-A polysulfone, d-PB - deuterated polybutadiene, d-PP - deuterated amorphous polypropylene, EB - ethylene-butene copolymer (with indicated butene content or MW = xx in kg/mol), HFPC - hexafluoro bisphenol-A polycarbonate, HFPSF - hexafluoro polysulfone, hhPP - head-to-head amorphous polypropylene, HMBIPSF - hexamethyl biphenol polysulfone, Li-SPS - Li-sulfonated PS, Mn-SPS - Mn-sulfonated PS, MXD - poly(m-xyleneadipamide), P1,4PE poly(1,4-phenylene ether), P4HS - poly(4-hydroxystyrene), P4VP - poly(4-vinylpyridine), Pα-MS - poly(α-methyl styrene), PAN - polyacrylonitrile, PBA - poly(1,4-butylene adipate), PC - bisphenol-A polycarbonate, PCEMA poly-1-chloroethyl methacrylate, PCDS - poly(1,4-cyclohexanedimethylene succinate, PCMMA poly-1-chloromethyl methacrylate, PCSt - poly-pchlorostyrene, PCZ - bisphenol-Z polycarbonate, PDPS - poly(2,2-dimethyl-1,3-propylene succinate, PDPA - poly(2,2-dimethyl1,3-propylene adipate, PEB - polyethylene-butene, PEP - polyethylene-propylene, PEE - polyester-ether aromatic block copolymer, PES - polyethersulfone, PES-C - phenolphthalein polyetherether sulfone, PHS - polyhexamethylene sebacate, PI - polyisoprene, PMS - poly(methylmethacrylate-co-pmethylstyrene), PNP - poly(methylmethacrylate-co-N-phenylmaleimide), Pp-MS - poly-paramethylstyrene, PPE - poly(2,6-dimethyl-1,4-phenylene ether), PSF - bisphenol-A polysulfone, PSiαMS - poly(sila-α-methylstyrene), SAN(25AN) - copolymer of styrene with 25% acrylonitrile, TCPC - tetrachloro bisphenol-A polycarbonate, TMHFPSF tetramethylhexafluoro polysulfone, TMPC - tetramethyl bisphenol-A polycarbonate, TMPSF - tetramethyl bisphenol-A polysulfone, TMPC-P tetramethyl bisphenol-P polycarbonate, TMPSF-P tetramethyl bisphenol-P polysulfone, Zn-SPS -Zn-sulfonated PS. References for Table 2.9: 1. Kim and Paul, 1992; 2. Callaghan and Paul, 1993 ; 3. Callaghan and Paul, 1994a ; 4. Callaghan and Paul, 1994b; 5. Papadopoulou and Kalfoglu, 1997; 6. Huo and Cebe, 1993; 7. Gallagher et al., 1993; 8. Londono and Wignall, 1997; 9. Alamo et al., 1997; 10. Sato et al. 1996; 11. Ogawa et al., 1986; 12. Gan and Paul, 1995; 13. Ikawa and Hosoda, 1991; 14. Peng et al., 1994; 15. Cowie et al., 1992a; 16. Cowie et al., 1992b; 17. Cowie and Harris, 1992; 18. Dinililuc et al., 1992; 19. Lu and Weiss, 1992; 20. Molnar and Eisenberg, 1992; 21. Goh and Siow, 1988; 22. Lu and Weiss, 1991; 23. Lu and Weiss, 1992; 24. Hasegawa et al., 1991; 25. Neo and Goh, 1992; 26. Woo et al. 1985; 27. Shah et al. 1986; 28. Woo et al. 1983; 29. Penning and Manley, 1996; 30. Takeno et al. 1996; 31. Maier et al. 1996; 32. Pinder, 1997; 33. Clarke et al. 1997; 34. Krishnamoorti; 35. Graessley et al. 1995; 36. Maiti and Nandi, 1995; 37. Schipp et al. 1996; 38. Lin et al. 1995; 39. Kundu et al. 1996; 40. Lezcano et al. 1996; 41. Jinghua et al. 1997; 42. Hong et al. 1997; 43. Martuscelli et al. 1984; 44. Iriarte et al. 1989; 45. Shibayama et al. 1995; 46. Goodwin and Simon, 1996; 47. Ohlsson and Törnell, 1996; 48. Aubin et al., 1983; 49. ten Brinke et al., 1983; 50. Kambour et al., 1980; 51. Plans et al., 1984; 52. Maconnachi et al., 1984; 53. Ziaee and Paul, 1996; 54. Ziaee and Paul, 1997; 55. Zheng et al., 1997.

In Table 2.10, besides Bij, two other measures of the binary interactions between segments, χij and ¨P*ij, are also given. The former dimensionless

parameter is directly proportional to Bij. The latter represents the binary interactions energy density calculated from S-L EoS. Since ¨P* is considered

153

Thermodynamics of Polymer Blends

proportional to χ12 or B [Sanchez, 1989], one may apply S-L theory to the experimental data and compute the segmental contributions, ¨P*ij, in analogy to Bij: *

*

*

*

∆P = ∆P12 φ′2 (φ′2 − φ′′3) + ∆P13φ′′3(φ′′3 − φ′2) + ∆P23φ′2φ′′3 (2.53) The segmental interaction parameters have been also used to explain the “miscibility window” or “miscibility chimney” in polymer/copolymer or copolymer/copolymer blends [Lath and Cowie, 1988]. These parameters have been found useful to predict miscibility of blends containing one component whose structure systematically varied, e.g., polyesters with either halogenated polymers or Phenoxy [Prud’homme, 1982; Harris et al., 1983; Woo et al., 1985; Woo et al., 1986], polyamide blends [Ellis, 1989], ternary blends [Shah et al., 1986] and other systems, viz. SAN/ PMMA, SAN/PC, polyethyloxazoline/polyester, PPE with a mixture of P(oClS) and P(pClS), PC/PCL/Phenoxy, and many more.

Ellis [1988-90] used the same approach to evaluate miscibility of polyamide blends. He treated their molecules as copolymers, viz. AxByC1 - x - y, where: A = -CH2-, B = -NH-CO-, and C = -C6H4is a phenyl unit. The analysis made it possible to systematize the experimental observations and predict conditions of miscibility for aliphatic and semi-aromatic polyamides. The method was applied to 1:1 composition of blends and χ12 values were assumed to be temperature independent. Examples of the segmental interaction parameters, χij are given in Table 2.11. For binary mixtures of aliphatic polyamides AxB1-x with AyB1-y (where A and B groups are defined above, and subscripts x,y refer to their mer volume fraction), the binary interaction parameter of the blend was expressed as: χ12 = 7.984(x - y)2 • 0. Thus, these polymers are expected to be immiscible. The miscibility, if observed, may be explained by either presence of hydrogen bonding, or trans-reaction (transamidation) resulting in rapid homogenization of the system (e.g., PA-6/PA-46 at 310°C homogeneous in 4 min) [Ellis, 1992].

Table 2.10. Binary segmental interaction parameters, Bij, ¨P*ij or χij Note: To convert cal/mL to Joule/m3, multiply the listed values by 4.187x106 Unit 1

Unit 2

-CH2-CH2-COO-CH2-CH2-CH2S S AN -CH2-CH2-CH(C6H5)-CH(C6H5)-CH(CN)-CHx-CHx-CHx-

Phenoxy -COOPhenoxy S AN -COOAN -COO-COO-CH(C6H5)-CH(CN)-CH(CN)-COO-COO(C6H5)-CN -COO-

Bij (cal/mL) 8.34 250.3 154.9 5.72 97.5 136.5 122 103 104 8.07 234 277 96.8 170 10.4 508 136

¨P*ij (cal/mL)

χij

Ref. 1 1 1 2 2 2 2 2 2 2 2 2 2 2 2 2 2

154

L. A. Utracki

Table 2.10. Continued Unit 1

Unit 2

(C6H5)(C6H5)-CN VAc VAc VAc VC VC S S MA MA MA MA MA AN AN AN AN S S S S α-MS α-MS TMPC TMPC HFPC HFPC HFPC S S MA S S MA S MA S MA MMA S MMA

-CN -COO-COOC(VC) VC -CH2C(VC) -CH2MMA MA MMA EMA nPNM S AN MMA EMA nPMA S MMA EMA nPMA MMA MMA S S MMA S MMA TMPC MMA MMA MMA MA AN AN TMPC TMPC PPE PPE VC VC S

Bij (cal/mL)

¨P*ij (cal/mL)

χij

579 93.0 351

0.22 7.32 9.30 10.4 11.9 14.9 - 0.381 5.00 5.33 5.85 8.14 0.181 - 0.0361 - 0.0309 0.26 0.26 0.018 - 0.15 0.26 1.5 - 0.07, - 0.73 0.25 ± 0.04 0.18, 0.26 0.22 7.18 10.7 6.8 - 0.31 - 0.14 11.5 ± 0.2 ” - 0.37 14.6 ± 0.5 0.54 2.85 2.93

0.15

0.324 - 0.061 - 0.151 0.32 - 0.084 0.018

0.23 7.28 10.9 7.37 - 0.36 - 0.17 11.7 ” - 0.42 15.0 ± 0.3

0.17 0.27 1.01 0.042 0.15 -

Ref. 2 2 2 3 3 3 3 3 4 4 5 5 5 5 5 5 5 5 5 5 5 5 6 6 6 7 7 7 7 7 7 8 8 8 8 8 8 8 8 8 9 9 9

155

Thermodynamics of Polymer Blends

Table 2.10. Continued Unit 1

Unit 2

AN TMPAr TMPAr TMPAr PAr PAr S -pC6H4-Ot-BMA t-BMA t-BMA t-BMA t-BMA t-BMA t-BMA t-BMA t-BMA S S S S S S S S S MMA MMA MMA CHMA CHMA CHMA CHMA CHMA CHMA S (140-170°C) S (140-170°C) S (140-170°C) AN (140-170°C) AN (140-170°C) MMA (130-200°C) MMA (130-200°C) iBMA (130-200°C)

S PAr S AN S AN AN -pC6H4-COS (130°C) S (150°C) S (180°C) MAN (130°C) MAN (150°C) MAN (180°C) MMN (130°C) MMN (150°C) MMN (180°C) MAN (130°C) MAN (150°C) MAN (180°C) MMA (130°C) MMA (150°C) MMA (180°C) CHMA (130°C) CHMA (150°C) CHMA (180°C) MAN (130°C) MAN (150°C) MAN (180°C) MAN (130°C) MAN (150°C) MAN (180°C) MMA (130°C) MMA (150°C) MMA (180°C) MMA AN Cl MMA Cl nBMA iBMA nBMA

Bij (cal/mL)

¨P*ij (cal/mL)

χij

95.52 5.36 2.18 99.04 10.52 89.09 117.89 1.44 0.34 0.50 3.03 1.47 2.29 7.40 0.76 1.00 4.37 1.20 2.02 2.15 0.221 0.229 0.242 - 0.15 - 0.43 - 0.10 0.28 0.84 0.91 1.40 2.65 2.11 0.75 0.64 0.50 -0.635+287/T -11.0+4940/T -0.913+412/T -4.44+2000/T -4.76+2140/T -0.216+95.7/T -0.157+73.1/T 0.0179 - 10 -7 (846 T - T2)

Ref. 9 10 10 10 10 10 10 11 12 12 12 12 12 12 12 12 12 12 12 12 12 12 12 12 12 12 12 12 12 12 12 12 12 12 12 13 13 13 13 13 14 14 14

156

L. A. Utracki

Table 2.10. Continued Unit 1

Unit 2

S (80-130°C) S (80-130°C) Cl (80-130°C) Cl S AN VCVAc-90 VCVAc-90 VCVAc-90 VCVAc-90 iBMA SAN (75 % S) MAN (90 % MMA) MAN (90 % MMA) SAN (80 % S) MAN (95 % MMA) MAN (95 % MMA) - CH2- CH2- CH2-COO-CHCl-COOS S DNS DNS

AN Cl AN DTC DTC DTC iBMA iBMA-nBMA50 iBMA-nBMA13 nBMA nBMA SMMA (90 % MMA) SMMA (90 % MMA) SAN (75 % S) SMMA (95 % MMA) SMMA (95 % MMA) SAN (80 % S) -NHCO-COO-CHCl--NHCO--NHCO--CHClVME DNS VME PPE

Bij (cal/mL)

¨P*ij (cal/mL)

χij

Ref.

0.12 0.0077 0.049 0.001 0.02 0.042

15 15 15 16 16 16 17 17 17 17 17 18 18 18 18 18 18 19 19 19 19 19 19 20 20 20 20

-5.49 -2.22 -0.11 -1.73 0.00002 - 0.0134 0.0083 - 0.0108 - 0.0154 0.0021 - 0.011 8.534 2.233 0.500 3.880 6.750 0.038 0.0167 1.79 1.50 2.936

Abbreviations for the polymeric units in Table 2.10: (C6H5)- - phenyl ring, α-MS - alpha-methyl styrene, AN, acrylonitrile, BMA - butylmethacrylate, CHMA - cyclohexylmethacrylate, Cl - caprolactone, C(VC) - unit of chlorinated PVC, DNS - 2,4-dinitrostyrene-co-styrene, DTC -2,2-dimethyltrimethylenecarbonate, HFPC - hexafluoro bisphenol-A polycarbonate, MA - maleic anhydride, MMA - methylmethacrylate, PAr - unit of polyarylate, Phenoxy - unit of polyhydroxy ether of bisphenol-A, PPE - unit of poly(2,6-dimethyl-1,4-phenylene ether), S - styrene, TMPAr - unit of tetramethyl bisphenol-A polyarylate, TMPC - unit of tetramethyl bisphenol-A polycarbonate, VAc - vinyl acetate, VC - vinyl chloride, VCVAc90 - VC-co-VAc copolymer with 90 wt% VC, VME - vinylmethylether. References for Table 2.10: 1. Paul and Barlow, 1984; 2. Fernandes et al., 1986; 3. Shiomi et al., 1986; 4. Kim et al., 1989; 5. Brannock et al., 1991; 6. Callaghan and Paul, 1993; 7. Takakuwa et al., 1994; 8. Gan and Paul, 1994; 9. Dompas et al., 1997; 10. Ahn et al., 1997; 11. Harris and Robeson, 1987; 12. Nishimoto et al., 1995; 13. Higashida et al., 1995; 14. Sato et al., 1996; 15. Schulze et al., 1993; 16. Kammer and Kumerlowe, 1996; 17. Sato et al., 1997; 18. Cowie et al., 1992c; 19. van Ekenstein et al., 1997; 20. Fernandez et al., 1997.

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Thermodynamics of Polymer Blends

Table 2.11. Segmental interaction parameters, χij, for polyamide blends [Ellis, 1990b; 1993; 1995; 1997] χij

Segment-1

Segment-2

-CH2-CH2-CH2-CH2-CH2-CH2-CH2-CH2-NHCO-NHCO-NHCO-NHCO-NHCO-NHCO-CH2-NHCO-CH2-CH2-NHCO-C6H4-(m)C6H4-(m)C6H4-

-NHCO-CH2-NHCO-CH2- COO-C6H4(m)-C6H4-NHCO-C6H4-NHCO-NHCO-(p)C6H4-NHCO-NHCO-(m)C6H4-NHCO-C6H4-COO-(m)C6H4-NHCO-C6H4-NHCO-NHCO-(p)C6H4-NHCO-NHCO-(m)C6H4-NHCO-NHCO-C6H4-NHCO-COO-COO-COO-COO-NHCO-(m)C6H4-NHCO-

It has been known since the early days that behavior of the aromatic polyamides (aramids or PARA) critically depends on the type of isomeric substitutions — para resulted in crystalline, while meta in amorphous resins [Kwolek et al., 1962]. Similarly, the two PARA’s: poly(m-xylene adipamide) and poly(hexamethylene isophthalamide), MXD6 and PA-6I, respectively, show different miscibility, e.g., with aliphatic PA’s. Clearly, blind application of the segmental interaction strategy to aromatic or semi-aromatic polyamides could lead to conflicts. However, the problem can be resolved considering p and m substituted phenyl as two different statistical segments [Ellis, 1995]. This idea is indeed evident in the segmental contributions listed in Table 2.11. The segmental interaction values of χij from Table 2.11 were used to compute the binary interaction parameters, χ12, for blends of a semiaromatic polyamide, PA-3Me6T (Trogamid™ T)

7.984; 8.534 1.479 2.233 - 0.288; - 0.308 0.1 1.571 1.680 1.693 7.460; 7.974 3.880 8.000 2.275 2.432 2.445 - 0.083 2.233 3.880 1.692 1.500 1.680

with aliphatic polyamides, PA-4 to PA-12. These values are listed in Table 2.12, along with the observations of miscibility — qualitatively, an agreement was found [Ellis, 1989]. The data of Table 2.12 were used to construct Figure 2.12. Similarly as it was the case for the miscibility of halogenated polymers with polyesters, here also the miscibility depends on the CH2/NHCO ratio — the best miscibility was found when the group ratio was 7, e.g., for PA-410. Later, this analysis was extended to PA blends with polyesters [Ellis, 1993, 1995; 1997]. Thus, in the simplest case of PA blends with aliphatic polyesters, AxB1-x with AyD1-y (where: A = -CH2-, B = -NH-CO-, and D = -COO — is a carboxyl units, and subscripts x,y refer to their mer volume fraction) [Ellis, 1993]: χ12 = (y-x)(1-x) χAB + (x-y)(1-y) χAD + (1-x)(1-y) χBD

(2.54)

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L. A. Utracki

Table 2.12. Binary interaction parameters, χ12, for PA-3Me6T blends with aliphatic PA’s [Ellis, 1989] Aliphatic PA

Observation

PA-4 PA-46 PA-6 PA-66 PA-69 PA-610 PA-612 PA-11 PA-12

immiscible immiscible miscible miscible miscible miscible miscible immiscible immiscible

χ12 0.0400 0.0065 - 0.0050 - 0.0050 - 0.0069 - 0.0058 - 0.0023 0.0061 0.0104

PA-3Me6T blends with aliphatic PA's.

c12

0,04

0,02

0

2

6

10

CH2/NHCO ratio in aliphatic PA Figure 2.12. Blends of PA-3Me6T with aliphatic polyamides — the binary interaction parameter as a function of the -CH2 -/-NHCOgroup ratio for the latter polymers. Data [Ellis, 1989]. See text.

After 20 years, the heat of mixing (or analog calorimetry) method for the determination of polymer-polymer miscibility is becoming increasingly sophisticated. The low molecular weight analogs are selected on the basis of detailed calculations of the electrostatic charges on the atoms and molecules, using molecular orbital theories. The following principles have been formulated [Sandler and Wu, 1991; Ziaee and Paul, 1996; 1997]:

* The geometry of a functional group (in the polymer and selected analog molecule) should be the same. * The electrostatic charge of each atom in a functional group should be approximately the same. * The functional group should be approximately electroneutral. * The functional group should be the smallest entity, identified by dividing the analog molecule into a collection of electroneutral groups.

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Thermodynamics of Polymer Blends

Ziaee and Paul [1996; 1997] computed the electrostatic charges for a series of polymers and potential analogs. The heat of mixing was determined at T = 80°C and at a composition 1:1. A careful analysis of the data made it possible to determine the binary interaction parameters, Bij, listed in Table 2.13. From Eq 2.52, using the Bij values of Table 2.13, the authors calculated the (enthalpic) binary interaction parameters for several polymer pairs (see Table 2.9). The calculated values were in good agreement with the measured ones [Ziaee and Paul, 1996; 1997]. In summary, the heat of mixing approach and the resulting from it the segmental binary interaction parameters, uses several simplifying assumptions. The numerical values of the parameters vary depending on the initial assumption, method of evaluation, selected system, concentration, temperature, etc. However, the method

has been found useful for the identification of potentially miscible systems and conditions. Furthermore, the approach provided valid tool for interpretation of the blends’ behavior at higher concentration, viz. 50:50 composition, where the interaction parameters are relatively insensitive to variability of concentration. Under these conditions the segmental binary interaction parameters can be used to describe: * miscibility in systems without strong specific interactions, * behavior of blends of a homopolymer with copolymer, * miscibility of polymer series (chemically similar polymers, e.g., PA’s), * window of miscibility in two or three component systems, * chimney of miscibility in two or three component systems.

Table 2.13. New binary interaction parameters, Bij, at 0.5 volume fraction [Ziaee and Paul, 1996; 1997] Group-1

Group-2

C6H5C6H5C6H5C6H5-CH2-C6H4-O-C6H4-OC6H5-CH2C6H5-CH2C6H5-CH2C6H5-OC6H5-CH2C6H5-CH2C6H5-CH2-

-CH2-CH2-CH2C6H5-OC6H5-OC 6H 5 -CH2-C6H2(CH3)2-O-C6H2(CH3)2-O-CO-CO-CO-CO-CO-C6H4-O-CO-O-C6H4-C6H4-O-CO-O-C6H4-C6H4(CH3)2O-CO-O-C6H4(CH3)2-C6H4(CH3)2O-CO-O-C6H4(CH3)2-CH2-CN -CH2-CN

T (°C)

Bij (cal/mL)

30 80 130 80 80 80 80 80 80 80 80 90 90 90 90 90 130 130 30 30

8.74 ± 0.16 7.74 ± 0.16 5.93 ± 0.20 - 2.11 ± 0.30 5.44 ± 0.15 - 0.598 ± 0.450 7.24 ± 0.26 - 0.912 ± 0.143 3.13 ± 0.18 - 36.9 ± 18.4 - 19.1 ± 15.0 15.1 41.4 - 8.97 0.55 ± 0.21 10.1 ± 0.2 0.40 ± 0.06 4.56 ± 0.06 19.1 ± 0.2 60.8 ± 0.5

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L. A. Utracki

2

The concept of the solubility parameter originates from Hildebrand’s work on enthalpy of regular solutions [Shinoda, 1978]. Accordingly, in a strict sense, the molecular interactions should be nonspecific, without forming associations or orientation, hence not of the hydrogen or polar type. Another fundamental assumption was that the intermolecular interactions 1-2, are geometric mean of the intramolecular interactions, 1-1 and 2-2: v

v

v

2∆E12 ∆E1  ∆E ∆Hm / Vφ1φ2 =  1 + = 1/2 1/2 V1 V1 V2 V2  v 1/2

∆E1   V 

v

−

1

∆E2   V  2

1/2

(2.55)

2

 = (δ − δ )2 ≥ 0 1 2 

v

where ∆E1 is the molal energy of vaporization of substance i, and δi is the solubility parameter. Comparing Eq 2.55 to Eq 2.39 makes it evident that in the solubility parameter terms the binary interaction parameter: χ′12 = (Vr / RT)(δ1 - δ2)2 ≥ 0

(2.56)

where the reference volume is usually taken as Vr = 100 (mL/mol). It is important to note that Eq 2.55 is limited to the enthalpic part of binary interaction parameter, i.e.: χ′12 = χS + χH = χS + (Vr / RT)(δ1 - δ2)2 ≥ 0 (2.57) The entropic term in Eq 2.57, χS, originates from the local configurational as well non-configurational effects. When Eq 2.56 is used, then χS must be accounted for by other means. For molecules without polar groups the solubility parameter, δi, may be determined from: 1. the definition (see Eq 2.55): 2

v

v

1/3

δi = 14νi / Vi

2.5.2.3 Solubility Parameter Approach

v

δ2 = ∆Ei / V = (∆Hi - PVi) / Vi ⯝ (∆Hi - RT) / Vi (2.58) 2. an empirical correlation with the surface tension coefficient, νi:

(2.59)

3. solving Eq 2.55 for δ1, knowing the experimental values of ∆Hm for material 1 in a series of solvents with known value of δi. For small molecules without strong interactions the values of the solubility parameter vary from 5.9 (for C6F14) to 14.1 (for I2) and 23.8 (for H2O). The standard error of these estimates is ± 0.2. The solubility parameter of a polymer is usually determined by measuring its behavior in a solvent, whose solubility parameter is known. The polymer can be crosslinked or not — in the first case degree of swelling (Ds), in the second the intrinsic viscosity ([η]), is usually measured. From the plot of either Ds or [η] vs. δsolvent the value of δpolymer is determined as that that corresponds to the peak location [Utracki, 1972]. For copolymers, the volume additivity of the segmental solubility parameters is assumed, i.e., δ copolymer = Sδ iφ iin copolymer. This assumption also follows from the group contribution method used to compute δ from the chemical and structural characteristics of polymeric chain [Grulke, 1989; van Krevelen, 1992]. Correlation between the solubility parameter theory SPT and the EoS based on the Flory model, was explored by Biros et al. [1971]. The authors demonstrated that δ = P*1/2 / V, where P* is the pressure-reducing parameter and V ≡ V / V*. The temperature and pressure gradients of the solubility parameter were given as: [∂ ln δ/∂Τ]P ≅ −αP; and [∂ ln δ/P]T ≅ +βΤ (2.60) where αP is the isobaric thermal expansion coefficient, and βT is the isothermal compressibility. These relations can be used to correct the value of δ (usually given at 298 K and ambient pressure) to any temperature and pressure of interest. The biggest drawback of the solubility parameter approach is omission of the entropic and specific interactions’ effects. Furthermore, the fundamental dependencies do not take into account either the structural (isomeric), orientation, or the neighboring group effects. However, since the contributions that are included in the solubility parameters are

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indeed detrimental to miscibility, minimizing their values must but help the miscibility. In Table 2.14 few examples of the solubility parameters for selected polymers are given. More extensive listing can be found in the source literature [Shinoda, 1978; Van Krevelen, 1976; Grulke, 1989; Coleman et al., 1990]. As evident by data in Table 2.14, one additional problem of the solubility parameter method is poor reproducibility of the measured values. Selection of different commercial resin or using different set of solvents may significantly change the value of measured δ. In the late 1980’s a new approach to the solubility parameter concept was developed [Painter et al., 1988, 1989, 1990; Coleman et al., 1988, 1989, 1990, 1991]. The authors start by recasting Huggins-Flory Eq 2.39 into the form [Painter et al., 1988]: ∆Gm / RT = ∆GH / RT + (V / Vr) {(φ1 / Ν1) ln φ1 + (φ2 / Ν2) ln φ2 + χ′12φ1φ2}

(2.39a)

where V is the total molar volume of the two components, and Vr = 100 (mL/mol) is the previously defined reference volume. Originally, the new term, ¨GH, was introduced to express the effects of hydrogen bonding in blends in which one polymer self-associates while the other does not (but it is capable to form hydrogen bonds with the first one). Since then, this term has been used to describe all specific interactions (hydrogen bonding, ion-ion, ion-dipole, charge transfer, ›ï›-electron interactions, etc.) that provide negative contribution to the free energy of mixing. The interactions of the van der Waals type are accounted for by the χ′12φ1φ2 • 0 term, with χ′12 given by Hildebrand’s solubility parameter Eq 2.56. The degree of polymerization dependence of miscibility was assumed to follow Eq 2.40. Eq 2.39a distinguishes three contributions to the free energy of mixing: the configurational entropy (given by the two logarithmic terms), the dispersive or van der Waals interactions (given by the χ′12φ1φ2 • 0 term), and the strong interaction term, ¨GH. For systems where ¨GH = 0 the polymer-polymer miscibility can only be expected if χ′12 is very

Table 2.14. Solubility parameters at 298 K Experimental δ (J/mL)1/2

Polymer

PTFE PDMS PE PP PPG PIB PS PVAc PMMA PVC PET PMAN CA PA-66 PAN

Shinoda, 1978

van Krevelen, 1976

12.69 14.94 16.17 16.58 18.63 19.24 19.44 19.65 21.90 21.90 22.31 27.84 31.52

12.7 15.8 16.8 - 18.8 15.4 - 20.3 16.0 - 16.6 17.4 - 19.0 19.1 - 22.6 18.6 - 26.2 19.2 - 22.1 19.9 - 21.9 21.9 27.8 25.6

Calculated δ ± 0.8 (J/mL)1/2 Grulke, 1989

14.90 - 15.59 15.76 - 17.09 18.80 - 19.20 19.20 14.50 - 16.47 17.40 - 19.80 18.00 - 22.61 18.58 - 26.27 19.20 - 22.10 21.54 22.30 - 23.22 22.87 - 23.37 25.60 - 31.50

Coleman et al., 1990 16.4 15.2 17.4 14.8 19.5 19.7 18.5 20.3 23.6 24.4 28.3

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L. A. Utracki

small, e.g., for N1 = N2 = 2000, χ′12 < χcrit = 0.001, or (see Eq 2.56) δ1 ï δ2 < 0.2 (J/mL)1/2. Thus, these blends must have only very weak binary interactions. On the opposite end are blends with strong specific interactions whose critical binary interaction parameter is two orders of magnitude higher, χcrit = 0.1, or δ1 ï δ2 < 2 (J/mL)1/2. Most blends will fall in the intermediate category where: 0.001 < χcrit < 0.1. The cited (see columns 3 and 4 in Table 2.14) ranges of these parameters far exceeds the magnitude of their critical difference, δ1 ï δ2 < 2 (J/mL)1/2. The calculated values of δ i (column 5 of Table 2.14) are claimed to be precise within ± 0.8 (J/mL)1/2 [Coleman et al., 1990]. However, even this precision is insufficient for the prediction of the phase behavior of polymer blends. The heart of the Coleman-Painter theory is the derivation of the free energy of mixing term for the hydrogen bonding blends [Painter et al., 1988]: ο /Γ2ο 1

o Γ ∆GH / RT = (φ1 / r) ln(φ01 / φ1) + φ2 ln[φ21 / φ21 φ2

]

+ φ2 {[(Γ1ο / Γ2ο) - (Γ1 / Γ2)] + [X(Γ1 / Γ2) / (1 + X)]} Γ1 = (1 - K2 / KΒ) + K2 / KΒ (1 − KΒ φ21); Γ1ο = (1 - K2 / KΒ) + K2 / KΒ (1 − KΒ φο21); Γ2 = (1 - K2 / KΒ) + K2 / KΒ (1 − KΒ φ21); Γ2ο = (1 - K2 / KΒ) + K2 / KΒ (1 − KΒ φο21)2; X = KAφ01 / r; r = V1 / V2 = V1 / Vr

(2.61)

Thus, to compute the thermodynamic contribution of strong interactions to the overall thermodynamic behavior of a blend, one must first determine the three principal constants: KA, KB, and K2. These are defined as association equilibrium constants for the formation of hydrogen bond between B and A units (of polymer 2 and 1), respectively, of self-association between many B-units, and formation of doublets of the B-units. The volume fractions, φ1 and φ2 represent the total volume fraction of polymer 1 and 2, respectively, in the blend, φ01 is the volume fraction of polymer 1 that remains unassociated, φ21 and φο21 represent the not-bonded units of polymer 2 in the blend and in the neat state, respectively.

For computation of an isobaric phase diagram the temperature dependence of the association constants has to be known: KA(T), KB(T), and K2(T). An Arrhenius-type dependence has been assumed: Ki = K°i exp{-(hi / R)[(1 / T) - (1 / T°)]}; i = A, B, 2

(2.62)

where hi is the enthalpy of formation of a hydrogen bond. When the pressure influence on blend miscibility is of interest, these constants must be evaluated within the appropriate range of pressures: KA(T, P), KB(T, P), and K2(T, P). However, since this treatment does not take into account the free volume contribution, attempts of extending it to higher pressures may not be logical. One of the basic assumption of the theory is that the three equilibrium association constants can be defined in terms of the chemical repeating units and thus can be measured by spectroscopic means [Painter et al., 1989]. Infrared spectroscopy is sensitive to hydrogen bonding, but selection of the most appropriate vibration band is essential. For example, instead of N-H or O-H stretching the carbonyl group (C=O) may be a better choice. Examples of the association constants for blends of poly(4-vinyl phenol), PVPh, with acrylic polymers are given in Table 2.15. The cited values were used to compute the phase diagrams for blends of styrene-co-vinylphenol (comprising 2 to 100% VPh), with either acrylic resins or polyethers [Serman et al., 1989, 1991; Xu et al., 1991]. The experimental data confirmed the theoretical prediction, thus corroborating validity of the model. As stated before, the solubility parameter concept was developed for non-polar, low molecular weight liquids at room temperature. For polar molecules the method did not provide consistent information. To avoid trouble, initially all liquids were divided into three categories for poorly, moderately and strongly interacting systems. Another route was taken by Hansen [1967], who postulated that all the intermolecular forces: • London dispersion forces between non-polar molecules • Repulsive forces between non-polar molecules • Culombic ion/ion interactions

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Thermodynamics of Polymer Blends

Table 2.15. Association equilibrium constants for PVPh blends at 25°C [Coleman et al., 1989; Xu et al., 1991] Polymer

N

PVPh STVPh [75% VPh] STVPh [43% VPh] STVPh [25% VPh] STVPh [8% VPh] STVPh [2% VPh] PMA PEA PVAc EVAc (70% VAc) PCL hi (kcal/mol)

60 371 223 131 37 11 350 700 3,000 3,000 3,000 -

KA

KB

K2

37.1 27.5 16.5 9.7 2.8 0.8 53.2 46.8 64.0 61.6 66.2 3.8

66.8 49.6 29.8 17.5 5.0 1.4 5.2

21.0 15.6 9.4 5.5 1.6 0.5 5.6

Note to Table 2.15: STVPh [XX% VPh] stands for poly(styrene-co-vinylphenol [XX])

• Dipole/dipole interactions between the permanent dipoles • Permanent dipole/ion interactions • Induced dipole/ion interactions • Permanent dipole/induced dipole interactions • Charge-transfer forces • Hydrogen bonding • Metallic bonding • Coordination bonding, etc.

forces play in the solubility problems. The concept is illustrated in Figure 2.13.

can be combined into three types of interactions: 2

2

2

2

δi = δid + δip + δih

(2.63)

where the subscripts d, p, and h represent the dispersive, polar and hydrogen bonding interactions, respectively. Accordingly, two substances would be miscible only when their solubility parameters place them within the critical radius of a spheroid, defined as [Hansen, 1967, 1995]: 2

2

2

2

Rcrit = Θ(δ1d - δ2d) + (δ1p - δ2p) + (δ1h - δ2h) = χ12 • 0 (2.64) where the semi-empirical parameter: Θ = 4 to 5, accounts for the dominant role the dispersion

Figure 2.13. Schematic representation of the Hansen’s miscibility sphere, Eq 2.62.

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L. A. Utracki

Examples of the numerical value of the Hansen’s parameters are given in Table 2.16 [Hansen, 1967, 1994, 1995; Hansen and Beerbower, 1971; Grulke, 1989; Luciani et al., 1996b]. As reported by Hansen [1995], values of these parameters may greatly vary from one commercial resin to another, reflecting diversity of molecular weights, molecular weight distribution, presence or absence of catalyst, and a great diversity of additives — caution is advised. The solubility method has the advantage that the values of the Hansen’s partial solubility parameters, δid, δip and δih, can be calculated from the molecular structure using the additive group

contribution method. Van Krevelen [1976] demonstrated usefulness of the method using contributions from individual atoms, structural groups, and configurations. The solubility parameters are insensitive to temperature. However, since the tabulated data for the group contributions are given for amorphous materials at room temperature, T = 25°C, while miscibility at 200-300°C is often of interest, it is often necessary to correct them for the temperature effects. For this purpose either Eq 2.25 or Eq 2.60 can be used. The calculated values of δid, δip and δih at 25 and 150°C for selected polymers are listed in Table 2.17.

Table 2.16. Hansen’s solubility parameters for selected polymers at 25°C Polymer

δ d (J/mL)1/2

ABS CA CR HDPE LLDPE PA-6 PA-66 PAN PB PIB PDMS PES PET PMA PnBA PMMA PP PPS PS PSF PVAc PVC PVDF PVP SBR

18.60 18.60 19.00 18.01 17.35 17.39 18.62 21.70 16.98 14.53 16.60 19.60 19.44 15.22 16.38 18.64 17.19 18.70 21.28 19.03 20.93 18.82 9.65 21.40 17.55

δ p (J/mL)1/2 8.80 12.73 8.00 12.71 5.11 14.10 2.52 1.90 10.80 3.48 11.54 8.97 10.52 5.30 5.75 11.27 10.03 5.87 11.60 3.36

δ h (J/mL)1/2

δ (J/mL)1/2

4.20 11.01 11.14 12.28 9.1 1.02 4.66 8.00 9.20 8.59 7.63 5.77 7.51 3.70 4.30 6.96 9.66 3.07 6.66 21.60 2.7

21.00 25.08 20.62 18.01 17.35 24.25 22.87 27.43 17.02 15.47 18.52 24.20 21.54 20.57 19.55 22.69 17.19 19.79 22.47 20.26 25.66 21.54 13.11 32.54 18.07

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Thermodynamics of Polymer Blends

Table 2.17. Calculated partial solubilities at 298 and 423 K [Luciani et al., 1996b] Resin

T*(K)

V298 /V423

δ d, 298

δ p, 298

δ h, 298

δ d, 423

δ p, 423

δ h, 423

HDPE LLDPE PVAc PMA PS PEA PnBA PiBA PtBA PMMA PEHA PCP PET PA-6 PA-66 PEG PTMG PP PnBMA PtBMA PVDF PDMS*

11560 9710 9389 10360 12680 9929 8590 8590 8590 11880 8160 10752 11740 15290 11980 10170 10300 11260 10080 13340 10440 7825

0.934 0.915 0.911 0.923 0.942 0.918 0.899 0.899 0.899 0.937 0.891 0.927 0.936 0.956 0.937 0.921 0.922 0.932 0.920 0.946 0.924 0.885

18.012 17.351 15.216 15.216 16.723 15.750 16.375 16.116 14.366 13.589 16.807 15.950 15.030 17.393 17.393 16.742 17.325 16.695 15.003 13.446 9.652 16.60

0.000 0.000 11.535 11.535 8.253 10.453 8.972 8.988 8.157 9.253 7.273 10.519 13.129 12.707 12.707 10.137 7.511 0.000 7.750 7.168 5.870 1.90

0.000 0.000 7.627 7.627 5.149 6.830 5.775 4.609 8.342 10.304 3.686 6.976 10.540 11.141 11.141 8.742 6.240 0.000 8.517 9.576 6.662 8.00

15.538 14.321 12.435 12.784 14.705 13.082 13.002 12.797 11.407 11.792 13.101 13.528 13.010 15.783 15.120 13.998 14.534 14.317 12.514 11.935 8.126 12.732

0.0000 0.0000 9.4268 9.6916 7.2571 8.6824 7.1241 7.1368 6.4769 8.0297 5.6693 8.9217 11.3640 11.5310 11.0470 8.4754 6.3011 0.0000 6.4641 6.3627 4.9418 1.4573

0.0000 0.0000 6.2331 6.4081 4.5277 5.6731 4.5855 3.6597 6.6238 8.9417 2.8733 5.9167 9.1233 10.1100 9.6852 7.3090 5.2348 0.0000 7.1038 8.5002 5.6085 6.1360

Recently, Eq 2.64 was used to calculate the interphasial tension coefficient, ν12, from the chemical structure of two polymers forming an immiscible blend. ν12 = n-1

2

2 n

2

k1(ρRT) {Θ(δ1d - δ2d) + (δ1p - δ2p) + (δ1h - δ2h) } 2

2

2

= k {Θ(δ1d - δ2d) + (δ1p - δ2p) + (δ1h - δ2h) } (2.65) where k, k1 are constants, and ρ is density. Good agreement was found between the computed and experimental values of the coefficient for 46 polymer blends (see Figure 2.14). The best correlation was found for the values 0.3 ” Θ ” 0.5. Thus, contrary to the dissolution processes domi-

nated by the dispersive forces (Θ = 4-5), for the interphasial phenomena the dispersive forces are less important than the polar ones [Luciani et al., 1996, 1997]. A pragmatic modification of the solubility parameter approach was recently proposed by Bush et al. [1996]. Recognizing that ¨Hm can be negative, the authors wrote: ∆Gm / V ≈ ∆Hm / V = φ1φ2[(δ1 - δ2)2 - eh] (2.66) where eh • 0 is the energy density that originates from the specific interactions. In other words, Eq 2.66 says that to maximize miscibility the solubility parameters of the two blend components should be the same. This approach has been used by others.

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L. A. Utracki

Figure 2.14. Interfacial tension coefficient at 150°C for 46 polymer blends plotted vs. the solubility parameter contributions. R is the correlation coefficient.

2.5.3

Polymer Reference Interaction Site Model, PRISM

Thermodynamic properties of a system can also be obtained from the atomistic considerations. Molecular dynamics or Monte Carlo methods have been successfully used to study polymers. The success stems from the fact that many properties can be projected from dynamics of relatively simple, oligomeric models. Unfortunately, miscibility strongly depends on the molecular weight and so far it cannot be examined by these methods. Another, a similar route that considers interactions between individual elements of a system, is the reference interaction site model, RISM. The theory involves computations of the system structure by means of the probability density function, which describes location of all N particles of the system. The binary interactions define the pair-density function: (2)

ρN ≡ [N! / (N - 2)!] œ PNdr(N-2) = ρ2g(r)

(2.67)

PN is the probability density of N particles, r is the particle coordinate, and g(r) is the radial distribution function. The potential energy of the system is given as a product of the potential energy of

a single particle with all others, multiplied by a factor N/2: ∞

U = (Nρ / 2) œ 4πr2 u(r)g(r)dr

(2.68)

0

Thus, the total description of the thermodynamic state hinges on g(r). Ornstein and Zernike suggested separating the influence of this parameter into direct and indirect parts. The first describes direct interaction of the reference particle 1 with particle 2, while the indirect one that of particle 1 with particle 3, which in turn interacts with particle 2. g(r12) - 1 = c(r12) + ρ œ c(r13) [g(r13) - 1]dr3 (2.69) where c(r) is the direct correlation function. The advantage of the Ornstein-Zernike equation is that it can be readily generalized to more complex systems with inter- and intra-molecular interactions. RISM has been applied to progressively more complex liquids, from mono-atomic to di-atomic, CCl4, C6H6, etc. [Chandler and Andersen, 1972]. In the late 1980’s RISM has been applied to polymers [Curro and Schweizer, 1987; Schweizer and Curro, 1989]. The PRISM requires less computing time than atomistic

Thermodynamics of Polymer Blends

simulations, but still makes it possible to incorporate structural details of polymeric molecules that the lattice models have been unable to account for. For example, bond lengths, their angles, chain conformation statistics and different interaction potentials can be included in the mathematical simulations [Honeycutt, 1992; Curro, 1994]. For polymer blends PRISM provides good correlation with the experimental data obtained by SANS and light scattering. The agreement very much depends on selection of the “closure” approximation. Initial selection of the most appropriate closure for a given system can be quite difficult. Nevertheless, the method gave encouraging results, predicting details of phase separation in PO blends. The deuteration effects as well as these related to structural differences between homopolymer and copolymers were well described [Schweizer, 1993]. 2.5.4

Summary of Theories

Three theoretical approaches to the thermodynamics of polymer blends were briefly discussed: (1) the lattice theories, including the newer equation of state theories, (2) of-lattice theories, and (3) the computational methods. The lattice theories are the oldest and most frequently used to interpret and to predict the thermodynamic properties of multicomponent systems containing polymers. The Huggins-Flory lattice theory is the best known. To use the theory one must know the temperature, pressure and concentration dependence of the enthalpic and entropic contributions to the binary interaction parameter, χ12 = χΗ(T, P, φ, ...) + χS(T, P, φ, ...) / T. Two types of extension of the H-F theory were discussed, the first that evolved from Paul and Barlow heat of mixing approach, and the second developed by Coleman and Painter. The first of these two makes it possible to treat a homopolymer as a copolymer composed of sub-units of the chain, thus to generalize and predict behavior of a great number of polymer blends. The second divides all thermodynamic influences into three groups: configurational entropy, the van der Waals interactions, and the specific interactions. The

167

novelty of this modification is the method of treatment of the latter interactions that combines the FTIR-measured kinetics of associations with the H-F theory. All these approaches suffer from the fundamental drawbacks of the H-F theory: inability to take into account the fine structure of polymeric chains, non-randomness, orientation, and free volume. The interaction parameters depend on many variables and the reported numerical values vary widely. The theories based on the equation of state are more versatile. The model developed by Simha and many of his collaborators is most useful. By contrast with the H-F theory it leads to two binary interaction parameters, one energetic the other volumetric, that are constant in the full range of independent variables. Furthermore, it has been found that the numerical values of these two parameters can be approximated by the geometric and algebraic averages, respectively. The nonrandom mixing can easily be incorporated into the theory. The non-lattice approaches, viz. strong interactions, heat of mixing, and solubility parameter approaches have more limited use. Their application should be confined to cases where the assumptions used in the derivations are fulfilled, i.e., the first for blends of strongly interacting polymers, or for systems dominated by the enthalpic contributions. The influence of pressure, P, on miscibility needs a comment. Since pressure reduces the effects of the free volume contributions, for most blends the miscibility increases with P [Walsh and Zoller, 1987]. The effect depends on the magnitude of the heat of mixing. For the ¨Hm < 0 the miscibility is enhanced by compression, whereas for those with ¨Hm > 0 it is reduced [Rostami and Walsh, 1984, 1985; Walsh and Rostami, 1985]. For PS solutions, the pressure gradient of the critical solution temperature was found to be a function of the molecular weight — the higher MW the more negative the gradient value [Stroeks and Nies, 1990]. For polymer blends the gradient is usually positive: for PPE blends with a random copolymer of o- and p-fluorostyrene the gradient d(UCST)/dP = 64 to 108°C/GPa [Maeda et al.,

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L. A. Utracki

1986], for PS/PVME d(LCST)/dP = 300°C/GPa [Hiramatsu et al., 1983], blends of PEA/PVF the LCST showed a complex dependence [Suzuki et al., 1982]. Prediction of the pressure effect of solubility poses great challenge to the precision of the theoretical description. It should be noted that according to the conditions for phase separation (see the following Part) the critical point is given by third partial derivative of the free energy of mixing, and its pressure gradient by the forth partial derivative [Nies et al., 1990].

2.6

Phase Separation

2.6.1

Thermodynamics of Phase Separation

The thermodynamic conditions for phase separation are given by: phase ond phase binodal: (∂∆Gm / ∂φi ) Pfirst = (∂∆Gm / ∂φi )sec P,T , ni ,T , ni

spinodal: (∂ 2 ∆Gm / ∂φ 22 ) P,T = 0 critical po int: (∂ 3 ∆Gm / ∂φ 23 ) P,T = 0 stability condition: (∂ 4 ∆Gm / ∂φ 24 ) P,T > 0

(2.70) Schematic of a phase diagram of a binary system is shown in Figure 2.15. There are three regions of different degree of miscibility: 1. The single-phase miscible region between the two binodals, 2. The four fragmented metastable regions between binodals and spinodals, and 3. The 2-phase separated “spinodal” regions of immiscibility, bordered by the spinodals. The diagram also shows two critical solution temperatures, the lower, LCST (at higher temperature!), and the upper, UCST (at lower temperature). The phase diagram with two critical points is a rule for mixtures with low molecular weight component(s), whereas the polymer blends usually show either LCST (most) or UCST. Few blends having UCST are exemplified by: PS with SBS, PoClS, PBrS, or poly(methyl-phenyl siloxane), BR with SBR, SAN with NBR, etc. [Utracki, 1989].

Figure 2.15. The phase diagram for liquid mixtures with the upper and lower critical solution temperature, UCST and LCST, respectively.

The origin of the critical point can be traced to the temperature effect on miscibility. Patterson [1982] observed that there are three principal contributions to the binary interaction parameter, χ12: the dispersive, free volume and specific interactions. As schematically illustrated in Figure 2.16, the temperature affects them differently. Thus, for low molecular weight systems where the dispersion and free volume interactions dominate, the sum of these two has a U-shape, intersecting the critical value of the binary interaction parameter in two places — hence two critical points, UCST and LCST. By contrast, most polymer blends derive their miscibility from the presence of specific interactions, characterized by a large negative value of the interaction parameter that increases with T. The system is also affected by the free volume contribution, as well as relatively unimportant in this case dispersion forces. The sum of the interactions reaches the critical value only at one temperature — LCST. To predict the phase behavior for a given system the following steps should be taken: 1. Select the most appropriate theoretical model for the free energy of mixing, ¨Gm. 2. Determine values of the characteristic material parameters required by the selected theory. 3. Solve Eqs 2.70 for the selected theory.

Thermodynamics of Polymer Blends

169

Figure 2.16. Interactions in polymer solutions usually comprise the dispersive forces and the free volume effects, whereas the specific interactions and the free volume contribution usually control those in polymer blends.

4. Readjust the fitting parameters of the theory to optimize the fit. 5. Make predictions of the thermodynamic behavior, then verify experimentally. The most important step is the selection of the theoretical model, i.e., the form of ¨Gm. The balance between the complexity of its form and the adequacy of the description of experimental behavior must be preserved. The fundamental question is whether the pressure dependence of the phase diagram is important or not. Most data that originate from the academic laboratories deal with the ambient pressure data. Most industrial “experiments,” viz. compounding and forming, are done at high pressures: extrusion P ” 70 MPa, injection molding P ” 500 MPa. The accuracy demanded from the theory that must account for large variations of pressure is very severe. It suffice to note that according to Eq 2.70 the critical point is given by the third derivative of the free energy of mixing — its pressure dependence is given by the fourth derivative! The sometimes complicated algebraic forms of the free energy may require that Eq 2.70 is solved numerically [Jain & Simha, 1984; Nies et al., 1990]. The latter authors described the computational procedure for the modified S-S theory, Eqs 2.45-2.46, with great details. S-S theory explicitly gives the Helmholtz free

energy that easily can be converted into Gibbs free energy, viz. G = F + PV, or ¨Gm = ¨Fm + P¨Vm, where the symbols with ¨ represent the mixing functions, defined as the function of the mixture less the contributions of pure components. Results of these computations are exemplified by the three isobaric phase diagrams of PE solutions in n-hexane, shown in Figure 2.17.

Figure 2.17. Calculated phase diagrams for n-hexane PE system at P = 1, 50, and 100 bar; x2 is the mole fraction of PE having MW = 8 kg/mol [Nies et al., 1990].

170

2.6.2

L. A. Utracki

Mechanisms of Phase Separation

The phase separation takes place when a singlephase system sees a change of either composition, T or P that forces it to enter either: (1) the metastable or (2) the spinodal region (see Figure 2.15). There is a drastic difference between the phase separation mechanisms that take place for the case (1) and the case (2). When the system enters from the single-phase region into the metastable region the phase separation occurs by the mechanism resembling crystallization — slow nucleation followed by growth of the phase separated domains. Thus, this process is known as the nucleation and growth, or NG for short. By contrast, when the system is forced to jump from a single phase into the spinodal region of immiscibility the phases separate spontaneously. The process starts with instantaneous segmental density fluctuation that progress in amplitude and later with wavelength. The process is known as the spinodal decomposition, or SD for short. The two processes are schematically illustrated in Figure 2.18, as composition vs. the local linear distance. Three stages of the phase separa-

tion process are distinguished: early, intermediate and late. The upper and lower limits of the concentration are given by the tie-line limits intersecting the binodal at T = constant (see Figure 2.15). Following the late stage of phase separation, the process continues into a series of coalescence steps that start with the Ostwald ripening (d ∝ t1/3; see Eq 2.71). This is followed by the surface tension controlled coarsening (d ∝ t; see Eq 2.72), and ends with the gravitational or divergent coalescence that lead to the formation of large size separate domains. n

dVd ∝ (ν12φcVDt / RT)dt; ∴ (d / do) c = 1 + Kct (2.71) where Vd is the drop volume, ν12 is the interfacial tension coefficient, φc is the equilibrium concentration of the phase separated system (binodal concentration), V is the molar volume of the drop fluid, nc is the coarsening exponent, and Kc is the coarsening rate constant. For a steady state coarsening nc = 3; for sheared system nc = 3/2 [Ratke and Thieringer, 1985].

Figure 2.18. Schematic representation of the density fluctuations during the spinodal decomposition mechanism (SD, left) and the nucleation and growth (NG, right). Three stages are shown: early, where in SD the wavelength is constant but the amplitude increases, intermediate, where both the wavelength and the amplitude change, and late, where the concentration amplitude is at maximum and the wavelength increases due to coarsening processes; after [Hashimoto, 1986].

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Thermodynamics of Polymer Blends

KBT / ν12 ” d2 ” ν12 / g ¨p;

d∝t

(2.72)

where g is the gravitational constant, and ¨ρ is the density difference of the two phases. During the fundamental studies of phase separation, to be certain that the phase separation will proceed by the SD mechanism, the composition near the critical point is usually selected. The blend is then quenched from the miscible to the spinodal region passing through or near LCST or UCST. For this reason the SD studies are often called “critical quenching.” Conversely, to study the NG mechanism (significantly less popular subject) the researchers are selecting compositions away from the critical — these studies are then known as “off critical.” Several reviews on phase separation have been published [Kwei and Wang, 1978; Olabisi et al., 1979; Herkt-Maetzky, 1984; Aifantis, 1986, 1987; Hashimoto, 1987; Nose, 1987; Binder, 1987; Hashimoto, 1988; Han, 1988; Utracki, 1989, 1994; etc.]. 2.6.2.1 Spinodal Decomposition (SD) Theories of the dynamics of phase separation are based on a mean field approach developed for metallurgical applications [Cahn and Hilliard, 1958; Cahn, 1978]. In the spinodal region the concentration fluctuations are delocalized, leading to long range spontaneous phase separation by SD. The mean-field theory is not adequate to describe the phase dynamics of small molecule liquids (especially near the critical point), but owing to the large chain dimension and slow diffusion rates, the theory has been successful in describing the phase separation in polymeric systems. The time-dependent probability distribution function for concentration can be determined directly by digital image analysis methods [Tanaka and Nishi, 1987)]. For SD three stages and three mechanisms of domain growth have been identified [Siggia, 1979]: diffusion, liquid flow and coalescence. The earliest diffusion stage follows the Ostwald Eq 2.71 — it is limited to period when do ” d ” 5do, where do is the initial diameter of the segregated region (do = 2 to 9 nm [Voigt-Martin et al., 1986]).

The flow region dominates when 5do ” d = 1 µm. Within these two regions the SD structure is regular and the growth can be observed by the scattering methods. At the coalescence stage of SD it becomes bimodal, then irregular. Thus, the kinetics of phase separation has been studied using the time-resolved scattering techniques, with light, neutrons or other irradiation sources. There is a direct relation between the virtual structure function, S(q, t), and the scattering intensity function, I(q, t): I(q, t) = Ib + KS2 (q, t) = (KS∞ + Ib) + K(So - S∞) exp{2R(q)t} = I∞ + (Io + I∞) exp{2R(q)t} (2.73) where Ib is the background scattering, K = constant, while So and S are the values of the virtual structure function at time t = 0 and at time t. To extract the concentration fluctuation function R(q) a semi-logarithmic plot of ln(I - I∞) vs. t should be used. However, in many cases, since the scattering intensity at equilibrium is low, the concentration fluctuation function is determined as: R(q) ≅ d[ln I(q, t) / Io(q,to)] / 2dt

(2.74)

As Eq 2.73 indicates, the light scattering intensity, I(q, t) is proportional to S2(q, t). For this reason the plot of I(q, t) vs. q (at constant decomposition time and temperature) already provides evidence of the dynamics of phase separation in polymer blends. The mechanism of phase separation is analyzed from the R vs. q dependence. The dynamics of phase separation within the SD domain starts with balance between the thermodynamics and material flux. The mean field theory of phase separation leads to the following simple form of the virtual structure function, S(q) [Cahn and Hilliard, 1958]: S(q) = So exp{R(q)t} R(q) = -Mq2 [(∂2G / ∂φ2) + 2q2κ] q = (4π / λ) sin (Θ / 2) = 2π / Λ

(2.75)

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where M is Onsager mobility factor, G is free energy for the homogenous system, φ is segmental volume fraction, and κ is the gradient energy coefficient arising from localized composition fluctuation. As shown in Eq 2.75, the wavevector, q, is a function of the wavelength, λ, and scattering angle, Θ. It has been shown that near spinodal ˜2G / ˜φ2 ∝ 1- (T/Ts), where Ts is the spinodal temperature, whereas the maximum wavelength for SD phase separation is λmax ∝ [1-(T/Ts)]-1/2 [van Aarsten, 1970]. According to Eq 2.75, the concentration fluctuation function, R(q), can be linearized by plotting R(q)/q2 vs. q2. The linearity provides evidence of the SD mechanism independently of the scale of the phase separation. From the intersection at q = 0 the mutual diffusion coefficient DM is obtained:

(

)

lim R(q ) / q 2 = DM = − M (∂ 2 G / ∂φ 2 )

q→0

DM = 2φ1φ 2 ( χ12, s − χ12 )( N1 D1φ 2 + N2 D2φ1 )

[

χ12, s = ( N1φ1 ) −1 + ( N2φ 2 ) −1

] (2.76)

Evidently, DM depends on the second derivative of the free energy, which in turn can be expressed in terms of the self-diffusion coefficients of polymers 1 and 2, as well as to the

distance from spinodal. The method of determining DM is presented in Figure 2.19, whereas its dependence on T across the spinodal is shown in Figure 2.20. From DM one may calculate the binary interaction parameter, hence determine conditions for the thermodynamic miscibility. Figure 2.21 show the χ12 dependence on T and φ. Coarsening of the structure continues through a series of steps that lead to a gross phase separation. However, under certain conditions, the coarsening progresses only to a certain stage, then the structure is fixed, or “pinned”. The pinning originates in a transition from the percolation to cluster formation then freezing-in of the molecular diffusion. In simple term, the generated structure is relatively regular, thus there is little energetic incentive for the molecules to diffuse from one drop, through “unfriendly” territory of the other phase, to another drop. An example of this has been given by Takenaka et al. [1989] for blends of PI containing 20, 30 and 50 wt% of SBR. Evidently, stirring the mixture will disrupt the fine balance of forces that make the pinning process possible, thus pinning is not expected to take place during processing. The phase separation was computer simulated using finite-difference in time and space Runge-

Figure 2.19. Determination of the mutual diffusion coefficient from the early data of spinodal decomposition, SD, in SBS/PBD blends. After [Izumitani and Hashimoto, 1985].

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173

Figure 2.20. Mutual diffusion coefficient in PS/PVME system vs. T. The data were obtained from the spinodal decomposition, SD, across the spinodal temperature, Ts = 151.2°C. Quenching into SD region, and reversed quenching from the SD region were performed [Sato and Han, 1988].

Figure 2.21. Binary interaction parameters vs. temperature in PVME blends with deuterated PS. After [Herkt-Maetzky and Schelten, 1983].

Kutta and Monte Carlo with a Hamiltonian methods [Petschek and Metiu, 1983; Meakin and Reich, 1982; Meakin et al., 1983]. Both methods were found equivalent, reproducing the observed pattern of phase separation in both, NG and SD regions. The unity of the phase separation

dynamics on both sides of the spinodal has been emphasized [Leibler, 1980; Yerukhimovich, 1982]. Most work on SD focuses on the effects of temperature and composition on phase equilibria in binary polymer mixtures. However, in industrial processes other variables may be of equal importance, e.g., the shear stress and pressure. It is known that these variables are important for miscibility, hence for the morphology and performance. For example, during extrusion of PC/PBT blends the LCST was increased by at least 60°C, causing miscibility. The blend coming out from the extruder phase separated by the SD mechanism. The co-continuity of phases resulted in excellent performance. Solvent casting of polymer blends and controlled evaporation can also lead to SD. This technology has been used for industrial production of semi-permeable, selective membranes. The product characterized by co-continuity of phases also showed excellent mechanical performance. The type of solvent, concentration, temperature and method of casting are used to control the blend morphology and its final performance [Inoue et al., 1985, 1987; Nauman et al., 1986].

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2.6.2.2 Nucleation and Growth (NG) As shown in Figure 2.22, there is a significant difference in the scattering pattern evolution for NG and SD mechanisms, especially during the early stages of phase separation — SD follows the semi-logarithmic dependence (see Eq 2.75), whereas NG a linear one.

The relation shows that the rate depends on the conditions for SD, the same for systems on both sides of the spinodal. The NG is an activated process with a linear growth rate [Matsuda, 1991]: Nucleation: dN / dt = kN exp{-∆EaN / kBT} ∆EaN = (4 / 3)∆Gmπr3N,cr + 4πr3N,crν12 = (16π / 3) ν312∆G2m Growth:

Figure 2.22. A semi-logarithmic plot of the scattering intensity function vs. time for PBD blends with PS at wavevector q = 2750 nm-1; filled circles are for SD and open circles are for ND [Hashimoto, 1987].

When the concentration of the minor phase is above 10 to 15%, the SD occurs by rapid growth of the regularly spaced concentration waves, while NG is a slower and random process. On the other hand, at low concentrations of the minor phase, neglecting the fine structure of the dispersed phase, the phase separation by NG and SD mechanism looks similar. Also the reverse quenching of SD and NG morphologies is similar, both being controlled by [Kumaki and Hashimoto, 1986]: (∂2G / ∂φ2)spinodal = −lim R(q) / Mq2 = constant q→0

(2.77)

∂φ / ∂t = D∇2φ

(2.78)

where rN,cr is the critical radius of the nucleating particle and D is the diffusion constant. Accordingly, the nucleation process requires an activation energy for nucleation. However, once nucleated, the phases separate and grow spontaneously. The nucleation is initiated by the local fluctuations of density. The activation energy, ¨EaN, depends on the local potential of the free energy of mixing and the interfacial tension coefficient. Once formed, the drops grow by diffusion of macromolecules into the nucleated domains, with the rate given by the Ostwald ripening. The diffusion stage is followed by coalescence [Hashimoto et al., 1986a; Hashimoto, 1988]. In principle, during the initial stage of phase separation the NG mechanism leads to the drop/ matrix morphology in the full range of concentration. However, the morphology at the later stages depends on the volume fraction of the dispersed phase as well as the method and extend of stabilization. Furthermore, since nucleation depends on the local density fluctuation and their amplitude depends on the distance from the critical conditions, near the spinodal the phase separation can occur either by the NG or SD mechanism [Langer, 1977, 1980]. This ambiguity is exacerbated by the compounding and processing stresses. 2.6.3

Phase Diagrams

There is a large body of literature on phase diagrams for pseudo-binary and pseudo-ternary polymer systems (pseudo- is to indicate polydispersity of molecular weights) with solvents [Flory, 1953;

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175

Table 2.18. Phase equilibrium studies for polymer/solvent(s) systems Polymer

Solvent

PE

di-phenylether, n-hexane, ethylene, nitrobenzene, amyl acetate, xylene, 1,2,4,5-tetrachlorobenzene (TCB), xylene/ glycol, toluene, petroleum ether, bromonaphthalene benzene, o-dichlorobenzene, isopropyl ether, n-heptane, tetralin/butyl carbitol, TCB oligo-isobutylene, carbontetrachloride/cyclohexane/methanol cyclohexane, acetone, tert-butyl acetate, ethyl acetate, ethyl oxalate, vinyl acetate, ethyl malonate, methyl acetate, methyl succinate, octene, polyvinyl acetal/chloroform, rubber/benzene, toluene/ethanol tetrahydrofuran/water, cyclohexanone/methanol, chlorobenzene, cyclohexane/acetone water/acetone acetone, water, methyl-isobutyl ketone, dioxane/isopropanol water, water/acetone, water/propanol benzene, toluene, 3-heptanone, n-propanol, p-cymene, MEK/iso-propanol, acetone/hexane, butanone/cyclohexane benzene dioxane/methanol water, chloroform/hexane water, isopropanol/water, isooctane di-isobutyl ketone, benzene/acetone, tri-methylpentene, 2-methylheptane, toluene/methanol benzene, benzene/acetone, benzene/methanol acetone/ethanol dimethylformamide, dimethyl sulfoxide/toluene methylenechloride/methanol phenol/cyclohexane, dimethylformamide, trifluoroacetic acid/chloroform m-cresol/cyclohexane, phenol, formic acid, acetone, butyl acetate, acetone/water, ethylacetate/heptane benzene

PP PDMS PS PVC PVP PVAc PVAl PMMA PBMA PCHMA PEG PPG PIB CR PIP PAN PC PET PA-6 NC SBR

Tompa, 1956; Cantow, 1967; Utracki, 1989]. Examples of such systems are listed in Table 2.18. The phase diagrams of polymer blends, the pseudo-binary polymer/polymer systems, are much scarcer. Furthermore, owing to the recognized difficulties in determination of the equilibrium properties, the diagrams are either partial, approximate, or built using low molecular weight polymers. Examples are listed in Table 2.19. In the Table, CST stands for critical solution temperature — L indicates lower CST, U indicates upper CST (see Figure 2.15).

2.7

Methods of Measurements

The thermodynamic properties of a polymer blend determine its performance. Many authors have focused on miscibility — this aspect has been extensively discussed in the literature [Olabisi et al., 1979; Solc, 1982; Utracki, 1989; Coleman et al., 1991] and some selected results are cited in this chapter. The methods of characterization of the thermodynamic behavior of polymer blends can be divided into four groups: 1. The thermodynamic studies not directly related to miscibility, e.g., the PVT measurements.

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Table 2.19. Phase equilibrium studies for polymer/polymer systems Polymer-1

Polymer-2

PS, Mw = 237

PVME, Mw = 13.3 TMPC, Mw = ? PoClS, Mw = ? P(S-co-BrS), Mw = 22 PMPS, Mw = 2.8 PIP, Mw = 2.7 PMMA, Mw = 92 PCL, Mw = 35 NBR, Mw = 297 PMMA, Mw = 60 NBR, Mw = 340 PEMA, Mw = 332 PMMA, Mw = 126 PMMA, Mw = 42 PEG, Mw = 20-200 SBR, Mw = 480 STVPh, Mn = 11 STVPh, Mn = 14 PFSt copolymer STVPh (0-0.4% -OH) N-TPI PS, Mw = 330 SAN, Mw = 160 SMMA, Mw = 160 PMMA, Mw = 100 PMMA, Mw = ? PS, Mw = 114 PS, Mw = 230 PMMA, Mw = 130 d-PB, Mw = 53 PHEDO, Mw = ? PMA, Mw = 257 PMA, Mw = 257 PBA, Mw = 14 LLDPE, Mw = 125 PMMA, Mw = 87 PPrA, Mw = 87 PBA, Mw = 383 PPeA, Mw = 415 SMMA, Mw = 281 SMMA, Mw = 106 SMMA, Mw = 81

PS, Mw = ? PS, Mw = 29 PS, Mw = 22 PS, Mw = 10 PS, Mw = 2.7 SAN, Mw = 223 SAN, Mw = 223 SAN, Mw = 194 PVC, Mw = 55 PVC, Mw = 160 PVDF, Mw = 100 αMSAN, Mw = 160 CPE, Mw = 190 PES, Mw = ? BR, Mw = 390 PHMA, Mn = 55 PHMA, Mn = 55 PPE, Mn = ? PVME PEK PC-co-TMPC, Mw = 72 PC-co-TMPC, Mw = 72 PC-co-TMPC, Mw = 72 P-n-BMA, Mw = 270 SMAN, Mw = 101 PVME, Mw = 99 PVME, Mw = 389 PEG, Mw = 300 PI, Mw = 101 PVDC, Mw = 101 PVDF, Mw = 428 PVDF, Mw = 736 PVDF, Mw = 140 PIB, Mw = 1,000 PC, Mw = 58 PVC, Mw = 170 PVC, Mw = 170 PVC, Mw = 170 TMPC, Mw = 33 TMPC, Mw = 33 PMMA, Mw = 94

CST (°C) L = 120 L < 220 L = 350 U = 218 U = 103 U = 173 L = 150 L = 90 L = 52 L = 190 ? L = 200 L = 185 L = 100 L = 80 U = 103 L = 159 L = 84 U = 270 L = 153 - 193 U = 445 L = 175 U = 290 U = 230 U = 160 L = 135 L = 119 L = 152 L = 227 L = 85 L = 174 L = 332 L = 300 L = 233 L = 85 L = 240 L = 129 L = 126 L = 107 L = 250 L = 264 L = 150

Ref. 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 17 18 19 20 21 21 21 22 23 24 25 26 27 28 29 29 30 31 32 33 33 33 34 34 35

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Table 2.19. Continued Polymer-1

Polymer-2

GMA/MMA, Mw = 471

SAN, Mw = 152 PαMS, Mw = 55 αMSAN, Mw = 55 SAN, Mw = 178 (20%AN)

PPE, Mw = 39 PVC, Mw = 213 PVC, Mw = 213

CST (°C)

Ref.

U = 350 L = 253 U = 261 U = 150

36 37 37 37

Abbreviations for Table 2.19: PHMA - poly-n-hexyl methacrylate, STVPh - polystyrene-co-vinylphenol, PFSt - poly(o-fluorostyreneco-p-fluorostyrene), P(S-co-BrS) - poly(styrene-co-4-bromostyrene), N-TPI - “new thermoplastic polyimide”, PPrA - poly-n-propyl acrylate, PPeA - poly-n-pentyl acrylate, References for Table 2.19: 1. Voigt-Martin et al., 1986; 2. Shaw, 1974; 3. Takahashi et al., 1985; 4. Strobl et al., 1986; 5. Nojima et al., 1982; 6. Koningsveld et al., 1974; 7. McBrierty et al., 1978; 8. McMaster and Olabisi, 1975; 9. Ougizawa and Inoue, 1986; 10. Jagger et al., 1983; 11. Inoue et al., 1985; 12. Saito et al., 1987; 13. Goh et al., 1982; 14. Walsh et al., 1982; 15. Walsh and Rostami, 1985; 16. Ougizawa et al., 1985; 17. Bhagwagar et al., 1994; 18. Kambour et al., 1980; 19. Hoy, 1970; 20. Sauer et al., 1996; 21. Kim and Paul, 1992; 22. Sato et al., 1996; 23. Nishimoto et al., 1995; 24. Ougizawa et al., 1991; 25. Han et al., 1988; 26. Fernandez et al., 1986; 27. Hasegawa et al., 1992; 28. Woo et al., 1986; 29. Maiti and Nandi, 1995; 30. Penning and St John Manley, 1996; 31. Krishnamoorti et al., 1995; 32. Kyu et al., 1991; 33. Sham and Walsh, 1987; 34. Kim and Paul, 1992; 35. Andradi and Hellmann, 1993; 36. Gan and Paul, 1994b; 37. Gan et al., 1994.

2. Determination of the interaction parameter (measured by SAXS, SANS, depression of the melting point, from the Hess cycle, etc.) 3. Studies of the phase equilibria (determined by the diverse scattering methods, viz. turbidity, PICS, SAXS, SANS, by fluorescence techniques, ultrasonic measurements, etc.). 4. Indirect methods of miscibility determination (for example, the glass transition temperature from either thermal, dielectric, or mechanical tests; NMR; spectroscopic methods; microscopy; etc.).

2.7.1

PVT and Related Measurements

The equilibrium thermodynamics controls the PVT behavior of any system, its thermal expansion coefficient, compressibility, bulk modulus, hardness, etc. The thermodynamic pressure can be defined as a partial derivative of the Helmholtz free energy (see Eq 2.5). For multicomponent blends the latter function comprises two interac-

tion parameters, <ε*>k ; k=2,4. These values can as easily be determined from the dilatometric measurements as from the phase diagram [Jain et al., 1982]. It is unfortunate that with the advances of other test methods the dilatometry as been neglected. It is still being used to characterize compressibility of neat resins, but rarely to study the behavior of polymeric blends [Plochocki, 1982, 1983, 1986; Zoller, 1989; Steller and Zuchowska, 1990; Zoller and Walsh, 1995]. The interaction parameters have been calculated from the PVT data on polymer blends [Jain et al., 1982; Privalko et al., 1985]. However, due to the need for subtraction of two large numbers, the computed values were erratic, with errors up to ± 6%. Such errors may be acceptable for some applications, but not for constructing phase diagrams. More straightforward could be computation of the reducing parameters, P*, V* and T*. Next, there are two possible approaches: (1) Compare these values with the theoretical predictions assuming an EoS and an averaging scheme, e.g., Eq 2.46. This approach was

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used to interpret the thermodynamic behavior of PS/PVME blends at T = 25-200°C and P = 0-200 MPa [Ougizawa et al., 1991]. For a series of polyolefins (e.g., EPR, aPP, PP, PEP, PEB, iPB, PIB, etc.) the solubility was studied using the PVT and SANS measurements [Krishnamoortie et al., 1996]. The results were interpreted using Hildebrand’s solubility parameter formalism. For regular blends a close agreement was found. (2) From P*, V* and T* calculate the molar attractive energy, ε12, and repulsion volume, v12, then compare these with the theoretically derived averages, viz. geometric mean for the former and algebraic mean for radius of the later (see Part 2.3.1). For miscible systems the dependence of these values on composition can be easily calculated. Usually these compare reasonably well with the experimental data; immiscibility causes large variations from the monotonic dependence An interesting observation of the exothermic effects associated with phase separation of the LCST-type polymer blends was reported by Natansohn [1985]. The author observed that the magnitude of this effect correlates with the interactions between the components, but χ12 was not calculated. There is also a correlation between the thermodynamic and transport properties (see Chapter 7). Thermodynamic properties of multiphase polymeric systems affect the flow, and vice-versa the flow affects the thermodynamics. As discussed in that chapter, the effects of stress can engender significant shift of the spinodal, ¨Ts ” 16°C. While at low stresses the effects can vary, i.e., the miscibility can either increase or decrease, at higher values an enhancement of miscibility is expected. The flow has also been used to establish whether the molten blends are miscible or not [Schlund and Utracki, 1987; Utracki and Schlund, 1987]. Thermodynamics also plays a dominant role in the interphasial phenomena, viz. the interfacial tension coefficient, thickness of the interphase, ¨l, the rheological properties of the interphase, the adhesion, etc. It is worth recalling that most

lattice theories of the interphase directly relate the magnitude of ν12 and ¨l to the magnitude of the binary interaction parameter, χ12 [Helfand and Tagami, 1971, 1972; Helfand, 1975; Helfand and Sapse, 1975; Roe, 1975; Joanny and Leibler, 1978; Broseta et al., 1990]. For example, the equilibrium interfacial thickness, ¨l, in PMMA blends with PS and SAN was determined by ellipsometry [Higashida et al., 1995]. From the values of ¨l, the authors computed the temperature dependence of the binary interaction parameter, χij(T), then the phase diagrams. However, since these topics are of prime importance for Chapter 4, they will not be discussed here. 2.7.2

Determination of the Interaction Parameters

2.7.2.1 Binary Systems All three types of radiation scattering techniques: light, X-ray and neutrons, have been used to measure the interaction parameters and study the phase equilibria. Using the relations derived for polymer solutions (see Eqs 2.31-2.36), the methods can be used to measure Mw, < s2o >1/2 and A2, thus χ12. The relation between the characteristic dimension of the scattering phase, dav, the wavelength of irradiation, λ, and the scattering angle Θ, is given by (see Eq 2.55): dav = Λ = λ /2 sin (Θ/2)

(2.79)

Depending on the value of λ and Θ the experimentally accessible values of dav vary. Approximate ranges of the dimensions of scattering phases for the light, neutron, and X-ray techniques are given in Table 2.20. Great majority of polymer blends have the domain size in the range of 50 nm to 5 µm. For this reason the LS and SAXS methods have limited use for studies of their morphology, while SANS is preferred. The latter method has been used to study macromolecular sizes, conformation and morphology, in a single or multicomponent system, in molten or solid state. Since the dis-

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Table 2.20. Approximate ranges of the dimensions of scattering phases Method Light scattering, LS Small angle neutron scattering, SANS Small angle X-ray scattering, SAXS Wide angle X-ray scattering, WAXS

Discrimination based on refractive index atomic number electron density electron density

crimination is based on the atomic number, it is customary to replace the hydrogen atoms in one polymer by deuterium. The deuterated polymer is dissolved in its hydrogenated homologue at a concentration usually ” 0.1%. The mixture is then used as one of the blend’s components. Several reviews of the SANS application for the characterization of polymer conformation and morphology have been published [Sperling, 1984; Rennie, 1992; Balsara et al., 1992; Lohse, 1994; Krishnamoorti et al., 1995; Takeno et al., 1996]. It should be noted that deuteration may change conformation of macromolecules and their solubility [Utracki, 1989]. Theory predicts phase separation (with UCST) for a blend of hydrogenated with deuterated polymer [Buckingham and Hentschel, 1980; Edwards, 1983; Honeycutt, 1992]; see also Figure 2.5. The phase separation data for poly-1,4- butadiene with its deuterated analog confirmed the prediction [Bates et al., 1985, 1986; Bates and Wignall, 1986; Bates and Wilzinius, 1989; Schelten et al., 1977; Yang et al., 1983; Atkin et al., 1984]. Similarly, blends of hydrogenated with deuterated PDMS show immiscibility [Lapp et al., 1985]. Apparently, there are three isotopic effects: (i) position of the isotope; (ii) conditions for SANS measurements and (iii) difference in segmental volume upon deuteration. In conclusion, SANS is an excellent method for determination of molecular size and intermolecular interaction in polymer blends provided that the isotopic effects are either absent or appropriately corrected for. The method is precise in evaluating the effects of molecular weight, molecular structure, chemical substituents and additives as well as the independent variables,

Scattering domain size 1 to 100 µm 0.02 to 3 µm 1 to 100 nm 0.1 to 1 nm

P, T, deformation, etc. In Table 2.21 examples of more recent studies of blends’ miscibility are summarized. Interested reader may consult an earlier publication for the prior data [Utracki, 1989]. SANS has been used to determine the phase diagram in PVME/deuterated polystyrene [Schwahn et al., 1987]. In the vicinity of the critical region the non-mean field behavior was observed. The SANS results were summarized as [Utracki, 1989]: • The Mw determined by the solution methods agreed with values obtained by SANS; the average ratio Mw (sol)/Mw (SANS) was found to be 1.02. • Increasing the blends’ temperature causes the second virial coefficient, A2, to decrease, thus LCST was a rule for the studied systems. • The radius of gyration, <s20>1/2, of the deuterated specimen decreased on increase of Mw in the matrix of the same chemical character. In most cases, the coil size of the probe molecules was found to be slightly increasing with A2. • Judging by the small, positive values of A2, the blends: PMMA/SAN, PVC/PMMA, PPE/PS and PVME/PS are miscible. However, in each system the value of A2 decreased with Mw, indicating worsening miscibility. In these systems <s20>1/2 ∝ Mnw, where the exponent n depended on the system and the measurements temperature. Thus, for PS/PVME n = 0.57 or 0.52 at 25 or 120°C respectively, for PPE/PS n = 0.55; for PS in PS and PMMA in PMMA n = 0.5; while for PMMA in PSAN-19 n = 0.60 and for PSAN-19 in PMMA n = 0.64.

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Table 2.21. Recent SANS studies of polymer blends (Note: “d” before an acronym indicated deuterated polymer) Blend

Comment

Reference

dPS/PpMS or their block copolymer

for the 1:1 blend: 1000χ12 = -(0.2 ± 0.1) + (2160 ± 60)/T

dPB/PI = 3/7, 5/5, 7/3 LDPE/HDPE PIB/model PP and EB PP/model PO’s Poly(methylbutylene) (PM), poly(ethylbutylene) (PE) & block copolymers dPS/PVME, dPB/PB, dPB/PI

χ12 and phase diagram were determined Studies of PE/PE miscibility Composition and T dependence of χ12(φ, T) Isotopic, φ and T effects on χ12 Miscibility in binary and ternary systems; determination of φ and T effects on χ12, viz. for PM/PE: χ12 = 0.0028-2.30/T + 584.45/T2 Three regions of T were defined: far above Tg, near Tg, and below Tg. where the structure is frozen. Difficulties of interpretation of SANS data Interpretation of SANS data included the compressibility effects by Sanchez-Lacombe theory Comparison between SANS and PVT data MW, φ and T dependence of χ12 Binary and ternary blends of head-to-head PP with head-to-tail PP and PI — T and composition effects LDPE/HDPE at 143°C were found immiscible: χ12 = +0.00056; dHDPE with deuterated PB had χ12 = +0.0004; blends of hydrogenated PB with deuterated PB gave: χ12 = +0.00053 at 130°C.

Jung and Fischer, 1988 Hasegawa et al., 1991 Londono et al., 1994 Krishnamoortie et al., 1995 Graessley et al., 1995 Lin et al., 1996

Linear PE/branched PE dPS/PS, PIB/model PP and EB PO’s blends: PP/EPR, etc. dPS/PpMS PP/PI LDPE, HDPE, deuterated and hydrogenated PB blends

Blends of PMMA with PVDF were studied by SAXS and SANS [Wendorff, 1980, 1982]. The binary interaction parameter χ12 was plotted vs. PVDF content. The isothermal data (T = 200°C) could be expressed by the linear dependence: -1 / (χ12 + 0.035) = 0.76φPVDF - 0.72

(2.80)

Several authors reported data of < s2o >1/2 and A2 for polymer blends [Hadziioannou et al., 1983; Ree, 1987; Maconnachie et al., 1984]. The binary interaction parameter χ12 can be extracted from the parameter A2. As discussed in Part 2.4.1, < s2o >1/2 and A2 are measured using the scattering methods, including light scattering (LSC). An innovative way of studying LSC involves ternary systems: polymer-1, polymer-2 and a solvent. It requires that either one of the polymers is iso-refractive with the solvent, or that the

Takeno et al., 1996 Schipp et al., 1996 Taylor et al., 1996 Krishnamoortie et al., 1996 Londono and Wignall, 1997 Reichart et al., 1997 Alamo et al., 1997

polymers have equal contrast [Pinder, 1997]. The method was successfully used to measure χ12 for PS/PMMA blends of different MW. Similarly, A2 was determined from the osmotic pressure measurements of a ternary system: polymer-1, polymer-2 and a solvent. Values of χ12(φ) were also determined for PVCVAc with acrylic copolymers in cyclohexanone [Sato et al., 1997]. SAXS has been mainly used to study morphology of the semi-crystalline blends, affected by composition, crystallization rate, compatibilization, additives, etc. However, it can also be used to study local structures in molten polymer blends, for example within the interphasial regions. The method have been used for liquid, glassy or crystalline systems, to determine the spinodal and binodal temperatures, as well as to measure χ12. A reasonable agreement between the values measured by different methods was obtained

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Thermodynamics of Polymer Blends

[Harris et al., 1983; Riedl and Prud’homme, 1984; Barlow and Paul, 1987]. The depression of the melting point, Tm, has also been used to determine χ12 (see Table 2.22). Development of the method is credited to Nishi and Wang [1975, 1977]. Tm depends on two factors: (1) the type and size of the crystals as well as dimensions of the unit cell, and (2) the thermodynamic interactions between the crystalline polymer and other ingredient(s). To determine χ12 from Tm it is important that all specimens

(representing different compositions) are identically treated, that there is no chemical reaction between polymers, and that the mutual solubility of low molecular weight fractions and additives is either small or it is independent of the blend composition. However, it is more important to ascertain that incorporation of other ingredients changes the crystallinity only by the thermodynamic interactions, while other effects on crystallinity are negligibly small. Blending affects crystallinity in diverse ways.

Table 2.22. Determination of χ12 from the melting point depression Blend

Comment

χ12 as a function of PCL content showed a maximum at ca. 30 wt%. Small T-dependence was reported The binary interaction parameter B(cal/mL) was negative for all PVDC/acrylate blends Miscible blends with χ12 < 0. The conclusion confirmed by FTIR and SAXS Review of the procedures used to calculate χ12 PAr depressed PBT’s Tm by 17°C. Calculated χ12 varied from-0.65 to -0.22 for 20 and 80 wt% of PAr. PBT with segmented The miscibility was found to depend on polyester-polyether block copolymer the copolymer composition PA-g/poly(m-xylene adipamide) Miscibility in amorphous phase concluded from the negative values of χ12 . The miscibility increased upon transamidation during aging. PP/SEBS/Oil or PP/EPR/Oil systems Miscibility was concluded for the system: PP/SEBS/Oil, χ12 = -0.043, and immiscibility for PP/EPR/Oil PCL/poly(4-hydroxy styrene) Miscibility concluded from the negative value of χ12 / V = -0.013, and single Tg CR/PEMA Miscibility concluded from the negative value: -χ12 (φ) = -0.030-0.122 and FTIR PVDF/PMA or Temperature dependence of χ12 (φ) was determined. PVF-VDF/PMA blends Upon increase of PVF content the miscibility was replaced by immiscibility. PVDF/Poly(1,4-butylene adipate) The system was miscible in the full range of composition. Value χ12 = -0.19, was calculated PET/PETG Miscibility concluded from the negative value: -χ12 = -0.122 at 280°C PCL/PVDC with PVDC-VC, PVDC-VAc or PVDC-AN PVDF with PMMA, PEMA, PMHA copolymers PA blends with Mnor Zn-sulfonated-PS PVDF/PMMA PBT/PAr

Reference Zhang and Prud’homme, 1987 Goh and Siow, 1988 Lu and Weiss, 1991; 1992 Runt and Gallagher, 1991 Huo and Cebe, 1993

Gallagher et al., 1993 Shibayama et al., 1995

Ohlsson and Törnell, 1996

Lezcano et al., 1996 Kundu et al., 1996 Maiti and Nandi, 1996

Penning and Manley, 1996 Papadopoulou and Kalfoglu, 1997

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L. A. Utracki

This is due to the effects on nucleation and growth rates. Thus, blending method may have serious effects on crystallizability and crystal size. Experimentally, the presence of a miscible, amorphous polymer in the blend usually slows down, or it even prevents, crystallization of the semicrystalline resin. The enhancement of crystallinity and increase in Tm on blending have also been reported [Harris and Robeson, 1987; Dumoulin et al., 1987]. As a result, the Tm method is far from being fool-proof and the obtained values of χ12 should be confirmed by other techniques [Utracki, 1989; Groeninckx et al., 1998]. The enthalpic interaction parameter determined for low molecular analogs via direct measurements of ¨Hm was already mentioned. The measurements have been performed using a commercial calorimeter, e.g., NBS batch-type, Setaram C80 or Tronac 458. The data were found to correlate well with χ12 determined by other methods [Barlow and Paul, 1987; Rana et al., 1996; Ziae and Paul, 1996, 1997]. The most important aspect is the selection of suitable analogs. In this task the consideration of the electrostatic charges on the atoms of a molecule can be used as guide [Ziae and Paul, 1996, 1997]. As discussed in Part 2.5.2.2, there are several disadvantages of the method, as for example inability to account for structural and/or polydispersity effects. A “half way” method is the microcalorimetric determination of ¨Hm using low viscosity oligomeric mixtures [Singh and Walsh, 1986; Sham and Walsh, 1987]. 2.7.2.2 Ternary Systems Containing Solvent The difficulties in calorimetric determination of the interaction parameters are caused by the high viscosity of commercially interesting polymers, slow diffusion rates, heat generation during their mixing, etc. These problems do not exist for solutions. In systems: 1-solvent, 2-polymer, and 3-polymer, any method that determines either ¨Gm or its derivatives should make it possible to calculate χ23. Thus, for example, osmotic pressure measurements were used to characterize PS/PVME blends

dissolved either in toluene or ethylbenzene [Shiomi et al., 1985]. The χ23 was found to depend on the blends’ composition. Elimination of the solvent effects gave: χ23/V1 = -104(7.41 - 11.01φ3). Thus, the system was expected to remain miscible up to the PVME volume fraction φ3 = 0.67. The osmotic pressure has also been used to determine χ23 = 0.070 for PS with poly-p-chlorostyrene in toluene, 2-butanone and cumene [Ogawa et al., 1986]. For the same system χ23 = 0.087 was calculated from the intrinsic viscosity measurements. Thus, the system is thermodynamically immiscible. More recently, the osmotic pressure measurements in cyclohexanone of a ternary system resulted in of χ23(φ) for poly(vinylchloride-co-vinylacetate) blends with a series of acrylic copolymers [Sato et al., 1997]. Vapor sorption of PS/poly-α-methylstyrene gave the value: χ23 = 0.504, varying with T and polymer concentration, indicating that this system is immiscible with UCST > 100°C [Saeki et al., 1981]. Light scattering measurements (LSC) of ternary systems: polymer-1, polymer-2 and solvent was also successfully used to determine polymer-polymer interaction parameter, χ23(φ). The method is particularly easy to use if either one of the two polymers is iso-refractive with the solvent, or if the polymers have equal contrast [Pinder, 1997]. The method was successfully used to measure χ23 for PS/PMMA blends of different MW. Over the years, several authors tried to correlate polymer-polymer miscibility with solution viscosity in a common solvent [Bohdanecky and Kovar, 1982]. An interesting report in this field was published by Chee [1990]. The author considered that the parameter b = kH[η]2 (kH is the Huggins constant of Eq 2.37, and [η] is the intrinsic viscosity) is a measure of the interactions between solvent and the polymeric species. For polymer blends, the author wrote: [η]blend = w2 [η]2 + w3 [η]3 bblend = w22 b22 + w32 b33 + 2 w2 w3b23

[

]

∴ µ ≡ b23 − (b22 + b33 ) / 2 / {[η]3 − [η]2 }

2

(2.81)

Thermodynamics of Polymer Blends

The parameter µ defined in Eq 2.81 is a measure of the polymer-polymer miscibility — negative values indicate immiscibility, positive the miscibility. Three series of blends were examined: (1) PVC/PMMA, (2) PiBMA/PMMA, and (3) PiBMA/PVC. In agreement with the calculated values of the parameter µ, the first of these three blends was found miscible, whereas the two other immiscible in the full range of composition. However, the method is, at best, qualitative. For example, the effect of solvent on the parameter µ was not investigated, but fundamentals of intermolecular interactions make it dubious that non-polar and strongly polar solvents will lead to the same value of the parameter µ. The author observed that the method breaks down for polymer pairs that form molecular associations. Intrinsic viscosity measurements were also used to evaluate intermolecular interactions in blends of cellulose diacetate with polyvinylpyrrolidone [Jinghua et al., 1997]. Another method is based on the principle that the change ∆ in any thermodynamic state function depends only on the initial and final states (Hess cycle). For example, in path I two polymers are dissolved separately in the same solvent, then mixed together. In path II, polymers are first blended together, then dissolved in the same solvent. From balance of the dissolution enthalpies the heat of mixing of two polymers, ∆Hm, at the corresponding temperature can be calculated. However, since ∆Hm is determined by subtracting one large value from another, the error of these measurements is large. Furthermore, χ23 is determined in dilute systems and extrapolation to melt is impractical [Koningsveld et al., 1974]. The ∆Hm measurements have been used to characterize PPE blends with either PS, halogenated polystyrenes or copolymers [Zacharius et al., 1983]. At 34.8°C the heat of mixing of PS/PPE blend was small and negative, indicating weak specific interactions. By contrast, in blends of PS with poly-2-chlorostyrene, ∆Hm was small and positive, dependent on the molecular weight and temperature. In agreement with the observations, for PS/poly(2-chlorostyreneco-4-chlorostyrene) ∆Hm = 0.31 J/g was found

183

in the full range of copolymer composition, confirming immiscibility. Size exclusion chromatography, SEC, has been used for the determination of polymer-polymer interaction coefficient in three component systems: solvent/polymer/polymer [Narasimhan et al., 1979, 1983, 1984]. The method was found precise and thermodynamically significant. Strong solvent concentration dependence of χ23 > 0 was reported from tests of toluene/PMMA/PS system [Lau et al., 1984, 1985]. In summary, it is important to note that the determination of χ23 is of dubious value for predicting polymer/polymer miscibility at the processing conditions. The “chi parameter” is a complex function of many variables. The solution methods requires high polymer dilution, and low temperatures, significantly lower than that used for compounding or forming polymer blends. Methods capable to extrapolate the solution data to the processing conditions do not exist. The above comments are pertinent to any of the ternary solvent methods of χ23 determination. 2.7.3

Phase Diagrams

2.7.3.1 Turbidity Measurements The method consists of preparation of a series of mixtures near the phase separation condition then causing the separation to occur. The onset of turbidity is observed either visually, using a photoelectric cell, or UV-visible spectrophotometer. The ensemble of the cloud points defines the cloud-point curves (CPC) that closely follows the binodal. The method can be extended to rigorous studies of the phase separation by measuring the light scattering intensity. Many refinements of the original turbidity method have been made [Utracki, 1989]. The scattered intensity of light due to concentration fluctuation, extrapolated to zero scattering angle, is inversely proportional to the second derivative of ∆Gm. Thus, it can be used to determine the location of a spinodal, i.e., the spinodal temperature, Ts. As Eq 2.36 indicates, LS makes it possible to determine the second

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virial coefficient and from it the binary interaction parameter. However, the technique is applicable only to homogeneous systems, i.e., at temperatures T ≤ Ts for systems having LCST, or at T ≥ Ts for those having UCST. As mentioned in Part 2.7.2.2, the LSC methods has been used mainly to study the phase equilibria in polymer solutions. Pulse-induced critical scattering, PICS, is an elegant method of LS measurements that makes it possible to extend the measurements closer to the spinodal. It uses a small mass of a homogenous liquid mixture very rapidly heated or cooled into the meta-stable region. The laser light scattering intensity is measured after the thermal equilibrium is reached, but before the system can phaseseparate, then the mixture is brought out into the homogeneous region and the cycle repeated [Gordon et al., 1973]. The temperature change can be accomplished in milliseconds, the time of one cycle is less than one minute. One of the most serious obstacles in the phase equilibrium studies of polymer blends is the viscosity of the system. At temperatures limited by the degradation the self-diffusion coefficient of macromolecules is of the order of magnitude 10-4 to 10-6 m2/s [Kausch and Tirrell, 1989]. As a result the phase separation is slow. To accelerate the process a low speed centrifuge, the “centrifugal homogenizer” (CH), with PICS has been used [Koningsveld et al., 1982]. In short, centrifugation within the immiscibility zone permits determination of binodal and critical points, while the use of PICS mode allows location of the spinodal. For the polymer blends the CPC can be determined by preparing films under conditions of miscibility, in a wide range of compositions. The films are then heated through the cloud point at a rate not exceeding 0.1°C/min. Depending on the rate, type of system, and polydispersity the difference between CPC on heating and cooling can be significant. Examples of blends whose phase diagrams were determined are listed in Table 2.23.

2.7.3.2 Scattering Methods Turbidity, light scattering, and PICS methods were discussed in the preceding paragraphs. They are based on the scattering of the light waves by liquid system with optical heterogeneities. These principles have been extended to other types of radiation, e.g., X-rays and neutrons. SAXS and SANS, have been used to study polymer blend structures. By contrast with the light scattering, SAXS uses the regularity of crystalline, or pseudo-crystalline arrays of atoms, whereas SANS that of different atomic mass. The data are treated via the modified relation derived for the conventional light scattering, Eq 2.36: Kc2/R(q) = A1/M2P(q) + 2A2c2 + ...

(2.82)

where q = 4R/λsinθ is the scattering vector, c2, M2, and A2 are respectively concentration, molecular weight, and second virial coefficient of the polymer component present in the lower concentration, R(q) is the ratio between the scattered and the incident intensities. P(q) is Debye’s one-particle scatter function, λ is the neutron wavelength and θ is the scattering half-angle. The constant K contains the scattering distances of hydrogen and deuterium as the most important quantities. The average scattering radius of the polymer is calculated from P(q). The phase equilibria in systems containing diblock poly(styrene-b-butadiene), SB, mixed either with a homopolymer or a random copolymer was established by plotting the reciprocal of the intensity of the main SAXS peak as a function of 1/T — the spinodal, Ts, and binodal, TB, temperatures were measured [Zin and Roe, 1984]. SAXS has been used to study the morphology of polymer blends in the solid state [Khambatta, 1976; Russel, 1979; Russel and Stein, 1982, 1983]. For example, in the interlamellar regions of PCL/ PVC blend the system is miscible on a molecular scale. Addition of PVC impeded crystallization of PCL. At high PVC concentration PCL remained in solution. The radius of gyration was larger than that under unperturbed conditions, in spite of the fact that at the same time the second virial coefficient, A2, was virtually zero. SAXS was used

185

Thermodynamics of Polymer Blends

Table 2.23. Blends with known phase diagram Blend

Comment

Reference

PS/PVME

A rare case of miscibility of two homopolymers. MW only slightly affects the critical concentration of LCST (about 10 wt% of PS), but strongly the level of LCST = 102-145°C. PMMA is immiscible with either PS or PAN. Miscibility with SAN having AN = 5.7-38.7 wt%, at T = 140-170°C. Interfacial thickness data. LCST Miscibility chimney dependent on the blend composition and AN content in SAN. LCST

Nishi & Kwei, 1975; Nishi et al., 1975; Reich, 1986; Radusch et al., 1996

SAN/PMMA

SAN/PCL

PS/4MPC PVDF/i-PEMA PMMA/ α -MSAN PMMA/CPE PMMA/PVC PMMA/Phenoxy PES/PEO PES/Phenoxy PVC/NBR/plasticizer PVC/Acrylates PS/PMPS SBR-45/BR NBR-40/SAN Poly(styrene-comethacrylonitrile)/ acrylic copolymers

LCST ≤ 220°C LCST = 220°C LCST = 185°C LCST = 100°C LCST = 190°C Phase diagram determined using turbidity measurements. LCST = 158°C and 30 wt% Phenoxy. LCST = 80°C LCST = 194°C at 57 wt% of Phenoxy. Miscibility only for the PVC/AN part LCST = 106, 127, and 131°C at about 50 wt% of, respectively, PPrA, PBA, and PPeA. UCST = 103°C UCST = 140°C UCST = 140°C Determination of CPC. Complex phase diagrams dependent on the composition of the two types of copolymers

to study the morphology of LDPE/HDPE blends [Reckinger et al., 1984-5]. It was found that during the crystallization, macromolecules segregate. This segregation was also observed during rapid quenching at about 100°C/min. At the high rates associated with the process, the segregation distance was comparable to coil dimension in the melt. 2.7.3.3 Fluorescence Techniques The excitation fluorescence is one of the “fluorescence techniques,” frequently used for studying

McMaster, 1975; McBrierty et al., 1978; Higashita et al., 1995 McMaster, 1973; Schulze et al., 1993; Kammer & Kummerloewe, 1996; Higashita et al., 1995 Shaw, 1974 Saito et al., 1987; Hahn et al., 1987 Goh et al., 1982 Walsh et al., 1982 Jagger et al., 1983 Chiou & Paul, 1991; Etxeberria et al., 1997 Walsh & Rostami, 1985 Walsh & Singh, 1986 Inoue et al., 1985 Sham & Walsh, 1987 Takahashi et al., 1986 Ougizawa et al., 1985 Ougizawa & Inoue, 1986 Nishimoto et al., 1995

polymer blends. The method comprises three steps: incorporation of an excimer, its excitation, and recording the excitation decay. The excimer can be an aromatic polymer component of the blend (viz. PS, polyvinyl-di-benzyl, polyvinylnaphthalene, an aromatic group can be grafted into the macromolecular chain) or added as “probe” molecules (e.g. anthracene). There are three possibilities for the aromatic rings to form excimers: intra-molecular adjacent and non-adjacent as well as inter-molecular types. Each of these is sensitive to different aspects of the chain conformation and

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L. A. Utracki

environment. The most important of these for studies of polymer blends is the intermolecular, usually identified from the concentration measurements [Winnik et al., 1988]. In the nonradiative energy transfer method, NRET, the energy is transferred from donor to receptor chromophores if the distance between them is of the order of 2 to 5 nm. Phase separation is concluded from a decrease of the chromophore energy transfer. The method as been used to study PVC miscibility with PMMA or with SAN; PS or poly-α-methylstyrene, PMS, with PS-MS copolymer, PS or PMS or poly-tertbutylstyrene, PBS, and PS + PBS with PS-BS copolymer [Morawetz, 1980, 1981, 1983; Albert et al., 1986].

to studies of blends. Since the sound is most frequently generated by debonding of two phases, there should be a drastic change in the acoustic activity in blends located on the two sides of spinodal. To quantify miscibility between PVC and EVAc, acoustic emission measurements during the peel test of a PVC/EVAc/PVC sandwich were carried out [Muniz et al., 1992]. The authors considered that the acoustic emissions at slow rates of peeling are related not to the viscoelastic dissipation processes, but rather to the work necessary to pull apart polymeric chains or break bonds. The highest emission was obtained for VAc content in EVAc of 18 and 29 wt%. 2.7.4

Others Indirect Methods for Determination of Polymer/Polymer Miscibility

2.7.3.4 Ultrasonic Velocity For homogenous systems, the ultrasonic velocity is related to the ratio of modulus to density. Thus, one may expect that any method that determines density changes with adequate precision can provide indication of miscibility [Singh and Singh, 1983]. The ultrasonics can also be used as a screening method for the optimization of processing and its parameters, e.g., for the online characterization of polymers blends [Piau and Verdier, 1993; Verdier and Piau, 1995; Gendron et al., 1995]. The compressive, ultrasonic velocity at 6 MHz at room temperature was used to study cast films of PMMA/PVAc, PMMA/PS, PVC/CR, and PS/EPDM blends [Singh and Singh, 1983; Shaw and Singh, 1987]. A linear correlation between the sound velocity and composition was observed for miscible blends, whereas immiscibility, viz. in PMMA/PS blends the dependence was irregular. The phase separation in PVC/CR was detected at w = 70 wt % of CR, as a sudden departure from the linear correlation. The ultrasonic absorption vs. composition gave even stronger evidence of immiscibility. Ultrasonics have been successfully used to study the phase behavior in polyurethanes [Volkova, 1981]. Acoustic emission has been frequently used in the studies of the fracture behavior of fiber-reinforced composites. The method was also adopted

These methods do not provide data for the binodal, spinodal, or the numerical value of the interaction parameter, but a general information about the polymer/polymer miscibility. However, the information can frequently be used, e.g., to construct a map of miscibility — a simplified phase diagram. 2.7.4.1 The Glass Transition Temperature (Tg) When polymer is cooled, either from the liquid or rubbery state, its molecular motion slows down, and eventually it undergoes a transition — either crystallization or vitrification. Depending on the cooling rate or the probing method, Tg is detected at different temperature. Thus, the glass formation is a non-equilibrium phenomenon. Nevertheless, most theoretical treatments consider the glass to be at a pseudo-equilibrium state, with Tg being a second-order transition. The thermodynamics of such a state demands knowledge of “order parameters”, zi: dG = (∂G / ∂T)P,z dT + (∂G / ∂P)T,z dP + i

i

n

∑(∂G / ∂zi)P,T,z dzi

i=1

j

(2.83)

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Thermodynamics of Polymer Blends

At equilibrium, the affinity coefficients, A ≡ (∂G / ∂P)T,z → 0, and the equilibrium Eq 2.14 i regains validity. It has been argued that liquids are collections of inherently non-crystallizable macromolecules that show equilibrium amorphous properties in all T-ranges. Furthermore, in the thermodynamic description of the glassy state, it was postulated that cooling a liquid causes the configurational entropy to decrease, attaining zero at Tg [DiMarzio and Gibbs, 1958; Dong and Fried, 1997]. This concept gives legitimacy to the pseudo- or semiequilibrium theories of the glassy state, viz. Couchman’s theories [Couchman, 1978, 1979]. It is widely accepted that when the test methods are slow enough (usually ” 1°C/min and/or 1 Hz) the glass behaves in a semi-equilibrium manner, and its behavior can be generalized. In polymers the glass transition is related to the cooperative segmental motion involving 50 to 100 backbone chain carbon atoms or 15 to 30 statistical segments, i.e., a domain of a size dd = 2-3 nm [Boyer, 1966; Warfield and Hartman, 1980]. However, the glass transition is not an iso-free volume phenomenon. The most common use of Tg in determination of polymer/polymer miscibility is based on the premise that a single Tg indicates that the domain size comparable to the macromolecular radius of gyration, i.e., 2 ” dd ” 15 nm. This approach has already been discussed in Part 2.4.2. It is important to recognize that a single Tg is not a measure of miscibility but rather of the state of dispersion. There are several equations relating Tg to composition [Utracki, 1989]. Couchman [1978] proposed the following relation for the miscible systems:  ln Tg =  

∑ w ∆C i

i

Pi

  ln Tgi  /   

liquid glass − CPi ∆CPi ≡ CPi

∑ w ∆C i

i



Pi 

 (2.84)

where wi, and Tgi are respectively weight fraction and glass transition temperature of polymer i in the blend. The parameter ¨CPi is a difference

of the isobaric heat capacity in the liquid and glass, assumed to be independent of T. From this relationship several empirical and semi-empirical formulas were derived, including Gordon-Taylor, as well as Fox equation. Note that these relations are valid only for miscible systems. The latter one:

∑ w (1 − T i

g

/ Tgi ) = 0

or:

1 / Tg =

i

∑w / T i

gi

i

(2.85) is particularly simple and religiously used even to calculate ... blends’ composition from the measured values of Tg. This use should be limited to situations where the Tg vs. composition was “calibrated” and found to follow Eq 2.85. Few years back, Lu and Weiss [1991, 1992] derived the dependence also for miscible blends: Tg =

w1Tg1 + kw2 Tg 2 w1 + kw2

−

χ12 R(Tg 2 − Tg1 )bw1w2 ∆CP1 ( w1 + kw2 )( w1 + bw2 )2

(2.86) In Eq 2.86 k and b are ratios of, respectively, ¨CP’s and densities of polymers 1 and 2, and χ12 is the thermodynamic binary interaction parameter. Thus, the relation makes it possible to compute the interaction parameter of miscible blends from Tg vs. composition dependence. From Eq 2.84, the following dependence was derived for binary blends [Utracki and Jukes, 1984]: wl ln (Tg / Tgl) + kw2 ln (Tg / Tg2) = 0

(2.87)

For miscible case, the parameter k = ¨Cp1/¨Cp2. However, it was found that relaxing this condition transforms the relation into a semi-empirical one, valid for either miscible or for immiscible systems. The dependence is symmetrical — it must be valid when the indices are exchanged. Thus, the miscibility requires that k = 1/k = 1. Larger is the difference between k and 1/k, larger is the immiscibility of the system. The dependence should not be used for strongly associating polymer blends where Tg may reach values higher than these observed for either component. These

188

L. A. Utracki

miscible, hydrogen bonded or donor-acceptor pairs, are well described by a single parameter relation [Utracki, 1989]: (2.88) where K* is a material parameter — the larger is its value, the stronger is the polymer-polymer association. Several methods of Tg determination make it possible to measure the width of the glass transition (TW). The value of TW can be more reliable in assessing the degree of miscibility than Tg. For example TW = 6°C was determined for neat polymers, TW = 10°C for miscible blends and TW ≥ 32°C for blends approaching immiscibility [Fried et al., 1978]. By measuring Tg and TW for samples annealed at different temperatures than quenched, one may be able to determine the level of miscibility, hence construct a simplified phase diagram. This has been done for numerous blends listed in Table 2.23 and others, e.g., for PS/PTMPC, PVC/poly(α-methylstyreneco-methylmethacrylate-co-acrylonitrile), or NBR/ EVAc [Casper and Morbitzer, 1977]. To construct the phase diagram, thin blend specimens should be prepared. The preferred method is to cast film from a common solvent. However, it has been observed that the blend thermograms depend on the type of solvent used for casting the film. For example, PVC/PHMT blends cast from toluene had a single Tg, whereas from 1,4-dioxane two Tg’s. Clearly caution is advised — preferably the procedure should be carried out using specimens prepared by different methods. The thin specimens should be annealed at temperatures located on both sides of the expected spinodal, then quenched. The most popular method for detecting Tg is the differential scanning calorimetry (DSC). However, much more information can be obtained from a dynamic test: dielectric, shear, bending or tensile, preferably at low and constant test frequency. Tg can also be determined by dilatometry. The method extends testing of the blend miscibility to the pressures expected during processing [Jain et al., 1982; Walsh and Zoller,

1987; Zoller and Walsh, 1995]. For PPE/PS system the pressure gradient of Tg was reported as dTg/dP = 4.3-8.2°C/GPa, depending on composition [Zoller and Hoehn, 1982]. For immiscible blends seldom two Tg’s can be detected for compositions containing less than 20 wt% of the dispersed phase. The experimental range of resolution depends on the difference between Tg’s of the two polymers: ∆Tg = Tg1 -Tg2. Since the width of the glass transition can be as large as 40°C, this method of assessment of miscibility should not be used for systems with ∆Tg ≤ TW/2 = 20°C. Table 2.24 provides examples of more recent studies of blend miscibility by means of Tg measurements. Older data can be found in [Utracki, 1989]. 2.7.4.2 Spectroscopic Methods For the studies of interactions in polymer blends the nuclear magnetic resonance (NMR), and Fourier transform infrared spectroscopy (FTIR) are of principal significance. The NMR methods for the studies of polymer blends are well described in several reviews [Olabisi et al., 1979; Robeson, 1980; Martuscelli et al., 1983; Kaplan, 1984; McBrierty and Packer, 1993]. The parameters used in determination of polymer/polymer miscibility are mainly the halflife of the spin relaxation times: T1, the spinlattice, T2, the spin-spin, and T1ρ, the spin-lattice in the rotating frame. The shorter is the relaxation time the broader is the line width — for solid samples the lines are broad, for liquids and solution they are narrow. The position of the lines, i.e., the precessional frequency of the nucleus, depends on its chemical environment, spatial configuration and interactions. The position of the peak or the so-called “chemical shift” (usually quoted as d in parts-per-million, ppm) is a reflection of the energetic state of the nucleus, while the line intensity is that of its population. For example, T1ρ was used to analyze interactions between PVC and polymethacrylates: poly(cyclohexyl methacrylate) (PCHMA), poly(phenyl methacrylate) (PPMA), and poly(benzyl methacrylate) (PBzMA)

189

Thermodynamics of Polymer Blends

Table 2.24. Tg Measurements of polymer blends Blend

Comment

Reference

PVDC/aliphatic polyesters

THF cast films. Tg measured in DSC on

Aubin et al., 1983

samples annealed at 460K Oligo(styrene-co-allyl alcohol)/ Specimens mechanically mixed at T = Tg + 70°C. aliphatic polyesters DSC at 20°C/min. UCST directly observed on a hot plate PS/poly(styrene-co-4-bromo Films cast from CHCl3 of CH2Cl2. DSC scans styrene) at 20°C/min using 10 mg samples. UCST reported. Poly(aryl ether ketone) blends Linear dependence of Tg on the mole fraction of ketone linkages PMMA/SAN/SMA, Samples were prepared either by casting from PEM/SAN/SMA, MEK or melt blended. DSC as used at 20°C/min. MAN/SAN/SMA Tg determined from the onset during the second heating cycle 1,2-polybutadiene/1,4-polyisoprene Polymers co-dissolved in benzene, then freeze-dried. DSC at 10°C/min. Tg and TW measured in duplicate or triplicate SAN/SMMA/MAN Ternary blends prepared in THF, precipitated by MeOH, then dried. DSC at 20°C/min within T = 310-430K PC/TMPC/SAN/SMMA Samples by solution casting from THF in a Petri dish DSC at 20°C/min. Tg taken at onset. Phase diagrams constructed SMMA/poly(butyl-co-hexafluoro Samples prepared by dissolution in CH 2Cl2, carbonate) then precipitation by MeOH and drying. Tg taken at onset during the second heating Poly(α-MSAN)/SAN TMPC, PVC, Samples either cast from THF or hot cast from DMF or PPE or PMMA; PMMA-GMA/ SAN, acetonitrile, dried at 150°C for 2 days. or TMPC DCS at 20°C/min. Diverse phase diagrams PC/poly(ethylene Samples prepared by casting from CHCl3 terephthalate-co-caprolactone) DSC measurements at 20°C/min T ” 530K Tg taken at half-height PMMA-GMA/PVDF Samples cast from DMF into Petri dish DSC at 20°C/min to 190°C. Miscibility only for the GMA < 35.7 wt% in PMMA-GMA PS/PC or oligomeric cyclic carbonate Samples hot cast from o-dichlorobenzene into Petri dish DSC at 20°C. Samples annealed at 200°C for 5 min PS/PCHMA THF solution precipitated by MeOH. DSC at 10°C/min Tg taken as mid-point in the inflection PBT/Poly(ester carbonate) Samples either cast, precipitated from solution, or melt mixed at 250°C. DSC at 20°C/min Tg from second scan. Solution cast samples gave two Tg, whereas the other methods only one PEI/PAr Melt mixing at 300°C. DSC at 20°C/min. Tg taken at onset

Woo et al., 1984

Strobl et al., 1986 Harris and Robeson, 1987 Brannock and Paul, 1990

Roovers and Toporowski, 1992 Cowie et al., 1992c Kim and Paul, 1992

Takakuwa et al., 1994

Gan et al., 1994; Gan and Paul, 1994b Dezhu et al., 1995

Gan and Paul, 1995

Nachlis et al., 1995 Friedrich et al., 1996 Rodriguez et al., 1996

Bastida et al., 1996

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Table 2.24. Tg measurements of polymer blends Blend

Comment

Reference

CR/PEMA PMMA/PBMA

Melt mixing at 100°C. DSC at 20°C/min Samples prepared by MeOH precipitation of acetone solutions. DSC (20 mg) at 10°C/min Sample melt mixed at 400°C, quenched in ice water into 0.2-0.4 mm sheets. DSC (10 mg) at 20°C/min. Miscible blends. Melt mixing of ca. 50 mg on a hot plate at T 400-455°C, between two sheets of Kapton™ repeated 10 times. DSC at 20°C/min Samples cast from DMF. DSC at 20°C/min. Phase diagram with LCST found. Samples by co-precipitation from p-xylene, THF or DMSO. DSC at 10°C. Tg onset and inflection point.

Kundu et al., 1996 Sato et al., 1996

Samples by solution casting. DSC at 10°C. Samples melt-mixed, DSC at 20°C/min. Immiscibility window found. Samples MeOH precipitated from CH2Cl2. DSC at 20°C/min. Miscibility map given. Samples melt mixed. Contrast in thermograms enhanced by physical aging for 46 hrs at 60°C. DSC scan at 20°C/min.

Jinghua et al., 1997 Hong et al., 1997

PEEK/PEI

Poly(aryl ether ketone)/aromatic thermoplastic polyimides Poly(styrene-co-2,4-dinitrostyrene)/ PVME or PPE PVC or C-PVC with poly(caprolactam-co-caprolactone) Cellulose diacetate/PVP PMMA/PEG/Phenoxy SAN/PAr-co-TMPAr PVC/SMMA

[Sankarapandian and Kishore, 1996]. The detected domain size was between 3 and 15 nm. The miscibility was also analyzed by measuring Tg. Both methods indicated miscibility of the PVC/ PCHMA blends, and immiscibility of the PVC/ PBzMA system. However, for the PVC/PPMA = 1:1 blends while Tg indicated miscibility, the T1ρ data showed that the system is immiscible. The modern methods of taking NMR spectra involve the use of very short radio frequency pulses (of variable duration from 1 to 200 ms) instead of a continuous signal as in older NMR. This requires full automation of the test, the Fourier transform analysis, data storage and multiple scan capability. With the scalar (low power, ca. 4 kHz) and dipolar (about 45 kHz) decoupling, “magic angle” spinning and cross polarization methods one can obtain spectra of solid samples with resolution similar to those known for liquids. The spectra provide precise information on the

Goodwin and Simon, 1996

Sauer et al., 1996

Fernandez et al., 1997 van Ekenstein et al., 1997

Ahn, T. O. et al., 1997 Dompas et al., 1997

local environment of selected nuclei, configuration, interactions and sample morphology. The interactions between PVDF and PMMA, PVAc, or PVME, were studied by dipolar-dipolar decoupling, cross-polarization, magic angle spinning, high resolution 13CNMR [Lin, 1983; Ward and Lin, 1984]. Clear peak assignment for each carbon was made and peak intensities measured. It was found that blending has little effect on the peak frequency but a significant one on its intensity. For example, the attenuation (At as % of the observed to expected intensities) varied with the method of blend preparation. For PVDF/PMMA = 1:1 blend cast from DMF At = 100% (immiscibility), while cast from MEK At = 60-75%, or extruded At = 26-49% (miscibility). Information on the short distance spatial proximity of different segments of molecules can be obtained using the proton spin-diffusion NMR. This is a particularly valuable method for the

Thermodynamics of Polymer Blends

characterization of polymer blends. For example, in the case of PS/PVME cast films, the method provided information on: blend composition, fraction of interacting groups (phenyl from PS with ether from PVME) and group mobility within each of the three domains (PS, PVME and PS-PVME). Again, it was found that different degree of dispersion is obtained casting films from different solvents [Caravatti et al., 1985, 1986]. Miscibility of PEEKK/PEI = 1:1 blend was analyzed using a new solid-state NMR method [Schmidt-Rohr et al., 1990]. The method involved selection of 1H magnetization based on chemical shift difference, 1H spin diffusion during the mixing and high resolution 13C detection. The sample was prepared by compounding in a TSE at 653 K. Intimate mixing on a molecular scale was concluded. Two-dimensional deuteron NMR was used to analyze miscibility in blends of poly-1,4polyisoprene with polyvinylethylene (PIP/PVE) [Arendt et al., 1994; Chung et al., 1994]. The blends were prepared by casting 3 wt% toluene solution. The rate of reorientation as a function of temperature near Tg was determined for both components. It was found that the system is miscible, but the glass transition is broad owing to the wide distribution of segmental motions arising from the differences in the rates of the two polymers. As a result, the PIP/PVE blends were found to be rheologically complex. In spite of miscibility the time-temperature superposition was found to be invalid. High resolution 13C NMR spectroscopy was used to analyze miscibility of POM with terphenol (MW = 600 g/mol). The size of the heterogeneity in the amorphous phase was estimated as 1 nm. The 1H spin diffusion analysis indicated a homogenous mixing on the molecular level [Egawa et al., 1996].

191

fundamental aspects as well as principles of experimentation using infrared dispersive double beam spectrophotometer (IR) or computerized Fourier transform interferometers (FTIR) were well described by Klopffer [1984]. FTIR has been extensively used to study hydrogen bonding in polymer blends [Ting, 1980; Cangelosi, 1982; Moskala, 1984; Pennacchia, 1986; Coleman et al., 1991]. These interactions not only affect the -OH absorption region (3500 to 3600 cm-1), but also the =CO stretching (1737 cm-1), the -CH2- symmetric stretching (2886 cm-1) as well as the finger-printing frequency region (1300 to 650 cm-1). As discussed in Part 2.5.2.3, FTIR has been used to calculate the strong interactions term of the free energy of mixing, ∆GH (see Eqs 2.39a-2.61) [Painter et al., 1988]. FTIR was also used to analyze the mechanism of interactions in blends of aliphatic polyesters with chlorinated polymers, viz. PVC, PVDC, or PVC-DC. In miscible blends the polyesters carbonyl stretching absorption frequency (1700 to 1775 cm-1) was shifted — the shift was absent in immiscible systems. The hydrogen bonding between C=O and α- or β-hydrogen were proposed. However, there are indications that the interaction mechanism may vary from system to system. Thus, the dipolar C=O with C-Cl interactions have been also identified [Coleman et al., 1983; Prud’homme, 1982; Garton et al., 1983; Morra and Stein, 1984; Albert et al., 1986]. A short summary of the use of FTIR to studies of polymer blends is presented in Table 2.25. There are a few publications on the use of other spectroscopic techniques such as: Brillouin scattering, photoacoustic, and Raman spectroscopy. The primary application of these has been to study the heterogeneities in polymer blends, viz. crystallization or phase separation. 2.7.4.4 Microscopy

2.7.4.3 Infrared Spectroscopy The use of infrared spectroscopy for the characterization of polymer blends is extensive [Olabisi et al., 1979; Coleman and Painter, 1984; Utracki, 1989; Coleman et al., 1991]. The applicability,

The microscopic methods can be divided into several categories: optical or light microscopy (OM), scanning electron microscopy (SEM), transmission electron microscopy (TEM), atomic force microscopy (AFM), and several modifications

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Table 2.25. FTIR measurements of polymer blends Blend

Comment

Reference

PVDF/PMMA

blending slightly affected the carbonyl stretching, near 1735 cm-1

PPE/PS

FTIR 1030/700 cm-1 peak ratio was used for the determination of composition Strongest interactions for 30 wt% PPE

Saito et al., 1987 Mukherji et al., 1980

PPE/PS PPE/PS PC/PCL PC/PBT PCL/Phenoxy PS/PVME

PF (Novolac) resins with PS, SAN, PEA, PVAc, PEMA, PMMA, PMPS, PC, and PVME PET/PC PET/PA-6 PU/poly(ethylene glycolco-propylene glycol) Poly(4-vinyl phenol) with PVAc, EVAc, PCL, PPL, PMA, PEA, PBA, or P2EHA Poly(styrene-co-4-vinyl phenol) with PBMA, PHMA or poly(tetrahydrofuran) (PTHF) PVC/SAN

PPE/PS

PVPh/PMMA

FTIR was used for characterizing macromolecular orientation in solid-state miscibility in the amorphous phase miscibility in the amorphous phase H-bonding between the -OH group of Phenoxy and C=O of PCL changes in molecular environment were easiest detected using the vibrations: in PS the out-of-plane CH, and in PVME the COCH3 at 700 and 1100 cm-1 frequency shifts in CO vibration from 1774 to 1752 cm-1 were assigned to hydrogen bonding in miscible blends

Koenig and Tovar-Rodriguez, 1981 Wang and Porter, 1983 Coleman et al., 1984 Birley and Chen, 1984 Garton, 1983, 1984 Garcia, 1987

Coleman and Varnell, 1982; Fahrenholtz, 1982

FTIR used to study transesterification studies of ester-amide exchange reactions N-H stretching region (3500-3100 cm-1) was found sensitive to interactions The amount of free and bonded C=O vibrations were determined fort the PVPh/hydrogen bonding polymer blends, using a curve fitting procedure

Huang and Wang, 1986 Pillon and Utracki, 1986 Coleman et al., 1988

The hydroxyl stretching (3100-3700 cm-1) and “fingerprinting” (600-1650 cm-1) regions were analyzed. The bands 3525 and 3100-3500 cm-1 were assigned, respectively, to free and hydrogen bonded structures Nitrile stretch vibration region (2260-2220 cm-1) and C-Cl absorption (660-580 cm-1) were used to characterize the interactions New method of IR-data treatment was proposed. Weight fraction of polymer-1 in the blend: x1 = a0 + a1R / (1 + a2R), where ai are parameters, and R = A1 / (A1 + A2) is the absorbency fraction Solid-state measurements of cast films shows that the extend of hydrogen bonding depends on the solvent used for casting

Xu et al., 1991; Serman et al., 1991

of these fundamental techniques. For example, the scanning-transmission electron microscopy (STEM), and low voltage scanning electron

Coleman et al., 1989

Kim et al., 1996

Cole et al., 1996

Dong and Ozaki, 1997

microscopy (LVSEM, at 0.1 to 2 kV accelerating voltage) are particularly useful for polymer blends [Vesely and Finch, 1988]. STEM uses ultrathin

Thermodynamics of Polymer Blends

stained films, providing images with a few nanometers resolution. LVSEM provides about tenfold increased image contrast (in comparison to the conventional SEM) with almost no charging problem [Berry, 1988]. Owing to shallow sampling depth and low energy of the secondary electrons, conductive coating is not needed. The method uses flat, ultramicrotomed specimens providing image quality comparable to that of TEM. Excellent review of the microscopic methods for blends’ characterization was recently published [Vesely, 1996]. In most cases some mode of sample preparation has to be used: viz. staining, swelling, fracturing or etching. These have been extensively used to characterize morphology of immiscible blends. There are numerous sources of possible errors and for introduction of artifacts in electron microscopy, e.g., metallization in SEM and OsO4 staining in TEM may introduce an artificial grain structure observed under greater magnification. However, even at the highest resolution it is difficult to have sufficient confidence to declare whether the blend is thermodynamically miscible. One of the better source of this information comes from studies of the diffusion rate of one polymer into another using any of the above mentioned techniques. Preparation of samples for the observation under TEM is more tedious and exacting than that used for SEM. The specimens have to be hardened and stained with Br2, OsO4, or RuO4, microtomed into ca. 20 µm slices, mounted on a grid and polymeric film support and measured. Frequently, the SEM and TEM methods are being used in parallel. In all cases microscopy is considered but one method of characterization of polymer blends [Karger-Kocsis and Kiss, 1987; Kyotani and Kanetsuna, 1987; Hsu and Geil, 1987]. For example, the effects of AN content on miscibility of SAN with PMMA was studied by measuring the thickness of the interphase [Higashida et al., 1995]. The effects of concentration, compatibilization and annealing for PA with either PS or PE (compatibilized by addition of 5 wt% of PP-MA or SMA) were studied by SEM [Chen et al., 1988]. Compatibilization reduced the diameter of dispersed phase by a factor of ten

193

and stabilized the system against coalescence at the annealing temperature (T = 200 to 230°C, for at least 1.5 h). Interesting studies of phase coarsening in PMMA/SMMA blends were followed using either interference contrast light microscopy or TEM [Andradi and Hellmann, 1993]. The films cast from toluene solutions were homogenized at low temperature, then brought to the spinodal region for phase separation and coarsening. Owing to difference in the refractive indices, good contrast was obtained without etching. The kinetics of phase coarsening in blends of PS with poly(silaα-methylstyrene) was followed under an optical microscope [Maier et al., 1996]. The blends have UCST that depends on MW. Annealing within the single-phase region, than jumping to the spinodal region causes SD and phase coarsening. Similarly, optical and SEM methods were used to study the phase separation in blends of PP with isotactic poly(1-butene) [Cham et al., 1994]. The system was found to have UCST. Reactive compatibilization of PA/SAN blends was followed with careful TEM [Mujumdar et al., 1994a, b]. Better contrast was obtained using phosphotungstic acid than RuO4. The binary interaction parameter, χ12, was calculated from the micelle spacing in microphase-separated PS-bPVP system [Clarke et al., 1997]. The spacing was determined using an AFM. Details of the polymer blends morphology and the methods of its characterization are discussed in Chapter 8 of this volume.

2.8

Summary and Conclusions

The information provided in this Chapter can be divided into four parts: 1. Introduction, 2. Thermodynamic theories of polymer blends, 3. Experimental methods, and 4. The characteristic thermodynamic parameters for polymer blends. Introduction presents the basic principles of the classical, equilibrium thermodynamics, describes behavior of the single component materials, then focuses on the twocomponent systems: solutions and polymer blends.

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The main focus of the second part is on the theories and related to them experimental parameters characterizing the thermodynamic behavior of polymer blends. Several theoretical approaches are presented, starting with the classical Huggins-Flory lattice theory and evolving from it the solubility parameter and analog calorimetry approaches. Next, the highly valuable to polymer processing, the equation of state (EoS) theories were summarized. Finally, the newest descriptions of great potential based on the atomistic considerations, in particular the polymer reference interaction site model, PRISM was outlined. This approach requires less computing time than direct atomistic simulations, but still makes it possible to incorporate structural details of polymeric molecules, e.g., bond lengths, their angles, chain conformation statistics and different interaction potentials. As the reported values of the thermodynamic parameters indicate, the largest pool of data is based on the Huggins-Flory relation. This is only to be expected since the theory and related to it concept of the binary interaction parameters, either B or χ12, was introduced to polymer science more than a half century ago, in 1941 to be precise. Even their well-recognized complexity of functional dependence, viz. B = B(T, P, φ, MW, MWD, molecular structure, stresses, ...), does not discourage efforts to continue using this approach against all odds. During the last few years, the EoS approach has been used with growing frequency to describe the thermodynamic behavior of multicomponent systems containing polymers. Such problems as the temperature and pressure dependent miscibility of gases, low molecular weight liquids or polymeric liquids in selected polymers have been successfully solved. For binary blend, the EoS uses two parameters the first related to the change of specific volume and the other to the energetic interactions. It has been shown that for many cases the values of these experimental constants can be well predicted using the algebraic and geometric means, respectively. Since during the polymer processing pressure plays a major role, today the EoS is potentially the most

valuable tool available to polymer researchers and engineers.

2.9

List of Abbreviations

Most of the abbreviations used in this Chapter are listed in Appendix 1. Pertinent ones are cited under Tables 2.9 and 2.10.

2.10

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C., Vlachopoulos, J., and Walters, K., Eds., Kluver Academic Pub., Dordrecht (1995). Utracki, L. A., Simha, R., and Yahsi, U., “Interrelationships between P-V-T and Flow Behavior of Hydrocarbons,” XII International Congress on Rheology, Quebec, QC, 1996.08.18-23. Utracki, L. A., and Simha, R., Free-volume application to foaming, NRCC/IMI symposium on Polymer Foaming, Boucherville, QC, Canada, January 28, 1997. van Aarsten, J. J., Eur. Polym. J., 6, 919 (1970). van Ekenstein, G. O. R. A., Deuring, H., ten Brinke, G., and Ellis, T. S., Polymer, 38, 3025 (1997). Van Krevelen, D. W., “Properties of polymers,” 2nd edition, Elsevier, Amsterdam (1976). Van Krevelen, D. W., in “Computational modeling of polymers,” Bicerano, J., Ed., Marcel Dekker, New York (1992). Verdier, C., and Piau, M., Recent Prog. Genie Procedes 1995, 9(38), (Genie des Procedes dans la Chaine des Polymeres et dans la Chaine Catalytique), 25-30 (1995). Vesely, D., Polym. Eng. Sci., 36, 1586 (1996). Vesely, D., and Finch, D. S., Makromol. Chem., Macromol. Symp., 16, 329 (1988). Voigt-Martin, I. G., Leister K.-H., Rosenau, R., and Koningsveld, R., J. Polym. Sci.: Part B: Polym. Phys., 24, 723 (1986). Waldram, J. R., “The Theory of Thermodynamics”, Cambridge University Press, Cambridge (1985). Walsh, D. J., Higgins, J. S., and Zhikuan, Ch., Polymer, 22, 1005 (1982). Walsh, D. J., and Rostami, S., Macromolecules, 18, 216 (1985). Walsh, D. J., and Zoller, P., Makromol. Chem., 188, 2193 (1987). Wang, L. H., and Porter, R. S., J. Polym. Sci., Polym. Phys. Ed., 21, 1815 (1983). Ward, T. C., and Lin, T. S., A. C. S., Adv. Chem. Ser., 206, 59 (1984). Warfield, R. W., and Hartmann, B., Polymer, 21, 31 (1980). Wolf, B. A., Makromol. Chem., Rapid Comun., 189, 1613 (1980); Macromolecules, 17, 615 (1984). Woo, E. M., Barlow, J. W., and Paul, D. R., J. Appl. Polym. Sci., 28, 1347 (1983); ibid., 29, 3837 (1984); ibid., 32, 3889 (1986); Polymer, 26, 763 (1985). Xie, H.-K., Simha, R., and Mouliné, P., FoamTech Meeting, NRCC/IMI, Boucherville, QC, Canada, 1997.12.09. Xie, H.-K., and Simha, R., private communication (1997). Xie, H.-K., Nies, E., Stroeks, A., and Simha, R., Polym. Eng. Sci., 32, 1654 (1992).

Thermodynamics of Polymer Blends

Xu, Y., Graf, J., Painter, P. C., and Coleman, M. M., Polymer, 32, 3103 (1991). Yahsi, U., Ph.D. Thesis, Case Western Reserve University, Department of Physics, Cleveland (1994). Yamakawa, H., Modern Theory of Polymer Solutions, Harper & Row Pub., New York (1971). Zhang, S., and Prud’homme, R., J. Polym. Sci.: Part B: Polym. Phys., 24, 723 (1987). Zheng, S., Huang, J., Li, Y., and Guo, Q., J. Polym. Sci.: Part B: Polym. Phys., 35, 1383 (1997). Ziaee, S., and Paul, D. R., J. Polym. Sci.: Part B: Polym. Phys., 34, 2641 (1996); ibid., 35, 489, 831 (1997). Zoller, P., in Polymer Handbook , 3rd Edition, Brandrup, J., and Immergut, E. H., eds., J. Wiley & Sons, New York (1989). Zoller, P., and Hoehn, H. H., J. Polym. Sci., Polym. Phys. Ed., 20, 1385 (1982). Zoller, P., and Walsh, D., “Standard Pressure-VolumeTemperature Data for Polymers,” Technomic Pub. Co., Lancaster-Basel (1995).

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CHAPTER 3

CRYSTALLIZATION, MORPHOLOGICAL STRUCTURE, AND MELTING OF POLYMER BLENDS

G. Groeninckx, M. Vanneste and V. Everaert

Catholic University of Leuven, Heverlee, Belgium

3.1

Abstract

When the melt of a crystalline polymer is cooled to a temperature between the glass-transition and the equilibrium melting point, the thermodynamic requirement for crystallization is fulfilled. In a crystallizable miscible blend, however, the presence of an amorphous component can either increase or decrease the tendency to crystallize depending on the effect of the composition of the blend on its glass-transition and on the equilibrium melting point of the crystallizable component. The type of segregation of the amorphous component, influenced by parameters such as crystallization conditions, chain microstructure, molecular weight and blend composition, determine to a large extent the crystalline morphology of a crystallizable binary blend. Separate, concurrent or co-crystallization can occur in a blend of two crystallizable components. The spherulite growth of the crystallizable component in miscible blends is influenced by the type and molecular weight of the amorphous component, the former affecting the intermolecular interactions between both components and the latter the diffusion of the amorphous component. The blend composition, the crystallization conditions, the degree of miscibility and the mobility of both blend components, the nucleation activity of the amorphous component are important factors with respect to the crystallization kinetics. The melting behavior of crystallizable miscible blends often reveals multiple DSC endotherms, which can be ascribed to recrystallization, secondary crystallization, or liquid-liquid phase separation. For the discussion of the crystallization and melting behavior in immiscible polymer blends, a division into 3 main classes is proposed. In blends with a crystallizable matrix and an amorphous dispersed phase, both the nucleation behavior and the spherulite growth rate of the matrix can be affected. Nucleation of the matrix always remains heterogeneous; however, the amount of nuclei can be altered due to migration of heterogeneous nuclei during melt-mixing. Blending can also influence the spherulite growth rate of the matrix. During their growth, the spherulites have to reject, occlude or deform the dispersed droplets. In general, the major influence of blending is a change in the spherulite size and semicrystalline morphology of the matrix.

L.A. Utracki (Ed.), Polymer Blends Handbook, 203-294. © 2003 Kluwer Academic Publishers. Printed in the Netherlands.

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A completely different behavior is reported for blends in which the crystallizable phase is dispersed. Fractionated crystallization of the dispersed droplets, associated with different degrees of undercooling and types of nuclei is the rule. The most important reason is a lack of primary heterogeneous nuclei within each crystallizable droplet. An important consequence of fractionated crystallization may be a drastic reduction in the degree of crystallinity. When two crystallizable components are blended, a more complex behavior due to the influence of both phases on each other is expected. In general, the discussion for matrix crystallization and droplet crystallization can be combined. However, crystallization of one of the phases can sometimes directly induce crystallization in the second phase. As a consequence, the discussion of blends of this type has been subdivided with respect to the physical state of the second phase during crystallization. The special case of ‘coincident crystallization’, in which the two phases crystallize at the same time, is discussed. Finally, the effect of compatibilization of crystalline/crystalline polymer blends is briefly reviewed.

3.2

General Introduction

The study of the processing-morphology-property relations of polymer blends has become a topic of major scientific importance during the past 15 years mainly because of intensified technological interest in this area. The science and technology of polymer blends has now acquired an important position in the area of development of new polymeric materials. Moreover, the application of polymer blends has increased significantly and is expected to continue to grow. Of the total consumption of engineering polymers, more than 20 percent is currently thought to be composed of blends with important and various applications in the automotive, electrical and electronic industry, in computer and business equipment housings, in medical components, etc. Annually about 4900 patents related to polymer blends are published world-wide.

There are various reasons for today’s focus on polymer blends. Design of new polymers with special properties by chemical synthesis is always more expensive than the costs of the constituent existing polymers and the blending operation. A proper selection and combination of polymeric components in a certain ratio might result in a blend material with optimal properties for a specific application. The resulting blend will be the more successful; the more of the desired properties of the components are expressed in its property profile. A remarkable broad spectrum of properties can often be achieved by blending. These properties include mostly mechanical strength and stiffness, toughness, processability, heat distortion temperature, chemical and weathering resistance, flame retardancy, thermal and dimensional stability, aging resistance, elongation, permeability, transparency and gloss. A fundamental question, which has to be addressed first about any blend system of interest, is of course whether the components are miscible or not. Polymer mixtures of chemically dissimilar polymers can be divided on the basis of the miscibility of their components being miscible, partially miscible or fully immiscible. While miscibility of polymers was considered as rather rare two decades ago, it is now recognized as an achievable phenomenon with probaably well over 500 noted miscible combinations. The conceptual key toward forming miscible polymer blends is to choose polymer pairs with chemical structures capable of specific interactions leading to exothermic heats of mixing. Miscibility studies on homopolymer/copolymer blends indicate that strong repulsive interactions between the segments of the copolymer larger than those between its segments and the homopolymer might also lead to miscibility. Miscible polymer blends behave similar to what is expected of a single phase system. Their properties are a combination of the properties of the pure components and in many cases they are intermediate between those of the components. The characteristics of the components affecting the properties of miscible blends are their chemical structure and molecular weight, their concentra-

Crystallization, Morphological Structure, and Melting of Polymer Blends

tion and the intermolecular interactions, including crystallizability. While miscible blend systems are of considerable scientific and practical interest, it should not be concluded that miscibility is always the preferred situation with respect to the properties. In fact, immiscibility leading to two or multiple phases during blending is desired in various cases since the property combinations that one seeks require essentially a system in which each phase can contribute its own characteristics to the blend material. For thermodynamic reasons, i.e., small entropy gain on mixing, most arbitrary selected polymer pairs are immiscible and, as a consequence, display a two-phase behavior. Melt-mixing of immiscible polymers can result in a variety of phase morphologies depending on the blend composition, the rheological characteristics of the components such as viscosity and elasticity, the interfacial tension between the phases, and the intensity and type of flow that is applied. In the case of immiscible polymer blends, important characteristics with respect to their properties are the chemical nature of the components, the blend composition, the phase morphology (size and shape), the degree of crystallinity and semicrystalline structure of the phases in the case of crystallizable components, and the interfacial interactions between the phases. A number of miscible polymer blends are only completely miscible and form one-phase systems over a limited concentration, temperature and pressure range. Under certain conditions of temperature, pressure and composition, miscible binary blends may phase separate into two liquid phases with different compositions. Important characteristics of this type of blends are the overall blend composition, the morphology and the composition of the different phases as well as the nature of the interface between the phases. A large number of polymer blends contain one or two crystallizable components. The crystallization behavior of a polymer component in a blend is expected to be altered by the presence of the second blend component, whether both are completely miscible, partially miscible or totally immiscible. Therefore, a profound scien-

205

tific understanding of the crystallization behavior and the resulting semicrystalline structure in polymer blends is necessary for effective manipulation and control of their properties. There are a number of important factors governing the change of the crystallization rate and semicrystalline structure of a polymer in blend systems. Those include the degree of miscibility of the constituent polymers, their concentration, their glass-transition and melting temperature, the phase morphology and the interface structure in the case of immiscible blends, etc. This chapter, related to the crystallization, morphological structure and melting of polymer blends has been divided into two main parts. The first part (section 3.1) deals with the crystallization kinetics, semicrystalline morphology and melting behavior of miscible polymer blends. The crystallization, morphological structure and melting properties of immiscible polymer blends are described in the second part of this chapter (section 3.4).

3.3

Crystallization, Morphological Structure, and Melting Behavior of Miscible Polymer Blends

The crystallization of miscible and immiscible polymer blends can differ remarkably from that of the neat crystallizable component(s). In the case of crystallizable miscible blends (discussed in this section), important polymer characteristics with respect to crystallization are the chemical nature and molecular mass of the components, their concentration in the blend, and the intermolecular interactions between the components. The thermodynamic requirement for crystallization in a miscible blend is that the blend exhibits a free energy on crystallization that is more negative than the free energy of the liquidliquid mixture. A liquid-solid phase separation can occur when the miscible melt is cooled to a temperature between the glass-transition of the blend and the equilibrium melting point of the crystallizable component(s) (section 3.3.1).

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G. Groeninckx, M. Vanneste and V. Everaert

The presence of an amorphous component in a crystallizable binary blend can either increase or decrease the tendency to crystallize depending on the effect of composition on the glass-transition of the blend and on the equilibrium melting point of the crystallizable component. The morphology of a semicrystalline polymer blend is largely determined by the type of segregation of the amorphous component (section 3.3.2.1). In the case of interspherulitic segregation of the amorphous component, where the spherulites of the crystalline component are imbedded in an amorphous matrix, the semicrystalline morphology will be influenced by to a lesser extent than when the amorphous component is located within the spherulites (interlamellar and interfibrillar segregation). The parameters determining the type of segregation are not fully understood. Studies [Defieuw, 1989] indicate that the crystallization conditions, blend composition, chain rigidity and structure, and molecular weight of the components are important. Blends consisting of two crystallizable components (3.3.2.2) can exhibit separate concurrent or co-crystallization. Spherulite growth of the crystallizable component in miscible blend (3.3.3) will be influenced by the type and molecular weight of the amorphous component (the former affecting the intermolecular interactions between both

components and the latter the diffusion of the amorphous component). The blend composition, the crystallization conditions, the degree of miscibility and the mobility of both blend components, the nucleation activity of the amorphous component are important factors with respect to the crystallization kinetics (3.3.4). The melting behavior of miscible crystallizable blends (section 3.3.5) is often complex, revealing multiple DSC endotherms, which can be ascribed to several causes such as recrystallization, secondary crystallization, liquid-liquid phase separation (3.3.6), etc. 3.3.1

Crystallization Temperature Range of Crystallizable Miscible Blends

The crystallization of a polymer can only proceed in a temperature range limited on the low temperature side by the glass-transition temperature (Tg) and on the high temperature side by the equilibrium melting point (Tm°). Below Tg the mobility of the polymer chains is hindered, while in the proximity of Tm° crystal nucleation is inhibited. When dealing with crystallizable miscible blends, the glass-transition is located in between those of the neat components (Figure 3.1). The presence of an amorphous component in a crystallizable miscible polymer blend can increase or decrease

Figure 3.1. Possible crystallization temperature ranges for a crystallizable miscible polymer blend (1 = crystallizable component, 2 = amorphous component) [Runt and Martynowicz, 1986].

207

Crystallization, Morphological Structure, and Melting of Polymer Blends

the tendency to crystallize depending on the Tg of the amorphous component with respect to that of the crystallizable one. If the Tg of the amorphous component is lower than that of the crystallizable one, the crystallization envelope (Tm° - Tg) is widened, and the crystallization is facilitated. In the opposite case, where the Tg of the amorphous component is higher than that of the crystallizable one, the blend Tg is increased and the temperature range over which crystallization can occur becomes smaller. A limiting case of this is the inhibition of crystallization due to the fact that the blend Tg is higher than the Tm° of the crystallizing component, a phenomenon that is often seen in blends with a high concentration of amorphous component. An even more Table 3.1.

complex situation is observed when two miscible components are crystallizable. Some examples are given in Table 3.1. In PCL/ CPE blends the PCL crystallization is enhanced when CPE is added [Defieuw and Groeninckx 1989a]. The crystallization range becomes narrower in blends such as PCL/PECH [Runt and Martynowicz, 1986], PEG/PEMA [Cimmino et al., 1989], PCL/SAN [Defieuw and Groeninckx, 1989d], and PBT/PAr [Iruin et al., 1989]. It should be noted that the PBT/PAr 10/90 blend does not show any tendency to crystallize although the blend’s Tg is located beneath the melting point of PBT. A possible explanation for this observation is that crystallization is too slow to be noticed within the observation time limit.

Influence of Tg of the amorphous component on the crystallization ability of the crystallizable component in miscible binary polymer blends

Polymer Blend a Blend Composition (wt/wt)

Tg and Tm° (°C) of Crystallizable Component

Tg (°C) of Amorphous Component

Blend Tg (°C) Crystallization c

References

PCL/CPE42 PCL/CPE49

50/50 e 50/50 e

-63 -63

60 60

1.5 18.5

34 44

y y

[Defieuw and Groeninckx, 1989a]

PCL/PECH

10/90

-60

60

-15

ns d

y

[Runt and Martynowicz, 1986]

PCL/SAN24

50/50 e

-63

60

110

30

y

[Defieuw and Groeninckx, 1989c]

PBT/PAr

70/30 50/50 30/70 10/90

31

219

179

43 51 87 147

y y y n

[Iruin et al., 1989]

PEG/PEMA

90/10 80/20 70/30

-55

75

71

-47b -39b -30b

y y y

[Cimmino et al., 1989]

a b c d e

The crystallizable component is listed first Calculated from the Fox-equation Ability to crystallize: yes (y) or no (n) ns: not stated More compositions are reported in the article

208

3.3.2

G. Groeninckx, M. Vanneste and V. Everaert

Crystallization Phenomena in Miscible Polymer Blends

When crystallized from the melt, most polymers show a spherulitic texture (Figure 3.2). The spherulites then consist of lamellar stacks of alternating crystalline and amorphous layers, radiating from the center (the primary nucleus). 3.3.2.1 Modes of Segregation of the Amorphous Component during Crystallization in Crystalline/Amorphous Blends In blends of a crystallizable polymer with an amorphous one, the morphology is largely determined by the type of segregation of the amorphous component. Crystallization in a miscible blend involves two types of polymer transport: diffusion of the crystallizable component towards the crystallization front and simultaneous rejection of the amorphous component. This latter phenomenon is called segregation; it can take place at three different levels: interspherulitic, interfibrillar and interlamellar (Figure 3.3). Interspherulitic segregation, in which the spherulites are imbedded in an amorphous matrix, can be distinguished from the other two types using optical microscopy. In the case of intraspherulitic segregation a volume-filling texture is observed;

Figure 3.2. Schematic representation of the spherulitic texture of a semicrystalline polymer [Hoffman et al., 1976].

the amorphous components can be located either between the lamellae (interlamellar) or between stacks of lamellae (interfibrillar). To find out whether or not interlamellar segregation occurs, small angle X-ray scattering (SAXS) can be used. The increase of the long spacing, which is the sum of the average thickness of the crystalline and amorphous layers, as well as the increase of the thickness of the amorphous layers between the crystalline lamellae, with increasing concentration of the amorphous component are parameters that are often used as indications for interlamellar segregation. In Table 3.2 some examples are presented together with the parameter and/or technique used to make conclusions about the type of segregation. Most studies concerning the segregation behavior of amorphous components in a miscible crystallizable blend deal with PCL as the crystallizable component. In their blends with PCL, PVC has shown to segregate interlamellar [Stein et al., 1978, 1981, Khambatta et al., 1976, Ong and Price, 1978a, and Russell and Stein, 1980, 1983], PC interfibrillar or interspherulitic [Vandermarliere, 1986, Cruz et al., 1979 and Fernandez et al., 1986], CPE either interfibrillar or interspherulitic (depending on the amount of amorphous component) [Defieuw et al., 1989a], SMA interlamellar [Defieuw, 1989, Defieuw et al., 1989b, 1989c, and Vanneste et al., 1995], SAN interlamellar [Defieuw et al., 1989c, Vanneste et al., 1995], and Phenoxy interlamellar/interfibrillar [Defieuw et al., 1989d, Vanneste, 1993].

Figure 3.3. Schematic representation of the different types of segregation of the amorphous component in crystallizable miscible polymer blends (full lines: crystallizable component, dotted lines: amorphous component).

Crystallization, Morphological Structure, and Melting of Polymer Blends

Table 3.2.

Type of segregation of the amorphous component in some crystallizable miscible polymer blends

Polymer Blend

Amorphous Comp. (wt%)

PCL/PVC

Type of a

PCL/CPE

f

Technique b References

long spacing

SAXS, OM

Segregation interlamellar

0-30% CPE e

interfibrillar or long spacing interspherulitic interfibrillar long spacing

SAXS, OM

> 30% CPE e 0-40% SMA

interspherulitic interlamellar

long spacing long spacing

SAXS, OM SAXS, OM

amorphous layer thickness long spacing amorphous layer thickness long spacing

SAXS

[Stein et al., 1978 and 1981; Khambatta et al., 1976; Ong and Price (1), 1978; Russell and Stein, 1980 and 1983] [Vandermarliere, 1986; Cruz et al., 1979; Fernandez et al., 1986] [Defieuw and Groeninckx, 1989a; Defieuw, 1989] [Defieuw and Groeninckx, 1989a] [Defieuw, 1989; Defieuw and Groeninckx, 1989b; 1989c] [Vanneste, 1993]

SAXS, OM SAXS

[Defieuw and Groeninckx, 1989c] [Vanneste and Groeninckx, 1995]

SAXS, OM

[Defieuw, 1989; Defieuw and Groeninckx 1989d] [Vanneste, 1993]

0-10% SMA14 interlamellar PCL/SANy g

0-50% SAN24 0-20% SAN15

PCL/Phenoxy

0-50% Phenoxy interlamellar/ interfibrillar 0-10% Phenoxy interlamellar

PEG/aPMMA PEG/sPMMA PEG/iPMMA

0-40% PMMA 0-40% PMMA 0-40% PMMA

PEG/PVAc

0-40% PVAc

PEG/PEMA PEG/EVA j

Parameter

0-50% PVC

PCL/PC

PCL/SMAx

209

0-20% EVA

aPHB/iPHB

interlamellar interlamellar

interlamellar interlamellar interfibrillar or interspherulitic interlamellar and interfibrillar interfibrillar intraspherulitic interlamellar

SAXS

amorphous SAXS layer thickness long spacing long spacing long spacing

SAXS SAXS SAXS

[Silvestre et al., 1987a] [Silvestre et al., 1987a] [Silvestre et al., 1987a]

amorphous thickness

SAXS, OM

[Silvestre et al., 1987b]

OM

[Kalfoglou et al., 1988] [Cimmino et al., 1989] [Cimmino et al., 1993]

interlamellar and interfibrillarl interlamellar

aPMMA/iPHB PVDF/PMMA

ns

PVDF/PMMA

0 and 40% PMMA

PVDF/PMMA

0-75% PMMA

interlamellar, amorphous SAXS, OM interfibrillar or layer thickness interspherulitic c partially interlaSAXS, OM mellar/partially interfibrillar interlamellar amorphous SAXS thickness layer

[Abe et al., 1994] [Canetti et al., 1994] [Morra et al., 1982 and 1984; Stein et al., 1981] [Saito and Stuhn, 1994]

[Ullmann and Wendorff, 1985]

210

Table 3.2.

G. Groeninckx, M. Vanneste and V. Everaert

Continued

Polymer Blend

Amorphous Comp. (wt%)

b c d

e f g h

i j k

l

Technique b References

Parameter

Segregation

PVDF/iPEMA d

0-50% iPEMA interlamellar

long spacing

SAXS

[Eshuis et al., 1982]

iPS/PPE iPS/PS

0-30% PPE 0-30% PS

interlamellar interfibrillar

lamellar thickness amorphous layer thickness

SAXS SAXS

[Wenig et al., 1975] [Warner et al., 1977; Stein et al., 1981; Russell and Stein, 1980]

interlamellar

long spacing

SAXS

[Defoor et al., 1993]

interfibrillar or interspherulitic interlamellar interfibrillar interspherulitic

long spacing

SAXS

[Crevecoeur and Groeninckx, 1991]

melting behavior long spacing long spacing

DSC SAXS SAXS

[Chen and Porter, 1994] [Hsiao and Sauer, 1993] [Hsiao and Sauer, 1993]

LLDPE/LLDPE k

a

Type of a

PEEK/PEI

0-50% PEI

PEEK/PEI PEEK/PEI h PEKK/PEI i

0-50% PEI < 50% PEI 0-50% PEI

ns: not specified SAXS: small angle X-ray scattering, OM: optical microscopy Depending on the crystallization temperature [Stein et al., 1981] PVDF/iPEMA shows an LCST phase behavior (LCST = 183°C), the experiments were performed on blends prepared and studied below this temperature. CPE containing 42.1 or 49.1wt% chlorine x = wt% MA in SMA, i.e., 14 or 25 wt% MA y = wt% AN in SAN, i.e., 15 or 24 wt% AN Rapid crystallization conditions, interfibrillar segregation occurred in blends with a PEI concentration below 50%; at higher contents of the amorphous component interspherulitic segregation was observed. Slow crystallization conditions PEG/EVA is only miscible for a vinyl acetate content of 56% and higher [Cimmino et al., 1993] A blend of 1-octene LLDPE fractions with different short chain branching contents was investigated, i.e., 3 and 33 methyl groups per 1000 carbon atoms Two different molecular weights were used

An intensively studied blend is the PEEK/PEI blend for which interlamellar [Chen and Porter, 1994], interfibrillar [Crevecoeur and Groeninckx, 1991, and Hsiao and Sauer, 1993], and interspherulitic segregations [Crevecoeur and Groeninckx, 1991] were reported. In PEKK/PEI blends PEI is segregated interspherulitic [Hsiao and Sauer, 1993]. Warner et al. [1977] have shown that in iPS/ PS blends the non-crystallizable atactic PS was mainly segregated between the fibrils inside the spherulites. On the other hand, Wenig et al. [1975]

determined the segregation of PPE to be interlamellar region in the iPS/PPE blends. The influence of the tacticity of PMMA on segregation in PEG/PMMA blends was investigated by Silvestre et al. [1987]. Atactic and syndiotatic PMMA were found located in between the lamellae of PEG, whereas isotactic PMMA was reported to segregate interfibrillar or interspherulitic. It should, however, be noted that a low molecular weight iPMMA was used in this study. In other PEG blends the amorphous component resided in the

Crystallization, Morphological Structure, and Melting of Polymer Blends

interlamellar regions (EVAc, [Cimmino et al., 1994]), interspherulitic (PEMA, [Cimmino et al., 1989], and interlamellar and interfibrillar (PVAc, Silvestre et al., 1987b], Kalfoglou et al., 1988]). Atactic PMMA [Canetti et al., 1994] and atactic polyhydroxybutyrate (PBH, [Abe et al., 1994]) were located between the lamellae in blends with iPHB. Interlamellar segregation was also reported in blends of 1-octene LLDPE fractions with different short chain branching contents [Defoor et al., 1993]. Blends of PVDF with PMMA have been studied by several authors. All three types of segregation were detected, which was attributed to variation of the crystallization temperature by Stein et al. [1981] and Morra et al. [1982]. Hahn [1987] reported the existence of a compositional interphase (a region of varying polymer composition) between the lamellae and the amorphous interlayer. The order-disorder interphase seemed to contain pure PMMA, while in the remaining interlamellar region a homogeneous mixture of PMMA and amorphous PVDF was located. The concept of a crystal-amorphous (also orderdisorder) interface was first proposed by Flory [1962] for binary semi-crystalline/amorphous blends. The order-disorder interphase was defined as the region of loss of crystalline order. Kumar and Yoon [1991] examined this interface and found that in blends the thickness of this transition zone was essentially independent of the interaction parameter between the two polymers (when χ12 varied from -1 to -0.005). Following the theoretical predictions the thickness of this region increases only slightly when stiffer chains are considered. Due to the higher degree of order of segments of the crystallizable component in this zone, the penetration of the amorphous component is limited. The compositional interphase, however, is influenced by the stiffness of both chains and by the interaction parameter (the interfacial thickness varies with the reciprocal of |χ12|1/2) This prediction seems to be confirmed by experiments. Blends of iPS and PS as well as HDPE/LDPE blends (at a temperature above the melting point for the latter blend) have a χ12 that is nearly zero, as a consequence, they will not

211

form a mixed phase in the interlamellar region the amorphous polymer will be excluded from the interlamellar zone. This seems to be in agreement with the experimental observations for iPS/PS [Warner et al., 1977] and HDPE/LDPE [Song et al., 1988]. The presence of a pure order-disorder interphase has been observed in PVDF/PMMA blends [Wenig et al., 1975] using small angle X-ray scattering and dielectric relaxation experiments. Jonas et al. estimated the spatial extension of the order-disorder interphase of PEEK in its blends with PEI. A check of the theoretical predictions of Kumar and Yoon can be made comparing several miscible crystallizable blends with components having a similar stiffness but exhibiting variable interactions (i.e., different values of χ12). Such experimental work was done by Runt et al. [1991] who examined blends of crystallizable PEG with three different amorphous components (PMMA, PVAc and polyhydroxystyrene, PHS). The first two amorphous polymers (PMMA and PVAc) exhibited a small interaction with PEG, while PHS (being able to form hydrogen bonds with PEG), displayed large interactions. A pure PEG interphase was found for the PEG blends with PMMA and PVAc, whereas a relaxation suggestive for the presence of a mixed interphase for the PEG/PHS blend was observed. Barron et al. [1992] studied strongly interacting PCL blends with PC, PVC, Phenoxy, etc. by means of dielectric relaxation measurements. The blends exhibit a dielectric relaxation in between the relaxation of the pure components, indicating the presence of a mixed amorphous interphase. The possibility to observe this transitional behavior depended on the frequency used; a frequency of 10Hz was used in this case. Therefore, it was impossible to study these transitions by means of dynamic mechanical experiments (DMA, usually 1Hz is used) or differential scanning calorimetry (DSC). For the PCL/PVC blends (PVC segregates interlamellar [Russell and Stein, 1983]) three transitions were noticed: 1) a pure amorphous PCL region (γ-relaxation), 2) a mixed amorphous phase, located at a higher temperature than the former and which shifts to higher temperatures

212

G. Groeninckx, M. Vanneste and V. Everaert

with increasing PVC content, and 3) an interphase transition that shifts to higher temperatures the stronger the interaction between both polymers. A frequency lower then 10Hz results in an overlap of the transitions associated with the mixed phase and with the interphase, while at higher frequencies the γ-relaxation merges with the interphase transition. 3.3.2.2 Modes of Crystallization in Crystalline/ Crystalline Blends When dealing with miscible blends containing two crystalline components, several modes of crystallization are possible: separate crystallization, concurrent crystallization, co-crystallization, etc. Only those blends in which both components are miscible in the melt are considered here (Table 3.3). PET/PBT blends were reported to be an example of separate crystallization [Escala and Stein, 1979; Stein et al., 1981]. A spherulitic crystallization was observed for the neat components as well as for blends with small amounts of one component, and the crystals of the minor component were included within the spherulites of the major component, which results in a coarsening of the spherulitic texture. Transesterification is, however, the reason for the homogenous amorphous phase. A simultaneous (or concurrent) crystallization can only occur when the crystallization temperature ranges overlap and if the crystallizability of both blend components is similar. Cocrystallization is only possible when the components are isomorphic or miscible in the amorphous as well as in the crystalline phase. In both cases mixed crystals can result, but in the case of concurrent crystallization no changes in crystal structure may be induced. Cocrystallization requires chemical compatibility, close matching of the chain conformations, lattice symmetry and comparable lattice dimensions [Olabisi et al., 1979]. Some examples of miscible polymer blends with two crystalline components are given in Table 3.3 together with the type of crystallization.

3.3.3

Spherulite Growth of the Crystallizable Component

3.3.3.1 Spherulite Growth Rate in Homopolymers In the case of homopolymers, the growth rate of a lamellar crystal is controlled by two processes: on the one hand by the ability of forming a surface nucleus (as determined by the degree of undercooling, ¨T = Tm° - Tg), and on the other hand by the ability of diffusion of the chain molecules towards the crystal growth front (determined by the difference between the crystallization temperature, Tc and the glass-transition temperature, Tg). Both processes are inversely dependent on temperature, a maximum rate of crystal growth is usually observed at temperatures close to Tmax § (Tg + Tm)/2. The growth rate kinetics of bulk semicrystalline homopolymers have been described in the past by Mandelkern et al. [1954] and Hoffman and Lauritzen [1976, 1973], using a modified version of the theory of nucleation of Turnbull and Fisher [1949]: G = G° exp[-¨E/R(Tc - To)] exp[-¨F*/kBT] (3.1) G° is a constant dependent on the regime of crystallization, independent of temperature, and inversely proportional to the polymer molecular weight [Van Antwerpen and Van Krevelen, 1972]; To is the temperature at which motions necessary for the transport of molecules through the liquidsolid boundary cease; Tc is the temperature of crystallization; and kB is the Boltzmann constant. The rate of growth of a crystal, G, is governed by two processes: the activation energy required to transport crystalline molecules across the solidliquid interface (¨E) and the work necessary to form a critical nucleus (¨F*). At low supercooling the growth rate is nucleation controlled, while at high supercooling it is diffusion controlled; as a consequence, Eq 3.1 produces a bell-shaped curve. Such a behavior for iPS is shown in Figure 3.4.

Crystallization, Morphological Structure, and Melting of Polymer Blends

Table 3.3.

213

Crystallization types of miscible polymer blends consisting of two crystallizable components

Polymer Blend

Crystallization Type a

References

LDPE/LLDPE

concurrent crystallization separate crystallization

[Hu et al., 1977] [Kyu et al., 1987]

UHMWPE/HDPE UHMWPE/LLDPE UHMWPE/LDPE LLDPE/LLDPE b LDPE/EPDM c HDPE/LLDPE i LLDPE/VLDPE LDPE/VLDPE HDPE/VLDPE DHDPE e/LLDPE f DHDPE/LLDPE g DHDPE/HDPE

concurrent crystallization concurrent crystallization cocrystallization/ separate k separate crystallization cocrystallization cocrystallization cocrystallization j cocrystallization partial cocrystallization cocrystallization cocrystallization m partial cocrystallization partial cocrystallization n

[Kyu and Vadhar, 1986] [Kyu and Vadhar, 1986] [Vadhar and Kyu, 1987] [Kyu and Vadhar, 1986 1987] [Rego Lopez and Gedde, 1988] [Starkweather, Jr., 1980] [Hu et al., 1987; Edward, 1986; Gupta et al., 1994] [Huang et al., 1990] [Huang et al., 1990] [Huang et al., 1990] [Tashiro et al., 1992; 1994] [Tashiro et al., 1994] [Tashiro et al., 1994]

PEEK/PEK PEEK/PEK PEEK/PEEEK PEEK/PEEKK PEEK/P(E)0.43(K)0.57 o PEEK/PEK/PEI

cocryst. / separate cryst. q cocrystallization cocrystallization cocrystallization cocrystallization cocryst. of PEEK and PEK

[Sham et al., 1988] [Harris and Robeson, 1987] [Harris and Robeson, 1987] [Harris and Robeson, 1987] [Harris and Robeson, 1987] [Harris and Robeson, 1988]

PVF/PVDF VDF-TFE/VDF-HFA iPS/iP(p-Me-S) p P(iPr-vinylether)/P(sec-But-vinyl ether) P(4-Me-pentene)/P(4-Me-hexene)

cocrystallization cocrystallization l cocrystallization cocrystallization cocrystallization

[Natta et al., 1965] [Cho et al., 1993] [Natta et al., 1961] [Wunderlich, 1973] [Wunderlich, 1973]

PBT/PEE h PET/PBT

cocrystallization separate crystallization

[Gallagher et al., 1993] [Stein et al., 1978]

iPMMA/sPMMA

cocrystallization

[Liquori et al., 1965]

PCL/PC

separate crystallization

[Vandermarliere, 1986]

PPE/iPS

separate crystallization

[Hammel et al., 1975]

PED/EVAc d

cocrystallization

[Clough et al., 1994]

214

Table 3.3. a

b c d

e f g h

i j k

l

m

n o p q

G. Groeninckx, M. Vanneste and V. Everaert

Continued.

It should be noted that not all authors use the same terminology concerning the type of crystallization. Especially the terms “cocrystallization” and “concurrent crystallization” are often confused. Since some authors did not examine whether the lattice parameters change are not, it is not possible to decide if they mean cocrystallization or concurrent crystallization. Different molecular weight fractions Ethylene/propylene/1,4-hexadiene with an ethylene/propylene ratio of 4.5 mol% EVAc has a molar ratio of ethylene to vinyl acetate of 7:1 and is amorphous, an increase of the lattice parameters was noticed when adding EVAc DHDPE: deuterated HDPE LLDPE with a branching content of ca. 17 ethyl groups/1000 carbons LLDPE with a branching content of ca. 41 ethyl groups/1000 carbons Miscibility of PBT/PEE depends on the copolymer composition of PEE, cocrystallization occurs under all crystallization conditions and is possible because the unit cell parameters of PBT and PEE are the same. To avoid interchain chemical reactions the blends were prepared by solvent casting. LLDPE: ethylene butene-1 copolymer, 18 branches /1000 C Valid as well for slowly as rapidly (quenched) cooled blends Depending on the blend preparation: cocrystallization when sequentially mixed and separate crystallization when simultaneously mixed The type of crystallization depends on the thermal treatment of the samples: cocrystallization takes place in samples that are quenched or annealed at 110°C for 6h; separate crystallization is observed when annealed at 100°C for 6h. This is due to the existence of an UCST phase behavior between 100° and 110°C. The lattice parameters vary continuously with composition of the blend and the cocrystallization process is ascribed to the closeness of the crystallization rate of both species The tendency to cocrystallize increases with increasing HDPE concentration P(E)0.43(K)0.57 is a random copolymer composed of phenyl ether and phenyl ketone units Copolymer of styrene and p-methyl styrene containing 30 mol% of the latter comonomer Cocrystallization occurs when the blends are quenched rapidly from the melt (= 100°C/min); separate crystallites are formed when isothermally crystallized, annealed at high temperatures, precipitated from solution or slowly cooled from the melt (1°C/min).

3.3.3.2 Spherulite Growth Rate in Miscible Polymer Blends When dealing with crystallizable miscible polymer blends containing a non-crystallizable component, some refinements had to be made. Some modifications were proposed by Alfonso and Russell [1986] and by Cimmino et al. [1989] for blends in which the amorphous component is segregated into the interlamellar region (see also section 3.1.2.1). First, the chemical potential of the liquid phase might be altered by the specific interactions, that are often responsible for the miscibility of polymers [Olabisi et al., 1979]. Such interactions may change the free energy required to form a critical nucleus

as well as the mobility of both the crystalline and amorphous components. Second, the non-crystallizable component has to diffuse away from the crystal growth front into the interlamellar region. Thus the rate at which the growth front progresses depends on the competition between the inherent capability of the crystal to grow and on the rate of rejection (segregation) of the amorphous component. The kinetics of crystal growth will ultimately be determined by the slower of these two phenomena. A direct consequence of this consideration is the dependence of the crystal growth rate on the molecular weight of both components. Third, the concentration of the crystallizable component at the growth front will decrease crystallization.

Crystallization, Morphological Structure, and Melting of Polymer Blends

215

φ2 is the volume fraction of the crystallizable component; Tc is the crystallization temperature; k1 is the rate of transport of the crystallizable molecules across the liquid-solid boundary: k1 = G° exp[-¨E/R(Tc - To’)]

(3.3)

To’ is the value of To in the blend and can be written in terms of the glass-transition and a constant C (associated with the WLF constant C2) [Rostami, 1990]; and k2 is the rate at which the amorphous component segregates: k2 = D/d = 2 E/L

Figure 3.4. Spherulitic growth rate in iPS and iPS/PS blends (the values represent the percentage of atactic PS present in the blend) [Keith and Padden, 1964].

And finally, the glass-transition temperature and the melting temperature can be influenced by addition of an amorphous polymer. As already mentioned in section 3.3.1, the Tg of miscible blend lies in between the glass-transition temperatures of the neat components, its value being a function of the blend’s composition. Depending on the Tg value of the non-crystallizable component (higher or lower than the Tg of the crystallizable component) the crystallization temperature range will be respectively, narrowed or widened. Incorporating the concepts discussed above, the equation describing the crystal growth rate in a miscible polymer blend can be expressed as: Gm = [(φ2 k1 k2)/(k1 + k2)] exp(-¨Fm*/kBTc)

(3.2)

(3.4)

d is the maximum distance over which the amorphous component has to diffuse away during crystallization (d = L/2, with L the crystal lamellae thickness) and D is the diffusion coefficient. Since a simultaneous diffusion of the amorphous and the crystalline component takes place, the diffusion coefficient of interest is the mutual diffusion coefficient, E; ¨Fm* is the free energy of nucleus formation (secondary nucleation) in the presence of a non-crystallizable component. The rate of crystal growth in a semicrystalline blend, Gm, will depend on the magnitude of k1, k2 and ¨Fm*. At low undercooling, ¨T = Tm° - Tc, ¨Fm* is high and hence Gm is small. However, if the blend Tg approaches or exceeds the melting point (Tm°) k2 can prohibit crystallization regardless the value of ¨Fm*. Table 3.4 refers to a number of crystallizable miscible polymer blends for which the spherulite growth rate as a function of the crystallization temperature has been investigated. For most blends only a part of the bell-shaped curve could be measured. In Figure 3.4 the complete bell-shaped spherulitic growth rate curve of iPS in iPS/PS blends containing 0, 15 and 30 wt% PS is shown. Due to the addition of impurity (e.g., the amorphous PS) a suppression of the growth rate is observed, which is greater the larger the concentration of the impurity added. Important parameters of the impurity added to the crystallizable component are the type, concentration, and molecular weight [Keith and Padden, 1964].

216

Table 3.4.

G. Groeninckx, M. Vanneste and V. Everaert

Spherulite growth rate measurements in miscible polymer blends (G versus Tc).

Polymer Blend

Amorphous Comp. (wt%)

Temperature Range Studied (Tc, °C)

Bell-shaped Curve

130-230

complete

References

iPS/PS

0-30% PS

PEG/PMMA PEG/PMMA a PEG/PMMA b PEG/PMMA PEG/PVAc

0-30% PMMA 0-30% PMMA 0-40% PMMA 0-40% PMMA 0-40% PVAc

40-55 10-60 35-55 44-58 45-55

part part part part part

[Cimmino et al., 1989] [Alfonso and Russell, 1986] [Martuscelli, 1984] [Calahorra et al., 1982] [Martuscelli, 1984]

PVDF/PMMA

0-50% PMMA

110-160

part

[Wang and Nishi, 1977]

20-35 30-41

part part

[Ong and Price, 1978b] [Nojima et al., 1986]

PCL/PVC

a b

25-50% PVC 0-10% PVC

[Keith and Padden, 1964]

Several PMMA polymers with different molecular weights were used Several PEG polymers with different molecular weights were used

By means of several optical techniques, viz. small angle laser light scattering (SALLS), optical microscopy, etc, the spherulite structure can be studied. From the photographic scattering pattern the spherulitic radius, R, can be calculated as a function of the crystallization time and/or blend composition [Stein, 1964]: R = 4.1λ/4π {1/sin(0.5 θm)}

(3.5)

θm represents the azimuthal angle of the intensity maximum; R the spherulite radius; and λ the light wavelength in the medium. A general observation is a decrease of the spherulitic radius with increasing content of the amorphous polymer when a same crystallization time is used (see Figure 3.5 and Table 3.5: PCL/PVC).

Figure 3.5. Spherulite growth of PCL/Phenoxy blends at Tc = 45°C [Defieuw et al., 1989d].

3.3.3.3 Determination of the Lateral and Fold Surface Free Energies from the Growth Rate

Gm = {φ2 G° exp[-¨E/R(Tc - To’)] 2 E/L}/

Alfonso and Russell [1986] related the different terms in Eq 3.2 to the measurable or characteristic properties of the blend, which resulted in the following relation for blends in which the amorphous component segregated into the interlamellar regions:

{G° exp[-¨E/R(Tc - To’)] + 2 E/L} exp{(-2bσσe)/ (kBT¨hu f [1 - Tc/Tm° - RTV2u χ (1 - φ2)2/ ¨hu f V1u] )}

(3.6)

where: b is the thickness of a monomolecular layer; σσe is the product of the lateral and fold

Crystallization, Morphological Structure, and Melting of Polymer Blends

Table 3.5.

Maximum spherulite radius, Rmax, as a function of crystallization time (tc) and blend composition

Polymer Blend

Composition

Tc

(wt%)

(°C)

100/0 90/10 80/20 70/30 60/40 50/50

30 30 30 30 30 30

21 b 36 b 31 b 26 b 19 b 10 b

c

75/25

20 25 30 35

33 32 60 63

7 9 20 33

[Ong and Price, 1978b]

90/10

33.2 35.1 37.8 39.2

136 157 150 122

8 12 19 31

[Nojima et al., 1986]

PCL/CPE 42.1 d

100/0 90/10 80/20

45 45 45

70 150 119

3 25 43

[Defieuw et al., 1989a]

PCL/Phenoxy

100/0 90/10 80/20 70/30

45 45 45 45

168 119 114 160

19 25 61 209

[Defieuw et al., 1989d]

PCL/SMA 14 e

100/0 90/10 80/20

45 45 45

165 169 111

16 77 187

[Defieuw et al. 1989b]

PEG/iPS

100/0 90/10 70/30

ns

ns

PCL/PVC

a b c d

e

217

Rmax measured (µm) a

tc (min) required to obtain Rmax a

References

[Khambatta et al., 1976]

c c c c c

[Wenig et al., 1975]

Extrapolated values from figures A mean value is given, obtained by various optical techniques The crystallization was allowed to proceed for more than five half-times of crystallization for each composition CPE with 42.1 wt% chlorine; PCL/CPE 42.1 shows an LCST-behavior (LCST = 147°C); the experiments were performed on specimens prepared below the LCST SMA with 14 wt% MA

218

G. Groeninckx, M. Vanneste and V. Everaert

surface free energies; Viu is the molar volume of component i; χ is the Huggins-Flory interaction parameter; and ¨hu is the heat of fusion per mole of monomer of the crystallizable component, the temperature dependence of which is taken into account by the parameter f: f = [2Tc / (Tc + Tm°)]

(3.7)

Both σσe and χ are assumed to be independent of temperature and composition. In Eq 3.6 the ratio E/G, a modified version of the δ-parameter [Keith and Padden, 1963, 1964] appears. This length, relative to the thickness of the crystalline lamellae (L), is critical for the consideration of the crystal growth in crystallizable miscible polymer blends. Eq 3.6 can be written as: α = - σσeiβ

α = lnGm - lnφ2 - lnG° + ¨E/R(Tc - To’) + (3.9)

or

α = lnG° - (Kg/Tc¨Tf)

(3.10)

with: Kg = (n b σσe Tm°)/(¨hu kB)

β = (2b / kBTc) {(¨huf ¨T/Tm°) - [RTV2u χ (1 - φ2)2]/V1u}-1 (3.14)

α contains parameters associated with the kinetic processes, while thermodynamic variables are met in β. If the product σσe is independent of the blend composition and temperature, then, according to Eq 3.8 a curve of α versus β should produce a straight line, regardless of the concentration and molecular weight of the amorphous component. The slope of such a plot is a measure of the product σσe. Another way to rewrite Eq 3.6 is [Cimmino et al., 1989]: α = lnGm - lnφ2 + C1 / [R(C2 + Tc - To] [(0.2 Tm° lnφ2) / (Tm° - Tc) ]

(3.15)

(3.8)

where:

ln{1 + [G°L exp[-¨E/R(Tc - To’)]/2 E]}

and:

(3.11)

where n is 2 or 4 depending on the regime of crystallization [Hoffman, 1982; Ong and Price, 1978b; Runt and Martynowicz, 1986]. The value of n = 2 refers to intermediate growth behavior (regime II), while the value of n = 4 corresponds with regime I and III in which low and high undercooling, respectively, is taking place. Furthermore, based on the WLF relation [Williams et al., 1955] the growth rate can be written as [Ong and Pric, 1978b]:

α = lnG° - C3 { Tm°/Tc (Tm °- Tc)} where:

(3.12)

C3 = (4 b σσe)/(kB ¨hu)

(3.13)

Note that K = C3 Tm when n = 4 (regime I or III).

Although equilibrium melting points, Tm° = Tc, should be used, generally the experimental Tm values are used. Considering the Eqs 3.10 and 3.12 a plot of α versus 1/(Tc¨Tf) and Tm°/(Tc¨T), respectively, should result in a straight line from which σσe can be obtained - see Figure 3.6 where a plot of α as a function of Tm°/(Tc¨T) is shown. The straight line in Figure 3.6 represents a fit of Eq 3.12 to the experimental values using the WLF constants, C1 = 17,250 cal/mol and C2 = 72 K (see Eq 3.15); the latter value being higher than the true WLF value of 51.6. Other authors, however, also had to use higher C2 values to fit their growth rate data [Hoffman and Weeks, 1962a; Magill, 1964; Boon et al., 1968]. The good fit in Figure 3.6 indicates that the temperature dependence of the spherulite growth rate of a crystallizable component in miscible blends is quite similar to that of homopolymers. It is also obvious from this figure that σe is independent of the concentration of the amorphous component (PMMA). Caution should be taken to generalize these data since: 1) the high concentration diluent was not investigated 2) the temperature range was near the melting point. The same observations were, however, made by Ong and Price [1978b] and by Wang and Nishi [1977].

Crystallization, Morphological Structure, and Melting of Polymer Blends

219

molecular weight of PEG used by Calahorra is much higher (Mw = 400 kg/mol) compared to the PEG used by Martuscelli (2 and 10 kg/mol). The value of σe was seen to depend on the molecular weight of PEG [Martuscelli, 1984], being smaller in the case of blends containing PEG with lower molecular weight (see Figure 3.8).

Figure 3.6. Plot of α versus Tm/(Tc¨T) for various compositions of a PEG/PMMA blend (triangles, 100/0; circles, 90/10; squares, 80/20; crosses, 70/30; filled circles, 60/40) [Calahorra et al., 1982].

3.3.3.4 Influence of the Molecular Weight of the Amorphous Component Alfonso and Russell [1986] found a significant curvature in the α versus β plots of PEG/PMMA blends (see Eq 3.8), while they were linear for neat PEG. The curvature could be due to an increase of σσe with a decreasing temperature. These authors also studied the influence of the molecular weight of the amorphous component (PMMA) on the spherulite growth rate of PEG. Noteworthy is the discrepancy seen at low undercooling for one of the blends containing PMMA with a molecular weight corresponding to the critical molecular weight for entanglement. Superposing all data for different molecular weights (above the critical value) results in a true master curve (see Figure 3.7), which shows that Eq 3.8 accounts quite well for the effect of molecular weight.

Figure 3.7. Master curve of α versus β for blends of PEG (145) with PMMA (125) and PMMA (525) (values between brackets refer to the molecular weight of the components in kg/mol) [Alfonso and Russell, 1986].

3.3.3.5 Influence of the Molecular Weight of the Crystallizable Component Martuscelli [1984] studied the influence of the molecular weight of the crystallizable component (PEG) on the spherulite growth rate of PEG/PMMA blend. In contrary to Calahorra et al. [1982] they found that the fold surface free energy, σe, decreases with increasing PMMA content in the blend. It should be mentioned, however, that the

Figure 3.8. Surface free energy of folding, σo, as a function of the PMMA content for PEG/PMMA blends using two PEG polymers differing in molecular weight (2 and 10 kg/mol) [Martuscelli, 1984].

220

G. Groeninckx, M. Vanneste and V. Everaert

Several authors [Ong and Price, 1978b; Alfonso and Russell, 1986, Runt and Martynowicz, 1986; Cimmino et al., 1989] used one of the equations mentioned above to calculate G°, σ, σo and/or σσo (see Table 3.6). The following empirical relationship [Thomas and Staveley, 1952; Geil, 1963; Vidotto et al., 1969] was developed:

σ = 0.1 b ¨hu

(3.16)

3.3.3.6 Influence of Copolymer Composition The influence of the SAN copolymer composition on the spherulitic growth rate of PCL has been studied at a fixed crystallization temperature by Kressler et al. [1992, 1993]. A minimum has been observed at about 20 wt% AN in SAN for several compositions (see Figure 3.9), due to a minimum in the value of the interaction parameter, χ, at the same copolymer composition that is responsible for a reduced chain mobility. The same authors also investigated the influence of the copolymer composition of SAN in PCL/SAN blends on G° and (σσe)1/2. The plot of (σσe)1/2 versus the acrylonitrile content in SAN shows a minimum (Figure 3.10), suggesting that the addition of SAN results in a stabilization of the growing PCL crystallites. This effect was more pronounced when the interactions between SAN and PCL, indicated by χ, are more favorable. Since G° is proportional to |χ - χs| [Saito et al., 1991], with χs the interaction parameter at the spinodal, a minimum was also noticed in the G° versus the copolymer composition of SAN (see Figure 3.11). 3.3.4

Figure 3.9. Dependence of the spherulite growth rate G on the copolymer composition of SAN in PCL/SAN blends at 45°C [Kressler et al., 1992, 1993].

Figure 3.10. Values of (σσe)1/2 versus the copolymer composition of SAN in PCL/SAN blends [Kressler et al., 1992, 1993].

Overall Crystallization Kinetics

3.3.4.1 General Aspects of the Avrami Theory Under Isothermal Conditions The overall crystallization kinetics of blends can often be described by the Avrami equation [Avrami, 1939]:

α = 1 - exp{-k tn}

(3.17)

Figure 3.11. The pre-exponential factor G° versus the copolymer composition of SAN in PCL/SAN blends [Kressler et al., 1992, 1993].

221

Crystallization, Morphological Structure, and Melting of Polymer Blends

Table 3.6.

Preexponential factor (G°), lateral (σ), and fold surface (σe) free energy and/or their product (σσe) as calculated from Eqs 3.8, 3.11 or 3.13

Polymer Blend

Composition

G° (x10-3 m/s)

PEG/PMMA

100/0 a 13 000 145 000 594 000 990 000

PEG/PMMA

0-40% PMMA

PEG/PMMA

100/0 90/10 80/20 70/30 60/40

PEG/PEMA

100/0 90/10 80/20 70/30

PCL/PVC

25-50 wt% PVC d

PCL/SAN19.5 e

90/10 80/20 70/30 60/40

PVDF/PMMA f

100/0 75/25 50/50

a b c d

e f g

h i

σσe (x 106 J2/m4)

σe (x103 J/m2)

σ (x103 J/m2)

References

[Alfonso and Russell, 1986] 3.6 2.2 8.0 48.3

363 496 627 724

4.1

186

18.8

9.9 h

58/60/57 i 43/48/26 i 38/39/27 i 36/39/22 i 36/37/— i

0.72

175

[Martuscelli et al., 1984]

58 28 24 14

b

27

6.5 c

464 495 396

[Cimmino et al., 1989]

b b b

567 507 518 491 0.182 0.131 0.012

[Calahorra et al., 1982]

[Ong and Price, 1978b] [Kressler et al., 1993]

47 51 40

9.76 g 9.76 g 9.76 g

[Wang and Nishi, 1977]

Several molecular weight PEG polymers were used σ was calculated by Eq 3.16, using b = 4.65A [Vidotto et al., 1969] σ was calculated by Eq 3.16, using b = 4.38A and ∆hu = 1.48x10+8 J/m2 Since all compositions could be superposed on one straight line in the α versus Tm/T∆T plot, all blend compositions have the same value for G°, σ, σe and σσe SAN with 19.5% AN; other copolymers have been used by the author leading to the same tendencies Other compositions have been studied too This parameter has been calculated using a modified version of Eq 3.16: σ = 0.1 (∆hu) (aobo)0.5, where ao = 5.43A, bo = 4.45A and ∆hu = 1.986x106 J/m2 This parameter has been calculated using Eq 3.16 where bo = 4.65A and ∆hu = 2.13x106 J/m2 [Van Krevelen, 1976] The first value is calculated from the spherulitic growth rate data of PEG(10)/PMMA. The second one from the overall crystallization data of PEG(10)/PMMA. The third one from the spherulitic growth rate data of PEG(2)/PMMA [Martuscelli and Demma, 1980] (the value between brackets refers to the molecular weight of PEG in kg/mol)

222

G. Groeninckx, M. Vanneste and V. Everaert

α is the weight fraction of crystallinity at time t, n is the Avrami index depending on the type of nucleation and the crystal growth geometry, and k is the Avrami constant related to the crystallization rate: k = ln(2/tn1/2 )

(3.18)

where tn1/2 is the half time of crystallization (the time for half the crystallinity to develop), which is often used as a measure for the overall rate of crystallization. The theory was applied to polymer systems, e.g., by Morgan [1954] and Mandelkern [1954]. In Figure 3.12 typical crystallization isotherms were obtained by plotting α versus the crystallization time for the PEG/PEMA 80/20 blend at different crystallization temperatures. From such curves, the half time of crystallization, tn1/2, can be deduced. Eq 3.17 can be rewritten as: log{-ln(1 - α)} = logk + n logt

(3.19)

Plotting the left part of this equation against logt should result in a straight line, from which both Avrami parameters, n (slope) and k (intercept), can be obtained. In Table 3.7 some literature data on the Avrami constants and the half time of crystallization are presented. Cimmino et al. [1989] calculated the half time of crystallization (t1/2) for some PEG blends, PEG/ PEMA, PEG/PVAc and PEG/PMMA, using the same blend composition and the crystallization temperature. Blends of PEG with PVAc had

the smallest t1/2, while the PEG/PEMA blends showed the highest values for the half time of crystallization. The type of amorphous component added to PEG seems to be important. The differences observed in t1/2 (and also in the values of G) depend on: • the degree of miscibility and mobility of the crystallizable and amorphous components. • the influence of the amorphous component on the nucleation of PEG. • influence of the noncrystallizable component on the secondary nucleation or the crystallization regime (neat PEG and PEG/PEMA crystallize in regime I, whereas PEG/PVAc and PEG/PMMA crystallize in regime II). Adding PMMA to PEG results in a decrease of k = ln2/tn1/2 (see Eq 3.18), an effect that is clearly seen in Figure 3.13 where 1/t1/2 is plotted against crystallization temperature [Martuscelli et al., 1984]. 3.3.4.2 Modified Avrami Expression It was often found that, contrary to the theoretical prediction, the value of n is non-integer [Avrami, 1939]. The Avrami model is based on several assumptions, such as constancy in shape of the growing crystal, constant rate of radial growth, lack of induction time, uniqueness of the nucleation mode, complete crystallinity of the sample, random distribution of nuclei, constant value of radial density, primary nucleation process (no secondary

Figure 3.12. Crystallization isotherms for the PEG/PEMA 80/20 blend crystallized at different Tc [Cimmino et al., 1989].

223

Crystallization, Morphological Structure, and Melting of Polymer Blends

Table 3.7.

Overall kinetic rate constant, k, Avrami index, n, and half time of crystallization, tn1/2, as function of the crystallization temperature and blend composition for some crystallizable miscible blends

Polymer Blend Composition PCL/PVC

70/30

65/35

60/40

50/50

PCL/PVC a

100/0

90/10

PEG/PMMA a

100/0

90/10

80/20

70/30

a b c

k (sec -n)

n

4.36 x 10-2 1.41 x 10-2 1.95 x 10-3 1.99 x 10-3 7.60 x 10-4 9.57 x 10-5 4.57 x 10-6 8.50 x 10-6 2.62 x 10-6 2.45 x 10-9 3.98 x 10-9 3.16 x 10-9

2.86 2.60 2.55 2.82 2.70 2.93 3.23 3.14 3.26 3.16 3.22 3.08

29.6 38.0 42.1 28.4 38.4 40.4

19.95 x 10-3 3.88 x 10-3 1.01 x 10-3 13.80 x 10-3 2.16 x 10-3 0.87 x 10-3

3.5 2.2 2.4 2.1 1.7 1.9

47 51 54 46 50 53 42 47 50 39 44 49

1.09 c 5.18 x 10-3 c 3.52 x 10-5 c 2.84 x 10-1 c 7.53 x 10-3 c 8.46 x 10-5 c 3.57 x 10-2 c 4.94 x 10-4 c 5.40 x 10-5 c 9.26 x 10-3 c 2.90 x 10-4 c 6.05 x 10-6 c

2.53 2.78 2.63 2.52 2.67 2.57 2.45 2.46 2.50 2.56 2.82 2.72

Tc (°C) 5 15 25 5 15 25 5 15 25 5 15 25

tn1/2 (min)

Reference

3.6 6.2 20 10 15 37 46 35 58 630 480 595

[Ong and Price, 1978b]

45 b 217 b 848 b 61 b 374 b 949 b 0.8 5.8 42.9 1.4 5.4 33.3 3.4 19.0 39.1 5.7 16.8 72.4

[Nojima et al., 1986]

[Martuscelli et al., 1984]

The authors investigated a wide range of crystallization temperatures, from which only a few are presented here The half time of crystallization is expressed in seconds The overall kinetic rate constant is expressed in min-n

nucleation) and absence of overlap between the growing crystallization fronts. These assumptions are often not met in polymer (blend) crystallization. Also, erroneous determination of the ‘zero’ time and an overestimation of the enthalpy of fusion of the polymer at a given time can lead to noninteger values for n [Grenier and Prud’homme, 1980].

Pérez-Cardenas et al. [1991] developed a modified Avrami expression, taking into account the secondary crystallization effects. The weight fraction of crystallinity, α, can be written as the sum of two terms:

α = αp + αs

(3.20)

224

G. Groeninckx, M. Vanneste and V. Everaert

Figure 3.13. Reciprocal of the half time of crystallization, t1/2, versus Tc for neat PEG (10) and PEG (10) blends with PMMA (the value between brackets refers to the molecular weight of PEG in kg/mol) [Martuscelli, 1984].

where the subscripts ‘p’ and ‘s’ refer to primary and secondary crystallization, respectively. The crystallization process is divided in three regions (Figure 3.14): (I) the initial primary crystallization region, (II) a region in which both primary and secondary crystallization takes place and finally (III) a region in which only secondary crystallization occurs. A parameter, ζ, was introduced, which is the weight fraction of the polymer crystallized by primary and secondary crystallization at the moment that the primary crystallization has ended (end of region II). The whole crystallization process is then described by two equations: 1 - α = exp (-kt n-k’t n’) [kn (1 - ζ) t

œ exp (kτ n + k’τ n’) τ n-1 dτ + 1]

(3.21)

1 - α = (1 - ζ) exp (k’t *n’) exp(-k’t n’)

(3.22)

o

Eq 3.21 is valid for α ≤ ζ and Eq 3.22 for α > ζ. Instead of two Avrami parameters, five parameters are required to describe the process. They have the following physical meaning: k and n (the primary crystallization parameters) depend on crystallization temperature, nature of primary nucleation and the fast growth; the secondary

Figure 3.14. Comparison between a typical experimental crystallization isotherm (solid line) and the Avrami equation (Eq 3.17, broken line). The three regions I, II and III correspond to primary, primary and secondary, and secondary crystallization, respectively. [Pérez-Cardenas et al., 1991].

crystallization parameters, k’ and n’ depend on the conditions under which the slow crystallization of remaining amorphous regions takes place; and a fifth parameter, ζ, indicates the weight fraction of material crystallized up to the moment the primary crystallization ends. t* is the moment at which the third region starts (e.g., pure secondary crystallization). Some literature data concerning isothermal crystallization experiments of linear PE at 128°C [Doremus et al., 1958] have been fitted using different values for the parameters in Eqs 3.21 and 3.22 (Figure 3.15). The most accurate fit was obtained using the following parameters: n = 4, k = 3.7 x 10-10, n’ = 2, k’ = 4 x 10-6 and ζ = 0.68. In the case of miscible polymer blends, the temperature dependence of the overall kinetic rate constant, k, can be calculated from [Boon and Azcue, 1968; Wunderlich, 1973; Hoffman, 1982]: 1/n lnk - lnφ2 + ¨E/{R(Tc - To)} - {(0.2 Tm° lnφ2)/¨T} = α2 = lnAn - {Kg /Tc¨Tf} (3.23) with Kg the same as in Eq 3.11.

Crystallization, Morphological Structure, and Melting of Polymer Blends

225

Figure 3.15. Theoretical isotherms (solid lines) using Eq 3.21 and 3.22 for three different sets of values of the five parameters [Pérez-Gardenas et al., 1991] fitted to the experimental values (points) of Doremus et al. [1958].

Figure 3.16. Surface free energy of folding, σe, versus the volume fraction of the crystallizable component, φ2, for blends of PEG (10) with PMMA from spherulite growth rate data (circles) and from overall rates of crystallization data (triangles) (the value between brackets refers to the molecular weight of PEG in kg/mol) [Martuscelli et al., 1984].

3.3.4.3 Determination of the Surface Free Energy of Folding from Overall Kinetic Data

obtained. In Table 3.8 and Figure 3.16 the free energy of folding, σe, for some PEG/PEMA and PEG/PMMA blends, respectively, derived from the overall kinetics of crystallization (Eq 3.23) is compared with the values obtained from the radial

A plot of α2 versus 1/(Tc¨Tf) results in a straight line and from the slope, values of σe can be Table 3.8.

Free energy of folding (σe) for some PEG/PEMA blends calculated using Eq 3.10 and Eq 3.23 Composition (wt%)

σe (x103J/m2) (Eq 3.10)

σe (x103J/m2) (Eq 3.23)

PEG/PEMA

100/0 90/10 80/20 70/30

58 28 24 14

75 42 34 29

[Cimmino et al., 1989]

PEG/PMMA

100/0 90/10 80/20 70/30 60/40

58 43 38 36 36

60 48 39 39 37

[Martuscelli et al., 1984]

Polymer Blend

References

226

G. Groeninckx, M. Vanneste and V. Everaert

growth rate data (Eq 3.10). The compositional dependence of σe derived from both methods is similar, although higher values were obtained using Eq 3.23 (overall kinetics of crystallization). The σe values obtained from both analyses almost coincide. The dependence of σe on the composition of the PEG/PMMA blends may be partly accounted for by the effect of concentration, since the concentration dependent part of σe is only a few joule per square meter [Martuscelli, 1984]. It is possible that PMMA molecules located in the interlamellar regions easily form entanglements with PEG molecules, favoring the formation of large loops on the surface of the PEG crystals. This can lead to an increase of both the surface enthalpy and the entropy of folding which contribute to σe (σe = He - TSe). The decrease of σe when adding PMMA suggests that the entropic term overwhelms the enthalpic one. 3.3.4.4 Non-Isothermal Kinetics The theory of Avrami is limited to isothermal processes. Since polymer processing is mostly performed under nonisothermal conditions, the theory has been extended [Ziabicki, 1967; Ozawa, 1971; Ziabicki, 1976]. According to Ozawa [1971] the crystallinity at any temperature is given by: - ln(1 - α) = C(t)/qn”

(3.24)

where q is the heating or cooling rate; C(t) is a cooling function of the process; and n” is the Ozawa exponent. Ziabicki [1967; 1976] based his analysis on the assumption that any nonisothermal process can be treated as a combination of several isothermal crystallization steps: E(t) = k tn = ln2 {0 œ t ds/t1/2 [α(s)]}n

(3.25)

This equation is an analogue of Eq 3.18. Under nonisothermal conditions, the ultimate Avrami parameters will be averages of the parameters of the subsequent steps. Gupta et al. [1994] used the following equation, based on the theory described by Ziabicki [1967; 1976], and by Kamal and Chu [1983], to

describe the nonisothermal behavior of cocrystallizing HDPE/LLDPE blends. [αp’ (Tp - Tons)]/[β (1 - αp)] = (n - 1) [E (Tp - Tons)]/[RTp2]

(3.26)

where α’=dα /dt = nkt(n-1) (1 - α); and β is the heating rate. In the case of a cooling experiment, β will be negative and β should be replaced by -β in the equation; Tons is the onset temperature and Tp is the temperature after time t (Tp - Tons = β t), where the subscript ‘p’ denotes the peak temperature; and E is the activation energy of the crystallization process. A plot of the left part of Eq 3.26 against (Tp - Tons)/Tp2 should be linear with slope and intercept equal to E/R and (n - 1), respectively (see Figure 3.17). The quantity α was evaluated from the ratio of the area under the crystallization peak per unit mass of the sample. αp is the extent of crystallization at the peak maximum and is determined by the fractional area under the exotherm from the onset temperature Tons to the peak temperature Tp relative to the total area under the exotherm. The Avrami exponent and the activation energy decrease with increasing LLDPE content (see Table 3.9). The authors suggested that the Avrami constant could be seen as the sum of two processes: a contribution due to nucleation and a contribution due to growth: n = nnucleation + ngrowth

(3.27)

Since both components cocrystallize [Edward, 1986; Hu et al., 1987; Gupta et al., 1994] the crystalline growth can be considered to be identical. As a consequence, the differences seen in the Avrami exponent, n, in the blends must be due to a difference in the nucleation behavior that depends on the blend composition. A value varying from 0 (instantaneous nucleation) to 1 (sporadic nucleation) was attributed to the contribution of the nucleation of LLDPE and HDPE, respectively, to the Avrami constant, n. The remaining part of n (e.g., 1.94 for HDPE

227

Crystallization, Morphological Structure, and Melting of Polymer Blends

crystalline boundary that exert repulsive forces. Cocrystallization enhances these forces due to the greater abundance of the bulky groups that results in a decrease of the activation energy. 3.3.5

Melting Behavior of Crystallizable Miscible Blends

3.3.5.1 The Equilibrium Melting Temperature in Miscible Blends - Hoffman-Weeks Plot In a semicrystalline homopolymer, the change in free energy of melting per mole of monomer unit is given by ¨Gu(T) = ¨Hu - T¨Su

(3.28)

where ¨Hu and ¨Su are the enthalpy and the entropy changes on melting, respectively. For blends, the difference in free energy of the crystalline unit can be written as [Sanchez and Di Marzio, 1971]: Figure 3.17. Plot of [αp’ (Tp - To )] / [β (1 - αp)] vs. (Tp - To ) / Tp2 for HDPE, LLDPE and the 50/50 blend [Gupta et al., 1994].

and 1.74 for LLDPE) represents the value for the growth process. The lower activation energy for LLDPE compared to HDPE seems to be due to a storage of thermal energy by the crystallites caused by the presence of bulky pendant groups at the

Table 3.9.

a

where ¨Gu(T) has the same meaning as in Eq 3.28, i.e., the heat of fusion of the crystalline component in the blend is assumed to be equal to that of the homopolymer. For athermal blends (¨hM = 0) Eq 3.29 becomes: ¨Gub(T) = ¨Hu - T¨Su - T¨sM = ¨Hu - T¨Sub

Avrami exponent, n, for HDPE, LLDPE and their blends

Blend Composition (wt%) a HDPE/LLDPE 100/0 75/25 50/50 20/80 0/100

¨Gub(T) = ¨Gu(T) + ¨gM = ¨Hu - T¨Su + ¨hM - T¨sM (3.29)

Avrami Exponent, n

Activation Energy, E (kJ/mol)

2.94 2.65 2.30 1.93 1.72

More compositions were mentioned in the article [Gupta et al., 1994]

121.67 99.44 85.83 59.04 49.16

(3.30)

228

G. Groeninckx, M. Vanneste and V. Everaert

For an infinitely thick crystal with an equilibrium melting temperature in the blend of Tmb°, ¨Gub(Tmb°) is equal to 0 and: Tmb° = ¨Hu /¨Sub

(3.31)

substituting Eq 3.31 into Eq 3.30 results in: ¨Gub(Tmb) = ¨Hu (1 - Tmb/ Tmb°) - Tmb¨sM (3.32) At Tm: ¨Gub(Tmb) = 2 (σeb/nb )

(3.33)

Combining Eqs 3.32 and 3.33 and rearranging, gives a relation between the experimental and equilibrium melting point in athermal polymer blends: 1/Tmb = {1/Tmb° + ¨sM/¨Hu}{1/[1 - (2σeb /¨Hunb)]} (3.34) α

φ1)βn

with σeb = φ1 σe and nb = (1 with α and β constants which need to be evaluated for each system [Cimmino et al., 1988]. For ¨sM → 0 the dependence can be simplified to: Tmb = Tmb°[1 - (2σeb /¨Hunb)]

(3.35)

Figure 3.18. Hoffman-Weeks plot for PCL rich PCL/PC blends (the data are displaced by 5°C to discern the different blend compositions) [Jonza and Porter, 1986].

1962b; Mandelkern, 1964]. If the crystals are perfect and of finite size and no recrystallization takes place during the melting, the Tm vs. Tc data can be described by [Nishi and Wang, 1975]:

If the heat of mixing is not ignored (¨hM ≠ 0), then the same treatment of Eq 3.29, as used to obtain Eq 3.35, results in:

Tm° - Tm = φ (Tm° - Tc)

Tmb = Tmb° {1 - [(2σeb /¨Hunb) + (¨gM/¨Hu]}

Tm = Tm° (1 - φ) + φTc

(3.36)

This equation is a general form of the relation between the experimental and equilibrium melting temperature of the blend. The equilibrium temperature of a polymer (blend) can experimentally be determined by a Hoffman-Weeks plot, which is a plot of the experimental melting point versus the crystallization temperature (Tm vs. Tc) as presented in Figure 3.18. Extrapolation from experimental data to the Tm = Tc line results in the value of Tm°. The influence of Tc on Tm is due to morphological contributions such as degree of perfection and the finite size of crystals [Hoffman and Weeks,

(3.37)

or: (3.38)

Tm° and Tm are the equilibrium and observed melting point, respectively; φ is the stability parameter that depends on the crystal thickness, and assumes values between 0 and 1; (sometimes φ is replaced by 1/γ, where γ is the thickening factor of the crystal). The value of φ = 0 implies Tm = Tm° for all Tc , whereas φ = 1 implies Tm = Tc. Therefore, the crystals are most stable at φ = 0 and inherently unstable at φ = 1. Nishi and Wang [1975] examined the polymer system PVDF/PMMA and found a value for φ = 0.2 for all compositions studied, which suggests that the crystals are fairly stable.

Crystallization, Morphological Structure, and Melting of Polymer Blends

A comparable value has been found for other polymer crystals [Hoffman and Weeks, 1962b]. The same polymer blend was studied by Stein et al. [1981; Morra and Stein, 1984]. Since PVDF crystallizes into several types of morphologies, different lines are shown in the Hoffman-Weeks plot (Figure 3.19). The curve representing the melting point of PVDF as a function of the crystallization temperature for the α modification shows a break that was associated with defect exclusion from the crystal [Stein et al., 1981] and by entrapment of head-to-head defects of the PVDF chains into the crystals during rapid crystallization at large undercooling [Morra and Stein, 1984]. Hoffman-Weeks plots have also been drawn for several other amorphous/crystalline miscible blends, such as PVDF/PEMA [Eshuis et al., 1982], PEG/PMMA [Martuscelli, 1984], PCL/SARAN [Zhang and Prud’homme, 1987], as well as for some miscible blends containing two semicrystalline components, PCL/PC [Jonza and Porter, 1986] and PCL/Penton [Guo, 1990]. Table 3.10 represents equilibrium melting points derived from Tm vs. Tc plots for some of these systems. It should, however, be noted that several blends do not show a linear Tm vs. Tc relation [Rim and Runt, 1984; Jonza and Porter, 1986]. The absence of linearity originates from several effects, such as recrystallization, crystal defects, etc.

Figure 3.19. Hoffman-Weeks plot for PVDF/PMMA blends with 50.1 vol. % of PVDF [Stein et al., 1981].

229

3.3.5.2 Melting Point Depression The melting behavior of a semicrystalline component in a miscible blend strongly depends on the blend composition. In several blends a depression of the melting point has been observed after addition of an amorphous polymer. This behavior results from the kinetic, morphological, and thermodynamic factors [Cimmino et al., 1989]. Kinetic effects originate from the crystal formation at temperatures below the equilibrium melting point. They can be avoided by using equilibrium values derived from HoffmanWeeks plots. Melting point depression, caused by morphological effects (see Eq 3.38), is associated with changes in crystal thickness, perfection and geometry, as well as with different thermal histories of the samples. When a miscible diluent is added to a semicrystalline polymer, the equilibrium melting point of the crystallizable component can be depressed due to interaction between both components. The free energy of the crystallizable component will decrease from Gm° to Gmb°, when the crystallites are surrounded by a mixed melt phase. The free energy of the crystalline phase, Gc, is not affected by mixing. The melting temperature, defined as the cross section of Gc and Gm (e.g., when ¨G = 0), may be depressed (Figure 3.20). The melting point depression resulting from thermodynamic effects can be described by the following equation [Flory, 1953; Nishi and Wang, 1975]:

Figure 3.20. Schematic diagram of the free energy of a crystalline phase (Gc) and the free energy of a melt phase of a homopolymer (Gm°) and a miscible blend (Gmb°) as function of the temperature.

230

G. Groeninckx, M. Vanneste and V. Everaert

Table 3.10. Equilibrium melting points derived from Hoffman-Weeks plots for several crystallizable miscible blends Polymer Blend

Composition a

PVDF/PMMA

PCL/PC b

Tm° (°C)

References

100/0

173.8

[Nishi and Wang, 1975]

50/50

165.2

100/0

71

10-40% PC

71 ± 2

100/0

58.1

50/50

55.4

PCL/P(VCl2-VA)

50/50

55.3

PCL/P(VCl2-AN)

50/50

53.6

PCL/Penton c

0/100

185

50/50

172.5

PCL/P(VCl2-VC)

PEEK/PEI

iPS/PPE a b

0-60% PEI

0-35% PPE

[Jonza and Porter, 1986]

[Zhang and Prud’homme, 1987]

[Guo, 1990]

384

[Chen and Porter, 1993]

389

[Lee and Porter, 1987]

240

[Plans et al., 1984]

Most authors studied more compositions than the ones presented here. Both polymers crystallize in this blend, but no Tm - Tc plots could be made for PC since the blends are reactive at the higher crystallization temperatures required for this component.

c

Both polymers crystallize in this blend.

(1/Tmb°) - (1/Tm°) = -[(RV2u)/(¨hu V1u)] [(ln φ2/m2) + (1/m2-1/m1)(1 - φ2) + χ12 (1 - φ2)2]

(3.39)

Tmb° and Tm° are the equilibrium melting point of the blend and the neat crystallizable component, respectively; Vu is the molar volume of the repeating unit of the components (1 = amorphous component and 2 = crystallizable component); ¨hu is the heat of fusion per mole of repeating unit; m is the number of units in the molecule, i.e., the degree of polymerization; φ is the volume fraction; and χ12 is the polymer-polymer interaction parameter. Since, for polymers m → ∞, Eq 3.39 can be reduced to: (1/Tmb°) - (1/Tm°) = -[(RV2u)/(¨hu V1u)] χ12 (1 - φ2)2 (3.40)

Instead of χ12, the interaction energy density, B12, can be used; both parameters are related by:

χ12 = B12 V1u /RT

(3.41)

Melting point depression data are often used to determine the Huggins-Flory interaction parameter, χ12 (see Table 3.11), that is a measure for the miscibility of the blend, i.e., χ12 is negative for a miscible blend. A lack of melting point depression means that χ12 is zero. Eq 3.39 is only valid for systems in which the crystalline morphology is not affected by the composition. Many authors [Hay, 1976; Kwei and Frisch, 1978; Rim and Runt, 1984; Plans et al., 1984; Alfonso and Russell, 1986], however, encountered difficulties when fitting Eq 3.40 to their experimental data, due to [Rostami, 1990]:

Crystallization, Morphological Structure, and Melting of Polymer Blends

Table 3.11.

Interaction parameters, χ12 and B12, derived from the melting point depression data

Polymer Blend

Interaction Parameter, χ12

c d e

f g

References

-0.295 (160°C) -0.34 (160°C)

-12.48 -11.94 -13.11

[Nishi and Wang, 1975] [Kwei et al., 1976] [Imken et al., 1976]

PEG(20 000)/PMMA a PEG(100 000)/PMMA

-1.93 (76°C) -0.35 (74°C)

-65.32 -11.93

[Martuscelli and Demma, 1980] [Martuscelli et al., 1984]

PEG/PVC g

-0.094 (65°C)

PCL/SAN 19.2 b

-0.18

-6.56 (65°C)

[Marco et al., 1993] [Kressler and Kammer, 1988]

PBA/ Phenoxy PEA/ Phenoxy PCL/Phenoxy

-16.20 (61°C) -9.67 (49°C) -10.09 (56°C)

[Harris et al., 1982] [Harris et al., 1982] [Harris et al., 1982]

PCL/Penton d

-15

[Guo, 1990]

-0.46 -0.02 -0.21 -0.53 -0.01 -0.28 -0.37

P(VCl2/VCl)/PDPS d P(VCl2/VCl)/PDPA d P(VCl2/VCl)/PCL d P(VCl2/VCl)/PCDS d

b

Interaction Parameter, B12 (x106J/m2)

PVDF/PMMA PVDF/PEMA

PCL/P(VCl2-VC) 80/20 c,d PCL rich P(VCl2-VC) rich PCL/P(VCl2-VA) 80/20 c,d PCL rich P(VCl2-VA) rich PCL/P(VCl2-AN) 80/20 c,d

a

231

[Zhang and Prud’homme, 1987] [Aubin et al., 1983] [Aubin et al., 1983] [Zhang and Prud’homme, 1987] [Aubin et al., 1983] [Aubin et al., 1983] [Zhang and Prud’homme, 1987] -0.84 -4.60 -8.37 -12.98

[Woo et al., 1983] [Woo et al., 1983] [Woo et al., 1983] [Woo et al., 1983]

PCL/PVDF d,e

-1.5

FVA/EVAc f PI/EVAc f PED/EVAc f

-0.06 -0.02 -0.38

PEEK/PEI

-0.40 (400°C)

iPS/PS (2 200) iPS/PS (50 000)

-0.002 -0.003

[Runt, 1981] [Runt, 1981]

iPS/PPE

0.17 -0.022

[Plans et al., 1984] [Runt, 1981]

[Jo et al., 1992] -15.07 -7.12 -115.56 -5.02

[Clough et al., 1994] [Clough et al., 1994] [Clough et al., 1994] [Chen and Porter, 1993]

The absolute value of χ12 is too large in comparison with the other literature data on miscible blends. The authors [Martuscelli et al., 1984] suggested that for this blend non-negligible entropic effects occur during mixing of the two polymers, non-compliance with the assumption inherent in the extrapolation of Tm (observed melting point) by using the Hoffman-Weeks plot and the inadequacies of the Huggins-Flory theory to describe the melting behavior of such polymer-polymer system. SAN containing 19.2 wt% AN More compositions have been investigated by the authors Both polymers are semi-crystalline This system shows an LCST behavior for PCL rich blends, while blends with a high concentration on PVDF are phase separated; the blends considered here are only the miscible ones below their LCST (blends with a content of PVDF less than 30 wt%) EVAc (molar ratio ethylene to vinyl acetate: 7:1) is amorphous The polymer blend PEG/PVC is only miscible when the concentration on PVC is ≥ 40%

232

G. Groeninckx, M. Vanneste and V. Everaert

• The use of observable melting temperatures instead of the thermodynamic equilibrium temperatures. • The plot of the left-hand side of Eq 3.40 versus the right-hand side has no zero-intercept. The intercept contains information about the crystalline morphology [Kwei and Frisch, 1978; Walsh et al., 1985], that has been ignored. Following Eq 3.39 the intercept should equal 1/m1. When high molecular weight polymers are used 1/m1 equals zero. This was observed for PVDF/PMMA en PVDF/PEMA blends [Kwei and Frisch, 1978]. • The concentration dependence of the interaction parameter adds a restriction on plotting the left-hand side of Eq 3.40 versus φ1 to obtain a single value for χ12 [Kwei and Frisch, 1978; Plans et al., 1984; Walsh et al., 1985]. A modified version of this equation has been used by some other authors [Kwei and Frisch., 1978; Walsh et al., 1985], who added a constant that is related to the morphology of the crystalline region: ¨hu (Tm° - Tmb°)/φ1 RTm° - Tmb°/m1 - φ1 Tmb°/2m2 = C/R - b φ1

(3.42)

where C is a constant taking into account the morphological contributions (which were assumed to be proportional to φ1) and b is a constant derived from the equation relating the interaction parameter with temperature: χ12 = a + b/T with a << b/T near the melting point. This approach was, however, not satisfactory [Walsh et al., 1985]. It should be noted that in Eq 3.42 the assumption of infinite molecular weights has not been included as was done in Eq 3.40. For some miscible blends, a melting point elevation has been reported with respect to that of the neat crystallizable component, both crystallized at the same temperature [Eshuis et al., 1982; Rim and Runt, 1983; 1984]. These observations may originate from recrystallization, enhanced crystal perfection, and increased crystal size. 3.3.5.3 Multiple Melting Endotherms The melting behavior of binary crystallizable blends often reveals multiple melting endotherms, that can be ascribed to recrystallization, secondary crystallization effects, phase separation, etc.

Recrystallization is a process in which the initial, rather imperfect, lamellae melt and recrystallize to produce thicker and more perfect lamellae that as a consequence melt at a higher temperature. As a result of this process a double melting behavior may be observed. Recrystallization has been observed for neat polymers and blends. A method that is often used to determine if the dual melting behavior is caused by recrystallization is the variation of the heating rate in DSC experiments. It is suggested that during the DSC run an annealing of the crystalline lamellae occurs (see Figure 3.21) [Rim and Runt, 1983]. At slow heating rates, the original crystals are given sufficient time to reorganize and the melting behavior is then mainly caused by lamellae originating from recrystallization (C) and melting of the recrystallized material (Mr). The resulting behavior is a composite of the peaks due to the melting of the original crystals (M), the recrystallization exotherm and the melting of the recrystallized material. As the heating rate is increased, the crystals have less time to reorganize, thus C and M decrease in magnitude. A plot of the observed melting point of both melting endotherms as function of the crystallization temperature (Hoffman-Weeks plot) is another method to detect whether recrystallization occurs. The melting point of recrystallized material is independent of Tc; thus a horizontal line is observed in the Tm vs. Tc plot. During crystallization of a miscible polymer blend, the composition of the amorphous phase changes, i.e., becomes poorer on the crystallizable component. In some cases, a liquid-liquid phase separation can take place as a result of the crystallization. This phenomenon will be discussed more in detail in the next section. A complex melting behavior is also observed when a semicrystalline polymer exhibits two different types of crystal structure. A second crystal structure can be introduced by a variation in temperature, pressure, elongation, etc. This phenomenon is known for neat PE. Adding an amorphous polymer to a crystallizable component can result in a change of the unit cell dimensions of the crystal structure. This has been observed for

Crystallization, Morphological Structure, and Melting of Polymer Blends

233

Figure 3.21. Schematic representation of the melting mechanism proposed to account for the heating rate dependence of recrystallizing material. The top of the figure shows the melting of the original crystals (M), recrystallization (C) and remelting (Mr). The bottom portion of the figure shows the resultant thermograms that are experimentally observed [Rim and Runt, 1983].

LDPE blended with EPDM [Starkweather, 1980], where the unit cell expanded in the a-direction (a raises from 7.515Å to 8.350Å) when increasing the amount of EPDM, while the b and c-directions remained almost unchanged. The composition of EPDM plays also an important role. EPDM with an ethylene/propylene mole ratio of 4.5 (EPDM-1) exhibits the behavior as mentioned above. Decreasing the ethylene content in the EPDM copolymer results in an amorphous polymer (designated EPDM-2 and EPDM-3 [Starkweather, 1980]) that do not alter the unit cell dimensions as much as EPDM-1 does. The latter copolymer is thought to cocrystallize (at least partially) with LDPE. Several blends prepared by co-precipitation followed by crystallization from the melt exhibit a double melting behavior, due to the occurrence of the secondary crystallization process. The amorphous component causes a retarded crystallization of some of the crystallizable chains, which form lamellae smaller than and located between the primary ones constituting the spherulites (see Figure 3.22). This is a phenomenon often observed in PCL blends. In DSC scans as a function of crystallization time (tc) a single melting behavior is observed

after short tc, while a second melting endotherm is noticed at long tc (see Figure 3.23). This second melting endotherm becomes more important the more amorphous component is added [Vanneste and Groeninckx, 1995] (see Figure 3.24). In Table 3.12 some blends are presented exhibiting a complex melting behavior due to one or more of the above-mentioned reasons. It should be mentioned that several homopolymers (of which polyethylene is probably the best known sample) also exhibit a complex melting behavior. Branched polyethylene’s (LDPE, LLDPE and VLDPE) show multiple melting endotherms, due to the presence of fractions with different branching content [Schouterden et al., 1985; Defoor et al., 1993]. This was clearly illustrated by Defoor et al. who fractionated LLDPE with respect to the short-chain branching content and blended the fractions with the highest and the lowest branching content. It was shown that they both crystallized and melted separately. Both fractions determined the spherulitic morphology in a co-operative way. Other examples are PPS (which shows a double melting behavior due to the obstructive effect of branching or crosslinking of the molecules on crystallization at high temperature [Mai et al., 1994]), and PEEK. Much controversy exists about

234

G. Groeninckx, M. Vanneste and V. Everaert

Figure 3.22. Schematic presentation illustrating the secondary crystallization process (thin lines: crystallizable component, thick lines: amorphous component).

Figure 3.23. Influence of the crystallization time on the melting behavior of the ternary PCL/Phenoxy/SAN 15 = 90/05/05 blend [Vanneste and Groeninckx, 1994].

the cause of the double melting behavior of PEEK — recrystallization or secondary crystallization. According to one group of authors [Prasad et al., 1991; Hudson et al., 1991; Bassett et al., 1992; Lattimer et al., 1992] PEEK that was crystallized from the melt contains crystals with two types of lamellar thickness. The ticker ones grow first, while the smaller ones grow later within the thicker lamellae. Thermal analysis, however, indicates that a process of melting, recrystallization and remelting occurs [Cheng et al., 1986; Lee and Porter, 1987; Lee et al., 1989; Crevecoeur and Groeninckx, 1991].

Figure 3.24. Influence of the concentration of the amorphous component on the amount of secondary crystallization in PCL/SAN 15/SMA 14 polymer blends [Vanneste and Groeninckx, 1995].

Crystallization, Morphological Structure, and Melting of Polymer Blends

235

Table 3.12. Examples of crystallizable miscible polymer blends exhibiting a complex melting behavior Polymer System

Type of Melting Behavior

PCL/SAN PCL/SAN

recrystallization dual melting behavior b

PCL/SAN PCL/Phenoxy/SAN 15 a

Required Conditions

References

melt crystallized

[Rim and Runt, 1983; 1984]

recrystallization recrystallization

melt crystallized melt crystallized short crystallization times

PEEK/PEI

recrystallization

melt crystallized

[Kressler and Kammer, 1988] [Vandermarliere, 1986] [Vanneste, 1993; Vanneste and Groeninckx, 1994] [Crevecoeur and Groeninckx, 1991]

LDPE/EPDM

different crystal types

ethylene/propylene ratio: 4.5

[Starkweather, 1980]

PCL/P(VCl2-VC) PCL/CPE c

secondary crystallization secondary crystallization

[Zhang and Prud’homme, 1987] [Defieuw and Groeninckx, 1989a]

PCL/SMA x d

secondary crystallization

PCL/Phenoxy

secondary crystallization

PCL/SAN x/SMA y f

secondary crystallization

PCL/Phenoxy/SAN 15 a

secondary crystallization

PEEK/PEI

secondary crystallization

high P(VCl2-VC) content e coprecipitation technique + melt crystallization coprecipitation technique + melt crystallization coprecipitation technique + melt crystallization coprecipitation technique + melt crystallization coprecipitation technique + melt crystallization g melt blended and crystallized

LLDPE/LLDPE h

secondary crystallization

a b c d e f g h

coprecipitation technique + melt crystallization

[Defieuw and Groeninckx, 1989b] [Defieuw and Groeninckx, 1989d] [Defieuw and Groeninckx, 1989c; Vanneste, 1993] [Vanneste and Groeninckx, 1994] [Bassett et al., 1992; Hsiao and Sauer, 1994] [Defoor et al., 1993]

SAN containing 15 wt% AN Reason for this dual melting behavior ‘is not completely clear’, but recrystallization is possibly occurring. PCL/CPE is only totally miscible for CPE containing 49.1 wt% chlorine x = 14 and 25 wt% Solution cast blends followed by melt crystallization x and y are 25 and 24 wt% and 15 and 14 w%, respectively in both references Only the blends with 90 wt% PCL are dealt with since only those combinations were found to be miscible A blend of 1-octene LLDPE fractions with different short chain branching content was investigated; i.e. 3 and 33 methyl groups per 1000 carbon atoms

3.3.6

Coupling of Demixing and Crystallization Phenomena

Tanaka and Nishi [1985] were the first to report the existence of coupling between crystallization and demixing in crystallizable blends. A competition between demixing and crystallization is seen in

binary blends of a semicrystalline and an amorphous polymer when the crystallization curve and the miscibility gap intersects. The morphology of blends exhibiting such behavior is determined by the ratio of the rate of crystallization and of demixing. Four important situations can be distinguished (Figure 3.25):

236

G. Groeninckx, M. Vanneste and V. Everaert

A/ Simultaneous spinodal decomposition and crystallization The blend is quenched into the unstable region of the miscibility gap and to a temperature below the crystallization/melt coexistence curve. B/ Simultaneous binodal decomposition and crystallization This type is similar to spinodal decomposition, but a composition is quenched into the metastable region of the miscibility gap. C/ Crystallization induced decomposition The blend is quenched outside the miscibility gap to a temperature below the crystallization/ melt coexistence curve. The concentration of the noncrystallizable component increases with crystallization until the miscibility gap is reached inducing demixing. D/ Decomposition induced crystallization The blend is quenched into the miscibility gap to a temperature that lies above the crystallization/melt coexistence curve for the actual composition, but lies below the crystallization curve for the binodal composition. When the blend is quenched, demixing occurs resulting in two coexisting phases of which one is able to crystallize. The demixing can result in spinodally as well as in binodally decomposed material.

Table 3.13.

Only few experimental studies have been performed on polymer blends exhibiting one or more of the phenomena described above (see Table 3.13). Routes A and C of Figure 3.25 were discussed by Tanaka and Nishi [Tanaka and Nishi, 1985; 1989] for a system consisting of PCL and PS. In case A coarse spherulite results including PS droplets, while in case C the spherulites are separated and show large droplets on their surface (see Figure 3.26). Li et al. [1991; 1993] investigated case B and C for the same system, i.e., PCL/low molecular weight PS, for which the phase diagram is presented in Figure 3.27. For route C three different regimes (Figure 3.28) can still be distinguished depending upon the rate of crystallization, G = dR/dt, and the rate of diffusion of the noncrystallizable component, vd = (D/tc)1/2 (D = diffusion constant and tc is the correlation time of the macromolecules). Parameters, vd and G, show a different dependence on temperature (see Figure 3.29). The growth rate, G, has a maximum between the melting temperature and the glass-transition temperature of the blend, whereas the diffusion rate of the amorphous component, vd, increases with temperature. 1/ vd << G: the noncrystallizable component is trapped within the growing crystals. Depending on the composition of the amorphous phase,

Polymer blends showing coupling of demixing and crystallization. The routes describing the type of coupling are explained in the text

Polymer Blend

Composition

Type of Coupling

40/60 60/40

route B route C

[Li et al., 1991]

PPE/PEG

12.1/87.9 46/54 93.7/6.3

route B route A route B

[Shibanov and Godovsky, 1989]

PEG/PPG

88/12 54/46 10.3/89.7

route B route A route B

[Shibanov and Godovsky, 1991]

PCL/PS

References

Crystallization, Morphological Structure, and Melting of Polymer Blends

Figure 3.25. Phase diagram of a binary polymer blend with miscibility gap (UCST) and intersecting crystal/melt coexistence curve. The Tm° curve is extrapolated into the miscibility gap. Quenching routes A to D are explained in the text. For routes B and C, the quenching-induced phase separation and crystallization are indicated. –––– binodal, ——- spinodal, -.-.-. crystal/melt coexistence curve [Li et al., 1991].

237

Figure 3.26. Phase separation behavior at the growth front of the spherulites of PCL during the crystallization process in the PCL/PS 70/30 blend at Tc = 50°C. These morphologies were observed at a) 1860 min, b) 2790 min, c) 3250 min, and d) 4230 min after quenching (bars: 20µm) [Tanaka and Nishi, 1989].

liquid-liquid demixing may occur resulting in droplets of non-crystallizing polymer inside the spherulites. 2/ vd § G: a part of the amorphous component is trapped and another part is segregated from the growing crystals. The concentration of this component increases with crystallization and finally demixing occurs resulting in the formation of droplets at the spherulite surface. 3/ vd >> G: the noncrystalline component is fully segregated into the bulk melt. When the miscibility gap is reached the melt phase separates homogeneously and binodally. The crystallization rate is retarded for all regimes, but the extent of hindrance increases from regime 1 to 3. It should be noted that the diffusion of the crystalline polymer occurs on a lamellar scale (about 10 nm), whereas the diffusion of the amorphous component, induced by demixing, takes place on a spherulitic scale (10-20µm). Under normal processing conditions,

Figure 3.27. Phase diagram of the PCL/PS blend. (triangles:

crystallization presumably takes place at a higher rate than the demixing.

238

G. Groeninckx, M. Vanneste and V. Everaert

Figure 3.28. Concentration distribution of PS around a spherulite at different times t and for different ratios G/vd. φo: start composition, φb: binodal composition and r: local co-ordinate a) regime 1, b) regime 2: the binodal composition is not exceeded, c) as in b) but the binodal composition is exceeded, d) superposition of the concentration distribution curves (dashed lines) according to the transition between regime 2 and 3 of two adjacent spherulites [Li et al., 1991].

Crystallization, Morphological Structure, and Melting of Polymer Blends

Figure 3.29. Temperature dependence of the diffusion driven displacement of the non-crystallizing component, vd, and the spherulitic growth rate, G [Li et al., 1991].

239

Figure 3.30. Morphology development in a PS blend with 60 wt% PCL at 44°C after a) 55 min, and b) 126 min. (bar: 50 µm) [Li et al., 1991].

The morphology resulting from the three regimes are presented in the Figure 3.30 (regime 1), Figure 3.31 (regime 2) and Figure 3.32 (transition from regime 2 to 3) for the PCL/PS system.

Figure 3.31.

Morphology development in a PS blend with 60 wt% PCL at 49°C after a) 92 h, and b) 142 h. (bar: 25 µm) [Li et al., 1991].

240

G. Groeninckx, M. Vanneste and V. Everaert

Figure 3.32.

Morphology development in a PS blend with 60 wt% PCL at 51°C after a) 91.5 h, b) 100 h, and c) 109 h [Li et al., 1991].

A demixing induced crystallization is shown in Figure 3.33 (route B in Figure 3.25) for the binary PCL/PS 40/60 blend. 3.3.7

Conclusions

Most of the fundamental and experimental aspects related to the crystallization phenomena occurring in miscible polymer blends are relatively well known. Much research has been done in the 1970’s and 1980’s, especially the development of the general theory concerning the crystallization process itself and the concomitant kinetics. These theories could be adapted to simple systems under quiescent conditions. Later, modifications have been made to the original concepts to take into account the effects occurring under processing conditions (for instance the non-isothermal Avrami theory), unusual phenomena not responding to the simple theory (for instance the nonlinearity of Hoffman-Weeks plots), coupling of the crystallization with demixing processes, etc. The addition of a second component to a crystallizable polymer has several profound consequences:

1) Depending on the glass-transition temperature of the added component, the crystallization window is widened or narrowed. 2) The type of added component is also important. Crystallization in the presence of an amorphous component gives rise to segregation. The segregation can occur into three regions: interspherulitic, interfibrillar and interlamellar, depending on the ratio of the diffusion rate of the amorphous component and of the crystallization rate of the crystallizable component. In blends of two crystallizable polymers the phenomena such as separate, concurrent and co-crystallization may take place. 3) The spherulite growth rate changes by blending due to: interactions between the components, the necessity of diffusion of both components, the concentration change in the amorphous phase during crystallization, the possible changes of the glass-transition and melting temperature. 4) The overall kinetics are strongly affected by the type of amorphous component, its influence on the nucleation of the crystallizable component, the degree of miscibility, the presence of secondary nucleation effects, the molecular weight of both components.

Crystallization, Morphological Structure, and Melting of Polymer Blends

Figure 3.33.

241

Phase separation followed by crystallization in a PS blend of 40 wt% PCL at 46°C after a) 2.5 h, b) 13.5 h, and c) 27 h [Li et al., 1991].

5) The melting behavior is often complex due to phenomena such as reorganization, secondary crystallization, demixing, etc. A depression of the equilibrium melting temperature is often observed.

3.4

Crystallization, Morphological Structure and Melting Behavior of Immiscible Polymer Blends

3.4.1

Introduction

From a commercial point of view, semicrystalline polymers are of prime importance. Among the four mostly used commodity plastics (PE, PS, PVC and PP), only PS is completely amorphous. The three semicrystalline polymers account for the largest volume of the commercial polymer blends. A majority of the polymer blends contains at least one crystalline component. Most polymer blends are immiscible. The immiscible semicrystalline polymer blends may be classified in terms of crystalline/crystalline systems in which both components are crystallizable, and crystalline/amorphous systems in which

only one component can crystallize, being either the matrix or the dispersed phase [Utracki, 1989]. Numerous authors have been investigating the crystallization behavior of immiscible blends. In Table 3.14 and Table 3.15, an overview is given of a number of important immiscible crystallizable blend systems. The properties of the finished articles made from immiscible blends are governed by the morphology created as a result of the interplay of processing conditions and inherent polymer characteristics, including crystallizability. Therefore, a scientific understanding of the crystallization behavior in immiscible polymer blends is necessary for the effective manipulation and control of properties by compounding and processing of these blends. In the following part, a discussion on the crystallization behavior in immiscible polymer blends is given, including the nucleation behavior, spherulite growth, overall crystallization kinetics, and final semicrystalline morphology. Each topic is illustrated with several examples from the literature, to allow the reader to find enough references on the discussed subject for further information.

242

G. Groeninckx, M. Vanneste and V. Everaert

Table 3.14. Thermal data on immiscible crystalline/amorphous blends (after Nadkarni and Jog [1991]) Comp. A Comp. B

Tg,A

(cryst.)

(amorph.)

(°C)

Tg,B (°C)

LDPE LDPE LDPE LLDPE

PS PMMA PC PS

-123 -123 -123 -123

HDPE HDPE

PS PC

PP

Tc,pure A (10°C/min)

Tm,pure A

Physical state of the amorph. comp. at Tc

References

103 100-123 147 103

101 101 101 102-110

115-123 115-123 115-123 125

± solidified solidified solidified viscous melt/±solid

[Baïltoul et al., 1981]

-123 -123

103 147

118 118

132-140 132-140

melt solidified

EPDM a) e)

-10

100

121

165

melt

PP

EPR a) e)

-10

±-45

121

165

melt

PP

PIB e

-10

---

121

165

melt

PP

PS

-10

103

121

165

melt

PP PP

PC TR (SBS)

-10 -10

147 ---

121 121

165 165

solidified melt

PET PET PET

PMMA PS PPE

70 70 70

100-123 103 215

150 150 150

257 257 257

melt melt solidified

PVDF

PVME

-45

140

172

[Müller et al., 1995], [Morales et al., 1995] [Aref-Azar et al., 1980] [Kunori and Geil, 1980] [Martuscelli, 1984],[Bartczak et al., 1984], [Martuscelli, 1985], [Martuscelli et al., 1983], [Karger-Kocsis et al., 1979], [Greco et al., 1987] [Martuscelli et al., 1982], [Coppola et al., 1987], [Bartczak et al., 1984], [Kalfoglou, 1985], [Karger-Kocsis et al., 1979], [Martuscelli, 1985] [Martuscelli, 1984], [Bartczak et al., 1984], [Martuscelli, 1985], [Martuscelli et al., 1982], [Martuscelli et al., 1983], [Bianchi et al., 1985], [Bartczak et al., 1987], [Wenig et al., 1990], [Wei-Berk, 1993], [Santana and Müller, 1994], [Han et al., 1977], [Hlavatá and Horák, 1994] [Ghijsels et al., 1982], [Karger-Kocsis et al., 1979] [Nadkarni and Jog, 1987] [Quirk et al., 1989] [Quirk et al., 1989], [Liang and Pan, 1994]

Crystallization, Morphological Structure, and Melting of Polymer Blends

243

Table 3.14. Thermal data on immiscible crystalline/amorphous blends (after Nadkarni and Jog [1991]) Comp. A Comp. B

Tg,A

(cryst.)

(amorph.)

(°C)

Tg,B (°C)

PEG

PS

-60

103

40

PEG

PI

-60

10

POM POM

PS PC

- 80 - 80

PA-6 PA-6

PC PS

PA-6 sPS PEMA a b c d e

Tc,pure A (10°C/min)

Tm,pure A

Physical state of the amorph. comp. at Tc

References

65

solidified

40

65

viscous melt

[Lotz and Kovacs, 1969] d, [O’Malley et al., 1969] d [Robitaille and Prud’homme, 1983] b

103 147

143 143

170 170

melt ± solidified

52 52

147 103

199 199

210 210

viscous melt melt

elastomer PVME c

52 95

---28

199 ± 245

210 270

melt

PMMA

67

105

[Chang et al., 1991]

[Ide and Hasegawa, 1974], [Chen et al., 1988] [Martuscelli, 1984] [Cimmino et al., 1991], [Cimmino et al., 1993] [Kwei et al., 1977]

EPR and EPDM are amorphous only if they are random copolymers Studies were done on liquid/liquid phase separated triblock copolymer PEG-PI-PEG sPS/PVME is immiscible as long as the PVME-content is higher than 10% Studies have been done on block copolymers of either PEG-PS-PEG or PEG-PS Small and defective PP molecules or amorphous PP chains are partially soluble in the elastomer phase; this can lead to a slightly deviating behavior [Martuscelli et al., 1983], [Kalfoglou, 1985]

3.4.2

Factors Affecting the Crystallization Behavior of Immiscible Polymer Blends

The discussion on the crystallization behavior of neat polymers would be expected to be applicable to immiscible polymer blends, where the crystallization takes place within domains of nearly neat component, largely unaffected by the presence of other polymers. However, although both phases are physically separated, they can exert a profound influence on each other. The presence of the second component can disturb the normal crystallization process, thus influencing crystallization kinetics, spherulite growth rate, semicrystalline morphology, etc.

Important factors are: • molecular structure and molecular mass of the components, • blend composition, • type and degree of dispersion of the phases in the melt state, • phase interactions (e.g., nature of the interface, migration of nuclei, etc.), • melt history (Tmelt, tmelt, etc.), • crystallization conditions (for example Tc, cooling rate, etc.), • physical crystallization conditions (surrounded by melt or solidified material).

-10

-123 -10

PP EPR c PA-6 PP

POM

PPS

PA-11

LLDPE

HDPE

LDPE

EPR c PA-6

PA-12 PEG

PVDF HDPE LLDPE PP PPS

HDPE

LDPE HDPE HDPE HDPE

HDPE

HDPE

PE

PP

PP

PP

PP PP

PP PP

PP PET PET PET PET

PPS

80

-10 70 70 70 70

-10 -10

-10

-10

-123

-123

-123

-123 -123 -123 -123

-123

PP

LLDPE

Tg,A

Comp. B (minor)

Comp.A (matrix)

-123

-45 -123 -123 -10 80

-60

±-30 52

-123

-123

-123

47

80

-80

-10 ±-40 52 -10

-10

Tg,B

256

121 150 150 150 150

121 121

121 121

121

121

121

100-120

118

101

101 118 118 118

102-110

Tc,A

118

140 118 102-110 121 256

154 38

199

101

118

102-110

256

143

199 121

121

121

Tc,B

280

165 257 257 257 257

165 165

165 165

165

165

165

115-140

132-140

115-123

115-123 132-140 132-140 132-140

125

Tm,A

132-140

172 132-140 125 165 280

171-179 65

±125 210

115-123

132-140

125

180

280

170

225 165

165

165

Tm,B

superheated melt

solid melt melt melt solidb

solid melt

melt solid

melt

supercooled melt

melt

solid

solid

solid

solid viscous melt solid solid

solid

[Zhou and Hay, 1993], [Long et al., 1991], [Müller et al., 1995], [Morales et al., 1995] [Teh, 1983] [Greco et al., 1987a] [Chen et al., 1988] [Bartczak and Galeski, 1986], [Martuscelli et al., 1980], [Rybnikâr, 1988] [Frensch et al., 1989], [Klemmer and Jungnickel, 1984] [Nadkarni et al., 1987], [Nadkarni and Jog, 1986], [Chen and Su, 1993] [Chen et al., 1988]

solid

[Nadkarni et al., 1987], [Nadkarni and Jog, 1986], [Jog et al., 1993]

[Long et al., 1991], [Flaris et al., 1993], [Zhou and Hay, 1993], [Plawky and Wenig, 1994], [Müller et al., 1995] solid [Bartczak et al., 1986], [Noel and Carley, 1984], [Teh et al., 1994a], [Lovinger and Williams, 1980], [Wenig and Meyer, 1980], [Teh et al., 1994b], [Gupta et al, 1982], [Martuscelli et al., 1980], [Rybnikâr, 1988] solid [Bartczak et al., 1984], [Martuscelli, 1985], [Galeski et al., 1984], [Teh, 1983], [Teh et al., 1994a] solid [Greco et al., 1987b], [Martuscelli, 1984] melt [Liang et al., 1983]a, [Ikkala et al., 1993], [Moon et al., 1994], [Park et al., 1990], [Grof et al., 1989]a, [Ide and Hasegawa, 1974] melt [Tang et al., 1994] solid [Bartczak and Galeski, 1986], [Tang and Huang, 1994b] supercooled melt [Shingankuli et al., 1988] solid [Wilfong et al., 1986] solid [Wilfong et al., 1986] solid [Wilfong et al., 1986] melt b [Shingankuli et al., 1988]

solid

melt

melt

melt

melt solid melt melt

melt

Physical condition for cryst. of A/B References

Table 3.15. Thermal data on immiscible crystalline/crystalline blends (after Nadkarni and Jog [1991])

244 G. Groeninckx, M. Vanneste and V. Everaert

HDPE LLDPE PVDF PVDF POM

PA-6

PA-66 PBT

PCL

PA-6 PA-6 PA-6 PA-66 PA-6

PVDF

PVDF PVDF

PVDF

c

b

-45

-45 -45

-45

52 52 52 60 52

52

80

Tg,A

-60

60 35

52

-123 -123 -45 -45 -80

-10

70

Tg,B

140

140 140

140

199 199 199 229 199

199

256

Tc,A

229 180

199

118 102-110 140 140 143

121

150

Tc,B

172

172 172

172

225 225 225 264 225

225

280

Tm,A

60

264 225

210

132-140 125 172 172 170

165

257

Tm,B

melt

solid

melt b melt b

melt b

solid b solid solid b

solid solid

solid solid

solid

solid

melt melt melt solid melt melt

melt

supercooled melt

[Frensch and Jungnickel, 1989], [Frensch et al., 1989] [Frensch and Jungnickel, 1991] [Frensch and Jungnickel, 1989], [Frensch et al., 1989]

[Ikkala et al., 1993], [Holsti-Mietinen et al., 1992], [Park et al., 1990], [Ide and Hasegawa, 1974], [Tang and Huang, 1994a] [Chen et al., 1988] [Frensch and Jungnickel, 1989], [Frensch and Jungnickel, 1991], [Frensch et al., 1989] [Frensch and Jungnickel, 1991]

[Shingankuli et al., 1988], [Jog et al., 1993]

Physical condition for cryst. of A/B References

Study done on fibers Sometimes coincident crystallization occurs in finely dispersed morphology EPDM and EPR copolymers are crystalline if block copolymer and depending on the ethylene/propylene ratio; in blends with PE or PP they can extract to some extent (depending on the ethylene/propylene-ratio) amorphous low molecular weight molecules from the PP or PE phase that slightly influence the crystallization and melting behavior

PP

PA-6

a

PET

PPS

Comp. A Comp. B (matrix) (minor)

Table 3.15. Continued.

Crystallization, Morphological Structure, and Melting of Polymer Blends

245

246

G. Groeninckx, M. Vanneste and V. Everaert

These factors influence the crystalline morphology development, resulting in changes of crystallization parameters such as: • nucleation density, N, • spherulite growth rate, G, • overall crystallization rate, K, • total degree of crystallinity, Xc, • semicrystalline morphology, i.e., shape, size and texture of the spherulites, interspherulitic boundaries, etc. To discuss these topics in a systematic way, a distinction will be made between three main blend categories, namely: 1) blends with a crystallizable matrix and an amorphous dispersed phase, 2) blends with an amorphous matrix and a crystallizable dispersed phase, 3) blends containing two crystallizable components. 3.4.3

Blends with a Crystallizable Matrix and an Amorphous Dispersed Phase

In immiscible blends, the phases are separated in the molten state, before crystallization of the matrix starts. The dispersed amorphous phase is assumed to be homogeneously distributed in the melt in droplet-like domains. 3.4.3.1 Nucleation Behavior of the Crystallizable Matrix 3.4.3.1.1 General Considerations Related to Heterogeneous Nucleation When a crystallizable component forms the matrix phase in a polymer blend, nucleation can occur via heterogeneous nucleation by heterogeneities in a similar way as in the pure component. The heterogeneities, available in the melt, can be residual catalysts, fillers, impurities, crystalline residues (due to incomplete melting), etc. Each type of ‘heterogeneity’ has its own typical activation energy for the formation of an ‘active nucleus of critical size’, corresponding to a certain degree of undercooling (T°m - Tc). When Tc,1 is reached during cooling from the melt, all heterogeneities of type 1 (which have the lowest activation energy)

become active and the nucleation of the crystallizable phase is induced. Once the crystallization is initiated by the primary nucleation, it can further spread over the whole available material via secondary nucleation, before any other type of heterogeneity can become active. Since in immiscible blends the phases are physically separated, the same heterogeneities that nucleate the homopolymer at Tc,pure may nucleate the crystallizable matrix. As a result, the crystallization temperature, Tc, of the blend during cooling from the melt will in general not differ that much from the Tc of the pure component. Some general principles governing the crystallization behavior of homopolymers also remain valid for immiscible polymer blends in which the crystallizable component forms the continuous phase. The premelting temperature, Tmelt, may have a profound influence on the crystallization temperature of the matrix, Tc , during cooling from the melt (Table 3.16). The higher the temperature at which the blend is kept in the melt prior to crystallization, the less residual crystalline parts (otherwise leading to self-seeded nucleation) remain in the melt. As a result, fewer nuclei are available to nucleate the melt phase, thus leading to the formation of fewer, but larger spherulites. Another, less important factor, is the isothermal crystallization temperature, Tc,iso, when the crystallization is carried out at a constant temperature (Table 3.16). When a crystallization experiment is performed at lower temperatures, the activation energy for nucleation of several types of heterogeneities can be overcome. At that Tc,iso, more nuclei become active, leading to the formation of a larger number of smaller spherulites. Although most principles for the crystallization of homopolymers remain valid for immiscible blends with a crystallizable matrix, the crystallization behavior can be altered by two phenomena, inherently correlated with immiscible two-phase systems, namely migration of impurities during melt-mixing and the nucleating activity of the interface between two phases.

Crystallization, Morphological Structure, and Melting of Polymer Blends

Table 3.16.

Influence of Tmelt and Tc,iso upon the nucleation behavior in crystalline/amorphous polymer blends

Blend System

Influence of Tmelt

Influence of Tc,iso

PP/EPDM PP/EPR x PP/PIB PP/PS

x

sPS/PVME a Remarks:

a

247

x x x x x x x x x

References [Martuscelli et al., 1983] [Martuscelli, 1985] [Martuscelli et al., 1982] [Martuscelli, 1985] [Bianchi et al., 1985] [Martuscelli, 1985] [Bartczak et al., 1987] [Wenig et al., 1990] [Cimmino et al., 1993]

sPS/PVME is only immiscible in those blends where the amount PVME exceeds 10%

3.4.3.1.2 Migration of Impurities During the Melt-Mixing Process During the melt-mixing process, heterogeneous impurities can migrate across the interface between both blend phases [Bartczak et al., 1986]. The driving force for this migration is the interfacial free energy of the impurity with respect to its melt phase, σi,1. If this interfacial free energy is higher than the interfacial free energy of that impurity within the second melt phase, σi,2, it is energetically more favorable for the impurity to move to the second phase. As soon as it has the ‘possibility’, it will migrate across the interface [Galeski et al., 1984]. Several factors determine the ‘possibility’ for the impurities to migrate from one phase to the other phase during the melt-mixing process. Because the migration of heterogeneities can only occur when they find themselves close enough to the interface, the melt-mixing conditions play an important role [Bartczak et al., 1987] (Figure 3.35). It must be clear that, the longer the mixing or the more intense the mixing, the higher the probability that nuclei find themselves somewhere at an interface, where they can easily migrate. Thus, the effect of migration on the

crystallization behavior will be more pronounced — migration of heterogeneities across the interface will not proceed in the absence of mechanical mixing [Bartczak et al., 1987]. Furthermore, the possibility for impurities to be located close enough to an interface stands in direct relation to the phase morphology generated during the melt-mixing [Bartczak et al., 1987]. As the relative amount of the phases changes, the amount of nuclei that can migrate varies, and the effect of these migrating nuclei on the crystallization behavior changes. This can be understood if one assumes an amorphous/crystalline blend system in which heterogeneities migrate from the crystallizable matrix towards the second phase. With increasing amount of the second phase, the total amount of available nuclei is lower, and they will migrate towards a larger volume of the second phase, what may lead to a more than proportional decrease of the nucleation density in the crystallizable phase. However, the melt-morphology also changes with varying content of the phases. By increasing the amount of the second phase, the dispersion becomes coarser due to coalescence of droplets. This implies that larger droplets are formed, and

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G. Groeninckx, M. Vanneste and V. Everaert

as a consequence a lower total interfacial contact area is available. Hence, less impurities will find themselves located close enough to the interface to be able to migrate. It should be remarked that a critical volume fraction of the second component could exist, which is able to absorb all active nuclei of the crystallizable matrix. Adding higher amounts of the second component will no longer decrease the number of active nuclei per volume unit of the crystallizable matrix. An example is given for the PP/LLDPE blend, where LLDPE is in the molten state during PP crystallization (and thus can be considered as an amorphous melt) (Figure 3.36). Finally, a factor that also may influence the degree of migration is found to be the interfacial free energy between both phases of the blend in the melt, σ1,2. If σ1,2 is high, due to a high degree of immiscibility between the phases, a sharper interface will be formed. Nuclei close to such a sharp interface are found to migrate fast and efficient [Bartczak et al., 1987], in contrary to partially miscible blends where no evidence could be found for such a fast migration [Galeski et al., 1984; Bartczak et al., 1986]. In general, the migration of heterogeneities from one phase to the other in blends with a crystallizable matrix, only slightly affects the crystallization temperature of the matrix during cooling from the melt [Bartczak et al., 1987]. More important should be the influence of migration on the final semicrystalline morphology. This aspect will be discussed in Part 3.4.3.4. 3.4.3.1.3 Nucleating Activity of the Interface The second phenomenon found to influence the crystallization behavior in immiscible polymer blends is the nucleating activity of the interface [Bartczak et al., 1987; Wenig et al., 1990; Wei-Berk, 1993]. In immiscible polymer blends with a high degree of immiscibility such as PP/PS, it has been shown that nucleation at the interface affects the crystallization behavior. Wenig et al. [1990] showed that, with increasing the amount of PS in a blend with PP, the nucleation shifted from preferentially thermal (related to the degree of undercooling) to more athermal. This was

explained by the effect of heterogeneous surface nucleation at PS-interfaces (Figure 3.35). However, not all interfaces can produce additional nucleating centers. For immiscible and highly incompatible polymer blends, since their interfacial tension is higher, the interface is very sharp [Helfand and Tagami, 1972]. Such interfaces can rarely induce new nuclei. Furthermore, on a molecularly smooth surface, a new layer can only be grown after secondary nucleation, a somewhat lower energy barrier is present, since the surface area which must be created is smaller [Hoffmann et al., 1992]. Only an interface which wets well with the crystallizable matrix, so that a crystalline chain can deposit on it, can cause heterogeneous nucleation [Turnbull, 1950; Geil, 1973]. The wetting ability between two melt phases can be calculated from the spreading coefficient F12. An example can be given by the immiscible polymer blend pair PP/PS [Bartczak et al., 1987] (Figure 3.35). Furthermore, the physical state of the second component at the time of matrix nucleation is of importance. It may be presumed that the mode of nucleation of a polymer in the presence of solidified domains of the second polymeric phase is heterogeneous, and therefore the nucleation rate should be higher than in the pure homopolymer. The effect of blending on the nucleation behavior is more subtle and complex in the presence of a molten second component. Factors such as miscibility, relative melt-viscosity, and inherent crystallizability all influence the formation of critical size nuclei [Nadkarni and Jog, 1991]. Nucleation by the interfaces contributes to the crystallization behavior proportionally to the total amount of interface in the blend system. The finer the amorphous droplets are dispersed, the larger the total interfacial contact surface, and thus the higher is the possibility of nucleation at these interfaces. The main factors determining the melt morphology are the blend composition, the difference in melt-viscosity between both phases, and the interfacial tension. Hence, the nucleation effect on the crystallization behavior should be more pronounced in blends containing a higher amount

Crystallization, Morphological Structure, and Melting of Polymer Blends

of the dispersed phase, or in blends composed of components with nearly equal melt-viscosities. It has to be noticed that, due to coalescence, upon increasing the amount of the amorphous component, larger domains are formed. As a

249

result, the total interfacial contact area may not increase proportionally, leading to a less-thanlinear increase of Tc with increasing amount of the amorphous component (see Figure 3.34 and Figure 3.35).

Figure 3.34. Influence of the amount of dispersed phase, mixing time, Tmelt and Tc,iso on the amount of nuclei per volume unit in the immiscible PP/PS blend a) after 2x mixing, b) after 3x mixing; Tmelt was set to 190°C (open symbols) or 220°C (filled symbols); Tc,iso for experiments was 125°C (circles) or 130°C (blocks) [Bartczak et al., 1987].

Figure 3.35. Influence of the amount of dispersed phase, mixing time and crystallization temperature, Tc, on the amount primary nuclei active for crystallization at Tc in a PP/PS blend. All samples have been molten up at 220°C; a) 2x mixing, b) 3x mixing; Tc,iso was set to 119°C (∇), 123°C (¨), 125°C (o) and 130°C (†) [Bartczak et al., 1987].

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G. Groeninckx, M. Vanneste and V. Everaert

Figure 3.36. Influence of LLDPE on the nucleation of PP at various temperatures: (a) measured as blend volume; and (b) calculated as PP volume fraction [Long et al., 1991].

An interesting application of the direct relationship between nucleating interfaces and the total amount of the interfacial contact surface can be found in compatibilized immiscible blends. In these systems, the dispersed phase size becomes much smaller, strongly increasing the total amount of interface at which nucleation can occur. Some authors reported that this could cause an upward shift in the Tc by up to 10°C [Wei-Berk, 1993]. However, other studies in which the crystallization behavior of a compatibilized blend was investigated did not always mention such a clear nucleating activity (Table 3.17). Finally, the degree of nucleation at the amorphous/ semicrystalline interfaces was found to be temperature dependent. When the crystallization temperature was raised, the nucleating efficiency of the interface was found to decrease [Bartczak et al., 1987]. In conclusion, the polymer interface can induce some limited number of nucleation events, but does not cause transcrystallinity, as some other crystal surfaces do. Consequently, the amorphous droplet surfaces, either in the solid or molten state, only act as a weak nucleating agent [Bartczak et al., 1987]. 3.4.3.1.4 Nucleation Behavior of Some Selected Polymer Blends (see Table 3.18) 3.4.3.2 Spherulite Growth of the Crystallizable Matrix

For homopolymers, the temperature dependence of the isothermal spherulite growth rate, G, is described by Eq 3.1 [Turnbull and Fischer, 1949]: G1 = G° exp[-¨E/kTc] exp[-¨F*/kTc]

(3.1)

In the case of immiscible blends with a crystallizable matrix, the spherulite growth can be disturbed to a certain degree by the presence of an amorphous phase component, dispersed in the crystallizable melt. 3.4.3.2.1 Phenomena Affecting the Spherulite Growth Rate; Energetic Considerations Prior to crystallization, the amorphous component exists in the form of droplet-like domains, which at Tc can either be in the molten or glassy state. During the spherulite growth of the crystallizable matrix, small domains may be rejected by the spherulitic growth front either completely to the amorphous interspherulitic zone, or only partly over some distance. Furthermore, somewhat larger domains can be occluded by the growing stacks of lamellae after which they eventually can be deformed [Martuscelli, 1984; Bartczak et al., 1984]. In most cases, a combination of the above described processes is observed; small droplets are rejected over some distance, coagulate at the growth front and are engulfed and/or deformed subsequently by the growing lamellar stacks. The presence of droplet-like domains along the path of the crystallizing growth front can markedly

PP/EPR (melt-mix)

PP/EPR (solvent-mix + melt-mix)

PP/EPR (melt-mix)

HDPE/PS (melt-mix)

90/10 80/20 (70/30) 95/5 90/10 80/20 70/30

- Formation of more α-type spherulites with [EPR] ↑ (WAXS, O. M.) - Tc ↑ with [EPR] ↑ (DSC) N/S in PP phase ↑ with [EPR] ↑, mainly at higher [EPR] ⇒ spherulite size ↓ (O. M.) - N/S ↑ with [EPR] ↑ (5 to 7x) - this effect is more pronounced when the% ethylene ↑ or the MW ↓ of the used EPR (O. M.)

EPR act as a nucleating agent, probably due to migration of nuclei from EPR to PP phase

[Martuscelli et al., 1982], [Coppola et al., 1987], [Kalfoglou, 1985] [Martuscelli, 1985]

[Karger-Kocsis et al., 1979]

- Spherulite size ↓ to less than 1/2 size of PP pure with [EPR]↑ (SALS, O. M.)

[Baïtoul et al., 1981]

[Aref-Azar et al., 1980] EPR is nucleating agent for the production of α-type spherulites (most stable and dense packed)

Explanation

No change

No change

LDPE/PS (melt-mix)

90/10 80/20 70/30 60/40 90/10 80/20 95/5 90/10 80/20 60/40

Observed nucleation behavior (technique used)

Blend system (preparation) Comp. a

Table 3.18. Examples of the nucleation behavior of the crystallizable matrix in crystalline/amorphous polymer blends Reference

[Wei-Berk, 1993]

Copolymer lowers the interfacial tension → finer dispersion → more surface available for nucleation at interface + more formation of α-phase PP crystals as Tc rises above 125°C

PP/PS + PP/PS block

[Martuscelli, 1984]

Strong nucleation effect of EPR-g-SA on the PP phase

Compatibilization ↓ spherulite size (which was not found in PA-6/EPR blends) + serious ↑ of the interfacial adhesion Tc increased (116 → 126°C) (DSC) along with copolymer content up to 20-25% of PS phase

PA-6/EPR+ g-SA

Reference

Explanation

Observations

Influence of compatibilizers on the nucleation behavior of the semicrystalline matrix in crystalline/amorphous polymer blends

Blend system

Table 3.17.

Crystallization, Morphological Structure, and Melting of Polymer Blends

251

90/10 80/20 70/30 60/40 50/50

PP/PS (solvent-mix)

Remarks:

b

a

[Santana and Müller, 1994]

[Cimmino et al., 1993]

- Formation of larger spherulites with Migration of impurities during mixing from [PVME] ↑ sPS to PVME phase (O. M.) - t0.5 ↑ during isothermal crystallization (DSC)

[Wei-Berk, 1993]

[Wenig et al., 1990]

[Bartczak et al., 1987]

[Karger-Kocsis, 1979]

[Martuscelli et al., 1983], [Martuscelli, 1985] [Ghijsels et al., 1982]

[Bianchi et al., 1985], [Martuscelli, 1985]

[Martuscelli et al., 1983]

Reference

No significant change in Tc visible

Nucleating activity of the PS interface on PP

- mainly migration of nuclei from PS phase to PP phase - some limited nucleating activity of PS droplets at interface because of their high interfacial tension - nucleating activity of PS at the interface - when [PS] > 20%, coalescence of droplets occurs ⇒ less ↑ interfacial area - nucleation changes from thermal to athermal due to nucleation at interfaces

SBS is a weak nucleating agent for PP

SBS is a weak nucleating agent for PP

PIB seems to be an effective nucleating agent

No additional nucleation effects

Explanation

Other compositions than the ones mentioned have been investigated sometimes PVME and sPS are immiscible as long as the concentration PVME remains larger than 10% of the blend

96/4 77/23 65/35 60/40 PP/PS (melt-mix) 90/10 80/20 70/30 50/50 80/20 sPS/PVME b (solvent-mix and molten up) 70/30 50/50

PP/PS (melt-mix)

PP/PS (melt-mix)

PP/SBS (melt-mix)

75/25 50/50 95/5 90/10 90/10 80/20 70/30

PP/SBS (melt-mixed)

N/S in PP phase ↑ with [EPDM] ↑, mainly at higher [EPDM] ⇒ spherulite size ↓ (O. M.) N/S is not affected by addition of PIB (O. M., thick sheet cross-section)

PP/EPDM (solvent cast film) 90/10 80/20 70/30 PP/PIB (melt-mix) 90/10 80/20 70/30 PP/PIB (solvent-mix) 90/10 N/S seriously ↑↑ with addition of PIB up to 30 times N/S in PP pure (O. M., thin film) Tc ↑ slightly (± 2°C) with addition of SBS (DSC) Spherulite radius ↓ with addition of SBS (O. M., SALS) - N/V ↑ with [PS]↑ (up to 7x N/V found in pure PP) - This effect is more pronounced with longer mixing times (O. M.) - N/S ↑ due to addition of PS up to 20% PS (4x) - this increase becomes less pronounced with further [PS] ↑ (O. M.) - Avrami exponent n ↓ from 3 to 2 (DSC) Tc ↑ from 121° (4%PS) to 127°C (40% PS) (DSC)

Observed nucleation behavior (technique used)

Blend system (preparation) Comp. a

Table 3.18. Continued.

252 G. Groeninckx, M. Vanneste and V. Everaert

Crystallization, Morphological Structure, and Melting of Polymer Blends

253

disturb the spherulite growth. The outlined processes require the growth front to perform work against the interfaces, thus dissipating energy. Such energies constitute new energy barriers, controlling the spherulite growth in immiscible blends. The spherulite growth rate depression is proportional to the type of energy barrier that has to be overcome, and can be quantitatively expressed by a modified equation of the spherulite growth rate [Martuscelli, 1984]:

oped by Cissé and Bolling [1971] and by Omenyi et al. [1981]. Bartczak et al. [1984] have modified these theories in order to apply them to the case of a crystalline polymeric front that grows according to a spherulite-like morphology, while in the melt non-crystallizable polymeric domains of spherical shape are present (Table 3.19). The driving force for rejection, occlusion, or deformation processes is equal to the difference of interfacial free energies [Martuscelli, 1984]:

G = G1 exp[-(E1 +E2 +E3 +E4)/kTc]

¨F = γPS - γPL

(3.43)

where G1 is the spherulite growth rate of the plain crystallizable polymer (theoretically described by the Turnbull-Fisher equation); E1 is the energy dissipated for rejection (proportional to the melt viscosity); E2 is the energy needed to overcome the inertia of the drops; E3 is the energy needed to form a new interface if drops are engulfed; and E4 is the energy dissipated for deformation of occluded particles. Theories for the description of these energies for a non-polymeric solidification front were devel-

(3.44)

where γPS is the interfacial free energy between crystallizing solid and the inclusions, and γPL is the interfacial free energy between the melt and the inclusions. When ¨F is positive, the particle droplet will be rejected [Wei-Berk, 1993]. Martuscelli [1984] and Bartczak et al. [1984] have calculated the energies dissipated by growing PP spherulites in a blend with dispersed rubber particles for all the above mentioned phenomena that may disturb the spherulite growth (Table 3.20). It can be concluded that mainly rejection of small

Table 3.19. Expressions for the dissipation energy terms and corresponding spherulite growth rates in a crystalline/amorphous polymer blend system [Martuscelli, 1984; Bartczak et al., 1984] Rejection of droplets by growing spherulites E1 = 1.5 (EGRsµMc/ρMr2) G = G1/[1+(3µMcEG1Rs/2ρMr2RT)] E2 = CµMρPG2/2ρM G = G1 exp[-CµMρPG2/2ρMRT]

Occlusion of droplets in growing lamellae

Deformation of occluded droplets *

E3 = 3CµM∆F/ρMr G = G1 exp[-3CµM∆F/ρMrRT]

E4 = U(K) [3CµMγPS/ρMr] G = G1 exp[-U(K).(3CµMγPS/ρMrRT)]

G1is the undisturbed spherulite growth rate µM is the molecular mass of the repetitive unit of the macromolecular chain of the crystallizable matrix γPS is the interfacial free energy between the crystallizing solid and the inclusions ρM and ρP is the density of the matrix and of the dispersed component Rs and r is the radius of spherulite and of the dispersed particles, respectively c is the volume concentration of the non-crystallizable component R is the gas constant E is the kinetic energy supply required to move the dispersed droplet along with the motion of the crystallizing front = 2/3 G2 Πr2 ρP * The energy of deformation is the sum of two terms: the first is related to change of the surface of particles and the second to deformation of viscoelastic material. U(K) is a complicated function of the coefficient of deformation K of the particles. In the expression of E4 only the change in surface is taken into consideration with reference to the case where ∆F > 0.

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G. Groeninckx, M. Vanneste and V. Everaert

particles, and to a lesser extend deformation of large engulfed droplets (requiring the formation of new surface boundaries) cause a depression in the spherulite growth rate. Table 3.20. Energy dissipated in PP/TR blends for rejection, occlusion and deformation of the TR droplets by the growing PP lamellae [Martuscelli, 1984; Bartczak et al., 1984]. Process

Energy (J/mole PP repeating units)

Rejection Kinetic energy of rejection Occlusion Deformation: surface change term Deformation: viscous term

101 - 104 10-15 - 10-14 10-2 - 10-1 10-1 - 100 10-7 - 10-6

3.4.3.2.2

Factors Influencing the Spherulite Growth Rate, G Several factors determine the amount of energy required by the growth front to be overcome in order to allow the crystallizable matrix to form spherulites The first and most important of these is the crystallization temperature, Tc. The higher the isothermal crystallization temperature above Tc,max, the slower the spherulites will grow. However, higher Tc also implies a lower melt-viscosity. In such case, small droplets will be rejected easier, consuming less energy. This is reflected in a spherulite growth rate, nearly independent on the total amount of small amorphous droplets to be rejected, while at lower Tc it could be clearly seen that the growth rate is much affected by the amount of fine droplets (Figure 3.37) The temperature dependence of the spherulitic growth rate has been theoretically treated [Wenig et al., 1990], for several blends composed of a PP matrix in which PS droplets were dispersed. This temperature dependence could be calculated based on the work done by Hoffmann [1983] and by Suzuki and Kovacs [1970], and is defined as follows (Figure 3.38a):

Figure 3.37. Theoretical estimation for spherulite growth rate depression in immiscible PP-based blends in the case of REJECTION of particles: influence of particle size, Tc,iso and volume concentration of the second component [Martuscelli, 1984].

for T < Tg - C2: G(T) = 0 for Tg-C2 < T < Tm°: G(T) = G° exp [-C1C2 / (C2 + T - Tg)] exp[(-C3) / T(Tm° - T)] (3.45) for T > Tm°: G(T) = 0 where Tg is the glass transition temperature of the crystallizable component; Tm° is the theoretical melting temperature of the crystalline component; G°, C1, C2 and C3 are parameters describing the growth rate behavior in the blends. For the crystalline component, the parameters from the WLF equation, C1 and C2 can be found from literature [Icenogle, 1985]. Tg and Tm° can be measured for pure crystallizable component. The parameters G° and C3 can be calculated

Crystallization, Morphological Structure, and Melting of Polymer Blends

from the experiments that give the spherulite growth rate G as a function of temperature T. By plotting the value ln G + C1C2/(C2+T-Tg) versus 1/[T(Tm°-T)] for the entire crystallization temperature range, a linear plot is obtained from which the values of G° (intercept) and C3 (slope) for all the blend compositions can be determined (Figure 3.38b). Once all these parameters are known, the growth rate can be estimated as a function of temperature for all blend compositions, according to Eq 3.45. Secondly, the blend composition is of importance as well. The finer the dispersion (i.e., at low content of the amorphous phase, nearby equal melt-viscosities of matrix and dispersed phase,

255

etc.) more droplets need to be rejected. This high energy consuming process reduces the spherulite growth rate (see Figure 3.37). 3.4.3.2.3 Spherulite Growth Rate Investigations in Some Typical Polymer Blends (see Table 3.21) 3.4.3.3 Overall Crystallization Kinetics The effect of blending on the overall crystallization rate is the net combined effect of the nucleation and spherulite growth. Martuscelli [1984] observed that in blends of PP with LDPE, crystallized at a Tc high enough to prevent any LDPE crystallization, the overall rate of crystallization of the PP matrix

Figure 3.38. a) Temperature dependence of the spherulite growth rate, G, for PP (experimental values were fitted using the function defined in Eq 3.45); b) Plot to determine the parameters G° and C3 (C1 = 25, C2 = 30K, Tg = 260K, Tm° = 460.5K) [Wenig et al., 1990].

Figure 3.39. Global crystallization kinetics in immiscible PP/LDPE blends; influence of the amount dispersed phase and the crystallization temperature, Tc, on the half time for crystallization, t1/2 [Martuscelli, 1984].

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G. Groeninckx, M. Vanneste and V. Everaert

phase (thus in the presence of the LDPE molten droplets) was progressively depressed with increasing content of LDPE (Figure 3.39). This can be seen in the plot of t1/2 (half-time of crystallization at a fixed Tc,iso ) versus blend composition. The observations agree very well with the findings that the growth rate of the PP spherulites is almost unaffected while the nuclei density decreases with increasing LDPE content due to impurity migration from PP to the LDPE phase. A different case has also been explored by Martuscelli [1984] for PA-6 blended with an EPR-rubber. As shown in Figure 3.40, t1/2 of the PA-6/EPR blend decreased (faster overall crystallization rate) as the content of the rubbery phase increased, especially at lower concentrations of the EPR phase. The reverse could be observed in a compatibilized blend. Because in these blends a serious decrease of the spherulite size was observed, the authors concluded that the compatibilizer

Figure 3.40. Variation of the half time of crystallization, t1/2, with the percent of added rubber component (EPR) and Tc for PA-6/Dutral and the compatibilized blend PA-6/Dutral-g-SA [Martuscelli, 1984].

acted as a nucleating agent for the PP phase. However, due to the increase of the melt-viscosity upon compatibilization, the overall crystallization kinetics was retarded. Additionally, they observed experimentally that ¨F*(free energy for the formation of a nucleus of critical size) and σe (surface free energy of folding) in compatibilized blends were larger than in PA-6 homopolymer. An opposite trend was observed for the physical PA-6/EPR blends. No further investigations have been done to elucidate this phenomenon. 3.4.3.4 Final Semicrystalline Morphology The addition of a second non-crystallizable component to a crystallizable matrix can cause drastic variations of important morphological and structural parameters of the semicrystalline phase, such as the shape, size, regularity of spherulites and interspherulitic boundary regions, lateral dimensions of the lamellae, etc. These factors may greatly influence the mechanical behavior and, in particular, the fracture mechanisms, and thus are of great importance, especially when the toughening of semicrystalline polymer blends is considered. The first important parameter determining the final crystalline morphology is the nucleation density, N, (see Part 3.4.3.1). An increase in the nucleation density (per volume unit of the crystallizable material) due to migration of nuclei from one phase towards the other, or due to a nucleating activity at the polymer/polymer interface, results in the formation of more numerous, but smaller spherulites. The spherulite growth rate, G, also plays a role. • At low G values, there is a higher probability that all dispersed particles can diffuse fast enough away from the growth front, and be pushed along until complete crystallization. The second phase component will then be found mainly in the interspherulitic regions. • At high G values, even small particles will not be rejected anymore. Hence, the homogeneously distributed droplets will be as such engulfed, rejected into newly formed boundaries behind occluded particles, and eventually deformed.

100/0 90/10 80/20

100/0 90/10 80/20 70/30

100/0 90/10 80/20 70/30 60/40

100/0 80/20 70/30 50/50

PP/PIB (thin film)

PP/PS

PP/PS

sPS/PVME

c

b

a

32 23 20 at 121°C

PP/EPDM (thin film) 100/0 90/10 80/20 60/40

Remarks:

15 11 12 13 at 125°C

100/0 90/10 80/20 70/30

PP/EPR(thin film)

x

x

x

x

x

x

Tc c

Droplets are occluded intra-spherulitical

PS droplets are occluded intraspherulitical. (no rejection or deformation)

- Most droplets (3-5µm) are occluded intraspherulitic - Sometimes droplets are rejected (for short distance or even interspherul.) . if: MW PIB low . [PIB] low (< 10%)

[Wenig et al., 1990]

- G is only slightly ↓ except for 90/10 composition (no explanation) - globally G is independent on composition

- G is ± independent on [Cimmino et al., 1991] composition - G↑ with [PVME] ↑ due to [Cimmino et al., 1993] facilitated transport of chains (PVME and PS chains are partially miscible ⇒ Tg↓)

[Bartczak et al., 1987]

G is not changed and is independent on blend composition

Clear depression of G on [Martuscelli et al., 1982] addition of PIB which is [Martuscelli et al., 1983] strongly dependent on MW PIB [Martuscelli, 1985] and compositionb

[Martuscelli et al., 1983] [Martuscelli, 1985]

- Droplets are occluded intraspherulitic. Only small ↓ of G independent - Droplets are aligned along radial direction of the blend composition due to coalescence of droplets upon rejection for short distance, followed by deformation

[Martuscelli et al., 1982]

Small ↓ of G which is composition dependent [Martuscelli, 1985]

Reference

Comments

- At low [EPR]: droplets are ejected for a short distance and then occluded intra-spherulitic.

- Some droplets are ejected interspherulitic.

Final semicrystalline morphology

Some of the listed data were extracted from plots presented in the article PIB is shown to be partially miscible with the amorphous phase of PP at low concentrations. Hence, the behavior can deviate and exhibit some typical characteristics as in miscible systems x indicates that the influence of different crystallization temperatures Tc has been investigated

2.5 6.0 7.0 --at 244°C

3.2 1.6 3.1 2.8 2.2 at 133°C

2.5 2.3 2.5 2.5 at 133°C

32 ±30 ±27 ±27 at 121°C

Comp.

Blend system

G(µm/min)a (from O. M.)

Table 3.21. Examples of the spherulite growth rate, G, of the crystallizable matrix in crystalline/amorphous polymer blends

Crystallization, Morphological Structure, and Melting of Polymer Blends

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This results in a radial-like distribution of the droplets within the spherulite (Figure 3.41). • At intermediate growth rates, the dispersed drops will first be pushed along, but due to an increase of the amount of droplets at the solidification front, they will coagulate and subsequently be engulfed. This results in a spherulite center consisting of pure crystalline material, and an outer layer in which dispersed particles are occluded. Another parameter strongly influencing the final crystalline morphology is the blend composition (Figure 3.42).

The higher the amount of the amorphous phase, the higher the chance to have a coarse melt-morphology containing lots of large, easily coalescing amorphous droplets. In such a case, the crystallizing growth front will mainly engulf and deform these droplets. The resulting crystalline morphology will be heavily disturbed by the second phase component. Some examples of the final semicrystalline morphology in several immiscible crystalline/ amorphous blend systems have already been given in Table 3.21 for the discussion of the spherulite growth rate (Part 3.4.3.2). Some more information about this topic can be found in the articles listed in Table 3.23. 3.4.3.5 Melting Behavior of the Crystalline Matrix in Crystalline/Amorphous Blends

Figure 3.42. Optical micrographs of isothermally (Tc = 135°C) crystallized thin films of PP/PIB(HM) blends with different compositions [Martuscelli et al., 1983].

The behavior of binary blends with only one crystallizable component has been studied by several authors, who have investigated different systems. The crystals of the crystallizable matrix have grown in equilibrium with their own melt phase. The presence of separate domains of noncrystallizable component, dispersed in the molten matrix during the crystallization process, (owing to the kinetic and morphological effects), may cause a depression of the observed melting tem-

Figure 3.41. Optical micrographs of melt crystallized films of PP/EPDM blends at Tc = 135°C; a) 90/10, crossed polarizers; b) 90/10, parallel polarizers; c) 60/40, crossed polarizers [Martuscelli et al., 1983].

Crystallization, Morphological Structure, and Melting of Polymer Blends

259

Table 3.22. Global crystallization kinetics of the crystallizable matrix in some crystalline/amorphous blend systems Blend system

Comp.

Parameter

HDPE/PS

100/0 90/10 80/20 100/0 90/10 80/20 60/40

Avrami exponent (DSC) t0.5 (DSC) Xc,iso (DSC)

100/0 90/10 80/20 70/30 100/0 90/10 80/20 70/30 60/40

Xc,iso (WAXS, DSC)

PP/EPR

PP/PIB

PP/PS

PA-6/EPR

sPS/PVME

a

100/0 90/10 80/20 70/30 100/0 80/20

Tc a

Comments

References

x

Unaffected cryst. kinetics (insensitive to blend morphology)

[Aref-Azar et al., 1980]

-

At [EPR] < 20%: slight ↓ of Xc due to limited miscibility of aPP and EPR → hindered crystal growth - At [EPR] > 20%: Xc ↑ with [EPR] ↑ due to nucleating activity of EPR Xc ↓ with [PIB] ↑

[Kalfoglou, 1985]

Avrami exp., n (DSC)

t0.5 (DSC)

x

t0.5 (DSC)

x

n ↓ from 3 to 2 with [PS] ↑ (due to surface nucleation at PS droplets), and G = cte ⇒ crystallization rate is enhanced and strongly dependent on blend composition Serious ↓ in t0.5, which is most pronounced at low conc. EPR ⇒ enhanced crystallization kinetics Seriously retarded kinetics of sPS phase (t0.5 ↑) which is composition dependent effect of N/S ↓ is larger than that of G ↑

[Bianchi et al., 1985] [Martuscelli, 1985]

[Wenig et al., 1990]

[Martuscelli, 1984]

[Cimmino et al., 1993]

x indicates that the influence of different Tc on the overall crystallization kinetics have been investigated in the article mentioned

perature, Tm’ [Martuscelli, 1984]. However, the changes in Tm’ will be only in the range of a few degrees C. Some binary systems do not show any depression at all, indicating that Tm’ and Tm do not depend on blend composition. This is found when the second dispersed phase does not influence the normal crystallization behavior of the matrix polymer: no nucleating activity, no influence on spherulite growth rate, etc. Some examples of the melting behavior in previously discussed blend systems are given in Table 3.24.

3.4.4

Blends with a Crystallizable Dispersed Phase in an Amorphous Matrix

In immiscible polymer blends, the minor component often forms the dispersed phase, whose shape and size are complex functions of the blend composition, the melt-viscosity of the dispersed phase and the matrix, the viscosity ratio, the interfacial tension, and the processing conditions [Utracki, 1989; Folkes and Hope, 1993]. The crystallization behavior of a dispersed meltphase, for example discrete melt droplets, in an amorphous matrix can be dramatically affected

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Table 3.23. Overview of literature in which the final semicrystalline morphology in immiscible crystalline/amorphous polymer blends has been studied Blend System

Reference

PP/EPR

[Martuscelli et al., 1982] [Coppola et al., 1987] [Kalfoglou, 1985] [Karger-Kocsis et al., 1979] [Martuscelli, 1985] [Martuscelli et al., 1983] [Martuscelli, 1985] [Martuscelli et al., 1982] [Martuscelli et al., 1983] [Bianchi et al., 1985] [Martuscelli, 1985] [Bartczak et al., 1987] [Lotz and Kovacs, 1969] [Cimmino et al., 1991] [Cimmino et al., 1993]

PP/EPDM PP/PIB

PP/PS PEG/PS sPS/PVME

a b c

Composition a

x x x x x x x x x x x x x x x

Growth Rate b (rejection ↔ occlusion)

Nucleation Density c (spherulite size) x x x x

x

x

x x x x

x x x x x x x

Influence of compositional variations on the semicrystalline morphology has been investigated Influence of different spherulite growth rates on semicrystalline morphology is discussed Final spherulite size has been evaluated

compared to that of the bulk polymer. It has been reported by several authors that crystallizable dispersed droplets can exhibit the phenomenon of “fractionated crystallization” originating from the primary nucleation of isolated melt particles by species with different nucleating activity (heterogeneities, local chain ordering) [Aref-Azar et al., 1980; Baïltoul et al., 1981; Ghijsels et al., 1982; Robitaille and Prud’homme, 1983; Frensch et al., 1989; Santana and Müller, 1994; Müller et al., 1995, Morales et al., 1995]. 3.4.4.1 The Phenomenon of Fractionated Crystallization of a Dispersed Phase Crystallization is a phase transition that is controlled by nucleation and growth [Wunderlich, 1976]. As it has been outlined in Part 3.4.3.1.1, crystallization during cooling from the melt in homopolymers is initiated by impurities (primary

heterogeneous nucleation), after which the crystallizing front spreads over the whole material via the secondary nucleation, before other heterogeneities, requiring a larger degree of undercooling, ¨Tc,i = Tmo - Tc,i, can become active. A single crystallization exotherm is generally observed in DSC-thermograms. So, the primary nucleation is the rate determining step of crystallization. The dynamics of the process depend for a given component only on the temperature. However, for polymer blends in which the crystallizable phase is dispersed into fine droplets in the matrix, crystallization upon cooling from the melt can sometimes occur in several steps (fractionated crystallization) that are initiated at different undercooling, often ending up with a crystallization at the homogeneous crystallization temperature Tc,hom [Aref-Azar et al., 1980; Baïltoul et al., 1981; Ghijsels et al., 1982; Santana and Müller, 1994].

Crystallization, Morphological Structure, and Melting of Polymer Blends

261

Table 3.24. Examples of the melting behavior of the crystallizable matrix in some crystalline/amorphous blend systems. Blend System

Tm ↓

Tm ↑

References

Tm ↑ with [EPR] ↑ [Martuscelli et al., 1982] effect is largest at high ethylene content [Martuscelli, 1984] and low MW of EPR - EPR is able to extract selectively defective PP chains ⇒ Tm ↑ and Xc ↑

PP/EPR

-

PP/EPR

Tm ↓ with [EPR] ↑ (2 à 3°C) due to Tc ↓ ⇒ formation of smaller lamellae which melt at slightly lower Tm

[Kalfoglou, 1985]

PP/EPR

Tm ↓ with [EPR] ↑ (3 à 4°C) due to the formation of smaller spherulites with lower Cp (although Tc ↑) Tm ↓, min. is observed at 20% EPDM EPDM is able to dissolve aPP at low conc. EPDM (§ diluent-effect) Tm ↓ with [PIB] ↑ This effect is largest at high conc. PIBLM or low conc. PIBHM PIB dissolves aPP in low amounts (§ dilution effect) Tm ↓ with [PS] ↑ (2 à 3°C) due to some specific interactions at the interface

[Karger-Kocsis et al., 1979]

PP/EPDM

PP/PIB

PP/PS

POM/PC PEG/PS

PET/PPE

[Martuscelli et al., 1983]

[Martuscelli et al., 1983]

[Mucha, 1986] [Wei-Berk, 1993]

Tm ↓ with [PC] ↑ (2 à 3°C) ∆Hm = cte (per g POM) Tm ↓ with [PS] ↑ (max. 8°C) due to formation of smaller lamellae if the [PS] in the block-copolymer ↑

[Chang et al., 1991] [O’Malley et al., 1969]

Tm ↑ upon addition of PPE to a PET matrix (§ ±2°C)

The first investigations concerning the crystallization in discrete droplets date from 1880; Van Riemsdyk reported that small gold melt droplets

[Liang and Pan, 1994]

solidify at much larger undercoolings than the bulk material [Van Riemsdyk, 1880]. Similar observations were made later for other metals

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Figure 3.43.

DSC cooling curves (10°C/min) for PP/PS blends; difference in the crystallization behavior in blends with PP as a matrix phase and as a dispersed phase [Santana and Müller, 1994].

[Perepezko and Paik, 1982], indicating this to be a basic crystallization phenomenon. The creation of sufficiently small polymer droplets as a stable suspension was much more difficult. It was therefore only first in 1959 that similar experiments have been reported for polymers [Frensch et al., 1989] (Table 3.25).

It was clearly observed that the phenomenon of delayed crystallization was directly related to the size of the dispersed droplets [Koutsky et al., 1967]. Only the smallest droplets showed crystallization at much larger undercooling; droplets having a sufficiently large diameter crystallized at temperatures approaching the bulk crystallization

Table 3.25. Crystallization behavior in finely dispersed crystalline droplets Polymer Dispersion Method PE PP PEG PE PP PEG POM iPS PA-6 a b

thermodyn. inert liquid thermodyn. inert liquid thermodyn. inert liquid suspended in silicon oil and sprayed on slides * * * * * *

Average Droplet Size (µm) ∆Tc,bulk (°C) some µm some µm 5 1-2

± 20 ± 50 ± 20 ± 20

1-2 1-2 1-2 1-2 1-2

± 50 ± 20 ± 30 ± ± 15

∆Tc, droplets (°C)

Reference

± 55 a ± 102 a ± 65 b 55 a

[Cormia et al., 1962] [Burns and Turnbull, 1966] [Cormia et al., 1962] [Koutsky et al., 1967]

100 a 65 b 84 a or b 102 b 100 b

[Koutsky et al., 1967] [Koutsky et al., 1967] [Koutsky et al., 1967] [Koutsky et al., 1967] [Koutsky et al., 1967]

Crystallization by homogeneous nucleation at Tc,hom Nucleating activity of the suspending medium prevents to detect the real undercooling needed to obtain a homogeneous crystallization

Crystallization, Morphological Structure, and Melting of Polymer Blends

temperature, Tc. The explanation for this behavior is obvious: the spectrum of undercoolings at which several crystallization steps occur reflects the difference in nucleating activity of the various heterogeneities available in the melt [Frensch et al., 1989]. It can be assumed that, if the dispersion of the polymer is so fine that not every droplet contains at least one heterogeneity of type 1, only those droplets containing the latter will crystallize at an undercooling ¨Tc,1. Since the droplets are physically not in contact with each other, further growth via secondary nucleation in other crystallizable droplets is impossible. During further cooling, heterogeneities of type 2 requiring the second lowest degree of undercooling, ¨Tc,2, can become active in some of the remaining droplets, resulting in a second crystallization exotherm. This process goes on until finally some very fine droplets that have not yet been nucleated by the heterogeneous species, will crystallize in a homogeneous mode. In isothermal experiments, the fractionated crystallization of finely dispersed crystallizable droplets is reflected by longer crystallization times before the same degree of crystallinity Xc is obtained. This has been illustrated clearly by Koutsky et al. [1967] in experiments (see Table 3.25) in which finely dispersed droplets of PE and PP in a suspension of silicon oil were crystallized at different undercoolings ¨Tc (Figure 3.44). It should be mentioned that the occurrence of a fractionated crystallization is related only to the

Figure 3.44.

263

number densities of dispersed polymer particles and primary heterogeneous nuclei. No direct physical relationship has been found with the number, or size of spherulites. These parameters are additionally influenced by the cooling rate and the crystallization temperature [Frensch et al., 1989]. Several factors can influence the fractionated crystallization behavior. An important parameter that has already been discussed in Part 3.4.2.1. is the thermal history of the sample. Crystallizable dispersed droplets that were submitted to premelting at higher temperatures or longer times generally display a shift in the heterogeneous nucleation spectrum to greater undercooling. The homogeneous crystallization temperature however is not displaced and thus independent of the thermal history [Koutsky et al., 1967]. This may become less evident for blends with unstable phase morphology (rapid phase coarsening upon annealing); long residence times in the melt will cause fine droplets to coarsen. Consequently, the newly formed larger droplets have a higher probability to crystallize close to the bulk crystallization temperature of the homopolymer. The degree of dispersion of the minor phase plays a crucial role. Important factors here are the blend composition, the interfacial tension between both components, the melt-viscosity of both components, the processing device and mixing conditions, the blend preparation method, etc. In this context, it is interesting to evaluate also the influence of compatibilization on the

Isothermal homogeneous crystallization of finely dispersed polymer droplets as a function of time (a) Linear PE, (b) PP [Koutsky et al., 1967].

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G. Groeninckx, M. Vanneste and V. Everaert

crystallization behavior of the dispersed phase. Since compatibilization reduces the droplet size of the minor phase even more drastically, it can be expected that this can lead to a serious shift of the crystallization temperature towards lower temperatures, resulting in more pronounced fractionated crystallization, or even in a homogeneous crystallization. However, this issue is more complex due to numerous other factors involved in the nucleation process. Some examples from the literature are listed in Table 3.26. They illustrate how differently the compatibilization can influence the crystallization behavior of the dispersed phase. 3.4.4.2 Theoretical Considerations of the Fractionated Crystallization In crystallizable dispersed droplets, several different nucleating heterogeneities (type i) can be

present, each having a typical free energy for the formation of a nucleus of critical size, ¨F*, at an undercooling ¨Tc,i: ¨F* § ¨ypn/(¨Tc,i)2

(3.46)

This free energy is proportional to the specific interfacial free energy difference ¨ypn defined as [Wunderlich, 1976]: ¨ypn = yp(m,c) - ypn(m) + ypn(c)

(3.47)

where the indices refer to polymer (p), melt (m), crystal (c) and nucleus (n); and ypn(m) is the interfacial energy between the nucleating species and the polymer melt; ypn(c) is the interfacial energy between the nucleating species and the polymer crystal; and yp(m,c) is the lateral surface free energy between the crystal and its own melt. In the case of a homogeneous nucleation, the expression for ¨ypn can be simplified to read: ¨ypn = 2 yp(m,c)

Table 3.26. Influence of compatibilization on the crystallization behavior of the dispersed phase in amorphous /crystalline polymer blends Blend system

Compositions

PP/PS +SBS

18/80/2

Matrix at Tc melt

Comments SBS did not ↓ particle size (bad compatibilizer because immisc. with PP)

References [Santana and Müller, 1994]

Nucleation density ↑ because homogeneous nucleation process becomes more heterogeneous (higher Tc) SBS transfers heterogeneities to PP

9/90/1

LDPE/PS + Kraton G

15/77.8/7.2

solid

Kraton enhances the formation of a finer dispersed PE phase Shift of multiple crystallization to lower temp.

[Bailtoul et al., 1981]

PET/PS +PET-b-PS

23.75/71.25/5

melt

Addition of block copolymers caused a serious ↓ of droplet size (± 5µm → 0.2 à 4 µm) Compatibilization caused large ↓ of Xc (§ -10%)

[Quirk et al., 1989]

PET/PPE +PET-b-PS

23.75/71.25/5

solid

Addition of block copolymers caused a ↓ of droplet size (± 5µm → 2 à 4 µm) Compatibilization caused Xc ↑ (§ - 10 à 20%)

[Quirk et al., 1989]

Crystallization, Morphological Structure, and Melting of Polymer Blends

If one assumes that for the onset of crystallization ¨F*/kT must be smaller than a certain critical value (i.e., a nucleus of critical size can be formed at the given temperature), independent of the material, and if one neglects that the crystallization also depends on the temperature dependent mobility of the crystallizable segments, the following expression for the relation between ¨y and the degree of undercooling for 2 heterogeneities of type 1 and type 2 can be given [Frensch et al., 1989]: ¨y1/¨y2 § (Tc,1/Tc,2).(¨Tc,1/¨Tc,2)2

(3.48)

where Tc,1 and Tc,2 represent the temperatures at which nucleation is induced by the heterogeneity of type 1 and 2 respectively. In the special case of a homogeneous nucleation, Eq 3.48 can be simplified to read: ¨Tc,hom = Tm° - Tc,hom = Tm°/5

(3.49)

From the latter (Eq 3.49), the homogeneous crystallization temperature for each polymer could be estimated in a simple way. Furthermore, from Eq 3.48 for heterogeneity of type 1 one may write: ¨y1/yp(m,c) § 62.5 (Tc,1/Tm°) (¨Tc,1/Tm°)2

265

number density of the nucleating active species. Among a large number of small polymer droplets, each having a volume VD, the fraction of droplets that contain exactly z heterogeneities of type 1 (inducing normally crystallization in the bulk polymer at Tc°) follows a Poisson distribution function [Pound and LaMer, 1952]: fz(1) = [(M(1)VD)z /z!] exp(-M(1)VD)

(3.51)

where M(1) is the concentration of heterogeneities of type 1, and M(1)VD is the mean number of heterogeneities of type 1 per droplet with volume VD. Hence, the fraction of droplets that contain at least one heterogeneity of type 1 can be given by: f (1)z>0 = 1 - f0(1) = 1 - exp(-M(1).VD)

(3.52)

Now considering that not all droplets have the same size, f (1)z>0 describes that fraction of the droplets (with average volume VD), that crystallize induced by heterogeneity of type 1. The other droplets will crystallize at a different crystallization step. From the relative intensities of the fractionated crystallization steps, one can estimate the concentration of the different heterogeneities, if the mean size of these droplets is known.

(3.50)

From this dependence, the relative values of ¨y for different heterogeneities can be calculated at the corresponding degrees of undercooling (Figure 3.45). Remark however that the relationship Tc,hom = 0,8 TE , used for Eq. 3.49, with TE being the melt/crystal equilibrium temperature of the material, was originally established for metals [Jackson, 1965]. Frensch et al. (1989) have taken over this empirical relationship for polymeric systems, assuming that TE is the equilibrium melting temperature Tm°. A comparison of all data available in the literature on homogeneous crystallization of polymers (see Table 3.25), however suggests that Tc,hom is rather related to the veritable melting temperature, Tm’, of the polymer. From the fractionated crystallization behavior and the blend morphology, one can determine the

Figure 3.45. Plots of the relative specific interfacial energy difference ¨y/y versus the relative undercooling ¨T/T°m at which a heterogeneity nucleates the polymer; (1) and (2): two different heterogeneous nucleations; (3) homogeneous nucleation [Frensch and Jungnickel, 1989].

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In the special case where the usual crystallization from heterogeneity of type 1 is completely suppressed, Eq 3.51 can be written as: M(1)VD << 1

(3.53)

3.4.4.3 Droplet Crystallization in the Presence of a Matrix Melt In most immiscible crystalline/amorphous polymer blends, the crystallization of the dispersed phase occurs in the presence of a molten matrix phase. In the following description, examples will be categorized according to the major classes, as listed in Table 3.14. 3.4.4.3.1 Polyethylene Blends Blends of PS/HDPE have been investigated by Aref-Azar et al. [1980]. Table 3.27 gives an overview of the crystallization behavior in the crystallizable dispersed phase. It should be noted that the crystallization kinetics is related to the size of the dispersed HDPE droplets and the nucleation density. An increase in the amount amorphous PS caused the HDPE phase to be dispersed into finer droplets that, as a result, exhibited a lower degree of crystallinity, Xc, when isothermally crystallized. Furthermore, a higher degree of undercooling was needed to reach the same Xc in blends where the HDPE

phase was dispersed into finer droplets, indicating that crystallization depends on the temperature. The melting behavior of the HDPE phase did not seem to be affected by blending. Recently, Müller et al. [1995], and Morales et al. [1995] have reported on the crystallization of LLDPE that was finely dispersed in a PS matrix. A good correlation was found between the size of the LLDPE phase and the tendency to crystallize in a fractionated way. The authors showed that the relationship is only sensitive to the volume of the dispersed crystallizable droplets and not to the shape of the droplets. 3.4.4.3.2 Polypropylene Blends Numerous studies have been performed on the crystallization behavior of PP in blends with an amorphous component. However, only few authors paid attention to the crystallization behavior of the PP phase when it formed the minor phase of the blend. Ghijsels et al. [1982] investigated the multiple crystallization behavior of blends in which the crystallizable PP phase was finely dispersed into a SBS-rubber (TR). In the case where the latter was finely dispersed, the authors found the PP phase to crystallize at much higher undercooling. A serious drop in the degree of crystallinity, Xc, was also reported. The melting behavior of the fractionated crystallized blend did not seem to be markedly affected, e.g., ¨Hm and Tm remained constant, independent of the amount TR added.

Table 3.27. Evaluation of the blend morphology and thermal behavior in immiscible PS/HDPE blends in which HDPE was the minor phase [Aref-Azar et al., 1980] wt% HDPE

1 5 10 20 a b

Size dispersed phase (µm) 0.1 - 0.3 0.3 - 0.5 2.0 - 3.0 5 - 10

Number droplets (cm-3)

∆Hm a (J/g)

Xc b (%)

1011 - 1012 1010 - 1011 107 - 108 106 - 107

--163 159 184

--55 54 63

∆Hm for pure HDPE is 293 J/g Degree of crystallinity, Xc, for HDPE homopolymer is 80%

Crystallization, Morphological Structure, and Melting of Polymer Blends

Wei-Berk [1993] reported on the crystallization behavior of PP droplets dispersed in a PS matrix. A slight drop in Tc,PP (as the PP phase became the minor phase) was observed. However, the author only investigated the behavior in blends containing more than 35% PP, and did not correlate the crystallization behavior with the blend morphology. Santana and Müller [1994] investigated the same polymer blend. These authors reported that droplets having a diameter of less than 6 µm crystallized at higher undercooling (Tc § 78°C), while the larger droplets crystallized at Tc § 105°C - the latter temperature corresponding to the bulk Tc for the PP used in the studies. The authors referred to the fractionated crystallization behavior caused by a lack of heterogeneities in some of the finely dispersed PP droplets, ending up with the appearance of the homogeneous crystallization peak. Hence, the nucleation mechanism was found to be strongly influenced by the blend morphology. 3.4.4.3.3 Polyester Blends Quirk et al. [1989] investigated the crystallization behavior of PET in a PET/PS 25/75 blend. The PET particle diameter as determined by SEM, was found to be in the range of 5 µm, being quite large due to the large difference in interfacial tension between both phases (hindering easy droplet break up during mixing). Suppression of the cold crystallization in the PET droplets quenched from the melt was observed, along with serious depression of the total degree of crystallinity with increasing content of the amorphous phase.

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3.4.4.3.4 Other Blends Robitaille and Prud’homme [1983] studied the crystallization in the liquid/liquid phase separated melt of the triblock copolymer PEG-PI-PEG having a minor amount of PEG. The authors reported a lower degree of crystallinity of the PEG domains along with a slight melting point depression. Due to the fine dispersion of PEG, the droplets only crystallized at much higher undercoolings (up to 60°C lower than the bulk Tc) and less perfect crystalline lamellae were formed. These lamellae consequently melted at lower temperatures than the usual Tm. The bulk Tc has disappeared completely. The authors related this behavior to the lack of heterogeneities available in the PEG microdomains, which are hence nucleated at much lower temperatures by a homogeneous nucleation mechanism (Figure 3.46). Tang and Huang [1994a] investigated the crystallization behavior of PA-6 is an EPDM matrix. The fractionated crystallization of the PA-6 to occur when the PA-6 content decreased below 15 wt%. Two crystallization peaks were observed, one around the bulk Tc,PA-6 and another at about 25°C lower, caused by the smaller PA-6 droplets having a lack in heterogeneous nuclei. The ratio of the α/γ crystalline form was not altered by the fractionated crystallization, indicating that the lower crystallizing droplets do not crystallize in another crystalline form as in the bulk.

Figure 3.46. DSC cooling and heating curves for the PEG-PI-PEG block copolymer, a pure PEG sample, and PI; the homogeneous crystallization of the PEG segment at much higher degrees of undercooling does not really influence its melting behavior [Robitaille and Prud’homme, 1983].

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3.4.4.4 Droplet Crystallization in the Presence of a Glassy Amorphous Matrix The crystallization of dispersed domains in the presence of a solidified matrix has not yet been a field of active research. Some examples are given below. 3.4.4.4.1 Polyethylene Blends The thermal behavior of PS/LDPE blends has been investigated by Baitoul et al. [1981]. A clear indication of the fractionated crystallization was deduced from the appearance of two additional crystallization peaks around 71 and 64°C in all blends in which LDPE was the dispersed phase. Furthermore, the crystallization kinetics was found to slow down severely when the content of PS was raised. Kunori and Geil [1980] investigated the melting behavior of the binary PC/HDPE blends, in which the weight percentage PE varied between 2 and 10%. The melting temperature of the HDPE droplets did not seem to be affected. 3.4.4.4.2 Polypropylene Blends The majority of published papers on the crystallization of PP in blends concerns those blends in which the Tg of the amorphous component falls below the crystallization temperature of PP. The crystallization of PP in the presence of a glassy amorphous matrix has seldom been reported. The well known blend PC/PP has been intensively investigated by Favis and his co-workers [1987, 1988, 1990, 1992]. Morphology development, rheology, compatibilization, etc. were studied, but no results on the thermal behavior of these blends were reported. 3.4.4.4.3 Polyester Blends Quirk et al. [1989] have reported the crystallization behavior in blends of PPE with PET. Since the blend components showed only a small difference in the interfacial tension, quite small dispersions could be obtained (d § 4µm). The authors reported that the glassy PPE matrix enhanced the cold crystallization of PET after being quenched. Decreasing the PPE content from 75 to 50 wt%

resulted in the nearly disappearance of the cold crystallization exotherm. This behavior was found to be opposite to that in a PS/PET blend (where PET crystallizes surrounded by a PS melt). No clear explanation has been given here. A similar result for PPE/PET blends has been reported by Liang and Pan [1994]. The authors found the cold crystallization temperature, Tc,cold, in the dispersed PET phase to be markedly lower than that of the virgin PET, indicating that PPE may partly act as nucleating agent to promote the nucleation of the PET component. 3.4.4.4.4 Other Blends O’Malley et al. [1969] described the thermal behavior of PEG/PS blends in which PEG was dispersed into fine droplets. A clear indication of fractionated crystallization combined with a simultaneous decrease in the total degree of crystallinity with increasing weight fraction of PS has been observed. Again, a slight decrease of the melting temperature, Tm, with about 2°C was detected, although ¨Hm remained unaffected. This was attributed by the authors to the formation of less perfect crystalline lamellae during the crystallization at higher undercooling. Chang et al. [1991] reported on the melting behavior in PC/POM blends. The blends were found to behave in a similar way as the above described PEG/PS blend. 3.4.4.5 Conclusions It can be stated that the crystallization behavior of a semicrystalline polymer phase, dispersed into an amorphous matrix, is characterized by: (i) Fractionated crystallization or homogeneous nucleation if the minor phase is finely dispersed. Annealing or large droplets resulted in the appearance of a crystallization peak close to the bulk Tc of the homopolymer. (ii) A decrease in the overall degree of crystallinity, Xc, after cooling from the melt, most pronounced in finely dispersed blend morphologies. (iii) A slight decrease of the melting temperature due to the formation of less perfect crystalline lamellae at higher undercoolings. A decrease

Crystallization, Morphological Structure, and Melting of Polymer Blends

of the overall melting enthalpy, ¨Ηm, could be observed clearly, only in blends where the crystallizable dispersed phase did not undergo recrystallization upon heating. 3.4.5

Binary Polymer Blends Containing Two Crystallizable Phases

A large number of polymer blends consists of two crystallizable phases (Table 3.15), hence more studies have been carried out on the thermal behavior of crystalline/crystalline polymer blends. The morphology of a polyblend consisting of two crystallizable polymers can vary depending on the processing conditions, the relative rates and temperature of crystallization of the constituent polymers. These can either crystallize at the same time (coincident crystallization, see further) or separately in a sequential manner, leading to different morphologies and hence different properties. As such, in blends of two semicrystalline polymers, the physical properties may be altered not only by the blend composition and the phase morphology, but also by changing their relative crystallization behavior. Therefore, it is important to study the effect of blending on the crystallization behavior of each component in the blend, to understand the structure development as influenced by melt-processing. Because the phases are physically separated in the melt, the theory concerning the crystallization behavior as discussed in Parts 3.4.3 (matrix crystallization) and 3.4.4 (dispersed droplet crystallization) can be combined to understand the crystallization and melting behavior of most crystalline/ crystalline polymer blends. In general, both crystallizable phases crystallize separately around their characteristic bulk Tc-value (as long as the minor phase is not dispersed into very fine droplets). The Tc-values can be somewhat shifted due to the migration of heterogeneities from one phase towards the other phase, or due to the nucleating activity of one -crystalline or crystallizing- phase at the interface with the second phase. However, changes in the nucleation density of both phases will be more clearly reflected in the spherulite size of each blend component with respect to

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the homopolymer. This can have important consequences for the final mechanical properties of the blend [Friedrich, 1978, 1979]. In the following overview, a survey of the most important topics concerning crystallization behavior in immiscible crystalline/crystalline polymer blends is given. Because the physical state of the second phase affects the crystallization mode of the phase under consideration, a distinction has been made for blends crystallizing in a melt environment and those crystallizing when the second phase has solidified. 3.4.5.1 Crystallization of the Matrix in the Presence of a Molten Dispersed Phase For most commonly studied polymer blends, crystallization of the matrix occurs in the presence of a molten dispersed phase. The crystallization behavior of the continuous phase can be compared to that found for crystalline/amorphous blend systems in which the dispersed amorphous phase was in the molten state. 3.4.5.1.1 Polyethylene Blends Because of the low crystallization temperature of the polyethylenes (HDPE, LDPE, LLDPE, etc.) (see Table 3.15), in most commonly used blends, the dispersed phase has already solidified before the PE matrix starts crystallizing. However, Greco et al. [1987a] studied the crystallization of HDPE (Tc § 118°C) in a 80/20 binary blend with EPR elastomers containing a different ethylene/ propylene ratio. The HDPE phase was reported to exhibit higher Tc values during cooling from the melt, indicating enhanced nucleation, due to the nucleating effect of the EPR copolymers on the HDPE matrix. Furthermore, the melting point, Tm, shifted to slightly higher temperatures relative to the homopolymer due to better crystal perfection as a result of the dissolution of some low molecular weight (‘defective’) HDPE molecules into the EPR copolymer phase during the melt-mixing process. The latter phenomenon was directly related to the ethylene content in the copolymer.

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3.4.5.1.2 Polypropylene Blends Blends of isotactic polypropylene, PP, with a polyethylene are immiscible and, owing to their commercial importance, have been the subject of intensive studies. In these blends, PP crystallization mostly takes place in the presence of molten PE droplets. Long et al. [1991] investigated the crystallization behavior in blends of PP with LLDPE. They found the crystallization temperature of the PP matrix, Tc, to decrease slightly upon the addition of LLDPE. However, the degree of crystallinity, Xc, and the spherulite growth rate, G, were not affected. The authors concluded that the overall crystallization rate of PP in the matrix decreased due to a decreasing primary nuclei density. The latter was confirmed in O. M. experiments by the increased size of the PP spherulites upon the addition of LLDPE. However, Zhou and Hay [1993] reported that with the addition of LLDPE to PP, the crystallization rate remained similar as for the PP homopolymer. Flaris et al. [1993] investigated also the same blend system and reported that blending had a pronounced effect on the lamellar morphology. Furthermore, the isothermal crystallization experiments indicated that the spherulite growth rate, G, and the nucleation density of the PP phase were enhanced. The authors suggested that these observations could be related to the formation of additional nucleation sites, which arise from the polymer-polymer interfaces created by the blending. Because three different observations were reported on the crystallization of the PP matrix in which LLDPE droplets are dispersed, no unambiguous conclusions on this matter can be given. A serious investigation of all factors playing a role here is necessary in the future. Blends composed of a PP matrix with LDPE as the minor dispersed phase have been intensively investigated by Teh [1983], Bartczak et al. [1984], Galeski et al. [1984], and by Teh et al. [1994a]. All authors found LDPE to act primarily as an efficient nucleating agent for the PP matrix, reducing the average PP spherulite size, and to induce the formation of some large β-form PP crystals at the interface with the LDPE phase, that melt at a lower temperature, Tm =155°C as

compared to normal α-form PP crystals with Tm = 165°C. Galeski et al. [1984] and Bartczak et al. [1984] revealed that the nucleating activity of the LDPE phase was mainly attributed to the migration of heterogeneous impurities from LDPE to PP during the melt-mixing process. Furthermore, Galeski et al. [1984] showed the spherulite growth rate of the PP matrix to be unaffected by the dispersed molten LDPE droplets, and showed that these droplets were not rejected by the growing PP spherulites. In the case of PP/HDPE blends, the influence of the HDPE component was more complex and dependent on the physical state of the dispersed HDPE droplets. At a Tc high enough to prevent any HDPE crystallization, the overall rate of crystallization of the PP matrix in isothermal crystallization was found to be strongly reduced by the addition of HDPE [Bartczak et al., 1986]. Since the spherulite growth rate of PP was found to be constant and independent of the blend composition [Teh et al., 1994a], this decrease has been attributed to a decrease in the nucleation density of the PP phase. Bartczak et al. [1986] related this to the migration of heterogeneous nuclei from the PP phase towards the HDPE melt during melt-blending. As a result, the PP spherulite size was found to increase [Bartczak et al., 1986; Teh et al., 1994b]. The same observations were reported for PP/HDPE blends cooled slowly from the melt [Plesek and Malac, 1986]. However, in the case of either an isothermal crystallization at temperatures below the crystallization temperature of HDPE, or crystallization at a higher cooling rate, there may have been migration of nuclei from the PP towards the HDPE phase, but the overall number of heterogeneous nuclei was increased due to the presence of HDPE crystallites that may have acted as additional nucleating centers for PP [Lovinger and Williams, 1980; Gupta et al., 1982; Bartczak et al., 1986; Plesek and Malac, 1986; Teh et al., 1994a; Teh et al., 1994b]. This results in a drastic reduction of the PP spherulite size [Noel and Carley, 1984; Lovinger and Williams, 1980; Plesek and Malac, 1986]. Moreover, Bartczak and Galeski [1986] reported that spherulitic crystallization of a poly-

Crystallization, Morphological Structure, and Melting of Polymer Blends

mer near the interface can cause its deformation, increasing the interfacial area, and can lead to an improvement of toughness and impact properties. Greco et al. [1987b] studied the crystallization in immiscible PP/EPR blends. The average spherulite size in the PP phase was smaller than in the homopolymer. The higher the PP contents (C-3) in EPR, the stronger the nucleating effect for the matrix. The authors experimentally showed that migration of impurities could not cause this effect and that the copolymer composition was the most important factor. An increase in the PP-content of the EPR caused a higher miscibility (defective PP molecules could be partially dissolved in the EPR phase), leading to more perfect PP crystallites melting at a higher Tm and also stronger nucleating effect. Pukansky et al. [1989] investigated both the crystallization and melting behavior, and the global blend morphology in PP/EPDM blends over the whole composition range. Blends quickly

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cooled from the melt did not show significant changes in the crystallization behavior of the PP matrix. However, blends crystallized at a fixed rate of 10°C/min behaved different. Thermograms of the blends containing between 5 and 50 vol% EPDM showed a second melting peak at lower temperature, corresponding to the melting of the β-form of PP. Furthermore, the authors reported that small amounts EPDM slightly increased the Tc, PP, but did not affect the degree of crystallinity. Dispersed EPDM droplets thus seem to promote the formation of the hexagonal β-form of PP. An overview of the effects affecting the primary nucleation in immiscible PP based blends is provided in Table 3.28. 3.4.5.1.3 Polyethylene Terephthalate Blends Wilfong et al. [1986] reported on the effects of blending low concentrations (1 to 10 wt%) polyolefin with PET on the crystallization and toughening behavior of the latter. The authors

Table 3.28. Overview of the phenomena influencing heterogeneous primary nucleation in polypropylene based immiscible blends (after Bartczak et al., 1995). Blend system(a)

Migration of impurities (b)

Crystallization of the second component(b)

Influence of the interface (b)

References

PP/LLDPE

--- (d) ↓↓

-

PP/LDPE

↑↑

------↑ (c)

PP/HDPE

↓↓↓

[Zhou and Hay, 1993] [Long et al., 1991] [Flaris et al., 1993] [Bartczak et al., 1984], [Teh, 1983], [Galeski et al., 1984], [Teh et al., 1994a] [Bartczak et al., 1986], [Teh et al., 1994b], [Teh et al., 1994a] [Lovinger and Williams, 1980], [Gupta et al., 1982], [Bartczak et al., 1986], [Teh et al., 1994] [Greco et al., 1987b]

PP/EPR (a) (b)

(c) (d)

↓↓↓ (c)

↑↑↑ (c)

---

---

↑↑ ↑

Data concern the crystallization of the matrix polymer (mentioned first in the blend code) ↑ indicates an increase of the nucleation density in the blend, ↓ indicates a decrease of the nucleation density (the number of arrows is related to the intensity of the effect) Found for samples crystallized non-isothermally --- indicates that the authors did not find evidence explicitly for the mentioned topic to influence the nucleation of PP in the blend system described

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studied blends of PET with LLDPE, HDPE, PP, and poly(4-methylpentene-1), all of them having a lower melting point than PET (Table 3.15). Polyolefin melts did not enhance the nucleation of PET, although the spherulite size of the PET matrix was found to be 2.5 to 3 times larger than for the homopolymer, with a broader spherulite size distribution. Both the crystallization rate and the degree of crystallinity were found to be reduced by blending. This was attributed to the expense of energy that was required by the crystallizing growth front to reject and deform the polyolefin dispersed molten droplets. Martuscelli [1984] and Bartczak et al. [1984] have calculated that the rejection and/or deformation of dispersed droplets by the crystallizing growth front can cause a marked depression of the spherulite growth rate, G. 3.4.5.1.4 Poly(phenylene Sulfide) Blends Poly(phenylene sulfide), PPS, is an expensive, high-performance, but brittle specialty resin. Blending can offer a good alternative both in toughness improvement and cost reduction [Nadkarni and Jog, 1991]. Shingankuli et al. [1988] and Jog et al. [1993] investigated the influence of blending PPS with PET on its thermal and crystallization behavior. Blending was found to enhance the PPS nucleation. Isothermal crystallization experiments revealed that the crystallization time of PPS decreases along with the crystallization induction time. Both parameters were found to depend on composition. Optical microscopy confirmed this and revealed that the size of PPS spherulites in PPS/PET blends was drastically reduced as compared to the homopolymer. Furthermore, the degree of crystallinity of the PPS phase decreased with increasing PET concentration. However, dynamic crystallization experiments showed a constant value of Tc,PPS. The authors have related the accelerated crystallization of PPS in a blend with PET to the nucleation at the interface of the PET droplets. Owing to its supercooled state, the PPS matrix consists of highly ordered chains. Nadkarni and Jog [1986], Nadkarni et al. [1987] and Jog et al. [1993] investigated the

crystallization in blends of PPS with three types of HDPE, having a different melt flow index. In contrast to the PPS/PET blends, PPS crystallizes in a superheated HDPE melt environment. From the dynamic cooling experiments, it was found that the presence of the HDPE melt suppresses the crystallization of PPS. The crystal growth rate, G, of PPS was found to remain unchanged, but its nucleation density was reduced as the concentration of HDPE in the blend increased, or, when the melt-viscosity of the HDPE phase decreased. As a consequence, the overall crystallization rate of PPS was found to be retarded. 3.4.5.1.5 Other Blends Chen et al. [1988] reported about blends of polyamides with a polyolefin. PA-11/LDPE blends and PA/HDPE blends both showed an increase of the melting temperature of the PA-11 matrix due to the addition of the polyolefin. No further attention was paid to this phenomenon. Holsti-Miettinen et al. [1992] and Ikkala et al. [1993] recently studied the crystallization behavior of PA-6 blended with PP. No shift of the crystallization temperature of the PA-6 matrix was observed in the blends; the dispersed PP droplets did not influence the crystallization behavior of the matrix. Frensch and Jungnickel [1989] investigated the influence of blend composition on the crystallization and melting behavior of PA-6/PVDF blends and PBT/PVDF blends. The crystallization of the PA-6 matrix and PBT matrix was promoted by the dispersed molten PVDF phase, as indicated by the rise in their Tc in the blends, while their relative crystallinity remained unaffected. The authors assigned this increase in Tc to migration of nucleating heterogeneities from the dispersed PVDF phase towards the matrix phase during melt-mixing of the blends. 3.4.5.2 Crystallization of the Matrix in the Presence of a Solidified Dispersed Phase The crystallization of a polymer in the presence of solidified domains of the second phase takes

Crystallization, Morphological Structure, and Melting of Polymer Blends

place through a heterogeneous nucleation process. Since the rate of heterogeneous nucleation is higher than that of homogeneous nucleation, and since primary nucleation is the rate-controlling step for polymer crystallization, the crystallization rate is expected to be higher in such blends when compared to homopolymers [Nadkarni and Jog, 1991]. 3.4.5.2.1 Polyethylene Blends On account of their commercial interest, the crystallization of HDPE, LDPE and LLDPE in blends with PP has been extensively investigated. In these systems, the PP phase solidified already before the PE matrix starts crystallizing. In the case of LDPE/PP blends, not much attention has been focused on the case where the LDPE phase forms the matrix. Teh [1983] reported no shift in the melting temperature of the LDPE matrix in the presence of solidified PP domains. Bartczak and Galeski [1986] observed that the LDPE crystallinity remained unaffected by blending. Zhou and Hay [1993] investigated the crystallization behavior in LLDPE/PP blends. The crystallization rate of the LLDPE matrix, measured from isothermal DSC experiments, was not affected by the dispersed PP domains. However, its degree of crystallinity slightly decreased with increasing PP content in the blend. According to the authors, this could be ascribed to the lower degree of perfection of the LLDPE crystals. More extensive investigations have been performed on HDPE/PP blends by Martuscelli et al. [1980] and Bartczak and Galeski [1986]. From the isothermal crystallization experiments, it was found that the rate of crystallization of the HDPE matrix was markedly reduced upon addition of small amounts of PP (10 wt%). The authors attributed this phenomenon to the increased meltviscosity of the sample caused by the presence of solidified PP domains. Moreover, Plesek and Malac [1986] have calculated from the surface tensions of the homopolymers at Tc, that PP crystallization will not cause the nucleation of the HDPE phase, while in the reverse case HDPE crystals will induce the nucleation of PP.

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Similar results were reported by Nadkarni and Jog [1986] and Nadkarni et al. [1987] for HDPE/PPS blends. The degree of crystallinity of HDPE in blends with a HDPE matrix was not affected by blending. The degree of supercooling required for initiating non-isothermal crystallization of HDPE was surprisingly not affected by the presence of solid PPS domains. However, isothermal crystallization half-times for HDPE in the blends containing more than 10 wt% PPS were longer than for the HDPE homopolymer. Again, this has been attributed by the authors to the increased melt-viscosity due to the presence of solidified PPS domains. Frensch et al. [1989] reported on the crystallization of HDPE in a blend with POM. The HDPE matrix crystallized in all samples at almost the same temperature and to the same extent, independent of the extrusion time. 3.4.5.2.2 Polypropylene Blends The majority of papers related to the crystallization of isotactic polypropylene, PP, based blends concern those where the PP matrix crystallizes in the presence of a molten dispersed phase of polyethylenes and olefinic elastomers. As a result, crystallization of a PP matrix in the presence of a solidified dispersed polymer has seldom been reported [Nadkarni and Jog, 1991]. Shingankuli [1990] studied the crystallization behavior of PP in the presence of solidified PVDF domains. A higher crystallization temperature of the PP matrix phase was observed, indicating an enhanced nucleation in the blends. The degree of crystallinity of PP was found to increase by about 30 to 40% with increasing PVDF content. Isothermal crystallization studies also confirmed the acceleration of the overall crystallization rate in terms of shorter crystallization half-times for PP. More efforts have recently been dedicated in understanding the crystallization behavior in PP/PA-6 blends. Holsti-Miettinen et al. [1992], Moon et al. [1994] and Ikkala et al. [1993] found that the crystallization temperature of the PP matrix by cooling from the melt rises by about 10°C by adding PA-6. Ikkala et al. [1993]

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observed that the largest temperature increase was caused at a PA-6 concentration of about 20 wt%; in this case, the PA-6 dispersion size was quite small (2.5 µm). Moon et al. [1994] have related this temperature shift to the migration of heterogeneous nuclei towards the PP matrix during the melt-mixing process, together with the nucleating agent-like behavior of the solidified PA-6 domains. No change of the melting peak has been noticed [Park et al., 1990]. Grof et al. [1989] performed some isothermal crystallization experiments on fibers of the PP/PA-6 blend. In accordance with the cited findings, the latter authors reported a decrease both in the crystallization half-time and the induction time for crystallization of PP in PP/PA-6 blends, while no change in the degree of crystallinity was observed. Tang and coworkers [1994] investigated briefly the crystallization behavior of PP in blends with PA-12. The melting point remained unaffected by blending. However, a slight shift of the crystallization peak (about 2.5°C) upon cooling from the melt, was reported for blends comprising 33 wt% PA, along with an increase of the height of the Tc peak. The PP matrix has been nucleated by the dispersed PA-12 domains. The authors related this to the fine morphology; at the interface of the phases epitaxial crystallization had also been observed. This was also the reason why the PA-12 phase in the blends only existed in the γ-form. However, it should be mentioned that these PP/PA-12 blends were prepared from solution. 3.4.5.2.3 Polyethylene Terephthalate Blends Only few papers related to the crystallization of a PET matrix in immiscible crystalline/crystalline blends have been published. Shingankuli et al. [1988] investigated the thermal behavior of PET blends with the glass-fiber reinforced polymer PPS. Dynamic crystallization experiments revealed that the PET crystallization behavior was significantly altered by blending. Upon the addition of PPS, both the onset temperature for crystallization and the peak value, Tc, showed a dramatic shift to higher temperatures (up to 20°C). Also, the degree of crystallinity significantly increased in the blends. The author

attributed the phenomenon to the heterogeneous nucleation induced by the glass-fibers in the PPS phase and the nucleating activity of the already solidified PPS domains. As a result, the PET matrix in the blends became more richer in heterogeneous nucleating sites as compared to virgin PET. Isothermal experiments confirmed these conclusions and showed that the crystallization half-time of PET decreased drastically in the blends (attributed to the enhanced nucleation). Furthermore, an increase of the onset of melting of the PET matrix (15°C) with increasing content of PPS in the blends has been observed. The melting behavior PET in the blends has been explained by the formation of larger and more perfect crystallites (due to the nucleation at higher temperatures) with a narrower size distribution, and by an increased degree of crystallinity. 3.4.5.2.4 Other Blends Frensch and Jungnickel [1989, 1991] and Frensch et al. [1989] have investigated the thermal behavior of polyvinylidene fluoride, PVDF, in blends with polyamides, in relation to the blend morphology. PA-6 droplets could be finely dispersed into the PVDF matrix. The crystallization temperature of the PVDF matrix did not seem to be affected in the blends. A similar behavior was observed in PVDF/PA-66 blends. Investigations on the crystallization behavior of PVDF in a blend with polybutylene-terephthalate, PBT, have been reported by Frensch and Jungnickel [1989] and Frensch et al. [1989]. PBT dispersed droplet size was found to be an order of magnitude larger than the dispersed PA droplets in PVDF blends. However, in this case, the Tc,PVDF displayed a shift to higher temperatures (2 to 8°C) upon blending with PBT, which was attributed to the nucleating efficiency of amorphous or crystallizing PBT domains (which subsequently crystallized coincidentaly with the PVDF matrix). 3.4.5.3 Crystallization of the Dispersed Phase in the Presence of a Matrix Melt Immiscible blends most often show a two phase morphology consisting of a continuous matrix

Crystallization, Morphological Structure, and Melting of Polymer Blends

and a droplet-like dispersed phase beyond the phase inversion region. From Section 3.4.4 it is clear that the crystallization behavior of droplets can be dramatically affected as compared to the homopolymer. In summary: (i) dispersed drops can have an altered nucleation density, caused by the migration of heterogeneous nuclei during the melt-mixing process, they can be nucleated by a crystallizing or solidified matrix, the interface can induce some additional nucleating centers, etc. (ii) The smallest dispersed droplets can suffer from the lack of heterogeneous impurities in each droplet, what may result in a fractionated crystallization. In some cases, this can give rise to the coincident crystallization of the dispersed phase with the (lower crystallizing) matrix (see Section 3.4.5.6). 3.4.5.3.1 Polyethylene as Dispersed Phase Because of the low crystallization temperature of all polyethylenes as compared to most other commonly used thermoplastics, crystallization will proceed most often in an already solidified matrix. No literature could be found on the crystallization behavior of PE in a molten matrix environment. 3.4.5.3.2 Polypropylene as Dispersed Phase Typical polymer blends with isotactic polypropylene, PP, are the PP/PE blends, in which PP is the first crystallizing component. Zhou and Hay [1993] investigated the crystallization in LLDPE/PP blends. They reported that the extent of crystallization in PP droplets is seriously hindered by the low nucleation density of PP, resulting in a serious drop of the degree of crystallinity during the isothermal measurements. From these experiments it could be predicted that cooling from the melt would result in a fractionated crystallization (30 wt% PP) or even homogeneous crystallization (10 wt% PP). Similar results had already been reported by Long et al. [1991], Pukanszky et al. [1989], and recently by Müller et al. [1995], and Morales et al., [1995]. The latter authors even mentioned that the retarded crystallization of PP droplets in some cases finally resulted in the coincident crystallization of PP

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with the LLDPE matrix. Furthermore, a partial change in the crystallographic form from α to the lower melting β-form was observed. Lovinger et al. [1977] reported that the β-form is nucleated at a lower rate than the α-form, and hence promoted on homogeneous nucleation. Teh [1983] reported only the melting behavior of LDPE/PP blends - no shift in Tm,PP was seen. An enhancement in the formation of the β-form PP spherulites in the LDPE melt was observed. Blends of HDPE with PP have been studied by several authors. However, not much attention has been focused on the crystallization behavior of the dispersed phase yet. 3.4.5.3.3 Polyamide as Dispersed phase Several blends with polyamides, crystallizing at high temperatures, have been studied. Chen et al. [1988] investigated the phase morphology and melting behavior of HDPE/PA-11 and LDPE/ PA-11 75/25 blends. The melting point of the dispersed PA-11 phase was found to be unaffected by blending. Several studies have been performed on the thermal behavior of PP/PA-6 blends. Park et al. [1990] reported a melting point depression for the dispersed PA-6 phase (about 4°C), having an average particle size of 2 to 5 µm at 25 wt% PA-6 in the blend. However, the relation between the crystallization phenomena and the blend morphology were not explored. Ikkala et al. [1993] have investigated the correlation between the blend morphology, crystallization and melting behavior of the minor component in PP/PA-6 blends. The PA-6 phase was reported to crystallize at its bulk temperature. However, compatibilization (resulting in the formation of a finer dispersion) did not show any crystallization exotherm around the bulk Tc,PA-6. This could be explained by the retarded crystallization caused by a lack of heterogeneous nuclei in the PA-6 droplets. Finally, the nucleating activity of both blend components on each other caused the coincidental crystallization of the PA-6 with the PP matrix. Moon et al. [1994] also investigated the thermal behavior of PP/PA-6 70/30 blends. The authors reported the Tc of the PA-6 droplets to rise remark-

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ably (by about 14°C) as compared to the Tc of the virgin PA-6. This rise in Tc,PA-6 was explained by analogy to findings of Khanna et al. [1988] on pure virgin PA-6 homopolymer, suggesting that melt extrusion of PA-6 would lead to a more ordered molecular arrangement that persisted in the molten state due to hydrogen bonding, and as such caused a faster crystallization. This has been confirmed by crystallization experiments on melt-extruded PA-6 homopolymer. The results of the blends as compared to melt-mixed pure PA-6 agree with those reported by Ikkala et al. [1993] - no shift in the Tc,PA-6 was caused by blending. Furthermore, they also reported that compatibilization of the blends caused a decrease of the dispersed phase size, leading to fractionated and subsequently coincident crystallization. Frensch and Jungnickel [1989] and Frensch et al. [1989] tried to elucidate the crystallization behavior of the minor phase in the binary PVDF/PA-6 blends, in relation to the final blend morphology. They reported that the crystallization of the PA-6 droplets was fractionated and/or retarded, depending on the number of mixing cycles and dispersion size. The smaller the PA-6 droplets, the more pronounced the retardation of the crystallization peak (¨Tc § 40°C). Nevertheless, the melting endotherm remained unaffected. They concluded that part or all of the PA-6 phase finally coincidentaly crystallized with the PVDF matrix due to the specific mutual nucleating efficiency of both components. A similar behavior has been reported by Frensch and Jungnickel [1991] for PVDF/PA-66 blends. In this case, the undercooling associated with the retarded crystallization was about 90°C higher than the one for the bulk crystallization! The size of the dispersed PA-66 droplets has been found to be only about 0.3 µm. The authors concluded that the appearance of fractionated and coincident crystallization is correlated with the low interfacial energies between the amorphous melt phases, providing a high level of dispersion, and between the crystalline phases, providing a nucleating efficiency.

3.4.5.3.4 Other Blends Shingankuli et al. [1988] reported on the crystallization of dispersed PPS domains in a PET matrix. The onset of crystallization of the dispersed PPS domains decreased (by about 7°C) with decreasing PPS content, together with the crystallization peak and the degree of crystallinity. The authors concluded that the PPS crystallization was retarded mainly when the PPS content in the blends was below 20 wt%. Furthermore, the onset of melting of the PPS fraction remained nearly unaffected, except for those blends containing less than 20 wt% PPS. In the latter case, the onset of melting seriously decreased (by about 30°C), whereas the melting peak temperature and heat of fusion remained constant. This can be attributed to the lower crystallization temperature of the PPS droplets leading to the formation of less perfect, lower melting crystallites. Klemmer and Jungnickel [1984] have reported on the fractionated crystallization of POM in an HDPE matrix. They found an additional crystallization peak of POM to occur 14°C lower than the bulk crystallization peak. This was attributed to the fractionated crystallization of POM, caused by an interface-induced additional inhomogeneous nucleation and crystallization. It was shown that this phenomenon only occurs in those blends where the number of the dispersed particles was higher than the number of available heterogeneous particles. Moreover, the preparation method clearly influenced the fractionation due to the change of the particle sizes - fractionated crystallization has been observed only in meltmixed blends. Frensch and Jungnickel [1989] and French et al. [1989] have investigated PVDF/PBT blends and related their thermal behavior with the blend morphology. Similar to PVDF/PA-6 blends, the PBT droplet crystallization was completely suppressed in a 85/15 blend and finally crystallized coincidentaly with the PVDF matrix. Again this phenomenon could be related to the fine dispersion of PBT droplets, in number exceeding the available nuclei. Shorter melt-mixing cycles caused a coarser dispersion leading only to a

Crystallization, Morphological Structure, and Melting of Polymer Blends

fractionated crystallization of PBT at Tc,bulk and at Tc,PVDF. 3.4.5.4 Crystallization of the Dispersed Phase in the Presence of an Already Solidified Matrix 3.4.5.4.1 Polyethylene as Dispersed Phase. PP/PE blends have been studied extensively by several authors. Zhou and Hay [1993] reported that the dispersed LLDPE droplets in PP/LLDPE blends showed problems in nucleating at the normally expected bulk crystallization temperature, Tc. Also, a serious decrease of the degree of crystallinity from isothermal measurements, as the LLDPE content decreased, could be observed. Contrary to these observations, Müller et al. [1995] stated that the LLDPE droplets do not exhibit fractionated crystallization when they are dispersed in a PP matrix (although they do in a PS matrix), because of the nucleating effect of the solidified PP matrix on the LLDPE droplets. Galeski et al. [1984] and Teh et al. [1983] have investigated PP/LDPE blends. No shift of the melting peak for LDPE has been observed. Both authors showed migration of the impurities during the melt-mixing process from the PP towards the LDPE phase. No further details on the crystallization behavior of the LDPE droplets themselves were reported. Nadkarni and Jog [1986] have reported on PPS/HDPE blends. The degree of crystallinity of HDPE was reduced when HDPE was the minor phase. Furthermore, the Tc,HDPE shifted to somewhat lower temperatures (by about 5°C) but only in those blends with a low HDPE content. Isothermal crystallization half-times for HDPE in its blends with PPS decreased as the HDPE content decreased, indicating an enhanced nucleation from the solidified PPS interfaces. Chen et al. [1988] have investigated the melting behavior of 75/25 PA-6/HDPE and PA/LDPE blends. No shift has been observed in the melting point. No attention has been focused to the crystallization of the PE droplets.

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3.4.5.4.2 Polypropylene Blends Blends of PA-6 with PP dispersed as fine droplets have been examined recently by several authors. Ikkala et al. [1993] investigated the thermal behavior and morphology of blends of PA-6 in which PP had been dispersed. In binary blends PP droplets crystallized even at somewhat higher temperature (by about 5°C) than the PP homopolymer, attributed to the nucleating activity of the solidified PA-6 matrix towards the dispersed PP phase. Morphological investigations revealed that the PP dispersion in the blends was quite coarse; so nearly every droplet contained the heterogeneities that usually nucleate PP. However, upon compatibilization this behavior changed. Compatibilizers that formed an immiscible interlayer between PA-6 and PP and caused a reduction of the dispersed particle size, gave rise to a retarded crystallization of the PP phase in a PA-6/PP 80/20 blend, decreasing the Tc,PP by 50°C! This behavior was directly caused by the small size of the dispersed phase and the prevented nucleation from the solidified matrix. Blends containing 40 wt% PP did not crystallize in a retarded way due to their coarser droplet size, but clearly were not nucleated by the PA-6 phase as seen from Tc,PA-6 = T°c,PA-6. Similar results have been presented by Holsti-Miettinen et al. [1992]. 3.4.5.4.3 Other Blends Tang and Huang [1994b] investigated the relation between blend morphology and crystallization behavior in PP/PEG blends, prepared by solution blending. They reported that the PEG phase crystallized fractionated at different degrees of undercooling, but was always nucleated heterogeneously. The authors related the different crystallized fractions to PEG droplets of different sizes; the largest droplets crystallized at the bulk crystallization temperature. Shingankuli et al. [1988] studied the crystallization behavior of dispersed PET droplets in a PPS matrix. A serious increase of the crystallization temperature of the dispersed PET phase (by about 20°C) during cooling experiments from the melt, was explained as a result of

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the nucleating activity of the glass-fibers in the PPS matrix, but also from the solidified PPS itself. As a result, the crystallization became more heterogeneous and the crystallization peak width decreased drastically. A corresponding increase in the onset of melting for PET (about 15°C) was attributed to the formation of thicker and more perfect PET crystallites in the blends. Frensch and Jungnickel [1989] and Frensch et al. [1989] have studied PA-6/PVDF blends. The authors reported that the finely dispersed PVDF droplets crystallized fractionated at different undercoolings. Again this could be directly related to the lack of heterogeneous nuclei in some of the smallest droplets. Increasing the blend composition or decreasing the mixing cycles caused the crystallization of the PVDF droplets to shift to higher temperatures, due to the formation of a coarser morphology. A similar behavior has been reported for PA-66/PVDF blends [Frensch and Jungnickel, 1991]. A reverse case however has been reported by the same authors for the crystallization of dispersed PVDF droplets in a solidified PBT matrix. In the latter case, Tc,PVDF even shifted to higher temperatures (by about 5°C) than for homopolymer crystallization. The shift seemed to become less pronounced as the number of mixing cycles increased. No explanation for this behavior was reported. The melting endotherm of the PBT droplets was not affected by the blending. 3.4.5.5 Coincident Crystallization in Crystalline/Crystalline Polymer Blends A few authors have observed coincident crystallization of both phases in crystalline/crystalline immiscible blends. This phenomenon was reported for blends in which the minor phase exhibits a higher degree of undercooling for crystallization due to its fine dispersion (see Section 3.4.4.) and the matrix phase crystallizes at its bulk Tc that is lower than that of the minor phase. An additional factor that should be taken into account is that a heterogeneous nucleation is promoted on surfaces with a high interfacial tension [Helfand and Sapse, 1975] (i.e., a crystal-

lizing phase boundary). This can lead to the ‘coincident crystallization’ of both phases, as it has been reported by Frensch and Jungnickel [1989, 1991] and by Frensch et al. [1989]. 3.4.5.5.1 Principle of Coincident Crystallization It has been observed that this phenomenon is connected with the phase dispersion of the minor component, and is enhanced when the dispersion becomes finer. Upon cooling from the melt, a finely dispersed phase can exhibit fractionated crystallization, what implies that none, or only part of the dispersed droplets crystallize at their bulk Tc. This type of crystallization is related to the lack of heterogeneities in the droplets, required for nucleation at the bulk Tc. When the blend is now further cooled, two possible ways of primary nucleation are possible. In a first case, the matrix phase is nucleated by heterogeneous species present in this phase and instantly, newly created crystals appear. Hence, the crystallization temperature of the matrix will be situated at its bulk Tc. A second possibility for coincident crystallization occurs in the case one finds again a single crystallization peak for the matrix phase, which however takes place above its bulk Tc. Some novel mutual nucleating mechanism was suggested in such blends; a molten component (minor phase) acts as nucleating substrate for the matrix, which instantaneously crystallizes [Frensch and Jungnickel, 1989]. For both cases, when the ¨y-value (see Section 3.4.4.2.) between these newly formed crystals and the melt of the minor phase is smaller than that of all other heterogeneities present in the minor phase (except probably the nuclei of ‘type 1’ normally nucleating around the bulk Tc,minor), its associated specific undercooling must be so small, that the crystals can induce the crystallization of that minor phase from the instant of their own creation [Frensch et al., 1989]. Consequently, a single coincident crystallization peak will be registered in DSC-thermograms. It is clear that this phenomenon is phase morphology-dependent. Only in those blends where the minor phase is dispersed into sufficiently fine droplets, this phase has the opportunity to exhibit

Crystallization, Morphological Structure, and Melting of Polymer Blends

fractionated crystallization. Hence, only at low blend compositions and/or good matching viscosities of both phases (where the capillary number Ca predicts droplet break-up being dominant above coalescence), the occurrence of coincident crystallization is possible. 3.4.5.5.2 Examples of Coincident Crystallization Frensch and Jungnickel [1989, 1991] and Frensch et al. [1989] have investigated the crystallization behavior of PVDF/PA-6, PVDF/PA-66 and PVDF/ PBT blends. The PVDF/PA-6 blends showed a composite droplet type morphology (finely dispersed matrix droplets encapsulated in the minor phase droplets) that disappeared after sufficiently long mixing cycles. Along with these observations, coincident crystallization was found in PVDF/PA-6 blends for a 85/15 and 75/25 composition only. The influence of morphological changes could be significant; after 4 mixing cycles, the dispersed PA-6 droplets became finer and did not contain small PVDF inclusions anymore. Along with this observation, only one single coincident crystallization peak could be found from DSC. The small exotherm at 184°C caused

279

by some larger PA-6 domains containing the PVDF inclusions with a small exotherm at 113°C had disappeared completely by the mixing (Figure 3.47). PVDF crystallization was found to be initiated by nucleation from heterogeneities of ‘type 1’ at the bulk Tc,PVDF. A similar behavior has been reported for PVDF/PA-66 blends by Frensch and Jungnickel [1991]. A second system investigated by the authors was the PVDF/PBT blend. Similar effects could be observed. However, coincident crystallization in the PVDF/PBT 85/15 blend occurred at a somewhat higher temperature than the bulk Tc,PVDF. It could be concluded that in this case, the PBT melt induced the crystallization of the PVDF matrix phase. Besides the cases of coincident crystallization reported previously, recent investigations on PP/PA-6 blends in which a compatibilizing agent had been used to obtain a finer and more homogeneous dispersed phase morphology also mentioned coincident crystallization of the PA-6 droplets with the PP matrix [Ikkala et al., 1993; Moon et al., 1994]. However, this has not been observed in the binary blend.

Figure 3.47. Retarded and/or fractionated crystallization causing coincident crystallization in PVDF/PA-6 and PVDF/PBT blends. Influence of the blend composition (a) and the number of extrusion cycles Z (b) [Frensch and Jungnickel, 1989].

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G. Groeninckx, M. Vanneste and V. Everaert

3.4.5.6 Effect of Compatibilization on the Crystallization Behavior in Crystalline/Crystalline Polymer Blends Blending offers an interesting means of tailoring product properties to specific applications. However, in the case of immiscible polymer pairs, the desired properties are not achieved readily without a compatibilizer, which enhances the phase dispersion and stability, as well as a good adhesion between the phases. This can be effectuated by physical or reactive methods [Folkes and Hope, 1993]. Compatibilization strongly affects the blend phase morphology and as such, it also may influence the crystallization behavior of the blend [Flaris et al., 1993]. Because both factors are related to the final properties of the blend, it is worth paying attention to these phenomena. Several authors have investigated the influence of compatibilization on the global blend morphology. However, only a few authors really tried to understand the effect of compatibilization in crystalline/crystalline polymer blends on the crystallization kinetics, melting behavior and semicrystalline morphology of the components. In Table 3.29 some recent results on this topic are summarized. From the data presented in this Table, it appears that in contrast to binary blends without a compatibilizer, the crystallization of the minor component in compatibilized blends cannot be solely explained by the size of the dispersion [Ikkala et al., 1993; Flaris et al., 1993; Tang and Huang, 1994a; Holsti-Miettinen et al., 1995]. Other factors affecting the crystallization are the type of compatibilizer and its degree of miscibility with one or both of the blend components, the amount of compatibilizer added, the amount of interface created, and other effects. The general influence of a compatibilizer on the crystallization behavior of an immiscible polymer blend system is still far from being well understood. However, abstract can be made between two main classes. A first class consists of compatibilizers that form a kind of ‘immiscible’ interlayer between the two phases. Examples are given by Holsti-Miettinen et al. [1992], and

Ikkala et al. [1993] for PP/PA-6 blends to which MAH-g-SEBS, FA-g-EBA and GMA-g-EEA have been added. The compatibilizer prevents direct nucleating effects from one phase on the other. As such, only the size of the dispersion relative to the nucleation density of the dispersed phase and the nucleating effect of the compatibilizing agent itself play a role in the crystallization behavior (Figure 3.48). Remark however that the size of the dispersion is often directly related to the concentration of the compatibilizer added [Moon et al., 1994]. A second class consists of compatibilizers that have an analogous chemical structure compared to one or two of the blend components. Here, the influence of a compatibilizer on the crystallization behavior of both phases is complex. Several factors have to be taken into account: nucleating effect of the matrix on the dispersed phase or from the dispersed phase on the matrix, the size of the dispersed phase relative to the nucleation density of that phase (and thus to the composition, content of the compatibilizer, etc.), nucleating effect of the compatibilizer itself, interactions of the compatibilizing agent and one or both phases which can impede the crystallization, co-crystallization of the compatibilizer with one of the phases, etc. An illustration is given in Figure 3.48 and Figure 3.49. Again, the concentra– tion of compatibilizer plays a crucial role [Moon et al., 1994]. Compatibilization seems to be of industrial interest in several ways: besides the improvement of the phase dispersion and adhesion, leading to superior mechanical properties, it also often can prevent the minor crystallizable dispersed phase from fractionated or retarded crystallization, which make faster production times and higher thermal stability of the products possible. 3.4.5.7 Conclusions on the Crystallization Behavior of Immiscible Crystalline/Crystalline Polymer Blends The scientific literature on crystallization in polymer blends clearly indicates that the crystallization behavior and the semicrystalline morphology

80/20 60/40

PP/PA-6 +MAH-g-SEBS +MAH-g-PP +FA-g-EBA +GMA-g-E EA

80/20 60/40

PA-6/PP +MAH-g-PP +MAH-g-SEBS +FA-g-EBA +GMA-g-E EA

77/23 50/50

PP-g-MA/PA-6

70/30

PP/PA-6 + MAH-g-PP

Crystallization

Explanation

composition type of compatibilizer

* PP dispersed 20% → depending on type of compatibilizer: . cryst. around Tc,bulk . retarded cryst. (∆T§ 50°C)

type of compatibilizer

* If SEBS, EBA or E EA => immiscible interlayer between phases ⇒ small droplets with lack of nuclei=> Tc ↓↓ * PP droplets are large (coalescence) and contain enough heterogeneities to cryst. at Tc,PP

* If PA-6/PP phases in direct contact (no comp or analog. comp. as MA-g-PP) => PA-6 nucleates PP droplets => no lack of nuclei in droplets => Tc § Tc,bulk

Compatibilization => interfacial tension ↓ => droplet size ↓ ⇒ lack of heterogeneity of type 1 → crystallization by nuclei of type 2 which in his turn induces PP-MA crystallization

* PP matrix: Tc § Tc,bulk but * PA-6 has nucleating effect on PP cryst. at interface if in direct contact slight ↑ in Tc (10°C) in case of binary blend and SEBS, EBA and E EA form immiscible interlayer MAH-g-PP between PA-6 and PP => no nucleation => Tc,PP § Tc,bulk * PA-6 dispersed: coincident * Compatibilization => serious reduction of PA-6 crystallization with PP droplet size => lack of nucleating species => retarded cryst.

* PP dispersed 40%: no large shift down of Tc

* PA-6 matrix: Tc § Tc,bulk

* PA-6 dispersed: fractionated crystallization => coincident cryst. with PP

* PP matrix: Tc § Tc,bulk

composition

composition

concentration compatibilizer PP matrix: Tc,PP matrix § cte (in phr) PA-6 droplets: Increase of phr MA-g-PP => interfacial tension ↓ Tc,PA-6 ↓ from 195°C (0phr) to 188°C (1.5 phr) => dispersion size PA-6 ↓↓ => lack of heterogeneities in each droplet Xc,PA-6 also ↓ From > 2 phr MA => => retarded crystallization coincident crystallization of PA-6 with PP matrix Constraint effect of grafted PA-6 chains is negligible (∆T = 60°C)

Blend System (a) +compatibilizer Studies on the Influence of

Table 3.29 Influence of a compatibilizing agent on the crystallization behavior of binary crystalline/crystalline polymer blends

[Holsti-Miettinen et al., 1992] [Ikkala et al., 1993]

[Ikkala et al., 1993]

[Holsti-Miettinen et al., 1992]

[Tang and Huang, 1994a]

[Moon et al., 1994]

Reference

Crystallization, Morphological Structure, and Melting of Polymer Blends

281

(a)

concentration of ionomer

type of compatibilizer

type of compatibilizer

Dramatic ↑ in crystallization halftimes . Tc,PA-6 ↓ with [ionomer] ↑ (by ± 4°C) . Relative Xc ↓ with [ionomer] ↑ . Much wider distribution of PA-6 crystallite sizes

Strong interactions between PA-6 and ionomer impeded the crystallization of PA-6

Rate of nucleation ↑ Tc,PP ↑ (3 to 4°C) when compat. is added: - peak temp. independent on type of compat. - onset temp. is dependent on type compat. Half-width of exotherm ↓ Crystal size distribution becomes more uniform strongly by addition of due to higher nucleation rate compat., independent of type used

Nucleation density at Compatibilization changes the interface morphology => additional nucleation sites Tc= 135°C: - ↑↑ with EP from interfaces - ↓ with SEBS Spherulite growth rate Finer dispersion of LLDPE can cause more ↑ upon compatibilization unidimensional growth of crystals Interfacial free energy of PP crystal surfaces ↓ with compatibilization

All blend systems listed have been prepared by melt-mixing, except when indicated with (a)

90/10 80/20 70/30 60/40

PA-6/ionomer

72/18/10

PP/LLDPE + EP blockpol. + Butyl-rubber + EPR + SEBS triblock

72/18/10

PP/LLDPE + EP blockpol. + SEBS triblock + SEBS-g-MA

75/25 50/50

* Isothermal cryst. rate of Mutual nucleation between PP and PEG both PP and PEG ↑ with compatibilizer Growth rate remains cte. Avrami exp. PP = cte. Avrami exp. PEG: 1 → 2. * Fractionated cryst. of PEG is reduced by adding the compatibilizer

composition

PP/PEG (a) + (PP-MA)-g-PEG

Explanation

Crystallization

Blend System (a) +compatibilizer Studies on the Influence of

Table 3.29. Continued.

[Willis et al., 1993]

[Plawky and Wenig, 1994]

[Flaris et al., 1993]

[Tang and Huang, 1994b]

Reference

282 G. Groeninckx, M. Vanneste and V. Everaert

Crystallization, Morphological Structure, and Melting of Polymer Blends

283

Figure 3.48. Influence of compatibilization (10 wt%) on the crystallization and melting behavior of PA-6/PP blends with various blend compositions. Compatibilizer types used were EBA-g-FA (I), PP-g-MAH (II), SEBS-g-MAH (III) and E EA-GMA (IV) [Ikkala et al., 1993].

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Figure 3.49. The effect of compatibilizers (SEBS) and (EP) on the nucleation of isotactic PP at various temperatures: (a) measured as blend volume; and (b) calculated as PP volume fraction [Long et al.,1995].

of a polymer are significantly modified by the presence of the second component even when both phases are physically separated due to their immiscibility. The presence of the second component, either in the molten or solid state can affect both nucleation and crystal growth of the crystallizing polymer. The effect of blending on the overall crystallization rate is the net combined effect on nucleation and growth. From the above literature survey, it is clear that the physical state of the second phase at the moment of crystallization is of utmost importance. The crystallization of a continuous matrix in which the dispersed phase is in the molten state can be influenced by several phenomena. One of the most important factors that play a role here is the possibility that impurities and nuclei migrate during the melt-mixing process,

hence altering the nucleation density of the components. Furthermore, the interface may enhance the nucleation, mostly due to highly ordered structures in supercooled melt droplets. It should however be mentioned that the crystal growth rate, G, is generally not affected. Only in some exceptional cases where the growing crystallizing front rejects and/or deforms finely dispersed melt droplets, a decrease of G has been reported. It can thus be concluded that the matrix always crystallizes around its bulk temperature. Migration of nuclei, nucleation effects, etc. result in a shift of the Tc,matrix by 5 to 10°C, on average. The melting behavior of the matrix remains in general unaffected. In the case of the crystallization of the matrix in the presence of already solidified or crystallizing particles, migration of nuclei still can play an important role. However, several other phenomena have to be taken into account. First of all, the solidified domains can act as efficient nucleators. Furthermore, retarded crystallization of finely dispersed droplets can nucleate the matrix and leads to coincident crystallization of both phases. Finally, it has been reported that epitaxial crystallization at the interfaces sporadically occurs. All these phenomena lead to an increased heterogeneous nucleation of the matrix phase. Although most often also here the crystal growth rate is not affected, some authors have reported that finely dispersed solidified domains can increase the melt-viscosity of the matrix in such a way that the crystallization rate becomes depressed. Again, the matrix component will crystallize around its bulk temperature. The above mentioned phenomena can eventually alter the spherulite size and shift the Tc of the matrix on average by 5 to 10°C. The melting behavior remains normally unaffected. The crystallization and melting behavior of a dispersed phase is highly different from the behavior of the continuous phase and much more sensitive for changes. Droplets crystallizing in a melt matrix can just crystallize at their bulk temperature or show shifts of their Tc as a result of migration of nuclei, as has been outlined for matrix crystallization in the melt.

Crystallization, Morphological Structure, and Melting of Polymer Blends

However, an important additional factor that plays a role here is the size of the dispersed phase. When the number of finely dispersed droplets exceeds the available heterogeneities of ‘type 1’, fractionated or even homogeneous crystallization will occur leading to shifts in the crystallization temperature by sometimes up to 100°C (as compared to the homopolymer). This can result in a change of the crystal polymorphic form, coincident crystallization with a lower crystallizing matrix component, etc. However, the melting peak in the latter case will only be slightly depressed (by 2 to 4°C) due to the formation of less perfect crystallites at lower temperatures. Additionally, it has been demonstrated that compatibilization can induce drastic changes in the blend phase morphology, and thus in the crystallization and melting behavior. In the case where dispersed droplets crystallize in an already solidified matrix, the same phenomena as in the previously described case can influence the thermal behavior of the dispersed phase. Additionally, nucleation from the already solidified matrix will play a distinguished role. An induction of heterogeneous nuclei often can reduce the fractionated crystallization or even bring the Tc back at its bulk temperature. 3.5

Nomenclature

3.5.1

Abbreviations

AN aPMMA aPS BR CPE DHDPE EBA EEA EPDM

EPR

Acrylonitrile Atactic poly(methyl methacrylate) Atactic polystyrene Butyl-rubber Chlorinated polyethylene Deuterated high density polyethylene Ethylene butylacrylate Elastomeric copolymer from ethylene and ethyl acrylate Elastomeric terpolymer from ethylene, propylene and a nonconjugated diene Elastomeric ethylene-propylene

285

copolymer Elastomeric ethylene-propylene copolymer grafted with styrene acrylonitrile EVAc Poly(ethylene-co-vinyl acetate) (random) FVA Poly(vinyl acetate-co-di-n-tetradecyl fumarate) (alternating) GMA Glycidyl methacrylate HDPE High density polyethylene iP(p-Me-S) Isotactic copolymer of styrene and p-methyl styrene iPEMA Isotactic poly(ethyl methacrylate) iPMMA Isotactic poly(methyl methacrylate) iPS Isotactic polystyrene LDPE Low density polyethylene LLDPE Linear low density polyethylene MA or MAH Maleic anhydride p(4-Me-pentene) Poly(4-methyl pentene) P(E)0.43(K)0.57 Random copolymer of phenyl ether and phenyl ketone P(iPr-vinylether) Poly(isopropyl-vinylether) P(sec-But-vinylether) Poly(sec-butyl vinylether) PA-11 Polyamide 11 PA-12 Polyamide 12 PA-6 Polyamide 6 PA-66 Polyamide 66 PAr Polyarylate PBA Poly(1,4.butylene adipate) PBT Polybutyleneterephthalate PC Bisphenol-A polycarbonate PCDS Poly(1,4-cyclohexane-dimethylene succinate) PCL Poly-ε-caprolactone PDPA Poly(2,2-dimethyl-1,3-propylene adipate) PDPS Poly(2,2-dimethyl-1,3-propylene succinate) PE Polyethylene PEA Poly(ethylene adipate) PECH Poly(epichlorohydrin) PED n-Dodecyl ester terminated poly(ethylene glycol) PEE Poly(ester-ether) segmented block copolymers PEEEK Poly(ether ether ether ketone) EPR-g-SA

286

PEEK PEEKK PEG PEI PEK PEKK PEMA Penton PET PET-b-PS Phenoxy PI PI PiB PMMA POM PP PPE PPG PPS PS PVAc PVC PVDF PVF PVME SAN SARAN

SBS SEBS SMA sPMMA sPS TR

G. Groeninckx, M. Vanneste and V. Everaert

Poly(ether ether ketone) Poly(ether ether ketone ketone) Polyethyleneglycol (also PEO) Poly(ether imid) Poly(ether ketone) Poly(ether ketone ketone) Polyethylmethacrylate Poly[3,3-bis(chloromethyl)oxetane] Polyethyleneterephthalate Block copolymer of PET and PS segments Poly(hydroxy ether of bisphenol A) Di-n-octadecyl ester of itaconic acid Polyisoprene Polyisobutene Polymethylmethacrylate Polyoxymethylene Isotactic polypropylene Poly(2,6-dimethyl 1,4-phenylene ether), tradename PPO (GE) Poly(propylene glycol) Poly(phenylenesulfide) Atactic polystyrene Poly(vinyl acetate) Polyvinyl chloride Poly(vinylidene fluoride) (sometimes expressed as PVF2) Poly(vinyl fluoride) Polyvinylmethylether Poly(styrene-co-acrylonitrile) P(VCl2-VC), P(VCl2-VA) or P(VCl2-AN) random copolymers of vinylidene chloride (VCl2) with vinyl chloride (VC), vinyl acetate (VA) and acrylonitrile (AN), respectively Elastomeric Styrene-ButadieneStyrene triblock polymer (also TR) Styrene-ethylene/butylene-styrene triblock polymer Poly(styrene-co-maleic anhydride) Syndiotactic poly(methyl methacrylate) Syndiotactic polystyrene Thermoplastic rubber (also SBS)

UHMWPE VDF-HFA VDF-TFE VLDPE 3.5.2

Notations

compat. conc. cryst. cte. DSC etc. exp. HM LCST O. M. [(polymer)] phr. SALS SAXS SEM temp. UCST WAXS WLF 3.5.3

Ultra heigh molecular weight polyethylene Copolymer of vinylidene fluoride and hexafluoro acetone Copolymer of vinylidene fluoride and tetrafluoro ethylene Very low density polyethylene

compatibilization, compatibilized, etc. concentration crystallization, crystalline, crystallize constant Differential Scanning Calorimetry etcetera exponent High molecular weight Lower Critical Solution Temperature Optical Microscopy (also OM) amount/concentration of the cited polymer parts per hundred resin Small Angle Light Scattering (also SALLS) Small Angle X-Ray Scattering Scanning Electron Microscopy temperature Upper Critical Solution Temperature Wide Angle X-Ray Scattering Williams Landel and Ferry

Symbols

3.5.3.1 Roman letters b B12 C1,C2, C3 C-2 C-3

thickness of a monomolecular crystal layer interaction density WLF constants carbon chain with 2 C-atoms; i.e. ethylene carbon chain with 3 C-atoms; i.e.

Crystallization, Morphological Structure, and Melting of Polymer Blends

Cp d

D D E1 E2 E3 E4

∆E f F12 fz(1)

∆F

∆F* G Go G1 Gc Gm° Gmb°

∆Gu

propylene heat capacity under constant pressure maximum distance over which the amorphous component has to diffuse away during crystallization coefficient of diffusion mutual diffusion coefficient energy dissipated for rejection of droplets during spherulite growth energy to overcome the inertia of droplets during spherulite growth energy required to form new interfaces when droplets are engulfed energy dissipated for deformation of occluded particles during spherulite growth activation free energy for the transport of chains through the liquid-solid interface factor taking into account the temperature dependence of ∆hv spreading coefficient fraction of dispersed droplets of volume VD that contain z heterogeneities of type 1 difference of interfacial energies; driving force for rejection, engulfing and/or deformation of dispersed droplets during spherulite growth free energy for the formation of a nucleus of critical size (∆Fbl*: the same parameter for a blend) isothermal spherulite growth rate theoretical spherulite growth rate undisturbed spherulite growth rate of the homopolymer described by the Turnbull-Fisher equation free energy of the crystalline phase free energy of the melt phase of a homopolymer free energy of the melt phase when the crystallites are surrounded by a mixed melt phase change in free energy of melting

He ∆hu

∆Hm K k k’ k1 k2 kB L m M(1) MW n n’ nnucleation ngrowth n” N N/S q r R R Se ∆Su tc t* t1/2 or t0.5 Tc

287

per mole of monomer unit surface enthalpy of folding heat of fusion per mole of monomer of the crystallizable component total melting enthalpy of the crystalline polymer fraction (also ∆Hu) overall crystallization rate Avrami constant for primary crystallization processes Avrami constant for secondary crystallization processes rate of transport of crystalline molecules through the amorphouscrystalline interfaze rate of segregation of the amorphous component Boltzman constant crystal thickness number of units in the molecule, e.g. degree of polymerzation concentration of heterogeneities of type 1 molecular weight Avrami exponent for primary crystallization processes Avrami exponent for secondary crystallization processes Avrami exponent due to nucleation Avrami exponent due to growth Ozawa exponent nucleation density nucleation density normalized per unit area heating or cooling rate (also β) local co-ordinate spherulite radius universal gas constant surface entropy of folding entropy changes on melting crystalization time time at which the pure secondary crystallization starts half time of crystallization at a fixed Tc,iso (also tn1/2) bulk crystallization temperature

288

Tco Tc,cold Tc,hom

∆Tc,hom Tc,i

∆Tc,i Tc,iso Tc,max Tg Tm Tmo T m’ Tmelt tmelt Tons vd VD vol% Viu wt% Xc yp (m,c) ypn (m) ypn (c)

G. Groeninckx, M. Vanneste and V. Everaert

upon cooling from the melt crystallization temperature of the bulk homopolymer cold crystallization temperature homogeneous crystallization temperature degree of undercooling required for homogeneous crystallization (also ∆T) crystallization temperature at which heterogeneities of type i become active degree of undercooling required before a heterogeneity of type i can become active isothermal crystallization temperature optimal isothermal crystallization temperature which yields the highest overall crystallization glass transition temperature (also To) measured melting temperature of the crystalline phase theoretical melting temperature for crystalline lamellae of infinite thickness observed melting temperature of the crystalline phase in blends premelting temperature time the polymer is kept in the melt onset temperature of crystallization rate of diffusion of the noncrystallizable component average volume of dispersed polymer droplets volume percentage molar volume of the repeating unit of component i weight percentage total degree of crystallinity lateral surface free energy between the crystal and its own melt interfacial energy between the nucleating species and the polymer melt interfacial energy between

∆yi ∆ypn z

the nucleating species and the polymer crystal specific interfacial energy difference between a nucleating species of type i and the polymer specific interfacial energy difference between a nucleating species and the polymer number of heterogeneities of type 1, inducing crystallization in the bulk polymer at Tco

3.5.3.2 Greek letters

α αp αs αpeak β ξ

δ φο φb φ2

φ γPS γPL σ0 σσe λ σ1,2 σi,1

fraction of crystallinity at time t fraction of crystallinity due to primary crystallization fraction of crystallinity due to secondary crystallization extent of crystallization at the peak maximum of the DSC exotherm heating rate (cooling rate = -β) weight fraction of the polymer crystallized by both the primary and the secondary crystallization at the moment the primary crystallization has ended Keith and Padden parameter start composition binodal composition volume fraction of the crystallizable component stability parameter (value between 0 and 1) interfacial free energy between the crystallizing solid and the inclusions interfacial free energy between the liquid polymer melt and the inclusions surface free energy op folding product of the lateral and fold surface free energy wavelength interfacial free energy between two phases of a blend in the melt interfacial free energy of an impurity

Crystallization, Morphological Structure, and Melting of Polymer Blends

σi,2 θm χ12

3.6

with respect to melt phase 1 interfacial free energy of an impurity with respect to melt phase 2 azimuthal angle of maximum intensity polymer-polymer interaction parameter

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Santana, O. O., and Müller, A. J., Polymer Bull., 32, 471 (1994). Schouterden, P., Groeninckx, G., Reynaers, H., Riekel, C., and Koch, M. H. J., Polym. Bull., 13, 533 (1985). Sham, C. K., Guerra, G., Karasz, F. E., and MacKnight, W. J., Polymer, 29, 1016 (1988). Shibanov, Y. D., and Godovsky, Y. K., Makromol. Chem., Macromol. Symp., 44, 61 (1991). Shibanov, Y. D., and Godovsky, Y. K., Polymer Science USSR, A25, 395 (1983); Shibanov, Y. D., and Godovsky, Y. K., Progr. Colloid Polym. Sci., 80, 110 (1989). Shingankuli, V. L., Jog, J. P., and Nadkarni, V. M., J. Appl. Polym. Sci., 36, 335 (1988). Shingankuli, V. L., PhD. thesis, Bombay University, India (1990). Siegmann, A., and Nir, Y., Polym. Eng. Sci., 27, 1182 (1987). Silvestre, C., Cimmino, S., Martuscelli, E., Karasz, F. E., and MacKnight, W. J., Polymer, 28, 1190 (1987a). Silvestre, C., Karasz, F. E., MacKnight, W. J., and Martuscelli, E., Eur. Polym. J., 23, 745 (1987b). Song, H. H., Wu, D.-Q., Ree, M., Stein, R. S, Phillips, J. C., LeGrand, L. and Chu, B., Macromolecules, 21, 1180 (1988) Starkweather, H. W., Jr., J. Appl. Polym. Sci., 25, 139 (1980). Stein, R. S., in ‘Newer Methods of Polymer Characterization’, Chapter 4, Ed. by Wiley, New York (1964). Stein, R. S., Khambatta, F. B., Warner, F. P., Russell, T., Escala, A., and Balizer, E., J. Polym. Sci., Polym. Symp., 63, 313 (1978). Stein, R. S., Russell, T. P., Morra B. S., Wai M., and Gilmer J., in ‘Structural Order in Polymers’, Ciardelli, F., and Giusti, P. (Eds.), Pergamon Press, New York, p 195 (1981). Suzuki, T., and Kovacs, A., Polymer J., 1, 82 (1970). Tanaka, H., Nishi, T., Phys. Rev. Letters, 55, 1102 (1985). Tanaka, H., Nishi, T., Phys. Rev. (A), 39, 783 (1989). Tang, T., and Huang, B., J. Appl. Polym. Sci., 53, 355 (1994a). Tang, T., and Huang, B., J. Polym. Sci., Polym. Phys. Ed., B32, 1991 (1994b). Tang, T., Li, H., and Huang, B., Macromol. Chem. Phys., 195, 2931 (1994). Tashiro, K., Izuchi, M., Kaneuchi, F., Jin, C., Kobayashi, M., and Stein, R. S., Macromolecules, 27, 1140 (1994). Tashiro, K., Izuchi, M., Kobayashi, M., Stein, R. S., Macromolecules, 27, 1121 (1994a); ibid., 27, 1128 (1994b); ibid., 27, 1134 (1994c). Tashiro, K., Satkowski, M. M., Stein, R. S., Li, Y., Chu, B., and Hsu, S. L., Macromolecules, 25, 1809 (1992). Tashiro, K., Stein, R. S., and Hsu, S. L., Macromolecules,

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CHAPTER 4 INTERPHASE AND COMPATIBILIZATION BY ADDITION OF A COMPATIBILIZER

Abdellah Ajji

National Research Council Canada, Industrial Materials Institute, Boucherville, QC, Canada

4.1

Introduction

4.1.1

General

Polymer blends are mixtures of at least two macromolecular species, polymers and/or copolymers. For practical reasons, the name blend is given to a system only when the minor component content exceeds 2 wt%. Depending on the sign of the free energy of mixing, blends are either miscible or immiscible. In a general sense, the polymer/polymer miscibility does not exist — it is always limited to a “miscibility window”, a range of independent variables, such as composition, molecular weight, temperature, pressure, etc. More than 1600 of these “miscibility windows” have been identified for two, three or four components’ blends. The immiscibility dominates the field [Utracki, 1989]. For more details on the thermodynamics of mixing and phase diagrams, the reader is referred to Chapter 2. Thermodynamics of Polymer Blends in this volume. 4.1.2

Definition of the Surface and Interface Tension Coefficients

The surface tension is the reversible work required to create a unit surface area at constant temperature (T), pressure (P) and composition (n) [Wu, 1982]: νi = (∂G / δA)T, P, n

(4.1)

where νi is the surface tension coefficient of the substance i, G is Gibbs’ free energy of the system, and A is the surface area. In immiscible liquids, interactions between components are located at the physical boundary creating the interface. The energy required to reversibly separate the two liquids is expressed as the work of adhesion: W= ν1 + ν2 - ν12

(4.2)

where ν1 and ν2 are surface tension coefficients of neat components and ν12 is the interfacial tension coefficient between the liquids 1 and 2.

L.A. Utracki (Ed.), Polymer Blends Handbook, 295-338. © 2003 Kluwer Academic Publishers. Printed in the Netherlands.

296

4.1.3

Abdellah Ajji

Importance of the Interfacial Properties in Polymer Blends

The structure and morphology of immiscible blends depends on many factors among which the flow history and the interfacial properties are the most important. At high dilution, and at low flow rates the morphology of polymer blends is controlled by three dimensionless microrheological parameters: (i) the viscosity ratio, λ = η1 / η2, where η1 is the viscosity of the dispersed liquid and η2 that of the matrix; (ii) the capillarity number, κ = σ12 d / ν12 , where σ12, and d are respectively the shear stress, and the drop diameter; and (iii) the reduced time, t* = t y / κ, where t is deformation time, and y is the rate of shear [Utracki, 1989, 1994]. Thus, the interfacial and rheological properties are the keys for the morphology development in polymer blends, which in turn is the controlling factor for their performance. To improve performance of immiscible blends, usually they need to be compatibilized. There are three aspects of compatibilization: (1) Reduction of the interfacial tension that facilitates fine dispersion, (2) Stabilization of morphology against its destructive modification during the subsequent high stress and strain processing (e.g., during the injection molding), and (3) Enhancement of adhesion between phases in the solid state, facilitating the stress transfer, hence improving the mechanical properties of the product. Compatibilization can be carried out either by adding a compatibilizer to a polymer blend, or prepared during the reactive processing or blending. During the latter process the compatibilizing species are chemically formed in situ, directly across the interface. In this chapter, compatibilization of polymer blends by means of addition of a compatibilizer will be discussed. First, the theories will be summarized of the: (i) interface, (ii) interphase, and (iii) compatibilization process. This brief summary is to provide a general framework for understanding the phenomena associated with compatibilization, and guidance for optimization of the process to gain maximum performance.

The theoretical part is followed by the experimental part, where the methods for the determination of interfacial properties are presented. Most of the chapter is dedicated to provide comprehensive information on the characteristic properties of blends compatibilized by addition of a compatibilizing agent.

4.2

Theoretical Aspects of the Interface

4.2.1

Binary Immiscible Polymer Blends

Mixing two polymers usually results in an immiscible system, characterized by a coarse, easy to alter morphology, and poor adhesion between the phases. These blends have large size domains of dispersed phase and poor adhesion between them. As a result, their performance is poor and irreproducible. In particular the impact strength, maximum strain at break, and the yield strength are affected. The irreproducibility originates from instability of morphology — blends’ structure developed during the compounding step is unstable and irreproducible. To be able to solve these three problems (degree of dispersion, stability of morphology, and adhesion between the phases in solid state) one must learn about the region between the two phases in binary polymer blends, the interface or rather the interphase. Let us consider a molten, immiscible, binary blend of polymers A and B, without compatibilizer. Helfand and Tagami [1971], Helfand [1975], Roe [1975], and Helfand and Sapse [1975] have developed a quantitative lattice theory of the interphase that twenty years later still provides good basis for understanding. Helfand and Tagami model is based on selfconsistent field that determines the configurational statistics of macromolecules in the interfacial region. At the interface, the interactions between statistic segments of polymers A and B are determined by the thermodynamic binary interaction parameter, χ12. Since the polymers are immiscible, there are “repulsive” enthalpic effects that must be balanced by the entropic ones that cause chains A and B to intermingle.

297

Interphase and Compatibilization by Addition of a Compatibilizer

In the first simplified mean-field approach: (i) the two homopolymers were assumed to have the same degree of polymerization, (ii) the complex set of equations derived for the segmental density profile, ρi (where i = A or B), was solved for infinitely long macromolecules, Mw → ’, (iii) the isothermal compressibility was assumed to be negligibly low, and (iv) there was no volume change upon blending (i.e., the attractive or repulsive forces between two polymers were assumed weak). The analytical solution of the interfacial composition profile was found to follow an exponential decay function [Helfand and Tagami, 1972]: ρi = ρi,O y2 / (1 + y2)

(4.3)

y ≡ exp {(6χAB)1/2 (x/b)}

n 12 (mN/m)

1

0.8

0.6

2

r/ r = y /(1

2

+

y ) ;

1 / 2

(x/b)}

o

y = exp {(6c)

0.4

Dl

0.2

0 -10

-6

-2

2

6

10

Interfacial Tension Coefficient,

SEGMENTAL DENSITY, r / r

o

where b is the lattice parameter, and i can be either A or B. A typical dependence is shown in Figure 4.1 [Utracki, 1994].

where kB is the Boltzmannn constant and T is the absolute temperature. Predictions of Eq 4.5 were compared with the experimental data for three polymeric blends: polystyrene/polymethylmethacrylate, PS/PMMA, polybutylmethacrylate/polyvinylacetate, PBMA/ PVAc, and PMMA/PBMA. It was found that ν’ agreed with the experimental value of ν12 determined for PBMA/PVAc. However, the agreement for PMMA/PBMA was not as good, while for PS/PMMA the difference was 50%. In consequence, the authors postulated that for large values of χ12 the thickness of the interphase is too small for the mean-field theory to be valid. Eq 4.5 also predicts that the interfacial tension coefficient is a linear function of the temperature. Furthermore, since to the first approximation: χAB = a + b/T, the slope: dν12 / dT should be negative. Figure 4.2 indeed confirms these predictions.

PE/PA-6

10

PP/PS115

6

PP/PS86.4 PP/PS19.4 LDPE/PS PP/PS4.8

PP/PS380 PVDF/EP

2 160

200

240

280

Temperature, T(°C)

DISTANCE ACROSS THE INTERFACE, x/b

Figure 4.1. Representation of the interface, with the definition of the interphase thickness, ∆l; χ and b are respectively the binary interaction and the lattice parameters [Helfand and Tagami, 1971].

Figure 4.2. Temperature dependence of the interfacial tension coefficient for PP/PS (data from Kamal et al., 1994) and for PE/PA, LDPE/PS, and PVDF/EP [Luciani et al., 1996].

Similarly, for the symmetrical polymers A and B whose Mw → ’, the interfacial thickness, ∆l’, and the interfacial tension coefficient, ν’, were derived as:

The Helfand-Tagami lattice theory predicts that there is reciprocity between the interfacial tension coefficient and the interfacial thickness, and the product, ν’ ∆1’, is independent of the thermodynamic binary interaction parameter, χ12. Furthermore, the theory led to the conclusions that: (i) the surface free energy is proportional to χ121/2 ; (ii) the chain-ends of both polymers con-

∆l’ = 2b(6χAB)-1/2

(4.4)

ν’ = bρTkB (χAB / 6)1/2

(4.5)

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Abdellah Ajji

centrate at the interface; (iii) any low molecular weight third component is repulsed to the interface; (iv) that the interfacial tension coefficient is a linear function of temperature (see Eq 4.5 and Figure 4.2), and (v) the interfacial tension coefficient increases with molecular weight to an asymptotic value, ν’ (see Figure 4.3): ν = ν’ - aoMn-2/3

(4.6)

Four years later, Helfand and Sapse removed the restriction of the original theory for the symmetric character of both polymers, obtaining for ν’ the following expression: ν’ = (2 / 3) kb Tα1/2 (β3A - β3B) / (β2A - β2B) α ≡ χAB (ρoAρoB)1/2

(4.7)

βi2 = ρoibi2 / 6

Interfacial Tension Coefficient,

n (mNm) 12

where bi2 = 〈Ri2〉 / Zi , and bi is the Kuhn statistical segment length. Substituting in Eq 4.7: ρoA = ρoB and βA = βB results in recovery of Eq 4.5.

5

3

n2 = noo - aM

-2/3 n

1 7

9

11

ln (M ) n

Figure 4.3. Verification of the molecular weight dependence of the interfacial tension coefficient, as predicted by the Helfand-Tagami theory (see Eq 4.6)

A generalized gradient theory of the interface was developed by Anastasiadis [1988]. The approach is based on the assumption that the composition

gradient is small compared to the reciprocity of the intermolecular distances. Under these circumstances the free energy density, g, can be written as a power series, truncated after the square term. In essence, the theory determines the difference in the density fluctuation per unit interfacial area between polymer mixture and a system in which the properties are homogenous. The theory predicts that: φβ

∫ κ ≡ −[∂ ν12 = _

φα 2

_  κ ∆g(φ )

1/ 2

dφ

] [

g / ∂φ∂∇ 2φ + ∂ 2 g / ∂ ∇φ

2

]

(4.8)

Thus, to calculate ν12 one has to select an appropriate expression for the energy density gradient and then integrate Eq 4.8 within the limit of composition in both phases, φA and φB. Huggins-Flory, and Cahn-Hilliard theories were used with good success to predict the temperature gradient, but poor as far as the effects of molecular weight were concerned. Roe [1975] developed a quasi-crystalline lattice model for conditions where χ12 >> χcr (where χcr is the critical value of the interaction parameter at the phase separation) and for χ12 - χcr << χ2cr. Under the first conditions (high immiscibility) the theory predicted a proportionality between ν12 and χ12, whereas under the second (near the phase sep3/ 4 aration) a proportionality between ν12 and χ12 was predicted. By contrast with the previously summarized Helfand and Tagami predictions, Roe’s theory indicates that the product ν12∆l 1/ 2 should be proportional to χ12 . Kammer [1977] considered the interface between two polymers from the basic thermodynamic point of view. He derived a simple relation: ν12 = ∆µB/S, where ∆µB is the excess chemical potential of polymer B in the mixture, and S is the molar area of the interface. Near the spinodal decomposition, using Cahn-Hilliard gradient theory, he calculated:

ν12 = ν o + RTφ 1B/ 2 ε 1/ 2 ε ≡ ( χ − χ cr ) / χ cr

(4.9)

Interphase and Compatibilization by Addition of a Compatibilizer

The following year, from similar assumptions, Joanny and Leibler [1978] derived:

ν12 = (2 / 3)k B Tb 2 Z B−1/ 2 ε 3 / 2 ∆l = (1 / 3)bZ 1B/ 2 ε −1/ 2

(4.10)

299

multi-block, linear, star-shaped, etc. Those based on styrene and butadiene, SB or SBS, are the earliest to be applied and studied, as well as the largest as far as the volume of production is concerned [Holden et al., 1967].

∴ν∆l = (2 k B T / 9b)ε

4.2.2.1 Block Copolymers — Fundamentals

where ZB = ZA is the degree of polymerization. The proportionality ν12 ~ ε3/2 and ∆l ~ ε -1/2 are in agreement with the classical approach to the interface [de Gennes, 1977]. Thus, it seems that prediction of Helfand-Tagami lattice theory that the product, ν12∆l, is independent of the thermodynamic binary interaction parameter, χ12, may not be correct in the whole range of variables. The authors observed that the relation breaks down in antagonistically immiscible blends where χ12 >> χcr. Nowadays, there is growing theoretical evidence that the prediction is not valid for χ12 → χcr. More recently, Eq 4.4 was modified to account for finite Mw [Broseta et al., 1990]. For blends of polystyrene with polyvinylpyridine, PS/PVP, the authors obtained:

To ascertain control of the molecular weight, structure and composition, block copolymers are usually synthesized in anionic polymerization. The block copolymers of commercial interest are specifically prepared from monomers that upon polymerization yield immiscible macromolecular blocks, a smaller one rigid and the other flexible. The rigid blocks form physical crosslinks that upon heating above the transition point make the copolymer to flow. Thus, these materials belong to the growing family of thermoplastic elastomers. There are two distinct differences between the phase diagram of a block copolymer and the one obtained for a mixture of two homopolymers. In block copolymers, owing to the chemical links between blocks, micro-domains instead of macroscopic phases are observed. The size of micro-domains can be controlled by varying the molecular weight and composition. Furthermore, since the type of morphology depends on the concentration as well as on the transition temperatures of the individual phases, the phase diagram of block copolymers shows comparable complexity to those of metallic alloys. The domain size and shape, as well as the interfacial thickness, depend on the following factors: (i) magnitude of the repulsive interactions between the A and B blocks, χAB; (ii) the conformation entropy loss necessary to maintain constant segment density; (iii) the localization entropy loss that causes the chemical links to be present at the interface; and (iv) the composition. These mutually compensating factors (i vs. ii + iii) depend on molecular weight of each block and the binary interaction parameter [Helfand, 1975; Helfand and Wasserman, 1976, 1978, 1980; Hashimoto et al., 1980; Inoue et al., 1969; Krause, 1980; Meir, 1969, 1987; Hashimoto et al., 1993].

¨l = 2b(6χ12)-1/2 [1+(ln 2)(1/ZPS + 1/ZPVP) / χ12] (4.11) where Z is the degree of polymerization. Eq 4.11 was corrected for entropic effects. For Mw → ’ Eq 4.4 is recovered. Polydispersed systems were also considered by assuming the same bimodal distribution of chain lengths for both A and B components. As expected, chains of higher Mw tended to localize in the bulk phases rather than near the interfacial region. The loss of translational entropy associated with the confinement of the chain into half the space is more easily accepted by large chains than by small ones. 4.2.2

Copolymers

Block copolymers are polymers constituted of at least two different monomers arranged in a specific manner — they could be di-block, tri-block,

300

Abdellah Ajji

Theories of block copolymers are usually complex, involving computation of the domain size, the interphase thickness between the blocks, the structure, and the order-disorder transitions. Helfand and Wasserman [1976, 1978, 1980] using the narrow interphase approximation, showed that Eq 4.4 is valid in the limit of infinitely immiscible blocks having Mw → ’ (i.e., the strong segregation limit, SSL). The authors’ approach was based on the confined chain statistics. Expressions for the free energy of different structure formation, interfacial tension, and interfacial thickness (in the case of an infinite Mw) were derived. For large χ12Zc values (Zc is the copolymer degree of polymerization), the narrow interface approximation was assumed valid, thus the boundary thickness should be similar to that in an A/B mixture. This interfacial thickness ∆l was expressed as:

[(

)

∆l = 2 β A2 + β B2 / 2 χ12

1/ 2

]

(4.12)

2 2 with: β i2 = ( ρoi bi2 ) / 6 and bi = < Ri > / Zi , where bi is the Kuhn statistical segment length, Zi is the degree of polymerization, ρoi is the density and is the radius of gyration of the block i. For identical chains and the lattice size, b2 = bi2ρi and Eq 4.12 converts to Eq 4.4. The thermodynamic properties of block copolymers in disordered state, have been studied by Leibler [1980]. Using the random phase approximation [de Gennes, 1979], the author developed a relation between the segmental density correlation function and the scattering vector. An order parameter, related to the reduced segmental density, was introduced. In the disordered state, this order parameter is zero whereas for the ordered phase, it is a periodic non zero-function. Leibler’s demonstrated that the critical condition for microphase separation in di-block copolymers is: χABZc = 10.5. The work was extended by Olvera de la Cruz and Sanchez [1986] to block copolymers with more complex architectures. For di-block copolymers the authors confirmed Leibler’s results,

χABZc = 10.5. The same value was also found for the star copolymers having the same number of arms A and B arms, each of the same length: ZA = ZB. Computations for tri-block copolymers gave χABZc = 8.86, indicating poorer miscibility than that computed for di-blocks. Interphase thickness, D l (nm)

4.2.2.2 Interphase in Block Copolymers

4 Equation 4.13

3

2

1

0

40

80

120

Block copolymer Molecular Weight, Mw (kg/mol)

Figure 4.4. Interphase thickness in styrene-isoprene block copolymers vs. total molecular weight. Points are experimental [Hashimoto et al., 1980; Richars and Thomason, 1983], the line was computed from Eq 4.13 using ∆lo = 1.9 nm and χ = 1000 χAB/Mo = 0.6.

More recently, a theory based on confined chain statistics (CCS) predicted that ∆l should be a decreasing function of Zc. Thus, modification of Eq 4.4 resulted in an expression valid for χABZc • 20 [Spontak and Zielinski, 1993]: ∆l = ∆l∞ [1 − (8 ln 2) / ( χ AB Zc )]

−1/ 2

(4.13)

Eqs 4.12 and 4.13 are similar, but differences in the range of 40% have been found in the low χABZc region. Prediction of Eq 4.13 is compared in Figure 4.4 with the small angle neutron scattering and small angle X-ray scattering data [Hashimoto et al., 1980; Richars and Thomason, 1983]. 4.2.3

Copolymer/homopolymer Blends

4.2.3.1 Blends of Block Copolymer with Homopolymer Information on the phase diagrams of copolymer/ homopolymer blends can be found in reviews by

301

Interphase and Compatibilization by Addition of a Compatibilizer

Ajji and Utracki [1996, 1997] and in Chapter 2 Thermodynamics of Polymer Blends in this handbook. The interface thickness of A/A-B mixtures was not theoretically calculated, but experimental measurements indicate that the presence of the homopolymer leaves the domain boundary unchanged [Bates et al., 1983; Hashimoto et al., 1990; Tanaka et al., 1991; Zin and Roe, 1984; Green et al., 1989]. On the other hand, the part of the phase diagram where the concentration of the block copolymer in the mixture is low, was studied in detail [Whitmore and Noolandi, 1985b; Leibler et al., 1983; Leibler, 1988]. The proposed models were similar. In both, the conditions for the formation of spherical micelles were investigated and expressions of the critical micelle conspherical centration, φ CMC , were derived. For example [Leibler, 1988]:

{

spherical φ CMC = exp 1.72( χ AB ZC )1/ 3 f 4 / 9 (1.74 f −1/ 3 − 1)1/ 3

− fχ AB ZC

}

(4.14) where f is the fraction of A-polymer in A-B copolymer. The predictions of φCMC were extended to cylindrical and lamellar micelle morphologies [Shull et al., 1991]. In all cases, geometric parameters, e.g., core radius and corona thickness, can be computed, assuming that the interfacial thickness between the core and the corona is equal to ∆l’. Experimental studies of micellar systems were carried out using scattering methods [Selb et al., 1983; Rigby and Roe, 1984, 1986; Kinning et al., 1990, 1991]. Theoretical simulations of the scattering curves have been based on the assumptions that either an infinitely sharp boundary thickness or a diffuse interfacial thickness is equal to ∆l’. In spite of these seemingly diverse principles the simulations were reasonably correct. The enthalpy associated with the addition of a homopolymer having high Mw to a block copolymer in the disordered state is not compensated by sufficient entropic gains, hence a mesophase is formed. This effect was observed by Cohen and

Torradas [1984]. This contrasts with other reports [Ptaszynski et al., 1975; Hashimoto et al., 1990; Winey et al., 1991] where upon addition of a homopolymer with molecular weight smaller than that of the corresponding block, the average area per junction point was found to increase. SAXS scattering patterns obtained from samples having lamellar morphology followed an idealized model, in which layers of styrene and butadiene (with randomly varying thickness and a diffuse interface) were parallel. Green and Russel [1991] theoretically and experimentally studied effects of addition of a low molecular weight symmetric block copolymer of P(S-b-MMA) on the interfacial tension in blends comprising a high molecular weight homopolymer, either PS or PMMA. The theory was based on a mean field argument (similar to that of Leibler), valid within the low compatibilizer concentration range. The theory predicted that:

[

]

ν = ν o − k B T ρφ C a exp {χ AB ZC } / ZC

(4.15)

where φc is the volume fraction of the copolymer. The calculations were consistent with experimental data on similar systems. The estimated value of χΑΒ compared well with that determined in neutron scattering experiments. 4.2.3.2 Blends of Graft Copolymer with Homopolymer There are only few papers dealing with systems comprising a graft copolymer and a homopolymer. For example, blends of graft A-B copolymer with homopolymer A (identical with the backbone of the copolymer), were found to have unusual morphologies [Eastmond et al., 1977; Eastmond and Phillips, 1979]. The most common of these were spherical structures (named onions) consisting of alternating concentric layers of A and B components. The conclusion was that the onions had their origin in the immiscibility of copolymer with homopolymer, even when the molecular weight of chemically identical blocks were comparable [Eastmond et al., 1987]. Other studies reported similar findings, even when the molecu-

302

Abdellah Ajji

lar weight of homopolymer was lower than that of the corresponding block [Jiang et al., 1985]. It was suggested that the immiscibility and the onion-type morphology originated in molecular structure of copolymer. 4.2.4

Blends of Two Homopolymers with a Compatibilizer

4.2.4.1 Blends of Two Polymers with a Copolymer Ternary blends that comprise two immiscible polymers and a copolymer are of a particular interest. They not only represent an ideal model for studying compatibilization of polymer blends, but also they have found direct commercial applications. Phase diagram information can be found in reviews by Ajji and Utracki [1996, 1997] and in Chapter 2 in this Handbook. The theories discussed in Part 4.2.1 dealt with binary polymer blends without compatibilizer. Hong and Noolandi [1980, 1981a, 1981b] developed a theory, similar to that of Helfand, for the interfacial region in three component polymeric systems comprising polymers A and B, and either a co-solvent or a block copolymer. The theory is based on the lattice model — it uses the mean-field approximation. It is formulated using reduced equation of state variables. Finite molecular weights and conformational entropy effects were considered, but the excluded volume effects were not (this aspect was later treated by Broseta et al. [1987]). The resulting system of equations can be solved numerically for the interfacial composition profile, interfacial tension coefficient, thickness of the interphase, etc. At low values of Mw the theory well predicts ν12, but for higher Mw the prediction was up to 20% too high. 4.2.4.1.1 Calculations of the Interfacial Tension Coefficient and CMC The emulsifying effects of a small quantity of a block copolymer, A-B, added to immiscible blend of homopolymers A and B were examined by Leibler [1988]. The theory predicted reduction

of the interfacial tension coefficient, ν12, caused by equilibrium adsorption of a copolymer at the interface. For well chosen compositions and molecular weights of the copolymer, low values of ν12 are to be expected. This suggest a possible existence of thermodynamically controlled stable droplet phase, in which the minor phase homopolymer drops are protected by an interfacial film of the copolymer, interfacing the matrix polymer. The size distribution of the droplets is expected to depend on the rigidity and spontaneous radius of curvature of the interfacial film that can be controlled by molecular structure of the copolymer. For long copolymer chains (in strongly immiscible, or the “wet brush case”), the reduction of the interfacial tension coefficient should follow the relation: ∆ν ≡ ν − ν o = −( k B T / a 2 )(3 / 4)1/ 3 ( Σ / a 2 ) −5 / 3 [ ZCA Z A−2 / 3 + ZCB Z B−2 / 3 ] (4.16) where ZCA and ZCB are respectively the number of A and B monomeric segments in copolymer, Zc = ZCA + ZCB is the total number of segments in copolymer, ZA and ZB are the degree of polymerization in homopolymer A and B respectively, a is the monomer length, Σ the interfacial area per copolymer joint (Σ = A/Q, where Q is the number of copolymer chains at the interface between A and B, and A is the total interface surface area). For short copolymer chains (Zc < Zi-2/3 and Σ < Zc1/2a2), Eq 4.16 can be simplified to read: ∆ν = −3ZC ( k B T / a 2 )( Σ / a 2 ) −3

(4.17)

Both Eqs 4.16 and 4.17 predict that when adsorption density (Σ/a2) is high, the interfacial tension coefficient is low. For the same surface area per chain, longer copolymer chains are predicted to be more efficient. The expressions of Σ/a2 can be obtained, for both “wet” and “dry” brushes, as a function of the copolymer chemical potential, µ. The ratio was found to depend on

303

Interphase and Compatibilization by Addition of a Compatibilizer

the logarithm of the copolymer concentration [Leibler, 1988]. From Eq 4.14, the critical micelle concentration, CMC, was derived: + φ CMC = exp{µ CMC − fχ AB ZC }

(4.18)

+

where φ CMC is the volume fraction of copolymer in B-rich phase, and µCMC is its chemical potential at CMC:

µ CMC = (3 / 2) 4 / 3 f 4 / 9 (1.74 f −1/ 3 − 1)1/ 3 ( χ AB ZC )1/ 3 (4.19) As copolymers are added ν12 decreases until + φ+ reaches the φ CMC value, at which the limiting value of the interfacial tension coefficient, νCMC, is obtained. A question arises whether the copolymer can saturate the interface so that: ∆ν = νo = (kBT/a2) (χ/6)1/2, and effectively the interfacial tension vanishes. This saturation may occur when µ = µsat = (3/42/3)(χN)1/3. For a flat interface, ν12 should vanish when f > 0.31. The theory also predicts that for f < 0.31 spherical micelles may appear. In the latter case, the limiting reduction of ν12 is given by: ∆ν / ν o = −31/ 2 f 2 / 3 (1.74 f −1/ 3 − 1)1/ 2

(4.20)

The above expressions are valid for flat interface. For interface with a curvature R, e.g., for droplets having a radius R >> ∆l, the following expressions were derived (dry film) [Leibler, 1988]:

ν ( µ, R) = ν ( µ ) − [ K ( µ ) / Rψ ( µ )] + K ( µ ) / 2 R 2 (4.23) ∆l = ZC a 2 / Σ (4.24) The droplet size is given by:

{ } }[1 − (1 − ν ( µ )ψ 2 ( µ ) / K ( µ ))1/ 2 ]

R = K ( µ ) / [ν ( µ )ψ ( µ )]

(4.25) For copolymer content close to the saturation of a flat interface, the R-value becomes: R = K ( µ ) / [ν ( µ )ψ ( µ )] = (3 / 2)ψ ( µ )

(4.26)

The limiting value is obtained for the saturation value of µ. In Eqs 4.23-4.26:

[

]

K ( µ ) = 2 Σ −5 7β ( f ) − 32ε 2 ZC3 1/ 2

2

Σ = a (3 / 2

1/ 2

)( ZC / µ )

[

ψ ( µ ) = ( L / 3ε ) 7β ( f ) − 32ε 2

]

(4.27)

3

Typically, for f = 0.33, 0.2 and 0.15, ∆ν/νo is expected to be 1.00, 0.83 and 0.42 respectively, hence, ∆ν = νo for f = 1/3. Thus, the more asymmetric is the chain, the less it is efficient as an interfacial agent. Since the theory does not differentiate between polymers A and B, evidently the most efficient copolymer composition must be: 1/3 < fopt < 2/3, or fopt = 1/2. For more symmetric chains f > 0.31, ν = 0 can be reached when:

{

}

+ φ sat = exp (3 / 4 2 / 3 )( χ AB ZC )1/ 3 − fχ AB ZC (4.21)

For a flat interface saturated by the copolymer, the copolymer film thickness was calculated as:

L = ZC a / Σ

β ( f ) = f 3 + (1 − f )3 ε = (1 / 2) − f Equations 4.23-4.26 are valid at equilibrium, without taking into account the kinetic effects. The latter effects may be important when comparing theoretical predictions with experimental data. There are other theories of the interface, some of which lead to different dependencies. For example, Noolandi [1984; 1985] considered a binary polymer system compatibilized by addition of a di-block copolymer. For χΑΒZcφp ” 2, he derived:

{

ν12 = ν o + ∆lφ C χ ABφ P / 2 + (1 / ZC ) ×

[1 − exp( χ AB ZCφ P / 2)]}

∆l = a ZC2 / 3 (ν o a 2 / 3k B T )1/ 3 (4.22)

∴ ∆ν12 ≅ − exp( χ AB ZCφ P / 2)

(4.28)

304

Abdellah Ajji

Where ao is a numerical parameter, while φc, φp, and Zc are respectively volume fraction of copolymer, polymer, and degree of polymerization of the copolymer. It is noteworthy that Leibler’s (see Eqs 4.20 and 4.22) and Noolandi’s theories (see Eq 4.28) predict that the product ν12∆l depends on the binary interaction parameter χ12. Thus, the reciprocity between ν12 and ∆l predicted by Helfand and Tagami for binary systems is not expected to exist in compatibilized binary blends. Two semi-empirical relations between the interfacial tension coefficient and compatibilizer concentration were derived. The first was obtained assuming an analogy between addition of block copolymer to a polymer blend and titration of an emulsion with surfactant [Utracki and Shi, 1992]:

ν12 = (φν CMC + φ meanν o ) / (φ + φ mean ) ν CMC ≡ ν12 (φ C = φ CMC )

(4.29)

φ mean ≡ (φ CMC + φ o ) / 2 The other was proposed by Tang and Huang [1994]:

ν = ν CMC + (ν o − ν CMC ) exp{− Kφ}

(4.30)

The latter authors proposed a similar relation for the particle’s radius: R = RCMC + ( Ro − RCMC ) exp{− Kφ}

(4.31)

In Eqs 4.30 and 4.31 K is treated as an adjustable parameter. Its functional dependence can be deduced from comparison with Noolandi’s Eq 4.28, since K ∝ χAB Zc. These dependencies were found useful in describing the experimental data. 4.2.4.1.2 Blends with an Arbitrary X-Y Block Copolymer Vilgis and Noolandi [1988] investigated, by means of statistical thermodynamics, the use of an arbitrary block copolymer X-Y in A/B blends. The aim was to predict ν12, ∆l, and the concentration profile across the interphase

of the blocks. This was achieved by generalizing the theory of the interfacial properties of immiscible polymer blends in the presence of a block copolymer. The diffusion equations for the density profiles were solved numerically. Despite the use of chemically different blocks, their addition resulted in reduction of ν12, increasing as Mw of the blocks increased. The computations were performed assuming different values of the binary interaction parameters and degree of polymerization. Strong localization effects of the XY emulsifier were observed when the interactions between the blocks and the homopolymers were increased. Thus the competitive interactions of the blocks with different homopolymers were shown to promote strong interfacial activity. As a result, a selective orientation of the XY blocks in the A/B polymer mixture is expected to be similar to that computed for the A/A-B/B system. A simplified analytical calculations for the case where the interaction parameters obey the assumed relationships, led to the following relation for the reduction of the interfacial tension coefficient in A/B blend upon addition of copolymer XY:

{

}

∆ν = −(1 / ZC ) exp ZC [( χ1 / 2) + χ 2 ]

χ1 ≡ χ BY = χ AX = χ XY = χ AB χ 2 ≡ χ AY = χ BX > χ1

(4.32)

This suggests that it is possible to design an universal compatibilizer operating on the principle of competitive repulsive interactions between the homopolymers and the different blocks of the copolymer. 4.2.4.1.3 Critical Compatibilizer Concentration Mathos [1993] showed that the critical concentration of interfacial agents is directly related to the interfacial area of the dispersed phase, thus related to interface saturation. The chemical structure played an important role in the emulsification ability of copolymers. Many block or graft copolymers were selected such that their segments were identical to those of the homopolymers. Alternatively, the blocks could be chemically

305

Interphase and Compatibilization by Addition of a Compatibilizer

different but miscible each with different homopolymer. However, complete miscibility of all blocks in a single phase should be avoided. The morphological analysis showed that di-block copolymers of the type polystyrene-b-hydrogenated-polybutadiene, P(HB-b-S), have a higher interfacial activity than triblock or graft copolymers. It is possible that owing to the steric restriction at the interface, graft and tri-block copolymers, form micelles in the homopolymeric phases. This adds to the complexity of the system morphology without fulfilling the basic functions of a copolymer. The di-block copolymer more readily interacts with the two homopolymeric phases, forming appropriate entanglements that result in reduction of the interfacial tension coefficient, and enhanced interphasial adhesion in the solid state. These observations have been confirmed by measuring the mechanical properties. In blends prepared in solution, Mw of the homopolymers should be lower than those of the corresponding blocks. This rule seemed to be less critical for blends prepared in the molten state [Fayt et al., 1986a, 1986b, 1989b] — only one copolymer was reported to follow this rule [Fayt et al., 1989a]. The amount of the interfacial agent required to saturate the interface, wcr, is related to its Mw, the total surface area of the interface and the specific cross sectional area of the copolymer macromolecule, a. Paul and Newman [1978] proposed the following relation for the critical amount of copolymer necessary to saturate the interface: Wcr = 3 φMw /(RaNAv)

(4.33)

where φ is the volume fraction of the dispersed phase, R is the radius of the dispersed drop, and NAv is Avogadro’s number. The authors also suggested (not derived from molecular parameters) that for a diblock copolymer: a • 5 nm2. Similarly, Mathos [1993] developed an expression for the minimum copolymer amount required to cover the surface of spherical particles. The model was based on the division of an outer shell surrounding the particle into pseudo-cubic elements, each containing randomly oriented

(random coil configuration) blocks of di-block copolymer:

[ / 9]

wcr = 9(3φ / R)( M A + M B ) / N Av r 2

[

3φMw / RN Av r 2

]= (4.34)

Where MA and MB are molecular weights of the A and B blocks respectively, MA + MB = Mw, and is square of the end-to-end distance of copolymer chain. Since is proportional to Mw, the above relation depends on the A/B ratio rather than on Mw’s. Eqs 4.33 and 4.34 are identical if a = / 9. The density of A-blocks around the particle was expressed as:

ρ A = 35 / 2 M A /  N Av r 2 

3/ 2

 

(4.35)

In the most general case, the periphery of a drop contains A chains of the matrix and the copolymer. Since is proportional to Mw, then density should be proportional to MA-1/2. Clearly, for long copolymer chains, the segmental density will tend to zero, thus implying higher contribution from the A chains. 4.2.4.2 Blends of Two Polymers with “Co-Solvent” Another way of compatibilizing immiscible polymer blends is by addition of a mutually miscible ingredient, a “co-solvent,” usually polymeric in nature. The objective here is not to generate a wholly miscible three component system, but to add just enough mutually miscible polymer. The co-solvent is to induce interactions between the immiscible polymers, thus compatibilize the blend, but preserving its two-phase structure. To help in the judicious selection of appropriate cosolvent for a given immiscible blend a partial list of miscible systems are given in Table 4.1. For more detailed information see Appendix 2 in this Handbook. In polymer blends, the thermodynamic miscibility depends primarily on specific interactions. For the entropic reasons, high molecular weight

306

Abdellah Ajji

homopolymers are expected to be immiscible when the specific interactions are absent. Many types of interactions may exist between two polymers. These include London dispersion forces between non-polar molecules, Coulombic ion/ion and ion/dipole interactions, dipole/dipole interactions between permanent or induced dipoles, charge-transfer forces, hydrogen bonding, etc. Polymer-polymer miscibility has been generally identified as mainly caused by hydrogen bonding (e.g., PPE/PS, PVME/PS, PVC/PCL, Phenoxy/ PCL, PEG/PAA), and for specific cases, by ionic and dipolar interactions. Some polymers have been found miscible with many other resins, or in other words: there are many immiscible blends whose components are miscible with the same polymer. Addition of this polymer can be used to partially homogenize the system, i.e., to compatibilize the blend. The added polymer is a co-solvent. Of particular interest are systems in which presence of a co-solvent makes it possible for the two immiscible components to form three-body interactions. In this case, the blend is indeed compatibilized, with the co-solvent being located in the interphase. For the thermodynamic reasons, mostly copolymers belong to this type of co-solvents. In the left hand side column of Table 4.1 there are polymers that may be used as co-solvents for pairs of resins listed in the other column. Some of the latter resins may show local miscibility (e.g., PS with styrenic copolymers), but the vast majority is immiscible. Note that according to the discussed strategy, the added polymer is to be miscible with the principal polymeric ingredients of the blend. On the other hand, it has been reported in the open and patent literature that addition of a small amount of a third immiscible polymer to a blend of two immiscible ones also improves the degree of dispersion. A “classical” example is a blend of PE with ”15 wt% PS — here addition of <2 wt% of PMMA reduced the domain size of PS by a factor of ten. In this case, the drop size reduction originates from hindered coalescence. The immiscible polymer preferentially migrates to the interface between the two principal resins,

providing protection to the dispersed phase against coalescence. This phenomenon has been commercially explored, e.g., in blends comprising PPE/SEBS drops dispersed in PBT matrix — an addition of immiscible PC improved the degree of dispersion and stabilized the morphology by reducing the risk of coalescence. In principle, any polymer that is miscible with two others can be used as co-solvent. As the data in Table 4.1 indicate, Phenoxy, PMMA, PPE, PC and PCL are miscible with several polymers, hence they are the best candidates for co-solvents of many systems. It is noteworthy that acrylic multipolymers are often used as additives to many blends. Their role is to enhance compatibilization as well as to toughen the blends. One of the most frequently used co-solvents is poly(hydroxy ether) of bisphenol-A, Phenoxy. Owing to the presence of phenolic group, the resin forms strong hydrogen bonding with numerous polymers. In 1970 Mitsubishi Chemical discovered that toughness of engineering resin blends could be improved by incorporation of a Phenoxy [Kimura, 1973]. This commenced a wide spread use of the resin. Phenoxy has been used as a compatibilizer in many blends (see Table 4.2). It is however to be noted that adding too much Phenoxy can make the blend miscible, what may reduce the mechanical performance, especially the impact strength. For example, Phenoxy was reported to lead to single-phase blends, viz. PBT/PEST or PC; PCL/PVME, PC or PEST, etc. 4.2.5

Conclusions and outlook

In the case of a binary A/B blends, Helfand and his co-authors provided the basic relations and the theoretical guidance to the properties of the interface, as well as to their modification by incorporation of a compatibilizer. The theory is based on strong, limiting assumptions (e.g., infinitely long macromolecules), thus one should not expect a quantitative agreement with experimental data. However, the theoretically predictions offer a powerful guidance for the best compatibilization strategies.

307

Interphase and Compatibilization by Addition of a Compatibilizer

Table 4.1.

Some examples of miscible polymers

Potential co-solvent

Polymers miscible with the potential co-solvent

PEEK PEG Phenoxy PMMA

PEK, PEI, PES, ... Phenoxy, PVC, TPU, EP, PAA, PES, EVAc, PVP, poly(meth)acrylates,... PES, PMMA, PVME, PVP, TPU, PBT, PET, ... PC, MSAN, EPA, poly(meth)acrylates, PF, NC, PPG, PVAc, PVC, PVC-VAc, CPVC, PVDF,... PS, CPS, SMA, P(α-MeS), SBS, SAN, ABS, ... PC(TM-BPA), PPE, CHMA, PVME, ... PVDF, halogenated-PMMA,... PCL, EVAc, PVAc, PBT, poly(meth)acrylates,... PCL, PVC, PPE, co-polymethacrylates,... PCL, PVME, methacrylates, vinyl butyral,.. PVP, PA’s, PEG,... SAN, (meth-)acrylates, PVAc Chloroprene, NC, PVC, PVCAc,... Phenoxy, PET, PBT, PARA, PVC,... PA’s, PET, PBT, PAr EPDM, PI,... PEST, Phenoxy, PVC, PVDC-VAc, LCP... Carbonates, adipates, (co-)polyesters, PCL, PMMA, SMA, ... Phenoxy, PC, PVC, PVB, CPE, NC, c-PP, PVDC-VAc, EPI, ...

PPE PS PVAc PVC SAN SMA Co-polyacrylic acid Epichlorohydrin NBR PAr PARA PB PBT, PET PC PCL

Table 4.2.

Examples of Phenoxy compatibilized blends

Blend

Comments

Reference

PA with SEBS

0.5-3.0 wt% Phenoxy significantly improved the tensile and impact strength above the values for PA 10 wt% Phenoxy compatibilized the 3-ingredients mixture. The materials showed excellent impact strength, ductility and solvent resistance 0.5-15 wt% Phenoxy increased elongation, tensile strength, Izod impact strength and uniformity Either Phenoxy, EVAc-GMA, SMA-copolymers or MBS were used as compatibilizers. The blends showed good processability, toughness, plateability, heat and water resistance Addition of Phenoxy™ was also found to provide good compatibilization

Freed, 1975

PC/PET/ Akryloid™ KM660 PA-6, PA-66, or PA-610, with either PET, PBT PSF/ABS

LCP/PEST

Liu and Giles, 1986

Robeson, 1988 Gaafar, 1990a; Orikasa & Sakazume, 1990; Golovoy and Cheung, 1994 Dashevsky et al., 1993, 1994

308

Abdellah Ajji

The interfacial thickness, ∆1’, and the interfacial tension coefficient, ν’, are both related to the square root of the thermodynamic binary interaction parameter, χAB — ∆1’ directly, whereas ν’ inversely, thus their product: , ∆1’. ν’, is to be independent of thermodynamic interactions. The latter conclusion may have limited validity, but the general tendency — the reciprocity between the interfacial tension coefficient and the interphase thickness — is correct. The theory correctly predicted the magnitude of the interphasial thickness, ∆1’= 1-4 nm. Note that the theory is for A/B binary systems, thus extending these predictions to compatibilized systems, where ∆1’ ” 65 nm may lead to erroneous expectations. For the latter system the reciprocity between ν’ and ∆1’ is not to be expected. For the strategies of compatibilization, Helfand’s theory provides three important conclusions: (1) the chain-ends of both polymers concentrate at the interface, (2) any low molecular weight third component is forced by the thermodynamic forces to the interface, and (3) the interfacial tension coefficient increases with molecular weight up to an asymptotic value. Noolandi et al. developed a theory for the interfacial region in three-component polymeric systems comprising di-block copolymer. There are two aspects to consider: the phase separation in block copolymer upon addition of one or two homopolymers, and the modification of the A/B blend properties upon addition of a block copolymer (either A-B or X-Y type). The second aspect is more pertinent for the polymer blend technology. In particular, the ternary blends comprising two homopolymers and a copolymer, either A/B/A-B, or A/B/X-Y are of industrial interest. Addition of a block copolymer, A-B, to immiscible blend of homopolymers A and B reduces the interfacial tension coefficient similarly as addition of a surfactant affects emulsions. Thus, the idea of the critical micelle concentration, CMC, and the limiting value of the interfacial tension coefficient, νCMC, can be applied to polymer blends. This suggests possible existence of a thermodynamically controlled stable droplet phase,

in which the minor phase homopolymer drops, protected by an interfacial film of copolymer, are dispersed within the major phase polymer. The theory predicts that CMC depends exponentially on the product χABZc — larger are the repulsive interactions and/or higher is the molecular weight of copolymer, lower is the concentration for saturation of the interface. Once the concentration of added copolymer exceeds CMC the micelles are formed. The interfacial tension coefficient exponentially decreases with -φcχΑΒZc, toward the limiting value at CMC, νCMC. Vilgis and Noolandi investigated the effects of addition of a block copolymer X-Y to blends of A and B polymers. Similar dependencies as those derived for A-B copolymer were found. The work suggested that it is possible to design universal compatibilizers based on the principle of competitive repulsive interactions between the homopolymers and copolymer’s blocks.

4.3

Determination of the Interfacial Parameters

4.3.1

Interfacial Tension Coefficient

As seen in Part 4.2, several theoretical approaches have been proposed for the description of the interfacial phenomena. The lattice theories by Helfand, Roe, Noolandi and their collaborators are based on the study of conformation and molecular environment. The derived relations are written in terms of the binary thermodynamic interaction parameter χ12 and the lattice constants. The theories do agree that the interfacial tension coefficient is a function of χ12, but the predicted functional dependencies are different: ν12 ∝ χn12, with exponent n = 1/2 to 3/2, depending on the assumptions. These and more recent theories can be considered as guides for the expected dependencies, but they can not be used directly to calculate either the interfacial tension coefficient or the interphase thickness. Since there is a significant disagreement between the theoretical relationships derived

309

Interphase and Compatibilization by Addition of a Compatibilizer

4.3.1.1 Calculations of the Interfacial Tension Coefficient 4.3.1.1.1 From Surface Tension Coefficients Among the different approaches used to calculate the surface tension coefficient, νi, the most useful seems to be the parachor method [Sugden, 1924; Van Krevelen, 1976; Wu, 1982]. This parameter was defined as [Sugden, 1924]: Pi ≡ Miν i1/ 4 / ∆ρi

ν12 = ν1 + ν 2 − 2(ν1dν 2d )1/ 2 − 2(ν1pν 2p )1/ 2

where Mi is the molecular weight of a liquid i and ∆ρi is its liquid-minus-vapor density difference. The value of Pi can be calculated for any chemical compound from the contributions of structural elements. Once the surface tensions are known, the second step for calculating the interfacial tension coefficient, ν12, requires a relationship between νi and ν12. Two dependencies have been proposed, first one by Girifalco and Good [1957]: ν12 = ν1 + ν2 - 2ϕ (ν1ν2)1/2

12 error bars = ± 20%

8

4 harmonic mean G&G Eq

0

0

4

8

n (experimental)

12

12

Figure 4.5. Comparison between calculated and measured values of the interfacial tension coefficient from Girifalco & Good Eq 4.37, and from the harmonic mean Eq 4.38, incorporating the dispersive and polar contributions.

(4.37)

(here ϕ is a semi-empirical interaction parameter) and the second by Wu [1982]. The latter one is based on the concept of fractional polarity that assumes the molecular forces originate from the polar and dispersive interactions: ν12 = νl + ν2 - 4ν1dν2d / (νld + ν2d) - 4νlpν2p / (νlp + ν2p) 2

(4.39)

In Figure 4.5, the interfacial tension coefficients calculated from Eqs. 4.37 and 4.38 are compared with the experimental values. For the selected pairs of polymers, the latter seems to provide a better correlation with the measured values of ν12.

12

(4.36)

equation. The idea can also be written in form of a geometric mean equation:

n (calculated)

for the interfacial tension coefficient, several attempts were made to calculate its approximate values from other physical quantities. Two methods have been considered, the first using the surface tension coefficients, νi, as an intermediate physical quantity, and the second the solubility parameters.

2

= (νld - ν2d) / (νld + ν2d) + (νld - ν2d) / (νlp + ν2p) (4.38) The subscript d and p refer to dispersive and polar components of the surface tension coefficients. Eq 4.38 is called the harmonic-mean

4.3.1.1.2 From the Solubility Parameters The cohesive properties of a material are related to the solubility parameter, δi, that originates from different types of interactions: dispersive or atomic, molecular of the type polar and hydrogen bonding, induced dipoles, metallic, etc. The first three types are most important, thus respectively:

δ i2 = δ id2 + δ ip2 + δ ih2

(4.40)

The values of the δi-components can be calculated for any chemical substance from the tabulated group and bond contributions. Once δi is known for both polymers in the blend, the Huggins-Flory binary thermodynamic interaction parameter, χ12, can be calculated from:

310

Abdellah Ajji

2

χ12 = (V / RT )[δ1 − δ 2 ]

(4.41)

from Eq 4.41 and predictions of the lattice theories, the interfacial tension coefficient can be written as [Luciani et al., 1996]: n ν12 = kRTχ12 = k1 ( ρRT ) n −1 1

{(δ

1, d

− δ 2, d )2 + (δ1, p − δ 2, p )2 + (δ1,h − δ 2,h )2

}

n

{

∴ν12 = k1 (T ) (δ1, d − δ 2, d )2 + (δ1, p − δ 2, p )2 + (δ1,h − δ 2,h )

4.3.1.2 Determination of the Interfacial Tension Coefficient For low viscosity Newtonian fluids, several methods for the measurements of ν12 have been developed [Wu, 1974]. However, the high viscosity of industrial polymer melts makes most of them irrelevant. The few remaining ones that can be used for determination of ν12 in polymer blends, can be divided into equilibrium and dynamic methods [Luciani et al., 1997]. 4.3.1.2.1 Equilibrium Methods

2 n

}

(4.42) To evaluate the validity of Eq. 4.42, the experimental data of ν12 for 46 polymer blends are plotted in Figure 4.6 as a function of the computed values of the bracketed sum [Brandrup and Immergut, 1989]. The straight line represents the least squares fit. It is noteworthy that the exponent n = 0.402 is close to 1/2, predicted by Helfand et al., lower than 3/4, predicted by Roe et al., and significantly lower than the value of 3/2 derived by Joanny and Leibler.

Pendant drop method The pendant drop method, schematically shown in Figure 4.7, was used in many studies of polymeric blends, usually with low molecular weight fractions of given polymer pairs [Wu, 1974]. The technique is based on the analysis of the drop shape of component-1 emerging from the extremity of a capillary, immersed in component-2 [Anastasiades et al., 1987, 1988, 1989; Owens et al., 1989a, 1989b].

A

B

102

Interfacial Tension Coefficient, n12

T = 150°C error bars ± 36% (calc) 10

1

ds d max

de = d max

100 2 0.402

n 1 2 = 0.881 [(D Dd ) ] R = 0.815 10-1 0 10

101

102

103

D d2 = D d 2 + D d 2 + D d 2 d p h Figure 4.6. Interfacial tension coefficient at 150°C for 46 polymer blends plotted vs. the solubility parameter contributions (see Eq 4.42).

i f i representation l i b h Figurei 4.7.f h Schematic of the interfacial tension measurements by the pendant drop method; de is the maximum drop diameter, and ds is the diameter located at the distance de from the drop apex.

The interfacial tension coefficient can be calculated from the drop shape using the relation: ν12 = g ∆ ρde / H(S)

(4.43)

Interphase and Compatibilization by Addition of a Compatibilizer

where g is the gravitational constant, ¨ρ the density difference between the two molten polymers, and H(S) is a function of the form factor S, defined as: S ≡ d s / de

(4.44)

where ds and de are defined in Figure 4.7. The values of H(S) are tabulated [Adamson, 1982]. Application of this method requires the knowledge of the polymer melts’ density at the processing temperature. Sessile drop method The sessile drop method is similar to the pendant drop one [Sakai, 1965]. The same scheme is used, but in this case the droplet is resting on a plane surface immersed in the second component (see Figure 4.8) — ν12 is calculated from analysis of the drop shape at equilibrium (characterized by the relative magnitude of the shape parameters, X and Z, defined in Figure 4.8), knowing the densities of the polymeric fluids at the temperature of measurement.

Figure 4.8. Schematic representation of the interfacial tension measurements by the sessile drop method; X is the equatorial radius, and Z is the distance from the equator to the drop apex.

Spinning drop method In the spinning drop method, a droplet of polymer-1 immersed in polymer-2 is submitted to extensional deformation through the use of centrifugal forces [Elmendorp and De Vos, 1986]. The extent of deformation is controlled by the rotational speed imposed on the system contained

311

in a cylindrical tube (see Figure 4.9). The rotational energy is balanced by the increased interfacial area. The interfacial tension coefficient is calculated from the equilibrium drop deformation. For drops with length to diameter ratio larger than 4 the following dependence was suggested [Vonnegut, 1942]: ν12 = ω2 d3 ¨ρ / 32

(4.45)

where d is diameter of the spinning drop, and ω is the angular velocity of the cylinder. The technical difficulties of constructing an apparatus suitable for handling highly viscous polymer melts make this method too elaborate to be commonly used.

Figure 4.9. Schematic representation of the interfacial tension measurements by the spinning drop method; ω is the angular velocity, and d the deformed drop diameter.

All these equilibrium techniques are rarely applicable for the measurements of the interfacial tension coefficient of common, industrial polymers. A long time is usually required to reach the equilibrium. Since the rate is determined by the ratio of the interfacial tension coefficient to viscosity, ν12/η, the time to reach an equilibrium can take either minutes or days — this may lead to the thermal degradation and uncertain validity of the measured ν12. As the degradation products are preferentially adsorbed at the interface, the determined interfacial tension coefficient may not reflect the blend behavior during compounding or forming. The influence of the equilibration time (thus degradation and contamination) on the ν12 value is expected to be significant [Grace, 1982]. The pendant and spinning drop methods were used to determine ν12 for different polymer pairs as well as the effect of temperature and compatibilizer [Garmabi and Kamal, 1998]. The authors

312

Abdellah Ajji

found a linear relationship between the temperature (T) and ν12 for PE/PA and PP/EVAl (see Figure 4.2), as well as strong effect of compatibilizers. The pendant drop method and T-effects on ν12 for different PS/PP blends were also described [Demarquette and Kamal, 1994]. 4.3.1.2.2 Dynamic Methods In dynamic methods, determination of ν12 is based on the time evolution of a fluid element shape, from a non-equilibrium to an equilibrium state. The evolution is driven by the interfacial tension and, depending on the initial shape of the element, it can follow different dependencies. Capillary breakup method The capillary breakup method is based on Tomotika’s theory [Tomotika, 1935, 1936]. The author was the first to investigate the development of Rayleigh’s instabilities in cylinders of one fluid imbedded in another (see Figs. 4.10 and 4.11) [Rayleigh, 1879]. The amplitude variations of the sinusoidal distortions, α, can be described by:

Table 4.3.

α ≡ (b − a) / 2 d0 = α 0 exp{qt}

(4.46)

with q = ν12Ω(x, λ) / ηmd0

(4.47)

where λ is the relative viscosity of the system (λ = ηd / ηm) and Ω(x, λ) is a complex function of λ and the observed wavelength, Λ, of the distortion, expressed as: x = πd0 / Λ. The values of Ω(x, λ) can be calculated from Tomotika’s equations. The linear dependence of ln(α) on the fiber disintegration time enables estimation of q and, consequently, ν12. It seems that Chappelear [1964] was the first who applied this technique to measure the interfacial tension coefficient of polymer blends. Further refinements have been published [Elemans, 1989; Elemans et al., 1990; Elmendorp, 1986]. The method is simple, not requiring special equipment, but the zero-shear viscosity of the investigated polymers at the processing temperature must be known. Typical results obtained this method are shown in Table 4.3.

Interfacial tension coefficient as determined using the capillary breakup method [Luciani et al., 1997]

Matrix

Fiber

T(°C)

ν12(mN/m)

PS 103-300 LDPE 1001 PE 300

PA-6 Z211 PA-6 Z211 PA-6 C316

PS 220

LDPE 1001

PS 220 PMMA V920 PMMA V920 EP

PMMA V920 LDPE 1001 PA-6 Z211 PVDF

EP

PA-6 C316

EP LDPE 1001+0.5% Irganox MD 1024 LDPE 1001+0.5% Irganox MD 1024

PA-6 C316 + Lotader PS 220 PA-6 Z211

230 230 220 250 280 200 230 200 200 230 220 260 220 240 220 200 230

6.3 8.9 11.7 10.9 9.2 6.4 5.4 1.8 5.9 6.9 3.4 2.5 10.3 9.4 2.8 6.4 14.4

Interphase and Compatibilization by Addition of a Compatibilizer

313

Figure 4.10. Schematic representation of the interfacial tension measurements by the capillary breakup method; d0 is the initial diameter of the fiber, b and a are the maximum and minimum diameters observed during the break-up process, respectively.

Deformed drop retraction method (DDRM) The deformed drop retraction method (DDRM) has been recently proposed. It makes it possible to determine ν12 from the time evolution of deformed ellipsoidal drops towards its spherical equilibrium form [Luciani et al., 1997]. Assuming an ellipsoidal shape for a drop of initial radius R0 deformed under simple shear, the following relation was found: ν12 (tc )   40(λ + 1)   D = Do exp −t   =   (2 λ + 3)(19λ + 16) ηm Ro   Do exp{−t / τ }

Figure 4.11. Optical micrographs of a PA-6 fiber imbedded in PS matrix at 230°C. The initial diameter of the fiber was 46 µm. The photographs were taken (from the top) after: t = 0, 10, 19, 23, 25, and 29 min.

(4.48)

with D = (L - B)/(L + B) (where L and B represents the long and short diameter of the prolate ellipsoid, respectively), λ is the viscosity ratio, ηm the matrix viscosity, ν12(tc) the interfacial tension at the polymer-polymer contact time tc, and D0 is the initial value of D at t = 0. In principle, this relation is valid only for Newtonian systems. However, the method can be used to characterize viscoelastic materials provided that:

314

Abdellah Ajji

1. The retraction rate is sufficiently slow to ensure that the material behave as Newtonian, and 2. The elastic relaxation of the materials after deformation is faster that the ellipsoidal droplet retraction, that is, a clear separation of the two mechanisms is achieved.

dynamically driven migration of chain-ends, low molecular weight fractions and additives, as well as thermal degradation. The contact time dependence of ν12 explained some of the differences reported for the data obtained using different measurement techniques, viz. pendant drop, capillary breakup, or ellipsoid retraction techniques [Luciani et al., 1997].

7

4.3.2

6

n1

2

(mN/m)

Interfacial tension coefficient vs. contact time for PS/LDPE at 200°C

error bars: ± 6% of value 5

0

100

200

Contact time (min) Figure 4.12. Interfacial tension coefficient in LDPE/PS system as a function of the contact time at 200°C. Vertical bars correspond to ± 6 % error [Luciani et al., 1997 ].

DDRM is particularly useful for the binary polymer blends. The dynamic interfacial tension coefficient, ν12, is determined from the time evolution of a distorted fluid drop toward its equilibrium form. Measurements of either low viscosity model systems or high viscosity industrial polymer mixtures led to a good agreement with values obtained from the widely used breaking thread method. DDRM enables to measure ν12 in polymeric blends of commercial interest — the high viscosity systems that frequently are impossible to characterize by other techniques. Furthermore, for the first time it is possible to follow the time dependence of ν12, thus unambiguously determine its dynamic and equilibrium values. For example, in LDPE/PS blends ν12 decreased with the polymer-polymer contact time, tc, from ν12 = 6.9 mN/m at tc = 12 min, to ν12 = 5.2 mN/m at tc 75 min — the latter may represent the true thermodynamic equilibrium value, ν12εθ. Such a large reduction of ν12 may be due to the thermo-

Interphasial Thickness

Much less information is available on the methods for the determination of the interphasial thickness in polymer blends, ∆l, than that of ν12. For binary systems, assuming that these two parameters are interrelated, one may estimate ∆l from ν12, the latter determined using one of the above-described methods. To determine the experimental value of ∆l in any system, diverse methods have been used, viz. electronic microscopy, X-ray scattering, ellipsometry [Inoue, 1991; Fayt, 1986], light scattering, etc. Results of these measurements are presented in the next section (see Table 4.6). 4.3.2.1 Microscopy Since the dimensions to be probed are of the order of few nanometers, the most useful microscopic method would be that of the transmission electron microscopy (TEM). Using this technique Fayt et al. [1986] observed location of P(S-b-HB) compatibilizer in a PE/PS blends. The authors inserted a short sequence of isoprene between the styrenic and hydrogenated butadiene blocks. After staining the isoprene double bonds with OsO4, the authors were able to observe presence of the copolymer at the interface between the matrix and dispersed phase. Thickness of the interphase could then be measured. The experiments also demonstrated presence of the added compatibilizer as dispersed micelles inside the PE phase. This technique is applicable, however, only when selective staining affects only the compatibilizer. 4.3.2.2 Ellipsometry The ellipsometric method has been developed by Yukioka and Inoue [1991, 1993]. The principles

315

Interphase and Compatibilization by Addition of a Compatibilizer

Rip exp(i∆ p ) Rip Rip ρ= s = s = s exp[i( ∆ p − ∆ s )] = Ri Ri exp(i∆ s ) Ri

of the technique and the model used for calculating the thickness the interphase are schematically illustrated in Figures 4.13 and 4.14, respectively. The retardation (∆) and reflection ratio (tan(ψ)) can be determined from the ellipsometric readings. The adopted model assumes existence of four layers: air, thin polymer-1, interphase, and thick polymer-2 (see Figure 4.14). In the interphase the refractive index is assumed to be an average: n3 = (n1 + n2) / 2. Thus, one can compute the best value of the interfacial thickness, d3, to fit the observed values of ∆ and tan(ψ). The following relations were derived for the computation of d3:

tan(ψ ) exp(i∆ ) j

where: Rij =

ri + Rij+1 exp( −iDi +1 ) 1 + ri j Rij+1 exp( −iDi +1 )

;

Di = 4πni di cos(θ i ) / λ ; and ri p =

ni +1 cos(θ i ) − ni cos(θ i +1 ) ; ni +1 cos(θ i ) + ni cos(θ i +1 ) ris =

ni cos(θ i ) − ni +1 cos(θ i +1 ) ni cos(θ i ) + ni +1 cos(θ i +1 )

n1 sin(θ1 ) = n2 sin(θ 2 ) = n3 sin(θ 3 ) = n4 sin(θ 4 )

(4.49) 1

2

7 3 4

6 7 0° 5

lock- in

ampl ifier

pulse mot or comput er

Figure 4.13. Experimental setup for the ellipsometric measurements: 1. is a He-Ne laser, 2. is a chopper, 3. is a polarizer, 4. is a compensator, 5. is a sample chamber, 6. is an analyzer, and 7. is a detector (e.g., a photodiode). After [Yukioka and Inoue, 1991].

Figure 4.14. The four-layer model used for the calculations of the interfacial thickness. ni designates the refractive index, θi the incidence angle, ri the Fresnel reflection coefficient, Ri the reflection coefficient in the incident plane and di the thickness of the layer.

316

Abdellah Ajji

In Eq 4.49: superscript p and s designate parallel and perpendicular polarization of light, and λ is the true interfacial thickness. The other quantities in these equations are defined in Figure 4.14. The interface thickness needs to be corrected using the composition profile. The true interfacial thickness is then determined as ∆l = λ = d3 / 1.7. For example, the ellipsometric technique was used to determine the interfacial thickness in SAN/PA blends, compatibilized by addition of SMA [Yukioka and Inoue, 1991, 1993]. A thin bi-layer film of SMA/SAN was prepared then mounted on a thick PA substrate. The interfacial thickness varied with the compatibilization time from ∆l = 2 to 30 nm. The method was also used to study the interphasial thickness variation in PCL blends with CTBN or CTBR. Without a reactive modifier (aminopropyltriethoxysilane, APS), the measured thickness was ∆l = 3 nm, upon addition increasing to 6 nm [Okamoto and Inoue, 1993]. 4.3.2.3 X-ray Scattering Small angle X-ray scattering (SAXS) has been used by many authors to determine the interfacial thickness. An excellent review of this subject can be found in Perrin and Prud’homme [1994]. Many methods of calculations can be used. One of these involves an analysis of the deviation from Porod’s law, in which the desmearing procedure is avoided. This procedure was applied to blends of PS with PMMA added with a P(S-b-MMA) block copolymer. Upon addition of copolymer the interface thickness changed from ∆l = 2 to 6 nm [Perrin and Prud’homme, 1994]. 4.3.2.4 Atomic Force Microscopy Another technique of interest is the atomic force microscopy (AFM). The method can be used in any system where properties of the separated phases significantly differ. For example, the method was used to determine thickness of the interphase in blends of PA with soft elastomer. In this case the rigidity scanning procedure was the most appropriate. Since the difference in moduli

between the dispersed phase and the matrix was large, scanning in this mode gave a rigidity profile across the thickness. The interfacial thickness was found to be about 20 nm [Champagne, 1996]. 4.3.2.5 Other Methods Several other methods, more or less tedious, have been described in the literature. They include welding cleavage experiments [Foster and Wool, 1991], neutron reflectivity [Fernandez et al., 1988], etc.

4.4

Compatibilization by Addition of a Compatibilizer

There are several strategies of compatibilization, e.g., (i) Addition of a small quantity of a third component that either is miscible with both phases (a co-solvent), or is a precisely tailored copolymer whose one part is miscible with one phase and another with another phase (0.5-2 wt%, usually block-type, less frequently graft one); (ii) Addition of a large quantity, ”35 wt%, of a core-shell copolymer(s) that behaves like a multi-purpose compatibilizer-cum-impact modifier; and (iii) Reactive compatibilization, designed to enhance the domain interactions and generate finer morphology by creating chemical bonds between the two homopolymers during the compounding or forming processes. The latter type of compatibilization is treated in Chapter 5 of this Handbook. In the following part aspects related to morphology of the blends with and without an added compatibilizer (from the nm to the µm scale) will be presented. These include information on the interfacial characteristics (interfacial tension coefficient and thickness of the interphase), morphology, crystallization and performance of the blends. This information will be presented mainly in a tabulated form, summarizing the main features from the referenced publications.

317

Interphase and Compatibilization by Addition of a Compatibilizer

12

Coefficient, n Tension

0.45

Interfacial

0.65

PS/PB/P(S-b-B) Tang & Huang, 1994

0.55

Utracki & Shi,1992

0.35

0.25

0

1 SB

2

3

(wt% dispersed phase)

Figure 4.15. Interfacial tension coefficient as a function of diblock copolymer in the minor phase. System: matrix — polybutadiene, dispersed phase (pendent drop technique) — polystyrene with styrene-b-butadiene di-block copolymer. Data: [Anastasiadis, 1988]. Dotted and solid lines were computed from Eqs 4.29 and 4.30, respectively.

(mN/m)

Experimentally, it has been shown that di-block copolymers, especially of the type polystyrene-bhydrogenated polybutadiene, P(S-b-HB), have a higher interfacial activity than triblock or graft copolymers. Di-blocks more readily interact with the homopolymer phases, forming appropriate entanglements that result in reduction of ν12 in the melt and enhanced interphasial adhesion in the solid state. When properly selected and added in optimum quantity, they were also found to stabilize the morphology against the shear degradation during abusive processing. Experimental data are listed in Tables 4.4-4.6. Addition of a graft copolymers is much less frequent. Some blends of this type were reported to have unusual onion-like morphology. This observation is not universally valid — the compatibilizing effect must depend on the structure and composition of the copolymer. Owing to complexity of the these structures the theoretical analysis has not been attempted. Effective compatibilization of binary polymer blends by addition of a copolymer reduces the dispersed particles’ size and ν12 [Anastasiadis et al., 1987; Wu, 1987; Patterson et al., 1971]. An illustration is shown on Figure 4.15. The effect of compatibilizer addition is similar to the emulsification of the classical emulsions. In the former systems, the compatibilizer effect on the drop size and ν12 follows the same behavior as the emulsion drop size reduction upon addition of a surfactant. The latter behavior is usually described as the “titration curve” that characterizes the surfactant efficiency. The shape of the titration curve depends on the type of emulsifier and the emulsification process, e.g., mixing time and equipment. However, the amount of emulsifier to saturate the interface also depends on the affinity of emulsifier to the dispersed phase, the size of the dispersion, the orientation of the emulsifier at the interface and its ability to prevent flocculation and coalescence [Djakovic et al., 1987]. A similar behavior is to be expected for polymer blends upon addition of a compatibilizer.

Recently, the effect of a reactively formed compatibilizer on ν12 was studied using the pendant drop method [Garmabi and Kamal, 1998]. Their results are shown in Figure 4.16. A small quantity of MAH drastically reduced the interfacial tension coefficient, but the amount of formed graft was not determined.

12

PA/PP/PP-MAH blends

1 2

Interfacial Characteristics

Interfacial Tension Coefficient, n

4.4.1

n

8 a a a

o

= a + a /(MAH + a ) o

1

2

Value

Error

0.74098

0.9194

1

0.07097

0.036909

2

0.0073112

0.0034922

R

4

12

0.97488

0 0

0.04

0.08

0.12

MAH content (%)

Figure 4.16. Interfacial tension as a function of added maleic anhydride for PE/PA-6 blend.

318

Table 4.4.

Abdellah Ajji

General interfacial characteristics of blends

Blend system

Techniques and Results

References

PE/PS with P(HB-b-S) or P(HB-b-I-S) PS/PMMA

Used TEM microscopy and other methods. Copolymer located at the interface Embedded fiber retraction technique at 190°C. ν12 = ν’- C1Mnz, z = 0.73±0.24.

Fayt, et al., 1986, 1989; Hobbs et al., 1983 Ellingson et al., 1994

For Mn exceeding 48 kg/mol, ν12 was independent of Mn with a mean value of 1.2 mN/m.

The efficiency of a block copolymer is limited by the formation of micelles in bulk phases and by the kinetic factors. Consequently, the block copolymer used as a compatibilizer should be designed by taking thermodynamic and kinetic parameters into account to achieve the desired effects. Thus the structure and transitions in copolymers and homopolymer/copolymer systems is of great interest. 4.4.2

Thickness of the Interphase

In contrast to measurements of the interfacial tension coefficient, only few measurements of the interphase thickness have been reported [Wlochowicz and Janicki, 1989; Janicki et al., 1986]. For example, domain boundary thickness’ were measured in PS/PMMA blends [Foster et al., 1990; Fernandez et al., 1988; Russel et al., 1991; Perrin and Prud’homme, 1994]. Generally, values in the range of 2-6 nm were reported for the interface thickness. 4.4.3

Morphology

As shown in Table 4.7, several methods have been used to characterize the morphology of polymer blends containing a block copolymer. In the pioneering works, transparency of films cast from solutions was used as a measure of the emulsifying agent efficiency [Riess et al., 1967; Banderet et al., 1967]. Other studies have focused on the observation of phase size reduction in the scanning

or transmission electron microscopy, SEM and TEM respectively [Ouhadi et al., 1986; Fayt and Jerome, 1990; Heuschen et al., 1990; Jo et al., 1991; Sakellariou, 1991]. For block copolymers, the spherical, cylindrical, lamellar, and many other morphologies (including newly discovered ordered bi-continuous and grain boundary structures) have been reported for the copolymers’ microphase separated state and for their blends with homopolymers. New structures were characterized by hyperbolic interfaces [Hasegawa et al., 1993]. 4.4.4

Crystallization

Another important aspect of compatibilization is the effect on blends micro-morphology, i.e., on the total crystallinity and the crystalline morphology [Utracki, 1989; Xavier, 1991; Nadkarni and Jog, 1991]. Crystallization may take place only within the temperature region limited on the upper side by the melting point, Tm, and on the lower side by the glass transition temperature, Tg. The crystallization behavior very much depends on the state of miscibility, as well as on the nature of other blend’s components. In miscible blends Tg is a monotonic function of the components’ Tg’s. When the second component is amorphous, its presence can either decrease or increase the tendency of the first resin to crystallize. The process depends on the blending effects on Tg and on the chain mobility (the free-volume effect). When Tg of the amorphous component is lower

319

Interphase and Compatibilization by Addition of a Compatibilizer

Table 4.5. Copolymers and homopolymer with copolymer systems System

Techniques and Results

References

Homo-polymer-copolymer

Solubilization of homopolymer by copolymer. The homopolymer A dissolves in micro-domains of A-blocks of the A-B copolymer, until a limit of miscibility is reached. Above this limit, excess of homopolymer forms macroscopic domains PS-block with Mw = 9.5 and PVP-block with Mw = 10.6 kg/mol.

Hashimoto et al., 1990; Tanaka and Hashimoto, 1991; Tanaka et al., 1991

PS with P(S-b-VP) copolymer

PS with P(S-b-I) copolymer Copolymer

PS/PMMA & symmetric P(S-b-MMA) copolymer

Copolymer

PS/SB/S-B

PPE, with SAN, SMMA, SAA or SMA

SAXS was used. Blends containing PS with Mw < 9.5 were transparent; those with Mw > 13.7 kg/mol were opaque. The solubility limit increased when MW of PS decreased. SAXS. Polymer dissolved in domains of block of the same nature as long as its Mw is lower than that of the block. Used X-ray scattering. Micro-domain measured repeat distance: D ∝ Mw0,643; theoretical D ∝ Mw2/3 Used SAXS and TEM. Stabilization of the interface, the formation of micelles, and the micro and macro phase separation. Same morphology was observed in solution and after solvent evaporation The effect of temperature, Mw, preparation history, etc., on D. The interphasial thickness, ∆l, was found to remain constant.

SALS, TEM and SAXS were used to study PS-solubilization in S-b-B copolymer (50 wt% styrene). The solubility of PS in the S-block was governed by the ratio of the homopolymer Mw to the block’s Mw. The lamellar repeat period increased linearly as more PS was added. The butadiene layer thickness remained constant. Miscibility of copolymers with PPE when the amount of comonomer is small. The interfacial energy between the blend components was significantly reduced by adding either a PS-b-PMMA, or PS-b-PEB-b-PMMA. The copolymers had a profound influence on morphology, phase adhesion and mechanical properties of the blend.

than that of the semicrystalline, the crystallization envelope is increased and the crystallization is facilitated. The effect is similar to crystallization from a solvent — both the Tg and chain mobility

Skoulios et al., 1971

Ptaszynski et al., 1975 Hashimoto et al., 1980; Matsushita et al., 1990 Ohta and Kawasaki, 1986 Lowenhaupt and Hellmann, 1990, 1991, 1994

Spontak and Zielinski, 1993; Ohta and Kawasaki, 1986; Roe et al., 1981; Anastasiadis et al., 1989, 1990; Bates et al., 1983; Hashimoto et al., 1977, Todo et al., 1977; Todo, Hashimoto and Kawai, 1978 Jeon and Roe, 1994

Gottschalk et al., 1994

tends to maximize the total crystallinity. Such blending may induce crystallization of polymers that are known to be amorphous, e.g., PC. In the opposite case, addition of an amorphous

320

Table 4.6.

Abdellah Ajji

Interfacial thickness in polymer blends

Polymer or blend

Copolymer

Techniques and Results

References

Polymer A

P(A-b-B)

The interfacial thickness did not change upon addition of homopolymer, ∆l ≅ 2 nm.

PS/PMMA

P(S-b-MMA)

PS

P(S-b-VP)

Neutron reflectivity using blends with high copolymer content and an equal amount of homopolymers. The micro-domain boundary thickness was found to increase by 25% from the pure copolymer to a mixture containing 17% homopolymer; ∆l ≅ 2 nm. Studied Mw effects of di-block copolymer

Ptaszynski et al., 1975; Bates et al., 1983; Zin and Roe, 1984; Hashimoto et al., 1990; Tanaka et al., 1991 Russel et al., 1991b

PS/PVP

poly(styrene-b-2vinyl pyridine)

PS/PMMA

PS/PMMA

P(S-b-MMA)

PCL with CTBN or CTBR PE/PP

APS

EPR

segregation at the interface using the segregation isotherms. A normalized interfacial thickness was found a universal function of that portion of the block copolymer chemical potential due to chain stretching. Forward recoil spectrometry, FRES, was used to study diffusion of copolymer randomly dispersed in PS layer of a PS/PVP bi-layer sample. Free copolymer chains were detected at the interface below φCMC, whereas, above it, copolymer chains were also found at the PS/air surface as well as micellar segregation at the interface was visually confirmed. Determined ∆l = 6 nm by the welding cleavage experiments. For spin coated samples, ∆l = 5 nm was determined by neutron reflectivity Similar experiments as above also gave ∆l = 5 nm. Increasing thickness of the P(S-b-MMA) layer between PS and PMMA to 23.9 nm caused a gradual increase of ∆l from 5 to 8.4 nm. Micellar segregation at the interface of samples was not observed by the elastic recoil detection method. SAXS method was used. Upon addition of the copolymer the interface thickness changed from ∆l = 2 to 6 nm Ellipsometric methods. ∆l = 3 nm without APS and 6 nm with APS (a reactively generated copolymer) The interfacial thickness was ∆l = 1.5 to 2.8 nm without copolymer. The blend was compatibilized by adding EPR.

Dai and Kramer, 1994

Shull et al., 1991

Foster and Wool, 1991 Fernandez et al., 1988 Anastasiades et al. 1989

Russel et al., 1991 Jeon and Roe, 1994; Perrin and Prud’homme, 1994

Okamoto and Inoue, 1993

Teh et al., 1994

321

Interphase and Compatibilization by Addition of a Compatibilizer

Table 4.6.

Continued.

Polymer or blend

Copolymer

Techniques and Results

PE/PS

P(HB-b-S) or P(HB-b-I-S)

Copolymer migrated to the interface. Observed Fayt, et al., 1986, 1989; a continuous layer around each dispersed particle, Hobbs et al., 1983 either of PE in PS or PS in PE. ∆l =10-12 nm was determined. 2 wt% of P(HB-b-S) was sufficient to reduce the ν12, the particles size as well as prevent coalescence. At 5 wt% the copolymer effect was stable.

Table 4.7.

References

Morphology in copolymer and homopolymer/copolymer blends

System

Techniques and Results

References

P(S-b-I) blends with PS PS with P(S-b-B), (K-resin) P(S-b-B-b-MMA) triblock copolymers and their hydrogenated analogs P(S-b-BMA) and blends with PPE

Lamellar-catenoid or double diamond morphologies as well as microstructural disorder have been reported. Mesh-and-strut morphology consisting mostly of hyperbolic interfaces, found between the lamellar and spherical phases were observed At 17% B-block a “cylinder at the wall” morphology was observed — here PB and PHB cylinders were located at the lamellar PS/PMMA interface. At 6% B-block, PB formed spheres at this interface. Hydrogenation of PB induced significant changes.

Spontak et al., 1993

From dilute 2-propanol solutions, 42 wt% copolymer formed micelles with PS in the core. From THF solutions, ordered cylindrical structures were obtained. Addition of 10 wt% PPE turned the cylinders into spheres with PPE cores insulated by swollen PS blocks from a PBMA matrix. In PPE/P(S-b-BMA) 50/50 blends multilayer vesicles and lamellar structures coexisted. In blends having less than 5% of P(S-b-BMA), micelles with PBMA cores were observed in PPE matrix. PMMA or PS The microphase separation was observed only when low Mw homopolymers with P(S-b-MMA) were used. Blends of a polymer C having strong specific interactions with A, did not show any Mw limit for the microphase segregation. PE with SB, SBS, In low-viscosity PE matrix, SB, SBS and statistical SB copolymers formed SI or SIS spheres or ellipsoids. Broken filaments have been observed for SBS at 240°C copolymers and for SIS at 270°C. When PE was blended with high-molecular-weight of either. SB or SIS, at 240 and 220°C respectively, co-continuous, elongated structures were observed.

Hashimoto et al., 1992 Stadler et al., 1995

Siqueira and Nunes, 1994

Lowenhaupt et al., 1994

Getlichermann and David, 1994

322

Table 4.8.

Abdellah Ajji

Blends of two homopolymers with or without a copolymer

Blend

Additive

Techniques and Results

References

PE/PP

EPR

Processing conditions, melt viscosity and elasticity ratios, Mw and MWD,

Teh et al., 1994

PAr/ PVDF

PP/PC

PPE/PE ionomer PBT/ PMMA PS/PP

PS/PVC

affect the melt and solid morphology, and hence the performance. PAr-bPVDF/PMMA is a miscible blend. Addition of PAr to the PAr/PVDF/PAr-b-PMMA PMMA resulted in reduction of PVDF Tm and an increase of PAr Tg. These effects were enhanced by addition of PAr-b-PMMA. Finer dispersion was obtained for higher block copolymer content. Contact angle measurements, showed that ν12 was greatly influenced by the presence of block copolymer. SEBS Distinct changes of morphology were observed for 10-20 wt% PC, when the PP/SEBS Kraton G ratio varied from 95/5 to 90/10. SEM studies showed composite PC droplets with SEBS envelopes, embedded in PP matrix. This suggests stronger interfacial interactions between SEBS and PP than between PC and PP. P(S-b-VP) Melt mixing at 290°C. Addition of a small amount of PS-b-PVP was found to reduce the domain size Phenoxy When a small amount of polyhydroxy ether of bisphenol-A (Phenoxy) was added to the PBT/PMMA blends, the morphology changed into more regular and finer. For >50 wt% Phenoxy the blend converted to single-phase. PSiS Excellent dispersions in PS with up to 5% of polydimethylsilylene-co-phenylmethylsilylene (PSiS) were obtained. In PP at >1 wt% PSiS, segregation was observed. Blending of PSiS with PS decreased surface resistance, increased hardness, and provided protection from degradation by sunlight. graft coA block-graft copolymer was prepared by grafting a di-block copolymer of styrene and polymer butadiene, with cyclohexyl methacrylate monomers. PS/PVC blends comprising small amount of the block-graft copolymer showed excellent behavior. The copolymer had strong effect on the blends’ morphology.

component increases Tg of the blend, reduces the crystallization envelope, and thus hinders crystallization. When the blends are immiscible in the molten state, the crystallinity is even more complex function of the ingredients’ properties, compatibilization method, processing parameters, and post-processing treatments. The following factors have been identified to play major role: (i) molecular constitution and Mw of the components; (ii) composition, (iii) the type of phase morphology and the degree of dispersion (iv) the interphase, thus interactions between the phases, nature of the interface, migration of nuclei from one phase to the other, etc.; (v) melt history, in

Ahn et al., 1994

Srinivasan and Gupta, 1994

Jo et al., 1994 Jo et al., 1994

Asuke et al., 1994

Braun et al., 1994

particular the time the polymers were exposed to T > Tm; (vi) crystallization conditions, Tc, cooling rates, annealing, etc.; (vii) physical conditions of crystallization, viz. presence of nucleating impurities, discontinuities, e.g., gas bubbles or filler particles, confinement by molten, crystalline, or glassy phase, etc. As predicted by Helfand and Tagami, the interface is the locus of low molecular weight impurities that have been shown to nucleate crystallization, as for example in PS/PP blends [Wening et al., 1990]. Compatibilization by addition of a third component may either reduce or enhance the tendency for crystallization. On the one hand, addition of compatibilizer increases the

323

Interphase and Compatibilization by Addition of a Compatibilizer

interfacial area, thus it increases the nucleation rate — an increase of Tc by 10°C was reported [Wei-Berk, 1993]. However, the compatibilization also may increase thickness of the interphase, ∆l, thus hinder diffusion of nucleating agent to the crystallizable phase. It has been reported that when the size of domains containing the crystallizable polymer falls below certain limit, the crystallization rate is drastically reduced. In blends having widely different drop diameters this may lead to fractionated crystallization with large under-cooling required for the smallest drops. The presence of under-cooling reflects the difference in nucleating activity of the heterogeneity’s existing in the dispersed phase. When the drop size is reduced beyond certain limit it may no longer have a nucleus for the heterogeneous crystallization — it may crystallize by the homogenous mechanism at much lower Tc [Frensch et al., 1989]. Table 4.9 presents data on selected semicrystalline blends. More details on crystallization can be found in Chapter 3. Crystallization, Morphological Structure and Melting of Polymer Blends.

Table 4.9.

4.4.5

Mechanical Performance

Since blends are mainly used as structural materials, the most important properties are mechanical, especially the impact strength, stiffness, and elongation. Historically, blending was developed to improve these properties in the early resins, viz. PS, PVC, PMMA, PET, PA, etc. When blending started to involve mixtures of polymers, the impact modification has progressively changed into compatibilization. However, even today, many blends profit from the simultaneous compatibilization and impact modification by addition of multicomponent modifiers. Many patents and papers in the open literature address this question, some of which will be presented in this section and the following one on patented blends. Two types of mechanical tests have been used: the low rate of deformation: tensile, compressive or bending tests; and the high speed impact tests. Immiscibility of polymers is reflected in both. For example, in tensile tests the maximum strain at break (or the maximum elongation), and the yield stress (or the maximum strength) can be

Crystallization with and without compatibilizer

Blend

Additive

Technique and Results

References

PA-6/ EPR

EPR-g-SA

Martuscelli, 1984

PP/PA-6

maleated-PP (PP-MA)

PP/PS

P(P-b-S)

PP/PS

SBS

PP/HIPS

SB

PET/PPE

P(ET-b-S)

Compatibilizer had strong nucleating capability, reducing the spherulite size and increasing adhesion between the phases. Thermal analysis and optical microscopy. One or two crystallization peaks were affected by PP-MA. Tc of PA-6 initially decreased then leveled off with PP-MA content, whereas that of PP was not affected. Blends with PP-MA showed concurrent crystallization at the Tc of PP. Tc increased from 116 to 126°C upon addition of PP-b-PS. Its presence lowered ν12, reduced PS-drop diameter, and thus produced higher surface for nucleation of the α−PP crystals. The drop size of PP did not change, but Tc increased, owing to immiscibility of PP with polybutadiene block. SBS is poor compatibilizer, but capable to transfer nucleating heterogeneity’s to PP. WAXS and SAXS methods. Addition of a SB slightly reduced crystallinity of α-PP. Addition of compatibilizer reduced PPE drop size (from d = 5 to 0.2 µm) as well as it increased the total crystallinity.

Moon et al., 1994

Wei-Berk, 1993

Santana and Müller, 1994 Hlavata and Horak, 1994 Quirk et al., 1995

324

Abdellah Ajji

dramatically decreased by poor adhesion between the phases in the solid state. Similarly, this lack of adhesion is responsible for low values of impact strength — the specimens are brittle. Polymeric systems can be classified as either brittle or pseudo-ductile. The first type tends to fail by the crazing mechanism, and have low crack initiation as well as propagation energy — typical examples: PS, PMMA and SAN. The second type tends to fail by yielding, have high crack initiation energy, and low crack propagation energy — typical examples: PA, PEST and PC. As usual, there are some polymers, e.g., POM and PVC, which show an intermediate behavior. The transition from brittle to ductile behavior depends on the intrinsic performance characteristics of the material as well as on external variables, such as the test temperature, rate of testing, geometry, loading mode, etc. It is a common practice to toughen brittle resins by addition of elastomeric particles. The effectiveness of the process depends on their diameter and concentration. It has been found that at constant concentration of the toughening agent its effectiveness, i.e., the plot of toughness vs. particle diameter, follows a bell-shape curve, defining the optimum particle diameter. As shown in Figure 4.17 the optimum does not change with concentration. 150

The optimum particle size, dopt, was found to depend on the entanglement density of the matrix resin (see Figure 4.18). However, the dependence can be used only as a general guide. Both, the fracturing and toughening mechanisms change from one resin to the next. Determination of the dopt also is ambiguous owing to the polydispersity of sizes as well as to presence of other macromolecular chains. For example, it has been accepted that to toughen PS into HIPS the optimum diameter is defined as a diameter of the elastomeric particles expanded by occlusion of the PS matrix. In PVC the diameter of the elastomer was defined as the diameter of the original butadiene latex, before grafting it with styrene and methylmethacrylate. In PC the “optimum” diameter was defined rather from the point of view of availability of the core-shell toughening agent, than true optimum performance — it is becoming increasingly difficult to find tougheners with elastomeric particles having diameter d < 100 nm. Furthermore, the strategy of preparation of polymer blends with stable morphology demands that blends have a thick interphase, ∆l ” 60 nm. Frequently it is impossible to decide how far the toughening by rubber core extends to the interphase. Many impact resistant engineering resin blends have been formulated using core-shell multicomponent with rigid core and elastomeric shell whose thickness and affinity to the matrix resin can be adjusted. 1

0

10 PVC

15 phr MBS 10 phr MBS 5 0

mm ) do p t (m

2

NIRT (kJ/m )

10 100

PS

SAN

0

10

-1

10 PC PMMA

Brittle resins 0

0

0.2

0.4

Semi-ductile resins -1

-2

10

10 -2

mm ) Elastomer Drop Size, d(m

Figure 4.17. Izod impact strength at room temperature as a function of diameter of elastomeric particles in methylmethacrylate-butadiene-styrene copolymer used for toughening polyvinylchloride resin [after Bertelo and Mori, 1994].

10

-1

10

0

ne

10

Figure 4.18. Optimum elastomeric drop diameter vs. entanglement density of the matrix resin for brittle resins (polystyrene, styrene-acrylonitrile and polymethylmethacrylate), and for semi-ductile resins (polyvinylchloride and polycarbonate).

325

Interphase and Compatibilization by Addition of a Compatibilizer

By contrast with brittle resins where dopt is independent of concentration, in pseudo-ductile one these two variables are related — when the concentration of the toughening agent decreases the elastomeric particle size must be reduced. In other words, in the latter systems it is the distance between the elastomeric particles that seems to control the fracture mechanism. Again, there are resins showing intermediate behavior between these two limits. The above discussion focused on the addition of discrete elastomeric particles. However, there have been reports that formation of co-continuous structures of brittle or pseudo-ductile resin and an elastomeric one can produce a quantum jump in toughness, without greatly affecting the key engineering properties of the high performance resin. Commercial blends of this type, e.g., comprising POM, PA, PC, or PEST, are available. When the blends are formulated with relatively large content of two resins it may be necessary to toughen both phases. This is often done using sequential reactive blending. Examples of mechanical properties of blends toughened by addition of block copolymers are given in Table 4.10. More information on this topic is presented in the section on patented blends. 4.4.6

Solvent and Chemical Resistance

In many application the chemical and solvent resistance is of paramount importance. To improve

these properties, blending with a resin having the required performance characteristics is used. By nature, the modifying resins must have very much different chemical composition, what leads to antagonistic immiscibility and a need for compatibilization. The examples listed in Table 4.11 illustrate some commercial solutions. 4.4.7

Electrostatic Dissipating Blends

Polymer blends have been often used as electrical insulating materials. Polymers do not readily conduct electrical current, are inexpensive in comparison to other known insulating materials, are sufficiently durable and heat resistant. However, in some applications, owing to the accumulation of surface charge that may discharge rapidly and cause damage to electronic components, cause fires or explosions, they may pose problems. A need has existed for electrostatic dissipating polymeric compositions, ESD. The ESD compositions must have sufficient resistivity to cause slow dissipation of a static charge, but not too low as to allow the charge to move too quickly through the material, thereby causing an arc or spark. The surface resistivity 105 ” R ” 1012 Ω cm is considered desirable. Attempts have been made to coat an electrostatic dissipative material onto plastic, add graphite, or metal particles, but coating can easily be wiped out and the additives are expensive, make

Table 4.10. Examples of compatibilized and impact modified polymer blends Blend

Additive

PE/PA-6 PS/PP = 1:1 PPE/PE-ionomer

SEBS-MA, ” 10 wt% 5 wt% of low MW SEBS PS-b-PVP

PS/PVC/CPE

PP + 0-30 wt% PC HIPS/PBT

Comments

reactive blending, reduced ν12, increased adhesion three-fold increase of the impact strength in PPE matrix blends tensile strength and elongation increased #1. SEB-GMA, compatibilization depends on GMA-content and #2. (SEB-GMA)3, degree of chlorination of the CP; best #1, #3. random S-r-EB-GMA worst #3 SEBS (5, 10, and 20 wt%) improved impact, reduce tensile & flexural moduli PS-GMA reactive processing to form PBT-b-PS

Reference Armat and Moet, 1993 Appleby et al., 1994 Jo et al., 1994 Koklas and Kalfoglou, 1994 Srinivasan and Gupta, 1994 Chang et al.., 1994

326

Abdellah Ajji

Table 4.11. Blends with enhanced chemical and solvent resistance Blend

Additive

Comments

HIPS/HDPE PA/ABS PPE/PA

SBS or SIS SAN-MA SBR-MA, SGMA, SMA, or EMAA PPE-MA SEBS PC/EGMA maleic, fumaric, acrylic, or methacrylic acid grafted PP or SEBS AES-GMA, ABS-MA or ABS-GMA PO-GMA or EEA-GMA

high impact strength and resistance to solvents excellent solvent resistance and toughness solvent and impact resistance

PPE/PA PPE/PC /PEST PPE/PBT PP/PA

PEST/ABS

PET/PO

Reference

Swartzmiller et al., 1993 Lavengood et al., 1988 Ueno & Maruyama, 1982; Fuji & Ting, 1987 reactive extrusion Akkapeddi et al., 1988 improved impact and solvent resistance Hobbs et al., 1992 high solvent, and temperature resistance, dimensional stability Yates, 1987 low water absorption and permeability Nishio et al., 1990

outstanding heat, chemical, and impact resistance

Yasue et al.., 1989; Hirai et al.., 1988

high heat and moisture absorption resistance

Mukohyama, 1993; Natarajan et al., 1993

processing difficult, and often lower the mechanical performance of the plastic part. Incorporation of organic semiconducting materials, or other low molecular weight antistatic agents, also is disadvantageous. These additives must migrate to the surface and pose similar problems as coatings. For ESD the surface resistivity should be: 105 ” R ” 1012 Ω cm. The best strategy for development of such materials is by incorporation of organic semiconducting polymers. Selected examples are listed in Table 4.12.

chloride or polycarbonate with acrylonitrile-butadiene-styrene copolymer, PVC/ABS or PC/ABS, the styrene-acrylonitrile copolymer, SAN, ascertains adequate compatibilization in the systems. Note that ABS went through a series of process and composition modifications to enhance performance in blends. In this brief description of selected patented and/or commercial blends, first the commodity then the engineering and specialty materials will be presented in a tabulated form [Utracki, 1998]. Further information on the topic can be found in Chapter 15. Commercial Blends in this Handbook.

4.5

4.5.1

Patented Blends with Added Compatibilizer

Commodity Resins Blends

4.5.1.1 Polyolefin Blends Patented or commercial polymer blends are in most cases multiphase, compatibilized systems. In the old, but still popular blends of polyvinyl-

4.5.1.3 Vinyl and Acrylics Blends

327

Interphase and Compatibilization by Addition of a Compatibilizer

Table 4.12. Examples of ESD blends Blend

Additive

Comments

Reference

ABS/PMMA

5-25 wt% EO-CHR epihalohydrin copolymer 2-50 wt% EO-ECH

ESD, reduced delamination and improved ductility excellent ESD and elongation

Gaggar et al.., 1988

PS/PCL

Giles and Vilasagar, 1994

Table 4.13. Typical polyolefin polymer blends and their characteristics Blend

Additive

Comments

Reference

PP and/or LLDPE PCO/PE

EPR

enhanced mechanical performance low melt viscosity thus good processability, high elongation at break, impact strength, toughness, hardness, and modulus had excellent impact strength, good solvent resistance and aging characteristics improved impact strength and flame resistance, as well as low permeation to gasoline

PE/POM

Block copolymers with Ethylene or Propylene and norborene blocks Styrenic copolymers (SB, SBS, SEBS) Block copolymer: n-butyl terephthalate/ ethyleneglycol and propyleneglycol, Hytrel™4074, and a mixture of two aromatic phosphites EVAc

Yamazaki and Fujimaki, 1970, 1972 Epple and Brekner, 1994

PE and PP

EPR

PP/SBR

PP grafted with Styrene

PE/PPE PE/PET

PE or PP 3-block copolymer of styrene with PS or HIPS and butadiene or isoprene PP/PC

EVAc

PP with PET or PBT PP/LCP

polyolefin-polyester graft copolymer maleated PP

abrasion resistant, glossy moldings, high elongation, excellent impact strength, dimensional stability and high environmental stress cracking good flow properties and a very good low temperature impact strength. useful as stand-alone structural materials. good processability, stress cracking resistance, impermeability to water vapor, and improved impact resistance good mechanical properties and excellent solvent resistance excellent impact strength, flexural modulus, HDT, and peel resistance compatibilization improved the performance well above that observed for two-component LCP/PP blends

Haff and Lee 1978 Abu-Isa and Graham, 1993

Ishida and Masamoto, 1974

Schreck et al., 1994 DeNicola and Conboy, 1994 Hoenl et al., 1993

Giles and Hirt, 1986 Fujita (Y.) et al., 1994 Baird and Datta, 1992

328

Abdellah Ajji

Table 4.14. Typical blends with styrenic polymers and their characteristics Blend

Additive

Comments

References

PS or HIPS with polyolefins PS/PP PS/LLDPE

SEBS

Holden and Gouw, 1979 Grancio, et al., 1983 Seelert et al., 1993

PS/EVAc

SB

45-70 wt% HIPS with 15-40 wt% HDPE

5-25 wt% SBS or SIS

2-30 wt% PP, and 5-20 wt% SEBS. Good impact, flexural strength, and tensile yield retention was reported Thermoformable blends 10-90 wt% PS, 10-90 wt% LLDPE, and 5-40 wt% SEBS — the system showed good resistance to impact and yellowing Improved elasticity, tensile strength, and reduced die build-up. Further improvement by replacing a part of the EVAc with a low crystallinity PO Thermoformable blends with good interfacial properties for refrigerator or freezer liners. They are easy to process into parts with high impact strength and solvent resistance

(SB)n block copolymer SEBS

Djiauw and Mitchell, 1990; Djiauw, 1994 Swartzmiller et al., 1993

Table 4.15. Typical blends of vinyl and acrylic polymers Blend

Additive

PVDC or PVC with 10-25 wt% NBR 70-20 wt% CR PVC with NBR 5-13 phr of SBR PVC MBS PVC or PVCAc

PVC/SAN

5-50 wt% PVC with 1-50% ABS PVC PVC/ABS

PVC/ABS

Comments

References

Transparent blends, thermoplastic, heat-sealable, printable, Signer and Beal, odorless and tasteless, suitable for a wide range of applications 1953 The alloys showed improved impact strength at low temperature Bataafse Petroleum, 1960 Butadiene methyl methacrylate styrene copolymer, MBS as Feuer, 1958 an impact modifier. poly(B-co-BA-co-S), 3-40 wt% P(S-co-BA-co-S). The alloys were transparent, with de and Deguchi, or poly(S-cogood weatherability and high impact strength. 1971; Ide and AN-co-MMA-gMiura, 1971; BA-g-MMA) Ide et al., 1972 ABS-type G 10 phr of SAN and 10 phr of ABS-type G. Give resins with Parks and Jennings, “low sensitivity to mechanical working.” For vacuum 1956 forming the blend performance could be improved by addition of ” 20% PVCAc post-chlorinated PVC The alloys showed improved processability Kojima et al., 1970 impact strength, and thermal stability. 5-30 wt% MABS PVC with MABS showed high impact strength. Himei et al., 1967 ABS-core and grafted Excellent mechanical properties, high impact Tanaka et al., 1971 on it AN-co-Etstrength, transparency, surface smoothness acrylate-co-styrene. and whitening resistance ABS (core) grafted PMMA improved PVC/ABS miscibility. Michel, 1969 with PMMA The alloys displayed excellent mechanical properties and impact strength.

329

Interphase and Compatibilization by Addition of a Compatibilizer

4.5.2

Engineering and Specialty Resins Blends

Table 4.16. Engineering and specialty polymer blends Blend

Additive

Comments

Reference

PPE/PEST

SBS

Lee, 1977

PC/ACM

styrene-acrylic acid copolymer, SAA SMMA-MA

Low viscosity, good impact, fire-resistance and tensile strength High impact strength and HDT Co-continuous morphology

Jones and Mendelson, 1983 Kress et al., 1985

PC/ABS PC/MeABS

PC/PBT/LLDPE PPE/PA

AN-co-α-methyl styrene-co-MMA, MeSAN MBA

PA/PAr or PA/PEST

styrene-butadiene radial copolymer (K-Resin™) with citric acid or chloro-epoxytriazine; or tapered block copolymer PA-co-PAr

PA/PC or PA/PAr PA/PVDF

EGMA a compatibilizer

PA/PEI

phenolic compounds, e.g., a nonyl-phenolic

4.6

Good compatibility, toughness, and thermal stability Impact and solvent resistance Izod impact strength was improved (in comparison to blends with HIPS) by a factor of three.

Good processability, mechanical properties, UV stability, heat, solvent, and stress-crack resistance. Impact modification Suitable for parts that are impact resistant and impermeable to gases Have reduced water absorption, had low swell, and good dimensional stability

Conclusions and Outlook

To select proper compatibilizer, it is imperative to know whether the copolymer is capable to (i) engender a fine dispersion during blending, (ii) preferentially migrate to the homopolymers’ interface, (iii) stabilize the morphology against segregation, and (iv) enhance the adhesion between the phases. It is only when all these conditions are satisfied that the idea of interfacial activity of classical emulsifiers can be applied to copolymers added to immiscible blends.

Henton, 1977, 1986

Boutni and Liu, 1984, 1986 Gianchandai et al., 1993; Yates, 1993.

Dean, 1990.

Yuichi and Suehiro, 1989 Hizasumi, et al., 1989 Gallucci, 1988; Gallucci and Rock, 1992

Effective compatibilization of binary polymer blends by addition of a copolymer should reduce the interfacial tension coefficient. Often, it also alters the molecular structure of the interface (as measured by the scattering methods). The process is similar to the emulsification in the classical emulsions. The emulsifier effect on the droplet size follows generally the same behavior as the interfacial tension. This behavior is described by the emulsification curves (evolution of the particle’s size with the emulsifier content) and characterizes the additives’ efficiency. The shape

330

Abdellah Ajji

of the emulsifying curve depends greatly on the type of emulsifier and the emulsification process. If the selected compatibilizer is miscible only with one component of the blend, its efficiency is negligible. The efficiency of a block copolymer is limited by the formation of micelles in bulk phases as well as by the kinetic factors. The morphology of blends may be complex. Addition of compatibilizer not only affects the size and shape of the separated phases (the macromorphology), but it may also affect the crystalline form, the size of crystalline entities as well as the total crystallinity (the micro-morphology). In a blend of two semi-crystalline polymers, e.g., PE/PP — four phases may coexist: PE-crystalline, PP-crystalline, PE-amorphous and PP-amorphous. The situation can be more complex if the polymorphism of PP (with four potential crystalline structures) is taken into account. The spherulites’ size, lamellar thickness, interlamellar amorphous zones, perfection of the crystallites, total crystallinity, and a possibility of epitaxial crystallization, are additional factors that affect the final morphology and performance. Processing conditions, such as the extend of shearing, the cooling rate, and annealing history, all affect the melt and solid morphology. The addition of a compatibilizer multiplies the factors that must be taken into account. Polymer blends must provide a variety of performance parameters. Usually it is a set of performance criteria that determines if the material can be used or not. For specific application more weight can be given to one or another material property. The most important properties of polymer blends are mechanical. Two type of tests have been used: the low rate of deformation — tensile, compressive or bending; and the high speed impact. Immiscibility affects primarily the maximum elongation at break, and the yield stress. Polymers are either brittle or pseudo-ductile. The first type tends to fail by the crazing mechanism, and have low crack initiation as well as propagation energy, while the second tends to fail by yielding, have high crack initiation energy, and low crack propagation energy. The transition from brittle to ductile behavior depends on the

intrinsic performance characteristics of the material as well as on external variables, such as test temperature, rate of testing, geometry, loading mode, etc. To toughen a brittle polymer, elastomeric particles with optimum diameter, dopt, should be used. It has been found that dopt depends on the entanglement density of the matrix resin, dopt ∝ νne, where the exponent n = -2 to -3. For pseudo-ductile polymers a similar dependence is valid, but since for them the entanglement density is significantly higher, the required size of the elastomeric zones is small, in order of nm. Furthermore, it has been found that not only the size but also the separating distance between the elastomeric zones is important. Thus, for these polymers the toughening strategy is based on creation of proper size elastomeric interphase, with rigid core that provide proper separation. The ultimate in the toughening method for pseudoductile polymers is generation of co-continuous structure having proper size of the phases. For brittle polymers the impact strength can be improved by engendering proper size heterogeneity in the matrix — rubber particles that may or may not be joined to the matrix or even gas bubbles. Toughening is accomplished by increasing the number of crazes, thus increasing the total energy of fracture. However, this strategy rarely leads to maximum improvement of impact strength and as a rule it decreases several other performance parameters. Thus, for good balance of properties, e.g., improved processability, dispersion, impact strength and other mechanical properties, it is imperative that the impact modifier is well bonded to the polymeric phases, i.e., that it plays broader role of compatibilizer. In PO blends the preferred compatibilizer has been EPR, while in styrenic blends SBS or SEBS maintain high visibility in spite of the price. Both compatibilizers can also improve the impact strength. However, excessive amount of elastomer can lower the modulus and strength of the alloy, thus the elastomeric particles’ size should be optimized. The PVC blends are usually compatibilized by addition of an acrylic copolymer, viz. ABS, MABS, MBS, or a core-shell-type copolymer (consisting of ABS-core and grafted on it acrylo-

Interphase and Compatibilization by Addition of a Compatibilizer

nitrile-ethyl acrylate-styrene) was introduced by Japanese Geon in 1966. Owing to miscibility, PPE has been compatibilized by adding a styrenic copolymer, capable of interacting with the other polymer, e.g., either PS or HIPS, SEBS, or SBS. Since the most interesting second polymer was either PA or PEST, in the early 1980’s the technology moved to reactive blending. The use of PC that preferentially migrated to the interface between the dispersed drops of PPE and SEBS, dispersed in PBT matrix, is an excellent example of the use of the third polymer to prevent coalescence, thus to stabilize the optimized morphology. In several engineering blends, e.g., ABS/PSF, the nearly universal co-solvent, Phenoxy, has been used. However, majority of modern commercial blends is prepared by reactive compatibilization.

4.7

References

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Yamazaki, I. and Fujimaki, T., Jap. Pat., 007,627, 03 Mar 1970, Appl. 24 Dec 1969; 007,141, 29 Feb 1972, Appl. 26 Nov 1969, to Showa Denko K. K. Yasue, K., Marutani, T., Fukushima, Y., and Ida, T., Europ. Pat. Appl., 301,663, 01 Feb 1989, Appl. 26 July 1988, Jap. Appl. 1987, to Stamicarbon B. V. Yates, J. B., III, PCT Int. Appl., 005,311, 11 Sep 1987; Europ. Pat. Appl., 303,077, 15 Feb 1989; Europ. Pat. Appl., 550,208, 07 Jul 1993; Europ. Pat. Appl., 550,210, 07 July 1993, Appl. 31 Dec 1991, to General Electric Company. Yuichi, O., and Suehiro, S., Europ. Pat. Appl., 304,041, 22 Feb 1989, Appl. 17 Aug 1988, Jap. Appl. 1987, to Nippon Petrochemicals Company, Limited. Yukioka , S. and Inoue, T., Polym. Commun., 32, 17 (1991); Polymer, 35, 1183 (1994) Zin, W. C. and Roe, R. J., Macromolecules, 17, 183 (1984).

∆l’ LCST MA , MB, Mv Me Mn, Mw, Mz N Av Nv Q R

4.8

Nomenclature

4.8.1

Notation S

A total interface surface area a, b, c, k, K, n, u, equation constants a, bi segment or monomer length or lattice parameter a1, a2 long and short axis of a deforming particle CMC critical micelles concentration D domain period in copolymer D drop deformability parameter d droplet diameter deq equilibrium droplet diameter ∆ΕS elastic energy f fraction of A component in copolymer f* transition concentration from microphase to macro-phase separation g energy density ∆Gm Gibbs free energy of mixing ∆Hm heat of mixing kB Boltzmann constant l mean length of statistical skeletal segment ∆l, ∆lo interfacial thickness, and

t , t* T, Tg, Tm Tcr, Ts UCST V wcr wi Zi χ, χAB φ1, φ2 φd, φm

initial interfacial thickness interfacial thickness for infinite molecular weight lower critical solubility temperature molecular weight of monomer-A, monomer-B, and a statistical segment entanglements molecular weight number, weight, and z-average molecular weight Avogadro number number of statistical skeletal segments number of copolymer chains at the interface ideal gas constant, droplet radius end-to-end distance of the polymer chains radius of gyration of chain i interfacial area per unit volume of the blend time, and dimensionless breakup time temperature, glass transition temperature, and melting point critical and spinodal temperature upper critical solubility temperature specific volume minimum amount of copolymer to cover surface of spherical particle weight fraction of i degree of polymerization of polymer i (or copolymer) binary thermodynamic interaction parameter volume fraction of the dispersed and matrix phase, respectively volume concentration of dispersed phase and the matrix

Interphase and Compatibilization by Addition of a Compatibilizer

φp γ η, η’, η* [η] η1, η2 ηd, ηm ηi, ηSi, ηEi ηo, ηr κ = σd/ν12 κcr λ = η1/η2 µ ∆µ ν

ν12 νe νo ν’ ρ Σ σ σ12 sm σy τ, ty ω ζo 4.8.2 ABS APS CTBN

polymer volume fraction. shear strain shear rate shear viscosity, dynamic and complex viscosity intrinsic viscosity viscosity of the dispersed and matrix phase, respectively. viscosity of the dispersed and matrix phase, respectively. interfacial, shear interfacial and extensional interfacial viscosity zero-shear viscosity and relative viscosity capillarity number critical capillarity number viscosity ratio chemical potential excess chemical potential surface tension coefficient interfacial tension coefficient between phase 1 and 2 entanglements density initial interfacial tension interfacial tension for infinite molecular weight density interfacial area per copolymer joint stress shear stress stress for the matrix phase yield stress relaxation or residence time, and characteristic relaxation time angular frequency Rouse friction coefficient

Abbreviations acrylonitrile-butadiene-styrene aminopropyltriethoxysilane butadiene acrylonitrile copolymer with α, ω−carboxyl groups

CTBR CPE DSC EGMA EPDM EPR EVOH, EVAl FRES HIPS LLDPE MA PA PAr PB PBMA PBT PC PCL PE PEI PES PEST PET PETG PHB Phenoxy PMMA PO POM PP PPE PS PSF PVAc PVAl PVC PVDF PVP SAN

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butadiene oligomer with α, ω−carboxyl groups chlorinated polyethylene differential scanning calorimetry ethylene-graft-maleic anhydride ethylene-propylene-diene terpolymer ethylene-propylene rubber ethylene vinyl alcohol copolymer forward recoil spectroscopy high impact polystyrene linear low density polyethylene maleic anhydride polyamide polyarylate polybutadiene polybutyl methacrylate polybutylene terephthalate polycarbonate poly(ε-caprolactone) polyethylene Polyether imide Polyether sulfone polyester polyethylene terephthalate polyethylene terephthalate glycol polyhydrogenated butadiene polyhydroxyether of bisphenol-A polymethylmethacrylate polyolefin polyoxymethylene polypropylene polyphenyleneether polystyrene polysulfone polyvinyl acetate polyvinyl alcohol polyvinyl chloride polyvinylidene fluoride polyvinyl pyridine styrene-acrylonitrile

338

SANS SAXS SB SBS SEBS SEM SI SIS SMA SMMA TEM THF

Abdellah Ajji

small angle neutron scattering small angle x-ray scattering styrene-butadiene copolymer styrene-butadiene-styrene three block copolymer styrene-ethylene/butenestyrene three block copolymer scanning electron microscopy styrene-isoprene copolymer styrene-isoprene-styrene three block copolymer styrene-maleic anhydride styrene-methyl methacrylate block copolymer transmission electron microscopy tetrahydrofuran

CHAPTER 5

REACTIVE COMPATIBILIZATION OF POLYMER BLENDS

S. Bruce Brown

General Electric Global Research Center, Niskayuna, NY, USA

5.1

Introduction and Purpose

An increasing number of commercial polymer products is derived from blending two or more polymers to achieve a favorable balance of physical properties. As described in Chapter 2 of this Handbook, from the thermodynamic point of view there are two basic types of polymer blends: miscible and immiscible. The vast majority of polymer pairs are immiscible with one another. There are only few commercially important polymer blends based on miscible and partially miscible (i.e., miscible within a low range of concentration) polymer pairs. It is seldom possible to mix two or more polymers and create a blend with useful properties. Instead, when preparing a new polymer blend from immiscible resins, it is necessary to devise a specific strategy for compatibilizing the mixture to provide for optimum physical performance and long-term stability. Although there do exist a very small number of commercial blends of immiscible polymers that are not compatibilized, most commercially available blends of immiscible polymers have been compatibilized by some specific mechanism. The majority of polymer blends containing elastomeric, thermoplastic, and/or liquid crystalline polymers are processed by melt extrusion at some point in their history. After melt extrusion with intensive mixing, the morphology of an immiscible polymer blend on a microscopic scale will often consist of a dispersed phase of the more viscous polymer in a continuous matrix of the less viscous polymer (depending upon the relative amounts and viscosities of the two polymers in the blend). A good analogy from every-day experience is a dispersed mixture of viscous oil in an immiscible water matrix. The formation of optimum dispersed phase particle size and the long-term stabilization of blend morphology are critical if the blend is to have optimum properties and in particular good mechanical properties. If this morphology is not stabilized, then the dispersed phase may coalesce during any subsequent heat and high stress treatment (such as injection molding). Coalescence may result in gross phase segregation and delamination on a macroscopic scale and/or brittleness or poor surface appearance in the final molded part. Good analogies from everyday experience would be the separation on standing of a not stabilized oil-in-water dispersion into two separate liquid phases, or churning butter. Therefore, an important aspect of all compatibilization strategies is the promotion of morphology stabilization. Morphology stabilization L.A. Utracki (Ed.), Polymer Blends Handbook, 339-415. © 2003 Kluwer Academic Publishers. Printed in the Netherlands.

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may be provided by sufficient interfacial adhesion and/or lowered interfacial tension between the two polymer phases. Of the various compatibilization strategies that have been devised, an increasingly common method is either to add a block, graft, or crosslinked copolymer of the two (or more) separate polymers in the blend, or to form such copolymers through covalent or ionic bond formation in situ during the Reactive Compatibilization step. The first of these processes was described in Chapter 4 of this Handbook, Interphase and Compatibilization by Addition of a Compatibilizer, while the second method is the topic of this Chapter. The copolymer can reside at the interface between the dispersed and matrix phases, acting as an emulsifying agent that effectively stabilizes dispersed drops against coalescence and promotes interfacial adhesion in a solid state. In this manner, such a copolymer can both promote optimum dispersed phase particle size formation during compounding and prevent phase coalescence of the dispersed phase during any subsequent high stress and/or heat treatment processes. Again, the analogy from everyday experience is the addition of soap or other emulsifying agent to stabilize an oil-in-water emulsion. Often, as little as 0.5-2.0 wt% copolymer is sufficient to achieve morphology stabilization of an immiscible polymer blend. However, frequently higher amounts, for example as much as 10-20 wt% copolymer, may be necessary to obtain optimum physical properties of the blend, e.g., impact strength. The majority of commercially important, immiscible polymer blends rely for compatibilization on the presence of a copolymer of the blended polymers. Nowadays, such a copolymer is almost never synthesized in a separate step and then added as a distinct entity to the blend of immiscible polymers. Instead, a compatibilizing copolymer is most economically formed simultaneously with generation of interphase morphology during extrusion processing, a process referred to as Reactive Compatibilization. The Reactive Compatibilization process is logically a sub-category of the broader class of Interchain Copolymer Formation reactions performed by Reactive Extrusion [Brown, 1992],

because there are other commercial reasons for preparing copolymers of immiscible polymers aside from using them as in situ-generated compatibilizing agents for immiscible blends Copolymer formation by Reactive Compatibilization is a heterogeneous reaction taking place across a melt-phase boundary. Often this process occurs by direct reaction between chemical functionalities on some fraction of each of the two polymers. In some cases a third reactive species may be added to the blend to promote copolymer formation by one of several mechanisms. Reactive Compatibilization has at least two advantages: 1. First, the copolymer is automatically formed at the interface between the two immiscible polymers where it is needed to stabilize morphology. In contrast, when a compatibilizing copolymer is added as a separate entity to a polymer blend, it must diffuse to the polymerpolymer interface to be effective for promoting morphology stabilization and interphase adhesion. However, the added copolymer may prefer to form micelles as a separate phase that is useless for compatibilization. 2. A second advantage of in situ copolymer formation is that the molecular weight of each of the two distinct polymeric segments in the copolymer is usually the same as that of the individual bulk polymer phase in which the segment must dissolve. Even approximate molecular weight matching between copolymer segment and bulk phase results in optimum copolymer/bulk phase interaction for maximum interfacial adhesion. Only a small number of ways have been devised to form a compatibilizing copolymer during extrusion processing. Therefore, another purpose of this chapter is to identify and give examples of the distinct chemical strategies that have been used to form block, graft, or crosslinked copolymers as compatibilizers for immiscible polymer blends. The emphasis is on copolymer formation during melt reaction occurring during development of morphology. With few exceptions, the examples are limited to processes that require mixing in the

Reactive Compatibilization of Polymer Blends

molten state. This includes processes run in single screw or twin screw extruders, or similar continuous or semi-continuous processing equipment, as well as in batch mixers. The references in this review are strictly limited to journal articles. Although the majority of Reactive Compatibilization examples are found in industrial research, these are documented mostly in patents. The majority of patents have so few details that often only an educated guess is possible concerning the compatibilization strategy employed. Numerous examples of industrial compatibilization methods have been provided in a recent book based on the patent literature [Utracki, 1998]. It is not the purpose of this chapter to describe “compatibilization” of layers of immiscible polymers in laminates. Strategies similar to those used to compatibilize intimately mixed polymer blends have also been used to prepare stable laminates and, in those cases where a chemical reaction takes place between laminate layers, similar types of chemical reaction have been used. Nevertheless, laminate macroscopic morphology is essentially fixed and formation of stable laminates is better treated as an adhesion problem. It is also not the purpose of this chapter to summarize examples of “compatible” polymer blends formed in a solution step involving dissolution of the polymer components. In some cases such blends are only “pseudo-stable”, since they may not have been processed above the Tg of one or both of the components. Also, mixing in solution followed by devolatilization is rarely economical for practice in industry, particularly since many commercially important compatibilized polymer blends comprise at least one semicrystalline component (e.g., PA) which is poorly soluble in common solvents. There are included in the Tables a small number of examples of solution blended polymer blends when these complement similar examples prepared by melt processing. It is also not the purpose of this chapter to describe examples of compatibilized polymer blends formed by polymerization of a monomer in the presence of a second polymer. In these cases the growing polymer chain reacts with functionality on the second polymer to form a

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certain fraction of compatibilizing copolymer. This chapter attempts to classify the specific copolymer architectures formed by the different chemical reactions that have been described in the open literature for Reactive Compatibilization processes. The emphasis is on illustrating the scope of these particular reactions, and not on presenting every known example of a particular compatibilization strategy.

5.2

Definitions of Compatibilization and Polymer Alloys

As defined in Appendix 5 compatibilization means “A process of modification of interfacial properties of an immiscible polymer blend, leading to creation of polymer alloy”. A polymer alloy in turn is defined as “An immiscible polymer blend having a modified interface and/or morphology”, whereas a polymer blend is simply “A mixture of at least two polymers or copolymers”. In other words, all polymer alloys are blends, but not all polymer blends are alloys. A somewhat more elaborate definition of a polymer alloy would describe a blend of at least two immiscible polymers stabilized either by covalent bond or ionic bond formation between phases, or by attractive intermolecular interaction, e.g., dipole-dipole, ion-dipole, charge-transfer, H-bonding, van der Waals forces, etc. Thermodynamic compatibility describes a miscible polymer blend that displays a single glass transition temperature, Tg, within the full range of composition of the two polymers. For the purposes of this chapter we will be more concerned with Technological Compatibility. This term describes a polymer blend that does not separate into its individual components and does not lose useful technological properties over the expected lifetime of a molded part (which has been estimated by the Society of Plastics Engineers to be about 10 years) [Gaylord, 1989; Rudin, 1982]. Taking a somewhat different view, Coran has defined compatibilization as a process for improving ultimate properties by making polymers in a blend less incompatible [Coran and Patel, 1983 b].

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5.3

S. B. Brown

Types of Polymer Blends

Table 5.1 shows North American demand for common commercial polymer blends for 1992 and projected volumes for 1997 [Freedonia, 1993]. For 1997 the total North American polymer blends market was projected at 545 kton (1.2 billion pounds). Table 5.1.

North American demand for common commercial polymer blends

Type

Demand kton (mil lb.) 1992

Polyphenylene ether (PPE) Polyvinylchloride (PVC) Polycarbonate (PC) Acrylonitrile-butadienestyrene copolymer (ABS) Polyester (PEST) Polyamide (PA) Polyolefin (PO) Other resins

Common polymer building blocks arranged in a hierarchy of price and performance are shown in Figure 11.1 of Chapter 11 Commercial Polymer Blends in this Handbook. As the price gets higher, one is typically paying for higher heat stability and higher modulus. High Performance Thermoplastics such as PPS, PEI, and LCP, and Engineering Thermoplastics such as PPE, PBT, and PC have high heat stability and are often designed to take the place of metals in typical applications. Lower modulus, commodity plastics such as PE, EPDM, and modified styrenics have lower heat stability and are often used in applications requiring high flexibility.

1997

120 (263) 47 (105) 44 (96) 36 (80)

141 (340) 64 (140) 67 (147) 54 (118)

35 (78) 29 (64) 27 (60) 50 (110)

53 (116) 44 (97) 36 (80) 69 (152)

It has been estimated that polymer blends and alloys consume about 30 wt% of all manufactured polymers with about 9 % per annum growth in sales volume. The principal market for all blends is the automotive industry, which accounts for about 60 % of world consumption with about 9.3 % yearly growth. Substitution of plastic for metal in cars results in weight savings that improves gas mileage. Other important markets for polymer blends include computer and other business machine housings, electrical components such as connectors, appliances, consumer products, recreational equipment, construction and industrial applications. Commercial activity is mirrored by technological activity. It was estimated that roughly 87,000 patents appeared worldwide on all aspects of polymer blends between 1970 and 1987 averaging almost 5,000 patents per year [Juliano, 1988]. The pace appears to have slowed little since then.

Figure 5.1 Potential effect on polymer blend properties as component concentration changes.

The goal of combining two or more polymers such as pairs from those categories described above (e.g., an engineering thermoplastic plus a commodity polyolefin) is to achieve in the blend a combination of favorable properties from each polymer. Figure 5.1 shows idealized expected property combinations from blending two polymers that are either miscible (solid center line), immiscible and uncompatibilized (bottom line), or immiscible and compatibilized (top line). In the case of polymers that are miscible in all propor-

Reactive Compatibilization of Polymer Blends

tions one can only hope to obtain in their blend an average of their physical properties depending upon the proportion of each polymer present. In a common example the Tg of the blend will vary linearly from that of the lower Tg polymer to that of the higher Tg polymer as the higher Tg polymer increases in proportion in the blend. When two immiscible polymers are blended without compatibilization, one generally obtains a mixture with physical properties worse than those of either individual polymer. Usually such a blend has poor structural integrity and poor heat stability since there is no mechanism for stabilizing a dispersion of one polymer in a matrix of the other. On a macroscopic scale the blend may appear heterogeneous and in the extreme case delaminated. When two immiscible polymers are blended with compatibilization, one may expect a synergistic combination of properties from each polymer. A common example is a blend of a thermoplastic (to provide high heat stability) with an immiscible, rubbery impact modifier (to provide impact resistance), e.g., a rubber-toughened PA. A second common example is a blend of a semi-crystalline thermoplastic with an amorphous thermoplastic. Because of their semi-crystalline nature, polymers such as PA, PBT, PPS, or PP often have high chemical solvent resistance but low ductility, low dimensional stability, and low Tg. In contrast, amorphous polymers such as PPE, PEI, PC, and PE have poor solvent resistance and stress crack sensitivity but higher ductility, dimensional stability, and Tg. Often a binary blend includes a third, lower modulus polymer to provide optimum impact strength. A good example is a rubber-toughened blend of PPE with PA. In commercial examples PPE and PA are combined in amounts such that PPE is the dispersed phase and PA is the continuous, matrix phase so that the blend possesses adequate solvent resistance, e.g., to common solvents used in automotive applications. Again, because of its inherent solvent resistance, this type of blend is prepared by melt processing and cannot economically be prepared by combining the components in solution.

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Quite generally, the goal in any blending of polymers is to obtain one or all of the following benefits: higher heat distortion temperature (HDT), improved variable temperature impact resistance, solvent resistance, dimensional tolerance, higher flow, utilization of recycle/regrind, and lower cost.

5.4

Characteristics of Immiscible Polymer Blends

The general characteristics of immiscible polymer blends have been described in a large number of references (listed in the following section on Compatibilization Strategies). Commercial polymer blends are most often prepared by some form of processing in a molten state, usually extrusion. Among the factors that determine which polymer will be the dispersed phase and which the continuous matrix are the relative volume proportions and relative viscosities of the two polymers. During intensive mixing in a twin screw or single screw extruder the less viscous molten polymer in a simple 1:1 mixture of two polymers will form an easily deformable matrix while the more viscous polymer with form a difficultly deformable dispersed phase. Generally, the more viscous polymer will form the dispersed phase even in some cases when it represents more than 50 vol% of the blend. Blend properties strongly depend on which polymer is the continuous phase. The majority of commercially important compatibilized blends of semi-crystalline polymers with amorphous polymers are prepared with compositions such that the semi-crystalline component is the matrix and the amorphous component is the dispersed phase. The blends show adequate solvent resistance since in this morphology the surface consists largely of the dominant, matrix phase,. The formation of optimum dispersed phase particle size and the stabilization of the resulting blend morphology are critical if the blend is to have optimum properties and in particular good mechanical properties. Figure 5.2 shows a morphology generated by processing an uncompatibilized

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S. B. Brown

Figure 5.2 Morphology of PPE dispersed phase in a PBT matrix (a) as extruded and (b) after molding

blend of PPE dispersed phase in a PBT matrix [Brown, 1988]. Figure 5.2A shows that a reasonably uniform dispersion of PPE may be formed simply by suitable degree of mixing during extrusion. Figure 5.2B, however, demonstrates phase coalescence of PPE particles to form large, irregularly shaped islands when the extrudate of the uncompatibilized blend is molded at normal processing temperature. In this blend there is no interfacial adhesion between the two phases and, hence, no mechanism for morphology stabilization. Even in the presence of an impact modifier the resulting molded parts are quite brittle since there are no uniform dispersed phase particles of proper size to dissipate impact energy. Table 5.2 shows further examples of dispersed phase coalescence in blends of PA dispersed phase in a less viscous PE or PS matrix. The data show that the mean PA particle size increases dramatically with simple heating under static conditions in the absence of any mechanism for morphology stabilization. The same coalescence can occur in molded parts of uncompatibilized polymer blends subjected to further thermal treatment after molding (e.g., in a paint drying oven). The mechanical properties of these blends are quite poor.

Table 5.2.

Change of dispersed phase dimensions in uncompatibilized polymer blends upon annealing [adapted from White and Min, 1989]

Blend

PS/PA-6 (60/40) LDPE/PA-6 (50/50) LDPE/PA-11 (50/50)

Mean dimension after annealing (µm) 30 min

60 min

90 min

90 251 162

140 314 275

310 319 303

In summary, a frequent goal in making a technologically compatible blend of immiscible polymers is to stabilize an appropriate morphology of the dispersed phase polymer in the matrix polymer by promotion of interphase adhesion and/or by lowering the interfacial tension. These conditions are critical for providing good mechanical properties, toughness, molded part dimensional integrity, and maximum solvent resistance in the blend.

Reactive Compatibilization of Polymer Blends

5.5

General Strategies for Compatibilization of Immiscible Polymer Blends

Numerous general reviews have been published, which discuss compatibilization of immiscible polymer blends [see for example: Datta and Lohse, 1996; Folkes and Hope, 1993; Bussink and van de Grampel, 1992; Elmendorp and Van der Vegt, 1991; Menges, 1989; Utracki, 1989, 1998; Paul et al., 1988; Teyssie et al., 1988; Sperling, 1987; Fox and Allen, 1985; Solc, 1981; Rudin, 1980; Paul and Newman, 1978; Bucknall, 1977]. The basic strategies for compatibilization of two-phase polymer blends can be divided into at least three major categories. 5.5.1

Co-crystallization of Two Phases

This particular strategy is limited to those cases in which an immiscible polymer blend contains two semi-crystalline polymers that can co-crystallize. Nadkarni and Jog [1989, 1991] have reviewed examples of this type of compatibilized blend. Co-crystallization may also occur as a secondary process in an intimately mixed blend containing a copolymer with concomitant effects on blend properties as shown in a few of the examples of this review. 5.5.2

In Situ Immobilization of One Phase: Dynamic Vulcanization

In these cases a dispersed phase of a crosslinkable rubber is vulcanized in the presence of a matrix of a second, immiscible, non-vulcanizable polymer during the residence time of melt processing. Examples have also been reported in which a mixture of two vulcanizable polymers has been employed. Coran [1995] has summarized five key requirements for preparing optimum compositions by dynamic vulcanization: 1. Good match between surface energies of the dispersed phase and the matrix. 2. Low entanglement molecular length (high entanglement density) of the rubber. 3. Crystalline plastic matrix.

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4. Stable rubber and plastic at blend processing temperatures. 5. Availability of appropriate curing system for rubber under desired processing conditions. Coran and others have reviewed work in this area [Coran, 1995, 1990, 1987; Coran and Patel, 1996, 1995; Abdou-Sabet et al., 1995]. Most examples of dynamically vulcanized blends do not involve covalent bond formation between the immiscible phases. However, more recent work has shown that covalent bond formation between phases in conjunction with dynamic vulcanization of one phase can lead to blends with improved properties in certain cases. This is particularly true in blends where certain of Coran’s five key requirements above are not met. Examples of dynamically vulcanized blends that also feature copolymer formation between the two immiscible phases are summarized in this review article under the appropriate categories [see also Coran and Patel, 1996]. 5.5.3

Addition of a Third Material as a Compatibilizing Agent

5.5.3.1 Addition of a Separate Compatibilizing Agent A separate compatibilizing agent may be a third material not derived from either of the two immiscible polymers. Representative examples include certain plasticizers, random copolymers, and block copolymers, which may lower the interfacial tension between the two immiscible polymer components. Examples also exist where the separate compatibilizing agent is a chemically unreactive analog of one (or both) of the two immiscible polymers that has an attractive interaction with each polymer. In any case this is often a semi-empirical compatibilization strategy since precedent may be the only basis for choosing a compatibilizer. Reviews on addition of this type of compatibilizing agent as a separate component to an immiscible polymer blend have appeared [Paul, 1978; Xanthos, 1988; Gaylord, 1989]. See also the preceding Chapter 3 in this Handbook.

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5.5.3.2 Addition of a Copolymer of the Two Immiscible Polymers A copolymer of the two immiscible polymers themselves would seem to be ideally suited to act as a compatibilizing agent for an immiscible blend. If the copolymer is at the interface of the two phases, then the segments of the copolymer dissolve in the respective bulk phases of the same identity. The copolymer acts as emulsifying agent for the blend resulting in reduced interfacial energy and improved interphase adhesion. Table 5.3 shows dramatic examples of the stabilization of dispersed phase morphology in the presence of a compatibilizing copolymer. In all examples essentially no change in dispersed phase particle size occurs after annealing under static conditions for up to 90 min. The data shown in this Table 5.3 should be compared with those presented in Table 5.2, where the dispersed phase mean dimensions were presented for similar, uncompatibilized blends. Table 5.3.

Change of dispersed phase dimensions in compatibilized polymer blends upon annealing [adapted from White and Min, 1989]

Blend

PS / PA-6 / SMA (57/38/5) PE / PA-6 / PP-MA (47/47/5) PE / PA-11 / PP-MA (47/47/5)

Mean dimension after annealing (µm) 30 min

60 min

90 min

4 5 5

5 7 6

5 13 4

There are two basic options for addition of a copolymer compatibilizer to a blend of immiscible polymers. First, the copolymer can be synthesized in a separate step followed by addition to the blend. One disadvantage is that this requires a new product synthesis with expensive and timeconsuming process development, and, hence, a significant number of years before profitability, since scale of copolymer manufacture will be initially low.

A second and more important disadvantage is that adding the copolymer as a separate species to the blend requires that the copolymer diffuse to the phase interface of the immiscible polymers to be effective as a compatibilizer. Diffusion to the interface may not be efficient within the residence time of a typical extrusion blending process (usually 2-5 min). In addition, high concentrations of added copolymer may form micelles as a third, distinct phase that does not contribute to compatibilization. A third, potential disadvantage is that for optimum interfacial interaction a copolymer synthesized in a separate reaction step must have carefully controlled segment lengths to best match the molecular weight of the bulk phase in which the segment must dissolve. It is often desired to offer a series of commercial blends containing the same two polymers but with different molecular weights for the polymers in each blend. Any copolymer synthesis process would have to be capable of producing a series of copolymers with a variety of controlled segment molecular weights for optimum compatibilization efficiency. Commonly, the most economical and efficient process for adding a copolymer to a blend of immiscible polymers is to form the copolymer in situ by a chemical reaction during the extrusion process during establishment of the immiscible phase morphology — the process known as Reactive Compatibilization. • An obvious advantage of this strategy is that copolymer is made only as needed and a separate copolymer commercialization process need not be developed. • A second and most important advantage is that the copolymer is formed directly at the phase interface where it can serve as a compatibilizer, and no diffusion process dependent on extruder residence time is involved. • A third advantage is that the copolymer automatically has segment molecular weights similar to the molecular weights of the bulk phases in which the segments must dissolve, which should promote optimum interaction between copolymer and bulk phases. This also facilitates commercialization of a series of blends con-

Reactive Compatibilization of Polymer Blends

taining polymers with quite different molecular weight since the copolymer formed in situ will always have ideal segment molecular weights. • A disadvantage of forming copolymer in situ is that such a process often requires that each of the immiscible polymers bear an appropriate chemical functionality for reaction across a melt phase boundary. Reactive Compatibilization has been discussed in earlier reviews [Brown and Orlando, 1988; Tzoganakis, 1989; Brown, 1992; Liu and Baker, 1992a]. As practiced commercially, Reactive Compatibilization is a continuous extruder process with material residence time usually 1-5 min. Such a process permits large scale preparation of a polymer blend as needed (“Just-In-Time” inventory control). Reactive Compatibilization involves a heterogeneous reaction across a phase boundary. Such a reaction is limited by the interfacial volume available at this phase boundary. Most often, twin screw extruders (having screw diameter from 20 to >120 mm) are employed. The screws are designed using an appropriate sequence of screw elements and auxiliary conditions to promote generation of a large interfacial area for the desired chemical reaction to form copolymer. There are five basic processes for achieving interchain copolymer formation between two polymers during Reactive Compatibilization in an extruder. Table 5.4 shows these five processes starting with idealized homopolymers “A” with structure AAAAAAAA and “B” with structure BBBBBBBB. Each process produces a specific type of copolymer compatibilizing agent by particular types of chemical reactions. When it is desired to form a graft, block, or crosslinked copolymer, there are two distinct processes available for copolymer formation: 1. Direct Process, where reactive functionalities on each of the two immiscible polymers react with each other across the melt phase boundary. 2. Addition of a Third, Reactive Species to effect or promote copolymer formation. This situation is a typical “three-body” reactive extrusion problem requiring that three chemical

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species (at least two of which are immiscible) react within the short residence time of extrusion processing. In some cases the third, reactive species is simply a catalyst that activates functionality on one polymer for reaction with functionality on the second polymer. Such cases produce results similar to the Direct Process. Frequently, however, the added reactant is a coupling agent capable of reacting with each of the polymers individually (as opposed to reacting with only one of the two polymers). If the coupling agent is preferentially soluble in one of the two phases (e.g. a polar coupling agent in the more polar polymer phase), it may give predominantly homogeneous reaction instead of promoting heterogeneous reaction (copolymer formation) across the melt phase boundary. Hu et al. [1997] have studied one type of polymer melt reaction (carboxylic acid + epoxide) in which the kinetic efficiency depends upon the partition coefficient of reactant between two immiscible polymer phases. Such considerations must be applied to all Reactive Compatibilization processes involving 3 or more reactive species. Coupling agents are further discussed in Section 5.11.1. Examples of individual processes for forming copolymer compatibilizing agent described in Table 5.4 are summarized in the following sections.

5.6

Compatibilization by Redistribution to Produce Block and Random Copolymer: Process #1

Blends are discussed in alphabetical order of the polymeric components. 5.6.1

Types of Redistribution Reactions and Characteristics

Redistribution reactions (sometimes referred to as “transreactions”) occur by chemical interchange of block segments of one polymer chain for corresponding segments of a second polymer chain. Such reactions may be homogeneous (self-reaction)

348

Table 5.4. Type

S. B. Brown

Copolymer formation processes in extruder reactors [adapted from Brown, 1992] Process

Type of Copolymer Obtained

1 1a

Redistribution (“Transreaction”) Block + Random by reactive end-groups of one polymer attacking main AAAAABBBBB + AABBBBBAAA + AABBAAABBB, etc. chain of second polymer 1b by chain cleavage/recombination involving (same as from process a) each polymer Characteristics: • In theory, all chains of each polymer participate in the redistribution reaction. The reacting polymers can be diluted with different, non-reactive polymers. In process 1a the extent of reaction depends on concentration of reactive end-groups. • In processes 1a and 1b the initial reaction product is a block copolymer, which often is an effective compatibilizer. Further reaction leads to random copolymer with loss of phase separation and loss of properties associated with uniform sequence distribution of each polymer (e.g., crystallinity). • The reaction may be catalyzed. Preferably, the catalyst may be easily quenched or thermally degraded before a significant amount of the random copolymer is formed. 2 2a

2b

2c

2d

Graft Copolymer Formation by direct reaction of end-group of the 1st polymer with pendent groups of the 2nd polymer

by reaction of end-group of the 1st polymer with pendent group of the 2nd polymer in the presence of a condensing agent by reaction of end-group of the 1st polymer with pendent group of the 2nd polymer in the presence of a coupling agent (“c”)

by reaction of pendent groups of the 1st polymer with main chain of the 2nd polymer in a degradative process

Graft A A - BBBBB A + A A (same as from process a)

A A - BBBBB A A - BBBBB A

A A -c- BBBBB A A A

+

A A -c- BBBBB A A -c- BBBBB A

+

A A - BBB A + BBB + BB A - BB A

A A - BBB A A A

etc

etc

Characteristics: • Only chains bearing reactive functionality participate in the copolymer-forming reactions. A small amount of crosslinked copolymer may be formed when the end-group functionalized polymer has functional groups at both ends • Graft copolymer formation between two polymers, each with different functionality, may occur either by direct reaction (e.g., A-acid end-group + B-pendent epoxy group), or by addition of a condensing agent that activates functionality of one polymer for reaction with functionality of the second polymer (e.g., A-acid + B-alcohol + triaryl phosphite). • Graft copolymer formation between two polymers, each with the same type of functionality, may be obtained using a coupling agent that links the two groups (e.g., A-acid + B-acid + di-epoxide). The reaction may be inefficient if the coupling agent segregates into one phase, instead of concentrating at the interface. Most examples in this class have appeared in the patent literature. The coupling agent is incorporated into the final copolymer as a linking group. • Graft copolymer formation may occur by a degradative process through reaction between a pendent functionality on one polymer and main chain linkages on the second polymer. Small amounts of crosslinked copolymer may be formed if degradable segments of the graft copolymer react further.

349

Reactive Compatibilization of Polymer Blends

Table 5.4. Type

Continued Process

Type of Copolymer Obtained

3 3a

Block Copolymer Formation Block by direct reaction of end-group of the 1st polymer AAAAA-BBBBB with end-group of the 2nd polymer 3b by reaction of end-group of the 1st polymer with (same as from process a) end-group of the 2nd polymer in the presence of a condensing agent 3c by reaction of end-group of the 1st polymer with AAAAA-c-BBBBB end-group of the 2nd polymer in the presence of a coupling agent (“c”) 3d by reaction of end-group of the 1st polymer with main AAAAA-BBB + BB chain of the 2nd polymer in a degradative process Characteristics: • Only chains that bear reactive functionality participate in the copolymer-forming reaction. A-B-A block copolymer may result from the reaction if at least one polymer is functional at both ends. • Block copolymer formation between two polymers, each with different functionality, may occur either by direct reaction (e.g., A-acid end-group + B-epoxy end-group), or by addition of a condensing agent that activates functionality on one polymer for efficient reaction with functionality on the second polymer (e.g., A-acid + B-alcohol + triaryl phosphite). • Block copolymer formation between two polymers, each with the same type of functionality, may occur by adding a coupling agent capable of linking the two end-groups (e.g., A-acid + B-acid + di-epoxide). The method may be inefficient if the coupling agent segregates to one phase instead of concentrating at the interface. The coupling agent is incorporated into the final copolymer as a linking group. • Block copolymer formation may occur by a degradative process through reaction between end-group functionality on one polymer and main chain linkages in the second polymer. 4 4a

Crosslinked Copolymer Formation covalent crosslinking by direct reaction of pendent functionality of 1st polymer with pendent functionality of 2nd polymer

4b

covalent crosslinking by reaction of a pendent functionality of the 1st polymer with pendent functionality of the 2nd polymer in the presence of a condensing agent main chain of the 1st polymer reacts with main chain of the 2nd polymer in the presence of a radical initiator covalent crosslinking by reaction of pendent functionality of the 1st polymer with pendent functionality of the 2nd polymer in the presence of a coupling agent (“c”)

4c 4d

Crosslinked A B A - - B A B or A - - B A B (same as from process a)

(same as from process a) A B A -c- B A B or A -c- B A B

A B A A B ----- A A ----- B A A B A A B A

A A -cA A A

B B B B -cB

A A A A A

Characteristics: • Only chains bearing reactive functionality participate in copolymer-forming reactions. • Crosslinked copolymer formation between two polymers, each with different functionality, may occur either by direct reaction (e.g., A-pendent acid + B-pendent epoxide), or by addition of a condensing agent that activates functionality on one polymer for efficient reaction with functionality on the second polymer (e.g., A-pendent acid + B-pendent alcohol + triaryl phosphite) (Process 4a and 4b).

350

S. B. Brown

Table 5.4. Continued

• Crosslinked copolymer formation may also occur via direct reaction of mechano-chemically-generated free radicals of each polymer in the absence of added radical initiator, or through radical trapping by a reactive site, such as olefin or acetylene groups in the second polymer (Process 4a). Alternatively, a radical initiator may be added to generate radical sites on one or both polymers (Process 4c). A crosslinked copolymer will be formed if the involved polymers are not degraded. • Crosslinked copolymer formation between two polymers, each with the same type of functionality, may occur by adding a coupling agent capable of linking the two pendent-groups (e.g., A-pendent acid + B-pendent acid + di-epoxide) (Process 4d). The method may be inefficient if the coupling agent segregates to one phase instead of concentrating at the interface. The coupling agent is incorporated into the final copolymer as a linking group. • In all the above cases the initial copolymer formed is a graft copolymer but this product may react further to form crosslinked copolymer as additional functionalities of one or both of the two segments of the copolymer react. Type

Process

Type of Copolymer Obtained

5 5a

Ionic Bond Formation ion-ion association mediated by metal cations as linking agents (“c”)

5b

ion-neutral donor group association mediated by metal cations ion-ion association mediated by interchain protonation of a basic polymer by an acidic polymer

Block, Graft or Crosslinked A B A B A -c- B A -c- B A B or A B A -c- B A B -cA B A B (similar to structure from process 5a)

5c

A A -A A -A

B B B B B

or

A A A A A

A B A A B ----- A A ----- B A A B A A B A

Characteristics: • Only chains bearing reactive functionality participate in copolymer-forming reactions. • In theory, it is possible to form block, graft, or crosslinked copolymers by ionic associations. However, in practice, telechelic polymers with ionic functionality at the chain ends are uncommon. Therefore, the majority of reported examples involve crosslinked copolymer formation between two immiscible polymers bearing pendent ionic groups. • Ionic groups include carboxylic, sulfonic, and (less frequently) phosphonic acid groups. The acidic groups may be at least partially neutralized with either monovalent, divalent, or trivalent metal cations, e.g., Na+1, Zn+2, Al+3. Ionic crosslinking of immiscible polymers bearing acidic groups may be mediated by such metal cations or by low molecular weight di-basic molecules such as diamines. Both these agents may link ions of different polymers. • Masked ionomeric groups may also be used. They generate ions during melt processing. Examples include phosphonate esters that form phosphonic acid salt in the melt by transesterification with a salt, e.g., zinc stearate.

or heterogeneous. In the homogeneous case the molecular weight distribution of a polymer may reach equilibrium. In the heterogeneous case redistribution reactions can form a copolymer between two different polymers. This type of reaction has rarely been used to form polymer alloys in a

Reactive Compatibilization process for reasons discussed below. Redistribution reactions can occur by several different processes. In one common example nucleophilic end-groups of one polymer react with electrophilic linkages in the main chain of

Reactive Compatibilization of Polymer Blends

a second polymer resulting in chain cleavage. The initial product is a block copolymer of the two polymers along with a lower molecular weight fragment of the second polymer — see Table 5.4 (Type 1). Since the initial block copolymer can participate in further redistribution reactions, the net product after sufficient time may be a random copolymer. The propensity to form random copolymer is further increased if both polymers have nucleophilic end-groups and also electrophilic linkages in the main chain that can participate in redistribution. Redistribution during thermal processing is a common self-reaction in condensation polymers such as PA, PEST, and PC that often contain nucleophilic amine, hydroxy, or phenolic endgroups, along with electrophilic groups such as amide, ester, or carbonate linking the individual monomer units. Essentially all of the polymer chains in each of the immiscible polymers are capable of participating in the copolymer-forming reaction by redistribution. This is in contrast to many other processes for in situ copolymer formation where only those few chains bearing reactive functionality participate. Unless the redistribution process is carefully controlled, it is difficult to stop the process to make stable, compatibilized polymer blends. If the reaction is thermally initiated, the blend processing temperatures and residence times must be strictly and reproducibly controlled within narrow limits to achieve reproducible properties. For prolonged reaction times at a temperature above that necessary to initiate the reaction, one may obtain a broad distribution of block lengths and eventually random copolymer. The random copolymer may not be as efficient a compatibilizer for the immiscible polymer blends as the block copolymer initially formed. More importantly, a high degree of random copolymer formation may destroy desirable properties in the polymer blend such as crystallinity in one of the polymers and, hence, solvent resistance in the final blend. In the extreme case phase separation is lost and the mixture may become homogeneous and transparent.

351

The problem of controlling the redistribution process does not stop at the manufacturing stage. After a compatibilized polymer blend leaves the manufacturer, it undergoes further thermal histories such as molding or paint-oven drying at the processing facilities of the final user. Continued redistribution reaction in the hands of a final user may cause deterioration and nonreproducibility in blend properties. When applicable, a common method for controlling a redistribution process is to initiate the reaction with a catalyst. Control may then be achieved by quenching the catalyst at the desired extent of reaction. Certain types of redistribution catalyst may thermally decompose under controlled processing conditions that make quenching unnecessary. In these cases, a predominance of block copolymer may be formed that serves as an effective compatibilizer for an immiscible polymer blend. Just as importantly, only a relatively small proportion of the polymer chains actually participates in the redistribution process so that phase separation and the properties attributable to the original sequence distribution are maintained. The redistribution process is a degradative process for making a compatibilizing copolymer. A common feature of all redistribution processes to form copolymer between two different polymers is that the molecular weight of at least one segment of the initially formed block copolymer is less than that of the bulk polymer phase from which it is derived. Therefore, even when the redistribution process is carefully controlled to give predominantly block copolymer, the copolymer may not be as efficient a compatibilizer as a similar type of block copolymer formed by an end-group/ end-group reaction (see Section 5.8). With a low molecular weight block segment one may have poor penetration into the corresponding bulk polymer phase and less than optimum interfacial adhesion between immiscible phases with copolymer at the interface. As discussed in Section 5.4, an optimum interfacial adhesion is usually obtained when the segmental molecular weights of the block copolymer are similar to the molecular weights of the individual bulk polymer phases.

352

S. B. Brown

A general review of “Interchange Reactions Involving Condensation Polymers” describes early work on redistribution reactions in the melt blends of polyesters, polyamides, and polyester + polyamide [Kotliar, 1981]. 5.6.2

Polyamide + Polyamide Binary Blends

Examples of copolymer formation by redistribution reactions in PA/PA blends are given in Table 5.5. In related work Liu and Donovan [1995] failed to find evidence for transamidation in PA-6 blends with an aromatic polyamide during molding and annealing. 5.6.3.

5.7

Compatibilization by Graft Copolymer Formation: Process #2

Polyamide + Polyester Binary Blends

Examples of copolymer formation by redistribution reactions in PA/PEST blends are given in Table 5.6. 5.6.4

lymer [Dezhu et al., 1997]; Polyarylate + PET [Eguiazabal et al., 1991, 1986]; PC + PET, Polyarylate + PET and Polyarylate + PC [Fakirov et al., 1996 a, b]; PEN + PET [Ihm et al., 1996]; Polyarylate + PET [Kimura et al., 1984]; PEN + PET [Lee, S. C., et al., 1997]; PBT + LCP [Paci et al., 1987]; PC + LCP [Wei and Ho, 1997b]; PBT + PC [Wings and Trafara, 1994]; PEN + PET [Yoon et al., 1997]; PC + PET-CL copolymer [Zhang et al., 1995]; and PC + PET [Zheng et al., 1993, 1994 a, b].

Blends are discussed in alphabetical order of the polymeric components. 5.7.1

Types of Grafting Reactions and Characteristics

Polyester + Polyester Binary Blends

Brief reviews covering redistribution reactions in polyester and in polycarbonate binary blends have been prepared by Porter et al. [1989] and Porter and Wang [1992]. Selected references for redistribution processes in PEST/PEST blends are listed in Table 5.7. Early studies of these processes focused on measuring the extent of redistribution under specific processing conditions rather than on producing compatibilized polymer blends with an attractive balance of properties. A number of more recent studies have reported the limits of miscibility for certain melt-mixed polyester pairs in the absence of transesterification — see for example the NMR study of PC/PET blends [Abis et al., 1994]. Table 5.7 omits references in which transesterification in PEST/PEST blends is brought about under static conditions either by annealing or heating in a DSC chamber. Blends in such examples were prepared by solution precipitation. Selected recent references in this class include: PC + polyester urethane [Ahn et al., 1997]; PBT + PET [Backson et al., 1995]; Polyarylate + PBT [Denchev et al., 1996]; PCL + PBT-CL copo-

Graft copolymer formation has been the most common method of forming a compatibilizing copolymer between two immiscible polymers during Reactive Compatibilization. In the Direct Process this involves reaction between one polymer containing reactive sites along its main chain and a second polymer with reactive sites only at end-groups. Depending upon stoichiometry and concentration of functional groups, copolymer structures of the general type shown in Table 5.4 (Type 2a and 2b) are obtained. In this particular type of graft copolymer formation the average molecular weight of the copolymer is the simple sum of the average molecular weights of the two reacting species. A second, less common process for forming graft copolymer, as shown in Table 5.4 is Type 2d. In this process multiple reactive sites pendent on one polymer chain can bite into the linkages of a second type of polymer chain. The reaction creates a copolymer with average molecular weight less than the sum of the two initial reacting species. This is a degradative method for forming a compatibilizing copolymer. A common example is the transesterification reaction between the

353

Reactive Compatibilization of Polymer Blends

Table 5.5.

Redistribution reactions in PA/PA blends

Polyamide / Polyamide

Characterization and Comments

References

PA-46 (70-0)/PA-6I (30-100)

mini-extruder at 295-325°C or TSE at 315°C / DSC / influence of processing conditions on extent of transamidation film extrusion using SSE at 275°C/DSC study of heat-aged films/transamidation vs. time SSE at 290°C / mechanical properties/DMTA/DSC/NMR/effects of annealing and of different transamidation levels on properties thermal redistribution in extruder at 215-280°C/DSC

Eersels and Groeninckx, 1997, 1996

PA-6 (100-0)/poly(m-xylene adipamide) (0-100) PA-6 (100-0) / poly(m-xylene adipamide) (0-100)

PA-6 (95-5) / PA-66 (5-95)

Table 5.6.

Shibayama et al., 1995

Takeda et al., 1992 a; Takeda and Paul, 1991

Schott and Sanderford, 1977

Redistribution reactions in PA/PEST blends

Polyamide/Polyester

Characterization and Comments

References

PA-6 (50) / PET (50)

extrusion in capillary rheometer at 280°C then annealing / SEM / selective solvent extraction / FTIR / DMTA / WAXS thermal redistribution in internal mixer at 260°C / NMR / TLC / DSC internal mixer at 240°C / torque rheometry / SEM / GPC / mechanical properties / effects of mixing time / effects of PA amine and acid end-group concentrations internal mixer at 240°C / selective solvent extraction / NMR / TGA / MS / IV / model studies / kinetics / also data from solution reactions thermal redistribution in internal mixer at 240°C or SSE at 250°C / DSC / selective solvent extraction / SEC / FTIR / SEM vs. time / rheology / mechanical properties / DMTA low degree of thermal redistribution in internal mixer at 240-250°C / torque vs. time, temperature / DSC / solvent extraction internal mixer / FTIR / NMR / DSC / mechanical properties vs. degree of redistribution / H bonding between phases / toluenesulfonic acid catalyst (0.2) internal mixer at 265-295°C or TSE at 290-370°C / NMR / 23% max. copolymer level / toluenesulfonic acid catalyst (0.2)

Evstatiev et al., 1996; Serhatkulu et al., 1995 Konyukhova et al., 1994

PA-6 (70-0) / PC (30-100) PA-6 (90-0) / PC (10-100)

PA-6 (80-20) / PC (20-80)

PA-6 (100-0) / PC (0-100)

PA-6 (75-25) / PC (25-75)

PA-66 (54-10) / PET (46-90)

PA-66 (0-54) / PET (100-46)

La Mantia and Valenza, 1994; Valenza et al., 1994 Montaudo et al., 1994

Gattiglia et al., 1992, 1990, 1989 a, b Eguiazabal and Nazabal, 1988

Pillon et al., 1987 a, b

Pillon and Utracki, 1984

354

Table 5.7.

S. B. Brown

Redistribution reactions in PEST/PEST blends

Polyester / Polyester

Characterization and Comments

References

PC (50) / PET (50)

internal mixer at 270°C / selective solvent extraction / NMR / GPC / lanthanide catalyst comparison TSE at 290°C / DSC / SEM / capillary rheometry / WAXD / NMR / fiber tensile properties / dibutyltin dilaurate catalyst (0-0.1 wt%) TSE at 270°C / DSC / NMR / FTIR / optical microscopy

Fiorini et al., 1997

PET (100-90) / LCP (0-10) (Unitika LC3000) PBT (70-30) / PC (30-70) PC (100-0) / PET (0-100)

PC (100-0) or PET / LCP (0-100) (Hoechst Vectra B950) Polyarylate (90-10) / PETG (10-90) PEN (40) / PET (60) PC (70-30) / PBT (30-70) PC (50) / LCP (50) (Unitika LC3000 or Eastman X-7G) amide-modified PBT (30) / PC (70) / TPPite (0-2.2%) PC (70-30) / LCP (30-70) PC (90-60) / LCP (10-40) (Unitika poly(oxy-benzoateco-ethylene-terephthalate)) PBT (50) / PC (50) / Ti(OR)4 catalyst Polyarylate (50-10) / PBT (50-90) / Ti(OBu)4 catalyst (0.005-0.2%) PBT (25) / PET (75) / Ti(OBu)4 catalyst (0-1050 ppm) PET (100-20) / LCP (0-80) (poly(oxybenzoate-coethyleneterephthalate))

internal mixer or TSE at 270-280°C / DSC / SEM / selective solvent extraction / NMR / SEC / rheology / torque rheometry / catalyst comparison TSE at 280°C / DSC / annealing studies internal mixer at 235°C / torque variation / DSC / NMR / FTIR / selective solvent extraction TSE at 280°C / NMR / time-resolved light scattering measurements during annealing TSE at 270°C / DSC / NMR / copolyester content vs. annealing time / miscibility vs. copolymer content internal mixer at 300°C / SEM / DMTA / selective solvent extraction / NMR / rheology / Ti(OBu)4 catalyst (0-0.3 wt%) mini-TSE at 270°C / DMA / DSC / NMR / amine end-group titration / tri-phenyl phosphite (TPPite) additive internal mixer at 310°C / DSC / NMR / FTIR (Hoechst poly(p-oxybenzoate-co-p-phenylene-isophthalate)) internal mixer at 240-290°C / DSC / NMR / SEM / use of copolymer as compatibilizer for PC + different LCP / transesterification inhibition with TPP or catalysis by Ti(OBu)4 internal mixer at 230°C / SEM / DSC / SAXS / WAXS Mini-Max molder at 260-280°C / mechanical properties / DSC / NMR including model compounds

internal mixer at 275-280°C / 13C NMR determination of transesterification level / stabilizing effect of TPPite (0-5%) internal mixer at 275-293°C / NMR / SEM / DSC / degree of reaction vs. mixing time and ratio of LCP monomer units

Hong et al., 1997

Hopfe et al., 1997; Pompe et al., 1996 Ignatov et al., 1997 a, b; 1996; Fiorini et al., 1995 Lin and Yee, 1997 Oh et al., 1997 Okamoto and Kotaka, 1997 Pompe and Haubler, 1997 Stachowski and Dibenedetto, 1997 Van Bennekom et al., 1997

Wei et al., 1997 Wei and Ho, 1997a; Wei et al., 1996; Wei and Su, 1996; Su and Wei, 1995 Wilkinson et al., 1997 Eguiazabal et al., 1996; Fernandez-Berridi et al., 1995; Espinosa et al., 1993; Valero et al., 1990 Jacques et al., 1996 c Ou and Lin, 1996 a, b

355

Reactive Compatibilization of Polymer Blends

Table 5.7.

Continued

Polyester / Polyester

Characterization and Comments

References

PBT (100-0) / PCE (42 or 70% ester) (0-100) Polyarylate (100-0) / PET (0-100) PBT (50) /PC (50)

internal mixer at 250°C / DSC / selective solvent extraction

Rodriguez et al., 1996

injection molded mixture / torque rheometry / DSC / NMR / SEM / DMTA / mechanical properties TSE at 290°C / TEM / time-resolved light scattering / effect of transesterification on phase separation internal mixer or SSE at 240-250°C / DSC / selective solvent extraction / FTIR / rheology / toluene sulfonic acid (0.1-2) or Ti(OBu)4 or dibutyltin dilaurate catalyst /

Martinez et al., 1994, 1992

PC (90-70) / PCL (10-30)

PC (91-39) or PBT or Polyarylate / PPVL (9-61) PC (50) / LCP (50) (poly(oxybenzoate-coethyleneterephthalate)) PC (75-25) / PBT (25-75) PEN (100-0) / PET (0-100) PC (80-20) / LCP (20-80) (poly(4,4-dioxy-diphenyl-coiso-or terephthalate)) PET (43) / PC (57)

other catalysts ineffective mini-TSE at 280°C / selective solvent extraction / FTIR / DSC / Ti(OBu)4 catalyst (0-0.5 wt%) internal mixer at 260°C / Ti(OBu) 4 catalyst (0-0.5 wt%) / FTIR / GPC / SEM / use of copolymer as compatibilizer for PC + LCP “melt mixing” or solution precipitation / TGA / mass spectroscopy / viscometry SSE at 275-315°C / effect of multiple extrusion passes / NMR / DSC / effect of residual titanium catalyst levels internal mixer at 290°C / rheology / selective solvent extraction / FTIR / NMR / TEM / DSC

internal mixer at 270°C / FTIR / NMR / DSC / TGA / DMA / IV / model reactions Polyarylate (50) /PBT (50) TSE at 260-320°C / DSC / NMR / Ti(OBu) 4 catalyst (60-270 ppm) poly(hexamethylene terephthalate) Mini-Max molder at 260°C / selective solvent extraction / (67-20) / LCP (80-33) DSC / NMR / SEM / WAXS / DMTA / effects of added (poly(oxybenzoatephosphite stabilizer co-ethyleneterephthalate)) Polyarylate (50) / thermal redistribution in SSE < 280-325°C / DSC / FTIR / PC (20) / PET (30) stabilizers to prevent thermal redistribution PC (50) or stabilizers to prevent thermal redistribution / 31 Polyarylate / PET (50) or PBT P NMR study of stabilizer fate PC (90-10) / PET (10-90) thermal redistribution / DSC / FTIR / optical microscopy Polyarylate (15-85) / PC (85-15) thermal redistribution in SSE vs. TSE 260-300°C / DSC / mechanical properties Polyarylate (50) / PBT (50) capillary rheometry at 280-300°C / selective solvent extraction / FTIR PC (40) / PET (60) thermal redistribution in internal mixer 260°C / FTIR / DSC / DMA

Okamoto and Inoue, 1994 Shuster et al., 1994

Tijsma et al., 1994, 1993 Amendola et al., 1993

Montaudo et al., 1993 Stewart et al., 1993a Belousov et al., 1992

Berti et al., 1992 a, b Miley and Runt, 1992 Croteau and Laivins, 1990; Laivins, 1989

Cheung et al., 1989 Golovoy et al., 1989 Suzuki et al., 1989 Golovoy et al., 1987 Arruabarrena et al., 1986 Huang and Wang, 1986

356

Table 5.7.

S. B. Brown

Continued

Polyester / Polyester

Characterization and Comments

References

Polyarylate / PC

thermal redistribution in internal mixer 250-300°C; torque vs. time, temperature / DSC / DMA stirred batch reactor at 275°C / selective solvent extraction / selective degradation / IV measurements / optional addition of Ti(OBu)4 catalyst

Mondragon and Nazabal, 1986, 1985 Pilati et al., 1985

SSE at 265-270°C / thermal redistribution. on molding 260-350°C; time, temp. effects / DMA / DSC / selective solvent extraction SSE at 275-300°C / IV / DSC / use of arsenic trioxide to quench residual titanate catalyst

Robeson, 1985

PC (50) /PET (50)

Polyarylate (100-0) / PET (0-100) or PETG or PC PC (100-0) / copolyester of cyclohexanedimethanol and iso-and terephthalic acids (0-100)

poly(hydroxy ether) of bisphenol A (a Phenoxy resin) and a polyester. Pendant hydroxy groups on Phenoxy resin can undergo transesterification with ester linkages in the polyester chains resulting in graft copolymer formation accompanied by lower molecular weight polyester fragments. Because the molecular weight of the grafted polymer species is less than that of the homopolymer from which it was derived, the grafted chain segments may be below optimum molecular weight necessary for most efficient chain entanglement with the remaining homopolymer phase. This may result in less than optimum physical properties. Furthermore, if there is a large number of pendent reactive sites, then degradation of the second polymer may reach the point where it has lost the physical properties that made it useful for blending. Consequently, formation of a graft copolymer compatibilizing agent by a degradative process is not a common method for immiscible blend compatibilization. Many commercial thermoplastics for high impact strength applications are two-phase blends in which a higher modulus thermoplastic matrix is toughened by the presence of a lower modulus, dispersed phase polyolefin. In the majority of cases, such blends are compatibilized by graft

Smith et al., 1981

copolymer formation between at least some fraction of the chains of the two immiscible polymers. Rubber-toughened PA’s are the most common examples. In most of these cases PA amine end-groups react with pendent anhydride groups along the main chain of an immiscible rubbery polyolefin to form sufficient copolymer to compatibilize a dispersed polyolefin phase in a matrix phase of PA. Rubber-toughened PESTs are also often compatibilized through graft copolymer formation formed through reaction between polyester acid end-groups and epoxy-functionalized polyolefins. In rubber toughened thermoplastic blends the efficiency of compatibilization depends among other things upon sufficient concentrations of both thermoplastic reactive end-groups and polyolefin reactive pendent groups to give adequate levels of copolymer under the mixing, temperature, and residence time protocol of the extrusion process. Concentrations of thermoplastic reactive endgroups are usually controlled during the manufacturing process, e.g., through control of stoichiometry in condensation polymerization and/or through addition of reactive or unfunctionalized chain-capping agents either during synthesis or in a subsequent processing step.

Reactive Compatibilization of Polymer Blends

Functionality in rubbery polyolefins is usually introduced in one of two ways: 1. Copolymerization of olefin monomer(s) with another functionalized monomer (e.g., poly(ethylene-co-acrylic acid), poly(ethyleneco-glycidyl methacrylate)) 2. Graft-functionalization of polyolefin in a separate processing step (e.g., polyethyleneg-maleic anhydride, polypropylene-g-maleic anhydride). Graft-functionalization may be performed either by reactive extrusion in the molten state [Brown, 1992], in solution, or by solid state processes. In these cases, the concentration of functionality is controlled (among other factors) by temperature, physical phase of the polymer substrate, stoichiometry of functionalization agent and (optional) catalyst. When graft-functionalization is performed by extrusion, the removal of unbound functionalization agent from functionalized PO is critical for success of subsequent copolymer formation with reactive thermoplastic end-groups. Unbound functionality (e.g., free maleic anhydride) in the PO phase may tie up reactive end-groups of the thermoplastic resin during Reactive Compatibilization, making these end-groups unavailable for copolymer formation. For functionalization during a separate extruder grafting reaction unbound functionalization agent is removed by efficient devolatilization; for solution functionalization, by solubilization of excess functionalization agent during isolation of solid, functionalized polymer. The different categories of polymer blends compatibilized by graft copolymer formation are grouped below in categories of pairs of polymers, and in sub-categories of specific chemical reaction used to form the compatibilizing copolymer linkage. In all examples the predominant reaction leading to a compatibilizer is believed to be graft copolymer formation.

5.7.2

357

Polyamide + Fluoropolymer Binary Blends

Copolymer Formation by Amine + Carboxylic Acid Reaction Mascia and Hashim [1997] have prepared compatibilized blends of PA with PVDF by using carboxylic acid-functionalized PVDF. In an example 20 parts PA-6 was combined with 80 parts PVDFg-MAA (10% MAA) in an internal mixer at 240°C. The graft copolymer-containing blend was characterized by SEM, FTIR, mechanical properties, selective solvent extraction, and rheology. The effects of adding zinc acetate were studied. 5.7.3

Polyamide + Polyesteramide + Polyolefin Ternary Blends

Copolymer Formation by Amine + Anhydride Reaction Seo [1997] prepared compatibilized PA blends with LCP polyesteramide (Hoechst Vectra B950) in the presence of anhydride-functionalized polyolefin. Specifically, 60 parts PA-6 was mixed with 25 parts LCP and 15 parts EPDM-g-MA in a TSE at 290°C. The blend was characterized by SEM, optical microscopy, Raman spectroscopy, mechanical properties, selective solvent extraction, and FTIR. 5.7.4

Polyamide + Polyolefin Binary Blends (excepting polypropylene)

Blends in these sections include those with either a single PA or a mixture of semi-crystalline PA’s or a mixture of amorphous and semi-crystalline PA’s. An early paper by Wu [1985] described PA/PO blends in which properties were compared for blends containing either functionalized or unfunctionalized PO at 25, 15, and 10 wt% loading. The type of functionality was not specified, making this paper impossible to classify. Blends were characterized by morphology (particle size distribution) and mechanical properties. The level of interfacial adhesion required for toughening was discussed.

358

S. B. Brown

5.7.4.1 Copolymer Formation by Amide-ester Exchange Graft copolymer compatibilizing agents have been prepared through amide-ester exchange reaction between polyamide amine end-groups and pendent ester groups on polyolefin copolymers such as EEA (Table 5.8; see also Section 5.7.6.1). In these cases a low molecular weight alcohol is generated as a by-product. 5.7.4.2 Copolymer Formation by Amine + Carboxylic Acid Reaction Compatibilized PA/PO blends have been prepared by graft copolymer formation between the amine end-groups of PA and PO pendent carboxylic acid groups to give a new amide linkage (Table 5.9). Water is the by-product of this reaction. The work by Aharoni is noteworthy in that it employs a condensing agent to effect copolymer formation between polyamide end-groups and pendent carboxylic acid groups on EAA. Condensing agents are discussed more fully in Sections: 5.8.1; 5.8.2; 5.8.3.1; 5.8.11.1 5.7.4.3 Copolymer Formation by Amine + Anhydride Reaction The most common method for compatibilizing PA/amorphous polyolefin blends involves graft copolymer formation by reaction of polyamide

Table 5.8.

amine end-groups with pendent anhydride groups on an appropriate polyolefin to form a compatibilizing copolymer linked through an imide bond. Anhydride groups may be incorporated into suitable PO’s through grafting or copolymerization with maleic anhydride, citraconic anhydride, itaconic anhydride and congeners, or through grafting or copolymerization with latent anhydrides such as vicinal di-acids or acid-esters including fumaric acid, maleic acid monoalkyl ester, and congeners. Selected examples are listed in Table 5.10. 5.7.4.4 Copolymer formation by carboxylic acid or 1° amine + epoxide reaction Aerdts et al. [1997] prepared stabilized PA blends with a functionalized core-shell impact modifier having grafted epoxide groups in the shell. Specifically, 75 parts PA-6 was extruded on a mini-extruder with 25 parts SB core/MMAGMA shell impact modifier at 240°C. Blends were characterized by TEM. Dispersed phase agglomeration was examined as a function of epoxide concentration on the shell surface. Koulouri et al. [1997] prepared blends of either PA-6 (100-0 parts) or PA-11 or PA-12 or PA-612 or PA-610 with PE-g-GMA (0-100 parts) (optionally + HDPE) in a custom mixer at 230°C and characterized the blend using selective solvent extraction, SEM, DSC, DMA, tensile properties, and FTIR.

Graft copolymer formation in binary PA/PO blends by amide-ester exchange

Polyamide / Polyolefin

Characterization and Comments

References

PA-6 (85-25) / EEA (18% EA) (15-75)

melt mixing at 230°C / optical microscopy / selective solvent extraction / SEM / DSC / DMA / mechanical properties / FTIR internal mixer coupled to SSE at 260°C / SEM / torque rheometry / mechanical properties / water absorption / hardness / effect of premixing LDPE + LDPE-g-BA

Koulouri et al., 1997

PA-6 (100-75) / LDPE (0-25) / LDPE-g-BA (18% BA) (0-2.4)

Raval et al., 1991

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Table 5.9.

Graft copolymer formation in binary PA/PO blends by amine + carboxylic acid reaction

Polyamide / Polyolefin

Characterization and Comments

References

PA-6 (9-6) / PE (90) / EAA (11% AA; 40% Zn neutralized) (1-4)

PA + EAA extruded on TSE at 240°C followed by blow molding with 90% PE / DSC / SEM / rheology / permeation test / mechanical properties / selective solvent extraction + FTIR for PA-EAA copolymer characterization SSE at 210°C / mechanical properties / SEM / DSC / WAXD internal mixer at 250°C / torque rheometry / SEM / dispersed phase particle size vs. interfacial modifier concentration / emulsification curves / effects of mixing protocol / also blends containing PP in place of HDPE SSE / interfacial tension measurements / morphology before and after annealing / mechanical properties TSE at 250°C / use of PE + 10% EMAA ionomer masterbatch / rheology / SEM / also ribbon extruded blends internal mixer at 250°C / torque rheometry / SEM / selective solvent extraction / FTIR / DSC TSE / mechanical properties / DSC / SEM / MFI / also blends containing glass fiber TSE at ≥ 270°C / FTIR / NMR / DTA / GPC selective solvent extraction / morphology /amine, acid quantification / % chains grafted TSE at 230-280°C / rheology / SEM / dispersed phase particle size vs. interfacial tension / effect of component viscosity ratio internal mixer at 250°C / mechanical properties and morphology as function of MAA content / FTIR / selective solvent extraction / DSC / role of H-bonding SSE at 265-315°C / selective solvent extraction/ model compound reactions / FTIR / 13C NMR / 31P NMR

Yeh and Fan-Chiang, 1997; Yeh et al., 1997

PA-1010 (75) / EP (5-25) / EP-g-AA (1% AA) (0-20) PA-6 (90-10) / HDPE (10-90) / EMAA ionomer (0-30 parts)

PA-6 (75-25) /HDPE (25-75) / EMAA ionomer (0-5) PA-6 (40-20) /HDPE (60-80) / EMAA ionomer (2) PA-6 (100-0) / EMAA ionomer (0-100) PA-6 (90) / EAA (10% AA) (10) PA-6 (80-90) / EAA (20-10)

PA-66 (85) / acid-functionalized EP (15)

PA-6 (90) / EMAA (10)

PA-6 (90-80) / EAA (10-20) or EMAA ionomer / TPPite (1) or trialkyl

Xiaomin et al., 1996 Favis, 1994; Willis and Favis, 1988

Chen and White, 1993 Gonzalez-Nunez et al., 1993 Willis et al., 1993 Crespy et al., 1992 Braun and Illig, 1987

Wu, 1987

MacKnight et al., 1985

Aharoni et al., 1984; Aharoni, 1983

phosphite condensing agent

5.7.4.5 Copolymer Formation by Miscellaneous Reactions Miscellaneous compatibilization methods belonging to this category are listed in Table 5.11. Park et al. [1997] have claimed that PA blends with PO may be compatibilized by copolymer

formed in reaction between PA amine end-groups and isocyanate groups grafted onto PO. Li et al. [1993] have shown that PA-PO copolymer may be formed through displacement reaction between PA amine end-groups and benzylic bromide groups pendent on brominated poly(isobutylene-co-pmethylstyrene]. Coran and Patel [1983] suggest

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that PA-PO copolymer may be formed by a displacement reaction between PA amine end-

groups and PO chloride groups, simultaneous with dynamic vulcanization of the PO phase.

Table 5.10. Graft copolymer formation in binary PA/PO blends by amine + anhydride reaction Polyamide / Polyolefin

Characterization and Comments

References

PA-6 (75) / EPDM-g-MA (25) or EP-g-MA or ULDPE-g-MA PA-66 (75) / EPDM-g-MA (0.4-1% MA) (25) PA-6 (9-6) / PE (90) / HDPE-g-MA (1-4)

TSE at 270°C / mechanical properties / SEM / rheology / comparison to SEBS-g-MA impact modifier TSE at 255°C / mechanical properties

Burgisi et al., 1997

PA + HDPE-MA extruded on TSE at 240°C followed by blow molding with 90% PE / DSC / SEM / rheology / permeation test / mechanical properties / selective solvent extraction + FTIR for PA-HDPE copolymer characterization SSE at 235°C / DSC / SEM / WAXD / tensile properties / selective solvent extraction / interfacial tension estimates / also ternary blends containing EP (30%) SSE at 240°C or TSE at 280°C / torque rheometry / TEM / titration of residual amine groups / effects of extruder type, and PA mol. wt. and amine end-group concentration on morphology, mechanical properties and ductile-brittle transition temperature / effects of di- vs. mono-functional PA TSE at 280-300°C / DSC / DMA / rheology / mechanical properties / MFI / TEM / selective solvent extraction / effects of different extrusion sequences internal mixer at 240°C / SEM / effects of adding phthalic anhydride to consume amine end-groups / model reactions to assess roles of amine- vs. amide-anhydride reactions TSE / mechanical properties / SAXS / SEM / TEM / effects of rubber level and functionality / factors affecting ductile-brittle transition temperature / analysis of fracture surfaces TSE at 290°C / blends made from PA + masterbatch of 20 wt% EP-g-MA / SEM / deformation and fracture behavior vs. strain retention and rubber content processing study with SSE and TSE / TEM / mechanical properties vs. extrusion conditions and morphology / effects of using PA + rubber masterbatch internal mixer at 200°C / torque rheometry / selective solvent extraction / SEM study of morphology development in reactive and in non-reactive blends

Yeh et al., 1997

PA-6 (70-0) / LDPE (30-100)/ E-BA-MA (0-6) or EVAc-g-MA PA-6 (80) / EP (0-20) / EP-g-MA (1.1% MA) (0-20)

PA-6 (100-0) / aromatic PA-(0-100) / EP-g-MA (0-20) PA-6 (70) / EP (0-30) / EP-g-MA (0-30) PA-66 (100-80) / EPDM-g-MA (0-20)

PA-6 (95-80) / EP-g-MA (0.7% MA) (5-20) PA-6 (100-80) or PA-66 / EP-g-MA (1.5% MA) (0-20) amorphous PA (80) / EP-g-MA (0.7% MA) (20)

Roberts et al., 1997

Beltrame et al., 1996

Oshinski et al., 1996 a, b, c, d

Xanthos et al., 1996

Marechal et al., 1995

Muratoglu et al., 1995 a, b

Dijkstra et al., 1994

Majumdar et al., 1994g

Scott and Macosko, 1994

361

Reactive Compatibilization of Polymer Blends

Table 5.10. Continued Polyamide / Polyolefin

Characterization and Comments

References

aromatic PA (50) / hydrogenated NBR (50) / NBR-co-MA (30% MA) (0-10)

miniature mixer at 250°C / morphology / ellipsometry / L101 (0.9 phr) added for vulcanization of rubber phase / effects of mixing protocol / other carboxylated rubbers also used TSE at 310°C / SEM / mechanical properties / also blends containing polysulfone and PA-polysulfone copolymer internal mixer or TSE at 210°C / torque rheometry / mechanical properties / SEM / selective solvent extraction / DSC / simultaneous addition of PA + PE + MA + radical initiator internal mixer at 260°C / FTIR / mechanical properties / SEM TSE / mechanical properties / DSC / SEM / MFI / also blends containing glass fiber double-extruded in SSE at 240 or 280°C / mechanical properties / torque rheometry / lap shear adhesion / DMA / effects of mixing protocol TSE at 275°C / mechanical properties / melt viscosity / rheology / SEM / effects of functionalized PE mol. wt. / MA grafted to mixture of HDPE + LLDPE

Bhowmick et al., 1993

PA-46 (90) / EP-g-MA (1% MA) (10) PA-11 (20) or PA-6 / HDPE(-g-MA) (80)

PA-6 (80) / EP-g-MA (1.3-4.4% MA) (20) PA-6 (90) / EEA-MA (1.5% MA) (10) PA-6 (100-0) or PA-66 / EP-g-MA (1.2% MA) (0-100) PA-6 (50) / (88 HDPE + 12 LLDPE) (0-50) / (88 HDPE + 12 LLDPE)-g-MA (0.27% MA) (0-50) PA-6 (100-0) /HDPE (0-100)/ HDPE-g-MA (0.8% MA) (0-5) PA-6 (30-10) /HDPE (70-90)/ EPDM-g-MA (1.8% MA) (1, 3, or 5 parts) PA-6 (90-10) or PA-11 / HDPE (10-90) or MDPE / E-MA (0-10) PA-66 (80) / EP-g-MA (20) PA-6 (90) / EPDM-g-MA (10) PA-6 (90) / EPDM-g-MA (10) or EP-g-MA or PE-g-MA PA-6 (100-80) / EPDM-g-MA (0-20) PA-6 (90-80) / EP-g-MA (0-10) / unfunctionalized EP (10-20)

Koning et al., 1993a Lambla and Seadan, 1993, 1992

Abbate et al., 1992 Crespy et al., 1992 Oshinski et al., 1992 a, b

Padwa, 1992

SSE at 245°C / SEM / DSC / capillary rheometry

Kim et al., 1991

internal mixer at 250°C / DSC / DMA / optical microscopy / SEM / also binary blends containing 95 PA + 5 EPDM-g-MA internal mixer at 200-300°C / torque rheometry / capillary rheometry / SEM / selective solvent extraction / FTIR TSE at 280°C / morphology vs. unfunctionalized EP use / DSC / DMA / EELS SSE / selective solvent extraction / %N / FTIR / mechanical properties vs. MA level SSE / selective solvent extraction / mechanical properties / morphology / dilatometry tests / also used unfunct ionalized. poly(etherester) TPE SSE / 0.4 wt % MA / morphology / mechanical properties correlation with interparticle distance internal mixer at 260°C / mechanical properties vs. morphology

Kim and Kim, 1991

Serpe et al., 1990

Ban et al., 1989, 1988 Borggreve and Gaymans, 1989 Borggreve et al., 1989 a, b

Borggreve et al., 1988 a, b D’Orazio et al., 1988

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Table 5.10. Continued Polyamide / Polyolefin

Characterization and Comments

References

PA-6 (100-80) / EP-g-MA (0-20) / unfunctionalized EP (0-20) PA-6 (100-70) / EP-g-MA (0-30) / unfunctionalized EP (0-20) PA-66 (100-96) / EPDM-g-MA (0-4) PA-6 (100-70) / EP-g-MA (10-30) / unfunctionalized EP (0-30) PA-6 (100-80) / EP-g-MA (0-20) / unfunctionalized EP (0-20) PA-66 (100-60) / PE-g-MA (0-40)

internal mixer at 260°C / 0.6-4.5 wt% MA / morphology / mechanical properties

Greco et al., 1987

internal mixer at 260°C / mechanical properties vs. morphology TSE at 204-347°C / FTIR / DSC / SEM /mechanical properties vs. MA level (0-10%) internal mixer at 260°C / crystallinity by DSC / SAXS / WAXS / SEM / optical microscopy

Cimmino et al., 1986

Table 5.11.

internal mixer at 260°C / 2.9% MA / morphology / mechanical properties vs. EP-g-MA fraction TSE at 270°C / 0.4% MA / mechanical properties / adverse effect of blocking anhydride groups / morphology

Crespy et al., 1986 Martuscelli et al., 1985

Cimmino et al., 1984 Hobbs et al., 1983

Graft copolymer formation in binary PA/PO blends by miscellaneous reactions

Polyamide / Polyolefin

Characterization and Comments

References

PA-6 (90-10) / LDPE-g-isocyanate (10-90)

internal mixer at 245°C / SEM / rheology / mechanical properties / comparison to blends made with unfunctionalized LDPE / FTIR / PE grafted with 2 wt% reaction product of 2-hydroxyethyl methacrylate and isophorone diisocyanate TSE / TEM / SEM / optical microscopy / DSC / mechanical properties / fracture mechanism

Park et al., 1997

internal mixer at 225°C / selective solvent extraction / FTIR / mechanical properties / also blends containing m-phenylene bis-maleimide or trimethylolpropane tri-acrylate + RI

Coran and Patel, 1983a

PA-6 (100-70) / brominated poly(isobutylene-cop-methylstyrene) (2.3 mol. % p-methylstyrene; 0.7 wt% Br) (0-30) PA-6 (100-0) or PA-6,66,610 terpolymer or PA-69 / chlorinated PE (0-100) / L101 RI

Li et al., 1993

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Reactive Compatibilization of Polymer Blends

Polyamide + Polyolefin + Polypropylene Ternary Blends

5.7.5.2 PA-PO Copolymer Formation by Amine + Anhydride Reaction

5.7.5.1 PA-PO Copolymer Formation by Amine + Carboxylic Acid Reaction

PA/PP blends have been compatibilized through graft copolymer formation between polyamide amine end-groups and pendent anhydride groups on a functionalized polyolefin (Table 5.12).

5.7.5

Favis [1994] and Willis and Favis [1988] prepared compatibilized PA blends with PP and carboxylic acid-functionalized EMAA ionomer. Blends containing 90-10 parts PA-6, 0-30 parts EMAA ionomer, and 10-90 parts PP were combined in an internal mixer at 250°C and characterized by torque rheometry and SEM. Dispersed phase particle size vs. interfacial modifier concentration was determined. Emulsification curves were constructed. Effects of mixing protocol on blend properties were studied. Blends were also prepared containing HDPE in place of PP.

5.7.5.3 PA-PO Copolymer Formation by Carboxylic Acid or 1° Amine + Epoxide Reaction As shown in Table 5.13, PA/PP blends can be compatibilized through graft copolymer formation between polyamide amine end-groups and pendent epoxide groups on a functionalized PO. Since there are multiple epoxide sites on the polyolefin, some crosslinked copolymer may result as

Table 5.12 Graft copolymer formation in ternary PA/PO/PP blends by amine + anhydride reaction Polyamide / Polyolefin / PP

Characterization and Comments

References

PA-6 (30) / EP-g-MA (0-20) / PP (50-70)

internal mixer at 240°C / TEM / mechanical properties vs. volume fraction of compatibilizer

PA-6 (100-0) / EP-g-MA(1.1% MA) (0-20)/ PP ( 0-100) PA-6 (100-0) / E-BA-g-FA (0.4% FA) (10) / PP (0-100)

SSE at 240°C / mechanical properties / ductile-brittle transition temperatures / DMA / tensile dilatometry / SEM / TEM / effects of PP and PA molecular weight / effects of adding PP-MA TSE at 245°C / DSC / SEM / mechanical properties

Rösch et al., 1996; Rösch, 1995; Rösch and Mülhaupt, 1994 Gonzalez-Montiel et al., 1995 a, b, c Ikkala et al., 1993; Holsti-Miettinen et al., 1992

Table 5.13 Graft copolymer formation in ternary PA/PO/PP blends by carboxylic acid or 1° amine + epoxide reaction Polyamide / Polyolefin / PP

Characterization and Comments

PA-6 (30) or PA-66 / PP (70) / TSE at 250-280°C / SEM/interfacial tension measurements by E-BA-GMA (2.5% GMA) (3.5) imbedded fiber retraction / comparison to blend without compatibilizer / masterbatch prepared from PP + 5% E-BA-GMA PA-6 (100-0) / PP (0-100) / TSE at 245°C / mechanical properties / DSC / SEM E-EA-GMA (8% GMA) (0-10)

References Kirjava et al., 1995

Ikkala et al., 1993; Holsti-Miettinen et al., 1992

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S. B. Brown

well if the PA is functionalized at both ends. The proportion of crosslinked copolymer formed also depends upon blend composition and processing conditions. 5.7.6

Polyamide + Polyolefin + Styrene Copolymer Ternary Blends

5.7.6.1 PA-Styrene Copolymer Formation by Amide-ester Exchange Horak et al. [1997] prepared compatibilized PA blends with ABS in the presence of polyacrylate copolymer bearing pendent ester groups. A graft copolymer is obtained when PA amine endgroups undergo transreaction with the pendent ester groups (see Section 5.7.4.1). In an example 50-45 parts PA-6 was mixed with 45-50 parts ABS and 0-10 parts MMA-co-AN (80% MMA) in an internal mixer at 240°C. The blend was characterized by SEM, WAXS, and mechanical properties. The effects of premixing PA + MMAco-AN or of adding dibutyl tin dilaurate or Ti(OBu)4 catalyst were examined. 5.7.6.2 Copolymer Formation by Amine + Anhydride Reaction PA/PO blends have been compatibilized through graft copolymer formation between polyamide amine end-groups and pendent anhydride groups on a functionalized styrene copolymer. Also, in one example [Lu et al., 1995] this ternary blend is compatibilized through copolymer formation between PA and anhydride-functionalized polyolefin (Table 5.14). 5.7.7

Polyamide + Polyphenylene Ether Binary or Higher Blends

PA/PPE blends in these sections include those containing PPE-miscible PS or a functionalizedPS. These blends often contain a rubbery impact modifier, such as SEBS, as well.

5.7.7.1 Blends Containing Anhydride-Functionalized PPE PA-PPE Copolymer Formation by Amine + Anhydride Reaction Anhydride groups are readily introduced into polyphenylene ethers such as poly(2,6-dimethylphenylene ether) (PPE) by extrusion with maleic anhydride (MA) or congeners such as fumaric acid (FA), itaconic acid, citraconic acid, and related compounds. Evidence has been presented that the anhydride groups are predominantly located along the PPE main chain and with a fraction also at chain ends [Glans and Akkapeddi, 1991]. This functionality distribution depends on the type and level of functionalization agent, and on the mixing protocol during extrusion processing. PPE that has been melt-functionalized with FA or MA reacts with PA to give predominantly graft copolymer by reaction of anhydride pendent groups on PPE with amine end-groups of PA. There is certainly a proportion of block copolymer formation as well since a proportion of anhydride functionality is present at the terminal sites on PPE. In the examples given in Table 5.15, PPE was functionalized in a separate extrusion step before extrusion with PA. 5.7.7.2 Blends Containing Unfunctionalized PPE + Functionalized PS PA-PS Copolymer Formation by Miscellaneous Reactions Unfunctionalized PPE may be compatibilized with immiscible PA by addition of functionalized polystyrene capable of forming copolymer with PA (Table 5.16). This is a common compatibilization strategy for PPE blends since both PS itself and functionalized polystyrenes with a relatively low level of functionality are miscible with PPE. The examples below include use of anhydride-, acid-, and epoxide-functionalized polystyrenes, all of which are capable of reacting with nucleophilic end-groups on PA to form a graft copolymer.

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Table 5.14. Graft copolymer formation in ternary PA/PO/Styrene copolymer blends by amine + anhydride reaction Polyamide / Polyolefin / Styrene Copolymer

Characterization and Comments

References

PA-6 (10) / HDPE (80-90) / SEBS-g-MA (0-10)

TSE at 230°C / rheology / mechanical properties / DSC / SEM SSE at 240°C / TEM / mechanical properties TSE at 250°C / SEM / observation of morphology development along extruder screw axis SSE at 240°C / mechanical properties / SEM / optical microscopy / DSC SSE / interfacial tension measurements / morphology before and after annealing / mechanical properties

Chandramouli and Jabarin, 1995 Lu et al., 1995

PA-6 (78-55) / EP-g-MA (1.1% MA) (0-22) / EPDM-g-SAN (0-45) or ABS PA-6 (25) / LDPE (75) / SEBS-g-MA (0-5)

PA-6 (75) / LDPE (25) / SEBS-g-MA (2% MA) (0-15) PA-6 (75-25) / HDPE (25-75) / SEBS-g-MA (0-5)

Lim and White, 1994

Armat and Moet, 1993 Chen and White, 1993

Table 5.15 Graft copolymer formation in PA/PPE blends by amine/anhydride reaction Polyamide / PPE

Characterization and Comments

References

PPE-g-MA (0-3% MA) (49) / PA-6 (41) / SEBS (10) PPE-g-FA (60) / PA-6 (40)

TSE at 280°C / selective solvent extraction / mechanical properties / ductile-brittle transition temperatures / SEM / TEM SSE at 275°C / X-ray diffraction for compositional analysis and crystallinity level / comparison to blend made with unfunctionalized PPE

Campbell et al., 1990 Murthy et al., 1990

Table 5.16. Graft copolymer formation in PA/PPE/PS blends by miscellaneous reactions Polyamide / PPE / Polystyrene

Characterization and Comments

References

PA-6 (70-30) / PPE (30-70) / SMA (8 % MA) (0-10) PA-6 (70-50) / PPE (30-50) / S-GMA (0-10) PA-6 (70-50) / PPE (30-50) / SAA (11-46% AA) (0-3)

TSE at 290°C / FTIR / SEM / capillary rheometry / DMA / mechanical properties TSE at 280°C / torque rheometry / capillary rheometry / SEM / DMA / mechanical properties internal mixer or TSE at 280°C / rheology / SEM / mechanical properties / MFI / effect of different AA contents in SAA

Chiang and Chang, 1997 Chiang and Chang, 1996 Jo and Kim, 1992

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5.7.8

S. B. Brown

Polyamide + Polypropylene Binary Blends

5.7.8.1 Copolymer Formation by Amine + Carboxylic Acid Reaction As shown in Table 5.17, PA/PP blends can be compatibilized through graft copolymer formation between PA amine end-groups and pendent acid groups grafted at more than one site along the PP chain or with acid groups along a poly(acrylic acid) segment grafted at a single PP site. 5.7.8.2 Copolymer Formation by Carboxylic Acid or 1° Amine + Epoxide Reaction

formed also depends upon blend composition and processing conditions. 5.7.9

PA-Styrene Copolymer Formation by Amine + Anhydride Reaction As shown by examples listed in Table 5.18, PA/PP blends have been compatibilized through graft copolymer formation between polyamide amine end-groups and pendent anhydride groups on a functionalized styrene copolymer. 5.7.10

Zhang et al. [1997, 1996] and Xiaomin et al. [1997] have prepared compatibilized PA/PP blends by adding epoxide-grafted PP. Graft copolymers result from reaction between PA amine end-groups and pendent epoxide groups on PP (or with epoxide groups along a poly(GMA) segment grafted at a single PP site). For example, 100-15 parts PA-1010 was mixed with 0-85 parts PP and 0-25 parts PP-g-GMA in either an SSE or TSE at 200-210°C. The resulting blends were characterized by selective solvent extraction, SEM, rheology, DSC, ESCA nitrogen analysis, FTIR, mechanical properties, and peel test. Since there may be multiple epoxide sites on the polypropylene, some crosslinked copolymer may result if the polyamide is functionalized at both ends. The proportion of crosslinked copolymer

Polyamide + Polypropylene + Styrene Copolymer Ternary Blends

Polyamide + Styrene Copolymer Binary or Higher Blends

5.7.10.1 Copolymer Formation by Amine + Carboxylic Acid Reaction Melt reactions between amine-terminated PA and carboxylic acid groups on styrene-co-acrylic acid copolymer have been demonstrated (Table 5.19). The initial reaction product is a graft copolymer, but longer reaction times may result in a crosslinking, since the polyamides (PA-66 and PA-69) can have two amine end-groups per chain [Kuphal et al., 1991]. Mono-amine-terminated polyamides were shown in the same study to exhibit miscibility with SAA through hydrogen-bonding depending upon AA content.

Table 5.17. Graft copolymer formation in binary PA/PP blends by amine + carboxylic acid reaction Polyamide / PP

Characterization and Comments

References

PA-1010 (75) / PP (5-25) / PP-g-AA (5% AA) (0-20) PA-6 (85) / PP (7.5-15) / PP-g-AA (6% AA) (0-7.5)

internal mixer at 205°C / SEM / rheology / mechanical properties / selective solvent extraction / ESCA nitrogen analysis internal mixer or TSE at 235-255°C / selective solvent extraction / morphology /mechanical and viscosity properties / reaction kinetics from torque measurements internal mixer at 240°C / rheology / mechanical properties / SEM / DSC / effects of processing conditions TSE at 240°C / torque rheometry / DMA / SEM

Zhang and Yin, 1997

PA-6 (100-0) / PP (0-100) / PP-g-AA (2 or 6% AA) (0-6) PA-11 (100-0) / PP (0-100) / PP-g-AA (6% AA) (0-100)

Dagli et al., 1994

La Mantia, 1993 Liang and Williams, 1992

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Reactive Compatibilization of Polymer Blends

Table 5.18. Graft copolymer formation in ternary PA/PP/Styrene copolymer blends by amine + anhydride reaction Polyamide / PP / Styrene Copolymer

Characterization and Comments

PA-6 (100-0) / PP (0-100) / SEBS-g-MA (2% MA)(0-10)

TSE at 245°C / TEM / SEM / mechanical and viscosity properties / DSC / DMTA / fracture mechanical study

PA-6 (30) / PP (50-70) / SEBS-g-MA (0-20) PA-6 (100-0) / SEBS-g-MA (0-1.8% MA) (0-20) / PP ( 0-100) PA-6 (100-0) / PP (0-100) / SEBS-g-MA (2% MA) (10)

References

Heino et al., 1997a; Holsti-Miettinen et al., 1994, 1992 internal mixer at 240°C / TEM / mechanical properties Rösch et al., 1996; Rösch, and dispersed phase domain size vs. volume fraction 1995; Rösch and of compatibilizer Mülhaupt, 1994, 1993 SSE at 240°C / mechanical properties / ductile-brittle Gonzalez-Montiel et al., transition temperatures / DMA / TEM / SEM / effects of 1995 a, b, c rubber functionality level / effects of adding PP-MA / tensile dilatometry TSE / DSC / SEM Ikkala et al., 1993

Table 5.19. Graft copolymer formation in binary PA/Styrene copolymer blends by amine + Carboxylic acid reaction Polyamide / Styrene copolymer

Characterization and comments

References

PA-6 (100-87.5) / SAA (8% AA) (0-12.5)

SSE at 240°C / torque rheometry / mechanical properties / TEM / effects of adding 20% BA core + MMA shell impact modifier internal mixer at 220 or 265°C / melt flow / DSC two-roll mill at 230°C / SEM / optical microscopy / selective solvent extraction

Lu et al., 1994

PA-66 (50) or PA-69 / SAA (20% AA) (50) PA-6 (70-50) / PS (0-90) / PS-MMA-MAA (3-11% MAA) (0-30)

Kuphal et al., 1991 Fayt and Teyssie, 1989

5.7.10.2 Copolymer Formation by Amine + Anhydride Reaction

5.7.10.3 Copolymer Formation by Amine + Oxazoline Reaction

Immiscible blends of PA and PS have been compatibilized through graft copolymer formation between PA amine end-groups and anhydridefunctionalized styrene copolymer to form a new imide linkage (Table 5.20).

As shown in Table 5.21, blends of PA and PS have been compatibilized through graft copolymer formation between PA amine or acid end-groups and oxazoline-functionalized styrene copolymer.

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S. B. Brown

Table 5.20. Graft copolymer formation in PA/Styrene copolymer blends by amine + anhydride reaction Polyamide / Styrene Copolymer

Characterization and Comments

References

PA-6 (75) / SMA (20% MA) (0-5) / SB core + MMA shell impact modifier (20-25) PA-6 (75) / SEBS-g-MA (25)

mini-extruder at 240°C / TEM / comparison to blends containing core-shell impact modifier with GMA grafted onto shell TSE at 270°C / mechanical properties / SEM / rheology / comparison to PO-g-MA impact modifiers custom mixer at 260°C / TEM / SEM / DMA / mechanical properties / PC dispersed phase encapsulation by SEBS SSE at 240°C / mechanical properties / TEM / effects of different MA levels internal mixer at 250°C / mechanical properties / SEM / WAXS / SAXS / blends with SMA pre-reacted with 0-100% octadecylamine based on MA content TSE at 250°C / rheology / SEM / selective solvent extraction / FTIR / effects of processing conditions

Aerdts et al., 1997

PA-6 (75-25) / PC (25-75) / SEBS (0-20) / SEBS-g-MA (2% MA) (0-20 parts) or PS-g-MA PA-6 (80) / SEBS (0-15) / SEBS-g-MA (0.5-1.8% MA) (0-20) PA-6 (95-60) / SMA (5-40)

PA-6 (20) /ABS (80) / poly(N-phenylmaleimide-S-MA) (3% MA) (0-20) PA-6 (85-75) / SAN (25% AN) (12-25) / imidized acrylate copolymer (56% methyl glutarimide, 40% MMA, 2% MAA, 3% glutaric anhydride) (0-8) PA-6 (80) / PS-g-MA (0.08-0.18% MA) (20) PA-6 (80-70) or PA-66 / SMA (0-10) / SB core + MMA shell impact modifier (20) or BA core + MMA shell PA-6 (80) / SEBS (0-20) / SEBS-g-MA (0.5-2% MA) (0-20)

PA-6 (80-20) / PS (20-80) / PS-g-MA (0.08-0.68% MA) (5) PA-6 (78-55) / EPDM-g-SAN (0-45) or ABS / SEBS-g-MA (1.8% MA) (0-22) amorphous PA (20) / SMA (17% MA) (80)

TSE at 240°C / torque rheometry / SEM / morphology development in extruder vs. screw design and processing conditions / also PA-6 and PA-66 blends with SEBS-g-MA / titration of residual amine end-groups internal mixer at 230°C / SEM / tensile properties / rheology / FTIR / NMR / selective solvent extraction SSE at 240-280°C / torque rheometry / mechanical properties / ductile-brittle transition temperatures / TEM / effects of mixing protocol SSE at 240°C or TSE at 280°C / torque rheometry / TEM / titration of residual amine groups / effects of extruder type, and PA mol. wt. and amine end-group concentration on morphology, mechanical properties and ductile-brittle transition temperature / effects of di- vs. mono-functional PA internal mixer at 230°C / SEM / tensile properties / rheology / laminate butt joint test / effect of PS mol. wt. in PS-g-MA / effect of mixing sequence SSE at 240°C / TEM / mechanical properties internal mixer at 200°C or TSE or internal mixers / selective solvent extraction / SEM study of morphology development in reactive and in non-reactive blends / comparison of mixer efficiencies

Burgisi et al., 1997 Horiuchi et al., 1997 a, b, c ; 1996 Kayano et al., 1997 Kelnar et al., 1997

Lee, C. W. et al., 1997

Majumdar et al., 1997

Jo et al., 1996 Lu et al., 1996, 1993

Oshinski et al., 1996 a, b, c ,d

Park et al., 1996

Lu et al., 1995 Scott and Macosko, 1995; Sundararaj et al., 1995

369

Reactive Compatibilization of Polymer Blends

Table 5.20. Continued Polyamide / Styrene Copolymer amorphous PA / SMA (8-23% MA) / SAN (25% AN)

Characterization and Comments

References

melt-pressed, spin-coated film layers / ellipsometer Yukioka and Inoue, 1994 measurement for interfacial area vs. MA content and annealing time PA-6 (95-60) / TSE at 245°C / TEM / mechanical and viscosity Holsti-Miettinen et al., SEBS-g-MA (28% S / 2% MA) (5-40) properties / DMTA 1994, 1992 PA-6 (50-45) or PA-66 or PA-11 or SSE at 240°C or two different types of TSE / Majumdar et al., 1994 a, PA-1212 or PA-612 or PA-610 or torque rheometry / TEM / DSC / mechanical PA-6-66 copolymer/ ABS (45-50) / properties vs. extrusion conditions and morphology / imidized acrylate copolymer (56% methyl ductile-brittle transition temperatures / effects of glutarimide, 40% MMA, 2% MAA, PA amine end-group concentration 1% glutaric anhydride) (0-10) PA-6 (100-0) / ABS (0-100) / imidized SSE at 240°C / TEM / torque rheometry / mechanical Majumdar et al., 1994 b acrylate copolymer with various acid + properties vs. compatibilizer content, functionalization anhydride concentrations (0-20) level, rubber conc. and mixing protocol / ductile-brittle transition temperatures PA-6 (60) / ABS (36-40) / double-extrusion on SSE / mechanical properties vs. Majumdar et al., 1994 c SMA ( 25% MA) (0-4) SMA content / TEM PA-6 (100-80) or PA-66 or PA-11 or SSE at 240°C / torque rheometry / mechanical Majumdar et al., 1994 d, e PA-12 or PA-1212 or PA-612 or PA-610 properties / TEM / ductile-brittle transition temperatures / or PA-69 or PA-46 /SEBS (0-20) / interfacial tension estimates / effects of PA amine SEBS-g-MA (1.8% MA) (0-20) end-group concentration on copolymer formation (titration before and after extrusion) / torque rheometry PA-6 (100-80) or PA-66 or PA-12 or processing study with SSE and two different types of Majumdar et al., 1994 f PA-1212 or PA-6-66 copolymer / TSE / TEM / mechanical properties vs. extrusion SEBS (0-20) / SEBS-g-MA conditions and morphology / effects of using PA + (29% S + 1.8 %MA) (0-20) rubber masterbatch PA-6 (75) or PA-66 or PA-6-66 SSE at 240°C / torque rheometry / DSC / TEM / Majumdar et al., 1994 g copolymer / SAN (6-40% AN) (0-25) / mechanical properties / effects of AN content / imidized acrylate copolymer with effects of acid + anhydride concentration / effects of various acid + anhydride PA amine end-group concentration concentrations (0-25) PA-6 (100-85) / SEBS (0-15) / TSE at 260°C / SEM / WAXS / DMA / Wu et al., 1994 SEBS-g-MA (8-13 % MA) (0-4) mechanical properties PA-6 (80-20) / PS (20-80) / SSE at 230°C / DSC / capillary rheometry and Chen and White, 1993; SEBS-g-MA (0-5) or SMA morphology before and after annealing / mechanical Chen et al., 1988 properties/ interfacial tension measurements / comparison to blends containing SAN in place of SMA PA-6 (100-0) / ABS (0-100) / internal mixer at 220-240°C / DMA / rheology / SEM / Kim et al., 1993 SMA (8% MA) (0-10) effect of preblending of PA + SMA or ABS + SMA PA-6 (70-0) / ABS (30-100) / SSE at 260°C / mechanical properties / SALS / Misra et al., 1993 S-MMA-MA (25% MA) (0-8) optical microscopy / SEM

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S. B. Brown

Table 5.20. Continued Polyamide / Styrene Copolymer

Characterization and Comments

References

PA-6 (80) or PA-66 /SEBS (0-16) / SEBS-g-MA (0.5-2% MA) (0-20) PA-6 (100-70) / SBS (0-10) / SBS-g-MA (0.5-1.5% MA) (0-30) PA-6 (85) / SEBS (11-15) / SEBS-g-MA (0-4) PA-6 (100-0) or PA-66 / SEBS (0-100) / SEBS-g-MA (1.8% MA) (0-100)

TSE at 250°C / mechanical properties / TEM

Modic and Pottick, 1993

TSE at 240°C / rheology / solvent swelling / SEM / TEM / mechanical properties TSE at 260°C / SEM / WAXS / mechanical properties

Seo et al., 1993

double-extruded on SSE at 240 or 280°C / TEM / mechanical properties / torque rheometry / lap shear adhesion / DMA / DSC / effects of mixing protocol PA-6 (80) / PS (10-16) / internal mixer at 240°C / torque rheometry / SEM / SMA (2% MA) (4-10) selective solvent extraction / DSC / morphological stability to annealing / lap shear adhesion / comparison to PA-PS block copolymer compatibilized blends PA-6 (100-0) + poly(m-xylene adipamide) SSE at 260°C / torque rheometry / mechanical properties / (0-100) / SEBS (0-20) / SEBS-g-MA ductile-brittle transition temp. / TEM / DMTA / DSC / (1.8% MA) (0-20) effects of mixing protocol / effects of PA-PA transamidation on properties PA-6 (100-80) or PA-66 or PA-11 or SSE at 260-280°C / torque rheometry / mechanical PA-12 or PA-1212 or properties / TEM / effects of di-vs. mono-functional poly(m-xylene adipamide) / PA / effects of pre-extrusion of PA-6 with 10 wt% SEBS-g-MA (1.8% MA) (0-20) of different PAs PA-66 (90-10) / PS (10-90) / TSE / rheology / DSC / SEM / TEM / selective solvent SMA (11 or 25% MA) (0-0.5) extraction / FTIR / mechanical properties / also binary blends of PA + SMA PA-6 (100-0) / SMA (8% MA) (0-100) TSE at 240°C / rheology / SEM / DSC / mechanical properties / HDT / FTIR PA-6 (100-0) / SAN (0-100) or ABS / double-extrusion on SSE / torque rheometry / SMA (14 or 25% MA) (0-10) or mechanical properties / selective solvent extraction / imidized acrylate copolymer with 1% MA SEM / lap shear adhesion / SAN + SMA content or SAN-co-IPO (1% oxazoline) double-extruded in separate step followed by extrusion with PA PA-6 (100-70) / SAN (0-30) / Mini-Max molder at 230°C / SEM / mechanical SMA (15 or 30% MA) (0-7.5) properties / selective solvent extraction / FTIR PA-6 (80-20) / S-MAA (20-80) extrusion molding at 230°C / selective solvent extraction / DSC / mechanical properties

5.7.10.4 Copolymer Formation by Carboxylic Acid or 1° Amine + Epoxide Reaction Chang and Hwu [1991] prepared compatibilized PA/PS blends through addition of epoxide-func-

Wu et al., 1993 Oshinski et al., 1992 a, b

Park et al., 1992

Takeda et al., 1992a

Takeda et al., 1992b

Chang and Hwu, 1991

Kim and Park, 1991 Triacca et al., 1991

Angola et al., 1988 Ide and Hasegawa, 1974

tionalized S-GMA copolymer. A graft copolymer results from reaction of pendent epoxide groups with either amine or acid end-groups on PA. A blend of 50 parts PA-66, 50 parts PS, and 0-0.5 parts S-GMA (3% GMA) was prepared

371

Reactive Compatibilization of Polymer Blends

Table 5.21. Graft copolymer formation in PA/Styrene copolymer blends by amine + oxazoline reaction Polyamide / Styrene Copolymer

Characterization and Comments

References

Amorphous PA (99-80) / S-IPO (1% IPO) (1-20) Amorphous PA (20) / S-IPO (1% IPO) (80)

internal mixer at 210°C / SEM / interfacial fracture toughness / flexural properties / selective solvent extraction / FTIR internal mixer at 200°C / torque rheometry / selective solvent extraction / SEM study of morphology development in reactive and in non-reactive blends

Tan et al., 1996

on a TSE and characterized by rheology, SEM, and mechanical properties. Since there are multiple epoxide sites on S-GMA, some crosslinked copolymer may result if the polyamide is functionalized at both ends. The proportion of crosslinked copolymer formed also depends upon blend composition and processing conditions. 5.7.11

Polyester + Polyether Binary Blends

Degradative Copolymer Formation by Transesterification A graft copolymer may be formed through transesterification between pendent hydroxy groups on Phenoxy polyether resin and ester linkages in the chains of an immiscible polyester phase (Table 5.22). Since the product is a graft copolymer accompanied by a low molecular weight fragment from the polyester, this is a degradative copolymer forming process. The initial product of the transreaction is a graft copolymer as the alcohol bites into the polyester chain. Longer reaction time may result in a crosslinked copolymer since the pendent polyester segment is capable of further reaction with OH on a different Phenoxy chain. These types of blends have also been prepared by solution casting followed by annealing. Blends of LCP with Phenoxy resin provide an example [Kodama, 1992].

5.7.12

Scott and Macosko, 1995

Polyester + Polyolefin Binary Blends (excepting polypropylene)

5.7.12.1 Copolymer Formation by Alcohol + Anhydride Reaction As shown in Table 5.23, PEST/PO blends have been compatibilized through copolymer formation between polyester alcohol end-groups and pendent anhydride functionality on a polyolefin such a maleic anhydride-grafted polyolefin. Because this alcohol-anhydride reaction is reversible with the equilibrium lying on the side of unreacted anhydride, only a relatively small amount of copolymer may be formed. Consequently, the dispersed polymer phase may not be well stabilized against coalescence upon further thermal treatment [Sun et al., 1996]. Alternatively, at least some copolymer may be formed by a degradative mechanism through transesterification between polyester main-chain linkages and a low concentration of pendent acid groups in anhydridefunctionalized polyolefin (see Section 5.7.12.5). 5.7.12.2 Copolymer Formation by Carboxylic Acid + Epoxide Reaction Table 5.24 shows examples of PEST/PO blends that have been compatibilized through graft copolymer formation by reaction of polyester carboxylic acid end-groups with pendent epoxide groups on an appropriate polyolefin. The copolymer is

372

S. B. Brown

Table 5.22. Degradative graft copolymer formation in binary PEST/Polyether blends by transesterification Polyester / Polyether

Characterization and Comments

References

LCP (Hoechst Vectra A950) (90-0) / Phenoxy

Mini-Max molder at 290°C / rheology / DSC / SEM / mechanical properties / selective solvent extraction / FTIR internal mixer at 240°C / torque rheometry / DSC / DMTA / selective solvent extraction / also PC + Phenoxy binary blends internal mixer at 230 or 250°C / torque rheometry / selective solvent extraction / FTIR / DSC / DMTA / mechanical properties internal mixer at 200-250°C; torque vs. time, temperature / DSC / DMA / improved modulus and tensile strength as copolymer forms internal mixer at 230, 250, or 270°C / torque rheometry / DSC / mechanical and thermal properties SSE at 265-270°C / thermal redistribution during molding at 270-320°C / DMA

Choi et al., 1995

PBT (30-0) /PC (0-100) / Phenoxy (0-100) PBT (50) / Phenoxy (50)

Phenoxy (100-0) / PC (0-100)

Polyarylate (100-0) / Phenoxy (0-100) Polyarylate (70-20) / Phenoxy (30-80)

Remiro and Nazabal, 1991 a, b

Eguiazabal and Nazabal, 1990

Mondragon et al., 1988, 1986; Mondragon and Nazabal, 1987 Mondragon et al., 1987 Robeson, 1985

Table 5.23. Graft copolymer formation in binary PEST/PO blends by alcohol + anhydride reaction Polyester / Polyolefin

Characterization and Comments

References

PBT (100-20) / EVAc-g-MA (0.8% MA) (0-80) or EMMA-g-MA PBT (60) / LCP (Hoechst Vectra A950 copolyester) (25) / EPDM-g-MA (15)

TSE at 260°C / GPC / SEM / DSC / mechanical properties vs. use of unfunctionalized EVAc or EMMA TSE at 290°C / SEM / optical microscopy / Raman spectroscopy / mechanical properties / selective solvent extraction / FTIR TSE at 270°C / rheology / interfacial tension measurements TSE at 280°C / DMA / DSC / mechanical properties / SEM / FTIR / optical microscopy / use of recycle PET TSE at 270°C / SEM / DSC / optical microscopy / mechanical properties internal mixer at 240°C / mechanical properties / SEM / rheology / DSC / selective solvent extraction

Kang et al., 1997

PET /HDPE / HDPE-g-MA (5 parts per 100 parts PET) or EVAc-g-MA PET (100-0) / HDPE (0-20) / EMAc-MA (1.4% MA) (0-100) PET (85-20) / HDPE (10-75) / PO-g-MA (0-5) PBT (80) / EP-g-MA (2.3% MA) (20)

Seo, 1997

Ihm and White, 1996

Kalfoglou et al., 1995 Sambaru and Jabarin, 1993 Cecere et al., 1990

373

Reactive Compatibilization of Polymer Blends

joined through a new ester linkage. Since there are multiple epoxide sites on the polyolefin, some crosslinked copolymer may result if the polyester is functionalized at both ends. The proportion of crosslinked copolymer formed also depends upon blend composition and processing conditions. When the new ester linkage is formed, a secondary alcohol group forms as well. The new secondary alcohol may also react with polyester mainchain linkages to form copolymer through a degradative transesterification process. Consequently, two very different types of copolymer may be formed in the blends of this section. If some frac-

tion of copolymer does indeed form by a degradative mechanism in these blends, there does not seem to be very much information about its relative amount or its relative efficiency for promoting compatibilization. Copolymer formation between PEST and PO may also be combined with subsequent dynamic vulcanization of the PO phase in either the same or in a separate processing step [Okamoto et al. 1994; Moffett and Dekkers, 1992]. Perret et al. [1996] have encapsulated a third polymer within a PO phase dispersed in a PEST matrix — the third polymer was PA-66, having higher Tm than the matrix.

Table 5.24. Graft copolymer formation in binary PEST/PO blends by acid + epoxide reaction Polyester / Polyolefin

Characterization and Comments

References

PBT (100-85) / EPDM-g-GMA (0-15)

internal mixer at 245°C / SEM / mechanical properties vs. use of unfunctionalized EPDM Buss ko-kneader / pre-compounding of PA + E-MAc-GMA at 270°C followed by compounding with PBT at 250°C / SEM / mechanical properties SSE or TSE / mechanical properties / TEM

Wang et al., 1997

TSE at 240°C / DSC / SEM / FTIR / rheology / mechanical properties / comparison vs. properties of ternary EEA-GMA blends with PBT or LCP + PP TSE at 280°C / DMA / DSC / mechanical properties / SEM / FTIR / optical microscopy / use of recycle PET TSE at 280°C / mechanical properties / SEM / rheology / properties comparison to unfunctionalized rubbers / effects of amine catalyst addition internal mixer or TSE at 270-290°C / DSC /selective solvent extraction / FTIR / rheology / mechanical properties / morphology vs. order of mixing TSE at 260°C / adipic acid (0.1) optionally added as rubber vulcanization agent / mechanical properties / TEM internal mixer at 270°C / torque rheometry / GPC / effects of residual levels of different catalysts used to synthesize polyesters TSE at 250°C / extrusion of PBT + EPDM -g-GMA followed by re-extrusion with vulcanization agent (either L130 or 6-aminohexyl carbamic acid) / mechanical properties / SEM / TEM

Holsti-Miettinen et al., 1995

PBT (80-60) / E-MAc-GMA (12-24) / PA-66 (8-16) PET (60) or PC + PE / EEA-GMA (25) or E-GMA / EEA (15) or EP PBT (100-0) or LCP (Hoechst Vectra A950) / EEA-GMA (0-100) PET (100-0) / HDPE (0-20) / E-GMA (0-100) or EEA-GMA PET (80-70) / EEA-GMA (20-30)

PET (20-18) / HDPE (73-78) / E-GMA (2-9) PBT (80-50) / E-GMA (3% GMA) (20-50) PET (50) or PETG / E-GMA (6% GMA) (50) PBT (50) / EPDM-g-GMA (1.5, 3 or 6% GMA) (50)

Perret et al., 1996

Akkapeddi et al., 1995

Kalfoglou et al., 1995 Penco et al., 1995

Dagli and Kamdar, 1994 Okamoto et al., 1994 Stewart et al., 1993b

Moffett and Dekkers, 1992

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In this ternary blend the epoxy-functionalized PO was capable of reacting with terminal functional groups on both of the two other polymers. The blend was formed either by pre-extrusion of PO with PA, followed by extrusion with PEST, or by feeding PO and PA to the feed-throat of an extruder, then adding PEST downstream. The morphology showed a PEST matrix in which shells of PO surrounded cores of PA. The blend was an impact modified PEST with higher impact strength than the corresponding blend containing epoxy-functionalized PO alone. One reason for the improvement may be that the volume fraction of PO impact modifier was effectively increased by inclusion of PA core.

5.7.12.3 Copolymer Formation by Carboxylic Acid + Oxazoline Reaction PEST/PO blends have been compatibilized through graft copolymer formation by reaction of polyester carboxylic acid end-groups with pendent oxazoline groups on an appropriate PO. The copolymer contains a new esteramide linkage. Wörner et al. [1997] have blended 0-20 parts oxazolinefunctionalized rubbers containing more than one oxazoline group per chain with 100-80 parts acidterminated PBT in an internal mixer at 240°C. Oxazoline-functionalized B-AN or E-B-AN was used. The copolymer structure was initially a graft copolymer. However, if additional oxazoline groups on the rubber react with additional PBT acid groups, then a crosslinked copolymer structure can arise. The blends of Wörner et al. were characterized by torque rheometry, mechanical properties, DMA, SEM, DSC, and level of oxazoline groups on the rubber (fraction of nitrile groups on AN-containing rubber that had been converted to oxazoline groups). Improved blend properties were obtained through addition of a bis-oxazoline chain extender for PBT. 5.7.12.4 Copolymer Formation by Diels-Alder Reaction The benzocyclobutene functionality can thermally ring open to give a highly reactive o-quinodimethane

intermediate that can function as the diene partner in a classical Diels-Alder ring-forming reaction with an olefin. When benzocyclobutene and olefin are located on two immiscible polymers, the reaction can in theory lead to copolymer formation. Dean [1993] has postulated that copolymer is formed when 50-15 parts benzocyclobutene-terminated polyarylate is mixed with olefin-containing EPDM in an internal mixer at 265°C. Blends were characterized by mechanical properties and DMTA. 5.7.12.5 Degradative Copolymer Formation by Transesterification Polyester-polyolefin copolymer compatibilizers have been made through catalyzed or thermal transesterification of polyester main chain ester linkages with pendent ester or acid groups in a polyolefin copolymer such as EVAc or EMAA (Table 5.25). In a separate example polyesterpolyolefin copolymer has been formed through transesterification between a carbonate ester linkage and an anhydride in a second polymer chain. All these examples represent degradative copolymer formation since the PEST chains are cleaved and the average molecular weight of the new copolymer is less than the sum of the average molecular weights of the two immiscible polymers. In Table 5.25 Debier et al. [1995, 1997] present evidence that copolymer formation occurs in PC + PMMA blends through transesterification between PC carbonate ester linkages and acid groups produced by hydrolysis on PMMA. The Table omits references in which transesterification in polyester-polyester blends is brought about under static conditions either by annealing or heating in a DSC chamber. Blends in such examples were prepared by solution precipitation. See for example PET + ethylacrylate-based rubber [Jha and Bhowmick, 1997], and PC + PMMA [Rabeony et al., 1992]. Kyu et al. [1993] found single Tg in annealed PC + PMMA blends and attributed this to thermooxidative crosslinking rather than to transesterification.

375

Reactive Compatibilization of Polymer Blends

Table 5.25. Degradative graft copolymer formation in binary PEST/PO blends by transesterification Polyester / Polyolefin

Characterization and Comments

PC (50-30) / imidized acrylate copolymer (77% glutarimide, 19% MMA, 3% MAA, 2% glutaric anhydride) (70-50)

internal mixer at 260-270°C or solution casting / FTIR / optical microscopy / SEM / TEM / SEC / model reactions / selective solvent extraction / comparison to blends with PMMA or with imidized acrylate copolymer containing different imide levels, or no acid or anhydride / details of reaction mechanism PBT (70) / EVAc (10-20) /PE (10-20) internal mixer at 230°C / SEM / dibutyl tin oxide catalyst (0-4%) / encapsulation of PE dispersed phase by PBT-EVAc copolymer PBT (50-10) / internal mixer at 230°C / selective solvent extraction / EVAc (9% or 28% VA) (50-90) NMR / FTIR / SEM / model reactions / dibutyl tin oxide catalyst (0-1%) / blends optionally + PE PETG (90-80) / EVAc (10-20) internal mixer at 210°C / SEM / NMR model study / rheology / DMA / interfacial tension measurements by breaking thread method / dibutyl tin oxide catalyst (0-1%) PET (100-0) / EEA (0-100) glass mixer at 260-290°C / DMA / SEM / mechanical properties / selective solvent extraction / FTIR / DSC PET (68-23) /LLDPE (23-68)/ extruded at 280°C / DMA / DSC / mechanical properties / EMAA ionomer (5-15) SEM / optical microscopy PETG (70-30) / EVAc (28% VA) (30-70) internal mixer at 180 or 210°C / SEM / rheology / DSC / mechanical properties / dibutyl tin oxide catalyst (0-4%) PC (95-50) / HDPE (0-20) / internal mixer at 220°C / torque rheometry / FTIR / EMAA ionomer (5-50) DSC / selective solvent extraction / mechanical properties / SEM poly(hydroxybutyrate) (80) / internal mixer at 180°C / mechanical properties / SEM EP-g-MA (4% MA) (20) or EP-g-dibutyl maleate or partially hydrolyzed EVAc

5.7.13

Polyester + Polyolefin + Polypropylene Ternary Blends

Copolymer Formation by Carboxylic Acid + Epoxide Reaction PEST/PP blends have been compatibilized through graft copolymer formation by reaction of polyester carboxylic acid end-groups with pendent epoxide groups on an appropriate PO with some

References Debier et al., 1997 a, b; 1995

Legros et al., 1997

Pesneau et al., 1997

Lacroix et al., 1996 a, b

Gravalos et al., 1995 Kalfoglou et al., 1994 Legros et al., 1994 Mekhilef et al., 1992

Abbate et al., 1991

miscibility with PP (Table 5.26). The copolymer is joined through a new ester linkage. Some crosslinked copolymer may also form as discussed in Section 5.7.12.2. When the new ester linkage is formed, a secondary alcohol group forms as well. The new secondary alcohol may also react with polyester main-chain linkages to form copolymer through a degradative transesterification process as discussed in Section 5.7.12.2.

376

5.7.14

S. B. Brown

Polyester + Polyolefin + Styrene Copolymer Ternary Blends

Copolymer Formation by Alcohol + Anhydride Reaction Studies have been made of PEST/PO/styrene copolymer blend compatibilization in which a copolymer may be formed between polyester alcohol end-groups and pendent anhydride functionality on a styrene copolymer (Table 5.27). Because the alcohol-anhydride reaction is reversible (with the equilibrium lying on the side of unreacted anhydride), only a relatively small amount of copolymer may be formed. In consequence, the dispersed phase polymer may not be well stabilized against coalescence upon further thermal treatment [Sun et al., 1996]. Alternatively, at least some copolymer may be formed by a degradative mechanism through transesterification between PEST main-chain linkages and a low concentration of pendent acid groups in anhydride-functionalized styrene copolymer (see Section 5.7.12.5). 5.7.15

Polyester + Polyphenylene Ether + Styrene Copolymer Ternary Blends

Blends Containing Unfunctionalized PPE + Various Functionalized PS; PEST-PS Copolymer Formation by Carboxylic Acid + Epoxide Reaction PEST/PPE blends have been compatibilized through graft copolymer formation between polyester and an epoxide-functionalized styrene copolymer (Table 5.28). The reaction is between polyester carboxylic acid end-groups and pendent epoxide groups on the styrene copolymer leading to a new ester linkage. When the new ester linkage is formed, a secondary alcohol group forms as well. The new secondary alcohol may also react with polyester main-chain linkages to form copolymer through a degradative transesterification process as discussed in Section 5.7.12.2. Since there are multiple epoxide sites on the styrene copolymer, some crosslinked copolymer may result as well if the polyester is functionalized at both ends. The proportion of crosslinked copolymer formed also depends

upon blend composition and processing conditions. 5.7.16

Polyester + Polypropylene Binary or Higher Blends

5.7.16.1 Copolymer Formation by Acid + Epoxide Reaction PEST/PP blends have been compatibilized through graft copolymer formation by reaction of polyester carboxylic acid end-groups with pendent epoxide groups on an appropriate PP or on a PO or styrene-polyolefin copolymer showing some miscibility with PP (Table 5.29). The copolymer is joined through a new ester linkage. When the new ester linkage is formed, a secondary alcohol group forms as well through ring-opening of the epoxide. It is theoretically possible that the new secondary alcohol may also react with polyester main-chain linkages to form copolymer through a degradative transesterification process as discussed in Section 5.7.12.2. Since there are multiple reactive sites on the epoxide-containing polymers, some crosslinked copolymer may result if the acid-containing polymer is functionalized at both ends. The proportion of crosslinked copolymer formed also depends upon blend composition and processing conditions. An example is also included in this section where a compatibilizing copolymer is postulated to form by reaction between acidic phenolic end-groups on polycarbonate and epoxide groups grafted to PP [Zhihui et al., 1997]. 5.7.16.2 Copolymer Formation by Carboxylic Acid + Oxazoline Reaction Vainio et al. [1997, 1996a] have compatibilized PEST/PP blends by graft copolymer formation between acid-terminated polyester and oxazoline-grafted PP. Specifically, 30 parts PBT was mixed with 0-70 parts PP and 0-70 parts PP-goxazoline in an internal mixer at 250°C or TSE at 240°C. Blends were characterized by SEM, torque rheometry, DMA, and DSC. Oxazolinefunctionalized PP was prepared by grafting PP

377

Reactive Compatibilization of Polymer Blends

Table 5.26. Graft copolymer formation in ternary PEST/PO/PP blends by acid + epoxide reaction Polyester / Polyolefin / Polypropylene

Characterization and Comments

References

PBT (70-30) / E-GMA (6% GMA) (0-20) / PP (30-70) PBT (100-0) or LCP (Hoechst Vectra A950) / EEA-GMA (5) / PP (0-100)

TSE at 255°C / torque rheometry / capillary rheometry / FTIR / SEM / TEM / mechanical properties / catalysis of acid + epoxide reaction using phosphonium salt TSE at 240°C / DSC / SEM / FTIR / rheology / mechanical properties / comparison vs. properties of binary EEA-GMA blends with PBT or LCP

Tsai and Chang, 1996

Holsti-Miettinen et al., 1995

Table 5.27. Graft copolymer formation in ternary PEST/PO/Steryne copolymer blends by alcohol + anhydride reaction Polyester / Polyolefin / Styrene Copolymer

Characterization and Comments

References

PET /HDPE / SEBS-g-MA (1-5 parts per 100 parts PET) PET (80) / HDPE (20) / SEBS-g-MA (0-20)

TSE at 270°C / rheology / interfacial tension measurements SSE at 270°C / rheology / SEM / mechanical properties / use of recycle PET TSE at 270°C / mechanical properties / SEM / DSC / WAXD

Ihm and White, 1996

PET (100-0) / HDPE (0-100) / SEBS-g-MA (5-20 phr)

La Mantia et al., 1994 Carte and Moet, 1993

Table 5.28. Graft copolymer formation in PEST/PPE/Styrene copolymer blends by miscellaneous reactions Polyester / PPE / Styrene copolymer

Characterization and Comments

References

PBT (60) / PPE (40) / SB core + MMA shell impact modifier with GMA grafted onto shell (5) PBT (70-50) / PPE (30-50) / S-co-GMA (2-15% GMA) (0-20 parts)

mini-extruder at 240°C / TEM / comparison of morphology with and without addition of epoxy-functionalized impact modifier TSE at 260°C / mechanical properties / SEM / DSC / torque rheometry / DMA / ethyltriphenyl phosphonium bromide catalyst (0-0.05 parts) TSE at 290°C / torque rheometry / HDT / capillary rheometry / SEM / FTIR / mechanical properties / ethyltriphenylphosphonium bromide catalyst (0-0.02 parts) TSE at 290°C / DSC / DMA / SEM / also addition of Phenoxy

Aerdts et al., 1997

LCP (Hoechst Vectra A950) (20-5)/ PPE (40-47) / PS (40-47) / S-co-GMA (2, 5, or 10% GMA) (0-5) PET (50-20) / PPE (50-80) / PS (0-20) / S-co-GMA (15% GMA) (0-10 parts)

Liu et al., 1996

Chang and Chang, 1995 Liang and Pan, 1994

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S. B. Brown

Table 5.29. Graft copolymer formation in binary or higher PEST/PP blends by acid + epoxide reaction Polyester / Polypropylene

Characterization and Comments

References

PET (80-17.5) /PP (17.5-80) / SEBS-g-GMA (0-5) or SEBS-g-MA PC (30-10) / PP (70-90) / PP-g-GMA (0.46 mol% GMA) (2.5-20) PBT(30) / PP-g-GMA (70)

TSE at 275°C / mechanical properties / SEM / rheology / DMTA

Heino et al., 1997

TSE at 250°C / mechanical properties / selective solvent extraction / SEM / DSC / WAXS TSE at 240°C / one-step grafting of GMA to PP followed by downstream feeding and copolymer formation with PBT / comparison to two-step process / effects of processing conditions / mechanical properties / SEM / rheology / failure of PP-MA or PP-g-AA to act as compatibilizer poly(butyl acrylate)-co-GMA TSE at 180°C / mechanical properties / DMTA / gel content / (2-5 mol% functional monomer) FTIR / TEM / poly(butyl acrylate) also copolymerized with (42) / PP (53-58) / PP-MA (0.4% MA) olefinic oxazoline, amine, carboxylic acid, hydroxyl

with ricinol oxazoline maleinate in presence of styrene monomer + RI. The inclusion of styrene monomer suppresses radical induced decomposition of PP. Some crosslinked copolymer may also form in this blend if the polyester is acidfunctionalized at both chain ends. 5.7.17

Polyester + Styrene Copolymer Binary or Higher Blends

5.7.17.1 Copolymer Formation by Alcohol + Anhydride Reaction Studies have been made of PEST/styrene copolymer compatibilization in which a copolymer is formed between polyester alcohol end-groups and pendent anhydride functionality on a styrene copolymer (Table 5.30). Because the alcoholanhydride reaction is reversible with the equilibrium lying on the side of unreacted anhydride, only a relatively small amount of copolymer is formed. Thus, the dispersed phase polymer may not be well stabilized against coalescence upon further thermal treatment [Sun et al., 1996]. Alternatively, at least some copolymer may be formed by a degradative mechanism through

Zhihui et al., 1997 Hu et al., 1996; Sun et al., 1996

Vainio et al., 1996 b

transesterification between polyester main-chain linkages and a low concentration of pendent acid groups in anhydride-functionalized styrene copolymer (see Section 5.7.12.5). 5.7.17.2 Copolymer Formation by Carboxylic Acid + Epoxide Reaction As the examples of Table 5.31 demonstrate, PEST/styrene copolymer blends may be compatibilized through graft copolymer formation by reaction of polyester carboxylic acid endgroups with pendent epoxide groups on an appropriate styrene copolymer. The copolymer is joined through a new ester linkage. When the new ester linkage is formed, a secondary alcohol group is formed as well. The new secondary alcohol may also react with polyester main-chain linkages to generate copolymer through a degradative transesterification process, discussed in Section 5.7.12.2. Since there are multiple epoxide sites on the polystyrene, some crosslinked copolymer may result if the polyester is functionalized at both ends. The proportion of crosslinked copolymer formed also depends upon blend composition and processing conditions.

379

Reactive Compatibilization of Polymer Blends

Table 5.30. Graft copolymer formation in PEST/Styrene copolymer blends by alcohol + anhydride reaction by acid + epoxide reaction Polyester / Styrene Copolymer

Characterization and Comments

References

PBT (70) / ABS (30) / SMA (2.5-7.5) PET (99-95) / SEBS-g-MA (0-4.5% MA) (1-5)

SSE at 255°C / mechanical properties TSE at 260-280°C / IV / SEM / DSC / selective solvent extraction / mechanical properties / rheology extruded at 280°C / DMA / DSC / mechanical properties / SEM / FTIR / optical microscopy TSE at 280°C / DMA / DSC / mechanical properties / SEM / FTIR / optical microscopy / use of recycle PET / ternary blends with HDPE also prepared

Basu and Banerjee, 1997 Tanrattanakul et al., 1997 a, b Kalfoglou et al., 1996

PET (100-0) / ABS (0-100) / ABS-g-MA (3.5% MA) (0-100) PET (100-0) / SEBS-MA (1.8% MA) (0-100)

Kalfoglou et al., 1995

Table 5.31. Graft copolymer formation in PEST/Styrene copolymer blends Polyester / Styrene Copolymer

Characterization and Comments

PBT (90-10) / PS (10-90) / S-GMA (0-9) LCP (75-10) (Hoechst Vectra A900) / PS (25-90) / S-GMA (5% GMA) (5)

internal mixer at 200°C / SEM / TEM / rheology TSE at 290°C / torque vs. time measurements / rheology / mechanical properties / FTIR / SEM / catalysis of acid + epoxide reaction using phosphonium salt PBT (75-25) / PS (25-75) / S-GMA (0-10 parts) internal mixer at 230-240°C / DSC / DMTA / rheology / SEM PBT (75-25) / ABS (25-75) / SAN-GMA (0-5) TSE / FTIR / morphology / mechanical properties vs. use of unfunctionalized SAN / melt flow / titration of residual acid groups / use of ethyl-triphenylphosphonium bromide catalyst /2, 4, and 10% GMA in copolymer PET (75-25) / PS (25-75) / TSE at 270°C / torque rheometry / mechanical S-GMA (5% GMA) (0-5) properties / SEM / melt viscosity / use of ethyltriphenyl-phosphonium bromide catalyst / 2, 5, and 10% GMA in copolymer PET (50) or PETG / S-GMA (5% GMA) (50) internal mixer at 270°C / torque rheometry / GPC / effects of residual levels of different catalysts used to synthesize polyesters

5.7.17.3 Degradative Copolymer Formation by Transesterification Landry et al. [1994] have compatibilized PEST with styrene copolymer by transesterification.

References Kim, J. K., et al., 1997 Chiou et al., 1996 a

Kim and Lee, 1996 Lee et al., 1994

Maa and Chang, 1993

Stewart et al., 1993 b

Specifically, 50 parts PBT (or PET) was extruded with 50 parts poly(vinylphenol) using a mini-SSE at 254 or 293°C, followed by annealing at 265 or 290°C. The blends were characterized by DSC. Copolymer formation evidently occurs through

380

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a degradative reaction between pendent phenolic groups on poly(vinylphenol) and ester linkages in the polyester main chain. A low molecular weight polyester fragment results from this process. 5.7.18

Polyolefin + Polypropylene Binary Blends

Copolymer formation by displacement reaction Solid-state chlorination of PP may lead to low levels of chlorine incorporation along the PP chain with minimal PP molecular weight degradation. Coran and Patel [1983 b] have blended (internal mixer, T = 190°C) 50 parts chlorinated PP with 45 parts NBR containing 5 parts amine-terminated NBR. The blend was characterized by mechanical properties vs. the use of unfunctionalized PP. Compatibilization resulted from copolymer formation through displacement reaction of chloride by amine groups. The blend was prepared in the presence of a vulcanization agent (3.75% of dimethylol phenol + 0.5% of SnCl2) to cause concomitant dynamic vulcanization of the rubber phase. 5.7.19

to an aromatic ring. The process typically proceeds by reaction between an alkyl cation precursor and an aromatic compound in the presence of a Lewis acid catalyst. In blends of 80-20 parts LLDPE with 20-80 parts PS and 0.3 parts aluminum chloride a copolymer was obtained. The proposed mechanism involved cation formation on PE through Lewis acid-catalyzed degradation, followed by attachment of PE cation to aromatic rings of PS [Sun and Baker, 1997]. Blends were prepared in an internal mixer at 180°C. They were characterized by selective solvent extraction, GPC, SEM, FTIR, and mechanical properties. The effects of added styrene monomer were also studied. This process is an example of degradative copolymer formation since the PE graft segment attached to the PS chain has lower molecular weight than the PE phase from which it was derived. Although the initial copolymer formed in this process is a graft copolymer, the grafted PE segments are still capable of reacting with catalyst with subsequent cation formation and attachment to a different PS chain leading potentially to a crosslinked copolymer.

Polyolefin + Polystyrene Binary Blends 5.7.20

5.7.19.1 Copolymer Formation by Transesterification Hu and Lambla [1995] have blended EMAc (90-65 parts) with mono-hydroxy-terminated PS (10-35 parts) in an internal mixer at 180-220°C in the presence of dibutyltin dilaurate or dibutyltin oxide catalyst. A compatibilizing copolymer arises from transesterification between pendent ester groups of EMAc and terminal hydroxy groups of PS. The effects on blend properties of PS molecular weight were reported. The effects of processing conditions and addition of solvent on conversion kinetics were studied. 5.7.19.2 Degradative Copolymer Formation by Friedel-Crafts Coupling The Friedel-Crafts alkylation reaction is the most widely used method for attaching an alkyl group

Polyphenylene Ether + Styrene Copolymer Blends

Blends Containing Unfunctionalized PPE + Functionalized Styrene Copolymers: Copolymer formation between different Styrene Copolymers by amine + anhydride reaction PPE is not miscible with SMA containing as much as 28% MA [Witteler, et al., 1993]. To compatibilize these two resins, Koning, et al. [1996, 1993b] have added a monoamineterminated PS that can form a graft copolymer with SMA. Since the amine-terminated PS is miscible with PPE, compatibilized PPE-SMA blends are obtained. Specifically, 30 parts of unfunctionalized PPE was blended (internal mixer at 220°C, mini-SSE at 280°C, or TSE at 326°C) with 56 parts SMA (28% MA) and 14 parts aminefunctionalized PS. The blend was characterized by TEM, SEM, mechanical and thermal properties, DMA, and GPC copolymer detection. The effect

Reactive Compatibilization of Polymer Blends

of pre-reacting amine-terminated PS with SMA was studied. The blend properties were compared to those for uncompatibilized blends. Blends were also made containing ABS + SEBS. 5.7.21

Polypropylene + Styrene Copolymer Binary Blends

5.7.21.1 Copolymer Formation by Amine + Anhydride Reaction Dharmarajan et al. [1995] have compatibilized PP/styrene copolymer by formation of a graft copolymer through reaction of secondary amineterminated PP (0.4 wt% amine) with SMA. The secondary amine-terminated PP was prepared by first extruding PP with MA to form anhydride-terminated PP followed by extrusion with N-methyl1,3-propanediamine to give the secondary amineterminated PP through reaction of the primary amine of the diamine with the PP-anhydride. Blends within the range PP to SMA from 0-100 to 100-0, containing 0-5 parts amine-terminated PP were prepared in an internal mixer at < 220°C. They were characterized by FTIR, DMTA, TEM, mechanical properties, rheology, lap shear adhesion, and paint adhesion. Some crosslinked copolymer may be present if the PP-MA contains more than one anhydride group. For related work see also Section 5.9.8. 5.7.21.2 Copolymer Formation by Carboxylic Acid + Oxazoline Reaction Sundararaj et al. [1995] have prepared blends containing oxazoline-functionalized PS and PP-MA. A graft copolymer may form through reaction between pendent oxazoline groups on PS and terminal acid groups (from some hydrolysis of anhydride groups) on PP. Specifically, 80 parts S-IPO (1% IPO) was blended with 20 parts PP-MA (0.1% MA) in either an internal mixer at 200°C or in a TSE. The blends were characterized by selective solvent extraction and SEM. Morphology development in the different mixing equipment was studied in both reactive and nonreactive blends (i.e., PS/PP).

5.8

381

Compatibilization by Block Copolymer Formation: Process #3

Blends are discussed in alphabetical order of polymeric components. 5.8.1

General Considerations

Compatibilized polymer blends have been prepared through block copolymer formation between immiscible polymers. In the Direct Process, during melt processing the functionalized end-groups on some fraction of chains in each of the polymers react across a melt phase boundary to form block copolymers. Depending upon stoichiometry, either A-B or A-B-A or both copolymer structures may be obtained as shown in Table 5.4 (Type 3a or 3b). The average molecular weight of the copolymer corresponds to the sum of the average molecular weights of the reacting polymers. Block copolymers may also be formed through reaction of the end-group on one polymer with a condensing agent that activates that end-group for reaction with a nucleophilic end-group on a second immiscible polymer. Typical condensing agents include phosphite esters that react with acid and hydroxy end-groups on condensation polymers. A by-product from the condensing agent is always formed in the copolymer reaction and is often removed by devolatilization of the blend melt. Since the condensing agent is not incorporated into the copolymer, the process is the same as that shown in Table 5.4 (Type 3a or 3b). Block copolymers may also be formed through reaction of end-groups on each of the immiscible polymers with a coupling agent. Typical coupling agents include multifunctional epoxy resins, oxazolines, carbodiimides, and isocyanates that react with nucleophilic end-groups of condensation polymers. The coupling agent is incorporated into the copolymer. The process is shown in Table 5.4 (Type 3c). When the coupling agent is an epoxide, a new secondary alcohol is formed when the epoxide ring is opened by a nucleophile. This alcohol may also be reactive to one or more polymeric components (particularly polyesters)

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similar to the reactivity of alcohol groups on Phenoxy resin (see Sections 5.8.3.2; 5.8.11.2; 5.8.13.1; and 5.8.13.2). General characteristics of processes involving coupling agents are also discussed in Section 5.5.3.2. Block copolymers may also be formed by a degradative process in which end-groups on one polymer undergo transreaction with linkages in the main chain of a second, immiscible polymer. A low molecular weight fragment of the second polymer is formed as by-product. The block copolymer has lower average molecular weight than the sum of the average molecular weights of the reactants. The process is shown in Table 5.4 (Type 3d). A typical copolymer architecture in this process is A-B block. An A-B-A block can only form if the degradable segment is further degraded through transreaction with another endgroup functionalized polymer. 5.8.2

Polyamide + Polyamide Binary Blends

Blends Containing a Condensing Agent; Copolymer Formation by Amine + Carboxylic Acid Reaction Aharoni et al. [1984] and Aharoni [1983] have shown that blends of immiscible polyamides may be compatibilized through copolymer formation mediated by addition of a phosphite condensing agent. Block copolymer results when the phosphite-activated end-group of one PA reacts at the phase interface with a nucleophilic end-group on the second PA. The reaction also produces a secondary phosphite by-product. The relative proportions of copolymer vs. simple chain-extended PA may depend upon the relative solubility of the condensing agent in each of the immiscible polymer phases. For example, blends of 95-50 parts PA-6 were extruded using a SSE at 265-315°C with 5-50 parts PA-11 (or PA-12 or PA-66 or PA-6T) in the presence of 0-1 part triphenyl phosphite or other trialkyl phosphite. Copolymercontaining blends were characterized by selective solvent extraction, FTIR, 13C NMR, and 31P NMR. Model compound studies were done to understand the mechanism of copolymer formation.

5.8.3

Polyamide + Polyester Binary Blends

5.8.3.1 Blends Containing a Condensing Agent Copolymer Formation by Amine + Carboxylic Acid Reaction PA/PEST blends have been compatibilized by Aharoni et al. [1984] and Aharoni [1983] through copolymer formation mediated by addition of a phosphite condensing agent. Block copolymer results when the phosphite-activated end-group of a PEST (or PA) reacts with a nucleophilic end-group on a PA (or PEST) at the phase interface with generation of secondary phosphite byproduct. The relative proportions of copolymer vs. simple chain-extended PA or PEST may depend upon the relative solubility of condensing agent in each of the immiscible polymer phases. For example, blends of 95-5 parts PA-6 (or PA-66 or PA-12) were extruded using a SSE at 265-315°C with 5-95 parts PET (or PBT or PCT) in the presence of 0-1 part triphenyl phosphite or other trialkyl phosphite. Copolymer-containing blends were characterized by morphology, viscosity, selective solvent extraction, FTIR, 13C NMR, and 31P NMR. Model compound studies were done to understand the mechanism of copolymer formation. Various other PEST and PA resins were also used. 5.8.3.2 Blends Containing a Coupling Agent Copolymer Formation by Carboxylic Acid + Epoxide Reaction As shown in Table 5.32, PA/PEST blends can be compatibilized through copolymer formation mediated by addition of a multifunctional epoxide coupling agent. The coupling agent may react with nucleophilic end-groups on each of the two immiscible polymers at the phase interface to give a block copolymer containing the coupling agent as linking group. The relative proportions of copolymer vs. simple chain-extended PA or PEST may depend upon the relative solubility of coupling agent in each of the immiscible polymer phases.

383

Reactive Compatibilization of Polymer Blends

5.8.4

Polyamide + Polyester + Polypropylene Ternary Blends

Copolymer Formation by Amine + Anhydride Reaction Tjong and Meng [1997] have described PA/LCP polyester blends with improved properties through addition of anhydride-terminated PP-MA. A block copolymer may form between PA amine endgroups and anhydride-terminated PP (see Section 5.8.8). For example, 86 parts PA-6 was mixed in an internal mixer with 14 parts PP-MA at 220°C followed by injection molding with 5-40 parts LCP (Hoechst Vectra A950). The blends were characterized by torque rheometry, mechanical properties, DMA, and SEM.

5.8.5

Polyamide + Polyolefin Binary Blends (excepting polypropylene)

Copolymer Formation by Amine + Carboxylic Acid Reaction As shown by examples listed in Table 5.33, PA/PO blends have been compatibilized through block copolymer formation between PA amine end-groups and terminal carboxylic acid groups of polyolefins. PA blends with oxidized PE fall into this category since carboxylic acid groups are located at PE chain ends [see El’darov et al., 1996].

Table 5.32. Block copolymer formation in PA/PEST blends by carboxylic acid + epoxide coupling agent Polyamide / Polyester

Characterization and Comments

References

PA-66 (70-50) / PBT (30-50) / mutifunctional epoxide resin (0-5 phr)

TSE at 275°C / SEM / TEM / extrudate swell / DSC / torque rheometry / capillary rheometry / effects of PA/PBT ratio and epoxy content on mechanical properties / blends optionally + core-shell elastomer TSE at 230-250°C / mechanical properties / SEM / DSC / DMA / WAXD

Huang and Chang, 1997 a, b

PA-6 (17) / PBT (83) / mutifunctional epoxide resin (0-12)

An et al., 1996

Table 5.33. Block copolymer formation in PA/PO blends by amine + acid reaction Polyamide / Polyolefin

Characterization and Comments

References

aromatic PA (50) / hydrogenated NBR (50) / carboxy-terminated NBR (0-5)

miniature mixer at 250°C / morphology / ellipsometry / L101 (0.9 phr) added for vulcanization of rubber phase / effects of mixing protocol / other carboxylated rubbers also used internal mixer at 240°C / morphology / mechanical and rheological properties vs. functionalization level

Bhowmick et al., 1993

PA-6 (75) / oxidized LDPE (25)

Curto et al., 1990

384

5.8.6

S. B. Brown

Polyamide + Polyolefin + Polypropylene Ternary Blends

PA-PP Copolymer Formation by Amine + Anhydride Reaction Chen and White [1993] and Chen et al. [1988] have reported properties for blends containing 75-25 parts PA-6 (or PA-11), 25-75 parts LDPE (or HDPE) and 0-5 parts PP-MA. A block copolymer may form between PA amine end-groups and anhydride-terminated PP (see Section 5.8.8). The blends were prepared in a SSE at 200-230°C. They were characterized by mechanical properties and DSC. Morphology and capillary rheometry were done before and after annealing. Interfacial tension measurements for the mixtures were also reported. 5.8.7

Polyamide + Polyphenylene Ether Binary or Higher Blends

Blends Containing Anhydride-Terminated PPE; PA-PPE copolymer formation by amine + anhydride reaction Campbell et al. [1990] have reported properties for compatibilized PA-PPE blends made using anhydride-terminated PPE. A block copolymer may form between PA amine end-groups and PPE-anhydride. For example, a blend containing 49 parts PPE-anhydride, 41 parts PA-6, and 10 parts SEBS impact modifier was extruded using a TSE at 280°C. The copolymer-containing blend was characterized by selective solvent extraction, mechanical properties, SEM, and TEM. Ductilebrittle transition temperatures were determined. Anhydride-terminated PPE was made by capping PPE phenolic end-groups with trimellitic anhydride chloride in solution. 5.8.8

Polyamide + Polypropylene Binary Blends

Copolymer Formation by Amine + Anhydride Reaction Immiscible blends of PA and PP have been compatibilized through copolymer formation between PA amine end-groups and maleic anhydride-

functionalized PP to form a new imide linkage (Table 5.34). The structure of maleic anhydridefunctionalized PP has been discussed by De Roover et al. [1995] and Sclavons et al. [1996]. The authors demonstrated that the free radical initiated maleation of PP in the molten state leads to anhydride groups locating mainly at PP chain ends. Consequently, unless PP has been functionalized by a process that suppresses PP chain degradation, the reaction product of amineterminated PA with MA-functionalized PP is predominantly a block copolymer. 5.8.9

Polyamide + Polyurethane Binary Blends

Blends Containing a Bis-Isocyanate Coupling Agent PA/TPU blends have been compatibilized by addition of a bis-isocyanate coupling agent that is capable of reacting with nucleophilic endgroups on both polymers to form a block copolymer. Franke et al. [1993] have extruded PA-6 (20-0 parts) with polyester-urethane (78-100 parts) using a TSE at 230°C in the presence of 0.5-2 parts diphenylmethane diisocyanate. The coupling agent was added downstream of the extruder feed-throat. The blends were characterized by TEM, SEC, DSC, DMA, and FTIR. 5.8.10

Polyamide + Styrene Copolymer Binary Blends

Copolymer Formation by Amine + Anhydride Reaction PA/PS blends have been compatibilized through block copolymer formation between amine-terminated PA and anhydride-terminated PS. Anhydride end-groups were introduced into PS through reaction of either anion-terminated PS or hydroxyterminated PS with trimellitic anhydride chloride. Park et al. [1992] blended 80 parts PA-6 with 10-16 parts PS and 4-10 parts anhydride-terminated PS in an internal mixer at 240°C. The blends were characterized by torque rheometry, SEM, selective solvent extraction, DSC, morphological stability to annealing, and lap shear adhesion. The effect

385

Reactive Compatibilization of Polymer Blends

Table 5.34. Block copolymer formation in binary PA/PP blends by amine + anhydride reaction Polyamide / PP

Characterization and Comments

References

amorphous PA (30) / PP (70-0) / PP-MA (0-70) PA-6 (30) / PP(60-70) / PP-MA (0.4% MA) (0-10) PA-66 (100-0) / PP (0-100) / PP-MA (0-5) PA-6 / PP-MA

Mini-Max molder at 240°C / laser light scattering / TEM / SEM / ellipsometry TSE at 230°C / rheology / DSC / WAXS / FTIR / FT-Raman / optical microscopy TSE at 255°C / mechanical properties

Li et al., 1997

interfacial fracture energies between molded plaques as function of temperature / video imaging, ESCA and SEM of fracture surfaces / DSC internal mixer at 240°C / TEM / mechanical properties and dispersed phase domain size vs. vol. fraction of compatibilizer / also blends containing EP rubber

Bidaux et al., 1996

PA-6 (30-25) or PA-12 / PP (60-70) / PP-MA (0-20)

PA-6 (100-25) / PP (25-100) / PP-MA (5.3% MA) (0-9) PA-6 (60) / PP-MA (40) PA-6 (100-0) / PP-MA (0.4-3.1% MA) (0-100) / PP ( 0-100) PA-66 (30-0) / PP (40-100) / PP-MA (0-30) PA-12 (25) / PP (0-75) / PP-MA (0-75) PA-6 (75-68) / PP (25-23) / PP-MA (10) PA-66 (75-25) / PP (25-75) / PP-MA (0.2-2.7% MA) (0-25) PA-66 (50-20) / PP-MA (0.9-2.1% MA) (50-80) PA-66 (76-19) / PP (19-76) / PP-MA (0.4% MA) (0-5) PA-6 (80-30) / PP (18-56) / PP-MA (0-14) PA-12 (80) /PP (20) / PP-MA (6% MA) (4) or PP-g-AA PA-6 (100-0) / PP (0-100) / PP-MA (0.2% MA) (0-10) PA-6 (100-0) / PP (0-100) / PP-MA (5% MA) (0-6)

internal mixer at 235°C / mechanical properties / DSC / water absorption / SEM / DMA SSE or TSE / mechanical and thermal properties / use of recycle PP containing 5-10% EVAl + 2-5% LDPE vs. virgin PP SSE or internal mixer at 240°C / torque rheometry / mechanical properties / TEM internal mixer at 240-270°C / SAXS / effects of extrusion temperature on dispersion solution precipitation followed by heat treatment at 180°C / SEM / DMA / DSC / effects of annealing / selective solvent extraction TSE at 240°C / SEM / mechanical and thermal properties / rheometry / effects of PA mol. wt. / also glass filled blends TSE at 285°C / morphology and mechanical properties vs. 0.2 or 2.7 wt% MA content PP / DSC crystallization behavior mechanical properties / rheology / SEM / copolymer analysis TSE at 280°C / morphology and mechanical properties vs. PP-PA viscosity ratio internal mixer or TSE at 240°C / SEM /mechanical and rheological properties / DMA / water absorption internal mixer at 210°C / torque rheometry / SEM / DSC / rheology / mechanical properties TSE at 245°C / DSC / SEM / mechanical properties

internal mixer at 240°C / rheology / mechanical properties / SEM / DSC / effects of processing conditions

Marco et al., 1997 Roberts et al., 1997

Rösch et al., 1996; Rösch, 1995; Rösch and Mülhaupt, 1995; 1993 Sathe et al., 1996 Akkapeddi et al., 1995 Gonzalez-Montiel et al., 1995 b, d Lin et al., 1995 Tang et al., 1995, 1996 Wu et al., 1995 Duvall et al., 1994 a, b, c Fritz et al., 1994 Hietaoja et al., 1994 Speroni, 1994 Valenza and Acierno, 1994 Ikkala et al., 1993; Holsti-Miettinen et al., 1992 La Mantia, 1993

386

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Table 5.34. Continued Polyamide / PP

Characterization and Comments

References

PA-6 (30) / PP (50-70) / PP-MA (0-20) PA-6 (60) / PP (40) / PP-MA (12)

internal mixer / TEM / SEM / creep measurements

Schlag et al., 1993

round-robin DMTA study of commercial blend supplied as extrudate fiber extrusion at 260°C / mechanical properties / SEM / SALS

Wippler, 1993 Grof et al., 1992

extruder at 245°C / SEM / DSC / rheology

Park et al., 1990

internal mixer at 220°C / PA + rubber vulcanization and PP + rubber performed in separate steps before blending / mechanical properties vs. use of unfunctionalized PP / comparison to use of functionalized PP in initial vulcanizate / PP-f-carboxymethyl maleamic acid extrusion-molding at 230°C / selective solvent extraction / DSC / morphology / mechanical properties vs. MA content / residual amine conc.

Coran et al., 1985

PA-6 (100-0) / PP (0-100) / PP-MA (0-5) PA-6 (100-10) / PP (0-75) / PP-MA (3% MA) (0-90) (PA-6-66 copolymer (35)-NBR (65) vulcanizate) (50) / (PP(50)-EPDM(50) vulcanizate) (50) / PP-MA (10) PA-6 / PP-MA

of mixing protocol on properties was studied. Properties were also compared to those for blends compatibilized by added PA-PS graft copolymer that had been synthesized in a separate step. 5.8.11

Polyester + Polyester Binary Blends

5.8.11.1 Blends Containing a Condensing Agent PEST-PEST Copolymer Formation by Alcohol + Carboxylic Acid Reaction As shown in Table 5.35, blends of immiscible polyesters may be compatibilized through copolymer formation mediated by addition of a phosphite condensing agent. Block copolymer results when the phosphite-activated end-group of one PEST reacts with a nucleophilic end-group of another PEST. The reaction takes place at the phase interface. A secondary phosphite is a byproduct. The relative proportions of copolymer vs. simple chain-extended PEST may depend upon the relative solubility of condensing agent in each of the immiscible polymer phases.

Ide and Hasegawa, 1974

5.8.11.2 Blends Containing a Coupling Agent PEST-PEST Copolymer Formation by Carboxylic Acid + Epoxide Reaction Chin and Chang [1997] and Chin et al. [1996] have compatibilized blends of immiscible PEST through addition of a coupling agent. A multifunctional epoxide coupling agent is capable of reacting with nucleophilic end-groups on each of the two immiscible PEST at the phase interface to give a block copolymer containing the coupling agent as linking group. In an example 100-85 parts PET was extruded using a TSE at 270-285°C with 0-15 parts LCP (Hoechst Vectra A900) and 0-2 parts tetrafunctional epoxy resin. The blends were characterized by torque rheometry, capillary rheometry, DSC, SEM, and FTIR. Mechanical properties were determined vs. composition and morphology. Ethyltriphenylphosphonium bromide was added as a catalyst to promote the reaction of polyester acid or alcohol end-groups with epoxy resin. The relative proportions of copolymer vs. simple chain-extended PEST may depend upon

387

Reactive Compatibilization of Polymer Blends

Table 5.35. Block copolymer formation in PEST/PEST blends by Alcohol + Carboxylic Acid + condensing agent Polyester / Polyester

Characterization and Comments

References

PET (75) / PBT (25) / various phosphite condensing agents (0-5%)

internal mixer at 275-280°C / GPC / DSC / torque rheometry / viscometry / selective solvent extraction / phosphorus analysis / effect of PET end-group concentrations / FTIR for end-groups concentration / effect of phosphite structure / model study with OH + COOH-terminated acrylic polymer / detailed mechanistic study internal mixer at 271°C / torque rheometry / mechanical properties / 31P NMR / GPC / DSC / environmental

Jacques et al., 1997; 1996 a, b; 1993

PET (80-70) / TPE polyester (20-30) / various phosphite condensing agents (0-3%)

stress cracking / use of recycle PET

the relative solubility of coupling agent and catalyst in each of the immiscible polymer phases. 5.8.12

Abu-Isa et al., 1996

Polyester + Polyolefin Binary Blends (excepting polypropylene)

5.8.12.1 Degradative Copolymer Formation by Transesterification A block copolymer may be formed through transesterification between nucleophilic end-groups of one polymer and ester linkages in the chains of an immiscible polyester phase. Sek and Kaczmarczyk [1997] and Minkova et al. [1996] used oxidized PE as the acid-terminated polymer and a LCP with ester linkages susceptible to transesterification with PE acid end-groups. Since the product is a block copolymer accompanied by low molecular weight fragments from the polyester, this type of copolymer forming reaction is degradative. The process is not redistributive (Section 5.6) since the PE chain cannot participate in redistribution. In an example 50 parts LCP polyester (a copolymer of sebacic acid, dihydroxy-biphenyl, and p-hydroxybenzoic acid) and 50 parts oxidized PE were mixed in an internal mixer at 240°C. The blend was characterized by selective solvent extraction, FTIR, DSC, and SEM. Optionally, Ti(OBu)4 catalyst was added to promote reaction.

5.8.12.2 Blends Containing a Coupling Agent Okamoto and Inoue [1993] have compatibilized an immiscible blend of terminally functionalized PEST and PO through block copolymer formation in the presence of a coupling agent. For example, 80 parts hydroxy-terminated PCL was mixed in a custom melt reactor at 120°C with 20 parts carboxyl-terminated butadiene oligomer (or carboxyl-terminated NBR) in the presence of 0-1 part aminopropyltriethoxysilane coupling agent. Morphology development could be followed through analysis of successive samples taken periodically from the reactor. Samples were characterized by light scattering photometry, ellipsometric analysis, and GPC. 5.8.13

Polyester + Polyphenylene Ether Binary or Higher Blends

5.8.13.1 Blends Containing Epoxide-Terminated PPE PEST-PPE Copolymer Formation by Carboxylic Acid + Epoxide Reaction Brown [1997] has reported that polyester-PPE blends may be compatibilized through block copolymer formation between epoxy-terminated PPE and carboxylic acid end-groups on such

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S. B. Brown

polyesters as PBT. For example, blends containing 49 parts PPE-epoxide, 41 parts PBT and 10 parts SEBS impact modifier were prepared using a TSE and characterized by selective solvent extraction, morphology, and mechanical properties. Model reactions were done in solution. PPE-epoxide was prepared by solution capping of PPE phenolic end-groups with 2-chloro-4glycidoxy-6-aryloxy-1,3,5-triazine. 5.8.13.2 Blends Containing Unfunctionalized PPE + Coupling Agent Lo et al. [1997] have presented evidence that a multifunctional epoxy resin may serve as a coupling agent for copolymer formation between PEST carboxylic acid end-groups and PPE phenolic end-groups. Compositions containing 100-0 parts PET and 0-100 parts PPE were extruded with 0-0.5 parts tetrafunctional epoxy resin using a TSE at 265-290°C. The blends were characterized by torque rheometry, FTIR, mechanical properties, SEM, and DSC. Ethyltriphenylphosphonium bromide was added as a catalyst to promote the copolymer-forming reaction. A property comparison was made to blends containing a difunctional epoxy resin. 5.8.14

Polyolefin + Polypropylene Binary Blends

Copolymer Formation by Amine + Anhydride or Amine + Carboxylic Acid Reaction Coran and Patel [1983] have shown that the mechanical properties of dynamically vulcanized

NBR-PP blends can be improved through copolymer formation between the two immiscible polymers concurrent with vulcanization. In the first example in Table 5.36, block copolymer resulted from reaction of amine-terminated NBR with anhydride-terminated PP. The latter was prepared through functionalization of PP with MA in the presence of radical initiator (see Section 5.8.8). In the second example, a block copolymer may have resulted from reaction of acid-terminated NBR with a primary amine-terminated PP. The latter was prepared in a prior reaction between maleic anhydride-terminated PP and triethylenetetramine. It is also possible that the block copolymer may be linked through ionomeric association resulting from protonation of PP-amine with NBR acid (see Section 5.12.4).

5.9

Compatibilization by Direct Covalent Crosslinking; Process #4a

Blends are discussed in alphabetical order of polymeric components. 5.9.1

Crosslinking Reactions. General Considerations

Immiscible polymer blends have been compatibilized by formation of crosslinked copolymers that act as compatibilizing agents. The crosslinking reactions have been performed by at least four methods: 1. Direct crosslinking by covalent bond formation between functionalities on each of the two

Table 5.36. Block copolymer formation in PO/PP blends Polyolefin / PP

Characterization and Comments

References

NBR (50-25) / amine-terminated NBR (0-10) / PP (45-50) / PP-MA (0-5)

internal mixer at 190°C / mechanical properties vs. omission of functional polymers / dimethylol phenol (3.75%) + SnCl2 (0.5%) added as vulcanization agent for NBR internal mixer at 190°C / mechanical properties vs. use of non-functionalized PP / dimethylol phenol (3.75%) + SnCl2 (0.5%) added as vulcanization agent for NBR

Coran and Patel, 1983 b

carboxyl-terminated NBR (50) / amine-terminated PP (50)

Coran and Patel, 1983 b

Reactive Compatibilization of Polymer Blends

immiscible polymers, without degradation of either polymer. The crosslinking is most often performed by reaction of pendent, nucleophilic sites of one multifunctional polymer with pendent, electrophilic sites of the second multifunctional polymer, as shown in Table 5.4 (Type 4a). Common examples include reactions of pendent acid or amine nucleophiles on one functionalized polymer with pendent electrophilic groups such as epoxide, oxazoline, or ortho ester on a second functionalized polymer. 2. Direct crosslinking by covalent bond formation arising from radical generation and recombination (mechano-chemical radical generation, e.g., during extrusion). The method is less frequently used than #1. When two immiscible polymers can both form radicals on their main chains in the absence of added radical initiator, then a copolymer may form by recombination at the phase interface of radical sites on the two polymers. If radical formation occurs without chain degradation, then the copolymer becomes crosslinked as multiple sites on each chain participate in the reaction (also Table 5.4; Type 4a). Alternatively, one radical-forming polymer can form a crosslinked copolymer with a second polymer containing a radical trap such as an unsaturated site (e.g. EPDM). This crosslinking process may be difficult to control since it only stops when thermal and/or shear conditions are below some threshold level. Some examples are given in Section 5.9.6.3. Casale and Porter [1975, 1978] and La Mantia and Valenza [1994] have briefly reviewed mechano-chemical radical generation and its use to form copolymers in immiscible blends. Ahn et al. [1995] have described radical generation and copolymer formation in immiscible polymer blends subjected to elastic strain pulverization in specially modified extruders building on earlier Russian work. Pulverization occurs in an extruder section kept below the polymer melting points. The compatibilizing copolymer formed in these cases may be a block, graft, or crosslinked copolymer depending upon the polymers involved. 3. Crosslinking mediated by a third, added reagent that serves as a coupling, condensing, or activat-

389

ing agent (discussed in Sections 5.10 and 5.11). 4. Ionic crosslinking (discussed in Section 5.12). 5.9.2

Polyester + Polyolefin Binary Blends (excepting polypropylene)

Copolymer formation by miscellaneous reactions Hourston et al. [1991] have prepared compositions of 60-0 parts PBT and 16-40 parts EPDM in the presence of 0-60 parts copolymer of PBT with maleate ester (3.5% maleate) using a TSE at 255°C. A compatibilizing copolymer resulted from the crosslinking reaction between maleate olefinic groups and EPDM olefinic groups. Blends were characterized by mechanical properties and TEM. Model studies were performed to understand the crosslinking process. Blends were also prepared using an internal mixer at 250°C. 5.9.3

Polyester + Polypropylene Binary Blends

Copolymer Formation by Acid + Epoxide Reaction Vainio et al. [1996b] have prepared (in a TSE at 180°C) blends of 42 parts poly(butyl acrylate-coGMA) (2-5 mol% GMA) with 53-58 parts PP and 0-5 parts PP-g-AA (6% AA) as a compatibilizer. A crosslinked copolymer may result through reaction of epoxide groups along the copolyester main chain with acid groups grafted at more than one site along the PP chain (or with acid groups along a poly[acrylic acid] segment grafted at a single PP site). The blends were characterized by DMTA, gel content, FTIR, and TEM. Functionalized poly(butylacrylate) copolymers for use in these blends were also prepared by copolymerization with olefinic oxazoline-, amine-, carboxylic acid-, and hydroxyl-containing monomers. 5.9.4

Polyether + Polyolefin Binary Blends (excepting polypropylene)

5.9.4.1 Copolymer Formation by Alcohol + Anhydride Reaction Compatibilizing copolymers have been formed by direct reaction between pendent alcohol groups of

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S. B. Brown

Phenoxy resin and pendent anhydride groups on MA-grafted polyolefins. Again, the initial product of the reaction was a graft copolymer, but longer reaction time resulted in a crosslinked copolymer since the pendent Phenoxy chain is capable of further reaction with MA on a different polyolefin chain. Mascia and Bellahdeb [1994 a, b] have blended 75-25 parts Phenoxy resin with 25-75 parts EP-g-MA (0.7% MA) in either an internal mixer or a TSE at 180-200°C. The blends were characterized by selective solvent extraction, SEM, DSC, and rheology. Increased crosslinking was observed with increasing MA-to-Phenoxy ratio and with the addition of either NaOEt or sodium benzoate (2%). The resulting copolymers were used to compatibilize PET + HDPE blends. 5.9.4.2 Copolymer formation by alcohol + ester transreaction Kim and Choi [1996b] have compatibilized blends of Phenoxy resin with PMMA. Compatibilizing crosslinked copolymer resulted from transreaction between Phenoxy pendent alcohol groups and pendent PMMA ester groups. Compositions of 100-0 parts Phenoxy resin and 0-100 parts PMMA were prepared in an internal mixer at 240°C. The blends were characterized by DSC, FTIR, mechanical properties, rheology, and DMA. 5.9.5

Polyether + Styrene Copolymer Binary or Higher Blends

Copolymer Formation by Alcohol + Epoxide Reaction Compatibilizing copolymers have been formed by direct reaction between pendent alcohol groups of Phenoxy resin and pendent epoxide groups on GMA-grafted styrene copolymer. Again, the initial product of the reaction was a graft copolymer, but longer reaction time resulted in a crosslinked copolymer since the pendent Phenoxy chain was capable of further reaction with GMA. Furthermore, the new hydroxy group arising from ring-opened epoxy groups may also participate in crosslinking reactions. It is also possible that some crosslinked copolymer resulted from trans-

esterification between pendent hydroxy groups on Phenoxy resin and the GMA ester groups. Chen and Chang [1994] have prepared blends of 75-45 parts Phenoxy resin and 25-50 parts ABS with addition of 0-10 parts SAN-GMA (5% GMA) copolymer in an internal mixer or a TSE at 230°C. The blends were characterized by torque rheometry, SEM, viscosity behavior, and mechanical properties. Sodium lauryl sulfonate (0.02-0.10 parts) was added as a catalyst to promote the copolymer-forming reaction. 5.9.6

Polyolefin + Polyolefin Binary or Higher Blends

5.9.6.1 Copolymer Formation by Acid + Epoxide Reaction As shown in Table 5.37, immiscible blends containing at least two polyolefins have been compatibilized through crosslinked copolymer formation between acid groups at multiple sites on one polyolefin and epoxide groups at multiple sites on another polyolefin. 5.9.6.2 Copolymer Formation by Transesterification of Pendent Ester Groups Immiscible polyacrylates have been compatibilized through transesterification between pendent ester groups on one polyacrylate with pendent ester groups on another polyacrylate. Selected examples are listed in Table 5.38. Dibutyltin oxide was used as transesterification catalyst. 5.9.6.3 Copolymer Formation Involving Mechano-chemical Radical Formation Radicals may be formed on polyolefins under high shear mixing conditions in a process termed mechano-chemical radical generation (see Section 5.9.1). Two immiscible polyolefins may both be subject to mechano-chemical radical generation during extrusion processing. In blends of POs that are not prone to degradation upon radical formation, a crosslinked copolymer may result

391

Reactive Compatibilization of Polymer Blends

Table 5.37. Crosslinked copolymer formation in binary PO/PO blends Polyolefin / Polyolefin

Characterization and Comments

References

poly(acrylic acid) (50) / ENR (50)

internal mixer at 180°C / rheology / solvent swelling / SEM / mechanical properties / DMA / carbon black filler (0-25 phr) / effects of mixing sequence internal mixer at 150°C / mechanical properties / DMA / solvent swelling / effects of adding ZnO or ZnSt to form ionomeric network / DSC / FTIR / UV / SEM internal mixer at 180°C /mechanical and thermal properties / FTIR internal mixer at 150°C /mechanical properties vs. omission of functionalized polymers / morphology

Mallick et al., 1997, 1993

PE-g-AA (6% AA) (100-0) / ENR (0-100) NBR-g-AA (71-0) / ENR (7-75) / PVC (14-75) PE (23) / PE-g-MA (3) / NR (58) / ENR (17)

Mohanty et al., 1997, 1996, 1995; Mohanty and Nando, 1997 Ramesh and De, 1993 Choudhury and Bhowmick, 1989

Table 5.38. Crosslinked copolymer formation in PO/PO blends Polyolefin / Polyolefin

Characterization and Comments

References

EVAc (48) / E-MAc (48)

TSE at 165°C optionally with on-line microwave treatment to effect transreaction, or TSE at 165-200°C / morphology development along screw axis vs. processing conditions in three different extruders / rheology / TGA-GC / selective solvent extraction / SEM / mechanical properties / use to fix morphology of EVAc + E-MAc dispersed phase in PP matrix / dibutyltin oxide catalyst (0-4) TSE at 165°C under which conditions no crosslinking occurred / crosslinked through heating between parallel plates in presence of dibutyltin oxide / swelling monitored by analysis for methyl acetate / rheology

De Loor et al., 1997, 1994; Cassagnau and Michel, 1994; Cassagnau et al., 1993

EVAc (90-5) / E-MAc (10-95)

from the recombination of radicals from each polymer. As a result, a compatibilized blend with improved physical properties is obtained. The examples listed in Table 5.39 show crosslinked copolymer formation in PVC blends with POs. Casale and Porter [1975] report that copolymer formation between NBR and PVC may occur via mechano-chemical radical generation on each polymer followed by recombination. The proposed mechanism was based on earlier studies by Akutin [1968]. More recently, another

Espinasse et al., 1994

mechanism was postulated for the same system [Manoj et al., 1993]. It involved hydrolysis of nitrile groups to amide or carboxylic acid followed by displacement of allylic chloride on PVC by amide -NH2 or acid -OH. 5.9.6.4 Copolymer Formation by Miscellaneous Reactions Unusual ester crosslinking reactions have been proposed as copolymer forming processes for

392

S. B. Brown

immiscible PO blends (Table 5.40). Future work may examine the possible contributions of simpler processes such as IPN formation, radical-radical coupling, or ionomeric crosslinking. 5.9.7

Polyolefin + Polypropylene Binary Blends

5.9.7.1 Copolymer Formation by Acid + Epoxide Reaction A compatibilizing copolymer may be formed through reaction between carboxylic acid groups grafted onto a PO chain and acrylate epoxide groups grafted onto a PP chain. In an internal mixer Liu et al. [1993] have prepared compositions comprising 20 parts NBR-g-AA, 0-75 parts PP and 0-25 parts PP-g-GMA (0.8% GMA). The blends were characterized by torque rheometry, SEM, FTIR, and mechanical properties vs. use of unfunctionalized NBR and vs. different GMA levels.

Although the exact structure of the copolymer is not known, it is convenient to consider that the reaction leads to a crosslinked copolymer. This is based on the assumption that the acrylate epoxide groups are truly grafted onto the PP chains and are not all at terminal sites. The latter structure could result from PP decomposition in the presence of radical initiator, followed by trapping the terminal radical by monomeric acrylate epoxide. 5.9.7.2 Copolymer Formation by Acid + Oxazoline Reaction A compatibilizing copolymer may be formed through reaction of carboxylic acid groups grafted onto a PO chain and olefinic oxazoline groups grafted onto a PP chain. Again, it is reasonable to consider that this copolymer may be crosslinked. This is based on the assumption that the oxazoline groups are truly grafted along the PP chains and are not all at terminal sites.

Table 5.39. Crosslinked copolymer formation in binary PO/PO blends Polyolefin / Polyolefin

Characterization and Comments

References

NBR (75-25) or hydrogenated NBR / PVC (25-75) LDPE (100-80) / PVC (0-20)

internal mixer at 180°C / rheometry / DMTA / DSC / FTIR / solvent swelling torque rheometer at 180°C / selective solvent extraction / FTIR / effects of PVC stabilizers

Manoj et al., 1993 a, b Ghaffar et al., 1980-81

Table 5.40. Crosslinked copolymer formation in binary PO/PO blends Polyolefin / Polyolefin

Characterization and Comments

References

chlorosulfonated PE (50) / carboxylated NBR (50)

two-roll mill / selective solvent extraction / DMA / FTIR / proposed carboxylate ester crosslink through displacement of chloride on PE by carboxylate two-roll mill / capillary rheometry / FTIR / 10-25 meq. per 100 g. sulfonate level / proposed sulfonate ester + ether crosslink rather than ionomer crosslink

Roychoudhury and De, 1997

EPDM-SO3 Zn ionomer (100-0) / ENR (0-100)

Manoj et al., 1994

Reactive Compatibilization of Polymer Blends

Liu and Baker [1994] and Liu et al. [1993] have prepared PO-PP blends containing 20 parts NBRg-AA, 0-80 parts PP, and 80-0 parts PP-g-IPO (0.2% grafted oxazoline) in an internal mixer. The blends were characterized by torque rheometry, SEM, FTIR, and mechanical properties vs. use of unfunctionalized NBR and vs. IPO level on PP. 5.9.8

Polyolefin + Polypropylene + Styrene Copolymer Ternary Blends

Copolymer Formation by Amine + Anhydride Reaction Dharmarajan et al. [1995] have prepared compatibilized blends of PO/PP/styrene copolymer. Blends of 100-0 parts PP, 0-100 parts SMA, 0-15 parts EP-g-1° amine (0.3 mol % amine), and 0-5 parts PP-2° amine (0.4 wt% amine) were combined in an internal mixer at 220°C. Blends were characterized by FTIR, DMTA, TEM, rheology, mechanical properties, lap shear adhesion, and paint adhesion. Properties were compared for blends containing either of the two aminefunctionalized polymers alone. Reaction of EP-g-1° amine with SMA results in a crosslinked copolymer because EP-g-1° amine contains a randomly distributed amine-functionalized comonomer in the backbone. PP-2° amine was prepared from PP-MA through extrusion with N-methyl-1,3-propanediamine. Dharmarajan et al. suggested that PP-MA is mostly end-group functionalized with a smaller portion of main chain-functionalized species as the result of trapping PP radicals before chain scission (see Section 5.8.8). Chains with combinations of end-group and main chain-functionalization should also be present since the authors report that their PP-MA contains a significant amount of highly functionalized PP-MA oligomer. Therefore, in these examples, the copolymers prepared from PP-MA and SMA may consist of a mixture of crosslinked and grafted species. For related works see Section 5.7.21.1.

5.9.9

393

Polyolefin + Polysiloxane Binary Blends

5.9.9.1 Copolymer Formation by Amine + Anhydride Reaction Kole et al. [1995] postulated the formation of a crosslinked copolymer through reaction between amine groups distributed along the main-chain of a polysiloxane and anhydride groups distributed along the main chain of a PO. A unique feature of this example is the use of acrylamide-grafted siloxane rubber as the source of amine groups. Amine groups alpha to carbonyl groups (as in acrylamide) are much less nucleophilic than typical amine functionality. In an example, 50/50 blends of EPDM-g-MA (1% MA) with acrylamide-grafted silicone rubber were mixed in an internal mixer at 35, 70 or 150°C. The blends were characterized by DMA, FTIR, mechanical properties, solvent swelling, and TGA. 5.9.9.2 Copolymer Formation Involving Mechano-chemical Radical Formation As shown in Table 5.41, copolymer formation was postulated between mechano-chemically generated radicals at tertiary sites on EMAc and vinyl groups on PDMS during melt processing [Santra et al., 1993]. The substitution pattern of vinyl groups on PDMS was not reported. Assuming that the vinyl groups are distributed along the PDMS chains (and not present only as end-groups), then the compatibilizing copolymer formed is a crosslinked copolymer. EMAc-PDMS copolymer formed in situ has also been used to compatibilize PDMS with thermoplastic polyurethane [Santra et al., 1995]. 5.9.10

Polyolefin + Styrene Copolymer Binary or Higher Blends

5.9.10.1 Copolymer Formation by Acid + Epoxide Reaction Kim, J. K. et al. [1997] and Kim, S. et al. [1997 a, b] have compatibilized PE/PS blends through blending 70-10 parts PE-g-AA with 30-90

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S. B. Brown

Table 5.41. Crosslinked copolymer formation in binary PO/Polysiloxane blends Polyolefin / Polysiloxane

Characterization and Comments

References

EMAc (100-0) / vinyl-functionalized PDMS (3.8% vinyl) (0-100) LDPE (50) / EMAc (0-10) / vinyl-functionalized PDMS (3.8% vinyl) (50)

internal mixer at 85-180°C / capillary rheometry / DMA / FTIR internal mixer at 180°C / mechanical properties / DMA / FTIR / SEM / WAXS / X-ray diffraction / lap shear adhesion

Bhattacharya et al., 1995; Santra et al., 1993a Santra et al., 1993b

parts PS in the presence of 0-3 parts S-GMA copolymer in an internal mixer at 200°C. Blends were characterized by SEM, TEM, rheology, and DMA. Effects of AA content on blend properties were studied. 5.9.10.2 Copolymer Formation by Acid + Oxazoline Reaction PO/styrene copolymer blends have been compatibilized through crosslinked copolymer formation between acid-functionalized PO and oxazoline-functionalized PS (Table 5.42). Anhydridefunctionalized PO is also effective in these blends since a proportion of acid groups is present from ring-opened anhydride. Oxazoline groups are most frequently introduced into PS by copolymerization of styrene with isopropenyl oxazoline. 5.9.10.3 Copolymer Formation by Alcohol + Anhydride Reaction Tselios et al. [1997] have compatibilized PO/styrene copolymer blends through crosslinked copolymer formation between PO alcohol groups and anhydride groups on styrene copolymer. Specifically, 50 parts EVAl (1.6-7.5% VAl) was mixed with 50 parts SMA (8.4-14.7 mol% MA) in an internal mixer at 200°C. The blends were characterized by torque rheometry, FTIR, DSC, TGA, selective solvent extraction, and mechanical properties as a function of mole ratio alcohol to anhydride. Blend properties were compared to those with EVAc in place of EVAl.

5.9.10.4 Copolymer Formation by Amine + Anhydride Reaction As shown in Table 5.43, PO-styrene copolymer blends have been compatibilized through crosslinked copolymer formation between amine-functionalized PO and anhydride-functionalized styrene copolymer.

5.10

Compatibilization by Covalent Crosslinking Mediated by an Activating Agent: Process #4b

Blends are discussed in alphabetical order of polymeric components. 5.10.1

Crosslinking Reactions Mediated by Addition of an Activating Agent. General Considerations

Covalent crosslinking reactions mediated by a third, added reagent might give the same type of copolymer structure as that which results from direct crosslinking reactions (see Table 5.4, Type 4b). In this case, the added reagent is a radical initiator or other type of activating agent such as a condensing agent. Such activating agents are usually not incorporated into the final copolymer. Radical initiators may promote radical formation on each of two immiscible polymers. A crosslinked copolymer results through radicalradical coupling between the two polymers at a melt phase interface. Self-coupling of each

395

Reactive Compatibilization of Polymer Blends

Table 5.42. Crosslinked copolymer formation in binary PO/Styrene copolymer blends Polyolefin / Styrene Copolymer

Characterization and Comments

References

EP-g-MA (0.7% MA) (20) / S-IPO (1% IPO) (80)

internal mixer at 200°C / torque rheometry / selective solvent extraction / SEM study of morphology development in reactive and in non-reactive blends internal mixer at 260°C / torque vs. time, temperature / SEM / FTIR / mechanical properties vs. use of unfunctionalized PS and oxazoline content in blend internal mixer at 240°C / FTIR / torque rheometry / capillary rheometry / also TSE with on-line rheometry and FTIR controls / optical microscopy internal mixer at 190°C / torque rheometry / FTIR / SEM internal mixer at 225°C / FTIR / SEM / selective solvent extraction / DSC / mechanical properties vs. processing conditions / effects of diluting with unfunctionalized LDPE and PS / effect of zinc chloride catalyst internal mixer at 185°C; torque vs. functionalization concentration / morphology / mechanical properties / detrimental effect of >5% IPO (too small dispersed phase particle size)

Scott and Macosko, 1994

NBR-g-AA (7% AA) (20) / S-IPO (1.7% IPO) (80) EAA (9% AA) (50-20) / PS (0-50) / S-IPO (1.2% IPO) (30-60) PE-g-MA (30) / S-IPO (1.7% IPO) (70) EAA (9% AA) (90-10) / S-IPO (1% IPO) (10-90)

B-AN-AA (20) / S-IPO (80) or combined with unfunctionalized PS

Liu and Baker, 1992b

Curry and Andersen, 1991/92 Liu et al., 1990 Saleem and Baker, 1990; Baker and Saleem, 1987 a, b Fowler and Baker, 1988

Table 5.43. Crosslinked copolymer formation in binary PO/Styrene copolymer blends Polyolefin / Styrene Copolymer

Characterization and Comments

References

amine-functionalized EP (0.3-0.5 mol% amine) (30-0) / HDPE (0-25) / SMA (8-14% MA) (65-100) amine-functionalized EP (0.3-0.5 mol% amine) (35-0) / SMA (8-14% MA) (65-100) amine-functionalized LLDPE (100-0) / SMA (6% MA) (0-100)

internal mixer at 180-220°C / mechanical properties / SEM / TEM / rheology / pre-reaction of functionalized EP and SMA / selective solvent extraction internal mixer at <220°C / mechanical properties / DSC / FTIR / SEM / rheology / selective solvent extraction internal mixer at 220°C / torque rheometry / SEM / mechanical properties / DSC / FTIR / MFI / PE grafted with dimethylamino-ethyl methacrylate or t-butyl-aminoethyl methacrylate

Datta et al., 1993

polymer may compete with cross-coupling as is the case with most 3-body reactions in which an

Dharmarajan and Datta, 1992 Song and Baker, 1992

added reagent is capable of reacting with each of the two immiscible polymers.

396

5.10.2

S. B. Brown

Polyester + Polyester Binary Blends

PEST-PEST Copolymer Formation by Radical Coupling Radical-radical coupling has been used to compatibilize blends of immiscible polyesters [Avella et al., 1996; Immirzi et al., 1994; Cavallaro et al., 1993]. Specifically, 70-30 parts poly(hydroxybutyrate-co-hydroxyvalerate) (4 mol% valerate) or polyhydroxybutyrate was mixed with 30-70 parts PCL in an internal mixer at 100° or 160°C in the presence of 0-0.5 parts DCP or DBP radical initiator. Blends were characterized by SEM, mechanical properties, selective solvent extraction, and FTIR. 5.10.3

Polyolefin + Polyolefin Binary or Higher Blends

PO-PO Copolymer Formation by Radical Coupling Immiscible blends of polyolefins have been compatibilized by copolymer formation brought about by radical coupling across a melt phase boundary, as shown in the examples listed in Table 5.44. The work of Kim et al. involved simultaneous copolymer formation between immiscible rubber and plastic phases with the dynamic vulcanization of the elastomeric phase.

5.10.4

Polyolefin + Polypropylene Binary Blends

PO-PP Copolymer Formation by Radical Coupling PO/PP blends have been compatibilized by addition of a radical initiator. Early examples were summarized in a review by Teh et al. [1994]; the more recent ones are listed in Table 5.45. PP undergoes chain-cleavage in the presence of radical initiator to give a terminal radical site. Coupling of this PP radical with a PE radical located on the PE chain initially yields a graft copolymer. Since the new copolymer may participate in further coupling reactions, a crosslinked copolymer may eventually be formed depending upon reaction time and radical concentration among other factors. 5.10.5

Polyolefin + Polyvinyl Chloride Binary or Higher Blends

PO-PVC Copolymer Formation by Radical Coupling As shown in Table 5.46, PO/PVC blends have been compatibilized by coupling in the presence of a radical initiator. In the work of Xu et al. [1997] the radical initiator may also serve to vulcanize the BR elastomeric phase.

Table 5.44. Crosslinked copolymer formation in binary PO/PO blends Polyolefin / Polyolefin

Characterization and Comments

References

EPDM (100-0) / LLDPE (0-100) / DCP RI (0.7) recycle mix of LDPE (45), HDPE (15), PVC (15), <10 PS, HIPS, PP, PET / L101 RI (0-2) LDPE (85) / EVAc (15) / DCP RI (1) LDPE (100-0) / LLDPE (0-100) / DCP RI (0-1) HDPE (100-0) / LLDPE (0-100) / DCP RI (0-1)

roll mill at 80°C / mechanical properties / rheology / DSC / SEM / X-ray diffraction TSE at 210°C melt T / morphology / selective solvent extraction / mechanical properties vs. RI level / rheology of mix and of individual resins vs. RI level (0-1%) internal mixer at 150-210°C / torque rheometry / effects of processing conditions on crosslinking / DSC / FTIR / DMA internal mixer at 160°C / torque vs. component ratio and DCP level / mechanical properties / selective solvent extraction internal mixer at 220°C / torque rheometry / mechanical properties / rheology

Kim et al., 1996 Vivier and Xanthos, 1994

Vogel and Heinze, 1993 Abraham et al., 1992 Kurian et al., 1992

397

Reactive Compatibilization of Polymer Blends

Table 5.45. Crosslinked copolymer formation in binary PO/PP blends Polyolefin / Polypropylene

Characterization and Comments

References

LDPE (100-0) / PP (0-100) / L130 RI (0-3) or t-butyl perbenzoate LDPE (100-0) / PP (0-100) / L101 RI (0-1)

internal mixer at 150 or 180°C / mechanical properties vs. RI and PP content / DSC / selective solvent extraction / optional addition of hydroquinone (0.75%) / also silica-filled blends internal mixer or TSE at 180°C / torque rheometry / rheology / SEM / mechanical properties / SEC / DSC / DMA / effects of PP-PE viscosity ratios / comparison to radiation induced crosslinking SSE at 185°C followed by static mixers at 200-220°C / mechanical properties/ SEM / DSC / SEC / TREF

Chodak et al., 1996; Chodak and Chorvath, 1993; Chodak et al., 1991 Yu et al., 1994, 1992, 1990

LLDPE (50) / PP (50) / L101 RI (0-0.25)

Cheung et al., 1990

Table 5.46. Crosslinked copolymer formation in binary PO/PVC blends Polyolefin / PVC

Characterization and Comments

References

LDPE (100-0)/ PVC (0-100) / butadiene rubber (0-5) / DCP RI (0-2) LDPE (40-38)/ PVC (57-60) / NBR (0-4) / DCP RI (0-0.02) LDPE (50) / PVC (50) / CHP RI (0-0.1)

two-roll mill at 155°C / SEM / optical microscopy / mechanical properties vs. BR presence two-roll mill at 155°C / morphology /mechanical properties vs. DCP level and NBR presence torque rheometer at 120-170°C / selective solvent extraction / FTIR / mechanical properties / DMTA / addition of copolymer to PE + PS

Xu et al., 1997

5.10.6

Polyolefin + Styrene Copolymer Binary or Higher Blends

Copolymer Formation by Radical Coupling As shown in Table 5.47, EPDM may be compatibilized with PS through addition of S-EP copolymer and a radical initiator [Crevecoeur et al., 1995]. It was proposed that a compatibilizing copolymer forms through cross-coupling of reactive sites on EPDM with radical sites on the PO portion of S-EP copolymer. Earlier work postulated direct copolymer formation between

Xu et al., 1993 Hajian et al., 1984

PE and PS when mixed at 170°C in the presence of radical initiator for about 50 minutes [Hajian et al., 1984]. Sulfonyl azides decompose thermally to give a nitrene intermediate that can react either through a singlet or triplet state. In the triplet state the nitrene may react by hydrogen abstraction to give a radical on a polymer chain. A mono-sulfonyl azide creates radical sites on both PO and PS that result in a compatibilizing copolymer through cross-coupling [Radusch et al., 1993].

398

S. B. Brown

Table 5.47. Crosslinked copolymer formation in PO/styrene copolymer blends Polyolefin /Styrene Copolymer

Characterization and Comments

References

EPDM (19) / S-EP (3) / PS (78) / various RI

TSE at 235°C / SEM / solvent swelling / mechanical properties vs. RI concentration internal mixer at 180°C / mechanical properties / FTIR / TEM torque rheometer at 120-170°C / selective solvent extraction / FTIR / mechanical properties / DMTA / addition of copolymer to PE + PS

Crevecoeur et al., 1995

PB (20) / PS (80) / aromatic sulfonyl azide (0.25-2) LDPE (50) / PS (50) / CHP RI (0-1)

5.11

Compatibilization by Covalent Crosslinking Mediated by a Coupling Agent: Process #4c

Blends are discussed in alphabetical order of polymeric components. 5.11.1

Crosslinking Reactions Mediated by Addition of a Coupling Agent. General Considerations

Covalent crosslinking reactions to form a compatibilizing copolymer may be performed by addition of a coupling agent (see also Section 5.5.3.2). Coupling agents react with the same type of functionality in each of the immiscible polymers and remain bound in the crosslinked product as linking agents. Examples include diepoxide reaction with pendent carboxylic acid groups on each of two immiscible polymers. In this case the structure of the crosslinked copolymer is that shown in Table 5.4 (Type 4C). Most commonly, coupling agents are multifunctional reagents with molecular weights less than about 1000. When immiscible polymer pairs are employed, each bearing multiple pendent nucleophilic groups (such as carboxylic acids), then low molecular weight coupling agents such as bis-, tris, tetra- and higher epoxides; bis-oxazolines, and other multifunctional electrophilic species are used. A low molecular weight coupling agent

Radusch et al., 1993 Hajian et al., 1984

bearing multiple olefinic sites may be used to promote crosslinking in blends containing PO’s. Common examples include commercially available tris-acrylates and triallyl isocyanurate. Often these crosslinking reactions are performed in the presence of radical initiator. Self-coupling of each polymer may compete with cross-coupling as is the case with most 3-body reactions in which an added reagent is capable of reacting with each of the two immiscible polymers. 5.11.2

Polyolefin + Polyolefin Binary Blends

Miscellaneous Coupling Agents Examples in Table 5.48 show use of coupling agents that can, in theory, react at sites along the main chain of each of two immiscible polyolefins. Both examples involve concomitant dynamic vulcanization of an elastomeric, dispersed phase with the added crosslinking agent. 5.11.3

Polyolefin + Polypropylene Binary or Higher Blends

5.11.3.1 Tris-acrylate + RI as Coupling Agent As shown in Table 5.49, PO/PP blends have been compatibilized by crosslinked copolymer formation mediated through addition of a tris-acrylate coupling agent and a radical initiator.

399

Reactive Compatibilization of Polymer Blends

Table 5.48. Crosslinked copolymer formation in PO/PO blends Polyolefin / Polyolefin

Characterization and Comments

References

EPDM (20) / EP (80) / m-phenylene bis-maleimide (0.075-0.6)

TSE at 210°C / mechanical and thermal properties / morphology / selective solvent extraction /evidence for EP-EPDM copolymer lab mill / SEM / swelling index / rheology / mechanical properties

Inoue, 1994 a, b

NR (100-0) / NBR-AA (0-100) / bis(diisopropyl) thiophosphoryl disulfide (3)

Naskar et al., 1994

Table 5.49. Crosslinked copolymer formation in PO/PP blends Polyolefin / PP

Characterization and Comments

References

PE (25-10) / PP (75-90) / Lupersol 101 RI (0-0.4) HDPE (10-0) / EPR (10-20) / PP (80) / RI (0-0.01)

TSE at 210°C / rheology / DSC / DMA / addition of trimethylolpropanetriacrylate TSE at 230°C / rheology / SEM / optical microscopy / DSC / mechanical properties / unidentified RI / unidentified methacrylate additive (0-0.3) TSE at 230°C / MFI / SEM / rheology / mechanical properties / use of PE-EPDM masterbatch / addition of trimethylolpropanetriacrylate TSE at 230°C / rheology / SEM / DSC / mechanical properties / optical microscopy / increased coupling efficiency + retardation of PP degradation by addition of trimethylolpropane triacrylate TSE at 230°C / DSC / capillary rheometry / SEM / mechanical properties / increased coupling efficiency + retardation of PP degradation by addition of trimethylolpropane triacrylate

Graebling et al., 1997

HDPE (8) / PP (80) / EPDM (12) / 1,3-bis(t-butylperoxy-isopropyl) benzene RI (0-0.3) EP (20) / PP (80) / 1,3-bis(t-butylperoxy-isopropyl) benzene RI (0-0.05%) NR (10) / PP (90) / 1,3-bis(t-butylperoxy)benzene RI (0-0.05%)

5.11.3.2 Miscellaneous Coupling Agents Examples in Table 5.50 illustrate the use of coupling agents that can react with both PO and PP in immiscible blends. Although PP may have at most one unsaturated site per chain [see Coran and Patel, 1983b], the immiscible rubber phase has multiple reactive sites. Depending upon stoichiometry of coupling agent and reactant polymers, the reaction may lead to crosslinking.

Do et al., 1996

Kim and Choi, 1996 a

Kim and Do, 1996

Yoon et al., 1995

In examples from Coran and coworkers an additional coupling agent could be used to dynamically vulcanize the elastomeric phase following copolymer formation with the PP matrix. 5.11.4

Polyolefin + Polyvinyl chloride Binary or Higher Blends

TAIC as Coupling Agent In Table 5.51 PO/PVC blends have been compati-

400

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Table 5.50. Crosslinked copolymer formation in PO/PP blends Polyolefin / PP

Characterization and Comments

References

EPDM (30-0) or PB / PP (70-100) / m-phenylene bis-maleimide (0.075-0.6) NBR (67.5-22.5) / dimethylol phenol-modified PP (25-75) / dimethylol phenol (0-10) / SnCl2 curing agent (0-1.1)

TSE at 210°C / mechanical and thermal properties / DMA / ductile-brittle transition temperatures / SEM / selective solvent extraction / DSC / MFR / evidence for PP-EPDM copolymer internal mixer at 190°C / mechanical properties vs. use of unfunctionalized PP / PP and dimethylol phenol(2-4%) premixed before addition of rubber / optional addition of additional coupling agent / addition of 0-7.5% amine-terminated NBR

Inoue and Suzuki, 1996, 1995; Ishikawa et al., 1996; Inoue, 1994 a, b Coran and Patel, 1983b

Table 5.51. Crosslinked copolymer formation in binary PO/PVC blends Polyolefin / PVC

Characterization and Comments

References

HDPE (50) / PVC (50) / NBR (0-15) / TAIC (0-2) / DCP RI (0-1.5) LDPE (80) / PVC (20) / TAIC (4) / DBP (0.5-1) or DCP RI

two-roll mill at 160°C / SEM / mechanical properties vs. NBR presence / selective solvent extraction internal mixer at 180°C / mechanical properties / SEM / 10% EA-MMA added as processing aid / also use of recycled PVC and PE

Fang et al., 1997 Ajji, 1995

Table 5.52. Crosslinked copolymer formation in PO/polystyrene or styrene copolymer blends Polyolefin / Styrene Copolymer

Characterization and Comments

References

LLDPE (90-10) / PS (10-90) / TAIC (0-0.35) + styrene monomer (0-7) / DCP RI (0-0.35) LLDPE (50) / styrene-co-vinylbenzaldehyde (50) / TAIC / DCP RI

TSE at 200°C / mechanical properties/ selective solvent extraction / DSC / SEC / FTIR / SEM / rheology

Teh and Rudin, 1992, 1991

TSE at 200°C / mechanical properties / selective solvent extraction (1-8% copolymer) / FTIR / morphology

Van Ballegooie and Rudin, 1988

bilized by crosslinked copolymer formation in the presence of a triallyl isocyanurate (TAIC) coupling agent and a radical initiator.

5.11.5

Polyolefin + Polystyrene or Styrene Copolymer Binary Blends

5.11.5.1 TAIC as Coupling Agent PO/styrene copolymer blends have been compatibilized by crosslinked copolymer formation in the

Reactive Compatibilization of Polymer Blends

401

presence of TAIC coupling agent and a radical initiator (Table 5.52).

5.12.1

5.11.5.2 Miscellaneous Coupling Agents

Immiscible polymer blends have been compatibilized through formation of a compatibilizing copolymer linked by ionic association instead of by covalent bonding. Although many examples have been published, most of these involve solutionmixing of the two immiscible polymers, viz. [Natansohn et al., 1990]. However, the examples given in this section describe only such polymer blends prepared by melt mixing. In theory, the possible architectures of a compatibilizing copolymer arising from ionic association may be the same as all the previously discussed architectures arising from covalent bond formation. However, in practice only a small number of copolymer architectures has been reported for compatibilizing agents arising from ionic association. In the most common examples (Table 5.4; Process #5a), ionizable groups such as carboxylic, sulfonic, or phosphonic acid are present in low concentrations (usually about 5%) on both polymers. The ionizable groups are at least partially neutralized by a mono-, di-, or tri-valent metal cation, such as Na+1, Zn+2, or Al+3. Multivalent cations may form a bridging linkage between the ionizable groups of the two immiscible polymers resulting in interchain copolymer formation by ion-ion association. Mono-valent cations such as Na+1 or K+1 may also be used to promote association through ion-dipole association. With either type of cation a morphology is formed in which there are concentrated domains of associated ionic species (ion clusters) in a matrix of the immiscible homopolymers. In the first type of ionic association (Process #5a), the ionizable functionalities of the two polymers are located in the pendent side groups. These polymers are prepared either through copolymerization with ion-containing monomers (or latent ion-containing monomers) or through subsequent grafting with such monomers. Consequently, the resulting compatibilizing agents are most often crosslinked copolymers with the structure shown in Table 5.4 as Type 5a. In a second type of ionic association (Process #5b) metal cations may mediate association between

Bromobutyl rubber and SMA blends have been compatibilized through addition of a bifunctional coupling agent [Willis et al., 1990]. The authors assumed that a low molecular weight aminoalcohol reacted with rubber bromo groups to form a quaternary ammonium salt. The resulting alcohol-functionalized rubber could form crosslinked copolymers through reaction of hydroxy groups with anhydride of SMA. Blends of 30-5 parts bromobutyl rubber with 70-95 parts SMA and 0-8 parts 2-dimethylaminoethanol coupling agent were prepared in an internal mixer at 150-200°C, followed by TSE processing at 200°C. The blends were characterized by FTIR, SEM, and mechanical properties. The effects of premixing rubber with amine and the effects of processing conditions were studied. Model reactions for the proposed copolymer-forming reaction were carried out in solution. 5.11.6

Polypropylene + Styrene Copolymer Binary Blends

Bis-maleimide as Coupling Agent PP/styrene copolymer blends have been compatibilized by addition of bis-maleimide coupling agent capable of reacting with both polymers. Inoue [1994 a, b] has prepared blends containing 80 parts PP and 20 parts SIS (or SBS) in the presence of 0-0.3 parts m-phenylene bis-maleimide coupling agent on a TSE at 210°C. The blends were characterized by mechanical properties, melt flow rate, DSC, morphology, and selective solvent extraction. Evidence was presented for PP-SIS copolymer formation.

5.12

Compatibilization by Ionic Interaction: Process #5

Blends are discussed in alphabetical order of polymeric components.

Types of Ionic Reactions and Characteristics

402

S. B. Brown

ionic groups on one polymer and neutral donor groups on a second immiscible polymer. Typical ionic groups are again carboxylic, sulfonic, and phosphonic acids. Neutral donor groups contain atoms, usually nitrogen or phosphorus, having unshared pairs of electrons capable of coordinating to metal cations. Such groups include pyridine, quinoline, and phenanthroline. These are usually introduced into a polymer via copolymerization with the vinyl analog, e.g., vinylpyridine. Again, the structure of the compatibilizing copolymer is usually crosslinked. In a third type of ionic association (Process #5c) acidic groups such as carboxylic acids bound to one polymer may mediate interchain copolymer formation by protonation of basic groups on a second, immiscible polymer. In this case the compatibilizing agent is again most often a crosslinked copolymer. General discussions of the properties of polymers containing ionizable groups (ionomers) have been published [Rees, 1986; Lundberg, 1987; Hara and Sauer, 1994]. Ionic crosslinks are usually thermally reversible, which may limit the usefulness of blends containing them in certain commercial applications. Since ionomers are initially selfassociated through ionic bonds, thermal reversibility of ionic crosslinks in the melt is necessary to overcome homogeneous, self-crosslinking within each homopolymer before heterogeneous, interchain crosslinking can occur. Often a high degree of plasticization of the ion-containing polymer melt is required so that high processing temperatures that might lead to polymer decomposition

need not be used. In some cases, polymers containing masked ionomeric functionality, i.e., chemical groups that form ionic species during extrusion, have been used to form copolymers during reactive processing. Use of masked ionomers may require lower energy during extrusion since ionic selfassociation does not have to be overcome before interchain copolymer formation can occur. 5.12.2

Ion-Ion Association Mediated by Metal Cations: Process #5a

5.12.2.1 Polyether + Polyolefin Binary Blends Mascia and Bellahdeb [1994 a, b] have compatibilized Phenoxy polyether resin with an acidfunctionalized polyacrylate resin through crosslinking. A base, capable of forming alkoxide from the polyether hydroxy groups, should be used. Thus, 75-25 parts Phenoxy resin and 25-75 parts E-tBA-AA terpolymer (or EMAA Na ionomer) were blended in an internal mixer or TSE at 180-200°C in the presence of 2 parts NaOEt. Blends were characterized by selective solvent extraction, SEM, DSC, and rheology. The resulting crosslinked copolymers were used to compatibilize PET + HDPE. 5.12.2.2 Polyolefin + Polyolefin Binary Blends Examples of some immiscible polyolefin blends that have been compatibilized through ionomeric crosslinking between ionomeric groups on each polyolefin are listed in Table 5.53.

Table 5.53. Ion-Ion association in PO/PO blends mediated by metal cations Polyolefin / Polyolefin

Characterization and Comments

References

EMAA Zn salt (100-0) / EPDM-g-MA (1% MA) (0-100) / zinc oxide (10) EMAA Zn salt (70-30) / EPDM-SO3 Zn salt (30-70)

internal mixer at 170°C / mechanical properties / DMA / FTIR / effects of multiple thermal histories

Datta et al., 1996

internal mixer at 170°C / mechanical properties / DMA / dielectric thermal analyses / FTIR / effects of multiple thermal histories

Kurian et al., 1996

403

Reactive Compatibilization of Polymer Blends

Table 5.54. Ion-Ion association in PO/PPE/styrene copolymer blends mediated by metal cations Polyolefin / PPE / Styrene copolymer

Characterization and Comments

References

EPDM-SO3 Zn salt (12) /

TSE at 270°C / mechanical properties vs. omission of PPE-SO3 Zn salt / replacement of EPDM-SO3 Zn salt with PDMS-SO3 or B-AN-AA Na salt

Golba and Seeger, 1987; Campbell et al., 1986

TSE at 240°C / mechanical properties vs. use of unfunctionalized rubber or unfunctionalized PS / model compound studies / EP-PO3Et forms EP-PO3 Zn salt in presence of ZnSt

Brown and McFay, 1986

PPE (52) / PPE-SO3 Zn salt (13) or PS-SO3 Zn salt / PS (23) / TPP (18) / ZnSt (12) EP-PO3 Et (9-0) / PPE (50) / PS (9-20) / PS-SO3 Zn salt (0-9) / ZnSt (0-9) / TPP (14)

5.12.2.3 Polyolefin + Polyphenylene Ether + Styrene Copolymer Ternary Blends Immiscible polyolefin impact modifiers have been compatibilized with PPE/PS blends by employing polyolefin sulfonate ionomers and including PS sulfonate ionomer in the blend (Table 5.54). In these compositions, the polyolefin ionomer can crosslink with PS ionomer. The latter is miscible with PS, which in turn is miscible with PPE. Simple ionomers have been used [Golba and Seeger, 1986; Campbell et al., 1986]. In similar blends, a masked ionomer, such as a polyolefin phosphonate ester, could also be used [Brown and McFay, 1986]. The masked ionomer is melt processable but generates a polyolefin phosphonate ionomer in situ during extrusion with zinc stearate. The polyolefin phosphonate ionomer can crosslink with PS sulfonate ionomer that is miscible with PS which in turn is miscible with PPE. 5.12.3

Ion-Neutral Donor Group Association Mediated by Metal Cations: Process #5b

5.12.3.1 Polyamide + Styrene Copolymer Binary Blends Parker and Hara [1997] have prepared compatibilized blends of poly(p-phenylene terephthalamide) potassium amide salt (5-0 parts) with poly-4-

vinylpyridine (95-100 parts) or poly-2-vinylpyridine by compression molding at 180-190°C. The blends were characterized by optical microscopy, DMA, and DSC. Evidently, the copolymer structure consists of poly(p-phenylene terephthalamide) anions associated through potassium cation with poly-2-vinylpyridine by ion-neutral donor group association. Although melt mixing was not used to prepare this blend, it is included as an instructive example since the compatibilization strategy may be applicable to other compositions that may be prepared by intensive melt mixing. 5.12.3.2 Polyolefin + Styrene Copolymer Binary Blends Compatibilized blends of 77 parts EPDM-SO3Zn salt and 9 parts S-co-4-vinylpyridine with 4 parts ZnSt plasticizer were prepared in an internal mixer at 200°C [Lundberg et al., 1988; Agarwal et al., 1987; Peiffer et al., 1986]. The blends were characterized by FTIR, DMA, melt viscosity, DSC, and SEM. Mechanical properties were compared to blends containing unfunctionalized PS or containing EPDM-SO3Na or Mg salts. A copolymer linked by ionomeric crosslinks may form associations between sulfonate anion and pyridine nitrogen mediated by Zn cation.

404

S. B. Brown

Table 5.55. Ionic association in immiscible blends mediated by interchain protonation Polymer / Polymer

Characterization and Comments

References

EAA (6.5 or 9% AA) (100-0) / poly(vinyl amine) (0-100) or poly(vinyl alcohol-co-vinyl amine) PA-6 (80-60) / PS-SO3H

internal mixer at 110°C / DMA / DSC / mechanical properties

Robeson et al., 1996

internal mixer at 250°C / torque rheometry / DMA / DSC / SEM / NMR / selective solvent extraction /SEC

Molnar and Eisenberg, 1992, 1991

(0-9.7 mol% acid)(20-40)

5.12.4

Ionic Association Mediated by Interchain Protonation: Process #5c

As shown in Table 5.55, copolymers have been formed through ionic association based on protonation of a basic polymer by an immiscible acidic polymer. See also [Coran and Patel, 1983b] in Section 5.8.14.

5.13

Summary

The papers cited in this chapter reach similar conclusions concerning the effects of compatibilizing copolymer formation on blend properties. In virtually all cases, the generated blend morphology shows a smaller dispersed phase particle size than that observed in the absence of copolymer formation. In the majority of cases, this morphology is stabilized against agglomeration and coalescence of the dispersed phase during subsequent thermal processing. The specific morphology size distribution and its stability result in improved mechanical properties, not observed in uncompatibilized blends. A few efforts have been made to characterize the phase interface and to correlate the interfacial thickness with the level of copolymer formed. However, in most papers the existence of copolymer is simply inferred from secondary evidence and its level is not quantified. Except for a few notable cases there was seldom consideration in these papers about the architectures of copolymers generated and their

possible effects on compatibilization efficiency and morphological stability. Furthermore, since morphological stability alone may not result in optimum physical properties, it is sometimes not possible to know if optimum physical properties have been obtained. In numerous papers clearly more than one type of copolymer architecture can form. For example, in acid-epoxide reactions, a new secondary alcohol may form that may possibly equilibrate into polyester constituting one blend component. Only a few papers discuss the consequences of using di- vs. mono-end-groupfunctionalized polymers and their possible effects on generating graft vs. crosslinked copolymers. In most papers the level and type of functionality on the reacting polymers are not specified. There are numerous cited examples of blends where at least one polymeric component is vulcanizable. Such systems are quite difficult to analyze after processing. In some cases, formation of an interpenetrating network (IPN) may be mistaken for covalently bonded copolymer. This is particularly so when no deliberate effort has been made to promote vulcanization and its occurrence was adventitious. Selective solvent extraction experiments to detect copolymer can be misinterpreted if some polymer is physically occluded in a matrix of the second polymer. Some polymers contain reactive functionality but are also themselves subject to mechanochemical radical generation. When such polymers are blended under high-shear mixing with a second functionalized polymer, the architectures

Reactive Compatibilization of Polymer Blends

of formed copolymer may derive both from the primary, expected reaction and also from an unexpected, radical-radical coupling process. Frequently, in the cited examples a third, multifunctional reagent is added that reacts with both polymers. The relative solubility of many such reagents in particular polymers is unknown. In immiscible blends the reagent may segregate into one phase and selectively couple or crosslink that phase, in competition with the desired reaction between both polymers at the phase interface. Again, an IPN formation may result, leading to confusion when interpreting selective solvent extraction data. For some types of reagents the problem of phase segregation can be solved by adding a stoichiometric excess of the reagent. The reagent caps the reactive functionality, but it does not couple or chain-extend the polymer in which it is selectively soluble. Thus, reaction of the newly functionalized polymer at a phase interface with a second polymer is facilitated. This strategy works when using a coupling agent and in certain cases when using a condensing agent but not always when using a non-selective activating agent. However, relative stoichiometries of reacting species are seldom considered in the cited examples. There have been very few studies of kinetics of copolymer-forming reactions under melt processing conditions because it is quite difficult to generate these data. In the papers cited in this chapter, there are many examples where not only the level but also the architecture of the copolymer strongly depends on the melt processing time. Many of the blends are “living” in the sense that more copolymer would be generated with additional time. In contrast, some types of copolymer linkage (e.g., the reaction product of an alcohol and a cyclic anhydride) may degrade and copolymer may be lost with additional processing time or at higher processing temperature. Different copolymer architecture may form with additional processing time, particularly when a redistribution reaction is possible as either the primary (Section 5.6) or as the secondary reaction (e.g., when a reactive alcohol is generated from epoxide ring-opening with acid).

405

It is usually difficult to isolate and characterize a copolymer from a melt-processed polymer blend. Model studies of copolymer formation between immiscible polymers have been performed either in solution (where there is unlimited interfacial volume for reaction) or using hotpressed films of the polymers (where the interfacial volume for reaction is strictly controlled at a fixed phase interface). Model studies using low molecular weight analogs of the reactive polymers are useful but their applicability to high molecular weight reacting systems is limited. It is to be hoped that future work on Reactive Compatibilization will combine the excellent materials science that has been done to date with additional considerations of the chemical processes occurring. Such knowledge of the chemistry, coupled to fluid mechanics and morphology development models, would provide a powerful tool for optimization of known and invention of new Reactive Compatibilization processes to prepare commercially valuable polymer blends.

5.14

References

The references below are listed in alphabetical order by first author, then by second author, etc. If the same authors are present, the reference is listed by date with most recent first. If similar authors and the same year, then the reference is listed by year followed by a letter in alphabetical order. Abbate, M., Di Liello, V., Martuscelli, E., Musto, P., and Ragosta, G., Polymer, 33, 2940 (1992). Abbate, M., Martuscelli, E., Ragosta, G., and Scarinzi, G., J. Mater. Sci., 26, 1119 (1991). Abdou-Sabet, S., Puydak, R. C., and Rader, P., Rubber Chem. Technol., 69, 476 (1996). Abis, L., Braglia, R., Camurati, I., Merlo, E., Natarajan, K. M., Elwood, D., and Mylonakis, S. G., J. Appl. Polym. Sci., 52, 1431 (1994). Abraham, D., George, K. E., and Francis, D. J., Angew. Makromol. Chem., 200, 15 (1992). Abu-Isa, I. A., Jaynes, C. B., and O’Gara, J. F., J. Appl. Polym. Sci., 59, 1957 (1996). Aerdts, A. M., Groeninckx, G., Zirkzee, H. F., van Aert, H. A. M.,

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Appl. Poly. Sci., 62, 67 (1996). Xiaomin, Z., Zhihui, Y., Tainhai, N., and Jinghua, Y., Polymer, 38, 5905 (1997). Xu, C., Fang, Z., and Zhong, J., Angew. Makromol. Chem., 212, 45 (1993); Polymer, 38, 155 (1997). Yeh, J. T., and Fan-Chiang, C. C., J. Appl. Poly. Sci., 66, 2517 (1997). Yeh, J. T., Fan-Chiang, C. C., and Yang, S. -S., J. Appl. Poly. Sci., 64, 1531 (1997). Yoon, K. H., Lee, S. C., Park, I. H., Lee, H. M., Park, O. O., and Son, T. W., Polymer, 38, 6079 (1997). Yu, D. W., Xanthos, M., and Gogos, C. G., Adv. Polym. Technol., 10, 163 (1990); ibid., 11, 295 (1992); J. Appl. Poly. Sci., 52, 99 (1994); ibid., 56, 239 (1995). Yukioka, S., and Inoue, T., Polymer, 35, 1182 (1994). Zhang, R., Luo, X., and Ma, D., J. Appl. Poly. Sci., 55, 455 (1995). Zhang, X., Li, X. L., Wang, D., Yin, Z., and Yin, J., J. Appl. Poly. Sci., 64, 1489 (1997). Zhang, X., and Yin, J., Polym. Eng. Sci., 37, 197 (1997). Zhang, X., Yin, Z., and Yin, J., J. Appl. Poly. Sci., 62, 893 (1996). Zheng, W. G., Qi, Z. -N., and Wang, F. S., Polymer International, 34, 307 (1994 a). Zheng, W. G., Wan, Z. H., Qi, Z. N., and Wang, F. S., Polymer, 34, 4982 (1993); Polymer International, 34, 301 (1994 b). Zhihui, Y., Xiaomin, Z., Yajie, Z., and Jinghua, Y., J. Appl. Poly. Sci., 63, 1857 (1997).

5.14 AA AN B BA tBA BPA CHP CL DBP DCP E EA EAA EB

Abbreviations acrylic acid acrylonitrile butadiene n-butyl acrylate t-butyl acrylate bisphenol A cumene hydroperoxide caprolactone dibenzoyl peroxide dicumyl peroxide ethylene ethylacrylate ethylene acrylic acid copolymer ethylene butene copolymer

Reactive Compatibilization of Polymer Blends

EEA EMAc EMAA EMMA ENR EP EVAc EVAl FA GMA HIPS I IA IM IPO L101 L130 LCP MA MAA MAc MMA NA NBR NR P-E-P PA PB PBT PC PCE PCL PCT PDMS PE PEST PET PETG Phenoxy

ethylene ethylacrylate copolymer ethylene methylacrylate copolymer ethylene methacrylic acid copolymer ethylene methyl methacrylate copolymer epoxidized natural rubber ethylene propylene copolymer ethylene vinyl acetate copolymer ethylene vinyl alcohol copolymer fumaric acid glycidyl methacrylate high-impact polystyrene isoprene isobutyl acrylate impact modifier isopropenyl oxazoline Lupersol 101 (2,5-dimethyl-2,5-dit-butylperoxy hexane) Luperox 130 (2,5-dimethyl-2,5-di-tbutylperoxy hexyne) liquid crystalline polymer (polyester type unless noted) maleic anhydride methacrylic acid methyl acrylate methyl methacrylate nadic anhydride nitrile-butadiene rubber natural rubber propylene-ethylene-propylene block copolymer polyamide(s) polybutadiene polybutylene terephthalate(s) bis-phenol A polycarbonate polycarbonate ester copolymer polycaprolactone poly(cyclohexanedimethanol terephthalate) polydimethylsiloxane polyethylene polyester(s) polyethylene terephthalate(s) poly(ethylene-co-1,4-cyclohexylene dimethylene terephthalate) copolymer of BPA and

PMMA PO PPE PPS PPVL PS PVAc PVAl RI S SAN SB SBS SEBS

SI SIS TAIC TBAB TPE TPP TPPite VTMS ZnSt

415

epichlorohydrin poly(methylmethacrylate) polyolefin polyphenylene ether polyphenylene sulfide polypivalolactone polystyrene polyvinyl acetate polyvinyl alcohol radical initiator styrene styrene-acrylonitrile copolymer styrene-butadiene copolymer styrene-butadiene-styrene terpolymer styrene(ethylene/butylene)-styrene copolymer styrene-isoprene copolymer styrene-isoprene-styrene terpolymer triallyl isocyanurate tetrabutylammonium bromide thermoplastic elastomer triphenyl phosphate triphenyl phosphite vinyl trimethoxysilane zinc stearate

CHAPTER 6

INTERPENETRATING POLYMER NETWORKS

L. H. Sperling1 and R. Hu2

1

Lehigh University, Bethlehem, PA, USA and

2

Elf Atochem North America, Inc., King of Prussia, PA

6.1

Introduction

An interpenetrating polymer network, IPN, is defined as a blend of two or more polymers in a network form, at least one of which is synthesized and/or crosslinked in the immediate presence of the other(s). An IPN can be distinguished from polymer blends, blocks, or grafts in two ways: (1) an IPN swells, but does not dissolve in solvents, and (2) creep and flow are suppressed. 6.1.1

Definitions

There are several kinds of IPNs: Sequential IPN. First, polymer network I is synthesized. Then, monomer II plus crosslinker and activator are swollen into network I and polymerized in situ, see Figure 6.1A. Simultaneous Interpenetrating Network, SIN. The monomers and/or prepolymers plus crosslinkers and activators of both components are mixed, followed by simultaneous polymerization via non-interfering reactions, see Figure 6.1B. Typical syntheses involve chain and step polymerization kinetics. While both polymerizations proceed simultaneously, the rates of the reactions are rarely identical. Latex IPN. The polymers are made in the form of latexes, each particle constituting a micro-IPN. Depending on the rates of monomer addition relative to the rates of polymerization, various degrees of interpenetration and/or core/shell morphologies may develop. There are several kinds of latex IPNs, see Section 6.4. Gradient IPN. In this case, the overall composition or crosslink density of the material varies from location to location on the macroscopic level. One way of preparing these materials involves partial swelling of polymer network I by the monomer II mix, followed by rapid polymerization before diffusional equilibrium takes place. Films can be made with polymer network I predominantly on one surface, and polymer network II predominantly on the other surface with a gradient composition existing throughout the interior. L.A. Utracki (Ed.), Polymer Blends Handbook, 417-447. © 2003 Kluwer Academic Publishers. Printed in the Netherlands.

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Thermoplastic IPN. When physical crosslinks rather than chemical crosslinks are utilized, the materials may flow at elevated temperatures. As such, they are hybrids between polymer blends and IPNs. Such crosslinks commonly involve block copolymers, ionomers, and/or semicrystallinity.

Figure 6.1. Two synthetic methods for preparing interpenetrating polymer networks. A, sequential IPNs, and B, simultaneous interpenetrating networks, SINs.

Table 6.1.

Semi-IPN. These are compositions in which one or more polymers are crosslinked, and one or more polymers are linear or branched. 6.1.2

History of IPNs

While the field of IPNs is considered younger than the corresponding field of polymer blends, blocks, and grafts, some of the early materials predated those commonly ascribed to the corresponding blend, block and graft fields. To complicate the issue, the patent literature reveals that the topology of IPNs was invented over and over again. A brief summary is given in Table 6.1. The first such invention was by Aylsworth [1914] (Table 6.1, No. 1), who was at the time the chief chemist of Thomas A. Edison in the latter’s West Orange, NJ laboratory. He was also an independent inventor, working part time in his own laboratory [Sperling, 1987]. In those years, Edison had switched from the cylinder-type phonograph records to the platter type. The latter ones were made of the new phenol-formaldehyde material, just invented by Leo Baekeland. The problem with the new material was that it was extremely brittle, hence the new platters needed to be very thick. Aylsworth’s solution to the problem was to mix in natural rubber and sulfur, which on heating forms a network. Since the phenol-formaldehyde compositions are all densely crosslinked, the overall composition was a simultaneous interpenetrating network.

Early IPN patents

No. Polymer 1

Polymer 2

Application

Inventor

Pat. No.

1

Phenol-formaldehyde

Natural rubber

J. W. Aylsworth

2

Natural rubber

PVC

Toughen phonograph records Plastic materials

3

Poly(methyl methacrylate) Positively charged network

Poly(methyl methacrylate) Negatively charged network

U.S. Pat., 1,111,284; 1914 Ger. Pat., 623,351; 1935 U.S. Pat., 2,539,377; 1951 Br. Pat., 728,508; 1955

4

Smooth surfaced plastics Ion exchange resin

H. Hopff J. J. P. Staudinger and H. M. Hutchinson G. S. Solt

Interpenetrating Polymer Networks

In 1935, Dr. Heinrich Hopff of I. G. Farbenindustrie patented two major advances in emulsion polymerization and latex formation (see No. 2): (1) Various vinyl monomers such as vinyl chloride were added to a natural rubber latex dispersion, and polymerized to make a coreshell structure. (2) Various vinyl latexes, previously polymerized, were mixed with natural rubber latex, both still in aqueous dispersion, and then coagulated. Thus, core-shell latex and latex blend materials were known very early. (3) The coagulated masses of both (1) and (2) above could be masticated with sulfur and vulcanized, as mentioned in the patent, producing semi-IPNs. Staudinger and Hutchinson (No. 3 in Table 6.1) were concerned with the surface unevenness of the new cast-polymerized plastics, as prepared between glass plates [Staudinger and Hutchinson, 1951]. While the plastics were clear, objects behind them appeared distorted. They started with either crosslinked PMMA or crosslinked PS, and added the same monomer mix as before, and polymerized in situ. This creates a homo-IPN. The intent was that by stretching polymer network I taut via the swelling action of monomer mix II, the surface waviness would be reduced, thus improving the visual characteristics for windows, etc. It must be remarked that J. J. P. Staudinger was Herman Staudinger’s son. The latter was the person who enunciated the now famous Macromolecular Hypothesis in 1920, stating that there was a certain class of colloids actually composed of long chains. In 1955, Solt discovered that IPNs with one network charged negatively and the other charged positively made superior ion exchange resins [Solt, 1955] see No.4. An important feature of these materials relates to having the two phases in juxtaposition on a very small scale. If the system is molecularly mixed, however, efficiency declines due to the interactions between the opposite polymer ions, ejecting the mobile salt ions.

419

The term interpenetrating polymer networks was coined by John Millar in 1960, who prepared homo-IPNs (an IPN with both polymers identical) of polystyrene [Millar, 1960]. Millar knew about Solt’s work, and his objective was to increase the size of suspension particles of polystyrene intended for ion-exchange applications. Starting in the late 1960s, the field of IPNs was examined by three teams of investigators. Papers by Frisch et al. on IENs, Sperling and Friedman on sequential IPNs, and Lipatov and Sergeeva on filled materials started systematic research to establish the field of IPNs. By the year 1979, there were approximately 75 patents and 125 papers in the field [Lipatov and Sergeeva, 1967; Sperling and Friedman, 1969; Frisch et al., 1969a]. Today, that many IPN patents and papers are produced each year [Kim and Sperling, 1997]. 6.1.3

A Brief Literature Survey

Some of the most recent IPN literature includes edited works by Klempner and Frisch, Klempner, et al., Sperling and Kim, and Kim and Sperling [Klempner and Frisch, 1993; Klempner et al., 1994; Sperling and Kim, 1996; Kim and Sperling, 1996]. A book by Sperling appeared in 1997 with a chapter on IPNs [Sperling, 1997]. Today, IPNs are widely accepted as a branch of multicomponent polymer materials, and as such are discussed in many general polymer textbooks [Elias, 1997] as well as more specialized monographs [Utracki, 1989] and edited works [Gergen et al., 1996].

6.2

Synthetic Methods

All interpenetrating polymer networks utilize two different polymers. The exception involves the homo-IPNs, where both polymers are identical [Millar, 1960; Siegfried et al., 1979]. While these polymers may be synthesized by any of the known methods of polymer synthesis, some methods clearly work better in given objectives than others. The principal kinetic methods used are chain and step polymerization.

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Many sequential IPNs have used two chain polymerizations, where monomer mix I is polymerized, monomer mix II is swelled in, followed by polymerization of monomer mix II. Typical examples involve poly(ethyl acrylate) and PS [Huelck et al., 1972], and SBR and PS [Curtius et al., 1972]. Table 6.2. Selected crosslinkers for network formation

A host of crosslinkers have been employed, see Table 6.2. The mixed functionality crosslinkers often serve as grafting sites between networks of quite different nature, perhaps one network chain polymerized, and the other step polymerized. Then, each type of functionality reacts with each type of monomer during polymerization.

Interpenetrating Polymer Networks

Where a step polymerization is used, almost always it is for the first polymer synthesized in a sequential IPN. The reasons involve the slow diffusion into a pre-existing network of most monomers used in step polymerization, and the relatively high glass transition of step polymerized polymers. The latter reason is important because in order for diffusion and concomitant polymerization to occur rapidly, polymer network I should be above its glass transition at the temperature of polymerization of monomer mix II. Table 6.2 presents glycerol as a simple trifunctional crosslinker for step polymerized materials, suitable for polyesters and polyurethanes. For simultaneous interpenetrating networks (SINs), two independent, non-interfering reactions are required. Thus, a chain and a step polymerization have been the method of choice for many such polymerizations. Typical examples have involved PS and polyurethanes [Hourston and Schafer, 1996; Mishra et al., 1995], and PMMA. A key factor in the kinetics of such polymerizations is to keep the system above the glass transition temperature of both components. If the glass transition of either the polymer network I or polymer network II rich phase vitrifies, the polymerization in that phase may slow dramatically. There are several interesting polymerization schemes intermediate between a sequential IPN and an SIN. For example, in in situ prepared sequential IPNs, both monomers are polymerized via free radical reaction [He et al., 1993; Rouf et al., 1994]. The two monomers must have quite different reactivities towards the free radicals. This situation arises with vinyl or acrylic double bonds and allylic double bonds. The allylic double bonds react about 100 times slower than acrylic or methacrylic bonds. Often, two initiators are used, one reacting at a lower temperature, and the other at a higher temperature. In one of the systems studied, based on methyl methacrylate and diallyl carbonate of bisphenol-A (DACBA), first, crosslinked PMMA was formed at moderate temperatures. Then, by just increasing the temperature after completion of the first polymerization, the synthesis of the allylic network followed.

421

In this case, the initiator azobisisobutyronitrile, AIBN, was used to polymerize the MMA at 60°C, and then the sample was heated to 95°C for the polymerization of the DACBA with t-butyl peroxy isononanate (TBPIN). Latex IPNs offer unique synthetic opportunities. Since an IPN double network, ideally, is contained in each sub-microscopic latex particle, special effects are possible. The simplest case involves a crosslinked seed latex particle that is polymerized first. Then, monomer mix II is added. There are two subclasses. First, all of the monomer mix II can be added at once, or at least far more rapidly than the polymerization takes place. In that case, the monomer will first swell the latex particle, and then the excess monomer forms a shell around the swollen core. If the monomer mix II is added slowly, or more slowly than the initiator can polymerize the material, little monomer can swell into the particle, and a better defined core/shell structure develops. There are other cases to consider, such as the polarity of the two polymers. Usually, the more polar polymer will prefer to be in contact with the aqueous phase (thermodynamics considerations). Without crosslinking, there can be a phase inversion from what was stated in the previous paragraphs. The presence of crosslinking, however, limits such inversion (where polymer network II forms the core), because of the swelling restrictions on polymer network I imposed by the crosslinks. Then, in the general case, one must ask what relative rates of polymerization are taking place, what level of crosslinkers are employed, etc., leading to many current research topics [Lovell and El-Aasser, 1997]

6.3

Morphology and Glass Transitions

The previous section showed how IPNs and related materials can be synthesized. The several synthetic methods, such as sequential, simultaneous, latex, and thermoplastic IPN formation, will result in different morphologies. One of the main advantages of IPN synthesis relates to the ease of promoting dual phase continuity, i.e., for a

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two-component system, the two components will be phase separated, and each will exhibit phase continuity throughout the material. Because most IPNs are phase separated, usually they exhibit multiple glass transitions, one for each phase. However, since the phases are small with large volumes of interphase material, and if significant mixing of the components takes place, the glass transitions may be broadened or moved towards each other. Another advantage of using IPNs involves its thermosetting characteristics. By definition, IPNs will not flow when heated. A partial exception is the thermoplastic IPNs, which behave crosslinked at ambient temperatures, but flow at elevated temperatures. While some IPNs are tough, impactresistant plastics, the crosslinking permits many other types of applications, such as sound and vibration damping, biomedical, adhesives and coatings uses, etc. (see Section 6.5). 6.3.1

tions of the polybutadiene and the polystyrene. The bottom left and right are the full sequential IPNs, with lower and high crosslink levels in the polybutadiene, as indicated.

Morphology via Electron Microscopy

A powerful method of examining the morphology of many multicomponent polymer materials utilizes transmission electron microscopy [Woodward, 1989]. If the two phases are nearly equal in electron density, staining with osmium tetroxide or other agents can be used. For more detailed discussion on the methods of morphology characterization, see Chapter 8. Donatelli et al. [Donatelli et al., 1976] examined the morphology of net-polybutadiene-internet-polystyrene, and compared these with those of the corresponding semi-IPNs and polymer blends, see Figure 6.2. In the absence of crosslinks, a solution graft copolymer develops. If the solution is stirred during the polymerization, then a phase inversion takes place, see Figure 6.2, upper left. This material is the well-known high-impact polystyrene, HIPS. On quiescent polymerization, the material fails to undergo phase inversion, and polybutadiene remains the continuous phase, see Figure 6.2, upper right. This last material was invented by Ostromislensky [Ostromislensky, 1927]. The middle left and right compositions are, respectively, the semi-I and semi-II composi-

Figure 6.2. Selected morphologies of IPNs based on SBR and polystyrene, SBR stained with osmium tetroxide. Upper left: Commercial high-impact polystyrene. Upper right, Ostromislensky’s material, with no phase inversion; middle left, semi-IPN, SBR crosslinked; middle right, semi-II IPN, PS crosslinked; lower left, full IPN; lower right, full IPN, higher crosslinking in the SBR.

The full IPNs shown here (as in numerous other cases) have dual phase continuity. The domains, as cut in thin section for transmission electron microscopy, appear to be ellipsoidal. Actually, they are more probably thin sections of cylinders, cut at various angles. Other studies show that both phases may be continuous. Spinodal decomposition kinetics, thought to apply in many such cases, results in interconnected cylinders [Utracki, 1994].

Interpenetrating Polymer Networks

The upper left has an Izod impact resistance of approximately 80 J/m, the middle left about 150 J/m, and the bottom left about 250 J/m. These early results encouraged the development of these materials. The domain size of sequential IPNs (such as shown in Figure 6.2) is controlled by several features: the interfacial tension coefficient between the two polymers, γ, the volume fraction of polymers 1 and 2, v1 and v2, respectively, effective network concentration of the two polymers ν1 and ν2 respectively, and the gas constant times the absolute temperature, RT. Basic equations were derived by Donatelli et al. and Yeo et al. [Donatelli et. al., 1977; Yeo et al., 1983]. Both assumed spheres of polymer 2 dispersed in polymer 1, although the spinodal decomposition model and much electron microscopy suggests that interconnected cylinders may be more prevalent, as discussed above. The Yeo et al., equation is the more general: 4γ RT ( Aν1 + Bν 2 )

D2 =

where A=

1 (3v11/ 3 − 3v14 / 3 − v1 ln v1 ) 2v2

(6.1)

and B=

1 (ln v2 − 3v2 2 / 3 + 3) 2

For lightly crosslinked systems of 50/50 composition, the value of D2 is of the order of several hundred Ångstroms. Even though spheres rather than cylinders were specified in the derivation, the numerical result is surprisingly accurate. 6.3.2

Phase Diagram Control of SIN Morphology

While the morphology of sequential IPNs is mainly controlled by the volume fractions of the two polymers and the thermodynamics of mixing, that of SINs depends primarily on the time order of three key events:

423

(1) The gelation of polymer 1 (2) The gelation of polymer 2 (3) Phase separation of the two polymers. If gelation of one or both polymers precedes phase separation, the polymer gelling first will tend to be more continuous. If phase separation precedes gelation, then a rather coarse morphology may develop. Of course, both gelation and phase separation are controlled by the kinetics of polymerization and the concentration of the crosslinkers, as well as the thermodynamics of mixing. To determine the time order of the three quantities indicated above, a metastable phase diagram was developed by Mishra et al. [1995]. These authors studied polyurethane-based SINs, with poly(methyl methacrylate) constituting the other polymer. The system has proved very popular for research, over 50 papers now being in the literature on PU/PMMA and PU/PS SINs [Mishra and Sperling, 1996]. The morphologies of these systems vary widely, from two well-defined glass transitions to one broad glass transition, or two glass transitions shifted inward toward each other. It appears that the differences in the morphology may have arisen through a different time sequence of the three critical events. (Mathematically, three factorial yields six different possibilities for the finished product.) As shown in Figure 6.3, Mishra et al. [1995] presented a tetrahedron as the spatial form of their metastable phase diagram, the corners of which represented: MMA, PMMA, urethane prepolymer “U”, and PU. The PMMA contained 0.5% tetraethylene glycol dimethacrylate, causing it to gel after about 8% conversion, see the G1-”U”-PU plane. The phase separation curve for the ternary system MMA-PMMA-“U” (front triangle) on polymerization of only the MMA is indicated by the points C-D-A-E. Similarly, the phase separation curve for the MMA-PMMA-PU system (see rear triangle, Figure 6.3), is represented by the points J-B-K-L. Thus, the entire tetrahedron volume is divided into two regions: one phaseseparated and the other single phased, separated by the curvilinear construction C-D-A-E-L-B-J.

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This surface exhibits a characteristic sail-like shape.

sion of the PU. The intersection of the two planes, G1-G2, illustrates the line of simultaneous gelation of the two polymers. Reactions passing to one side or the other of this line will have one polymer or the other gelling first. It must be noted that the line G1-G2 also intersects the line A-B of Figure 6.3, not shown. The intersection of these two curves expresses the presence of a triple critical point, where both polymers simultaneously gel and phase separate. While this triple critical point represents the ideal SIN synthesis condition, it would not, in general, produce the best physical or mechanical properties.

Figure 6.3. The metastable phase diagram for the polyurethane-poly(methyl methacrylate) SIN at 60°C. The PMMA gels at and above the plane G1-“U”-PU. The intersection of the PMMA gelation plane and the surface along the curve A represents the condition of simultaneous phase separation and PMMA gelation.

The inset in Figure 6.3 illustrates the actual experimentally determined curve, C-D-A-E of the tetrahedron. To the left of the diagram, the results were one-phased; to the right, two-phased. The curve down the middle of the triangle indicates where phase separation was first noted. The intersection of the PMMA gelation plane with that of the phase separation sail-like surface, the curvilinear line A-B, represents the critical line along which simultaneous gelation of PMMA and phase separation of the PU from the PMMA exists. Reactions passing to the left of this curve will have the PMMA gel before phase separation, while reactions to the right of A-B will phase separate before gelation takes place. The gelation plane of the polyurethane is illustrated in Figure 6.4, see G2-MMA-PMMA. For completeness, the gelation plane G1 of the PMMA from Figure 6.3 is also shown. The gelation plane G2 occurred after about 67% conver-

Figure 6.4. The metastable phase diagram, same compositions as Figure 6.3, illustrating the PU gelation plane, G2MMA-PMMA.

The tetrahedron construction represents the synthesis of sequential IPNs as well as the SINs. It must be noted that in real SIN syntheses, it is almost impossible to have the kinetics of polymerization of both polymers proceeding at identical rates. Sequential IPN polymerization represents the other extreme, where first one polymer is completely synthesized, and then the other follows sequentially. Both sequential IPN and SIN polymerizations begin somewhere along the edge of the line

425

Interpenetrating Polymer Networks

connecting the two monomers, MMA-“U” in the present case. However, the polymerization lines need not be straight, or even continue as single lines. When the polymerization line intersects the sail-like phase separation surface illustrated in Figure 6.3, the line divides into two portions. Both these lines must conclude somewhere along the line connecting the two polymers, in this case the line PMMA-PU. Since it is impossible to have complete phase separation, the lines cannot be exactly at either apex. However, the more complete the phase separation will be, the closer to the apices the ends of the polymerization will be. Sophiea et al. published the first classical composition-temperature phase diagram, working with the semi-IPN net-polyurethane-inter-poly(vinyl chloride) [Sophiea et al., 1994b]. They found a lower critical solution temperature, LCST ≅ 120°C; below this temperature the system was onephased, and above, two-phased. Such behavior is now known to be characteristic of most polymer blends (see Chapter 2). On the other hand, DuPrez et al., investigating the behavior of a poly(ethylene oxide)-inter-netpoly(methyl methacrylate) semi-II IPN found what appears to be the lower temperature portion of an upper critical solution temperature, UCST [DuPrez et al., 1996]. However, since the temperatures were all below any possible critical temperature, conclusions are difficult to draw. A TTT cure diagram was developed by Kim et al. for the system poly(ether sulfone)-internet-epoxy semi-II IPN [Kim et al., 1993]. They showed that the transformation of the monomers to the polymers involved five steps: onset of phase separation, gelation of the epoxy component, a fixation of the domain size and shape morphology, end of phase separation, and (in this case) vitrification of the epoxy component. In their system, the onset of phase separation preceded gelation. Thermoplastic IPNs are defined above as being hybrid materials between the IPNs and the mechanical blends. Many of these materials consist of polypropylene, a semi-crystalline polymer (the crystallinity contributing to the physical crosslinking of the system), and EPDM. The EPDM is

lightly crosslinked during the mechanical blending such that a dual phase continuity is developed with the EPDM forming cylindrical structures inside of the polypropylene matrix. The basic physical requirements for dual phase continuity in such blends were expressed by Paul and Barlow, and by Jordhamo et al. [ Paul and Barlow, 1980; Jordhamo et al., 1986]:

VI ηII × =X VII ηI where if

(6.2)

X > 1, phase I is continuous X ≅ 1, dual phase continuity or phase inversion is likely X < 1, phase II is continuous.

In Eq. 6.2, the quantity V is the volume fraction, and η represents the melt viscosity of phase I or phase II, depending on the subscript. Equation 6.2 represents the limiting case for zero shear. Utracki has proposed more complete relationships for finite shear stresses [Utracki, 1989]. However, although the forces involved in melt blending are large, the shear rates in blending are considerably smaller than on mixing lower viscosity fluids, such as oil and water. More information on this topic can be found in Chapter 7 Rheology in this Handbook. Davison and Gergen developed thermoplastic IPNs based on SEBS triblock copolymers (where the S hard blocks provide physical crosslinking), and semi-crystalline polymers [Davison and Gergen, 1977; Gergen and Davison, 1978; Gergen et al., 1996]. Figure 6.5 illustrates the dual phase characteristics of a thermoplastic IPN of SEBS and PA-11, the SEBS having been extracted [Gergen et al., 1996]. The important points developed in this section are that sequential IPN synthesis tends to make dual phase continuous materials. For both sequential and simultaneous syntheses, a metastable phase diagram can be developed to study the kinetics of phase separation and gelation, so that better control of the morphology can be attained. The thermoplastic IPNs depend on equal volume and viscosity ratios to attain the dual phase continuity.

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Figure 6.5. An SEM fracture surface of an extracted thermoplastic IPN, 25/75 PA-11/SEBS. The SEBS phase was extracted.

6.3.3

Glass Transition Behavior

Due to the small size of the domains, the glass transitions tend to be either broadened or moved towards one another on the temperature or frequency axis, see Table 6.3. This is a direct consequence of the greater extent of mixing, and/or greater interphase volume fraction encountered in these materials. The glass transition temperatures depend on the crosslinker level, that influences the extent of mixing and/or interpenetration of the chains [Nemirovski et al., 1996]. Of course, the glass transition temperature of any polymer increases slightly with increasing crosslink level. Table 6.3.

In this case, the polymer with the higher glass transition temperature may actually decrease as it is mixed with the other polymer. In such cases, within each phase, the Fox equation, often used for statistical copolymers, may be applied. The volume fraction of the two polymers also plays critical roles, see Figure 6.6 [Hsieh et al., 1996]. This behavior is similar to many polymer blends, where the intensity of the transition depends on the volume fraction. In Figure 6.6, an inward shift may also be observed. The glass transition temperature of three-polymer latex IPNs will be treated in Section 6.4.4.

Dynamic mechanical behavior of IPNs

No.

Observation

Comments

1 2 3 4 5

Two sharp Tg’s near homopolymer values Tg’s moved toward each other One broad Tg, spans both original Tg ranges One sharp Tg One or both Tg’s increased

Little or no mixing Molecular mixing in interphase and/or inside domains Extensive but incomplete mixing Possibly miscible ; Tg following the Fox equation Crosslinking causes the Tg to increase. Can be part of No. 2.

Interpenetrating Polymer Networks

427

Figure 6.6. Loss modulus-temperature behavior of PU(PTMO2000)/UPMI(PEG600) IPNs at various PU/UBMI ratios: ______ 100/0, — — — 90/10, __ .. __ .. __ .. 70/30, …… 0/100.

6.4

Latex-Based Materials

Latex-based polymer materials can be either nature-made, as natural rubber [Stern, 1967; White, 1995], or synthetically made. The synthetically made latexes are commonly based on recipes of monomer, water, surfactant, and free radical initiator to induce chain polymerization [Lovell and El-Aasser, 1997; Wickson, 1993]. However, recipes based on step polymerization are also well known, often resulting in crosslinked films [Walker and Shaffer, 1996]. The resulting latex material consists of small particles, usually spherical, of 50-500 nm in diameter, dispersed in water. Alternately, polymers are sometimes emulsified after polymerization (direct emulsification, the product sometimes called artificial latexes) via agitation of a melt in the presence of water and surfactant (emulsifier), and sometimes organic solvent or plasticizer [Piirma, 1989]. 6.4.1

Definitions

Latexes have long been used in IPNs and related materials. Some of the material types described in the literature are: Latex Blends. These are combinations of two or more kinds of latexes, differing in chemical

makeup. Blends of the same kind of latex, but differing in such characteristics as particle size are of great interest, but will not be considered here [Gilicinski and Hegedus, 1996b]. A major new type of latex blend used in coatings involves hard and soft latexes, serving to increase the modulus of the coating, as well as to decrease block resistance (the ability to resist adhesion between two such films) [ASTM, 1989; Friel, 1992]. Core-Shell Latexes. First a seed latex of polymer 1 is synthesized. Then, a second monomer is added to the system, usually with no added surfactant. Often, a starved polymerization route is employed, i.e., the rate of polymerization equals or exceeds the rate of monomer addition. This reduces the swelling of the seed latex by monomer 2, producing a two-layer latex having a spherical core, and an overlaying shell. Obviously, multiple shells can be added. Latex IPNs. A crosslinked seed latex of polymer 1 is synthesized first. Then, monomer 2, plus crosslinker and initiator are added, usually without new surfactant. If the monomer 2 mix is added either all at once or rapidly, then swelling of polymer 1 by monomer 2 is encouraged, with subsequent greater interpenetration. Latex IENs. Latex interpenetrating elastomer

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networks, latex IENs, are latex blends that have been crosslinked after film formation. They are named after the early works of Frisch and coworkers, who called these materials interpenetrating elastomer networks [Frisch et al., 1969b]. Many latex blends, as used in coatings especially, are crosslinked as finally used in service. Pre-crosslinked Latex Blends. In these materials the individual latexes are crosslinked during synthesis, then blended, and a film is formed. Because of limited deformation and/or interdiffusion capabilities, such films tend to be weak, and only used for special purposes [Zosel and Lay, 1993; Lesko and Sperry, 1997]. However, light crosslinking, as occurs in SBR latexes, may be tolerated. Pre-crosslinked latex blends materials are actually not IPNs, because the definition requires that at least one of the polymers be polymerized and/or crosslinked in the immediate presence of the other. An application of precrosslinked suspension-polymerized blends, in anionic and cationic form, is as ion-exchange resins. In suspensions, the particles are larger, usually of the order of 10-200 µm. 6.4.2

History of Latex Blends and Latex IPNs

The early research on core-shell and latex blend materials has already been mentioned, see Table 6.1. Hopff’s materials were intended to be used vulcanized, in semi-IPN form. A still earlier patent by I. G. Farbenindustrie prepared core/shell latexes from nitrocellulose and poly(ethyl acrylate) [Farbenindustrie, 1931]. Example 6 of the patent shows the level of sophistication already obtained 10 years after Staudinger’s Macromolecular Hypothesis: 20 parts of nitrocellulose are dissolved in 80 parts of acrylic ethyl ester and emulsified in 200 parts of water, which contain 0.3 % of Marseilles soap (saponified from olive oil) and 0.5 % of sodium isopropylnaphthalene sulphonate and which has been adjusted with the aid of N/100 aqueous sulphuric acid to a hydrogen ion concentration corresponding to a pH value of 6.2. 2.5 parts of 30 per cent. aqueous hydrogen peroxide are added to the emulsion. After heating

for 2 hours at from 80 to 90°C, the emulsion is freed from any non-polymerized ester with the aid of steam, and coagulated with the aid of aqueous 5 per cent. hydrochloric acid. The product is readily soluble in organic solvents and yields very elastic films. While no crosslinker is mentioned in the patent, poly(ethyl acrylate) usually gels. That it was “readily soluble” suggests a low molecular weight. These early patents set the stage for today’s latex blends and IPNs. 6.4.3

Types of Crosslinks

Klein and Daniels point out three basic categories of crosslinking in latexes [Klein and Daniels, 1997]: homogeneous (intraparticle crosslinking), interfacial crosslinking, and interstitial crosslinking. If two different polymer latexes are involved, the interfacial crosslinking actually constitutes a type of grafting. In interstitially crosslinked systems, crosslinking occurs in the aqueous phase, where water-soluble or water swellable polymers undergo crosslinking (or grafting) reactions. These, for example, may involve melamine-formaldehyde resins. The principal types of crosslinking (or grafting) systems include: carboxyl groups, hydroxyl groups, chloride (or other halogen) moieties, dimethyl meta-isopropyl benzyl isocyanate, TMI® or related materials, derivatives of acrylamide (such as N-methyl ethers), epoxy groups, and sulfonate groups. Ureido-functional monomers can also be used [Merger et al., 1988]. Hydrogen and ionic bonding constitute effective physical crosslinking and or grafting methods. A widely used self-condensable monomer is N-methylol acrylamide. This monomer undergoes a self-condensation reaction as well as reactions with hydroxyl and carboxyl moieties. The methylol functional groups can also react with epoxides or other methylol or methylol ether groups. Typically, N-methylol acrylamide is employed at concentrations of 3-7 wt% in latex coatings formulations [Daniels and Klein, 1991].

Interpenetrating Polymer Networks

6.4.4

Strategies for Low-VOC Latex Coatings

Feng et al. discussed some of the possibilities for producing low volatile organic compound (VOC) latex coatings, addressing one of the more important coatings technological problems of modern times [Feng et al., 1995]. They pointed out that low Tg latex materials produce low Tg films, tacky to touch. The ability to resist adhesion (between two such films) is called block resistance in the coatings industry [ASTM, 1989]. Historically, people added organic solvents, transient plasticizers, etc., promoting particle deformation and interdiffusion, which on evaporation increased block resistance. Reactive plasticizers that polymerize after film formation, can also be used. They form an independent network (sometimes topologically resembling bailing wire) or crosslink the polymer. Thus, the reaction results in a nonvolatile IPN. The chemistry of the latex surface region can be arranged such that it can be plasticized with water. This may be satisfactory for situations where the surface chemistry will be altered later, or that water will not reach the final film in deleterious quantities. A newer alternative relates to the blending of hard and soft latex particles. The soft particles flow to form a film, while the hard ones tend to raise the modulus, and, if on the surface, increase scratch resistance, etc. Another alternate involves the post-drying of the film via crosslinking. Any number of reactions can be used, including double bonds pendent to the chain, or condensable groups. Heat, radiation, or high temperature peroxides might be used. If two different latexes, or core-shell structures are present, then the product will be an IPN. 6.4.5

Selected Patents in Latex Blends and Crosslinked Systems

The field of latex blends and their crosslinked counterparts has been especially active, see Tables 6.4 and 6.5. Even in Table 6.4, where no crosslinkers are specified in the claims, it is likely that in many instances the final product will

429

contain some crosslinking. No. 4 in Table 6.4, for example, mentions a gelling agent. All the materials in Table 6.5 have a crosslinker in either one or both of the latexes. Both the N-methylol acrylamides and the ureido-functional groups described above appear in Table 6.5. The reader should note the use of both large and small latex pairs, as well as hard and soft latex pairs, in terms of making useful materials. As defined in Section 6.1 and above, the compositions described in Table 6.5 are all latex IENs, where two latexes are blended, then crosslinked. Latex IPNs are described in Section 6.4.7. Patents rarely include the description of the system morphology. Patel et al. [Patel et al., 1996] examined blends of hard and soft particles via atomic force microscopy, finding that when the soft component is present in amounts larger than 40%, smoothed bumps were observed that appeared larger than either the hard or soft particles alone. The smoothness of each bump, supported by other evidence, suggests that the soft particles have coalesced into a virtual continuum at the surface. The overall surface unevenness was thought to be indicative of underlying hard particles. The hard particles were thought to be submersed because of the lower surface energy of the soft polymer. Gilicinski and Hegedus [Gilicinski and Hegedus, 1996a] also used atomic force microscopy to identify hard and soft domains by supplementing topographic data with mechanical property maps. A third such study by Geurts et al. examined the effect of blending two sizes of monodisperse poly(butyl methacrylate) latexes via minimum film formation techniques [Geurts et al., 1996]. Usually, the addition of increasing amounts of small particles beyond a certain concentration level leads to a disruption of the particle matrix until the point at which the matrix changes from a continuous matrix of large particles into a continuous matrix of small particles. While these model materials were not crosslinked, it is easy to imagine the effects of crosslinking such materials, as indicated in Table 6.5.

SBR SBR PVC (small diameter) ABS Acrylic binder (soft) Aromatic vinyl monomer based Acrylic acid based Vinylidene chloride-butadiene copolymer Rubber grafted with vinyl cyanides

3 4 5

Poly(vinyl acetate) PB

Poly(vinyl chloride)

PI

Vinyl esters

Acrylic ester

SBR

Poly(vinyl halide)

Acrylic

Acrylic (soft)

1

2

3

4

5

6

7

8

Acrylic (hard)

Heterocycles

Acrylic

PVC

PVC

Resorcinol

Polymer 2

Polymer 3: melamine formaldehyde; N-methylol acrylamide N-methylol acrylamide in SBR-based latex N-methylol acrylamide in acrylic Ureido-functional monomers Ureido-functional monomers

N-methylol acrylamide

Hydroxyl-methyl diacetone acrylamide Vulcanizing rubber compound

Paint dispersions and leather assists Polymeric binder for aqueous coatings

Dielectric sealing

Heat sealable PVC films

Water-based coating

Woodworking adhesives

Wet primed paint adhesion Golfball thread rubber

Nakajima et al.

Friel

Merger et al.

Yannich and Katz

Wietsma and Stam

Leeson and Ludwig

Mudge et al.

Hamada

U.S. Pat., 5,166,269, 1992 U.S. Pat., 4,684,689, 1989 U.S. Pat., 4,777,265, 1988 Eur. Pat. Appl., 0,466,409 A1, 1992

Rohm and Haas

National Starch BASF

BASF

National Starch Goodrich

Sumitomo

Bordon

Company

Toray Ind.

3M Im. Adad. Lebedeva

Nippon Zeon

U.S. Rubber Co. Congoleum

Goodyear Polysar Goodrich

Xerox Standard Oil

Company

U.S. Pat., 3,935,151, 1976 U.S. Pat., 5,340,112, 1994 U.S. Pat., 5,434,216, 1995 U.S. Pat., 4,007,147, 1977

Patent No.

JP 08169999, 1994

EP 741005, 1995 RU 2055843, 1992

JP 08269274, 1995

U.S. Pat. 3,624,183, 1971 U.S. Pat. 4,609,431, 1986

U.S. Pat. 4,271213, 1981 U.S. Pat. 4,55,591, 1985 U.S. Pat. 4,461,869, 1984

U.S. Pat. 5,496,676, 1995 U.S. Pat. 4,510,204, 1985

Pat. No.

Nickerson et al.

Inventor

Heiner and Coates Kalinina et al.

Mori et al.

Leach and Murray Grose and Carlson

Grimm and Gunnerson Peltier et al. Yang

Croucher et al. Duke et al.

Inventor

Application

Impact resistant thermoplastics

Adhesive resistant to fuels Protection of wood

Adhesives for palm fibers, bedding

High surface gloss Dimensionally stable backings

Partitioning agent for crumb rubber Foam rubber backings Bimodal vinyl dispersion resins

Toner aggregation processes Corrosion resistant coatings

Application

Comments/Crosslinker

Selected patents in latex IEN applications

No. Polymer 1

Table 6.5.

11

9 10

8

ABS

SBR Olefinically unsaturated mers PMMA NBR Core-shell acrylic (large diameter) SAN Acrylic binder (hard) Conjugated diene based Urethane based PVC

Styrene/acrylic Nitrile based

1 2

6 7

Polymer II

Selected patents in latex blend applications

No. Polymer I

Table 6.4.

430 L. H. Sperling and R. Hu

Interpenetrating Polymer Networks

6.4.6

Core/Shell latexes

Core/shell latexes refer to systems with a submicroscopic particle morphology of one polymer forming the center part (the core) and the other polymer covering the core (the shell layer). Core/ shell latexes are made via two consecutive emulsion polymerization stages, usually forming a particle structure with the initially polymerized material at the center and the later-formed polymer as the outer layer. If more than two stages are employed in the emulsion polymerization process, latex particles with multilayered morphology can be obtained. The morphology of latex particles is controlled by the thermodynamic and kinetic factors. The thermodynamic factors determine the ultimate stability of the multiphase system, inherent in the production of a composite latex particle, while the kinetic factors determine the ease with which such a thermodynamically favored state can be achieved. The parameters affecting the thermodynamics of the system include the particle surface polarity, the relative phase volumes, and the core particle size. The parameters affecting the kinetics of the morphological development include the mode of monomer addition (monomer starved or batch) and the use of crosslinking agents. Of course, crosslinked core/shell latexes constitute IPNs, see Section 6.4.1. The core/shell morphology is usually more favorable if the latex contains a hydrophobic/ hydrophilic polymer pair polymerized in that order, the volume fraction of the second stage polymer is higher, the core particle size is larger (in the range from 100 to 900 nm), semi-continuous polymerization process (monomer starvedfeeding mode) for polymer 2 is employed, and crosslinking agents in both core and shell phases are used [Lee and Ishikawa, 1983; Cho and Lee, 1986, Berg et al., 1986; Lee and Rudin, 1989; Jonsson et al., 1991; Xu et al., 1991; Lee, 1981; Min et al., 1983; Hourston et al., 1986a; Lee and Rudin, 1992; Sundberg et al., 1990; Winzor and Sunberg, 1992; Durant and Sundberg, 1995; Kong et al., 1996].

431

To a large degree, the physical and mechanical properties of the final product depend on the particle morphology. By properly selecting polymer pairs for the core and the shell stage, latex particles can be designed to have unique properties for a wide range of potential applications, from paints, coatings, adhesives, organic opacifiers, to impact modifiers for plastics or carriers for biomolecules [Kowalski et al., 1984; Dong et al., 1995; Qian et al., 1995; Segall et al., 1995; Hu et al., 1995; Hu et al., 1997c; Lindemann and Deacon, 1993; Lee et al., 1986] 6.4.7

Latex IPNs (LIPNs)

LIPNs are also prepared through a two stage emulsion polymerization [Kim and Sperling, 1997; Sperling et al., 1972; Sperling et al., 1973; Hourston and Satgurunathan, 1984]. First, a crosslinked seed latex is synthesized. For the second stage, one variation consists of swelling crosslinked latex particles (as the seed) with the second monomer, crosslinker, and initiator/catalyst, followed by polymerization. Another method is to feed the second monomer (together with a crosslinking agent and initiator/catalyst) slowly in a monomer starved condition during the polymerization [Sperling, 1997]. The morphology of LIPNs is controlled by many factors, as will be discussed in Sections 6.4.7.2 and 6.4.5. The LIPNs are a special type of IPNs that combine both networks in a single latex particle, which differs from bulk IPNs, because the thermosetting character is limited to the size of the latex particles (of the order of several tens to a few hundreds nanometers). Such IPNs have unique features of morphology, as well as in physical and mechanical properties [Zhong and Zhu, 1991; Shu et al., 1990b; Zhang et al., 1991]. 6.4.7.1 Crosslinking and Grafting For an ideal IPN, the two or more polymer networks are at least partially interlaced on a molecular scale but not covalently bonded to each other and cannot be separated unless chemical bonds are broken [Mita and Akiyama, 1997].

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The crosslinking of each polymer network is simply controlled by the type and concentration of the crosslinker employed in the system. The higher the crosslinker level (especially in polymer network I), the finer the morphology. For a real IPN, however, especially for LIPNs, crosslinking, or grafting of the polymer 1 with polymer 2 is very likely to occur, based on the following three reasons: 1) The LIPNs are typically prepared through a two-stage emulsion polymerization process by a free radical polymerization mechanism. The crosslinking agents usually have two or more double bonds. Depending on the polymerization conditions, part of the double bonds from crosslinker I could be left during the first stage polymerization, subject during the second stage polymerization to act as graft sites; 2) If polymer 1 has residual double bonds, such as polybutadiene, part of the residual double bonds may be reacted in the second stage free radical polymerization process, forming grafting sites for polymer 2; and 3) If polymer 1 has an active α-hydrogen in the molecular structure, such as poly(butyl acrylate), hydrogen-abstraction easily occurs in the second stage polymerization. Actually, in many cases, to improve the compatibility between two networks, a graft site is intentionally added in polymer network I to enhance the grafting of polymer 2 on polymer 1. The grafting reaction between two polymers can usually be minimized if the two polymer networks are formed through different polymerization mechanisms, such as polycondensation and free radical polymerization. In this case, a grafting agent may be needed in polymer 1 to improve the compatibility between the two networks. Again, the higher the graft site level, the higher the grafting efficiency of polymer 2 on polymer 1, the better the compatibility of the two phases. With very high graft levels, an AB-crosslinked copolymer may be created [Sperling, 1997]. However, for most purposes, the graft site concentration should be low enough to prevent miscibility of the two polymers.

6.4.7.2 Morphology The morphology of latex IPNs refers to the degree of phase separation, the number of the phases, the shape and size of the domains, the degree of phase continuity, and the structure of the interface. On film formation, the ensemble of latex particles presents further unique morphologies. For LIPNs, the morphology is complicated because the swelling, polymerization, crosslinking, and phase separation are all taking place within the submicron polymer particles in an aqueous phase. Different LIPNs have different morphologies. Even for LIPNs with the same chemical composition, the morphology can be different when different synthesis process are used [Peng and Li, 1995]. Since an LIPN is typically prepared by sequential emulsion polymerization, a core/shell morphology often forms, similar to that of the latex particles synthesized by a two (or more) stage emulsion polymerization process [Sperling and Sarge, 1972]. However, if the crosslinking density of the first prepared polymer and the monomer 2 concentration are controlled properly, a cellular type morphology can also be obtained [Silverstein et al., 1989; Hu et al., 1995; Sionakidis et al., 1979]. Similarly as in the core/shell latex particles, the LIPNs morphology is controlled by miscibility of the polymers, the volume fraction of each polymer, the crosslinking density of the polymers, and the sequence or the order of the synthesis (which polymer is synthesized first) [Sionakidis et al., 1979; Hourston and Satgurunathan, 1987]. The effects of those factors on the LIPN morphology are summarized as follows: (1) The more miscible the two polymer compositions, the less the phase separation between them, and/or the smaller the domain size for each of them. The miscibility is partially indicated by the solubility parameters (δ) of the two polymers. The closer the two δ’s, the more miscible the system [Yin et al., 1991]; (2) Effect of the volume fraction of each phase in LIPN morphology depends on the nature of the two compositions. Phase separation

Interpenetrating Polymer Networks

can be reduced with increasing the content of the first component [Krause, 1972], or with increasing the volume fraction of the second component [Geng, 1992]; i.e., by staying away from 50/50 compositions. If polymer 1 is hydrophobic (synthesized first) and polymer 2 is hydrophilic (synthesized later), a normal core/shell morphology is more easily formed; vice versa, an inverted core/shell morphology may result [Narkis et al., 1985; Narkis et al., 1986]. In addition, the longer the swelling time for the monomer 2 in polymer 1, the more nearly thermodynamic equilibrium is achieved for the system [Hourston et al., 1987]. (3) Increasing the crosslinking density of polymer 1 (the seed) reduces the phase size [Sheu et al., 1990a]. However, stronger phase separation can occur when the crosslinking density of polymer 1 is high enough to reduce the diffusion of the monomer 2 into polymer 1. The higher the crosslinking density of polymer 2, the more efficient for the molecular interlocking, the less phase separation results [Sheu et al., 1990; Yin et al., 1988]. The morphology of LIPNs can be investigated optically by either TEM or SEM and mechanically via the glass transition behavior. Nemirovski et al. studied elastomeric LIPNs by a two-stage emulsion polymerization of styrene [with a comonomer to lightly crosslink the PS] upon a lightly crosslinked polyacrylate (PAcr) seed latex (75:25 PAcr:PS) [Nemirovski et al., 1996]. The influence of the sequence and composition upon the morphology is schematically illustrated in Figure 6.7 [Silverstein and Narkis, 1989]. They proposed a schematic diagram of the molded LIPN showing both the inter-particle and the intra-particle microdomains, and the bicontinuous IPN structure inside the particles, see Figure 6.8 [Silverstein et al., 1989]. Wu and Zhao studied LIPN systems by a two stage emulsion polymerization technique [Wu and Zhao, 1995]. A latex seed (polymer 1) was synthesized first by a semi-continuous emulsion polymerization process, swollen by the second

433

Figure 6.7. Schematic crossection of LIPNs at 100%, 75%, 50%, and 25% seed polymer: (a) hydrophobic seed; (b) hydrophilic crosslinked seed, both polymers crosslinked; (c) hydrophilic seed, second stage crosslinked.

Figure 6.8. Schematic crossection of molded LIPN showing intraparticle and interparticle PS nanodomains and an interwoven network structure.

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stage monomer or monomer mixture (forming polymer 2), followed by polymerization to form IPN materials. Six kinds of monomers were used: acrylonitrile (AN), vinyl acetate (VAc), n-butyl acrylate (nBA), methyl methacrylate (MMA), ethyl methacrylate (EMA), and ethyl acrylate (EA). The effect of composition, crosslinking level, feeding sequence of polymer 1 and polymer 2 on the IPN miscibility and mechanical behavior were investigated via TEM, equilibrium swelling ratio, density, and DMS measurements. They found that the latex IPN particles had a core/shell type morphology. Introduction of a third component as a comonomer (which had good miscibility with the original two polymers) improved the IPN miscibility. Their IPNs with higher crosslinking density had a higher phase miscibility. They also found that the IPNs with a hard core and a soft shell morphology, called an inverted IPN, had a higher miscibility than the IPNs with the opposite morphology, called normal IPNs (soft core and hard shell), as shown in Figure 6.9.

Figure 6.9. Dynamic mechanical spectroscopy curve of the LIPN with different synthetic orders: C-1 = P(AN-co-EMA) / PnBA and C-5 = PnBA/PAN-co-EMA).

Another kind of LIPN material, P(MMA-stat-S) / P(BA-stat-AA), with a core/shell structure, was synthesized by Chen and Wang [Chen and Wang, 1995a; Chen and Wang, 1995b]. P(MMA-stat-S) seed was synthesized first by batch emulsion polymerization, and then PBA-PAA was semicontinuously polymerized. The effects of emulsifier concentration, crosslinking agent, and initiator on the final latex films were investigated. They found that the core was a hard MMA-S statistical copolymer, while PBA located in the inner layer of the shell and the PAA located in the outer layer of the shell (due to a poor miscibility of PAA with P(MMA-stat-S) and the hydrogen bonding between PAA and water). An interpenetrated structure was formed at the interface of the core and the shell. These 0.1-0.2 µm latex particles were initially dispersed in the aqueous phase. After the evaporation of water, these latex particles formed a film, the particles coalescing with the soft shell becoming the continuous phase sticking these particles together. This kind of IPN can be used in coatings and adhesives. It was also shown that on increasing the core/shell ratio (hard to soft monomer weight ratio) from 2/8 to 8/2, the state of the film changed from an adhesive to one cracked into a powder. At the 1/1 core/shell ratio, increasing the AA concentration to 2.5%, the film was still uniform and clear; but at higher AA levels, the latex stability decreases in the presence of electrolyte (due to a thinner electronic double layer) and higher water absorbance (due to the hydrophilicity of the -COOH). From DSC, the core/shell LIPN latex particle showed two broad Tg’s at -36 and 103°C and one small Tg at 10°C. Besides two-component LIPNs, three-component LIPNs have also been studied through three stage emulsion polymerization processes [Zhang et al., 1991; Isao et al., 1992; Zhang et al., 1994]. These authors synthesized poly(n-butyl acrylate) crosslinked with ethylene glycol dimethyacrylate as the seed latex. Styrene and divinyl benzene were added at the second stage. The third stage was linear poly(methyl methacrylate). Starved polymerization conditions resulted in more regular shaped latex particles than batch addition of monomer.

Interpenetrating Polymer Networks

Another interesting three-component LIPN consists of crosslinked polyorganosiloxane (tetraethoxysilane as crosslinker) as the first stage, swollen by butyl acrylate monomer (allyl methacrylate as crosslinker) and then polymerized to form the second component, and finally, poly(methyl methacrylate) was grafted onto the IPN core as the shell layer [Isao et al., 1992]. Such kind of LIPNs are found to be a good impact modifiers in thermoplastics; i.e., they are toughened by the addition of finely divided low Tg phase. 6.4.7.3 Damping Polymers with sound and vibration damping properties are finding numerous applications, especially in the aircraft, automotive, tall buildings, and appliance industries [Aklonis and MacKnight, 1983; Thurman and Miller, 1986; Corsaro and Sperling, 1990; Sophiea et al., 1994a]. The damping properties of polymers are dominated by their glass transitions. Usually, homopolymers or statistical copolymers possess efficient damping in a temperature range of only 20-30°C around the Tg [Aklonis and MacKnight, 1983]. For most outdoor or machinery applications, however, good damping materials should exhibit a high loss factor (tan δ > 0.3) over a temperature range of at least 60-80°C [Yak, 1994; Yao et al., 1991]. Why use IPNs as damping materials? Because the introduction of crosslinks in IPN’s restricts the domain size to the order of 10-20 nm, enhancing the formation of a micro- heterogeneous morphology. This results in broad glass transitions, making them effective as broad temperature or frequency damping materials [Wang et al., 1996; Fay et al., 1991; Sperling and Fay, 1991]. The damping behavior of LIPNs can be improved by proper design of the chemical composition in the system and effective control of the IPN morphology. The same factors should be considered as controlling core/shell LIPN morphology, such as the miscibility of the system, chemical structure of the polymer molecule, glass transition temperatures, crosslinking density of the two networks, and the feeding sequence in the polymerization. But here, to develop a LIPN with broad tem-

435

perature damping behavior, a more miscible, or microheterogeneous morphology is needed instead of a core/shell type morphology. By contrast, if only a narrow temperature range is needed, such as an indoor, room temperature application, a homopolymer or simple statistical copolymer, with a sharp Tg will serve better. An example would be slightly plasticized poly(vinyl acetate), Tg = 25°C. Based on damping theory [Kim and Sperling, 1997; Sophiea et al., 1994a], the damping behavior of a polymer can be evaluated from its dynamic mechanical behavior by expressing it as the area under the tan δ versus temperature curve [Keskkula et al., 1971] or the area under the linear loss modulus versus temperature curve [Fradkin et al., 1986a; Fradkin et al, 1986b; Chang et al., 1987]. The broad temperature damping performance improves if the polymer system is more miscible, the glass transition peak is broader, and the tan δ or loss modulus peak intensity is higher. The damping performance can also be predicted and improved via group contribution analysis [Fradkin et al., 1986b; Chang et al., 1987; Chang et al., 1988; Fay et al., 1990; Foster et al., 1987]. However, the group contribution analysis may be unable to predict the partially miscible two-peaked (microheterogeneous) IPNs [Lipatov, 1989], because the analysis is based only on molecular structure, not considering the interaction between the polymer phases. Sperling et al. first developed a prototype “Silent Paint”, which is a two-layer coating system, PEMA/PnBA LIPN, capable of attenuating noise and vibration over - 30°C to + 100°C [Sperling et al., 1974]. Later, they investigated P(EMAEA) / P(BA-EA), PVC/P(B-AN) LIPNs for more effective damping applications [Grates et al., 1975; Lorenz et al., 1976; Sperling et al., 1975]. The damping capability was controlled by varying the chemical composition of the LIPN systems. Liu et al. also studied LIPN systems for damping control in coating applications [Liu et al., 1995]. A polystyrene (PS)/polyacrylate (PAcr) latex IPN was synthesized in a two-stage emulsion polymerization. Crosslinked PS was synthesized first as the seed polymer by a semi-continuous process,

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followed by a nBA-MA-BMA monomer mix with crosslinking agent, again feeding semi-continuously into the PS seed to polymerize, forming an IPN latex. The effect of PS and PAcr ratio and the addition of inorganic fillers, such as, mica, graphite, CaCO3, and TiO2 on the damping of the final coating film were studied. They found that the damping behavior of PS/PAcr IPN at room temperature was improved with increasing PS amount in the system, as shown in Table 6.6. This is because the presence of the benzene ring (from PS) enhances the friction of molecular chain movement. The introduction of a monomer with polar group to the PAcr backbone improved the miscibility of the IPN system and increased the adhesion of the coating film. This is because the polarity can increase the interchain attraction, favoring the permeation of the molecular chains, enhancing the mutual entanglement between the networks, and increasing the friction between the polymer molecular chains. Consequently, the damping property was improved. The addition of inorganic fillers increased the damping effective value, as well as expanded the temperature range for effective damping, with flake structure fillers, e.g., mica, showing the best result. This was confirmed by the DMS measurement, with the broadest tan δ peak and the highest valley value obtained. Table 6.6.

Effect of PS/PAcr ratio on damping

PS/PA 25/75 Noise Reduce, dB/mm 3.0

35/65 3.5

45/55 3.9

50/50 5.3

Peng and Li synthesized a series of PS/P(EAstat-nBA) LIPNs by two stage emulsion polymerization method, to investigate the effect on the damping properties of: composition, crosslinker content, chain transfer agent, polar group, and feeding sequence [Peng and Li, 1995]. The results show that the PS/P(EA-stat-nBA) LIPNs are a partially miscible system. The effect of crosslinker content of network I on the damping properties is not the same as that of network II.

The damping properties of the PS/P(EA-statnBA) LIPN (hard component synthesized first) are superior to that of P(EA-stat-nBA)/PS LIPN (soft component synthesized first). The introduction of polar groups into the system can increase the damping properties by increasing the tan δ and broadening the damping range; while the addition of a chain transfer agent can broaden the damping temperature range. 6.4.7.4 New Development of LIPNs Recently, a new type of three component latex with IPN cores as impact and damping improvers was prepared by three-stage emulsion polymerization [Hu et al., 1995; Hu et al., 1997a]. The IPN cores were composed of one impact part (polybutadiene based, with Tg around -75°C) and one damping part (acrylic based, with Tg around +10°C). The synthesis procedure was as follows: First, component A, a seed latex of poly(butadienestat-styrene), P(Bd/S) 90/10, was prepared at both 95% and 40% gel fractions. After stripping the residual monomer in the seed latex, a mixture of acrylate monomer (butyl acrylate or ethylhexyl methacrylate) and styrene or methyl methacrylate monomer together with tetraethylene glycol dimethacrylate and initiator was added and swollen for 24 hrs at room temperature, then polymerized to form an IPN hybrid core, A+B. Finally, the third component, C, poly(styrene-statacrylonitrile) (72/28) copolymer, was synthesized under monomer starved conditions to form a uniform glassy shell layer covering on the IPN cores. This last made the latex particles compatible with a polycarbonate matrix [Tanrattanakul et al., 1996]. A core/shell type phase separation was observed for IPN particles when using P(Bd/S) with 95% gel fraction, while a cellular IPN morphology with a distribution of polyacrylate based copolymer in P(Bd/S) phase was formed with 40% gel fraction P(Bd/S). Also, the domain size of the second polymer was larger when increasing the volume fraction of the second polymer. In order to study the effect of interfacial interaction between the two polymer components in an IPN particle (polymer A and polymer B),

Interpenetrating Polymer Networks

437

separate core/shell particles and multilayered structured particles were prepared, as shown in Figure 6.10 [Hu et al., 1997b].

Figure 6.11. DMS results for IPN core/shell latex film: (0.5 + 0.5)/1 (A + B)/C, A = P(Bd/S), B = P(EHMA/S), and C = SAN.

Figure 6.10. Models of different types of core-shell modifiers: (a) traditional core/shell; (b) IPN core/shell; (c) blend of separate core/shell; and (d) multilayered core/shell. Possible polymer compositions are: polybutadiene copolymer as polymer A, acrylic copolymer as polymer B, and SAN as polymer C.

Dynamic mechanical spectroscopy (DMS) was used to characterize the glass transitions of the above LIPN particles. It was found that poly(butyl acrylate) based copolymer provides higher miscibility with P(Bd/S) than poly(ethylhexyl methacrylate) based copolymer, which has a longer side chain in the molecular structure. As a comparison to the IPN core/shell system, a blend of separate core/shell latex particles, A/C and B/C latexes, was also synthesized and evaluated. Figure 6.11 [Hu et al., 1997a] shows three glass transitions for the IPN core/shell latex particles. Figure 6.12 [Hu et al., 1997a] shows a better developed middle glass transition for component B, and a lower temperature P(Bd/S) glass transition. The polymer B appears more miscible with polymer A in the IPN form. In the same work, the damping behavior of those three component LIPNs was evaluated from the integrals of the linear loss shear modulus vs. temperature (loss area, LA) and linear tan δ vs. temperature (tan δ area, TA) curves measured by

Figure 6.12. Dynamic mechanical spectroscopy results for blend of separate core/shell latex film: 1/1 A/C and 1/1 B/C at a 1/1 blend ratio: A = P(Bd/S), B = P(EHMA/S), and C = SAN. More distinctive glass transitions are shown here, particularly for polymer B.

DMS [Hu et al., 1997b]. Baselines were used for subtraction, as is common in spectroscopy calculations [Fay et al., 1991]. The IPN core/shell particles showed the highest damping, comparing with separate and multilayered core/shell particles. Also, normal synthesized (rubbery polymer was synthesized first) and inverted synthesized

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(glassy polymer was synthesized first) semi-I IPN and semi-II IPN latexes were compared for their damping performance. The inverted synthesized ones provided much higher LA and TA values, due to their finer morphology. Table 6.7 and Table 6.8 summarizes selected LIPN systems developed for two- and three-stage systems, respectively. For the latter, polymer 3 is usually a linear overcoat. Of great interest throughout are morphology and concomitant behavior. Applications for coatings, damping, and impact improvement are discussed in many of the works.

6.5

Actual and Proposed Applications of IPNs

While the field of IPNs is relatively young compared with this of polymer blends, the number of applications are growing rapidly. In 1979, the total number of scientific papers was about 125; there were about 75 patents. Today, about that many are published each year. Major advances include thermoplastic IPNs, renewable resource materials, biomedical materials, and non-linear optical materials.

Table 6.7.

Newly developed two-stage LIPN systems

Polymer 1

Polymer 2

Property investigated

References

Polyacrylate (PAcr) Polystyrene (PS) PAcr PS Acrylic Natural rubber PAcr PAcr PAcr PS PAcr

Poly(n-butyl acrylate) (PnBA) Polyacrylate (PAcr) MAA, PMMA, P(MMA/S) P(EA-nBA) Acrylic PBA PS or PMMA PS PAcr PAcr P(S/MMA) copolymer

Damping behavior Damping coating Mechanical behavior Damping Coating behavior Impact modifier in PC Film formation Structure and properties Powder paint Damping Film formation

Wu and Zhao 1996 Wang et al., 1996 Nagarajan et al.,1996 a-d Li et al., 1995 Rearick et al., 1996 Schneider et al., 1996a Nagarajan et al., 1996e Nemirovski et al., 1996 Mizoguchi et al., 1996 Liu et al., 1997 Nagarajan et al., 1997

Table 6.8.

Newly developed three-stage LIPN systems

Polymer 1

Polymer 2

Polymer 3

Property investigated

References

PBd based copolymer

PBA based copolymer

SAN

Morphology and damping

Natural rubber PBA PU

PS PS PBMA

PMMA PMMA PS

Morphology Impact modifier Miscibility, damping

Hu et al., 1995; Hu et al., 1997a; Tanrattanakul et al., 1996; Hu et al., 1997b Schneider et al., 1996 Zhang et al., 1996 Hourston et al., 1996

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Interpenetrating Polymer Networks

6.5.1

Thermoplastic IPNs

The thermoplastic IPNs utilize physical crosslinks, rather than chemical crosslinks. Usually, these materials will flow when heated to sufficiently high temperature (hence the terminology thermoplastic), but behave as thermosets at ambient temperature, with IPN properties, often possessing dual phase continuity. Most often, physical crosslinks are based on triblock copolymers (thermoplastic elastomers being the leading material), ionomers, or semi-crystalline materials. Sometimes one of the components is chemically crosslinked. Usually, such materials undergo dynamic vulcanization, i.e., chemical crosslinking during melt shearing. The resulting action yields a dual-phase continuity product, where the chemically crosslinked component forms cyl-

Table 6.9.

inders within the matrix of the other component. Coran and Patel [1996] list several examples of the improvements possible: a. Reduced permanent set b. Improved ultimate mechanical behavior c. Improved fatigue resistance d. Improved fluid resistance, i.e., to hot oils e. Improved high-temperature stability f. Greater stability of phase morphology in the melt (i.e., the phase domain size does not increase) g. Greater melt strength (i.e., green strength) h. More reliable thermoplastic fabricability. Some materials that have achieved importance include EPDM/PP, NBR/PA, PU/PA, SEBS/PA, EPDM/PBT, and Epichlorohydrin rubber/PA. Usually, these materials are made more processable through plasticizing oils. Holden [1996]

Selected plastic-elastomer thermoplastic IPNs by trade name#

Trade Name (Manufacturer)

Synthetic Method

Polymer

Elastomer

Notes

Ren-FlexTM (Dexter); HifaxR (Himont); PolytropeR (Schulman); TelcarR (Teknor Apex); FerroflexR (Ferro); FlexotheneR (Quantum)

Blend

Polypropylene

EPDM or EPR

Relatively hard, low density, not highly filled

SantopreneR (Advanced Elastomer Systems); SarlinkR (Novacor, DSM); UnipreneR (Teknor Apex); HifaxR (Himont) TrefsinR (Advanced Elastomer Systems); Sarlink 2000R (Novacor, DSM) VyramR (Advanced Elastomer Systems) GeolastR (Advanced Elastomer Systems) AlcrynR (DuPont);

Dynamic vulcanizate*

Polypropylene

EPDM

Dynamic vulcanizate

Polypropylene

Butyl rubber Natural rubber Nitrile rubber Ethylene reactive blend Nitrile rubber

Better oil resistance, low compression set, softer Low permeability, high damping Low cost Oil resistant Single phase, soft, oil-resistant Oil-resistant

Silicone rubber SEBS

Medical applications Automotive

SarlinkR 1000 (Novacor, DSM) ChemigumR (Goodyear); ApexR N (Teknor Apex); ElastarR (Nippon Zeon) RimplastR (Petrarch Systems) Kraton IPNR (Shell)

Blend Dynamic vulcanizate Blend Blend Blend

Chlorinated polyolefin Poly(vinyl chloride)

PA Polyester

Notes; *Dynamic vulcanizate: usually, the EPDM is crosslinked via free radical methods during a shearing action of the blend, often accompanied by a partial phase inversion to dual phase continuity; #Based on G. Holden, Ch. 16, Table 16.10, in G. Holden, N. R. Legge, R. Quirk, and H. E. Schroeder, Eds., Thermoplastic Elastomers, 2nd Ed., Hanser, Munich, 1996; AES = Advanced Elastomer Systems

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provided trade names of selected compositions, emphasizing those based on hard polymer-elastomer combinations, see Table 6.9.

6.5.4

Non-linear Optical Materials

Renewable resources are those agricultural products that can be grown repeatedly. A closely related term is natural products. However, the latter may not be renewable, i.e., coal or petroleum oil. Renewable resource materials have long been based on cellulose, triglyceride oils, and natural rubber, among other materials. Some IPNs based on renewable resources are summarized in Table 6.10. These materials suffer from two drawbacks in the current era: research on them is passé and rather poorly funded. Secondly, the price of many renewable resources, such as castor oil and cellulose, is actually above the corresponding prices for petroleum based products.

Non-linear optics is concerned with the manipulation of electromagnetic energy, particularly into different wavelengths than the original. For second-order nonlinear optics, an asymmetric molecule is required. While such molecules clearly do not need to be polymeric, the use of asymmetric polymer chains and side chains produces materials with spatial and temporal stability. Such asymmetric groups have the capability of holding one photon until another arrives, and then the energies may be added, with a single photon of twice the frequency produced, or other algebraic manipulations may be undertaken. As may be expected, a wider range of mechanical behavior patterns arise through the application of IPN technology, see Table 6.12. A leader in the field is Tripathy, with a new patent [Tripathy et al., 1996] in non-linear optics applications for IPNs.

6.5.3

6.5.5

6.5.2

Renewable Resource IPNs

Biomedical IPNs

It has been said that people working on new applications are more willing to consider new materials than people working on established applications. This is certainly a truism when considering the biomedical applications of polymeric materials, see Table 6.11. The emphasis on soft contact lenses is remarkable. The first item, showing a gradient IPN being used to form a gradient refractive index material is the most interesting. Of course, there are two ways to bend light either to have a curved surface, the basis of most lenses, or to have a variable refractive index. This invention combines the two. The last item in Table 6.11 addresses burn dressings. In this case, a polytetrafluoroethylene film is filled with silicone rubber to prepare a dual phase continuity structure. The silicone rubber provides high permeability allowing moisture to escape, while the polytetrafluoroethylene provides mechanical strength. Together, they are transparent, allowing better observation of the healing burn wound. A similar type of material is also used for scar treatments [Dillon, 1997].

Sound and Vibration Damping

One of the major effects of having extensive but incomplete mixing in two polymers, whether crosslinked or not, is that a very broad glass transition region may result, one that spans the temperature range between the two original glass transitions. Compared to their non-crosslinked counterparts, it is easier to prepare such materials with IPN technology, see Table 6.13. While the evidence is debatable whether such IPNs have more total damping (it seems to depend on the composition), having substantially the same damping level over a broad temperature range is good for outdoor and machinery applications. For example, a car door will be equally damped in winter and summer, and over broad frequency ranges, as well as at the standard 25°C, 600 Hz standard conditions. However, the vinyl-phenolic composition [Yamamoto and Takahashi, 1990] (in Table 6.13) damps specifically at the elevated temperatures at which machinery operates, but only weakly at room temperature.

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Table 6.10. Renewable resource IPNs Polymer 1

Polymer 2

Application

References

Castor oil-urethanes

Acrylics

Coatings for hydroelectric dams and iron, tough plastics

PEG Natural rubber

Dextran PMMA

Multistimuli-responsive drug Rubber-toughened plastics

Allyl cellulose cinnamonate Hyaluronic acid

Styrenics and acrylics Synthetic polymer

Improved thermal stability

Tan, 1994 Xie et al., 1993 Sperling et al., 1981 Kurisawa et al., 1995 Natural Producers, 1997 Schneider et al., 1996 Hourston and Romaine, 1990 Kamath et al., 1996

Biomedical and sanitary fields

Giusti and Callegaro, 1997

Table 6.11. Biomedical applications of IPNs Polymer 1

Polymer 2

Application

References

Poly(vinyl pyrrolidone)

HEMA

Calderara et al., 1996

Poly(dimethyl siloxane) Acrylic Poly(ε-caprolactone) PU Polyurea

Poly(acrylic acid) PEO-PPE PEG Malimide-PU Poly(2-hydroxyethyl methacrylate) Dextran Poly(methyl methacrylate) Polyurethanes Poly(dimethyl siloxane) Polyurethane or poly(acrylic acid)

Gradient-refractive index soft contact lenses High-permeability soft contact lenses Controlled transdermal drug delivery Bioerodible drug release Blood compatibility Contact lenses via reaction injection molding Drug delivery Artificial teeth Steam sterilizable tubing Burn dressing Drug release

Kurisawa et al., 1995 Roemer and Tateosian, 1984 Anon., 1983 Dillon, 1989 Lim et al.,1997

PEG Poly(methyl methacrylate) Silicone rubber Polytetrafluoroethylene Poly(N-isopropyl acrylamide)

Robert et al., 1995 Cho et al.,1995 Kim et al., 1995 Hsieh et al., 1997 Hill et al., 1996

Table 6.12. IPNs for non-linear optical applications Polymer 1

Polymer 2

Application

References

Polyurethane Epoxy Epoxy Polycarbonate

Polyacrylate Phenoxysilicon polymer Phenoxy Acrylic

Good temporal stability Frequency doubling Optical wave guide Second-order nonlinear optics

Xie et al., 1997 Tripathy et al., 1995 Martununkul et al., 1995 Tripathy et al., 1996

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Table 6.13. Sound and vibration damping materials via IPN formation Polymer 1

Polymer 2

Comments

References

Polystyrene Polyacrylate PU SBR Vinyl Polyurethane

Polyacrylate Polystyrene Poly(butyl methacrylate) Acrylic Phenolic Acrylic

Latex IPNs Latex IPNs Good damping Latex IPN, Polymer 3: SAN Damps machinery at high temperatures Full and semi-IPNs

Wang et al., 1996 Liu et al., 1997 Brovko et al.,1995 Hu et al., 1997 Yamamoto and Takahashi, 1990 Hourston et al., 1986b

6.5.6

A Few More Exciting Materials

One interesting invention is a poly(2,6-dimethylphenylene oxide)-poly(styrene-HEMA) based IPN that is useful for solar heating [Eck et al, 1995]. The IPN has an adjustable lower critical solution temperature, so that when coated onto the glass panels, clouds up if the temperatures inside get too hot. Polymer blends will not work so effectively for this application, because when they phase separate the phases may continue to grow; thus the time to come back to a one-phased material may be long. The IPN structure restrains phase domain growth, so that the material, acting as a smart material, clouds and clears relatively rapidly, reacting to the temperature of the structure it is protecting. Other areas of IPN application cover circuit boards [Tate and Varnell, 1989; Takahashi et al., 1996], and via gradient IPN technology, improved optical fibers [Bukhbinder and Kosjakov, 1996]. This last works by having a gradient refractive index from the axis to the surface, curving the light back away from the surface and reducing the fraction of light leaving the fiber for the exterior. Another area of interest involves coatings and adhesives. A recent patent describes IPNs based on epoxy and polyolefin resins for such materials [Ylitalo et al., 1997]. In summary, many of the IPN applications do not depend on high modulus or impact resistance (but some do!); these materials often show their best properties in the leathery, rubbery, or swollen states. Thus, it is seen that the applications for

IPNs covers a broad range today; the reader is directed to reviews by Sperling [1997] as well as by others for further reading.

6.6

Summary

The field of IPNs is simultaneously one of the oldest in multicomponent polymer literature, and one of its newest and fastest growing fields. With IPNs, it is relatively easy to prepare very small domain sizes and/or materials with dual phase continuity. IPNs can be made via a multitude of ways: sequential, simultaneous, latex, gradient, and thermoplastic, to name some of the more prominent materials. While IPNs can be and have been made extremely tough and impact resistant, many of the proposed applications involve such diverse fields and sound and vibration damping, biomedical materials, and non-linear optics. This is because the presence of crosslinks in both polymers reduces creep and flow, allowing relatively stable materials with a wide range of moduli to be prepared. Thus, those materials with leathery mechanical behavior, combinations of elastomers and plastics, are especially interesting to scientists, inventors, and engineers.

6.7

Abbreviations and Acronyms

AA — acrylic acid Bd — butadiene DACBA — diallyl carbonate of bisphenol-A

Interpenetrating Polymer Networks

DMS EA EMA EPDM

— — — —

dynamic mechanical spectroscopy ethyl acrylate ethyl methacrylate ethylene - propylene-dienemonomer hydroxyethyl methacrylate interpenetrating elastomer networks interpenetrating polymer network loss area latex interpenetrating polymer network n-butyl acrylate nitrile butadiene rubber polyamide poly(acrylic acid) polyacrylate polybutadiene polyisoprene poly(methyl methacrylate) polypropylene poly (2,6 - dimethyl phenylene oxide) polystyrene polyurethane poly(vinyl chloride) styrene styrene-butadiene rubber styrene-ethylene butylene-styrene simultaneous interpenetrating network tan δ area t-butyl peroxyisononate urethane vinyl acetate volatile organic compounds

HEMA IENs IPN LA LIPN

— — — — —

nBA NBR PA PAA PAcr PB PI PMMA PP PPE

— — — — — — — — — —

PS PU PVC S SBR SEBS SIN

— — — — — — —

TA TBPIN U Vac VOC

— — — — —

6.8

References

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CHAPTER 7

THE RHEOLOGY OF POLYMER ALLOYS AND BLENDS

L. A. Utracki1 and M. R. Kamal2

1

National Research Council Canada, Industrials Materials Institute, Boucherville, QC, Canada 2

McGill University, Montreal, QC, Canada

7.1

Introduction

7.1.1

Rheology of Multiphase Systems

Rheology is part of continuum mechanics. Thus, the basic principles of continuity, homogeneity and isotropy are incorporated into the basic rheological relations. The continuity principle requires that there is no discontinuity of material properties from one mathematical point to another; homogeneity demands that there is no concentration gradient, and isotropy implies that the flow does not impose orientation on the flow elements. In multiphase systems that comprise polymer alloys and blends, these three principles are seldom obeyed [Utracki, 1995]. The rheology of the multiphase systems follows its own sub-set of principles, extending the use of the general rheological dependencies. Obviously, the basic definitions of rheological functions, e.g., viscosity, η, dynamic shear moduli, G’ and G”, dynamic shear compliance, J’ and J”, etc., are identical. However, owing to the numerous influences, viz., concentration, morphology, flow geometry, time scale, type of flow field, thermodynamic interactions between the phases, and many others, it is difficult to relate the measured rheological functions to the intrinsic physical properties of the fluid. The rheological measurements of a multiphase system should be carried out in such a way that the length-scale of the flow is significantly larger than the size of the flow element. This makes it possible to treat the multiphase system as a homogeneous one, having an average, “specific” rheological behavior. For example, Brenner [1970] showed that relative viscosity, ηr, of diluted spherical suspensions, as measured in capillary flows, depends on the (d/D)2 factor, where d is the sphere diameter and D is the diameter of the capillary — for D ≅ 10d, the error in ηr, was 1%. Thus, if 1% error is the acceptable limit, the size of the dispersion should be at least 10 times smaller than the characteristic dimension of the measuring device, viz. radius of a capillary in capillary viscometers, distance between stationary and rotating cylinders or L.A. Utracki (Ed.), Polymer Blends Handbook, 449-546. © 2003 Kluwer Academic Publishers. Printed in the Netherlands.

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tion gradient and orientation of domains. Both of these depend on the magnitude of strain imposed on the fluid during the measurements, thus on the type of flow. Three types of flow are mainly used in the rheological measurements: steady state shearing, dynamic shearing, and elongation. The three can be classified according to the strain, γ, vorticity, as well as uniformity of stress, σ, and strain within the measuring space (see Table 7.1). The flow characteristics listed in Table 7.1 indicate that the steady-state flows strongly affect the morphology, whereas the dynamic flows have less influence. The extensional flows are characterized by uniform deformation and lack of vorticity, thus they are the most effective in changing the morphology and orientation of the system. In multiphase systems, there are many possible configurations of the interacting phases. Following the concepts of statistical mechanics, the rheological functions must be volume-averaged [Hashin, 1964]. The averaged quantities are sometimes known as bulk quantities. For example, the bulk rate of strain tensor, 〈˙γij〉, is expressed as:

plates in, respectively, the Couette or Weissenberg rheometer, etc. However, for many systems of industrial interest, the requirement of factor of 10 may imply a custom-made instrument. Instead, the data are usually generated with a smaller factor. As a result, the obtained data may be useful for comparative purposes (e.g., between different formulations, for quality evaluation, or automatic process control), but they do not have fundamental significance. Another aspect of the multiphase rheometry is related to the interrelations between the flow field and system morphology. In this text the term “morphology” will refer to the overall physical form or shape of the physical structure of a material, usually described as either a dispersed phase (particles or domains), co-continuous lamellae, fibrils or spherulites. Furthermore, morphology considers distribution and orientation of the phases, the interfacial area, the volume of the interphase, etc. However, the term must be distinguished from micro-morphology, which describes structures of the crystalline phase. Flow may induce two modifications of morphology that may complicate interpretation of data: the concentra-

Table 7.1. No.

Rheological models for miscible and immiscible blends

Type

γ

Vorticity

Uniformity of σ

1. Steady-state shear 1.1. Sliding plate, and rotational cone-&-plate 1.2. Poiseuille (capillary or slit), Couette, and rotational parallel plates 2. Dynamic shear 2.1. Cone-&-plate 2.2. Parallel plates 2.3. Couette 3. Extensional flows 3.1. Uniaxial 3.2. Biaxial

Comment γ

large

yes

homogeneous

homogeneous

large

yes

functions of spatial coordinates

functions of spatial coordinates

small small small

yes yes yes

homogeneous linear variable

homogeneous linear variable

mid mid

no no

homogeneous homogeneous

homogeneous homogeneous

for small gap, or for cone angle < 4° for laminar flows small measuring thickness is required for cone angle < 4° gap 0.8-2.0 mm gap 0.2-0.5 mm

The Rheology of Polymer Alloys and Blends

(7.1)

The stress tensor, 〈σij〉, in multiphase systems, is given by:

These works were reviewed by Barthes-Bisel [1988]. Within the higher concentration range, where the particle-particle interactions must be taken into account, Eq 7.2 is often approximated by a second-order polynomial. However, even for the hard-sphere suspensions, the theoretical extension of Eq 7.3 has been found difficult: (7.4)

(7.2) In Eqs 7.1 and 7.2 vi is local velocity, xi is local coordinate, ∆V is an elementary volume, p is pressure, δij is the unit tensor, ηo is viscosity of the continuous phase, while Sij and Fi represent hydrodynamic and non-hydrodynamic forces acting on a particle. These two functionals are usually coupled, as the thermodynamic interactions affect the hydrodynamic forces and vice versa. The first two terms on the right-hand-side (rhs) of Eq 7.2 are identical to those of a homogeneous fluid. For a multiphase system, these represent the stress tensor of the matrix liquid, while the third term describes the perturbing influences of the dispersed phase [Batchelor, 1974, 1977]. Owing to difficulties in deriving exact forms of the Sij and Fi functions in the full range of concentrations, Eq 7.2 is usually written as a power series in volume fraction, φ, of the suspended particles. The rheological behavior of multiphase systems within the linear, dilute region (φ < 0.05) is relatively well described. For example, for dilute suspensions of spherical particles in Newtonian liquids, Eq 7.2 reduces to Einstein’s formula for the relative viscosity, ηr: (7.3)

Equation 7.2 has been also solved for dilute suspension of anisometric particles [Hinch and Leal, 1972], elastic spheres [Goddard and Miller, 1967; Roscoe, 1967], and emulsions [Oldroyd, 1953, 1955; Barthes-Bisel and Chhim, 1981].

451

where the second-order coefficient was calculated as K = 5.2-7.6. These theoretical predictions should be compared with the experimental results. Thomas [1965] compiled the relative viscosity data, ηr vs. φ, measured in 16 laboratories for different types of hard-sphere suspensions, e.g., pollen in water, steel balls in oil, etc. After correcting the data (e.g., for immobilized adsorbed layer of the suspending liquid) the results superimposed and were fitted to the following relation (valid within the experimentally explored range of concentration, φ ≤ 0.6):

(7.5) From expanding Eq 7.5, the second-order coefficient of Eq 7.4 can be calculated as K = 10.43. Owing to difficulties in deriving general constitutive equations for multiphase systems, rheologists had to resort to simplified theoretical or semi-empirical dependencies derived for specific types of rheological tests and/or for specific multiphase systems. These, experimentally well established relations, constitute the basic tools for the rheological data interpretation of multiphase systems. They will be discussed in the following parts of the text. 7.1.2

Basic Concepts of Polymer Blends

In 1992, the world production of plastics reached 102 million m3/y with a value of over US$ 300 billion, while production of steel was 50 million m3/y

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with a value of about US$ 125 billion. Furthermore, between 1980 and 1990 plastics production increased by 62% while that of steel decreased by 21%. By the end of the Millennium, plastics production is expected to reach 151 million tons/year — polymers are the fastest growing structural materials. Within the plastics industry the most rapid progress has been observed for polymer blends. Polymer alloys and blends constitute over 30 wt% of polymer consumption, and with an annual growth rate of about 9.3% that has remained constant for the last ten years (i.e., four times the growth rate of the plastics industry as a whole), their role can only increase. In the text, the following standard definitions will be used [Utracki, 1989; 1991; see also Nomenclature in Chapter 1 of this Handbook]. 7.1.2.1 Definitions 1. Polymer blend is a mixture of two or more polymers and/or copolymers containing at least 2 wt% of the dispersed phase. 2. Miscible blend is a blend whose domain size is comparable to the dimension of a macromolecular statistical segment, or in other words, whose free energy of mixing is negative, ∆Gm < 0, and its second derivative of concentration, is positive: ∂2∆Gm / ∂φ2 > 0. Usually, miscibility is restricted to a relatively narrow range of independent variables, viz., molecular weight, composition, temperature, pressure, etc. — immiscibility dominates. 3. Polymer alloy is an immiscible, compatibilized blend, with modified interphase and morphology.

Table 7.2.

Alloying involves several operations that must result in blends showing stable and reproducible properties. These processes comprise: compatibilization, mixing and stabilization. Compatibilization may be accomplished either by addition of a compatibilizer or by reactive processing. Its role is to facilitate dispersion, stabilization of the morphology, and enhancement of the interaction between phases in the solid state. Today’s commercial alloys may comprise up to six polymeric ingredients. Development of such an alloy is complex, requiring knowledge of thermodynamics, rheology, processing and their influences on morphology, thus performance. If the rheology of suspensions and emulsions is difficult to describe theoretically and to determine experimentally, in the case of polymer blends these difficulties reach another order of complication. It suffices to note that in blends both phases are viscoelastic, the viscosity ratio varies over a wide range, and morphology can be very complex. To understand the rheological behavior of blends, it is useful to refer to simpler systems that can offer important insight. The following systems (Table 7.2) are commonly considered and will be treated in the following discussion. In addition, different types of the rheological analyses performed on polymer alloys and blends are listed in Appendix 7.1 — the table provides information on systems studied, conditions and assumptions, as well as it gives basic references to the topic. Appendices 7.2, 7.3, and 7.4 summarize the rheological information of blends comprising, respectively, commodity, engineering, and specialty resins.

Rheological models for miscible and immiscible blends

1. Miscible Blends

2.

Immiscible Blends

1.1. Solutions 1.2. Homologous Polymer Blends

2.1. Suspensions 2.2. Emulsions 2.3. Block copolymers

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The Rheology of Polymer Alloys and Blends

7.1.2.2 Phase Co-continuity Before discussing the rheological performance of polymer blends, it is important to recall the basic features of morphology in immiscible polymer blends. Addition of small quantity of a second polymer to the first one results in blends composed of a matrix (the first polymer) and the dispersed phase (the second polymer). When concentration of the dispersed phase is increased, the morphology may change from discontinuous dispersion of nearly spherical drops to progressively interconnected drops, then rods, fibers, and sheets. At a certain concentration, labeled as the phase inversion volume fraction, φI, the distinction between the disperse and matrix phases vanishes — the system morphology becomes co-continuous. It is important to recognize that the phase co-continuity is one of the most important aspects of blend morphology [Lyngaae-Jørgensen et al., 1999]. Since the morphology is strongly affected by large strain flow, it is to be expected that the method of specimen preparation influences the co-continuity. Both, the phase inversion concentration and stability of the co-continuous phase structure, depend on the strain and thermal history. It has been reported that the onset of cocontinuity occurs at an average volume fraction, φonset = 0.19 ± 0.09. In many branches of physics, the concept of percolation has been found useful. For example, when the concentration of conductive spheres in non-conductive medium exceeds the percolation threshold volume fraction, φperc, there is a sudden increase of electrical conductivity. For a three-dimensional case, 3D, the theory predicts that φperc = 0.156, for 1D it is φperc = 0.019. It was postulated that the observed changes of morphology in polymer blends belong to the group of percolation phenomena [Lyngaae-Jørgensen and Utracki, 1991]. In Figure 7.1, observed variation of phase co-continuity in blends of high density polyethylene with polystyrene, HDPE/PS, is shown. The data (obtained by selective extraction of the matrix phase) indicate that the onset of phase co-continuity occurred at φ1perc = 0.16 and φ2perc = 0.15, whereas φI = 0.64.

The co-continuity contributes to synergism of properties, e.g., advantageous combination of high modulus and high impact strength in commercial blends. Two expressions for predicting the phase inversion concentration from the viscosity ratio were proposed: φ I1 / φ I 2 = η1 / η2 ≡ λ φ I 2 = (1 + F(λ )λ ) where

or

φ I 2 = (1 + λ )

−1

(7.6)

−1

F(λ ) = 1 + 2.25 log λ + 1.81(log λ )

2

(7.7)

Note that φ1I = 1 - φ2I and φ2I is the volume fraction of liquid 1 and 2, respectively, at the phase inversion. Equation 7.6 is empirical, proposed by Paul and Barlow [1980] as a generalization of the experimental observations reported by Avgeropoulos et al. [1976]. Equation 7.7 was derived from the filament instability equation by Metelkin and Blekht [1984]. These relations are applicable to systems prepared at low stresses, thus in these equations the viscosity ratio, λ, should correspond not to the ratio of the zero-shear viscosities, but to its value at the shear stress used to prepare the blends. The relations were found to describe the phase inversion for systems with nearly equal polymer viscosities, where λ → 1. As the viscosity ratio increases, these equations predict too rapid change of φ2I .

Figure 7.1. For immiscible blends the onset of phase co-continuity should coincide with the percolation threshold. Theoretically, φperc = 0.156 for 3D flow of immiscible system. Experimentally, φ2perc = 0.19 ± 0.09 was found [LyngaaeJørgensen and Utracki, 1991].

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To derive a more general relation for the phase inversion concentration, one may start by computing: η(A) = ηoBηr (φA) and η(B) = ηoAηr (φB), where ηr is the relative viscosity. The latter dependence can be expressed as [Krieger and Dougherty, 1959]: (7.8) In Eq 7.8 φm is the maximum packing volume fraction, and [η] is the intrinsic viscosity. The computed curves are shown in Figure 7.2. To calculate these dependencies φm = 0.8 and [η] = 2 were assumed. The six points of intersection represent the iso-viscous conditions for the dispersions of liquid 1 in 2 and liquid 2 in 1, or in other words, the conditions for phase inversion.

Equation 7.9, can be expanded into MacLaurin’s series, then truncated after the second term to give a simplified version, valid within the range: -1 < φIi /φm < 1: (7.10) In Figure 7.3 experimental dependence of λ vs. φ2I for thermoplastics polymer blends is presented. The straight line represents Eq 7.10. For most polymer blends, the following values of parameters in Eq 7.9: [η] ≅ 1.9; φm ≅ 1 - φperc = 0.84, provide good approximation. It should be noted that the steady state viscosity ratio should be taken at a constant stress (not deformation rate). The “sharpness” of the phase inversion peak very much depends on the distribution of stresses within the mixing device, as well as on the absolute magnitude of polymer viscosity — the wider the distribution of stresses and/or the higher the viscosity, the wider the range of concentrations at which the phase inversion takes place. Since many experiments are conducted using an internal mixer known to possess a wide range of flow conditions, instead of a single point, usually a range of concentrations for the phase co-continuity has been reported (see Figure 7.3).

Figure 7.2. Concentration dependence of emulsion viscosity. Solid line represents η = η1(φ2) while the other lines the same dependence for η = η2•ηr(φ1). To calculate these dependencies [η] = 2 and φm = 0.8 were assumed. The intercepts correspond to the iso-viscous conditions defining the phase inversion concentration, φ2I = 1 - φ1I [Utracki, 1991].

In consequence, from Eq 7.8 the iso-viscous point can be expressed as:

(7.9)

Figure 7.3. Experimental φ2I vs. λ dependence for mechanically prepared thermoplastics blends. The dotted, broken and solid lines represent Eq 7.6, 7.9 and Eq 7.10, respectively; the values: [η] = 1.9 and φm = 0.84 were used [Utracki, 1991].

The Rheology of Polymer Alloys and Blends

455

Recently, another relation was proposed for predicting the phase inversion. It was derived from a similar assumption than that used by Metelkin and Blekht [1984] that at the phase inversion the morphology of both phases is fibrillar, and that the rate of fiber disintegration is the same for both components [Luciani, 1993]:

(7.11) The significance of the function Ω (Λ, λ) will be discussed in the Part 7.3.1.2, dedicated to the emulsion microrheology.

Figure 7.4. Theoretical representation of the interface, with the definition of the interphase thickness; χ and b are respectively the binary interaction and the lattice parameter [Helfand and Tagami, 1971].

7.1.2.3 The Interphase Lattice theory predicts that the density profile across the interface follows an exponential decay function [Helfand and Tagami, 1971, 1972]:

(7.12)

In Eq 7.12, χAB is the thermodynamic binary interaction between polymers A and B, and b is a lattice parameter. The dependence is shown in Figure 7.4. The intercept of the tangential line at the place of the steepest decline (or incline for the other component) defines thickness of the interphase, ∆l. The lattice theory of the interface predicts that there is a reciprocity between the interfacial tension coefficient, ν12, and the interphasial thickness, ∆l [Helfand and Sapse, 1975]:

independent of the thermodynamic binary interaction parameter, χΑΒ. The theory leads to the conclusions that: (i) surface free energy is proportional to square root of χΑΒ, (ii) the chain-ends of both polymers concentrate at the interface, (iii) any small molecular weight third component will be repulsed to the interface, and (iv) interfacial tension coefficient increases with molecular weight to an asymptotic value: ν12 = ν∞ - aoMa-2/3 . These conclusions were found to offer good guidance for development of the compatibilization strategies. There are several other theories of the interface, some of which lead to quantitatively different results [Ajji and Utracki, 1995, 1996, 1997]. For example, Noolandi [1984] considered a binary system compatibilized by addition of a block copolymer. For χΑΒNcφp ≤ 2 he derived:

(7.13)

(7.14)

where a, b, and m are lattice parameters. Notice that according to this theory, the product, ν12∆l, is

where ao is a numerical parameter, while φc, φp, and Nc are respectively volume fraction of copo-

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L. A. Utracki and M. R. Kamal

lymer, of polymer, and degree of polymerization of the copolymer. A semi-empirical dependence of the interfacial tension coefficient on compatibilizer concentration can be derived from an analogy to titration of an emulsion with surfactants [Utracki, 1992]:

(7.15) where νo is the initial interfacial tension coefficient at zero concentration of copolymer, νCMC, is the interfacial tension at saturation of the interface, and φ is the copolymer concentration. Eq 7.15 adequately described the interfacial tension coefficient in the system polystyrene/polybutadiene compatibilized by addition of styrene-butadiene block copolymer, see Figure 7.5 [Anastasiadis et al., 1988, 1989]. Recently an exponential decay relation was proposed [Tang and Huang, 1994]: (7.16) where K is a parameter — from Eq 7.14, its value should be proportional to: K ∝ χΑΒNc. There have been several efforts to provide means for computation of the interfacial tension coefficient from characteristic parameters of the two fluids [Luciani et al., 1996a]. The most interesting relation was that found between the interfacial tension coefficient and the solubility parameter contributions, that are calculable from the group contributions. The relation makes it possible to estimate the interfacial tension coefficient from the unit structure of macromolecules at any temperature. The correlation between the experimental and calculated data for 46 polymer blends was found to be good — the correlation coefficient R = 0.815 — especially when the computational and experimental errors are taken into account.

Figure 7.5. Interfacial tension coefficient vs. concentration of compatibilizer for polystyrene blends with polybutadiene, compatibilized with styrene-butadiene block copolymer. Data points [Anastasiadis and Koberstein, 1988], line computed from Eq 7.15.

There are several methods for measuring the interfacial tension coefficient for low-viscosity liquids, e.g., spherical shape recovery after slight deformation, liquid thread break-up, rotating bubble or drop, pendant drop, sessile bubble or drop, du Nuouy ring, or light scattering. For high viscosity polymeric melts, they can be used with decreasing reliability. The most recent and highly successful method involves spherical shape recovery of a drop deformed by about 15% either in shear or (preferably) in elongation. Since the drop can be repetitively deformed and its shape recovery followed this method is the only one that makes it possible to follow the time evolution of the interfacial tension coefficient. Furthermore, the method also makes it possible to examine whether, for a given polymer pair, the interfacial energy is symmetrical, i.e., if ναβ = νβα [Luciani et al., 1996a, b]. There are fewer methods available to measure the interphase thickness, e.g., ellipsometry, microscopy, and scattering. A summary of the measured ∆l is given in Table 7.3. The temperature dependence of ∆l in PMMA/SAN and PMMA/PS blends is presented in Figure 7.6. More information on the fundamentals, methods of measurements, and numerical values of the interfacial ten-

The Rheology of Polymer Alloys and Blends

sion coefficient or the thickness of the interphase can be found in Chapter 4 Interphase and Compatibilization by Addition of a Compatibilizer. Table 7.3.

Interphase thickness

Type of blend

In the first case one can learn about the effects of the thermodynamic interactions between chemically different components on the flow behavior. In the second case, it is the effect of molecular weight and molecular weight distribution that can be observed.

Thickness (nm)

7.2.1 Immiscible blend Block copolymer interphase Polymer/Copolymer Reactive Compatibilization Radius of gyration, 〈Rg2〉1/2

457

2 4 to 6 30 30 to 60 5 to 35

Solutions

For solutions [Glasstone et al., 1941]:

(7.17) where V is the specific volume, and xi is the mole fraction. For miscible blends, ∆Hm < 0 and the above relations predict a positive deviation from the log-additivity rule, PDB. The latter rule, the log-additivity, was formulated by Arrhenius [1887]: (7.18) with the excess viscosity term, lnηE → 0. There are several other blending rules for solution viscosity, e.g., [McAllister, 1960]:

Figure 7.6. Interphase thickness vs. temperature for polymethylmethacrylate blends with (from top) styrene-acrylonitrile copolymer and polystyrene [Kressler et al., 1993].

7.2

Rheological Models for Miscible Blends

By definition, miscible polymer blends are singlephase mixtures. Miscibility depends on the molecular weight, concentration, temperature, pressure, deformation rate, etc. Flow of these systems can be compared to that of solutions of low molecular weight, miscible components, or to flow of mixtures of polymeric fractions. Both models are far from perfect, but they serve to illustrate the basic behavior of miscible systems.

(7.19)

where ηk,i indicates the kinematic viscosity, Mi is the molecular weight, and the two kinematic viscosities with double subscripts are the empirical interaction viscosities. Equation 7.19 was derived from a three-body model of a miscible mixture comprising two low molecular weight liquids with two interaction viscosities. 7.2.2

Homologous Polymer Blends

Homologous macromolecular blends are simply mixtures of fractions of the same polymer having

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the same molecular constituents. On the one hand, any commercial polymer may be treated as a homologous macromolecular blend, and on the other, blending narrow molecular weight distribution fractions provides important information on the rheological behavior of commercial materials. Since the zero-shear viscosity for narrow molecular weight distribution samples can be expressed as: ηo = KMwα ; thus

Mw = (ηo / K )1/ α

(7.20)

(where K and α = 1 or 3.4 are parameters), but since: Mw =

∑w M i

(7.21)

i

i

then it follows that [Friedman and Porter, 1975]:  ηo =  

∑ i



wiη1o/,iα  

α

(7.22)

For binary mixtures, Equation 7.22 predicts that viscosity should show a positive deviation from the log-additivity rule, PDB. There is a mounting evidence that PDB is not a rule for miscible polymer blends. Depending on the system and method of preparation, polymer blends can show either a positive deviation, negative deviation, or additivity. Note that miscibility in polymeric systems requires strong specific interactions, which in turn affect the free volume, thus the rheological behavior. It has been demonstrated that Newtonian viscosity can be described by the relation [Utracki, 1983; 1985; 1986]: (7.23) where ai are equation parameters: ao ≡ 1nηo* with ηo* being the iso-free volume viscosity, a1 is a function of the molecular architecture and polydispersity (a1 = 0.79 was found for all paraffin’s and their mixtures), and a2 = 0.07 is the linearization parameter. The key to Eq 7.23 is the free volume fraction, f, computed from Simha’s statistical theory [Simha and Somcynsky,

1969; Simha and Jain, 1984]. This approach was successful in describing pressure, temperature, and concentration dependence of the viscosities of solvents and polymer melts [Utracki, 1983; 1985; 1986].

7.3

Model Systems for Immiscible Blends

Most polymer blends are immiscible. Their flow is complex not only due to the presence of several phases having different rheological properties (as it will be demonstrated later, even in blends of two polymers the third phase, the interphase, must be taken into account), but also due to strain sensitivity of blend morphology. Such complexity of flow behavior can be best put in perspective by comparing it to flow of better understood systems, suspensions, emulsions, and block copolymers. Flow of suspensions of solid particles in Newtonian liquids is relatively well understood, and these systems provide good model for flow of polymer blends, where the viscosity of dispersed polymer is much higher than that of the matrix polymer. Flow of emulsions provides the best model for polymer blends, where the viscosity of both polymers is comparable. The microrheology of emulsions provides the best, predictive approach to morphological changes that take place during flow of polymer blends. The effect of emulsifiers on the drop size and its stability in emulsions has direct equivalence in the compatibilization effects in polymer blends. Finally, the rheological behavior of block copolymers serves as a model for well compatibilized blends, with perfect adhesion between the phases. The copolymers provide important insight into the effects of the chemical nature of the two components, and the origin of the yield phenomena. 7.3.1

Suspensions

There are two reasons for discussing the solid-inliquid dispersions in the chapter dedicated to flow of polymer blends [Utracki, 1995]. Historically, the first systematically studied multiphase systems

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The Rheology of Polymer Alloys and Blends

were suspensions in Newtonian liquids, first at infinite dilution [Einstein, 1906, 1911], then at increasingly concentrated limits [Simha, 1952]. Knowledge of these derivations is fundamental to understanding the energy dissipation during flow in any multiphase system. Furthermore, suspensions in viscoelastic matrix are good models for polymer blends having viscous polymer dispersed in a significantly lass viscous matrix polymer. 7.3.1.1 Suspensions in Newtonian Liquids The following assumptions are often used: (i) the size of a rigid particle is large in comparison to the suspending medium molecules, but small compared to the smallest characteristic diameter of the flow channel, so the continuum theories are applicable; (ii) the flow is steady state, without inertia or sedimentation; (iii) the suspending medium perfectly adheres to the particles. Depending on the system (as well as the author), additional assumptions may be made, e.g., regarding interparticle interactions, orientation, etc. Denisov et al. [1985] as well as Brady and Bossis [1985] reported on numerical simulation of suspension rheology. The first authors used the 6-12 Lennard-Jones potential with the usual meaning of ε* and σ* characteristic constants (with dimensions of energy and length, respectively) of the interacting species. Taking Ro as a measure of distance from the center of the particle at which action of the potential begins, the necessary conditions for dilatant behavior were: (i) Ro ≥ σ* and (ii) particle concentration exceeding a critical value dependent on the system. The Stokesian dynamic’s method was used by the other authors. The simulation provided valuable information on the influence of various microstructural elements on the macroscopic viscosity. The relative velocity of two particles in suspension provided the most important contribution to energy loss. As φ increased, the correlation of interparticle motion also increased. Hydrodynamic lubrication resulted in an increased number of particles acting as a single agglomerate. The maximum packing volume fraction, φm, takes on a meaning as a percolation-like threshold for the viscosity to

increase to infinity owing to the formation of infinite clusters. Microstructural theories of suspensions appear to be particularly well suited to solve problems associated with time-dependent flows, thixotropy and rheopexy (anti-thixotropy) [Russel, 1983; Utracki, 1989, 1995]. 7.3.1.1.1 Relative Viscosity of Suspensions One of the most interesting derivations of the η vs. φ dependence (covering the full range of concentration) was published by Simha [1952]. He considered the effects of concentration on the hydrodynamic interactions between suspended particles of finite size. (Note that previously the particles were simply considered point centers of force that decayed with cube of the distance.) Simha adopted a cage model, placing each solid, spherical particle of radius a inside a spherical enclosure of radius b. At distances x < b, the presence of other particles does not influence flow around the central sphere and the Stokes relation is satisfied. This assumption leads to a modified Einstein [1906, 1911] relation: (7.24) where λ(y) is the modifying (or shielding) function of the relative cage size, y ≡ a/b: λ ( y) = with

4(1 − y 7 ) 4(1 + y ) − 25 y 3 (1 + y 4 ) + 42 y 5 10

[

]

y = 2(1 / φ˜ )1/ 3 − 1

(7.25)

−1

In Eq 7.25 φm is the maximum packing volume fraction. Thus, the magnitude of the shielding function λ(y) depends on the reduced volume fraction, o ≡ φ / φm. At low concentration, o → 0, the shielding factor vanishes and Einstein’s relation is recovered. However, at high concentration, o → 1, the shielding function and relative viscosity both go to infinity, λ(y), ηr → ∞. Substituting Eq 7.25 into Eq 7.24 and expanding it into power series make it possible to write simplified versions, valid respectively within the low (viz. Eq 7.26) and high (viz. Eq 7.27) concentration range:

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L. A. Utracki and M. R. Kamal

25 21 625 ˜ 2 ηsp ≡ ηr − 1 = (5φ / 2)1 + φ˜ − φ˜ 5 / 3 + φ + ... 64 128  32 

(7.26) and:

[

lim ηr = 27φ / 20 φ˜ (1 − φ˜ ) −3 φ →φ m

]

(7.27)

Two others semi-empirical relations have been often used to describe the concentration dependence of suspension viscosity. The first, was derived for the first time by Mooney [1951]: ln ηr = [η] / (1 − φ˜ )

(7.28)

where [η]s is the intrinsic viscosity. The subscript s indicates that the parameter refers to solid particles. The intrinsic viscosity is defined as: [η] = lim (η − ηo ) / ηoφ = lim d (ηr − 1) / dφ = .

φ ,γ → 0

lim d ln ηr / dφ

.

φ ,γ → 0

Figure 7.7. Relative viscosity of hard-sphere suspension in Newtonian fluid as a function of the volume fraction. Thomas curve represents the generalized behavior of suspensions as measured in 19 laboratories. The remaining curves were computed from Simha’s, Mooney’s and Krieger-Dougherty’s relations assuming Einstein value for intrinsic viscosity of hard spheres, [η] = 2.5, but different values for the maximum packing volume fraction, φm = 0.78, 0.91, and 0.62 respectively.

(7.29)

.

φ ,γ →0

The second dependence, is the already cited Equation 7.8, derived by Krieger and Dougherty [1959]. The relation belongs to a large group of dependencies of the type, discussed in detail a few years back [Utracki, 1989]: (7.30) In Eq 7.30, the values of the semi-empirical parameters, α, and β, are usually constant, e.g., respectively 2.5 and 1, or -2.5 and -1, or -1.73 and -2.0, etc. However, in the Krieger-Dougherty relation, these two parameters depend on the system: α = 1/φm, and β = [η]φm. In Figure 7.7, the plots of ηr vs. φ calculated from Simha’s Eq 7.24, Mooney’s Eq 7.28, and Krieger-Dougherty’s Eq 7.8 are compared with the empirical curve-fitted relation, Eq 7.5. For all the relations, the intrinsic viscosity [η]s = 2.5 was used. However, to optimize the fit, different values for the maximum packing volume fraction, φm = 0.78, 0.91, and 0.62 respectively, had to be used. Detailed analysis of Thomas’ data made it possible to conclude

that Simha’s relations provide the best fit with more realistic values of the physical parameters [Utracki and Fisa, 1982]. To summarize, the dependence of relative viscosity on the volume fraction of suspended particles can be expressed by any of several theoretical or semi-empirical relations. These can be written in terms of the two parameters, [η] and φm. Thus ηr = ηr([η], φ/φm). As it will be shown, the generality of this dependence extends beyond the monodispersed hard sphere suspensions. The relationships between ηr and φ have been derived for suspensions of monodispersed hard spheres in Newtonian liquids. However, most real systems are polydispersed in size, and do not necessarily consist of spherical particles. It has been found that here also Simha’s Eq 7.24, Mooney’s Eq 7.28, or Krieger-Dougherty’s Eq 7.8 are useful, provided that the intrinsic viscosity and the maximum packing volume fraction are defined as functions of particle shape and size polydispersity. For example, by allowing φm to vary with composition, it was possible to describe the ηr vs. φ variation for bimodal suspensions [Chang and Powell, 1994]. Similarly, after values

The Rheology of Polymer Alloys and Blends

of [η] and φm were experimentally determined, Eq 7.24 provided good description for the ηr vs. φ dependence of several multiphase systems, e.g., PVC emulsions and plastisols, mica-reinforced polyolefins, sealant formulations [Utracki, 1988; 1989]. The problem of packing a maximum volume of solids into a given space is common to numerous branches of physics and technology. It suffices to note that the relative viscosity of suspensions is a function of the reduced volume fraction, o ≡ φ / φm, to realize the importance of φm. Experimentally, it was demonstrated that φm calculated from dry packing of solid particles agrees well with the value determined for a suspension. Theoretically and experimentally, polydispersity increases the φm value, for example from 0.62 observed for random packing of uniform spheres to values exceeding 0.9. An interesting recipe for φm maximization requires four generations of nearly spherical particles with the diameter ratios: 1:3:9:17. Blending them at the volume ratios: 4:1:1:4 result in φm = 0.78. However, what was important, the suspensions were found to be non-sedimenting, and when dried they gave a solid bed with uniform porosity [Ritter, 1971; Lord, 1971]. In industrial practice it may be important to use mixtures of filler particles not only of spheroidal shape (as discussed above) but also of different shapes, e.g., filling and reinforcing polymer with CaCO3 particles and glass fibers. The theoretical basis for optimization of such systems was developed by Wieckowski and Strek [1966], and later by Milewski [Milewski, 1973, 1977, 1978; Milewski and Katz, 1987]. These studies are also important for polymer blends where, at concentrations exceeding the percolation threshold, the morphology is complex, comprising spheres, fibers and lamellas. For anisometric particles it is useful to use the particle aspect ratio, p, defined as a ratio of two orthogonal axes. For prolate ellipsoids (fibers) p > 1 is the length-to-diameter ratio, whereas for oblate ellipsoids (plates) p < 1 is the thickness divided by the largest dimension of the plate. It was observed that both, the intrinsic viscosity, [η], and the inverse of the maximum packing

461

volume fraction, 1/φm, increase linearly with p. Thus, the relative viscosity of suspensions of anisometric particles is higher than that observed for spheres. For example, Doi and Edwards predicted [1978] that for rods ηr ∝ φ3. In the extensional, irrotational field, under steady state conditions, the particles remain oriented in the direction of stress. In uniaxial flow, they align with the main axis in the flow direction, while in biaxial they lie on the stretch plane [Batchelor, 1970, 1971]. For dilute spherical suspensions in Newtonian liquid the extensional viscosity follows the Trouton rule, i.e., ηE ≅ 3η. However, for anisometric particles the Trouton ratio, ηE / η is a strong function of p. For example, at φ = 0.01 extensional viscosity of rods with aspect ratio p = 1000 is 1000 times higher than that for suspension of spheres. 7.3.1.1.2 Particles Orientation in Flow The orientation of particles in flow is of particular interest to microrheology. To predict the macroscopic rheological properties of a multiphase system, a detailed description of each phase behavior is required. In this field, contributions from the Pulp and Paper Research Institute of Canada by Mason et al., and later by van de Ven and his co-authors are particularly valuable. The earlier results were summarized by Goldsmith and Mason [1967], the latter by Van de Ven [1989]. The microrheology has been particularly well developed for infinitely dilute systems in Newtonian matrix — either solid particles or liquid drops. In the present part, only the former system will be summarized. More extensive discussion of microrheology of the liquid/liquid systems will be presented later, while considering the rheological behavior of polymer blends. For suspension of solid particles in a liquid, the theoretical and experimental works indicate that the angle of orientation of a spheroid can be expressed as: φ1 = arctan [p tan (2 π t / tp)]

(7.31)

where the period of rotation of a particle with an aspect ratio, p, is given by:

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(7.32) Accordingly, for rods, the maximum velocity of rotation occurs at t/tp = 0, 1/2, 3/4, 5/4,... For spheres with p = 1, φ1 = 2πt/tp, i.e., a constant rotational velocity. In non-uniform shear fields, such as that observed during flow through a capillary (Poiseuille flow), the particles rotate with velocity predicted by Eqs 7.31 and 7.32, according to the value of the shear rate existing at the radial location of the sphere in the capillary. Near the wall, for finite diameter spheres, the immobile layer of the suspending medium causes a reduction of rotational and translational velocity. The effect scales with the square of the sphere diameter. The wall also causes a geometric exclusion effect, i.e., a lower-than-average concentration of particles near the wall and a retardation of their motion. The phenomena are complicated by the axial migration of particles, dependent on the Reynolds number, Re = ρyd2 / ηοwhere ρ and d is the particle density and diameter, respectively. 7.3.1.1.3 Shear-induced Particle Migration There are at least two possible mechanisms for the particle migration during shear flow, inhomogeneity of the stress field and strong inter-particle interactions [Graham et al., 1991]. In the first case, the particles tend to migrate to low shear stress regions, while in the second case, the situation is more complex involving a coupled relationship between the thermodynamic and hydrodynamic forces. The Newtonian behavior of suspensions in Newtonian liquids is limited to low concentrations. An exception seems to be the extensional flow of anisometric particles (irrotational flow field) where the rate of strain independent region extends to concentrations where strong non-Newtonian behavior would be expected in shear. These rate of deformation dependent phenomena will be summarized below. During the capillary flow of concentrated suspensions, the difference in velocities of particles located at different radial positions results in

the formation of transient multiplets or stacks, behaving similarly to rods. Under these circumstances, the rate of axial migration is accelerated, and the flow profile flattens. For example, experimentally, for φ = 1/3 suspensions of spheres flowing through a tube at the Reynolds number Re ≡ 2ρQ/πR1η = 0.056 a partial plug flow was observed. However, when Re reached the value of 0.112, a complete plug flow was observed — the flow was no longer Newtonian [Karnis et al., 1966; Vadas et al., 1973]. Matsumoto et al. [1986] reported that in the cone-and-plate geometry, the storage, G’, and loss, G”, shear moduli of uniform, non-rigid spheres decrease monotonically with test time (or number of shearing cycles). G’ and G” were observed to decrease by 4 decades, but steady state shearing for 15 seconds returned them to the initial values. Since the phenomenon depended on the rigidity as well as on the uniformity of shape and size, development of a structure during the dynamic test must be postulated. In Couette flow, spheres migrate toward the outer cylinder. In shearing, a shear fractionation of spherical particles has been observed. For example, Giesekus [1981] observed that, during torsional shearing of binary sphere suspensions, the larger and the smaller spheres separated into two different annular volumes, i.e., for each sphere size a critical equilibrium radial distance had to be postulated. On the other hand, Prieve et al. [1984, 1985] reported that for each sphere diameter and speed of rotation, there is a critical radius, rc. In the parallel plate rheometer, a particle located at r < rc was observed to migrate inward, whereas that placed at r > rc migrated outward. There is no theoretical explanation for either observation. In a wide-gap Couette rheometer, migration of spheres was followed by nuclear magnetic resonance imaging [Abbott et al., 1991]. Migration to the low shear rate region was found to be determined by the total strain, proportional to the shear rate and square of the particle diameter, but independent of the (Newtonian) viscosity of the matrix liquid. More recently, similar studies were undertaken for suspensions of rods with

The Rheology of Polymer Alloys and Blends

p = 2-18 and φ = 0.3 or 0.4 [Mondy et al., 1994]. At the same φ, the composition gradient of rods of different aspect ratios was indistinguishable, the same as the one earlier reported for spheres. The rate of migration was found to increase with concentration. Owing to the periodically accelerated rotation of fibers in a shear field, alignment of fibers in Couette flow is to be expected. Theory indicates that the shear field is about half as efficient at causing fiber alignment as extension. However, the shear field is rarely homogeneous, and during the flow fibers undergo breaking, bending or coiling, which cause further reduction of alignment efficiency. Further details on various modes of orientation behavior of flowing suspensions can be found in reviews by Cox and Mason [1971], Batchelor [1974], and Leal [1980]. The evidence accumulated so far indicates that there is a full spectrum of structures, from a liquid-like where the yield stress, σy = 0, to a solid-like with large σy. For anisometric particles at φ > 1/p, yield may originate in mechanical interlocking of particles, but for spheres it stems from the interparticle interactions. When these interactions are weak σy → 0 is observed, with the arrow indicating the time effect. If the experiment is conducted at low rates of shear, no yield behavior would be noted. In uniaxial extensional (convergent) flow, there is evidence of spherical particles moving toward the center of the stream. Convergent flow of a dilute suspension of glass fibers, p = 200-800, in Newtonian liquids was studied by Murty and Modlen [1977]. The fiber orientation angle (defined as an average angle between the fiber axis and flow direction) changed from 45° (random) to about 15°. The orientation started upstream from the convergence. For low viscosity liquids, jamming at the entrance region was responsible for as much as 60% of fibers being “filtered out.” At higher fiber loading, φp > 1, the rheological responses of aligned fiber suspensions resemble those of liquid crystals. Becraft and Metzner [1992] analyzed the rheological behavior and orientation of glass fibers, GF, in polyethylene, PE, and polypropylene, PP. The experimental data

463

were interpreted using a modified Doi theory for liquid crystalline fluids, LCF [Doraiswamy and Metzner, 1986]. The kinetics of the distribution function, f, is given by: Df f∇V = ∇·Dr∇f + ∇·Dr - ∇·(uf) (7.33) Dt kBT where E is the rotational diffusivity of the rods, and u denotes a unit vector corresponding to rod orientation. The first term on the rhs of Eq 7.33 accounts for the contribution of Brownian motion to the orientation distribution function, the second for the effects of the liquid crystalline interaction potential between the rods, while the third term for the effects of flow. Doraiswamy and Metzner noted that use of the LCF approach is permissible at concentrations above that which would correspond to the transition from isotropic to aligned morphology, φ > 8/p. The theory provided fair description of the stressstrain dependence for systems containing 10 wt% GF and excellent agreement for those with 40 wt% GF. Also, the approach gave good predictions of the diagonal terms of the second-order orientation tensor. 7.3.1.1.4 Aggregation and Yield Stress One of the fundamental assumptions of the continuum theories is stability of structure (Newtonian behavior) or, alternatively, a well-defined process of structural changes (non-Newtonian behavior). However, as it was already mentioned, orientation effects in sheared layers of suspensions are responsible for either dilatant or pseudoplastic behavior, while strong interparticle interactions may lead to yield stress or a transient behavior. In short, there is an intimate relation between the liquid structure and its rheological response; change in one causes a corresponding change in the other. Some of these changes have been theoretically treated, viz. Eq 7.32. Aggregation, agglomeration and flocculation are structural phenomena ranging from transient rotating doublets observed within dilute region, to a pseudo-solid-like behavior of flocculated suspensions with yield stress. Aggregation can occur

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due to either thermodynamic interparticle interactions, chemical bonding, or geometric crowding. The latter type prevails in shear flows of suspensions of anisometric particles. There are numerous theories based on structural models of suspensions [Mikami, 1980]. Wildemuth and Williams [1984] considered that the maximum packing volume fraction, φm, is a function of normalized shear stress, σ˜ 12 ≡ σ 12 / M , where M is a numerical parameter. The authors derived the relation: m m φ m = (φ m,o + φ m,∞σ˜ 12 ) / (1 + σ˜ 12 )

(7.34)

where m = 1.00-1.17 is an experimental constant, while φm,o and φm,∞ are values of the maximum packing volume fraction at σ˜ 12 = 0 and σ˜ 12 = ∞ , respectively. Defining the yield stress as σ˜ 12 when φ = φm, one can rearrange Eq 7.34 to read:

σ˜ y = (φ − φ m,o ) / (φ m,∞ − φ )

(7.35)

Hoffman [1972, 1974] also reported that at low rates of shear and high solid content, φ > 0.54, the power law index n = dlnσ12 /dlny approaches zero. For the system: styrene-acrylonitrile (SAN) latex in ethylene glycol, addition of salt decreased n to zero. The experiment was performed to demonstrate that increased interparticle interaction causes the onset of dilatation to move to higher rate of shear. However, the work also demonstrated that at these high concentrations there is a yield stress, σy. Onogi and Matsumoto [1981] reported that in PS suspension with particles having strong attractive forces, a yield phenomenon was observed, while suspensions of PS-particles having repulsive forces behaved like Newtonian liquids. Thus, the yield stress is associated with formation of a three-dimensional structure by interacting particles, resulting in a behavior similar to an elastic solid. Similarly, impact modification of PMMA by incorporation of 0-50 wt% of core-shell latex particles of poly(butylacrylateco-styrene), demonstrated that the particles form a co-continuous network at φ ≤ 0.2 that result

in a low frequency rubbery plateau. The high frequency data were found independent of composition [Bousmina and Muller, 1992]. There are several methods for determining σy. Among these is the modified Casson equation [Utracki, 1982]: (7.36) where F may be any rheological function (viz. shear stress, σ12, elongational stress, σ11, shear loss modulus, G”, etc.), Fy indicates the yield value of F, Fm is the F-value of the matrix liquid at the same deformation rate as F, and a is a measure of the relative value of F. Another method requires a simultaneous fit of experimental data to a constitutive equation in which a parameter or parameters are related to σy [Utracki, 1987]. Measurements of creep and elastic recovery also provide sensitive, direct means of detecting yield stress, either by simultaneous fit of time dependent strain, γ(τ), at a constant stress, σ12, to the compliance equation: (7.37) [where Ψ(t) is the retardation function], or by plotting the recoverable strain versus stress. In the latter case, the maximum value of stress below which the Hookean behavior is obtained gives the value of σy. Polymer lattices and suspensions of carbon black in linseed oil and clay or calcium carbonate in aqueous media provide examples [Amari and Watanabe, 1983]. The values of σy determined from creep and those from shear viscosity were found to be in good agreement. There are several direct methods of measurement of yield stress. The constant stress rheometer is most frequently used to determine σy value in shear. Dzuy and Boger [1983, 1985] used a rotational vane viscometer. Yield stresses in compression can be calculated from the unrelaxed stress values in parallel plate geometry. Its value in elongation has been directly measured as the critical stress value below which no sample deformation was observed during 30 minutes of straining in an extensional rheometer.

The Rheology of Polymer Alloys and Blends

For unoriented particle systems, the von Mises criterion for plastic flow of solids should be obeyed. The yield stress in elongation and compression should be equal to each other, and larger by the factor of √3 than the yield stress in shear, σy. However, for highly concentrated suspensions of anisometric particles, von Mises criterion should not be used. For suspensions, the concentration dependence of σy was found to follow either one of the following two dependencies:

and

(7.38)

where ai are adjustable parameters. The exponent a2 depends on the particle geometry as well as the inter-particle interactions. For human blood σy = 26.87φ3 (mPa) was reported [Picart et al., 1998]. It has been observed that for many systems the value of yield stress depends on the time scale of the measurements. Setting all controversies aside, pragmatically it is advantageous to consider that in these systems, there are aggregates of different size, characterized by the dynamic interparticle interactions. For a given system, these interactions have specific strength, , and the aggregates have a characteristic relaxation time, τy. This model leads to the following relation:

[

{

σ y (ξ ) = σ y∞ 1 − exp −τ yξ

}]

u

.

;

for ξ = ω , γ

.

or ε

(7.39) where u = 0.2-1.0 characterizes polydispersity of the aggregates. Equation 7.39 was found to be easy to use, and the parameters computed from curve fitting of the experimental data, σapparent = σy + σtrue, agreed quite well with the independently determined values. Leonov [1994] introduced kinetics of interactions into his rheological equation of state. The new relation can describe systems with a dynamic yield stress, without resorting to a priori introducing the yield stress as a model parameter (as it has been done in earlier models).

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7.3.1.1.5 Time-dependent Flows Two types of flow are recognized: thixotropy, defined as a decrease of apparent viscosity under shear stress, followed by a gradual recovery when the stress is removed, and its opposite, anti-thixotropy, or rheopexy. Both are related to molecular or macroscopic changes in associations. In thixotropic liquids, the aggregate bonding must be weak enough to be broken by flowinduced hydrodynamic forces. If dispersion is fine, even slight interactions may produce thixotropic effects. When the dispersion coarsens, larger forces are required to engender the same effects. In the case of suspensions of anisometric particles, the interactions are particularly strong, while for spheres, the effect can be controlled by changing the type and concentration of ionic groups on the surface. Similarly, in polymer blends, the inter-domains interactions can be controlled by addition of a compatibilizer — its presence enhances the interphasial interactions. Breakup and recreation of the associated structure follow exponential decay kinetics. The simplest, single-exponential relation representing the thixotropic behavior is: (7.40) where t is the shearing time, and are values of shear viscosity at t = 0 and ∞, respectively, and τ∗ is the relaxation time of the system. Time dependency also enters into the consideration of the rheological response of any viscoelastic system. In the steady-state testing of such materials as molten polymers, the selected time scale should be sufficiently long for the system to reach equilibrium. Frequently, the required period, t > 104 sec, is comparable to that in thixotropic experiments. More direct distinctions between these two types of flow are the usual lack of elastic effects and larger strain values at equilibrium observed for the thixotropic materials (see Table 7.4). There is a correlation between these two phenomena, and theories of viscoelasticity based on thixotropic models have been formulated by Leonov [1972, 1994]. Inherent to the concept of thixotropy is the yield stress.

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Both the microstructural and continuum theories postulate that the material behaves as a Bingham body at stresses below a critical value. 7.3.1.1.6 Steady State Flows There are three types of melt behavior in a simple shear flow: dilatant, D (shear thickening); Newtonian, N, and pseudoplastic, P (shear thinning). Similarly, in an extensional flow, the liquids may be stress hardening (SH), Troutonian (T), or stress softening (SS). By definition, the response considered here is taken at sufficiently long times to ensure steady state, and the yield effect, Y, is subtracted. In consequence, within the experimental range of stress or deformation rate, several types of behavior may be observed. There exist a great variety of flow curves observed for different materials. 7.3.1.1.6.1 Pseudoplastic Flows For suspensions, the most common type is a pseudoplastic flow curve with the so called upper, ηo, and lower, η∞, Newtonian plateau [Cross, 1965, 1970, 1973]: (7.41) In this relation ao is the parameter describing how fast the viscosity changes between the two plateaux. In viscoelastic systems, the lower plateau is several orders of magnitude smaller than the upper one, η∞ << ηo, and it is frequently neglected.

Table 7.4.

2.

In Eq 7.42 τ is the relaxation time and m1 and m2 are polydispersity parameters, with a bound: n = 1 - m1 m2, where n is the power-law exponent in the relation:

(7.43)

Equation 7.42 well describes the flow behavior of polymeric systems, and it was found useful for polymer blends. It should be stressed that Equations 7.41-7.43 described the flow behavior of fluids without yield stress or thixotropicity. 7.3.1.1.6.2 Dilatant Flows Krieger and Choi [1984] studied the viscosity behavior of sterically stabilized PMMA spheres in silicone oil. In high viscosity oils, thixotropy and yield stress was observed. The former was well described by Eq 7.41. The magnitude of σy was found to depend on φ, the oil viscosity, and temperature. In most systems, the lower Newtonian plateau was observed for the reduced shear stress value: σr ≡ σ12d3 / RT > 3 (d is the

Thixotropic Systems

Yield stress defined by the conditions: y = 0 when σ12 < σy Initial slope of the flow curve: lim ∂ln ηapp / ∂ln y = 1 y→0

3. 4. 5.

(7.42)

Definitions of viscoelastic and thixotropic systems

No. Viscoelastic Systems 1.

Equation 7.41 resembles the one derived by Carreau [1972] for monodispersed polymer melts, which later was generalized for polydispersed systems [Utracki, 1984, 1989]:

Elastic effects are present After step-increase of y the shear stress, σ∗12 increases Smaller equilibrium strains than those for thixotropic systems

Critical stress defined by the conditions: y(t) = 0 when σ12 < σcrit(t) Initial slope of the flow curve: ∂ln ηapp / ∂ln y ≤ 1 y→0

Lack of elastic effects After step-increase of y the shear stress, σ∗12 decreases Larger equilibrium strains than those for viscoelastic systems

The Rheology of Polymer Alloys and Blends

sphere diameter, R is the gas constant, and T is the absolute temperature). However, when shear stress was further increased, dilatant behavior was observed. Dilatancy was found to depend on d, T, and silicone oil viscosity. The authors reported small and erratic normal stresses. To describe the above behavior, the following relation was derived [Utracki, 1989]:

(7.44) where ai are equation parameters. Excepting the assumptions that η∞ ≠ 0 and insertion of the middle square bracket on the rhs of Eq 7.44, the dependence is the same as Eq 7.42. Hoffman [1972, 1974], Strivens [1976], van de Ven [1984, 1985], Tomita et al. [1982, 1984], and Otsubo [1994] reported pseudoplastic/dilatant flow of concentrated suspensions of uniform and polydispersed spheres. A dramatic change in light diffraction pattern was systematically observed at the shear rate corresponding to the onset of dilatancy. Van de Ven and his collaborators demonstrated that, depending on concentration and shear rate, the distance between the sliding layers of uniform spheres in a parallel plate rheometer can vary by as much as 10%. The dilatant behavior of binary sphere suspensions in capillary flow was reported by Goto and Kuno [1982, 1984]. At constant loading, the dilatancy was observed only within a relatively narrow range of composition, 0.714 < x < 0.976, where x represents the fraction of larger spheres. Suspensions, even in Newtonian liquids, may show elasticity. Hinch and Leal [1972] derived relations expressing the particle stresses in dilute suspensions with small Peclet number, Pe = y/Dr ~ 1 (Dr is the rotary diffusion coefficient) and small aspect ratio. The origin of elastic effect lies in the anisometry of particles or their aggregates. Rotation of asymmetric entities provides a mechanism for energy storage, Brownian motion for its recovery. For suspensions of spheres, this mechanism does not exist,

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and N1 is expected to vanish. However, when at higher φ, the spherical particles aggregate into anisometric clusters the system may and it does show viscoelastic behavior. Indeed, large N1 [Kitano and Kataoka, 1981], Weissenberg rodclimbing [Newab and Mason, 1958], and large capillary entrance-exit pressure drops were reported [Goto et al., 1986]. On the other hand, owing to the yield stress, no extrudate swell was observed in suspensions of anisometric particles in Newtonian liquids [Roberts, 1973]. Theoretically, interparticle interactions contribute directly to the elastic stress component of spherical suspensions as well as by modification of the microstructure [Batchelor, 1977]: (7.45) where N is the number of particles, rij centerto-center separation of i and j particles with pairwise interparticle interaction force Fij. GadalaMaria [1979] reported that, for suspensions of PS spheres in silicone oil, N1 linearly increased with σ12. Newer theories have been discussed by Van Arsdale [1982], Bibbo et al. [1985], Brady [1993], Becraft and Metzner [1994], and many others. Dynamic-mechanical testing of suspensions is particularly suitable for studying systems with anisometric particles with well-defined structures [Ganani and Powell, 1985]. The authors studied the dynamic behavior of spheres in a Newtonian liquid. They reported that dynamic viscosity, η’, behaves similarly to the steady state viscosity, η, while the storage modulus G’ ≅ N1 ≅ 0. 7.3.1.1.7 Transient Effects In system where the structure changes with time upon imposition of stress, the transient effects are important. For example, semi-concentrated fiber suspensions in shear and extension show large transient peaks in the first and the second normal stress difference [Dinh and Armstrong, 1984; Bibbo et al., 1985]. It is interesting that the peaks appear at different times, first for N2, then for N1, and finally for σ12.

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7.3.1.2 Suspensions in Non-Newtonian Liquids This category mainly comprises filled and reinforced polymer melts. There are numerous reviews on the topic [Chaffey, 1983; Goettler, 1984; Metzner, 1985; Utracki, 1987, 1988; Utracki and Vu-Khanh, 1992]. There is particularly strong interest in flow of polymeric composites filled with anisometric, reinforcing, particles, with properties that strongly depend on the flow induced morphology and distribution of residual stresses. In the absence of interlayer slip, addition of a second phase leads to an increase of viscosity. The simplest way to treat the system is to consider that the relative viscosity as a function of the solids’ volume fraction, φ, particle aspect ratio and orientation. There is no difference between the flow of suspensions in Newtonian liquids and that of polymeric composites, when the focus is on the Newtonian behavior. The non-Newtonian behavior of suspensions originates either from the non-Newtonian behavior of the medium or from the presence of filler particles. The problems associated with this behavior can originate in inter-particle interactions (viz. yield stress), and orientation in flow [Leonov, 1990; Mutel and Kamal, 1991; Vincent and Agassant, 1991; Shikata and Pearson, 1994]. 7.3.1.3 Flow-induced Orientation The most efficient orientation fields are extensional. Using convergent and divergent flow one may control orientation of anisometric particles. Most of the work in this area has been done with fiber-filled materials but the effects are equally important for flow of neat semicrystalline polymer melts or liquid crystal polymers [Goettler and Shen, 1983; Goettler, 1984]. There is less information on the flow-induced orientation of platelets. In extensional flow, these particles are less susceptible to orientation. Two-stage orientation mechanism was observed in converging flow [Utracki, 1988].

The orientation profoundly affects flow, hence processability, as well as the product performance. It plays an important role in extrusion or injection molding where the anisometric particles may become oriented in a complex manner. Layered structures, weld lines, splice lines, swirls and surface blemishes are well known. Mold geometry (e.g., inserts) and transient effects make predictions difficult. It has been theoretically and experimentally shown that, when designing a mold for composites with anisometric particles, the principles developed for single-phase melts do not apply [Crowson and Folkes, 1980; Crowson et al., 1980, 1981; Folkes, 1982; Vincent and Agassant, 1986, 1991]. 7.3.1.3.1 Yield Stress Yield occurs as a result of structure formation due to physical crowding of particles, interparticle interactions, or steric-elastic effects of the medium. Depending on the stability of the structure, true or apparent (i.e., time-dependent) yield stress can be obtained. As a consequence, the magnitude of yield stress increases with aspect ratio of the particles, their rigidity, and concentration. The phenomenon is visible in steady state shear, dynamic, or extensional flow, especially at low rates of deformation, where the slope of the flow curve, logη vs. logy, is often: ∂logη / ∂logy = -1 (time-independent yield). Neglecting the yield stress may have serious consequences on interpretation of elasticity. Yield stress and plug flow are interrelated. The viscous loss energy is dissipated in a relatively small volume of material, where the concentration of solids differs from average. This may lead to excessive shear heating (effects as large as ∆T ≥ 80°C have been observed), degradation of polymeric matrix, strong change of skin morphology during polymer blends extrusion, as well as to attrition of anisometric particles, fibers or flakes. Thus, the skin layer may not only have different concentration, but different chemical and physical composition as well. At high flow rates this situation may lead to slip at the wall.

The Rheology of Polymer Alloys and Blends

In capillary flow, the slip velocity at the wall, s, can be calculated from [Reiner, 1930, 1931]:

(7.46)

where the first expression on the rhs of Eq 7.46 is the well-known Rabinowitsch correction, and the second expression provides for the contribution from the slip. Here s is the slip velocity, R is the radius of the capillary, and si are parameters. Experimentally, it was observed that the slip velocity depends on the difference of shear stresses, σ12 - σy. Exponent values as large as s1 = 6.3 was determined for rigid PVC. Slip may occur in any large strain flow, in capillary or between cone-and-plate or parallel plates [Kalyon et al., 1993, 1998]. Further consequences of the yield stress [i.e., the plug flow] are: (i) a drastic reduction of the extrudate swell, B ≡ d/do (d is diameter of the extrudate, do that of the die) [see, e.g., Crowson and Folkes, 1980; Utracki et al., 1984], and (ii) significant increase of the entrance-exit pressure drop, Pe (also known as Bagley correction). For single-phase fluids, these parameters have been related to elasticity by molecular mechanisms [Tanner, 1970; Cogswell, 1972; Laun and Schuch, 1989]. However, in multiphase systems, both B and Pe depend primarily on the inter-domain interactions and morphology, not on deformation of the macromolecular coils. Thus, in multiphase systems [i.e., blends, filled systems, or composites], only direct measures of elasticity, such as that of N1, N2 or G’ should be used. It is customary to plot the measure of the elastic component versus that of the shear components, viz., N1 vs. σ12, or G’ vs. G”, etc. For rheologically-simple systems, the relationships are independent of temperature, but for the multiphasial ones they are not — the viscoelastic time-temperature principle does not hold.

469

More recently, a new, viscoelastic-plastic model for suspension of small particles in polymer melts was proposed [Sobhanie et al., 1997]. The basic assumption is that the total stress is divided into that in the matrix and immersed in it network of interacting particles. Consequently, the model leads to non-linear viscoelastic relations with yield function. The latter is defined in terms of structure rupture and restoration. Derived steady state and dynamic functions were compared with the experimental data. 7.3.1.3.2 Extensional Flows The yield stress also occurs in extensional flows [Kamal et al., 1984; Utracki, 1988]. Yield stress is apparent in two related dependencies: (i) as a vertical displacement in the stress growth function at decreasing strain rates, contrasting with the normal linear viscoelastic behavior of single-phase polymeric melts, and (ii) as a deviation from the . theory that In qualitative agreement with the von Mises criterion σ11,y /σ12,y = 1.3 - 2.0 was reported [Utracki, 1984]. The Trouton ratio: (7.47) was found to decrease by half, as the concentration of glass beads in SAN increased (φ ≤ 0.37) [Martischius, 1982]. It was argued that in extensional flow, only the liquid undergoes deformation, thus both the extensional strains and viscosities should be corrected for the “diluting” effect of the filler volume [Nicodemo et al., 1975]. The relative extensional viscosity of suspensions in a power-law liquid can be expressed as [Goddard, 1978]:

(7.48) Equation 7.48 described well the extensional viscosity behavior of a PE/mica system. However,

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L. A. Utracki and M. R. Kamal

the yield stress had to be subtracted using Casson’s equation [Utracki and Lara, 1984]. 7.3.2

ponents. The plot of the emulsion viscosity as a function of the dimensionless viscosity ratio, Λ, is shown in Figure 7.8.

Emulsion Rheology

The liquid-in-liquid systems can be divided into three categories: those in which both liquids are Newtonian, those in which both phases are viscoelastic, and systems comprising one Newtonian and one viscoelastic liquid. The first of these categories covers emulsions, E, the second polymer blends, B, and the third undefined class of systems is usually used as models, M, to gain an insight into the effects of elasticity on the flow and morphology. Some polymer blends may also be classified as M. Several reviews on emulsion rheology were published [Sherman, 1963, 1968; Barry, 1977; Nielsen, 1977; Utracki, 1988; Pal, 1996]. Emulsions containing high volume fraction of the dispersed phase, φ > 0.74, have also been reviewed [Cameron and Sherrington, 1996]. Emulsions as models for flow of polymer blends and alloys have been discussed as well [Utracki, 1989]. 7.3.2.1 Newtonian Flow Einstein’s treatment of suspensions was extended to emulsions by Taylor [1932, 1934] who derived the following expression for the relative viscosity of emulsions: ηr = 1 + (5/2)σ [λ + 2/5]/[λ + 1] ∴ [η]e = (5/2) [λ + 2/5]/[λ + 1] where λ ≡ ηd/ηµ

(7.49)

Oldroyd [1953, 1955] modified this theory by incorporating effects of the interface:

Figure 7.8. Intrinsic viscosity of emulsion vs. the viscosity ratio (defined in the Figure) [Oldroyd, 1953, 1955].

Note that the upper bound of the emulsion intrinsic viscosity is the Einstein’s value for hard sphere suspension [η] = 2.5. This limit is observed for Λ > 100. For Λ = 1 (solutions), [η] = 1.75. The lower values are expected for emulsions of low viscosity liquids in highly viscous one — for Λ < 0.01, [η] ≈ 1. Equation 7.50 was found valid in a wide range of 1.3 < Λ < ∞. Oldroyd [1953, 1955], Choi and Schowalter [1975], Oosterbroek et al. [1980, 1981] and many others considered the interphase between the dispersed phase and the matrix liquid to be a physical, three-dimensional entity endowed with its own specific rheological properties. These considerations led to calculations of two relaxation times for Newtonian emulsions [Choi and Schowalter, 1975]:

(7.50)

In Equation 7.50, the interfacial viscosity, ηi, is expressed in terms of the interfacial shear (subscript Si) and extensional (subscript Ei) com-

(7.51)

The Rheology of Polymer Alloys and Blends

For the relative viscosity of emulsions, in the absence of deformation and coalescence, Eqs 7.247.30 may also be used, provided that the intrinsic viscosity is calculated from Eq 7.50 and that the maximum packing volume fraction is treated as an adjustable parameter, dependent on the interphase. This pragmatic approach has been successfully used to describe [η] vs. φ variation for such complex systems as industrial lattices (at various stages of conversion), plastisols and organosols. Industrial emulsions are usually prepared as concentrated systems, containing φm ≤ 0.94. Owing to interface interactions and deformability of droplets, these systems behave rather like elastic, soft solids without any sign of Newtonian behavior. Between the highly concentrated and dilute regions, there is a wide zone of structural change reflected in a spectrum of non-Newtonian behavior. 7.3.2.2 Emulsion Microrheology 7.3.2.2.1 Drop Deformability When a neutrally buoyant, initially spherical droplet is suspended in another liquid and subjected to shear or extensional stress, it deforms and then breaks up into smaller droplets. Taylor [1932, 1934] extended the work of Einstein [1906, 1911] on dilute suspension of solid spheres in a Newtonian liquid to dispersion of single Newtonian liquid droplet in another Newtonian liquid, subjected to a well-defined deformational field. Taylor noted that at low deformation rates in both uniform shear and planar hyperbolic fields, the sphere deforms into a spheroid (Figure 7.9). At low stress in steady uniform shearing flow, the deformation can be expressed by means of three dimensionless parameters: the viscosity ratio, the capillarity number, and the reduced time:

λ ≡ η1 / η2 ;

κ ≡ σ ij d / ν12 ;

and t * = γ / κ

(7.52) where: η1 and η2 are viscosities of the dispersed and the matrix phase, σij is the stress (either in

471

Figure 7.9. Deformation of drops in shear (a) and extensional (b) flow field [Taylor, 1932, 1934].

shear ij = 12, or in extension ij = 11), d is the drop diameter, ν12 is the interfacial tension coefficient between two phases, and γ is the generated strain. During shear or uniaxial extensional flow, the initially spherical drop deforms into a prolate ellipsoid with the long axis, a1, and two orthogonal short axes, a2. It is convenient to define the drop deformability parameter, D, as:

(7.53) where γ and ε are shear and uniaxial extensional strains, respectively. According to Taylor, the equilibrium deformability of drops is a complex function, which at two limits has simple solution. On the one hand, at low stresses, when the interfacial tension effects dominate the viscous ones (low value of λ), the deformability D and the orientation

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angle α (see Figure 7.9) of the droplet can be expressed as:

{

D(t ) = Do exp −[80(λ + 1) / (2 λ + 3)(19λ + 16)](

(tν12 / η2 d )} = = Do exp {−t / τ ret }

and

α=π/4

(7.54a)

Since, for λ = 0 to ∞, the quantity in the square bracket ranges from 1.00 to 1.18, the drop deformability D ≅ 0.55κ. Thus, a small deformation of Newtonian drops in Newtonian matrix varies linearly with the capillarity number. This proportionality was indeed demonstrated in Couette-type rheometer for a series of corn syrup/silicon oil emulsions [Elemans, 1989]. On the other hand, when the interfacial tension is negligibly small in comparison to viscosity (high value of λ): D=5κ/8

and

α=π/2

(7.54b)

Taylor predicted that droplet breakup will occur at D ≥ Dcrit = 0.5. Cox [1969] extended Taylor’s theory to systems with the full range of viscosity ratios:

where: τ ret = ηeq d / ν12 ;

ηeq = η2 [(2 λ + 3)(19λ + 16) / 80(λ + 1)] (7.56) where t is the retraction time and the relaxation time, τret, is expressed as a ratio of the equivalent viscosity, ηeq, divided by the interfacial tension coefficient scaled by drop diameter. Thus, knowing the time evolution of D and the material viscosity one can calculate the interfacial tension coefficient. In principle, Eq 7.56 is valid for Newtonian systems, but the method can also be used to characterize viscoelastic materials provided that the following two conditions are valid: (1) the retraction rate is sufficiently slow to ensure that materials behave as Newtonian, and (2) the elastic relaxation of the materials after deformation is faster than the ellipsoidal droplet retraction, i.e., a clear separation of the two mechanisms is possible [Luciani et al., 1996b].

D = (κ / 2)[(19λ + 16) / (16λ + 16)] / 1/ 2

[(19λκ / 40) + 1] 2

α = (π / 4) + (1 / 2)arctan{19λκ / 20} (7.55) The above relations are valid for Newtonian systems undergoing small, linear deformation, smaller than that, which would lead to breakup. Furthermore, experimental data indicate that it takes time to reach the equilibrium deformation. It is convenient to use the reduced time scale (see Eq 7.52) — to reach the equilibrium deformation t*d ≅ 25 is required [Elemans, 1989]. Taylor’s theory makes it also possible to predict the retraction of slightly deformed drops toward an equilibrium spherical form:

Figure 7.10. Drop deformability vs. shear rate [Silberberg and Kuhn, 1954].

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The Rheology of Polymer Alloys and Blends

The deformability of drops in a three-component Newtonian system (comprising 1.1% PS and 1.7% ethylcellulose dissolved in benzene) was studied by Silberberg and Kuhn [1952, 1954]. The authors reported that as the rate of shear increased, the spherical drops changed shape into prolate ellipsoids, with the long axis, a1, and two orthogonal short semiaxes, a2. The data are presented in Figure 7.10. The observed maximum on the a1 curve may be associated with the rheological effect on solubility. Owing to elasticity of the interphase, the first normal stress difference and the relaxation time could be calculated as [Schowalter et al., 1968]:

Note that in shear for λ = 1, the critical capillary number κcr = 1, whereas for λ > 1, κcr increases with λ and becomes infinite for λ > 3.8. This means that the breakup of the dispersed phase in pure shear flow becomes impossible for λ > 3.8. This limitation does not exist in extensional flows. The deformation of dispersed drops in immiscible polymer blends with the viscosity ratio λ = 0.005-13 during extensional flow was studied by Delaby et al. [1994, 1995]. In the latter paper, the time-dependent drop deformation during a start-up flow at constant deformation rate was derived. The model is restricted to small drop deformations.

(7.57)

7.3.2.2.2 Drop Breakup From the point of view of the drop deformation and breakup there are four regions of the reduced capillarity numbers κ∗, both in shear and elongation:

It is convenient to express the capillarity number in its reduced form κ∗ ≡ κ / κcr, where the critical capillary number, κcr, is defined as the minimum capillarity number sufficient to cause breakup of the deformed drop. Many experimental studies have been carried out to establish dependency of κcr on λ. For simple shear and uniaxial extensional flow, De Bruijn [1989] found that droplets break most easily when 0.1 < λ < 1, but do not break for λ > 4:

(7.58)

for 0.1 > κ∗ for 0.1 < κ∗ < 1 for 1 < κ∗ < 2 for κ∗ > 2

droplets do not deform, droplets deform, but they do not break, droplets deform then split into two primary droplets, droplets deform into stable filaments.

When values of the capillarity number and the reduced time are within the region of drop breakup, the mechanism of breakup depends on the viscosity ratio, λ. In shear, four regions have been identified [Goldsmith and Mason, 1967]:

Parameters of Equation 7.58 are listed in Table 7.5 (see also Figure 7.34).

Table 7.5.

Flow Shear Elongational

Parameters of the critical capillarity number for drop burst in shear and extension in Newtonian systems [R. A. de Bruijn, 1989]. 1000c1 -506.0 -648.5

1000c2 -99.4 -24.42

1000c3 124.0 22.21

1000c4 -115.0 -0.56

1000c5 -611.0 -6.45

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(7.59) where Λ is the distortion wavelength, Ω(Λ, λ) is a function tabulated by Tomotika, and Ro is the initial fiber radius. The hydrodynamic instability is characterized by a maximum for the dominant wavelength Λm that leads to thread breakup. For 0.01 ≤ λ ≤ 10: (7.60) Figure 7.11. Effect of the viscosity ratio on the critical time to break [Huneault et al., 1994].

where bo = -2.588, b1 = -1.154, b2 = 0.03987, b3 = 0.0889, and b4 = 0.01154. The distortion amplitude α grows exponentially with time, t:

for 0.1 >> λ

α = αo exp {qt}

small droplets are shed from sigmoidal drops — tip spinning, for 0.1 < λ < 1 drop breaks into two principal and odd number of satellite droplets, for 1 < λ < 3.8 drop deforms into fiber, which then disintegrates into small droplets, for λ > 3.8 drops may deform, but they do not break.

7.3.2.2.2 Drop Fibrillation and Breakup In addition to the previously discussed drop breakup into two principal drops (and odd number of small satellite droplets) there is another mechanism for dispersing one liquid in another. This is based on the “capillarity instability principle” of long cylindrical bodies. For κ* > 2, drops deform affinely with the matrix into long fibers. When subsequently the deforming stress decreases, causing the reduced capillarity number to fall below two, κ* < 2, the fibers disintegrate under the influence of the interfacial tension. The problem was theoretically treated by Rayleigh [1879], Taylor [1932, 1934] and Tomotika [1935, 1936]. The latter author showed that the degree of instability can be described by the growth rate parameter of a sinusoidal distortion:

(7.61)

where αo is the distortion at t = 0. Assuming that the initial distortion is due only to thermal fluctuations, Kuhn [1953] estimated that: αo = (21 kBT) (8π3/2 ν12 )-1/2

(7.62)

where kB is Boltzmann constant and T is the absolute temperature. The thread breakup occurs when α = R ≅ 0.81Ro. The time required to reach this stage can be expressed as:

(7.63) Thus, tb* is an important parameter describing the breakup process for fibers subjected to lower stresses than those required for fibrillation, i.e., κ* < 2. In practice, one of the most serious obstacles for quantitative use of Timotika’s theory is estimation of the initial distortion, αo. The time corresponding to the complete breakup, tb, was measured [Grace, 1982; Elemans, 1989]. The data are presented in Figure 7.11. Numerically they can be expressed as: (7.64)

The Rheology of Polymer Alloys and Blends

7.3.2.2.3 Coalescence During mixing, the dispersed phase progressively breaks down until a minimum drop diameter is reached. As the drop diameter decreases, further breakup becomes increasingly difficult. For emulsions, the size of the smallest drop that can be broken can be calculated from Taylor’s theory, but experiments have shown that in most cases the equilibrium droplet size is larger than predicted. Furthermore, the deviation increases with concentration of the dispersed phase, φd - φo, where experimentally the smallest value for which the deviation occurs, φo ≈ 0.005 [Utracki and Shi, 1992]. Two types of coalescence must be recognized, the first being determined by equilibrium thermodynamics (e.g., liquid-liquid miscibility, interfacial tension coefficient, rheological conditions of the interphase, etc.), the second, dynamic one, being also affected by the rheology. In the following text, only the second type will be discussed. Utracki [1973] studied steady-state shear coagulation of PVC lattices for a wide range of variables. Assuming that the locus of coagulation is at the particle-matrix interface and that the rate of coagulation depends on the frequency of particle collisions, the critical time for coalescence was calculated as:

(7.65) where ao is a numerical parameter; E+ is the threshold energy of coagulation, No and N+ are the number of coagulating drops, respectively, initially and at t = to, Vx/V is the volume fraction of emulsion undergoing uniform shearing, and ρd is the drop density. Validity of these relations was confirmed on many lattices. In particular, tc ∝ 1/(y)2 and dtc / dφd < 0 (thus, coagulation rate increases with φd) were experimentally confirmed. Since the coagulation was assumed to be related to the projected area of the drop, d ∝ φd2/3 was implicitly assumed.

475

Fortelny et al. [1988, 1990] assumed that the Brownian motion is the principal driving force for coalescence in polymer blends. Applying Smoluchowski’s theory, the authors obtained:

(7.66) where So and S are the interface areas in the unit volume of the blend having monodispersed spherical particles, respectively before and after coalescence. The experimental data confirmed the linearity of (S/So)-3/2 as a function of t1/2. However, it is to be noted that Brownian motion affects particles that usually are below the size of dispersions in polymer blends. Furthermore, coalescence is supposed to be independent of the stress field intensity and magnitude of the interfacial energy. The efficiency of particle collisions due to Brownian motion leading to coalescence should also have been taken into account. Coagulation is a result of collision between two spherical drops of diameter, d, that approach each other with certain velocity gradient to a distance smaller than their radius. The coalescence can happen only if the liquid remaining between two flattened drops will be able to be removed sooner than the global velocity field will force the drops to separate. The instabilities in the layer of the entrapped film will break it when the separation between drops is smaller than the critical separation distance, hc ≅ 5 nm. The critical coalescence time for systems with mobile interface is expected to follow the relation: tc = 3κ[ln(d/4hc)]/4y 2 [Elmendorp, 1986]. The relation was derived for isolated pair of drops and as such it does not take into account the concentration effect. It predicts that, in shear, the coalescence of two isolated drops is proportional to the exerted stress, and it is easier for larger spheres with high surface tension energy. In accord with Eq 7.65, tc ∝ y -2 was also found. Experimentally, the coalescence probability was found to rapidly decrease with increasing κ and d. The work was further extended by Chesters [1991].

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7.3.2.3.1 Non-Newtonian Flows Only dilute emulsion or systems undergoing slow deformation show Newtonian, deformation-rate independent flow. As the concentration and deformation rate increase, the flow progressively changes into pseudoplastic. Since the rheological response is a reflection of the inner structure of the material, modifications of the emulsion morphology are expected to accompany such a change. 7.3.2.3.1.1 Morphology Even in the dilute region, individual emulsion droplets rarely exist. In most cases, droplets are polydispersed in size, forming doublets, chain structures, or aggregates. Two types of emulsion morphology can be distinguished: (A) formed by the shear field (e.g., skin-core structures developed during flows through long tubes); and (B) formed by particle-particle interaction. Knowledge of type-A structures is important for proper interpretation of flow phenomena. Their formation is influenced by the difference between ingredients’ flow behavior, as well as by the shear coagulation. Knowledge of type-B structures is the key for utilization of suspension rheology in processing. Since the effective volume fraction of dispersed particles increases with increase of association, the relative viscosity of the emulsion is strongly affected. By contrast with polymer blends (see Part 7.5.1.), emulsions are prepared by carefully designing the interface system and by sequential addition of ingredients. Both elements are essential when 96 vol% of one liquid must be dispersed in 4 vol% of another. If, due to interactions of emulsifiers, the continuous phase becomes viscoelastic, the emulsion has high consistency or a ‘body.’ There is gradual passage of structures, from rotating doublets in dilute systems, to entrapment of the dispersed phase in a continuous network of interacting interfaces. Consequently, emulsions can show Newtonian character as well as a complex thixotropic and viscoelastic one [Nielsen, 1977].

7.3.2.3.1.2 Theoretical Treatment The theoretical treatment of two-phase flow was reviewed by Cox and Mason [1971], Leal [1980], and Barthes-Biesel [1988]. As indicated before, dispersions of one Newtonian liquid in another result in systems that are characterized by elasticity and relaxation times, e.g., viz. Eq 7.57. For dilute emulsions, with neither hydrodynamic interactions nor interfacial effects, Fröhlich and Sack [1946] developed the following timedependent constitutive equation:

(7.67) where G is the Hookian modulus of the elastic, dispersed spheres, while λi is the relaxation time of the emulsion. Thus, the theory considered viscoelasticity of dilute emulsions to originate in elastic deformability of the dispersed phase. Nearly a decade later, Oldroyd [1953, 1955] proposed a constitutive model similar to that of Fröhlich and Sack, valid at small deformations. The model considered low concentration of monodispersed drops of one Newtonian liquid in another. The interfacial tension and the viscoelastic properties of the interfacial film were incorporated by means of convected derivatives. The model provided the following relation for the complex modulus:

(7.68) where ν12 is the interfacial tension coefficient, R is the drop radius, Gi(ω) is the complex modulus, and subscripts i = m, d indicate matrix or disperse phase, respectively.

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Oldroyd’s model was extended by Palierne [990] to emulsions with polydispersed spherical drops. The model considered viscoelastic liquids. The concentration range was extended up to that at which drop-drop interactions start complicating the flow field. However, the drops must be spherical, undergoing small deformation, and the interfacial tension coefficient was considered constant and independent of stress and the interfacial area. The following relation was derived for the complex modulus: n n     G* = Gm* 1 + 3∑ φi Hi  / 1 − 2 ∑ φi Hi  i =1 i =1    

Hi =

( 4 / Ri′)(2Gm* + 5Gd* ) + (Gd* − Gm* )(16Gm* + 19Gd* ) ( 40 / Ri′)(Gm* + Gd* ) + (2Gd* + 3Gm* )(16Gm* + 19Gd* )

Gi* = Gi* (ω ) ; Ri′ = Ri / ν12 ; thus: Hi = Hi (ω ; R′)

(7.69) Here φi and Ri are, respectively, the volume fraction and the drop radius. The main feature of this model is the inclusion of a contribution from the interfacial tension to the viscoelastic properties and the inclusion of the effect of particles size polydispersity. For example, knowing Gi*(ω) of the two main components of the blend, one can predict the dynamic moduli of the emulsion (as well as dilute polymer blends) from the knowledge of the interfacial tension coefficient and distribution of drop size. Note that the theory is applicable to low strains, and to the concentration range where the yield stress is absent [Graebling and Muller, 990, 99; Graebling et al., 993]. Low viscosity mixtures of PDMS and PI, with λ = 0.55, 0.825, and 4.02 were studied at room temperature [Kitade et al., 997]. The dynamic data were analyzed using Eq 7.69. Good agreement was found. However, for a λ = 4.02 system the drops were insensitive to the flow field — they neither broke nor coalesced. Similar observations were reported for PDMS/PIB system [Vinckier et al., 996]. The latter authors also observed that agreement with Palierne’s model worsens for blends pre-sheared at higher shear rate, i.e., blends with finer drop dispersion.

In 990, Honerkamp and Weese published a seminal paper on the use of Tikhonov’s regularization for the determination of material functions. The developed method of data treatment was found particularly useful for the computation of the relaxation and retardation spectra [Elster et al., 99; Honerkamp and Weese, 993]. It has also been used to compute the sphere-size distribution of the dispersed phase in binary blends [Gleinser et al., 994a], as well as the ratio of the dispersed drop diameter divided by the interfacial tension coefficient, d/ν12 [Gleinser et al., 994b]. Friedrich et al. [995] modified Palierne’s Eq 7.69 by a continuous function:

[

]

∞

G* (ω ) = GM* (ω ) 1 + 3∫ H (ω , R′)u( R′) R′d ln R′ / −∞

[1 − 2∫

∞

−∞

H (ω , R′)u( R′) R′d ln R′

]

(7.70) where R’ = R/ν12, and u(R’) = ν12v(R) is scaled, volume-weighted distribution of sphere sizes. Using the Tikhonov’s regularization method, the distribution function, u(R’), can be computed. The experimental data (storage and loss shear moduli, G’ and G”, respectively, within six decades of frequency, and transmission electron microscopy, TEM) were determined for 2, 5, 0 and 20 wt% of PS in PMMA matrix. From the dynamic viscoelastic data of the neat components and the blends, the monomodal distribution of the distribution function u(R’) vs. log R’ = log(R/ν12) was computed and compared with data obtained from TEM. Excellent agreement was found for blends containing 2 and 5 wt%, fair for 0 wt% and poor for 20 wt% PS. In the latter case, TEM shows a bimodal distribution (possibly resulting from coalescence) whereas a monomodal distribution was obtained from the rheological data. The interfacial tension coefficient computed from these results varied from ν12 = 2.08 to 3.0 mN/m. The average value, ν12 = 2.5 mN/m at 90°C, is comparable with the literature data: ν12 = 0.8 to .8 mN/m at 200°C [Luciani et al., 996]. For infinitely diluted viscoelastic emulsions, the shear dependence of inherent viscosity was derived as [Barthes-Biesel and Acrivos, 973]:

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L. A. Utracki and M. R. Kamal

(7.71) where f(λ) > 0 is a rational function of λ. The relation predicts that in dilute emulsions subjected to small deformations, the ηinh should decrease with square of the shear stress. The effect of stress is moderated by interphase elasticity expressed as Eint. The theory was experimentally verified. Note the similarity of the first term on the rhs of Eq. 7.71 to the expression derived by Oldroyd for the intrinsic viscosity of emulsions, Eq 7.50. Accordingly, Eq 7.71 may be modified replacing λ by Λ as derived by Oldroyd. Semi-concentrated emulsions were examined theoretically and experimentally only within the linear viscoelastic region [Oosterbroek and Mellema, 1981; Oosterbroek et al., 1980, 1981; Eshuis and Mellema, 1984]. Recognizing that the interphase has a final thickness (sometimes the total volume of interphase is comparable to, or even exceeds, the volume of the dispersed phase) the authors postulated that the interphase should have two interfacial coefficients, ν’ and ν” facing the two principal polymer domains. Next, two models of the interphase were evaluated: (i) a two-dimensional viscoelastic film, and (ii) the interphase of final thickness. Both led to at least two relaxation times:

(7.72) The experimental data of dynamic testing in the kHz region for ionic emulsions could be equally well described using either model. The emulsion elasticity was found to originate in droplet deformation. For non-ionic emulsions, only one relaxation time was observed. The data were interpreted in terms of the second Oldroyd’s model, in which the interfacial tension is more important than the viscoelasticity of the interphase.

The steady state viscosities of both ionic and nonionic systems at the volume fraction φ ≤ 0.2 were found to follow Simha’s Eq 7.24. For concentrated emulsions and foams, Princen [1983, 1985] proposed a stress-strain theory based on a two-dimensional cell model. Consider a steady state shearing of such a system. Initially, at small values of strain, the stress increases linearly as in elastic body. As the strain value increases, the stress reaches its yield value, then at still higher deformation, it catastrophically drops to the negative values. The reason for the latter behavior is the creation of unstable cell structure that provides the recoil mechanism. The predicted dependencies for modulus and the yield stress were expressed as:

(7.73) is the concentrationwhere the function dependent, dimensionless contribution to stress per single drop. The theory was evaluated using a concentrated oil-in-water system. Since the early 1980’s, Princen’s work was continued by several other authors, e.g., by Reinelt [1993]. The latter author considered theoretical aspects of shearing three-dimensional, highly concentrated foams and emulsions. Initially, the structure is an assembly of interlocked tetrakaidecahedra (which have six square surfaces and eight hexagonal ones). An explicit relation for stress tensor up to the elastic limit was derived. When the elastic limit is exceeded, the stressstrain dependence is discontinuous, made of a series of increasing parts of the dependence, displaced with a period of γ = 23/2. 7.3.2.3.1.3 Experimental data Experimentally, there are three concentration regions of emulsion flow: (i) dilute for φ < 0.3, characterized by nearly Newtonian behavior; semiconcentrated at 0.3 < φ < φm with mainly pseudoplastic character; and concentrated at φm < φ < 1.0, showing solid-like properties with modulus and yield.

The Rheology of Polymer Alloys and Blends

479

The necessary condition for non-Newtonian flow to occur is droplet deformation, expressed either by the deformability parameter, D = (a - a2)/(a + a2), or by the aspect ratio, p = a/a2. Figure 7.2 illustrates the effect of shear rate, initial drop diameter, and the viscosity ratio on the droplet aspect ratio, p. For low and high values of λ, pseudoplastic dependence has been observed [Talstoguzov et al., 974]. The shear viscosity of polymer-thickened oilin-water emulsions was studied by Pal [992]. Addition of poly(ethylene glycol), PEG, made it possible to vary the matrix liquid viscosity. The flow was pseudoplastic, following Ellis dependence: (7.74)

where A and α are equation parameters. The zero-shear viscosity, ηo, was found to follow a modified Mooney dependence (see Eq 7.28): (7.75)

where Ki are equation parameters. It was observed that ηr of the thickened emulsions is lower than that for the emulsions without PEG. In liquid-liquid systems, upon increase of concentration of the dispersed phase, at a certain concentration suddenly the dispersed and continuous liquids exchange roles. This is known as a phase inversion. Salager et al. [983] and Minana-Perez et al. [986] reported two types of phase transition in ionic emulsions — in the first, viscosity goes through a minimum, whereas in the second it goes through a maximum. The first type of transition (normal) is associated with a decrease of the interfacial tension coefficient and formation of a micro-emulsion. The second (catastrophic) transition is associated with an inversion of unstable structure to a stable one. Recently, the extensional viscosity of nonNewtonian emulsions, ηE, at the volume fraction of the dispersed phase φ = 0.3-0.8, was measured

Figure 7.12. Effect of the shear rate (top), the initial drop diameter (middle), and the viscosity ratio (bottom) on the drop aspect ratio for the systems: water-gelatin-dextran (W-G-D, circles) and water-gelatin-polyvinyl alcohol (W-G-P, squares) [Talstoguzov et al., 974].

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L. A. Utracki and M. R. Kamal

using the opposed nozzles’ configuration. For more diluted emulsions, φ < 0.6, the elongational viscosity, ηE, was found to decrease with the rate of elongation, e, mimicking the flow curves in shear field, η vs. y. Furthermore, the Trouton rule, ηE ≅ 3η, was found to be reasonably obeyed. However, for more concentrated emulsions, φ • 0.7, owing to the presence of yield stress, ηE was found to depend on the test geometry, viz. nozzle diameter and their separation [Anklam et al., 994]. Since in more concentrated emulsions, the structure is engendered by close packing of interacting spheres, evidently its destruction must depend on the type of imposed deformation as well as on strain. In consequence, the lack of correlation between the shear flow and extensional flow data was to be expected. Compliance, J, of concentrated oil-in-water emulsions at 0.4 ” φ ” 0 7 was found to follow a two-retardation time process: (7.76)

where Jo and Ji are, respectively, instantaneous compliance and retarded values, all three decreasing with concentration of the dispersed phase [Gladwell et al., 986]. For a similar system, the shear viscosity was found to follow the power law model with yield [Pal et al., 986]. Owing to the presence of the yield stress, the flow of concentrated emulsion was found to be facilitated by superposition of 0 Hz oscillation on the steady state shear flow — up to 40 % energy saving was reported [Jezequel et al., 985]. More recently, the relative viscosity of emulsions was described in terms of scaling parameters [Pal, 997]. The author incorporated the ten principal variables into six dimensionless groups: λ, κ, reduced time, tr = t/(ηmd3/8kBT), relative density, ρr = ρd/ρm, Peclet number, Pe = ηmyd3/8kBT, and Reynolds number, Re = ρmyd2/4ηm. For the steady-state flow of well-stabilized emulsions, the author argued that the relative viscosity of emulsions should depend only on two parameters: volume

fraction of the dispersed phase and Re, i.e., ηr = f(φ, Re). At constant composition, the experimental data for coarse and fine oil-inwater emulsions plotted vs. the deformation rate, y, showed different dependencies (higher ηr for finer dispersion), but when plotted vs. Re, a single dependence was found. 7.3.3

Melt Flow of Block Copolymers

Block copolymers, BC, are macromolecular species in which long chains of one polymer are joined to long chains of another polymer. Thus, BC’s are made of at least two chemically different chains arranged linearly, in form of multi-branch stars, combs, etc. Linear block copolymers are the most common — diblock, AB, triblock, ABA, or multiblock, A(BA)n. Commercial BC’s are prepared from monomers that upon polymerization yield immiscible macromolecular blocks, one rigid and the other flexible, that separate into a two-phase system with ‘rigid’ and ‘soft’ domains. The concentration and molecular weights provide control of the size of the separated domains, thus morphology and the interconnection between the domains. The existence of a dispersed rigid phase in an elastomeric matrix is responsible for its ‘thermoplastic elastomer’ behavior. For symmetric block copolymers, Leibler [980] showed that a sufficient condition for microphase separation is (χΑΒN) = 0.5, where χΑΒ is the binary thermodynamic interaction parameter, and N is the degree of BC polymerization [Folkes, 985]. As in polymer blends, also in BC the phase transition is affected by flow. Theoretically it was predicted that a homogeneous melt can be obtained at T < UCST, provided that the stress field exceeds the critical value for the phase demixing, σ2 > σ2, crit [Lyngaae-Jørgensen, 989]. For most BC the phase diagram is characterized by the presence of an upper critical solution temperature, UCST, also known as an orderdisorder transition temperature or a microphase separation temperature. Below UCST the block copolymers phase separate, while above it, an isotropic melt is obtained. Owing to the chemical

The Rheology of Polymer Alloys and Blends

link between the blocks, during phase separation in BC, micro-domains instead of macroscopic phases are usually obtained. Furthermore, since the micro-morphology depends on the concentration as well as on the temperature, the phase diagram is complex, similar to those of metallic alloys. Under thermodynamic equilibrium conditions, depending on the composition, magnitude of the interaction parameter, and temperature, spherical, cylindrical, lamellar or some other structures are formed. There are three elements to BC morphology: domain size, domain shape, and the interfacial thickness — they lead to a wide variety of rheological responses [Inoue et al., 969; Meir, 969; Hashimoto et al., 980; Krause, 980]. For A-B block copolymers, the thickness of the interphase, ∆l, was theoretically derived by Helfand and Wasserman [976, 978, 980]: ∆l = 2

(β

2 A

48

Considering melt flow of BC, it is usually assumed that the test temperature is: UCST > T > Tgc, where Tgc stands for glass transition temperature of the continuous phase. However, at Tgc < T < Tgd (Tgd is Tg of the dispersed phase) the system behaves as a crosslinked rubber with strong viscoelastic character. At UCST > T > Tgd, the viscosity of BC is much greater than would be expected from its composition. The reason for this behavior is the need to deform the domain structure and pull filaments of one polymer through domains of the other. Viscosity increases with increase of the interaction parameter between the BC components in a similar way as an increase of the interfacial tension coefficient in concentrated emulsions causes viscosity to rise [Henderson and Williams, 979]. In shear, block copolymers exhibit time-dependent flow with yield stress [Liu et al., 983]:

)

+ β B2 / 2 χ AB

where: β i2 = ( ρoi bi2 ) / 6 ; and bi = < Ri2 > / Zi

(7.77) where bi is length of the Kuhn’s statistical segment, Zi is degree of polymerization, ρoi is density and is the radius of gyration of the block. For immiscible systems χΑΒ • 0, thus as the “antipathy” of the two types of blocks towards each other decreases and χAB → 0, the interphase thickness increases. Recently, another theoretical expression for ∆l was derived for symmetric diblock copolymer with NA = NB = N/2 in a lamellar morphology [Spontak and Zielinski, 993]: ∆l = ∆l∞ [1 − (8 ln 2) / ( χ AB N )]

−1/ 2

(7.78)

where N is the degree of BC polymerization and ∆l∞ is the interface thickness when N → ∞. The dependence should only be used for χABN • 20. The theory predicts that as χABN decreases, the thickness of the interphase increases — nearly three times more rapidly for block copolymers than for homopolymer blends.

(7.79) where β represents the relative residual viscous dissipation parameter, b and m are parameters originating from the structural breakdown and reformation of structure, while k is the loss-rate constant. The relation is capable of describing multiple phenomena: yield, upper and lower Newtonian plateaus, pseudoplasticity, stress growth and overshoot, thixotropy, hysteresis, etc. The multiplicity of rheological phenomena observed in BC is related to sensitivity of the melt structure to independent molecular and rheological variables. For example, for styrene-butadienestyrene (SBS), the activation energy of flow ∆Eη = 80 or 60 kJ/mol for compositions containing less or more that 3 vol% of styrene. The difference originates in the structure; it is dispersed below 3% and interconnected above [Arnold and Meier, 970].

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Block polymers, owing to the tendency for formation of regular structures tailored by molecular design, are ideal models for compatibilized, two-phase polymer blends or alloys. Blends do show similar rheological behavior, e.g., yield, pseudoplasticity, thixotropy, structural rearrangements, but since the morphology is more difficult to control, the interpretation of data could present serious difficulties.

7.4

Rheology of Miscible Blends

7.4.1

General Observations

Miscible polymer blends are less common that immiscible ones. The miscibility is usually confined to a specific range of independent variables, such as chain configuration, molecular weight, composition (viz. for alternate copolymers), temperature, pressure, etc. Nevertheless, Krause reported that 1680 two-, three-, or four-component polymeric mixtures were identified as miscible in 780 publications [Krause, 1980]. More detailed listing is provided in Appendix 2 of this Handbook. Unfortunately, the rheological studies of miscible systems are relatively rare. It is noteworthy that even in miscible polymers of similar molecular structure, viz. 1,4-polyisoprene with 1,2-polybutadiene, the time-temperature superposition fails. The polymers having glass transition temperatures separated by 60°C preserve their different dynamics in the blends [Kannan and Kornfield, 1994]. Thus, even miscible systems can be rheologically complex. The rheological behaviors of blends in the vicinity of the phase separation are of great fundamental importance. They will be discussed in Part 7.4.3. The PPE/PS mixtures are considered classical examples of miscible polymer blends. Within the accessible range of temperatures, single phase melts have been observed with the size of homogeneity below 20 nm. Dynamic flow behavior of PPE/PS blends, with the molecular weight ratio MW(PS)/MW(PPE) ≅ 1, was studied in a wide range of temperatures and compositions [Prest and Porter, 1972]. The authors assumed additivity

of the components’ free volume and characterized the blend flow behavior under iso-free volume conditions. Increasing the PPE content resulted in higher values of the storage and loss shear moduli, G’ and G”. The same blends, but with MW(PS)/ MW(PPE) = 100 were studied by Araujo and Stadler [1988]. Blends of atactic poly(methyl methacrylate) with poly(ethylene glycol), PMMA/PEG, were reported miscible [Colby, 1989]. Their rheology, PMMA/PEG = 50/50 and 80/20 at T = 160-210°C, was studied in a dynamic shear field [Booij and Palmen, 1992]. By contrast with homopolymers, the blends did not follow the time-temperature superposition. The deviation was particularly large at low temperatures. The reason for the deviation is most likely based on the different temperature dependence of the relaxation functions. The authors concluded that in miscible blends, the temperature dependence of the relaxation times of individual macromolecules depends on composition. This leads to different degrees of mutual entanglement and hence the rubber plateau moduli. By contrast with PPE/PS blends, those of PS with PVME are known to have the lower critical solution temperature, LCST, in the middle of the accessible to experimentation temperature range, LCST = 100-180°C (depending on composition and MW of the components). Rheology of these systems was studied within the miscible, immiscible as well as across the phase separation region. Within the miscible region, addition of PVME was reported to plasticize PS, thus shifting the terminal zone of G’ and G” to higher frequencies [Schneider and Brekner, 1985; Brekner et al., 1985; Yang et al., 1986]. It was also reported that the time-temperature superposition principle for the blends breaks down as the temperature approaches the glass transition temperature, Tg [Cavaille et al., 1987]. For the PS/PVME blends, with the molecular weight ratio MW(PS)/ MW(PVME) ≅ 40, separate relaxation times were found in the entanglement region [Stadler et al., 1988] (in homologous polymer blends having significantly different molecular weights, the relaxation spectra also show separate relaxation

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The Rheology of Polymer Alloys and Blends

times for the components). The latter authors also measured the dynamic shear moduli of the PS/PVME blends, across the phase separation temperature. The time-temperature superposition was obtained up to T = LCST + 40°C [Stadler et al., 988]. For the PS/PVME blends, with the molecular weight ratio MW(PS)/MW(PVME) ≅ 2 phase separation was seen to increase G’ but not G”, thus the time-temperature superposition breaks down [Ajji et al., 989; Ajji and Choplin, 99]. For similar blends, with MW(PS)/MW(PVME) ≅ 0.8 total breakdown of the time-temperature superposition principle was reported for the phase separated region. Large increases in both, G’ and G” were observed [Mani et al., 992]. Blends of tetramethyl polycarbonate, TMPC, with PS were reported miscible [Wisniewsky et al., 984]. Couchman [996] demonstrated that the concentration dependence of ηo at 230°C can be predicted from the characteristic parameters of the two components, assuming absence of specific interactions. In miscible blends, the flow behavior depends on free volume, entanglements, and specific interactions. From Doolittle’s equation, assuming additivity of the occupied volume and non-additivity of the free volume, Steller and Zuchowska [990] derived:

Figure 7.13. The experimental (points) and computed from Equation 7.80 relation between viscosity and composition for polypropylene blends with styrene-ethylene-butene-styrene block copolymer [Steller and Zuchowska, 990].

GNo =

∑G

o Niφ i

; ηo =

∑η

oiφ i

∑J

; Jeoηo2 =

o 2 eiηoiφ i

(7.8) where GoN is the plateau modulus, ηo is the zero-shear viscosity, and Joe is the recoverable shear compliance. This “single reptation model” neglects the thermodynamic interactions and constraint release. Viscoelasticity of miscible polymer blends was also analyzed by Tsenoglou [988]. The “double reptation model” resulted in the following mixing rules for the miscible blends: GNo

 2  o 1/ 2 =  (GNi ) φi   i =1 

2

∑

ηo = ηo1φ12 + ηo 2φ 22 + 4(GNo1GNo 2 )1/ 2 φ1φ 2 (7.80) where wi and Vi are, respectively, weight fraction and specific volume of i-th component, while Vo and Vf is the occupied and free volume, respectively. Since the parameter δ can be experimentally determined (from the temperature gradient of the viscosity and specific volume — see Eq 7.80) the dependence is fully predictive, as shown in Figure 7.3. The reptation model provides simple mixing rules for miscible systems [Doi and Edwards, 986]:

[(G

o N1

) (

/ ηo1 + GNo 2 / ηo 2

)]

−1

 Jeoηo2 = Jeo1ηo21φ12 + Jeo2ηo22φ 22 + 8φ1φ 2    1/ 4 −1 o  G o / G o 1/ 4  GN 2 / GNo1 N 1 N 2   +  J o η 2 1/ 2 o 2 1/ 2  η J e1 o1 e2 o2  

(

)

(

)

(

)

(

)

(7.82)

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L. A. Utracki and M. R. Kamal

Validity of Eqs 7.8 and 7.82 was examined for mixtures of entangled, nearly monodispersed blends of poly(ethylene-alt-propylene) with headto-head PP [Gell et al., 997]. The viscoelastic properties compared at constant distance from the glass transition temperature of each system were found to obey the time-temperature superposition principle. The data agreed better with the predictions of Eq 7.82 than Eq 7.8. However, for blends of linear and branched PE the relations 7.82 were found valid only when MW and rheological properties of the two components were similar [Groves et al., 996]. 7.4.2

Relaxation Spectrum and Linear Viscoelasticity

Since the Gross frequency relaxation spectrum can be computed from η’, i.e., from the loss modulus, G” = η’ω, the agreement between the computed and measured G’ values provides good means of verifying both the computational and experimental procedures. It has been found that Eqs 7.83 and 7.84 are useful to evaluate the rheological performance of systems that obey the linear viscoelastic principles. Owing to definition of the reduced relaxation spectrum, the integral: (7.85)

Thus, the coordinates of the maximum,

Substituting Eq 7.42 into the Gross frequency relaxation spectrum, HG, results in an expression:

(7.83) Thus, once the four parameters of Eq 7.42 are known, the relaxation spectrum, and then any linear viscoelastic function can be calculated. For example, the experimental data of the dynamic storage and loss shear moduli, respectively G’ and G”, or the linear viscoelastic stress growth function in shear or uniaxial elongation can be computed from the dependencies [Utracki and Schlund, 987]:

(7.84)

,

are related respectively to the system polydispersity and molecular weight. However, if the system is miscible, these functions are fully predictable from the composition, polydispersity and molecular weight of the components. Note that in miscible blends, the general relation between the relaxation spectrum of a mixture and its composition follows the third order blending rule: (7.86) The dependence can be significantly simplified when all fractions are either entangled or not (the situation that exists in most polymer blends): (7.87) Thus, for miscible polymer blends, the relaxation spectrum is a linear function of the relaxation spectra of the components and their weight fractions, wi, hence one may use rheological functions to detect miscibility/immiscibility of polymer blends. An example is presented in Figure 7.4 [Utracki and Schlund, 987]. Two principles can be used for the rheological methods of miscibility detection: . effect of polydispersity on the rheological functions, 2. effect of the inherent nature of the two-phase flow.

The Rheology of Polymer Alloys and Blends

485

The first principle makes it possible to draw conclusions about miscibility from: • Coordinates of the relaxation spectrum maximum [Utracki and Schlund, 1987] • Cross-point coordinates (Gx, ωx) [Zeichner and Patel, 1981] • Free volume gradient of viscosity: α = d(ln η)/df • Initial slope of the stress growth function: S ≡ dlnη∗Ε/dlnt, [Schlund and Utracki, 1987] • The power-law exponent n ≡ dln σ12/dlny ≅ 1 - S, etc. The second principle involves evaluation of: • Extrudate swell parameter, B = D/Do • Strain (form) recovery • Yield stress, etc. These effects, associated with immiscibility, will be discussed in a later part of this chapter. 7.4.3

Phase Separation and Flow

For most polymer blends, the phase diagram is characterized by the presence of the lower critical solution temperature, LCST. Thus, as the temperature increases, the miscible polymer blends may phase-separate. Theoretically, the miscibility region extends up to the binodal. However, as the system approaches the binodal, there is strong mutual interaction between the rheology and thermodynamics [Rangel-Nafaile et al., 1984; Larson, 1992]. The interaction between stress and composition in single-phase polymer solutions and blends is of growing theoretical and experimental concern. The flow-induced encapsulation has been known for a long time. Recently, the focus has shifted to miscible systems of non-uniform compositions placed in a non-uniform stress field. Two mechanisms have been proposed: the first postulating that long chains migrate to lower stress regions to decrease the elastic energy stored by deformation of the macromolecular coil [Metzner et al., 1979], the second assumes that long chains can support stress more than the short ones, creating imbalance of stresses and relative motion of the components [Doi and Onuki, 1992]. The latter theory makes it

Figure 7.14. Coordinates of the maximum of the Gross frequency relaxation function, HG, max (top) and ωmax (bottom), vs., respectively, the polydispersity index, Mw/Mn, and the zero shear viscosity, ηo, a measure of the molecular weight [Utracki and Schlund, 1987].

possible to calculate the concentration gradients in sheared blends. For example, in cone-and-plate geometry, the theory predicts migration of the high molecular weight fractions toward the center. Phenomenologically, the effect may be considered to originate from the hoop stress created in shearing of the larger polymeric chains that force them to migrate toward the center, engendering what is known as Weissenberg effect. The effect is related to the normal stress, σN = N1 + 2N2, and the osmotic pressure gradient, dπ/dφ, while the diffusion time is determined by the ratio (l /Dm), where l is the diffusion length scale and Dm is the mutual diffusion coefficient — the process is rather slow.

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L. A. Utracki and M. R. Kamal

The phase behavior of polymer blends under flow was reviewed by Kammer et al. [993]. 7.4.3.1 Influence of Thermodynamics on Rheology For linear viscoelastic functions near phase separation at low strains, Larson and Fredrickson [987] derived:

(7.88) where ∆ = 2[(χN)s - (χN)] is a measure of thermodynamic distance from the spinodal. The theory indicates that at the spinodal, the linear viscoelastic functions go to infinity, or in other words, the system becomes rheologically nonlinear. Numerically, near the spinodal, the theory predicts that the ratio Ψ2/Ψ ≅ -.35, instead of the usual Ψ2/Ψ = -0.05 to -0.20. All the data for diverse liquid systems indicate that when approaching phase separation, the viscosity should increase with the correlation length, but the rate of the increase and the absolute magnitude may vary from one system to another. The effect depends on the deformation rate and is more pronounced in high molecular weight systems. This prediction was found qualitatively valid for blends of low density polyethylene, LDPE, with linear low density polyethylene, LLDPE. At about 20 wt% of LDPE a sharp peak in the plot of a≅ d ln η /d ln Y (where Y = /(f + 0.07) with f being the free volume fraction) vs. composition. This behavior was associated with the phase separation in the blends [Utracki and Schlund, 987]. In contrast to the predictions of Eq 7.88, the effective viscosity at the spinodal was described by the fluidity additivity relation [Onuki, 994, 997]: .

1 / ηeff = φ1 / η1 + φ 2 / η2 ; N1 ∝ γ (φ1η2 / φ 2 ) at SD: φ1 / η1 ≈ φ 2 / η2 ; where φ1 = 1 − φ 2 (7.89)

Similarly, the first normal stress difference, N, also is predicted to be proportional to the shear stress. The phase co-continuity condition was derived from the equal shear stress principle between two phase-separating phases. The dependence is the same as empirically derived by Paul and Barlow — see Eq 7.6. Effect of phase separation on the dynamic shear flow of PS/PVME blends was frequently studied [Schneider and Brekner, 985; Brekner et al., 985; Yang et al., 986; Cavaille et al., 987; Stadler et al., 988; Ajji et al., 989; Ajji and Choplin, 99; Mani et al., 992]. The results seem to differ depending on the relative magnitude on the molecular weight, MW(PS)/MW(PVME). When the ratio was large, PVME acts as a plasticizer. When the ratio was about 2, a passage through the phase separation region affected the dynamic storage shear modulus, G’, but not the loss one, G”. Finally, for ratios less than one, it was reported that neither the storage nor loss dynamic shear modulus, G’ and G”, respectively, indicated any significant change near the phase separation region, but upon entering the phase separation region, both functions increased. By measuring the fluorescence intensity, the authors mapped the phase separation region. The timetemperature superposition principle was found to be valid only within the miscible blends region [Mani et al., 992]. The dynamic shear flow within the linear viscoelastic region was used to determine binodal and spinodal temperatures (Tb and Ts, respectively) in LCST-type blends [Vlassopoulos, 996]. The system of interest was PMMA (containing 2 wt% of ethyl acrylate) with 0-60 wt% of SMA. The demixing temperature was determined in temperature sweeps, by plotting log G’ vs. T — departure from a straight line was taken as Tb. Determination of Ts involved plotting [G’’(ω)]2/[TG’ (ω)]2/3 vs. /T. Again, linear extrapolation was used to determine Ts. The rheologically determined data were found to correspond reasonably well to those determined by turbidity and light scattering. The viscoelastic properties of model blends with UCST were studied in dynamic and steadystate shearing [Vlassopoulos et al., 997].

487

The Rheology of Polymer Alloys and Blends

Low molecular weight PS and poly(methyl phenyl siloxane), PMPS, was used — the neat resins showed Newtonian behavior. The equilibrium phase diagram was determined by optical means. Within the miscible region, blends viscosity followed the log-additivity rule, provided that the concentration was corrected for difference in the surfaces [Mertsch and Wolf, 994]:  ςw w  ln η = w1 ln η1 + w2 ln η2 + ln(η1 / η2 ) 1 2  1 + ςw1   where: ς ( Bondi ) =

S1 / S2 V1 / V2

−1

(7.90)

Bondi

where wi is the weight fraction of polymer i. The correction factor, ς, can either be treated as a fitting parameter (ς = -0.54 was determined for PS/PMPS system) or it can be calculated using Bondi’s values of the surface, Si, and volume, Vi, for each component (ς = -0.5 was calculated). The phase separation resulted in rheologically complex behavior. However, the predicted by Eq 7.88 rapid increases of the rheological functions near SD were not observed. 7.4.3.2 Influence of Rheology on Thermodynamics The response of heterogeneous systems to a stress field allows them to be placed in two categories: (i) those in which stress induces irreversible changes (e.g., precipitation, denaturation of protein, crystallization, etc.) and (ii) those in which the changes are reversible. The classification is not perfect, as the type and magnitude of stress field can be crucial, but it provides a guide: in most cases, miscibility in systems (i) is reduced by stress, while in systems (ii) it is increased. In other words, if a system can be irreversibly modified by rheological means, its solubility will be reduced. An excellent review on phase transition in shear flow was recently published [Onuki, 997]. Microrheology indicates that drops burst when the capillarity parameter κ ≈  (see Eq 7.52). Thus, in shear the equilibrium drop diameter d ∝ ν2/σ2 — the higher is the shear stress

in the matrix or the lower is the interfacial tension coefficient, the smaller is the drop size. In other words, it is natural to expect that shearing improves dispersion. When the drop diameter becomes comparable to the radius of gyration of the macromolecules, miscibility is achieved [Silberberg and Kuhn, 952; Wolf, 980, 984]. The above argument is certainly correct considering infinitely diluted systems. In the practical case at finite concentration, drop coalescence may limit the dispersion process. However, when shearing takes place near the critical point, phase separation can only occur when the rate of shear is smaller than /τc, where τc, is the thermodynamic relaxation time for concentration fluctuations. Strain compatibilization at low, steady state stress was considered by Lyngaae-Jørgensen [985]: 2 σ cr ≅ ao T (Ts − T ) ; T ≤ Ts

(7.9)

where ao is a material parameter. Subjecting block copolymers to above-critical stresses at low deformation rates, made it possible to change Ts by ∆T = 29°C. For systems with lower critical solution temperature, LCST, the spinodal was shifted to higher temperatures. The values ao = 0.26 and 0.53 (kPa/K)2 were calculated, for block copolymers and poly(styrene-co-acrylonitrile)/ poly(methyl methacrylate) blends, respectively. Equation 7.9 is in qualitative agreement with experimental observations that in polystyrene/ polybutadiene/dioctylphthalate systems the critical point shifted as: ∆Tc(y) ≈ aoTcy /2 where ao is a numerical factor [Hashimoto et al., 990]. A new thermodynamic theory of strain demixing was proposed [Horst and Wolf, 99, 992, 994]. The authors postulated that the Gibbs free energy of mixing for flowing blends can be expressed as a sum of the equilibrium thermodynamic free energy of mixing, ∆Gm, and the flow-induced stored energy term ∆ES:  n  . 2n   ∆G . = ∆Gm − ∆Es =  xi Vi  Jeo  ηo γ  γ    i =1 

∑

(7.92)

where the averaged values of the zero-shear viscosity and the steady-state shear compliance

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L. A. Utracki and M. R. Kamal

can be calculated from, respectively:

ηo Jeo

 n  = wiηi1/ 3.4   i =1 

3.4

∑

ηo

4.4 / 3.4

(7.93)

n

=

∑

wi Jeioηi4.4 / 3.4

i =1

The theory was found to predict complex behavior near the phase separation conditions. As the rate of shear increases, first, the system undergoes homogenization, then demixing, followed by another homogenization and demixing. At high rates of shear, the system should behave similarly as in a quiescent state. These predictions were found to be in qualitative agreement with experimental data, e.g., for blends of ethylenevinylacetate copolymer with chlorinated polyethylene, EVAc/CPE, or polystyrene with poly(vinyl methyl ether), PS/PVME [Hindawi et al., 992; Fernandez et al., 993, 995]. The first observation of shear-induced increase of the LCST was reported for PS/PVME by Mazich and Carr [983]. The authors concluded that shear stress can enhance miscibility by 2-7°C. Larger effects, ∆T ” 2°C, were reported for the same system in hyperbolic flow [Katsaros et al., 986]. In a planar extensional flow at e = 0.02 - 26 s- the phase separated PS/PVME was homogenized at temperatures 3 to 6°C above Ts. The critical parameter of homogenization was found to be the extensional strain, εc = etc = 44±4 where tc is the critical time to achieve miscibility at various levels of φ, T and e. The constancy of εc indicates that the main mechanism of flow induced miscibility is related to deformation; after cessation of flow, the deformation dissipates and the homogenized blend phase separates within 20 to 70 s. By contrast, large stresses can cause demixing in colloidal (e.g., denaturation of proteins) and polymeric systems. In the latter case, precipitation from poor solvent solution, shear crystallization, and stress-related phase separation, are all known. For example, PS/PVME under planar stresses at σ < 0 MPa shows the previously discussed strain compatibilization,

whereas at σ • 30 MPa it exhibits stress demixing [Katsaros et al., 986]. The demixing may be related to differences in the rheological behavior of the two blend components. For true shifting of the critical point to occur, the suppression of long-range concentration fluctuation must be anisotropic. This has not been observed in PS/PVME blends. Using neutron scattering, it was demonstrated that shearing suppresses the fluctuations only parallel to the flow, leaving the concentration gradients in other directions unchanged [Nakatani et al., 990]. Viscoelastic effects (caused by the presence of a high-MW polymer) during the early stage of SD were recently discussed [Clarke et al., 997]. The data were verified using PVME blended with PS having MW = ,60 kg/mol. Good agreement was observed. The relation between rheology and morphology during late stages of SD in PS/PVME blends was investigated by means of several techniques [Polios et al., 997]. The results were interpreted using Doi-Ohta [99] theory. Shear-induced mixing was reported for polystyrene/polyisobutylene, PS/PIB, blends [Wu et al., 992]. Optical measurements indicated that shearing within the miscible blend region did not cause demixing, while shearing within the two-phase region reduced turbidity. The latter observation was interpreted as most probably resulting from the shear-induced mixing of the blends. Blends of PI with PB were dynamically sheared at large amplitude (γo = 0.8) and frequency ω = 0.63 and 6.3 rad/s [Matsuzaka et al., 997]. After a temperature jump, the spinodal decomposition (SD) was in-situ observed at the lower frequency, but not at the higher. In the latter case, after stopping the oscillation, a modified SD pattern emerged. The authors postulated that the dynamic flow induced a structure in miscible system, quite different from that that exists in the non-sheared specimens.

The Rheology of Polymer Alloys and Blends

7.5

Rheology of Immiscible Blends

7.5.1

Rheological Equation of State

With the knowledge of the flow behavior of simpler systems, viz. suspensions, emulsions, block copolymers, as well as that of the mutual interactions between the rheology and thermodynamics near the phase separation, one may consider the flow of more complex systems where all these elements may play a role. Evidently, any constitutive equation that may attempt to describe flow of immiscible polymer blends should combine three elements: the stress-induced effects on the concentration gradient; an orientation function; and the stress-strain description of the systems, including the flow-generated morphology. Such a comprehensive description still remains to be formulated. The first steps toward such a theory of blend flow behavior were proposed by Helfand and Fredrickson [1989], then by Doi and Onuki [1992]. A greatly simplified constitutive equation for immiscible 1:1 mixture of two Newtonian fluids having the same viscosity and density was also derived [Doi and Ohta, 1991]. The derivation considered time evolution of the area and orientation of the interface in flow, as well as the interfacial tension effects. The relation predicted scaling behavior for the stress and the velocity gradient tensors:

489

recovery of the overshoot (or that of undershoot for step-down experiments), as well as the shear thinning, were not predicted by it. However, the theory was able to predict the extra stress arising from the interfacial tension. Also the transients (η and N1) at the startup of steady-state flow agreed qualitatively with the theory. Doi-Ohta’s theory was also compared to the experimental data of semi-concentrated mixtures of PIB in PDMS [Vinckier et al., 1997]. The theory described reasonably well the transient effects at the startup of steady state shearing. The scaling laws were also obeyed by these slightly viscoelastic blends. Following Doi and Ohta’s work, a more general theory was derived for immiscible polymer blends by Lee and Park [1994]. A constitutive equation for immiscible blends was proposed. The model and the implied blending laws were verified by comparison with dynamic shear data of PS/LLDPE blends in oscillatory shear flow. This new approach considered the influence of morphology in determining the rheological behavior in a given flow field. Thus, instead of formulating a single droplet problem, for example, as microrheology does for the dilute dispersions, the authors considered the complex interfaces formed between two phases of the immiscible fluids created by deformations, breakup, and the coalescence of drops (caused by flow and interfacial tension). A semi-phenomenological kinetic equation was derived that described the time evolution of the interfacial area per unit volume, Q, and its anisotropy in a given flow field, qij:

(7.94) Experimental verification of Eqs 7.94 indicated that the scaling relationships are valid, but the shape of experimental transient stress curves, after step-change of shear rate, did not agree with Doi-Ohta’s theory [Takahashi et al., 1994]. Similar conclusions were reported for PA-66 blends with 25 wt% PET [Guenther and Baird, 1996]. For steady shear flow the agreement was poor, even when the strain-rate dependence of the component viscosities was incorporated. Similarly, the

(7.95) where ni denotes the unit normal vector to the interfaces, V the total system volume, and dS an interface element. The time evolution of Q and qij is affected by the flow that deforms the interface to an anisotropic state:

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L. A. Utracki and M. R. Kamal

dQ / dt = − dij qij dqij / dt = − qik dkj − q jk dik + (2 / 3)dij dlm qlm − − (Q / 3)( dij + d ji ) + (qlm dlm ) / Qqij

(7.96) where dij = ∂ui/∂xj is the macroscopic velocity gradient tensor. For a mixture of fluids with equal viscosity, the stress tensor may be expressed as:

where the dimensionless parameters, ci, are respectively: the total relaxation, the size relaxation, and the breakup and shape relaxation. They all depend on the volume fraction φ. The macroscopic stress tensor for the two-phase fluid can be expressed as: σij = (pressure term) + ηm(dij+ dji) + (viscosity ratio term) + (morphologydependent term), or respectively:

σ ij = − Pδ ij +

[1 + 3φ (λ − 1) / 5(λ + 1)]ηm (dij + d ji ) − ν12 qij

σ ij = ηm ( dij + d ji ) − ν12 qij − Pdij ; dij ≡ dui / dx j

(7.0)

(7.97)

To complete the constitutive equation, contributions originating from flow must be incorporated. These are expressed as:

where ηm is the matrix viscosity and ν12 is the interfacial tension coefficient. In Eq 7.97, the excess shear stress is proportional to the spatial anisotropy of the interfaces, qij, and the structure of the interface is distorted by the competition between flow and interfacial tension. The constitutive equation can be used not only for arbitrary volume fractions but also for arbitrary flow fields. It is advantageous to consider that the time evolution of Q and qij, originates in the external flow as well as in the interfacial tension: dQ / dt = (dQ / dt)flow + (dQ / dt)interf.tens. dqij / dt = (dqij / dt)flow + (dqij / dt)interf.tens. (7.98) For concentrated systems, dimensional analysis of the retraction caused by the interfacial tension makes it possible to express the second terms in Eq 7.98 as: dQ / dt = − c1c2 (ν12 / ηm )Q 2 − c1c3 (ν12 / ηm )qij qij (7.99) dqij / dt = − c1 (ν12 / ηm )Qqij − c1c3 (ν12 / ηm )(qlm qlm / Q)qij (7.00)

dQ / dt = − dij qij

(7.02)

dqij / dt = − qik dkj − q jk dki + (2 / 3)δ ij dlm qlm − (Q / 3)( dij + d ji ) + (qlm dlm / Q)qij (7.03) Substitution of Eqs 7.99, 7.00, 7.02 and 7.03 into the two dependencies in Eq 7.98 provides two relations that, combined with Eq 7.0, form the rheological equation of state. Note that at t = 0: Q = Qo and qij = qijo. For dynamic oscillatory flow, the relationships between the complex shear moduli, G* = iωη* can be written as:

(7.04) where Gb*, Gm*, and Gi* are the complex moduli of the blend, the matrix, and the dispersed phase, respectively. The term Gint* is the complex modulus attributed to the interfaces. Imposition of sinusoidally varying strain, γ(ω) = γosin(ωt) results in frequency dependent stress at the interface: ν12qij. The latter also sinusoidally changes with the stress amplitude, δint, and is out of phase with the strain. The interfacial stress amplitude σint, and its

The Rheology of Polymer Alloys and Blends

phase lag by δint can be calculated by solving Eqs 7.02-7.03 assuming the initial values of Qo, qijo as well as of the parameters, ci. The interfacial moduli, Gint*, G’int, and G”int can be expressed as: (7.05) (7.06) (7.07) where γo is the strain amplitude. The effects of shear flow on PS/LLDPE morphology were investigated by observing the structure of quenched samples under the scanning electron microscope, SEM. Predictions based on the constitutive equations were compared with observations from the dynamic shear experiments at 200°C (see Figure 7.5). The frequency variations of Gb’, Gb”, and ηb* were found to be in good agreement with computations based on Eqs 7.0-7.03. However, to get such agreement, seven parameters (viz. ν2, φ, initial value of the anisotropy parameter, qijo, initial size of the dispersion, and three dimensionless equation parameters) were required [Lee and Park, 994].

49

It should be noted that the Doi and Ohta theory predicts only an enhancement of viscosity, the so called “emulsion-like behavior” that results in positive deviation from the log-additivity rule, PDB. However, the theory does not have a mechanism that may generate an opposite behavior that may result in a negative deviation from the logadditivity rule, NDB. The latter deviation has been reported for the viscosity vs. concentration dependencies of PET/PA-66 blends [Utracki et al., 982]. The NDB deviation was introduced into the viscosity-concentration dependence of immiscible polymer blends in the form of interlayer slip caused by steady-state shearing at large strains that modify the morphology [Utracki, 99]. More recently, a complete set of governing relationships was derived from the requirements of the compatibility of dynamics and thermodynamics [Grmela and Aït-Kadi, 994, 998; Grmela et al., 998, 200]. The authors developed a set of equations governing the time evolution of the functions Q and qij (see Eqs 7.95), as well as the extra stress tensor expressed in their terms. The rheological and morphological behavior was expressed as controlled by two potentials: thermodynamic and dissipative. Under specific conditions for these potentials, Lee and Park formalism can be recovered.

Figure 7.15. Comparison of model predictions with experimental results for (a) 0 wt%, and (b) 30 wt% of PS in LLDPE [after: Lee and Park, 994].

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L. A. Utracki and M. R. Kamal

Lacroix et al. [997, 998] attempted to evaluate the extend the three approaches (that of Palierne, Lee and Park, and Grmela and AïtKadi) are capable to describe experimental data of different types of polymer blends, viz. PS/PE, PETG/EVAc, PP/EVAc, and PP with EVAc and poly(ethylmethacrylate) (PEMA). Since EVAc is miscible with PEMA, the latter blend is also a two-phase blend with PP being the matrix. All blends were prepared within the concentration range that assured dispersed morphology. In the first publication, the Palierne model was found to well describe the linear viscoelastic behavior, whereas that of Lee and Park was found useful for describing the rheological behavior under large strains. In the second cited publication, the overshoot at the startup of shearing was well described using either Lee and Park, or Grmela and Aït-Kadi models. 7.5.2

Morphology of Immiscible Blends

In immiscible blends, the properties are related to the interface as well as to the size and shape of the dispersed phase. The morphology is controlled by the equilibrium and non-equilibrium thermodynamics as well as by the flow. As discussed in Part 7..2.2, at equilibrium and within the region of low volume fraction of dispersed phase, φ < φperc = 0.6, droplets are expected, while at φ > φperc = 0.6 a co-continuous morphology, e.g., fibers or lamellae, are usually found. When the polymers are miscible under one set of conditions (for example, within a specified range of concentration, φ, pressure, P, and temperature, T), but immiscible under another, the nonequilibrium morphology depends on the quench depth and time scale. Shallow quenching into the meta-stable region (between the binodal and spinodal curves) results in nucleation of the dispersed drops, followed by their growth. The mechanism of this phase separation is appropriately called nucleation and growth, NG. By contrast, deep quenching into the spinodal region results in spinodal decomposition, SD. Here there is an instantaneous generation of regularly spaced co-continuous structures, with progressive increase

of the concentration difference between the two adjacent regions and increased spacing. The cocontinuity of structures has been reported for scales varying from few nano-meters to hundreds of micro-meters. Both, the NG and SD morphologies are transient, progressively coarsening. At a late stage, both NG and SD mechanisms follow similar ripening patterns, leading to appropriate equilibrium morphology. Without compatibilization, the two phases may totally separate into two layers. In the case of well-compatibilized blends, the action of a compatibilizing agent is similar to that of surfactants in emulsions — one may assign specific surface area coverage per single molecule of the compatibilizer. However, dimensions of the dispersed, compatibilized phase do change with time after cessation of flow. By contrast, addition of a stabilizing agent (e.g., a third polymer immiscible in the two others) may prevent coalescence, preserving the degree of dispersion (but not the orientation) generated during the flow. 7.5.3

Microrheology of Polymer Blends

In this part, the breakup of polymer drops will be discussed, initially dealing with dilute systems (isolated drops) and subsequently with concentrated dispersions where coalescence is of equal importance. Dispersion in Newtonian systems was discussed in 7.3.2.2. Emulsion microrheology. The mechanisms governing deformation and breakup of drops in Newtonian liquid systems are well understood. The viscosity ratio, λ, critical capillary number, κcrit, and the reduced time, t*, are the controlling parameters. Within the entire range of λ, it was found that elongational flow is more efficient than shear flow for breaking the drops. 7.5.3.1 Deformation and Breakup of Viscoelastic Drops The shear deformation of viscoelastic drops in a Newtonian medium has been the subject of several studies. Gauthier et al. [97] found higher values of the critical capillary number than

The Rheology of Polymer Alloys and Blends

those determined for Newtonian drops. Prabodh and Stroeve [99] observed that, during shearing, some drops are greatly extended and only break when the flow is stopped. The authors concluded that at λ < 0.5, the drop elasticity has a stabilizing effect, but for λ > 0.5, the opposite is true. To make things more interesting, the experimental observations of De Bruijn [989] seem to have contradicted the latter conclusion. The author found that the critical capillary number for viscoelastic droplets is always higher (sometimes much higher) than for Newtonian ones, whatever the λ-value; De Bruijn concluded that drop elasticity always hinders drop breakup. For Newtonian drops suspended in viscoelastic fluid, Flumerfelt [972] reported the existence of a minimum drop size below which breakup cannot be achieved. The author pointed out that the elasticity of the medium tends to increase this minimum value for breakup, that is, to stabilize the droplets. In the case when both the droplets and the suspending medium are viscoelastic liquids Wu [987] reported that drops can break up during extrusion even when λ > 4. However, owing to the complex nature of the deformation during flow through an extruder, it was difficult to even speculate on the origin of this phenomenon. Van Oene [978] studied the mechanisms of two-phase formation in a mixture of two viscoelastic fluids. He pointed out that, besides the viscosity ratio and the equilibrium interfacial tension of the two liquids, the elasticity of the liquids plays an important role in deformability of drops. Thermodynamic considerations led to the following relation for the dynamic interfacial tension coefficient:

[

o v12 = v12 + ( do / 12) (σ 11 − σ 22 ) d − (σ 11 − σ 22 ) m

]

(7.08) where νo2 is the interfacial tension in a quiescent polymer blend, do is the initial diameter of the dispersed drop, and (σ-σ22)i is the first normal stress difference of the dispersed (i = d)

493

and of the matrix (i = m) phase, respectively. For (σ-σ22)d > (σ-σ22)m the dependence predicts that higher elasticity of the dispersed than the continuous phase result in more stable drops. On the other hand, for (σ-σ22)d < (σ-σ22)m Eq 7.08 predicts that ν2 < νo2, thus the flow tends to enhance the dispersing process (flow compatibilization). Note that ν2 cannot be negative; for large differences of the normal stress difference and for large drop diameters (thus, for higher concentration of the dispersed phase) this translates into co-continuous morphology, for which the above relation is no longer valid. Since flow affects miscibility of blends near the spinodal, the interfacial tension coefficient must also change with the flow conditions. Theory leads to: ν2(y)= νo2[ - aκ/3b]2b, where a and b are parameters [Onuki, 986]. Han and Funatsu [978] studied droplet deformation and breakup for viscoelastic liquid systems in extensional and non-uniform shear flow. The authors found that viscoelastic droplets are more stable than the Newtonian ones; in both Newtonian and viscoelastic media they require higher shear stress for breaking. The critical shear rate for droplet breakup was found to depend on the viscosity ratio; it was lower for λ <  than for λ > . In a steady extensional flow field, the viscoelastic droplets were also found less deformable than the Newtonian ones. In the viscoelastic matrix, elongation led to large deformation of droplets [Chin and Han, 979]. Bousfield et al. [986] studied the surface tension driven breakup of Newtonian and viscoelastic filaments. The authors found that disturbances grow more rapidly in viscoelastic filaments than in the Newtonian ones, but that there is a retardation of the growth and stabilization at long times, resulting from large extensional stresses (see Figure 7.6). The formation of satellite drops was found to be retarded by the elasticity. The authors analyzed the problem by means of numerical analysis using the Galerkin finite element method, as well as a one-dimensional theory for viscoelastic filaments. Their findings were successfully used to interpret existing experimental data on Newtonian and viscoelastic jet

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disintegration, where the initial disturbance was imposed by nozzle vibration. For viscoelastic jets, an asymptotic solution was offered for the later stages of the process: 1n( R / R0 ) = a0 − t / 3τ ;

for

t >> τ

(7.09) where ao and τ are, respectively, a numerical constant and the relaxation time. This work should have direct bearing on disintegration of viscoelastic filaments in a Newtonian matrix, but application of these findings to polymer blends is more difficult.

Figure 7.16. Computed radius at the neck for disintegrating jet stream of Newtonian (solid line) and Maxwell fluid (dashed line) [Bousfield et al., 986].

dN S / dt = Nb RLb − 2 RSSC − RLSC dN L / dt = RSSC − RLb − RLSC

(7.0)

da2 / dt = ( a2 / dt )V , flow + ( a2 / dt )coalescence

where dNS/dt, dNL/dt and da2/dt are the rate of change of the number of small drops and large drops and the rate of change of drop dimensions correspondingly; RLb is the rate at which the number of small drops, Nb, are produced by breakage of the large drops; RLSC and RSSC are the number rates at which small drops are destroyed by coalescence with large and small drops, respectively; RLSC is the number rate of coalescence between large drops. The first term on the right side of the last relation in Eq 7.0 describes the contribution due to the flow process, and the second reflects that the average thickness of the large domains increases during coalescence. The theory makes it possible to compute the drop aspect ratio, p = a/a2, a parameter that can be directly measured in either transient or steadystate flows. Following the derivation by Hinch and Acrivos [980] the flow-induced changes to the drop aspect ratio were assumed to be proportional to the first normal stress difference coefficient of the matrix fluid. The coalescence was assumed to follow the Silberberg and Kuhn [954] mechanism. These assumptions substituted into Eq 7.0 gave a simple dependence for the aspect ratio: .

p = 4φ d prηm γ / πa1 (σ 11 − σ 12 ) m − a2

(7.) Lyngaae-Jørgensen et al. [993] developed a predictive model of morphology variation during simple shear flow of dilute polymer blends. The model considers the balance between the rate of breakup and the rate of drop coalescence. It was assumed that: (i) The viscosity and elasticity of the dispersed phase are significantly lower than those of the matrix. (ii) Only the cylindrical, large drops (defined by the long and short semiaxes a and a2) are able to break and form small drops. (iii) The coalescence can occur between all types of the dispersed entities. The dynamics of drop formation and breakup can be described by:

where ai are adjustable parameters. A special unit, equipped for light scattering measurements, was attached to a Rheometrics Mechanical Spectrometer with cone and plate to follow the transitional events during shearing of polymer blend melts. The predictions of p obtained by the proposed model were found to be in a reasonable agreement with the experimental observations for poly(methyl methacrylate) blends with either 8 or 0 wt% polystyrene, PMMA/PS. The most interesting finding that came out of this work was that, both theoretically and experimentally, under steady-state flow condi-

The Rheology of Polymer Alloys and Blends

σ12 (Pa) Figure 7.17. Deformability of PS drops in PMMA during steady-state shear flow at 80°C. The points are experimental, the line is to guide an eye [Lyngaae-Jørgensen et al., 993].

tions, the aspect ratio plotted vs. shear stress showed a sharp peak at the stresses corresponding to transition of PMMA viscosity from the Newtonian plateau to the power-law flow, i.e., to the onset of the elastic behavior. The dependence is presented in Figure 7.7. In summary, there is no theory capable to describe the deformability and break-up of viscoelastic drops in viscoelastic media, or even which could provide an explanation for the divergent observations of behavior. The complexity of the problem involves three-dimensionality, free surface, non-stationary flow leading to complex constitutive equations, etc. Furthermore, the role of fluid elasticity in the breakup process is not well understood. In most experiments, elasticity stabilized the deformed drops, thus making the dispersing process more difficult. However, there are also reports on opposite effects. The reported microrheological behavior of two-phase liquid systems may provide but a general guidance for the drop deformation and breakup in polymer blends. For blends of industrial interest, the phenomenon is further complicated by higher concentration of the dispersed phase as well as by presence of numerous additives that affect the interfacial phenomena. The dynamic mechanical shear behavior of several blends, viz. PS with PMMA, PDMS with PEG, PS with PEMA, were studied by Graebling

495

et al. [989, 993]. The linear viscoelastic behavior of these blends with the volume fraction of the dispersed phase φ ” 0.5 was found to follow predictions of Palierne’s emulsion model, which makes use of the viscoelastic behavior of component polymers and a single parameter that characterizes the interface, i.e., the ratio of the interfacial tension coefficient and drop radius, ν12/R. The values of the interfacial tension coefficient determined from the viscoelastic measurements were found to be in good agreement with results obtained from the pendant drop method. However, the theory seems to break down for polymer blends with φ • 0.2. The observed agreement between the experimental data and the theory means that the emulsion model can indeed be used for interpretation of the viscoelastic behavior of polymer blends. The noted deviations at the higher concentration range are not in conflict with the basic premises of the approach. They originate from the imposed limitations of the model (see Part 7.3.2.3.2). 7.5.3.2 Coalescence of Viscoelastic Drops For dilute Newtonian systems, the size of the smallest drop that can be broken is calculable from Taylor’s theory. However for polymer systems, many studies have shown that the equilibrium drop size is usually larger than predicted, and the deviation increases with concentration of the dispersed phase, φ - φo, where φ is the volume fraction of the dispersed phase, and φo § 0.005 is the smallest concentration for which the deviation occurs. Roland and Bohm [984] studied the shearinduced coalescence in two-phase polymeric fluids by small-angle neutron scattering. The coalescence rate was high, dependent on the rheological properties of the two phases and the flow field. Coalescence occurs in shear as well as in quiescent systems. In the latter case the effect can be caused by molecular diffusion to regions of lower free energy, by Brownian motion, dynamics of concentration fluctuation, etc. Diffusion is the mechanism responsible for coalescence known as “Ostwald ripening”. The process involves diffusion from smaller drops (high interfacial

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energy) to the larger ones. Shear flow enhances the process [Ratke and Thieringer, 985]: (d/d0)n =  + a0t,

n = n(y) = 3/2 to 3

(7.2)

where do is the drop diameter at the moment of imposition of stress and ao is a constant. The exponent n decreases from the classical value of three, for quiescent systems, to 3/2 at high shear rates. Flow-induced coalescence is accelerated by the same factors that favor drop breakup, e.g., higher shear rates, reduced dispersed-phase viscosity, etc. Most theories start with calculation of probabilities for the drops to collide, for the liquid separating them to be squeezed out, and for the new enlarged drop to survive the parallel process of drop breakup. As a result, at dynamic equilibrium, the relations between drop diameter and independent variables can be derived. Tokita [977] calculated the total number of collisions per unit volume and time. The author assumed that coalescence is proportional to it and to the number of particles. The latter was assumed to increase with mixing time, being proportional to the shearing energy, y σ2, and inversely proportional to the interfacial tension coefficient, ν12. At equilibrium, the rates of coalescence and breakup are equal. Thus, the equilibrium drop size can be expressed as: .   d = (24 / π ) pr v12φ1 / σ 12 γ − ( 4 / π ) pr Eφ1   

(7.3) where pr is the probability of collision, and E is the macroscopic bulk breaking energy. In agreement with experimental findings, the relation predicts that the equilibrium drop diameter increases with concentration and the interfacial tension coefficient, but it decreases with shear stress. At the low concentration limit, φ ” φo, Eq 7.3 also agrees with the conclusions of Taylor’s theory, but for φd → 0, it predicts an unrealistic limit, d → 0. The idea that, under steady state flow condition, the morphology is fully defined by the dynamic breakup and coalescence processes is an attractive one — one may say, a natural one.

However, behind it is a silent assumption that the flow conditions are strong enough to erase the initial morphology. The presence of the critical value of shear rate, ycr, was recently documented [Minale et al., 997]. The authors reported that unique morphology was observed only above ycr. Below this limit, multiple pseudo-steady state structures were observed for the model PDMS/ PIB system. No attempt was made to generalize this observation. In principle, the phenomenon should be related to the critical value of the capillarity number, κcr, and a ratio of the polymer(s) relaxation times to the rate of shearing. Presence of ycr can also be used to explain recent observations on morphology evolution of PDMS/PIB blends [Grizzuti and Bifulco, 997]. Following a procedure similar to that of Tokita [977], for equilibrium drop diameter in steady simple shear flow, the following dependence was proposed [Fortelny et al., 988, 990]: d = dT + [(ν12 prφ1 ) / η2 f (κ )]

(7.4)

where dT is Taylor’s equilibrium diameter (e.g., calculable from Eq 7.52) and f(κ) is a function of the capillarity number and the rheological properties of the system. Equation 7.4 predicts that as φ → 0, the drop diameter is determined by the Taylor breakup conditions. As the concentration increases, d becomes proportional to the expression in square brackets. The authors reported that in the system PP/EPDM coalescence was more intense than predicted by the dependence. A more recent theory for the dynamic equilibrium drop diameter also started from separate calculations of the drop breakup and coalescence during the steady state shearing. The rate of particle generation was taken to be determined by microrheology, viz. Eq 7.52, [Huneault et al., 995]: .

(dNd / dt )break −up = γ Nd / κ cr tb*

(7.5)

Since the dispersed phase volume is constant, the number of drops, Nd, can be related to the volume fraction of the dispersed phase, φ, and to the drop diameter, Nd = 6φV / πd3.

The Rheology of Polymer Alloys and Blends

The coalescence rate is a function of the collision probability and the dynamics of the collision process. From Equation 7.65 coalescence rate can be written as: .

(dNd / dt )coalescence = −C γ Ndφ 8/ 3 / d 2

(7.6)

where C is a coalescence constant. At equilibrium, the diameter rate of change is zero. From Eqs 7.5 and 7.6, the dynamic drop diameter is: deq = doeq + (6Cκcrt*bφ8/3)/2

(7.7)

where deq is the equilibrium drop diameter (at steady state shearing), in a blend with the volume fraction of the dispersed phase, φ, mixed under a given set of processing conditions, while doeq = dT is its value extrapolated to zero concentration. The only unknown in Equation 7.7 is the coalescence constant, C. Its value can be determined from a plot of deq vs. φ (Figure 7.8.)

Figure 7.18. Equilibrium drop diameter as a function of polyethylene volume fraction in polystyrene matrix. The blend was compounded for 5 min in an internal mixer at 200°C — line is theoretical, Eq 7.7, the points are experimental.

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7.5.3.3 Predicting Drop Size Changes During Processing Mohr et al. [957] analyzed the degree of mixing in a single screw extruder (SSE), using the concept of striation thickness suggested by Spencer and Wiley [957]. The amount of shear strain experienced by an element of fluid in the extruder screw channel was calculated for a number of flow paths. Decreased helix angle, increased ratio of pressure flow to drag flow, and an increased flight height were predicted to improve mixing. The ratio of the viscosities of the minor and matrix phases significantly influenced the degree of mixing. Schrenk et al. [963] analyzed the degree of mixing in a simple annular mixer, which might be helpful for understanding mixing in a SSE. To evaluate the mixedness of the two component polymers, the striation thickness was measured, when the inner shaft was rotated and the outer cylinder was stationary. Near the shaft, the thickness was substantially reduced, but only slightly near the external cylinder. Bigg and Middleman [974] studied the transverse flow in a rectangular cavity, similar to that in a SSE. They used the Marker and Cell technique to calculate the degree of mixing, which was described by the interfacial perimeter per cavity width. As the viscosity ratio decreased, the degree of mixing was enhanced. Chella and Ottino [985] theoretically studied the degree of mixing in a SSE by applying the kinematics of mixing. They evaluated the degree of mixing as a function of the ratio of screw length to height of flight, helix angle, the ratio of pressure flow and drag flow, and the direction of the shearing plane. The stretch of the minor phase increased with axial distance. Mixing was relatively insensitive to the initial feed conditions. The results of the studies on the dependence of mixing on extruder dimensions and operating conditions were in qualitative agreement with Mohr’s analysis [Mohr et al., 957]. The initial morphology generated during the melting and mixing stages in an extruder is important in the development of final morphol-

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L. A. Utracki and M. R. Kamal

ogy of the extrudate. Lindt and Ghosh [992] suggested that an abrupt morphological change occurs during the simultaneous melting and striation formation in the melting zone in a SSE. Within a fraction of a second, the scale of mixing drops by several orders of magnitude. High stress in the thin molten film in the melting zone causes a reduction of striation thickness of the minor phase. The lamellar layers may be developed when the minor component pellets melt at the interface between the melt film and the solid bed. The layers could become threads as they undergo breakup. Finally, the threads change into droplets, as they are broken. Scott and Macosko [99] proposed a mechanism of the morphology development based on experiments carried out in a batch mixer. When the minor component pellet melts, sheets or ribbons of the dispersed phase are formed due to dragging of the pellets on the hot surface of the mixing equipment. Next, holes are formed in the sheets or ribbons of the dispersed phase, as the interfacial instability starts, and, sheet or ribbon morphology changes into a lace structure. Then, the lace breaks into irregularly shaped pieces with diameters equal to the ultimate sphere morphology. The above two proposed mechanisms incorporate concepts involving distributive and dispersive mixing. Layer or sheet morphology development is mainly due to distributive mixing. Distributive mixing refers to the physical process of blending two fluids such that the physical separation distances are reduced to a scale where diffusion or a chemical reaction can occur [Bigio and Conner, 995]. Breakup of layers into threads, laces or spheres could be attributed to dispersive mixing which is related to instability of the minor phases. During the last few years, an attempt was made to develop a model describing morphology evolution during polymer blending in a twin screw extruder. The first model [Shi and Utracki, 992] was based on a simplified flow analysis and microrheological considerations of the dispersedphase drop disintegration. The effects of coalescence were neglected.

The second generation model comprised more refined flow analysis, two mechanisms of dispersion (the fibrillation mechanism and a drop splitting mechanism for low supercritical capillary numbers, with the choice of break-up mechanism based on locally computed microrheological criteria), as well as coalescence effects [Huneault et al., 995a]. The latter effects were taken into account by determining the coalescence constant in Equation 7.7 from the plot shown in Figure 7.8. Thus, the developed model was selfconsistent, fully predictive, without any adjustable parameters. The validity of the theoretical assumptions was evaluated by comparing the two models’ predictions with the experimentally measured drop diameter at different axial positions in the twin screw extruder. Experimentally, after the extrusion reached steady state, the screw rotation was stopped and the molten blend was quenched within a specially designed extruder barrel. It was estimated that the PS/PE blends were quenched within 7-0 sec. The second model predicted well the morphology evolution of non-compatibilized blends of PS in PE, and their inverse of PE in PS (see Figure 9.6 in Chapter 9). The microrheology makes it possible to expect that: (i) The drop size is influenced by the following variables: viscosity and elasticity ratios, dynamic interfacial tension coefficient, critical capillarity number, composition, flow field type, and flow field intensity; (ii) In Newtonian liquid systems subjected to a simple shear field, the drop breaks the easiest when the viscosity ratio falls within the range 0.3 < λ < .5, while drops having λ • 3.8 can not be broken in shear; (iii) The droplet breakup is easier in elongational flow fields than in shear flow fields; the relative efficiency of the elongational field dramatically increases for large values of λ • ; (iv) Drop deformation and breakup in viscoelastic systems seems to be more difficult than that observed for Newtonian systems; (v) When the concentration of the minor phase exceeds a critical value, φd > φo ≅ 0.005, the effect of coalescence must be taken into account; (vi) Even when the theoretical predictions of droplet deformation and breakup

The Rheology of Polymer Alloys and Blends

are limited to infinitely diluted, monodispersed Newtonian systems, they can be successfully used for predicting the development of blend morphology during compounding in twin-screw extruders. More recently, experiments were conducted in a corotating, intermeshing twin-screw extruder using the same PE/PS system as described above [Huneault et al., 995b]. The screw geometry consisted of five zones: melting, melt conveying (no pressure), mixing (kneading), pumping, and flow through a die. The specimens were scooped from three ports and quenched within a second. After dissolution of the matrix, the dispersed phase was divided into fibers and droplets, characterized separately. Immediately after melting the dispersed phase formed into fibers and droplets, both with diameters below 0 µm. Contrary to the previous model assumptions, fibers did not break in the unfilled conveying region that followed the melting section, instead, they were mainly destroyed in the kneading section. Fibers were present after melting even at concentration of the dispersed phase as low as 2 wt%. The effect of concentration, other than increasing the final diameter of droplets, was to increase the fiber content. The observations also seem to indicate that coalescence is not limited to drops. However, near the die, the average drop diameter did decrease to about d ≅  µm range (as observed earlier). Clearly, the microrheology of polymeric systems is more complex than the classical microrheology of Newtonian, low viscosity liquids would suggest. Near the liquefaction point (either Tm or Tg) the viscosity is of the order of 02 Pas and the relaxation time of the order of 00 s [Angell, 997]. As temperature increases along the barrel, these magnitudes decrease according to the Vogel-Tamman-Fulcher relation: η = ηo exp {B / (T − To )} τ = τ o exp {B / (T − To )}

(7.8)

where: To ≈ Tg + 50

These long relaxation times are responsible for non-equilibrium structures, generated by the

499

mechanical action of the compounding equipment that are not taken into account by the microrheology. The microrheological model provided good agreement with the experimental data obtained after 7-0 s quenching. However, these data were on purpose collected from the second half of the TSE barrel, where the temperature was reasonably stable (isothermal model). Evidently, the evolution of blend morphology is more complex than a simple, “steady-state” model can predict. The rapid variations of morphology will be particularly important for computations of reactive compatibilization. Different type of morphological changes during blending was reported by Sundararaj et al. [992, 995]. For either reactive or non-reactive blends in an extruder, internal mixer, or a miniature cup-androtor mixer, similar morphological features were observed. Initially, during melting, the polymers were stretched into sheets and ribbons, which broke into fibers, then in turn into drops. One reason that may explain the reported differences in morphology is the concentration of the dispersed phase — 5 vol% was used by the first authors, whereas • 20 wt% by the latter. More detailed information of this, as well as on the topic discussed in the following part, will be found in Chapter 9: Compounding Polymer Blends. 7.5.3.4 Mixing and Blending in Extensional Flow Field Most works on liquid mixing in the extensional flow field considered convergent flow of a Newtonian liquid from a reservoir to a capillary [Tsebrenko et al., 974, 976; Ablazova et al., 975; Krasnikova et al., 978; Han and Funatsu, 978; Chin and Han, 979, 980; Han, 98; Han and Yu, 98; Suzaka, 982; Vinogradov et al., 982; Utracki et al., 986]. Recently, a device capable of mixing polymeric liquids (having widely ranging viscosity ratios) in an extensional field, was constructed [Nguyen and Utracki, 995]. The extensional flow mixer, EFM, was designed incorporating the following principles, based on microrheological analysis:

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L. A. Utracki and M. R. Kamal

. The blend must be exposed to the extensional flow fields and to semi-quiescent zones. 2. The convergence’s and divergence’s should be of progressively increasing intensity. 3. The convergent-divergent flow should be generated in the radial not axial direction. 4. To reduce the pressure drop, and to prevent blockage, slit restrictions should be used. 5. The extensional flow mixer must be adjustable. 6. The rate of flow, up-stream from the plates, should be approximately constant. In EFM, the material flows from the rim between two circular, convergent-divergent plates with ridges facing each other, toward the opening in the center of the lower plate. To assess relative merit of the extensional mixing, EFM was attached to a SSE. For comparison, the blends were also prepared in a co-rotating, intermeshing TSE. The same temperature profiles were used for SSE+EFM as for TSE. In all cases, the dispersed phase was significantly more viscous than the matrix, λ • 4. The efficiency was judged considering: . The degree of dispersion in PS/PE blends of PS with either 5 or 0 wt% HDPE. At the exit from SSE+EFM, fibrillar and/or nodular morphology was observed. The number average fiber diameter decreased with pressure across the c-d plates from dn = .2-0.7 µm to dn = 0.2-3 µm, at, respectively, P = 0.3 to 8.6 MPa. The blends prepared in a TSE showed much coarser morphology, containing mainly infinitely long HDPE fibers with diameter varying from dn =  to 0 µm. 2. The results of PP impact strength improvement by incorporation of EPR are summarized in Figure 7.9. As evident, impact strength at room temperature increased with EPR content. Clearly, SSE+EFM compounding resulted in higher impact strength (points) than that obtained from TSE (broken line). 3. The ability to homogenize resins with widely different molecular weight can be exemplified by UHMWPE/HDPE blends. Addition of high MW polymer is expected to increase G’, G”, η, and the first normal stress coefficient, Ψ1. For the linear polymers, these parameters at

low deformation rates: ηo and Ψ1ο, are proportional to Mw3.5, and Mw7, respectively. Thus, the elasticity is more sensitive to the high MW fractions. For this reason, the frequency dependence of the storage modulus ratio: G’(blend)/ G’(PE), at 200°C, for HDPE and its blends with 3 wt% UHMWPE was measured. The blends prepared in TSE had the worst performance G’ at ω = 0.0 rad/s increased by 90%, while in EFM+SSE the increase was up to 20% at P = 8.6 MPa. 4. Elimination of gel particles in the reactor powder. The gel particles can form during polymerization of EVAc or TPO. Since they may be considered very high molecular weight fractions of the same resins, SSE+EFM was used to eliminate or reduce the blemishes and improve the mechanical performance. On both counts, the performance was found at least equivalent to that obtained using a TSE. One may calculate the pressure drop, ∆P, in EFM starting with well-known expressions [Cogswell, 972; Binding, 988; Tremblay, 989]. An expression derived from Binding’s theory was found to provide excellent prediction (no adjustable parameters) of the pressure across EFM. For more details see Part 9..6.4 in Chapter 9 of this Handbook. Historically, the counter-rotating TSE was known as the “calendering” extruder, with high stresses existing between the two screws and low stresses outside this region. Owing to the high calendering pressures the screws could rub against the barrel causing premature wear. Thus, slower speeds (up to 50 rpm), and large intermeshing gaps were recommended. One of the advantages of these counter-rotating machines has been the presence of the elongational flow field within the calendering zone. The machines have been successfully used in numerous applications requiring high dispersive stresses. During the last few years, there has been renewed interest in upgrading the TSE performance. This was spurred by the advances in the theory and methods of analysis of the screw elements’ performance. The analysis resulted in

The Rheology of Polymer Alloys and Blends

50

weight fractions. On the other hand, the modern TSE, designed for maximum extensional flow, provides superior mixing capabilities over the older models. 7.5.4

Figure 7.19. Effect of EPR addition on PP’s notched impact strength at room temperature. The specimens were prepared either in a TSE (horizontal lines) or in a SSE/EFM. In the latter case the results depended on the pressure drop across EFM [Utracki and Luciani, 996].

modification of the screw profiles (higher free volume of the process), as well as in development of new mixing (or kneading) elements. The increased free volume (thus slender screw profile) resulted in lowering the average shear rate, thus the screw speeds needed to be increased. New kneading and mixing elements have been designed to improve either the distributive or dispersive mixing. The new kneading blocks (mono- and bi-lobal, to be used either in coor counter-rotating TSE, the tri- and hexa-lobal only in counter-rotating TSE) were designed to maximize the extensional flow field within the lobal pools, and reduce the shear field in the intermeshing and overflight regions. Initial reports indicated greatly improved mixing capabilities, even for polymeric systems having large differences in the rheological flow parameters [Thiele, 995]. In conclusion, mixing in extensional flow fields was found to provide better mixing capabilities. On the one hand, the data demonstrate that a singlescrew extruder equipped with EFM, may offer an inexpensive alternative to a twin-screw extruder, TSE, for the mixing and compounding applications. The unit outperformed TSE in: generating finer dispersion, improving the blend toughness, or facilitating dissolution of very high molecular

Flow Imposed Morphology

Several aspects of blend morphology have been already discussed, viz. in Part 7..2.2 the equilibrium morphology in 3D and D, e.g., the concentration dependent changes of morphology from dispersed to co-continuous at the phase inversion concentration, φI; in Part 7..2.3. the interphase thickness, and its variation with compatibilization; in Part 7.3.2.2. flow effects on the drop deformation and break-up in Newtonian system; in Part 7.4.2 the phase separation and flow, in Part 7.5.2. an introduction to morphology of immiscible blends, and in Part 7.5.3. the microrheology of polymer blends, including mathematical modeling of morphological changes during compounding dilute, non-compatibilized blends. In this Part, focus will be on the changes of morphology imposed by different flow fields. Since the chapter is not as much concerned with blend morphology as with the flow and rheological characterization of polymer blends, the main interest is in the reversed relation — how the variations in morphology affect the flow behavior. After reviewing the flow of suspensions and emulsions as well as discussing microrheology of immiscible polymer blends, it must be evident that the degree of dispersion and the type of structure strongly relate to the type and intensity of imposed stresses during flow. Note that both concentrated suspensions and emulsions show yield stress and time-dependent flow. These macroscopic observations are related to the structural changes occurring on the micro-scale. Similar behavior of polymer blends is to be expected. One has to keep in mind that during polymer processing, neither the thermodynamic miscibility, the macromolecular configuration (e.g., entanglement), nor the morphology is in an equilibrium state. Most models of the morphological changes in polymer blends assume that an average response (e.g., an average size drop is being broken, or aver-

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age size drops coalesce) provides good representation of the whole system. This assumption should be reasonably correct for blends with narrow distribution of drop sizes. However, there are reports that (e.g., during the initial stages of blending in a twin-screw extruder) the domain sizes may differ by three orders of magnitude. Here the “average size” response may not be valid. Recently, a kinetic theory of the structure development in moderately concentrated polymer blends was proposed [Patlazhan and Lindt, 996]. The breakup and coalescence in a steady-state shearing were considered, assuming a temporal population balance. This development provides a framework for incorporation of the elementary phenomena of drop breakup and coalescence to an overall model. The effect of flow on miscibility of polymer blends is another area of industrial importance. There is evidence that during processing, the imposed stresses can change the critical temperature by at least 60°C, causing miscibility inside the processing equipment. The blends, upon release of pressure may undergo spinodal decomposition that results in superior performance [Inoue, 993]. Flow may also result in mechano-chemical degradation processes that generate reactive sites, viz., radicals, peroxides, acids, etc. Furthermore, trans-esterification and ester-amide exchange reactions are well documented. These reactions affect the phase equilibrium as well as the regularity of the chain structure, thus dispersion in the blend and its crystallinity. Blend morphology refers to the spatial arrangement of the blend components forming either a disperse, a stratified (e.g., lamellar or a sandwichtype) or a co-continuous structure. Generation of morphology depends on the viscosity and elasticity ratio of the polymeric blend components (at constant stress). Both ratios vary with the type and the intensity of the flow field. While the viscosity ratio seems to control the ease of dispersing the component, thus the degree of dispersion, the elasticity ratio contributes to shaping the phases — the type of morphology. Two others pertinent parameters are the concentration and the level of interfacial interactions. To modify the

interfacial energy, blends are usually compatibilized either by addition of a compatibilizer, or by reactive blending. Once formed, the morphology needs to be stabilized against possible destruction during the forming steps. The rheological properties of a two-phase system depend not only on the rheological behavior of the components, but also on the size, size distribution, and the shape of the discrete phase droplets dispersed in the continuous matrix phase. Flow affects morphology in two different ways. * It changes the degree and type of dispersion on a local level, viz. drop breakup and coalescence. * It causes migration of the dispersed phase, thus imposing global changes of morphology in the formed parts, viz., skin-core structures, weld lines, blush lines, etc. In consequence, the flow imposed morphologies can be classified as: (i) Dispersion (mechanical compatibilization), (ii) Fibrillation, (iii) Lamellae formation, (iv) Coalescence, (v) Interlayer slip, (vi) Encapsulation, etc. These types will be discussed below under appropriate headings. There is a reciprocal relation between morphology and flow behavior. Plochocki [978, 983] defined the ‘particular rheological composition’, (PRC), most frequently observed in polyolefin blends. At PRC the η = η(φ) function reaches a local maximum or minimum. The existence of the maximum is related to a change of the dispersed phase, e.g., from spherical to fibrillar or from dispersed to co-continuous, while that of the minimum is related to a reciprocal change and/or to variation of the specific volume. Table 7.6 provides a partial reference to studies on the effects of flow on the morphology of polymer blends [Lohfink, 990; Walling, 995]. Dispersed phase morphology development has been mainly studied in a capillary flow. To explain the fibrillation processes, not only the viscosity ratio, but also the elasticity effects and the interfacial properties had to be considered. In agreement with the microrheology of Newtonian systems, an upper bound for the viscosity ratio, λ, has also been reported for polymer blends — above certain value of λ (which could be significantly larger than the

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The Rheology of Polymer Alloys and Blends

Newtonian value of 3.8) the dispersed phase could not be deformed. By contrast, lower bounds of λ were not established for polymer blends. Incorporation of compatibilizers (a third phase) into immiscible blends improved the adhesion between blend phases and helped to achieve

Table 7.6.

mechanical properties comparable to those of homopolymers. The formation of lamellar structures with specifically designed arrangement of the dispersed phase in the matrix phase could provide barrier properties comparable to those achieved in multi-layer parts.

Studies of flow field effects on polymer blends morphology

Flow Type

Blend

Observations

Reference

. Theory 2. Uniform shear field

Viscoelastic fluids PS/PE

elastic free energy approach particle size distribution for: λ > , coarse; λ ” , fine 0.5 < λ < 2.0, composition dependent: PS — droplet breakup; PMMA — elongated droplets λ < , long PE fibers; λ > , long PE fibers maximum aspect ratio at the transition from the Newtonian to power-law flow region λ < , long PE fibers λ > , long PS fibers λ < , PP fibers, high shear stress dependent length. λ ” , POM fibers, shear stress dependent shape.

Van Oene, 972 Starita, 972

PMMA/PS

LLDPE/PS PMMA/PS 3. Capillary flow

HDPE/PS PS/PP PP/EP POM/CPA

PP/PS POM/EVAc

PE/PP HDPE/PS

4. Annular and slit flow 5. Convergent flow

EVAl/PP HDPE/PA-6 model fluids model fluids

6. Flow in mixing devices. PP/PC PA-66/EPR PET/EP EVAl/PP

λ < , PP fibers, relaxation dependent length. λ ” , POM films & fibers ; λ =.32, POM microfibers; λ = 4.3, POM fibers & particles λ > , continuous fibers; λ < , breakup, small droplets λ < 0.7, fibers; 0.7< λ < .7, undulant fibers & rods; λ > 2.2, undeformed droplets λ > , EVAl fibers PA-6 platelets & lamellas permeability barrier single drop deformation in axisymmetric convergence/divergence single drop deformation in a slit convergence/divergence PC drop size depends on viscosity and λ EP particle size depends on ν2, κ, and λ EVAl particle size dependence on concentration

Chuang and Han, 984

Dreval et al., 983 Lyngaae-Jørgensen et al., 993 Han and Yu, 97, 972 Han, Kim and Chen, 975 Danesi and Porter, 978 Ablazova et al., 975 Tsebrenko et al., 976 Tsebrenko, 978 Krasnikova et al., 978 Tsebrenko et al., 980

Alle and Lyngaae-Jørgensen, 980; Alle et al., 98 Min et al., 984 Lepoutre, 989 Subramanian, 985, 987 Mighri et al., 997 Bourry et al., 998 Favis and Chalifoux, 987, 984 Wu, 987 Lepoutre, 989

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Table 7.6.

L. A. Utracki and M. R. Kamal

Continued

Flow Type

Blend

Observations

Reference

7. Flow in mixing devices with compatibilization (interfacial tension modification)

PA-6/PP LDPE/PS PA-6/PE PA-6/ PE/EVAc

Maleic anhydride grafted PP (PP-MA) surface active compounds chemically modified dispersed phase chemically modified dispersed phases

LDPE/PS LLDPE/PS HDPE/PS PVF/PE LDPE/ABS PP/EVAl HDPE/PA

hydrogenated butadiene-b-styrene diblock copolymer (HPB-b-PS)

Ide and Hasegawa, 974 Heikens and Barensten, 977 Chuang and Han, 984 Chuang and Han, 985; Han and Chuang, 985 Fayt et al., 98, 982, 986

8. Slit flow & compatibilization

PP/EVAl PP-MA/ EVAl HDPE/PA-6

9. Flows in injection molding

PP-MA/EVAl; PE-MA/PA-6 ABS/ rubber reinforced PP/EPDM PP/PE/ EPDM PA-6/ EVAc

poly(hydrogenated butadiene-bmethylmethacrylate) PP-MA Modified PA: platelet formation, permeability barrier maleic anhydride grafted PP: lamellar formation, permeability barrier methacrylic acid/isobutyl acrylate: lamellas, post extrusion calendering/ elongation maleic anhydride grafted PP & HDPE: lamellas for permeability barrier delamination layer of rubber particles arranged in rows skin: 350 to 400 µm; thin, elongated minor phase; core: isotropic spherical inclusions

skin: no other distinct layer; EVAc concentrated near core PP/TPO skin: major deformation; core: dispersed spherical drops POM - rubberized skin: semi-crystalline and rubber sheets; core: tough rod shaped rubber particles aligned in flow direction PA-6/PE maximum anisotropy at intermediate position near the mold wall PP/EPDM maximum particle deformation 00 µm under surface PP/EVAl skin: small rectangular platelets; shear zone: lamellar morphology core: undeformed EVAl particles

Ouhadi et al., 986 Fayt & Teyssie, 989 Lepoutre, 989 Subramanian, 985, 987 Lohfink, 990 Lohfink and Kamal, 993 Gonzalez-Nunez et al., 993 Kamal et al., 995; Garmabi and Kamal, 995 Kato, 968 Ho and Salovey, 98

D’Orazio et al., 986, 987 Karger-Kocsis, 987 Percorini, 990

Ghiam and White, 99 Michaeli et al., 993 Walling, 995

The Rheology of Polymer Alloys and Blends

7.5.4.1 Dispersion Microrheology can be used to express the time dependence of the drop diameter, d, in steady-state shearing, without coalescence. From Eq 7.5, assuming that drops breakup occurs only if the shearing time at each appropriate shear stress exceeds the required time to break, t > tb, (for shearing times t < tb, the average drop remains unchanged) one can express the relative change of drop diameter as:

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of the specimen. The data in Figure 7.20 were obtained from capillary flow. The morphology has been affected by the extensional flow field upstream from the die, shear flow (and flow-induced encapsulation) inside the die, and slower cooling in the center than at the core of the extrudate. After such complex morphological changes, empirically the drop diameter decreases with logarithm of the deformation rate, d / do =  - aolny, where ao ≅ 0.3-0.6, is a material parameter. 7.5.4.2 Fibrillation

.

ln ( do / d ) = γ (t − tb ) / 3κ cr tb* = [(γ / γ b ) − 1] / 3 sin ce: γ b ≅ 2 → ∴ d = do exp {(1 − γ / 2) / 3}

(7.9) Since derivation of this relation considered only the drop-splitting mechanism and neglected coalescence, its validity may be limited to small capillarity numbers, κ* = -2, and low concentrations.

Figure 7.20. Reduced drop diameter vs. rate of shear at 250°C. The blend comprising 0 wt% PA-6 in HDPE was extruded through a capillary with the L/D = 40 [Utracki et al., 986].

The drop diameter usually decreases with an increase of shear rate (see Figure 7.20). However, one must be careful with the data analysis. As microrheology indicates, there are different mechanisms operating in different flow types (e.g., shear and elongation) or at different field intensity. Furthermore, there is usually a difference in the quench time between the outer layer and the core

The term “fibrillation” seems to suggest that the mechanism responsible for formation of fibers or fibrils is extensional, e.g., at the entrance to capillary. Once inside the capillary, the blend undergoes shear flow, with intensity dependent on the radial position. A study was made of the evolution of morphology of PMMA blends with core-shell elastomeric latex particles, poly(butylacrylate-co-styrene) [Bousmina and Muller, 996]. It was found that within the plug flow region, the particles were randomly distributed, but in the outer part of the extrudate, the particles were aligned into straight pearl strings. Had coalescence been possible, these would form fibers. The authors proposed a mechanism of string formation within the steady-state shear zone of the flow field. It seems reasonable to postulate that there are at least two possible fibrillation mechanisms: the extensional flow at the entrance to capillary, and the “stringing” process described by Bousmina and Muller. Stress-induced fibrillation occurs in steady-state shearing or elongation when κ > 2. Under these conditions, the dispersed phase is co-deformational with the matrix. Since the capillarity parameter is proportional to diameter (viz., Eq 7.52), it is easier to fibrillate coarser dispersions at concentrations exceeding a limited value: φlimit ≈ aλb, where the numerical value of the a, b equation parameters depend on the blend [Krasnikova et al., 984]. Fibrillation is also affected by the presence of compatibilizer. From the perspective of the capillarity parameter, κ, addition of a surface tension modifier has two effects: it lowers the

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since the diameter of the fibrils remains virtually constant (similarly as d = a2 in Figure 7.22), low viscosity ratios result in short fibers. On the other hand, for λ • .7, again drops and platelets were detected. In short, for the best results λ ≅  is preferred.

Figure 7.21. Stress-dependent values of the orthogonal axis of deformed polypropylene drops in shear at 250°C. The blend comprised PP/PA-6 = :5, with and without an acrylic compatibilizer, AA [Søndergaard et al., 992].

interfacial tension coefficient (thus increasing κ) and decreasing the initial drop diameter (thus decreasing κ) — the net result is difficult to predict. An illustration is provided in Figure 7.2. Here PP drops in PA-6 matrix were observed during shearing in a cone-and-plate geometry, without and with an acrylic compatibilizer, AA. For both systems the dimensions (long and short axes of a prolate ellipsoid) were reaching a plateau level at shear stress σ2 • 0 kPa. Evidently, the plateau value of the long axis, a = L, was higher for the system without AA than with it. However, the rate of elongation indicates that AA facilitated the fibrillation process [Søndergaard et al., 992]. Tsebrenko et al. [976] reported on fibrillation of POM in a copolyamide (CPA) matrix, as a result of flow through a capillary. Fine fibrils with diameters of about 20 µm and length 3.2 mm were obtained during extrusion at T = Tm(POM) + 6°C. The low extrusion temperature facilitated stress-induced crystallization of the POM fibers, preserving the morphology engendered at the entrance to the spinneret. As evident from data in Table 7.7, fibrillation of POM in poly(ethylene-co-vinyl acetate), PEVAc, strongly depended on the viscosity ratio, λ [Tsebrenko et al., 982]. These data also indicate that for low viscosity dispersed phase, the coalescence that results in formation of plate-like objects, complicates the blend morphology. Furthermore,

Table 7.7.

Effect of viscosity ratio on fibrillation of POM in POM/PEVAc = 20/80 [Tsebrenko et al., 982]

λ = ηd/ηm

d±σ (µm)

0.35 0.9 .05 .70 4.0

5.3 ± 2.5 4.2 ± .8 5.5 ± 3.6 6.2 ± 3.6 7.3 ± 5.8

Number POM dispersion form (wt%) of fibrils drops fibers plates 6,500 3,200 6,800 4,300 4,400

0 0 0 0 48

83 00 00 80 50

7 0 0 20 2

Recently, polymer blends comprising a polyetheresteramide block copolymer, PEBA, with liquid crystalline copolymers, LCP, were prepared [Champagne et al., 996]. The minor component was deformed into fibrillar-type morphology that enhanced the mechanical properties in the draw direction in a manner comparable to unidirectional continuous-fiber reinforced composites. Films prepared using a single screw extruder were melt drawn on calendering rolls. The storage modulus of blends containing 30 wt% LCP increased with draw ratio, DR ” 2, nearly 50-fold in comparison to neat PEBA (from ” 8 MPa to almost  GPa). The blend morphology was characterized by dissolving the PEBA matrix, followed by gravimetric and microscopic analysis of the LCP phase. As expected, the average fiber diameter decreased as a function of DR-0.5. It was noted that only relatively large drops were deformed into fibers, leaving nearly 50% of LCP in the form of small dispersed nodules. The fiber content as a function of DR followed a trend parallel to that of the mechanical properties. Longitudinal and trans-

The Rheology of Polymer Alloys and Blends

verse moduli followed the Halpin-Tsai predictions for unidirectional composites. Properties of compression molded specimens prepared from these blends compared favorably with glass fiber composites. Drop deformation in shear that leads to fibrillation was recently examined using microscopy, light scattering and fluorescence [Kim et al., 997]. The authors selected to work with systems near the critical conditions of miscibility, thus where the flow affects miscibility and reduces the value of ν2. The drop aspect ratio, p, plotted as a function of the capillarity number, κ, showed two distinct regimes. For κ < κcr, p was directly proportional to κ, whereas for κ > κcr, p followed more complex behavior, with an asymptote that corresponds to flow-induced homogenization. 7.5.4.3 Lamellar Morphology Lamellar morphology occurs in flow regimes where the dispersed phase undergoes two dimensional stretching with the formation of multilayers. In immiscible blends, such a structure may enhance barrier properties when the dispersed phase is a barrier material (e.g. PA, EVAl) and the matrix phase is a commodity polymer (e.g., PE, PP, etc.). Well-developed lamellae increase the length of the pathway for permeants diffusing through the blend. The longer path causes a lowering of the concentration gradient across the blend material, thus reducing the mass flux or permeability. Subramanian [985,987] was the first to develop a method for generating lamellar morphologies in polymer blends during melt processing. The method has been used to impart permeability barriers to low cost polyolefins, PO, using small amounts of a barrier polymer. For example, blending, under controlled conditions, HDPE and a modified polyamide, either PA-6 or PA-66, led to compositions that during the subsequent blow molding or film blowing, generated lamellar PA dispersions. In particular, the lamellar morphology blends of PE and PA-6, were produced in a blow-molding machine. The product exhibited good barrier properties.

507

The optimum performance was obtained using 8 wt% of PA. The work resulted in commercialization of the Selar™ technology. PP/EVAl blends with lamellar morphology were produced either in a single screw extruder with a specially designed die [Lohfink and Kamal, 993], using the injection molding machine [Walling, 995; Walling and Kamal, 996], or in an extruder with an annular blown film die [Lee and Kim, 997]. To produce PP/EVAl sheets with lamellar morphology, Lohfink and Kamal [993] designed and constructed a biaxially stretching slit die, which had converging and diverging sections to achieve the desirable extensional flow. In PP matrix, EVAl lamellae were formed in the sheet core. The optimum barrier performance for oxygen transmission was obtained using 25 wt% EVAl. The barrier properties of the blends were superior to those obtained later in the injection molding process [Walling and Kamal, 996]. During the injection molding of PP/EVAl blends, a complex morphology was obtained. In the core region, small relatively undeformed EVAl particles were found. By contrast, in the high shear zone near the skin, lamellae were present. Formation of the lamellar structure was enhanced by increasing EVAl concentration, compatibilization, and reduced mold thickness [Walling, 995; Walling and Kamal, 996]. Morphology of blends is strongly influenced by the mixing mechanism. Well-developed lamellar morphology is produced when deformation of the minor phase is high, and its breakup is minimized. Coalescence of the deformed minor phase could also contribute to lamellar morphology [Lohfink and Kamal, 993]. Thus, from the microrheological point of view, the best results are to be expected from systems where: () the domain size of the dispersed phase is relatively homogenous, with d ≅ 50 µm, (2) the viscosity of the dispersed phase is lower than that of the matrix, i.e., λ < , (3) the dispersed phase shows a strain hardening behavior. Breakup of the minor phase has been discussed in detail in the former sections. Important aspects relating to the production of lamellar morphology blends will be discussed below.

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L. A. Utracki and M. R. Kamal

7.5.4.3.1 Kinematics of Mixing Spencer and Wiley [957] have found that the deformation of an interface, subject to large unidirectional shear, is proportional to the imposed shear, and that the proportionality factor depends on the orientation of the surface prior to deformation. Erwin [978] developed an expression, which described the stretch of area under deformation. The stretch ratio (i.e., deformed area to initial area) is a function of the principal values of the strain tensor and the orientation of the fluid. Deformation of a plane in a fluid is a transient phenomenon. So, the Eulerian frame of deformation that is traditionally used in fluid mechanical analysis is not suitable for the general analysis of deformation of a plane, and a local Lagrangian frame is more convenient [Chella, 994]. A general equation for the kinematics of distributive mixing was developed in a Lagrangian frame. The degree of mixing was described in terms of inter-material area density, or striation thickness, which could be obtained experimentally. Using an ideal laminar mixing model, the thickness of an individual particle of the minor phase was expressed as δd = 2φ / av, where φ is the volume fraction of the minor phase and av is the interfacial area density [Ottino et al., 98]. As a result of deformation, the lamellar thickness as well as the interfacial area density change with the local strain, γl, viz. δd = δοd / γ; aoνγ = aν, where the symbols with upperscript “o” indicate the initial conditions. For simple shear flow deformation, when the deforming interface has the same direction as the flow, the local area strain is related to the linear strain of the flow field, γ = ty: . 2  δ d = δ do 1 +  t γ      

−1/ 2

(7.20)

The above equation could be used for the interpretation of lamellar morphology development, when breakup of the minor phase is excluded, and the interfacial tension coefficient is vanishingly small.

7.5.4.3.2 Parameters Determining Lamellar Morphology Development Distributive lamellar mixing depends on the deformation rate, deformation time and the initial direction of the interface. The degree of mixing increases as the deformation rate and time increase. Initial direction of interface, favorable for maximum mixing, is also needed for achieving a high degree of mixing [Ottino et al., 98]. Lohfink and Kamal [993] observed that an increase of die gap size yielded fewer, but thicker layers. On the other hand, a smaller die gap size resulted in an increased number of stacked thin layers. Higher screw rpm produced a more pronounced lamellar structure. In real mixing equipment, such as a SSE, deformation time (residence time) could be limited by the deformation rate (screw speed). Lee and Kim [997] reported that an increase in screw speed reduced the degree of mixing, because the residence time decreased and the minor phase melting was insufficient. The viscosity ratio, λ, is one of the major parameters in determining the deformation of the minor phase layer. When the viscosity ratio varies from zero to infinity, Taylor’s Eq 7.49 predicts that, deformability of a small drop would change from .0λ to .8λ. For viscoelastic systems, Gonzalez et al. [993] and Lee and Kim [997] obtained higher deformation for lower viscosity ratio, when they changed the viscosity ratio by changing the viscosity of matrix material. Lamellar morphology development also depends on the volume fraction of the minor phase. The individual thickness of EVAl phase in PP matrix phase decreased when the concentration of EVAl phase decreased from 30 to 20 vol% [Kamal et al., 995]. In the experiment, for the processing conditions, the viscosity ratio was λ > 1. The minor phase layers become thinner as interfacial tension coefficient, ν2, decreases [Kamal et al., 995; Lee and Kim, 997]. This confirms that the decrease of ν2 results in a more efficient transfer of stress from the matrix to the minor phase layer [Gopalakrishnan, 995]. As the ν2 is reduced further, the layers of minor phase transform into fibers [Kamal et al., 995]. These results are in agreement with the discussed morphology develop-

The Rheology of Polymer Alloys and Blends

ment mechanisms [Ottino et al., 1981; Lindt and Ghosh, 1992; Scott and Macosko, 1991]. 7.5.4.4 Coalescence Droplet-droplet coalescence was already discussed in Part 7.4.2.2. Here, the effects of coalescence on morphology will be summarized. Under normal circumstances, there is a dynamic equilibrium between coalescence and dispersion processes, thus it is difficult to assign a particular effect as due to coalescence. However, during flow at temperatures near the melting point, the effects of coalescence dominate the final morphology. For example, blends of HDPE with up to 30 wt% of PA-6 were extruded using a capillary viscometer at T = 150, 200 and 250°C. All the extrudates contained PA-6 fibrils, independently at T below or above the melting point of PA-6, Tm = 219°C [Utracki et al., 1986]. Judging by the diameter of the resulting PA-6 domains and their internal structure, the fibrillation originated mainly from the flow-induced coalescence. For capillary flow at 150°C, the extensional stress, σ11 = 50-800 kPa, at the entrance to capillary was calculated from [Cogswell, 1972]: η/ηΕ = 2 tan-2α; σ12γ˙ = 2σ12ε˙

α = arctan (2ε˙ /γ˙); (7.121)

Since the tensile yield stress for “solid” PA-6 at 150°C was determined as σy = 15 kPa, independent of the rate of straining, the extensional stress in the capillary entrance was more than sufficient to deform the amorphous part of PA-6. Owing to crystallization, the elongated structures, once created, could neither disintegrate nor elastically retract to spherical shapes. Similarly, at T > Tm, coalescence of semicrystalline dispersed domains combined with stressinduced crystallization leads to formation of long fibers. This effect has been explored for the performance improvement of blends comprising liquid crystal polymers, LCP [La Mantia, 1993].

509

7.5.4.5 Interlayer Slip The interlayer slip originates in low entanglement density at the interface [Helfand and Tagami, 1971, 1972]. Thermodynamics causes preponderance of chain ends and of the low molecular weight species in the interphase. This leads to the low viscosity — in binary PS/PMMA blends, the interphase viscosity was determined as: ηInterph
90 Pas, hence three orders of magnitude smaller than viscosities of the two comprised polymers [Valenza et al., 1991]. The net result of interlayer slip is a drastic reduction of viscosity for mixtures of two immiscible liquids. The phenomenon, first observed for mixtures of low molecular weight liquids, was empirically described using the fluidity-additivity equation. The latter dependence was first derived by Bingham [1922]:

(7.122)

where wi and ηi are the volume or weight fraction and the viscosity of the component i, respectively (see Figure 7.22). The dependence was re-derived 42 years later for a telescopic flow of two polymers through a pipe [Heitmiller et al., 1964]. The two liquids formed a large number of concentric layers, each of the same crossectional areas. The fundamental condition that leads to the fluidity additivity relation was the continuity of the shear stress across the multi-stratified structure. Lin [1979] followed this derivation with an additional assumption that the shear stress of each layer can be modified by the presence of an additional frictional stress, Z = (β-1)(R∆P/2L), where R is the capillary radius, ∆P is the pressure drop, and β is a characteristic material parameter in: (7.123) For a mixture of two liquids having the same viscosities, η1 = η2, Eq 7.122 predicts additivity,

50

L. A. Utracki and M. R. Kamal

while Eq 7.23 with β > 0 predicts a negative deviation form additivity (NDB). For β =  Bingham’s relation is recovered. However, as shown recently, there are serious reservations about the fundamental consequences of the frictional extra stress Z [Bousmina et al., 999]. In a rigorous derivation for a telescopic flow with the interfacial slip, the following dependence was obtained: 1 / η = ( w1 / η1 ) + ( w2 / η2 ) + θ φ1φ 2 where: θ ∝ 1 / ∆lη Interphase

(7.24)

Equations 7.23 and 7.24 predict a negative deviation from the log-additivity rule. The material parameter θ in Eq 7.24 governs the NDB behavior. It was shown that its value is inversely proportional to the thickness of the interphase, ∆l, and its viscosity, ηInterphase [Bousmina et al., 999]. Theoretically, the same molecular mechanism should be responsible for both factors, viz. better miscibility, better inter-diffusion, thus higher ∆l and ηInterphase. However, the low molecular weight components of the blend, that are forced by the thermodynamics to diffuse to the interphase, may not change much the former parameter, but drastically reduce the latter. For the immiscible blends, ∆l is small, typically 2-6 nm, thus θ is large and the interlayer slip takes place. For compatibilized blends, the macromolecules of the two phases interact and interlace, which increases both factors, and thus the slip effects are negligible. Measured or calculated values of the interphase viscosity are listed in Table 7.8.

Table 7.8.

Figure 7.22. Concentration dependence of shear viscosity of PP/LCP blends; dotted line represents the fluidity equation, Eq 7.22. Data from [Ye et al., 99].

The interlayer slip creates a tree-ring structure in extrusion, e.g., observed in samples containing 30 wt% PA-6 in HDPE matrix, extruded at T = 250°C. The HDPE/PA-6 capillary viscosities at 250°C followed Lin’s Eq 7.23 [Utracki et al., 986]. The simplest fluidity equation, Eq 7.22, may be useful in describing steady-state viscosity of antagonistically immiscible polymer blends, such as PP/LCP shown in Figure 7.22. When the volume of the interphase is known, the general form of the Bingham formula in Eq 7.22 can be used to calculate the interphasial viscosity. This indeed has been done in the case of shear flow of a multilayer PS/PMMA sandwich [Lyngaae-Jørgensen et al., 988].

Calculated viscosity of the interphase

Blend (1/2)

η1 (Pas)

η2 (Pas)

ηInterphase(Pas)

POM/CPA PP/PS PE/PS PS/PMMA PS/PMMA

349 24 4000 72 3400

583 693 582 460 5500

0.69 .64 25.22 3.3 90.00

Reference Bousmina et al., 999 Bousmina et al., 999 Bousmina et al., 999 Bousmina et al., 999 Valenza et al., 99

The Rheology of Polymer Alloys and Blends

7.5.4.6 Encapsulation Shear-induced segregation of polymer domains is related to a difference in magnitude between the rheological properties of blend components. During large strain flow, the segregation takes place not only in immiscible blends, in which the viscosities and elasticity of the two phases differ, but also in miscible blends comprising components of different MW. In the latter case, it is the difference in chain lengths that causes an imbalance of stresses and relative motion of the components [Doi and Onuki, 992]. Migration of the low viscosity component toward the high stress regions results in flow encapsulation of one phase by another. The effect has been well documented and successfully explored in polymer processing. For example, this mechanism is responsible for the lubricated, high throughput flow of PO’s upon addition of either fluoro- or siloxane polymers. Similarly, enhancement of flow of engineering and specialty resins by incorporation of low viscosity (when molten) LCP is based on the flow segregation [Utracki, 987, 988, 989]. 7.5.5

Shear Flows

The easiest way to discuss flow of polymer blends is to compare them to simpler, low molecular weight homologues, viz. Part 7.3. For immiscible blends, the best model is that of emulsions. Like blends, the emulsions comprise one liquid dispersed in another. The emulsion morphology is stabilized by addition of a surfactant or an emulsifier, similarly as immiscible blend is stabilized by addition of a compatibilizer. Both systems, emulsions and blends, show phase inversion, viz. Part 7.3.2. In emulsions, the phase inversion concentration, φΙ1, depends mainly on the type and concentration of emulsifier, while in blends it is dominated by relative rheological properties of the two polymers. In emulsion technology, by carefully selecting surfactants and the sequence of liquid addition, it is possible to generate (at the same concentration) two emulsions having different morphologies, viscosities and other properties.

5

Such a duality of behavior in polymeric mixtures has not been explored so far [Utracki, 989]. 7.5.5.1 Concentration Dependence of Viscosity It is instructive to generalize the observed flow behavior of polymer blends into specific categories. There is a correlation between the thermodynamic and rheological functions. In miscible blends where the free energy of mixing is negative, ∆Gm < 0, the experimental data indicate that in most systems either log-additivity rule (see Eq 7.8) or small positive deviations from it have been observed. Near the phase separation, where ∆Gm ≈ 0, the rheological response is complex as the free energy of mixing is precariously balanced by the term describing the energy input by the flow. Finally, in immiscible systems, where ∆Gm > 0, five different types of behavior have been identified. In Figure 7.23, curves -5 represent, respectively: . positively deviating blend, PDB, 2. negatively deviating blends, NDB, 3. log-additivity, 4. PNDB, and 5. NPDB [Utracki, 99]. To understand the origins of this behavior, it is necessary first to evaluate morphology and flowimposed morphology in polymer blends. From the discussion of phase inversion in Part 7..2, the emulsion model predicts that immiscible blends should show positive deviation, PDB, from the log-additivity rule: lnη = ∑ wilnηi. However, while PDB has been found in about 60% of blends, the remaining four types (see Figure 7.23) must also be accounted for. This means that at least one other mechanism must be considered when modeling the viscosity-concentration dependence of polymer blends. This second mechanism should lead to the opposite effect, that is to the negative deviation from the log-additivity rule, NDB. The simplest mechanism that explains the NDB behavior is the interlayer slip, which led to derivation of Eqs 7.23 and Eq 7.24. One may postulate that at constant stress, the net η vs. φ dependence can be written as a sum of two contributions: the interlayer slip, expressed by ηL (calculated from either Eq 7.23 or Eq 7.24), and the emulsion-like viscosity enhancement given by

52

L. A. Utracki and M. R. Kamal

an excess term, ∆log ηΕ [Utracki, 99]: ln η = ln η L + ∆ ln η E

{ [

2

(

∆ ln η E = ηmax 1 − (φ1 − φ1I ) / φ1φ 22I + φ 2φ12I

)]} (7.25)

Figure 7.23. Five types of the relation between shear viscosity and concentration for immiscible polymer blends: . PDB, 2. NDB, 3. additivity, 4. PNDB, and 5. NPDB [Utracki, 99]

As mentioned in Part 7., for polymer blends the relation between the steady-state shear viscosity and concentration can be quite complex. In the following discussion, the constant stress (not the constant rate) viscosity, corrected for the yield and time effects, will be considered. To illustrate flexibility of Equation 7.25 to describe (and thus to facilitate interpretation of the rheological results) η vs. φ dependence examples of computations are shown in Figures 7.24-7.3. The numerical values of the phase inversion concentration, φ2I, as well as the two material parameters that enter Eq 7.25 are listed in Table 7.9. 7.5.5.2 Dynamic Flow

Figure 7.24. Concentration dependence of blend viscosity at five levels of shear stress, (from top: σ12 = 0 to 05) indicating a gradual change of dominant flow mechanism from emulsion-type to interlayer slip.

Figure 7.25. Concentration dependence of blend viscosity for polymer-/polymer-2 blends; three different molecular weight grades of polymer-2 were used. For the lowest molecular weight PDB, whereas for the highest NDB behavior is to be expected.

Owing to the variability of blend structure with flow conditions, the rheological responses must be sensitive to the way that they are measured. Since modification of structure is related to strain, it is to be expected that responses measured at high and low strain values will be different. For this reason, the selected type of test procedure should reflect the final use of the data. When simulation of flow through a die is attempted, large strain capillary flow is useful; but, if the material characterization is important, low strain dynamic testing should be used. Because of morphology sensitivity to test conditions, there is serious a disagreement between predictions of the continuumbased theories and experiments. This is summarized in Table 7.0. The dynamic testing of polymer blends at small amplitude is a relatively simple and reliable procedure. The resulting storage and loss shear moduli, G’ and G”, respectively, should be first corrected for yield stress then the loss data can be fitted to Equation 7.42, to determine the value of the four parameters, ηo, τ, m, and m2.

The Rheology of Polymer Alloys and Blends

Figure 7.26. Concentration dependence of shear viscosity of PS/PMMA blends. Points are experimental [LyngaaeJørgensen, 1983], while the lines were computed from Equation 7.125.

Figure 7.28. Concentration dependence of η’ and η, shear viscosities of LLDPE/PC blends at 245°C and at constant stresses: G” = 1, 10 and 100 kPa and σ12 = 100 kPa, respectively. Points — experimental; error bars of measurements ± 2% [Utracki and Sammut, 1990].

Once these parameters are known, the Gross frequency relaxation spectrum can be calculated (see Eqs 7.85-7.87) and as a result all linear viscoelastic functions. Over the years, there has been an effort to describe the rheological functions of liquid mixtures from those of neat ingredients and their content. The theoretical treatment has been discussed, to some extent, in Part 7.3.2.3. A summary of these efforts is given in Table 7..

53

Figure 7.27. Concentration dependence of zero shear viscosity of polypropylene blends with two linear low density polyethylenes at 190°C. Points are experimental with error bars indicating the standard deviation [Dumoulin, 1988]. Lines are computed from Eq 7.125.

Figure 7.29 Concentration dependence of zero shear viscosity at 260, 280 and 300°C of polyethyleneterephthalate blended with polyamide-6,6. Points are experimental with the error bars indicating the error of measurements ± 2% [Utracki et al., 1982].

The last entry in Table 7. refers to a recent theory by Palierne [990]. The theory is based on the following assumptions: (i) the system consists of two viscoelastic liquids, (ii) the concentration of the dispersed phase is moderate, (iii) the drops are spherical, polydispersed and deformable, (iv) the drop deformation is small, so the blend behavior is linear viscoelastic, and (v) the interfacial tension coefficient, ν2, is constant, independent of stress and interfacial area.

54

L. A. Utracki and M. R. Kamal

Figure 7.30. Concentration dependence of zero shear viscosity of PS/PMMA blends at 180 and 210°C. Points are experimental [Valenza et al., 1991], solid lines are computed from Eq 7.125.

Table 7.9.

Figure 7.31. Concentration dependence of zero shear viscosity at 200°C of PS/LDPE blends without (solid line) and with (broken line) 5 wt% SEB. Points — experimental; error bars of measurement ±3% [Austin et al., 1987].

Parameters used for curve fitting of viscosity vs. concentration data to Equations 7.9, 7.23 and 7.25 [Utracki, 99]

System

PS/PMMA PS/PMMA PP/LLDPE- PP/LLDPE-2 LLDPE/PC LLDPE/PC LLDPE/PC LLDPE/PC PA-66/PET PA-66/PET PA-66/PET LDPE/LLDPE-I LDPE/LLDPE-II PS/PMMA PS/PMMA PE/PMMA PE/PMMA PS/PMMA PS/PMMA PS/LDPE PS/LDPE

Conditions

80°C 20°C 90°C 90°C G” =  kPa, 245°C G” = 0 kPa, 245°C G” = 00 kPa, 245°C Capillary flow, 245°C 260°C 280°C 300°C 90°C 90°C without compatibilizer with compatibilizer h’ at G” =  kPa, 60°C h at s2 =  kPa, 60°C 200°C at s2 = 50 kPa 200°C at s2 = 00 kPa with SEB without SEB

φI2 from Eq 7.9

0.886 0.820 0.998 0.093 0.306 0.47 0.68 0.68 0.928 0.882 0.796 0.956 0.288 0.7 0.888 0.93 0.888 0.69 0.804 0.827 0.827

Eqs 7.123 and 7.125 Parameters ηmax

β

ρ2

0.5597 0.4882 .738 .4890 0.9523 0.52 0.0365 0.64 0.5897 0.637 0.966 0.432 0.2972 .672 0.9822 0.473 0.3723 0.4495 .365 .933 .483

.6797 2.336 0.022 5.4634 2.4706 0.0 0.000 .8245 2.0757 4.044 8.6473 0.06 0.000 39.9755 6.057 0.000 0.000 4.66 4.6292 22.2605 .8930

0.9956 0.9882 0.9789 0.9804 0.9855 0.9929 0.9897 0.9903 0.9996 0.999 0.9968 0.9999 0.9825 ----0.9922 0.9965 0.9476 0.999 0.9754 0.9806

55

The Rheology of Polymer Alloys and Blends

Table 7.10. Comparison of continuum-based predictions for simple fluid with experimental observations for polymer blends Rheological function

Simple Fluid

Polymer Blend

Viscosity at vanishing

η(y) = η′(ω) = ηΕ(e)/3

η(y) ≠ η′(ω) ≠ ηΕ(e)/3

ηΕ = ηΕ(Cogswell)

ηΕ ≠ ηΕ(Cogswell)

N = N(Β-swell); Tanner, 970

N ≠ N(Β-swell)

deformation rates Extensional viscosity (from entrance effects) First normal stress difference (from extrudate swell)

Table 7.11. Blend Viscoelasticity from Emulsion Models [Graebling and Miller, 99] Author

Concentration

Liquids

Drops

Results

Taylor, 932 Oldroyd, 953 Palierne, 990

Dilute Dilute Moderately concentrate

Newtonian Newtonian Viscoelastic

Undeformable Deformable Deformable, Polydispersed

η = η(λ, φ) G’, G”, η = f(λ, φ, ν) G*, H = f(G*m, G*d, ν2/d, φ)

The theoretical analysis leads to Eq 7.70. Note that the ratio ν12/d is the only parameter of the equation. The model was found to provide good description of the dynamic behavior for several blends, supporting the idea that the long relaxation times in blends originate from geometrical relaxation of droplets [Graebling et al., 989, 993]. 7.5.5.3 Compatibilization Effects Most immiscible polymer blends require compatibilization, to reduce the interfacial tension, thus increase the degree of dispersion, to stabilize the morphology developed during compounding against extensive damage during high stress and strain processing (e.g., during injection molding), and to enhance adhesion between the phases in the solid state, facilitating the stress transfer and improving the mechanical properties of the

product. Compatibilization is achieved either by addition of a small quantity, 0.5-2 wt%, of a precisely tailored (usually) block copolymer; an addition of a multi-purpose core-shell copolymer, ” 35 wt%, that also improves toughness of the blend; or by reactive processing. Compatibilization strategy for either addition or reactive blending requires that the copolymer migrates to the interface, thus, on the one hand lowering the thermodynamic immiscibility barrier between the two phases, and, on the other hand, engendering formation of the third phase, the interphase. From the point of view of blend morphology in the molten state, compatibilization enhances the dispersion, increases the total apparent volume of the dispersed phase, rigidifies the interface, and increases interaction, not only between the two phases, but also between the dispersed drops. Furthermore, reactive compatibilization may involve

56

L. A. Utracki and M. R. Kamal

chemical bonding between the two polymer macromolecules, resulting in significant increase of the molecular weight at the interface. The rheological consequences of these changes can be predicted from a model system. The emulsion model indicates that making the interface more rigid causes the intrinsic viscosity of the emulsion to increase (see Eq 7.50). Similarly, an increase of the apparent volume of the dispersed phase causes the relative viscosity to increase (see Eqs 7.24-7.25). Furthermore, enhanced interactions between the phases will reduce the possibility of the interlayer slip, and increase formation of associative network formation, which may result in the yield stress. In short, compatibilization is expected to increase melt viscosity, elasticity and the yield stress. There are two mechanisms that may invalidate the above prediction: () In spite of the best efforts of researchers and technologist, the added copolymer may prefer to form micelles inside one of the polymeric phases than to migrate to the interphase. This has been frequently observed in blends with block copolymers, e.g., for blends of PS with PE, “compatibilized” by addition of a hydrogenated styrene-butadiene block copolymer, SEBS [Utracki and Sammut, 988, 990]. (2) Depending on the blend composition, the addition of compatibilizer may affect the total free volume of the system. These changes are difficult to predict. An increase of the free volume (evidenced by reduction of melt density) is expected to result in increased fluidity of the system. Effects of addition of hydrogenated styrenebutadiene di-block copolymers (one strictly diblock and the other tapered) on properties of HDPE/HIPS blends were investigated [Brahimi et al., 99]. The rheological behavior of the blends, especially in the low frequency region, was sensitive to the copolymer content. However, at high frequencies, the copolymer only slightly affected the flow. Furthermore, addition of a small amount of compatibilizer reduced the zero-shear viscosity, ηo, while higher loading had an opposite effect. This behavior was interpreted as due to the change in the copolymer state in the blend, i.e., saturation of the interface followed by micelle formation.

The results were compared with prediction of Palierne model, viz. Eq 7.39. For diluted, uncompatibilized blends, PS/PE = 0/90 or 90/0, relatively good agreement was found. The agreement was poor for blends containing 3 wt% of the tapered di-block copolymer. In the latter case, the reduction of storage and loss shear moduli, especially at low frequencies, could not be explained by the emulsion model. The effect of the interfacial tension and particle size over a relatively wide range, did not significantly affect model predictions. These observations confirmed an earlier report for the same system [Aït-Kadi et al., 992]. Since, at low frequencies, slip is less likely to occur than at high, low values of G’ and G” could not be explained by this mechanism. The most likely explanation is an increase of the free volume by incorporation of the copolymer. The non-linear variation of specific volume as a function of composition has been frequently observed for systems with limited miscibility. Blends of PS with LDPE were compatibilized by addition of di- and tri-block copolymers, Kraton™ G 1701 and G 1605 [Pascault et al., 994]. A continuous decrease of PS-drop size, and an increase of shear viscosity with addition of copolymer was reported. Blends of PP with a polyamide (PA-6, PA-66, or PA-2), have been the object of intensive studies. Linear viscoelastic shear moduli were measured for PP/PA-6 blends comprising different amount of PP-g-PA-6 copolymer. It was reported that, in spite of the expected reduction of the particle size with increase of the compatibilizer content, no qualitative effect of the flow was observed [Scholz et al., 989]. In a thorough study, PP, PA-6, and their compatibilized blend were studied at 225-250°C, in steady state and dynamic shear, as well as extensional flow conditions [Utracki and Sammut, 992]. The dynamic flow curves for the blend were significantly higher than what could be predicted from the component flow behavior. The blends showed regular, pseudoplastic flow behavior, without the yield stress. By contrast, capillary flow was found to be insensitive to temperature, suggesting a major modification of

The Rheology of Polymer Alloys and Blends

morphology during these large strain tests. Similarly, the extensional viscosity of the blends was one order of magnitude higher than what could be expected from the component polymer behavior. During extensional flow, the blends showed strainhardening, absent for either PP or PA-6. This could be explained by postulating that the reactive compatibilization significantly increased the molecular weight of the system — strain hardening is to be expected for highly entangled, high MW systems. It was also reported that the measured elongational viscosity, ηE, for two homopolymers agreed quite well with the value calculated from the entrance pressure drop in capillary flow, Pe. However, for the blend, the calculated value of ηΕ was one order of magnitude higher than measured. Reactive compatibilization in a specially designed twin-screw extruder was carried out compounding maleated polypropylene, PP-MA (0-0.4 wt% MA, with PA-6. During the reaction a di-block copolymer was formed at the interface. As the copolymer content increased from zero to 20 wt%, the number average diameter of PA-6 drops decreased from the initial value do = 20, to d = 0.4 µm at 20 wt% copolymer. The concentration dependence of shear viscosity also changed with compatibilization from negative deviation from the log additivity rule, NDB, to positive deviation, PDB [Nishio et al., 992]. The effect of compatibilization on the shear flow of PP/PA-6 and PP/PA-2 blends was also studied [Germain et al., 994]. Here the copolymer flow curve was one order of magnitude lower than that of PA. The authors reported that, at low deformation rates the emulsion-type morphology dominated the flow, whereas at higher rates concentric layered-type morphology, with appropriate flow behavior, has been observed. At low shear rates, the blend viscosity was higher than the viscosity of the matrix, while at high shear rates, the contrary was observed. The low shear rate behavior was analyzed by means of Palierne theory, assuming that the copolymer is located at the interface — good agreement was obtained for the low-concentration blends. For high stress (and strain) deformation a model of lamellar telescopic

57

flow for power-law fluids was derived. The skincore model represented the flow of blend well. Thus, it can be postulated that, in blends with low MW copolymeric compatibilizers, the stress may cause its disentanglement from at least one phase, formation of a layered morphology, resulting in flow lubricated by the presence of the low molecular weight compatibilizer. Effects of addition of a compatibilizing block copolymer, poly(styrene-b-methyl methacrylate), P(S-b-MMA) on the rheological behavior of an immiscible blend of PS with SAN were studied by dynamic mechanical spectroscopy [Gleisner et al., 994]. Upon addition of the compatibilizer, the average diameter of PS particles decreased from d ≅ 400 to 20 nm. The data were analyzed using weighted relaxation-time spectra. A modified emulsion model, originally proposed by Choi and Schowalter [975], made it possible to correlate the particle size and the interfacial tension coefficient with the compatibilizer concentration. It was reported that the particle size reduction and the reduction of ν12 occur at different blockcopolymer concentrations. This brief summary of compatibilizer effects on the flow behavior of polymer blends substantiate that many aspects of the behavior remain unclear. For example, the effects of limited miscibility of copolymer, the copolymer structure and molecular weight, micelle formations and inter-particle interactions are still beyond the capability of the proposed models. Similarly, little information has been published on copolymer/homopolymer blends. 7.5.5.4 Time-temperature Superposition The time-temperature superposition principle, t-T, has been a cornerstone of viscoelastometry. It has been invariably used to determine the viscoelastic properties of materials over the required 0 to 5 decades of reduced frequency, ωaT [Ferry, 980]. Measuring the rheological properties at several levels of temperature, T, over the experimentally accessible frequency range (usually two to four decades wide), then using the t-T shifting, made it possible to construct the complete isothermal function.

58

L. A. Utracki and M. R. Kamal

As demonstrated before, the shifting involves three shift factors, one horizontal, usually expressed as: aT = bTη0(T)/η0(T0), where bT = ρ0T0/ρT is the first vertical shift factor that originates in the thermal expansion of the system (ρ is density). The subscript o indicates reference conditions, defined by the selected reference temperature To, usually taken in the middle of the explored T-range. For homopolymer melts, as well as for amorphous resins, the two shift factors, aT and bT, are sufficient. However, for semi-crystalline polymers, a second vertical factor, vT has been found necessary — it accounts for variation of the crystallinity content during frequency scans at different temperatures [Ninomiya and Ferry, 967; Dumoulin, 988]. Only when all the relaxation times in a given system are multiplied by the same factor when the temperature is changed, the t-T principle can be observed. In single-phase homologous polymer blends, the relaxations are mainly controlled by the segmental mobility. Thus, t-T superposition has been observed in a wide range of conditions. Similarly, for polymers filled with high modulus particles, the filler is responsible for enhancement of modulus without affecting the relaxation spectrum, and as a consequence, the t-T superposition principle is obeyed. However, in rheologically complex heterogeneous systems, individual polymeric components contribute to the relaxation and since their activation energies are usually different, a change of temperature affects them differently — resulting in poor t-T superposition. Fesko and Tschoegl [97] demonstrated that the simplest form of the time-temperature superposition relation for a function G(t, T) = G[t, a(T)], is:

[∂G(t, T ) / ∂T ]t = {[∂G(t, T )] / [∂lnt ]T }{d[lna(t, T )] / dT } d[lna(t, T )] / dt = ∑ Ni (t )d[lnai (T )] / dT ; i = 1,.., n i

Ni (t ) = φi Li (t, To ) / ∑ φi Li (t, To ) i

Li (t, To ) = ∂Gi [(lnt ), To ] / ∂lnt

(7.26)

where the summation extends over every species in the system. Equation 7.26 assumes that the effects of time and temperature can be separated, G(t, T) = G[t•a(t, T)]. A similar derivation was published by Goldman et al. [977]. There is growing evidence that t-T superposition is not valid even in miscible blends well above the glass transition temperature. For example, Cavaille et al. [987] reported lack of superposition for the classical miscible blends — PS/PVME. The deviation was particularly evident in the loss tangent vs. frequency plot. Lack of t-T superposition was also observed in PI/PB systems [Roovers and Toporowski, 992]. By contrast, mixtures of entangled, nearly monodispersed blends of poly(ethylene-alt-propylene) with head-to-head PP were evaluated at constant distance from the glass transition temperature of each system, homopolymer or blend [Gell et al., 997]. The viscoelastic properties were best described by the “double reptation model”, viz. Eq 7.82. The data were found to obey the timetemperature superposition principle. The explanation proposed by Ngai and Plazek [990], was based on the postulate that the number of couplings between the macromolecules varies with concentration and temperature of the blend. The number of couplings, n, can be calculated from the shift factor, aT = [ζo(T)/ζo(To)]/(-n), where ς0(T) is the Rouse friction coefficient. Thus, in miscible, single phase systems, as either the concentration or temperature changes, the chain mobility changes and relaxation spectra of polymeric components in the blends show different temperature dependence, i.e., the t-T principle cannot be obeyed. Similar conclusions were reached from a postulate that the deviation originates from different temperature dependence of the relaxation functions of the blend components [Booij and Palmen, 992]. In immiscible blends, the t-T principle does not hold. For immiscible amorphous blends it was postulated that two processes must be taken into account: the t-T superposition, and the aging time [Maurer et al., 985]. On the other hand, in immiscible blends, at the test temperature, the polymeric components are at different distance

519

The Rheology of Polymer Alloys and Blends

from their respective glass transition temperatures, T - Tg1 ≠ T - Tg2. In blends of semicrystalline polymers, such as PE/PP, the superposition is limited to the molten state, within a narrow, high temperature range [Dumoulin, 1988]. As an alternative to the t-T superposition, a plot of the elastic stress tensor component as a function of the viscous one has been used, e.g., (σ 11 - σ 22) vs. σ12, or G’ vs. G”. For systems in which the t-T is obeyed, such plots provide a temperature-independent master curve, without the need for data shifting and calculating the three shift factors. Indeed, from the Doi and Edwards tube model, the following relation was derived: ln G’ = 2ln G” + ln (6Me / 5ρRT)

Chuang and Han [1984] reported that for miscible and immiscible blends at constant composition, the plots of N1 vs. σ12 and G’ vs. G” are independent of T. However, while for single phase systems the two dependencies are approximately parallel, for immiscible blends, such as PS/PMMA, the steady state relation may be quite different from the dynamic one. The agreement can be improved by means of the Sprigg’s theory [1965]. The general theory leads to the conclusion that:

(7.127)

where Me is the entanglement molecular weight [Han and Kim, 1993]. The dependence suggests that a plot of G’ vs. G” should be insensitive to temperature. Indeed, good superposition was obtained for several blends where the structure remained unchanged within the range of independent variables, e.g., in such miscible systems as PS/PVME and PEO/PMMA, or even in some immiscible blends whose components have similar glass transition temperature, viz. PS/PMMA. However, there was lack of superposition in other systems where the structure did change, viz., PS/ PVME heated across the binodal, block copolymer across the microphase separation temperature, LCP across the nematic transition temperature, etc. 7.5.5.5 Steady State vs. Dynamic Viscosities For most blends, the morphology changes with the imposed strain. Thus, it is expected that the dynamic low strain data will not follow the pattern observed for the steady-state flow. One may formulate it more strongly: in polymer blends, the material morphology and the flow behavior depend on the deformation field, thus under different flow conditions, different materials are being tested. Even if low strain dynamic data can be generalized using the t-T principle, those determined in steady state will not follow the pattern.

(7.128) 2

1/2

where: C ≡ ω / γ˙ = [(2 - 2e - e ) / 3] and e is a model parameter. For a series of PMMA/ABS blends the plot of C vs. composition was non-linear, with C = 1 found only for PMMA homopolymer. Variation of this structural parameter seems to be related to differences of morphology existing in dynamic and steady-state flow fields [Utracki, 1989]. Capillary flow of EPDM with poly(vinylideneco-hexafluoropropylene), Viton™, showed a sixfold reduction of shear viscosity upon addition of about 2% of the other component [Shih, 1976, 1979], whereas in dynamic tests the complex viscosity behavior of EPDM and EPDM with 5% Viton™ was similar over a wide range of frequency and strain [Kanu and Shaw, 1982]. The latter authors postulated accumulation of the second component at the capillary entrance, which periodically feeds into the capillary, lubricating the main stream by a sort of roll bearing effect. In this particular case, the difference is related not only to material properties but also to a flow segregation enhanced by the geometry of the measuring device. Since the effect is strongly affected by flow geometry, the data obtained in capillary flow have little value for process design requirements. The phenomenon of flow segregation has been commercially explored. For example, a high vis-

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cosity engineering resin that has poor resistance to solvents, e.g., polycarbonate or polyetheretherketone (PC or PEEK, respectively), can be blended with low melt viscosity liquid crystal polymer, LCP. Extruding such a blend through die with long enough land, forces the LCP to migrate toward the high stress surface, thus lubricating the die flow, and in addition engendering a protective layer on the surface of PC or PEEK. The first effect, the reduction of viscosity of polymer melt upon addition of LCP, was originally described in 979, in a patent deposition from ICI [Cogswell et al., 98; 983; 984]. The rheological behavior of LCP blends with poly(ether imide), Ultem™ was studied by Nobile et al. [990] in steady state capillary and dynamic mechanical modes of deformation. The flow of LCP was reported to be sensitive to pressure, thus the flow of blends was carried out using a short capillary. Another, more common commercial use of the phenomenon is the addition of fluoropolymers to polyolefins. In this case, a small amount of fluoropolymer progressively migrates to the die surface, reducing the die pressure drop and making it possible to extrude the resin at high throughput without the melt fracture. It has been shown that this approach also works for other polymers, viz. PEEK. Thus, blends of PEEK with polytetrafluoroethylene, -5 wt% PTFE, were extruded. The pressure drop across the die was reported to decrease with time to an equilibrium value, Plim. The value of Plim depended on PTFE content, whereas the time to reach it depended on the rate of extrusion — the higher was the rate, the shorter was the saturation time [Chan et al., 992]. Over the years, dynamic testing has become the preferred method of testing the rheological behavior of the multiphase systems. For example, Nishi et al. [98] carried out careful studies on the dynamic behavior of PS/PVME. The specimens were cast at temperatures either below or above the lower critical solution temperature, LCST ≅ 95°C. While those prepared at T < LCST (single-phase system) showed superposition of dynamic data onto a master curve, the ones that were cast at T > LCST did not.

7.5.5.6 Blend Elasticity Four measures of melt elasticity are commonly used: in steady state shearing, the first normal stress difference, N, in dynamic tests, the storage modulus, G’, and the two indirect and controversial ones, namely entrance-exit pressure drop (Bagley correction), Pe, and the extrudate swell, B. In homogeneous melts, the four measures are in qualitative agreement. More complex behavior is expected for blends. If the blend can be regarded as an emulsion, without interlayer slip, PDB behavior for the elastic measurements is to be expected. On the other hand, in systems where the dispersed phase is difficult to deform (as in suspensions), the extrudate swell should be small. Deformation-and-recovery of the dispersed phase shape provides a potent mechanism for the elastic energy storage that results in large elastic response (this does not have anything to do with the molecular energy storage). Direct measurements of N and σ2 indicate a parallel dependence of both these functions plotted vs. φ, even when these have a sigmoidal form. Considering the steady shear flow of a twophase system, it is generally accepted that the rate of deformation may be discontinuous at the interface, and it is more appropriate to consider variation of the rheological functions at constant stress than at constant rate, i.e., N = N(σ2). Using a similar argument for the dynamic functions, it should be concluded that G’ = G’(G”) should be used. Note that, as discussed above, the steady-state and dynamic data for polymer blends rarely superimpose. Another method for estimating the elasticity contribution is through the Bagley entrance-exit pressure drop correction, Pe. For single-phase systems, the plot of Pe versus σ2 is independent of capillary diameter, temperature and molecular weight, but rather sensitive to changes in flow profile [Utracki, 985]. The plot was found to be useful for interpretation of the stress, temperature and composition dependent morphological changes in immiscible polymer blends [Dumoulin et al., 985]. However, it could not be used to estimate the elasticity of blends.

The Rheology of Polymer Alloys and Blends

Extrudate swell, B, has been used to calculate the recoverable shear strain, γR, for single-phase materials [Utracki et al., 975]. Introduction of the interface negates the basic theoretical assumptions on which the calculation of γR was based. In addition, presence of the yield stress, frequently observed in multiphase systems, prevents B from reaching its equilibrium value required to calculate γR and then N. Nevertheless, B is used as a qualitative measure of blend elasticity. Note that the presence of the dispersed, deformable phase leads to form recovery, i.e., shrinkage of the prolate ellipsoids motivated by the interfacial energy, which results in unduly large enhancement of B. The phenomenon has little to do with deformability of macromolecular coils — the postulated mechanism of swelling in singlephase polymeric systems. The main origin of blend swelling is the elastic recovery of domains extended during the convergent flow in the capillary entrance. The observed form swelling of blends has been large, giving strain recovery, γR ” 6.5. It is worth pointing out that strain recovery can be non-symmetrical as far as the blend composition is concerned. In the case when the viscosity ratio at low deformation rate exceeds the limiting value λcr = 3.8, there is a significant difference in the mechanism of dispersion. During compounding in a twin-screw or twin-shaft instrument, the material undergoes stretching in shear and extensional flow fields. Depending on the composition, the stretched forms undergo different recovery. For blends with λ > λcr the prolate ellipsoids cannot be broken, thus they slowly retract into large spheres. For blends with λ < /λcr the prolate ellipsoids will disintegrate by the capillarity instability mechanism [El Khadi et al., 995]. Similar observations were recently reported for PMMA/PS blends [Gramespacher and Meissner, 995]. The elastic creep compliance for PMMA/ PS = 6/84 behaved regularly, similar to what has been observed for single-phase polymers. However, when the composition was reversed, i.e., PMMA/PS = 84/6, the recovery creep compliance showed a maximum at which the recovery direction was reversed. The authors attributed the

52

dissymmetry of behavior to different retardation times of the blend components. 7.5.6

Elongational Flows

Owing to experimental difficulties, there are but few publications on uniaxial deformation of blends. To prepare specimens for testing, samples usually are transfer molded and relaxed, both operations requiring relatively long heating time, during which only well stabilized blends will not coarsen. It is convenient to distinguish two contributions to the tensile stress growth function, η+Ε, one due to the linear viscoelastic response, η+ΕL, and the other originating in the structural change of the specimen during deformation, η+ΕS. The first can be calculated from any linear viscoelastic response, while the second (which originates in either intermolecular interactions or entanglements), depends on both the total strain, ε = et, and either strain rate e or straining time, t. Owing to the industrial importance of strain hardening (SH), a large body of literature focuses on the optimization of blend composition to maximize SH. Since SH depends on entanglement, blending branched polymers usually affects SH even in the low concentration range. Most of the work on uniaxial extensional flow of immiscible polymer blends has focused on the behavior of systems containing PE. The main reasons are the need for better, easier-to-process film resins, and for a stable polyolefin blend morphology. Film blowing conceptually involves two different engineering operations, extrusion and blowing. For most production lines, the latter limits productivity. For the low density LDPE resin, strain hardening provides a self-regulating, self-healing mechanism. For HDPE and LLDPE, only small SH can be obtained for the high MW and narrow MWD resins. As a result, most LLDPE resins on the market are blends with LDPE, rubbers, copolymers, or another type of LLDPE. SH was also found to be an important resin characteristic for wire coating. Here the surface finish and uniformity of the deposited layer were superior for blends with high strain hardening and low shear viscosity [Utracki, 988].

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Blends of LLDPE/PP = 50:50, with or without compatibilizing ethylene-propylene copolymer, EPR, were studied by Dumoulin et al. [984]. In spite of the expected immiscibility, the blends showed additivity of properties with good superposition of the stress growth functions in shear and elongation, as well as with the zero deformation rate Trouton ratio, RT ≅ . In earlier work, blends of medium density PE (MDPE) with small quantities of ultra-high molecular weight polyethylene (UHMWPE) were studied in shear and extension. Again, SH and RT ≅  were observed. It has been shown that the stress growth function in uniaxial extension provides three important pieces of information on the polymer. The initial slope of the stress growth function:

Si = lim d ln η E+ / d ln t

(7.29)

t→o

was found to correlate with polydispersity of the molecular weights, Mz/Mn, where Mz and Mn are respectively, z- and number average molecular weights [Schlund and Utracki, 987]. The observation agrees with Gleissle’s principle [980]: . η γ  = η + (t )

∴

for

.

γ =1/ t

. . lim ∂lnη + (t ) / ∂lnt = lim ∂lnη γ  / ∂ln γ = n − 1 . t →o

Figure 7.32. Relative deformation of the dispersed phase vs. the viscosity ratio [Delaby et al. 994].

γ →∞

(7.30) The plateau or equilibrium value provides information on the weight average molecular weight, Mw, and the stress hardening part, η+ΕS(t) on the entanglements, i.e., branching, association, etc. The Si provides means for detection of blend miscibility. Solubility usually broadens the width of the MW distribution, causing Si to increase. By contrast, immiscibility causes separation of high molecular weight fractions, narrowing MWD and decreasing its value. The miscibility can also be reflected in a maximum strain at break, εb. In ‘antagonistically’ immiscible blends of PA-6 in LLDPE, a sharp decrease of εb was observed. However, in blends with co-continuous morphology εb may increase to an average value, with negative deviations on both sides [Min, 984].

Li et al. [990] studied the elongational viscosity of specially prepared blends of styreneacrylonitrile copolymer, SAN, with different loading of crosslinked polybutadiene, BR, particles having diameter, d = 70 nm. At higher rates of extension, the SAN showed strong strain hardening behavior. As the volume of rubber particles in the system increased, the strain hardening became less evident. By contrast, the low deformation rate elongational viscosity was shown to increase with BR loading, and as its content increased, the system progressively showed increased sensitivity to strain. Delaby et al. [994] attempted to experimentally examine the relation between droplet deformability and the extensional viscosity ratio, DE = DE(λΕ). As the authors demonstrated, the theories proposed by Taylor [932, 934] and Palierne [990] predict the same dependence:

(γ d − 1) / (γ m − 1) = 5 / (2λ E + 3)

(7.3)

where γd and γm are respectively strain of the dispersed phase (defined as a ratio of the long axis to the original drop diameter) and the matrix. Co-deformation of drop is to be expected only for λE = . The dependence is illustrated in Figure 7.32.

The Rheology of Polymer Alloys and Blends

The experiments show good agreement with the behavior predicted by Eq 7.3 in the full range of the viscosity ratios used in the studies, λE = 0.005 to 3. Drop deformability computed in 2D, using the boundary element method, resulted in higher value than 5/3, given for 3D by Eq 7.3; independently, two research teams found the limit for λE → 0 to be 2 [Khayat et al., 996; Stradins and Osswald, 996].

523

different behavior under different extensional flow field conditions. To close the part on extensional flow, it seems pertinent to stress its importance for the dispersion process. As the microrheology indicates, the minimum of the κ vs. λ curve is narrow for shear flow, but very broad (and lower) for extensional flow (see Figure 7.34). This suggests that it should be easier to disperse fluids in extensional than in shear flow fields. In particular, the benefits are obvious for high and low values of the viscosity ratio. It has been demonstrated theoretically and experimentally that drops can be deformed but not broken by shear flow if λ > 3.8. This is not the case for the inverse ratio, but the drop spinning mechanism in this region is slow, which makes the dispersion process inefficient. These theoretical findings have been confirmed by constructing an extensional flow mixer, EFM, and performing experiments with polymer blends having large difference of viscosity [Luciani and Utracki, 996; Utracki and Luciani, 996].

Figure 7.33. Entrance-exit pressure drop in capillary flow of polypropylene, polyamide-6, and commercial PP/PA-6 blend, Orgalloy at 230°C — line computed from Eq 7.32.

Convergent flow at the die entrance provides strong elongational flow. In 989, Laun and Schuch, discussing the approximate methods of ηE determination, derived the following relation between the entrance pressure drop in capillary flow and the shear stress at the capillary wall: 1

Pe = 1.89(η E / η) 2 σ 12 / (n + 1)

(7.32)

For Newtonian liquids Eq 7.32 predicts Pe § .64σ2. This proportionality is shown in Figure 7.33 along with the experimental values determined for ORGALLOY™, PP and PA-6. Evidently, Eq 7.32 provides satisfactory approximation for the homopolymers, but for the blend, the prediction is about one decade too low. Similarly, as for the shear flows, here also the elongational properties of blends show

Figure 7.34. Critical capillarity number vs. viscosity ratio is shear flow (solid lines) and extension (dash line).

7.6

Conclusions

An attempt has been made to underline the common features of multiphase flows. It is shown that, to obtain non-Newtonian behavior, it is not necessary to have systems containing non-Newtonian

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liquids. Strong interparticle interactions, generated by interface modification or concentration, provide sufficient conditions for the complex rheological behavior of suspensions or emulsions based on simple fluids, e.g., apparent yield stresses exist not only in suspensions but also in emulsions and foams. These simpler systems provide an introduction to the rheology of polymer alloys, blends and composites. Here, the flow is not only controlled by the interphase conditions, but also by the viscoelastic properties of each phase. In multiphase systems, there is a reciprocal relation between morphology and rheology. During flow, strains change the morphology, engendering diverse rheological responses, such as normal stress in Newtonian liquid emulsions. The most efficient method of structure modification is extensional flow, but large strains in Couette or Poiseuille flow also affect the morphology. To study material behavior with well-defined structure, only low strain dynamic methods provide the answer. For generating data to be used in modeling or interpretation of die flow, one must simulate not only flow conditions but also the system history as well as the geometry to scale. Since miscibility changes with temperature and pressure, flow of some blends may be sensitive to these variables.

7.7

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Van Arsdale, W. E., J. Rheol., 26, 477 (982). van de Ven, T. G. M., “Colloidal Hydrodynamics”, Academic Press, New York (989). van de Ven, T. G. M., Polym. Compos., 6, 209 (985). Van Gisbergen, J. G. M., Hoeben, W. F. L. M. and Meijer, H. E.H., Polym. Eng. Sci., 31, 539 (99). Van Oene, H., J. Colloid Interf. Sci., 40, 448 (972). Vincent, M. and Agassant, J. F., Polym. Compos., 7, 76 (986); in Two-Phase Polymer Systems, Utracki, L. A., Ed., Progr. Polym. Process., Vol. 2, Hanser, Munich (99). Vincent, M., Dr. Eng. Thesis, Ecole Nationale Superieure des Mines de Paris, Sophia Antipolis (984). Vinckier, I., Moldenaers, P., and Mewis, J., J. Rheol., 40, 63 (996); ibid., 41, 705 (997). Vinogradov, G. V., Krasnikova, N. P., Dreval, V. E., Kotova, E. V., and Plotnikova, E. P., Int. J. Polym. Mat., 9, 87 (982). Vlassopoulos, D., Koumoutsakos, A., Anastasiadis, S. H., Hatzikiriakos, S. G., and Englezos, P., J. Rheol., 41, 739 (997). Vlassopoulos, D., Rheol. Acta, 35, 556 (996). Wakita, N., Polym. Eng. Sci., 33, 78 (993). Walling, N., M. Eng. Thesis, Chemical Engineering Dept., McGill University (995). Walling, N., and Kamal, M. R., Adv. Polym. Technol., 5, 269 (996). Wang, K. J., and Lee, J. L., J. Appl. Polym. Sci., 33, 43 (987). White, J. L., Czarnecki, L., and Tanaka, H., Rubber Chem. Technol., 5 823 (980). Wieckowski, A., and Streg, F., Chemia Stosowana, 1B, 95 (966). Wildemuth, C. R., and Williams, M. C., Rheol. Acta, 23, 627 (984). Willis, J. M., and Favis, B. D., Polym. Eng. Sci., 28, 46 (988). Wisniewsky, C., Marin, G., and Monge, Ph., Eur. Polym. J., 21, 479 (985). Wolf, B. A., Makromol. Chem., Rapid Commun., 189, 63 (980); Macromolecules, 17, 65 (984). Wu, D., Tang, N., Gu, D., and Wen, W., Intern. Polym. Process., 5, 47 (990). Wu, R., Shaw, M. T., and Weiss, R. A., J. Rheol., 36, 605 (992). Wu, S., J. Polym. Sci., Part B, 25, 557 (987); Polym. Eng. Sci., 27, 335 (987); Polymer, 28, 44 (987). Xanthos, M., Young, M. W. and Biesenberger, J. A., Polym.

The Rheology of Polymer Alloys and Blends

Eng. Sci., 30, 493 (990). Yakovlev, K. V., Zhitomirets, R. I., Romankevich, O. V., Zabello, S. E., and Yudin, A. V., Khim. Tekhnol. (Kiev), 5, 4 (984). Yalmazer, U., and Kalyon, D. M., J. Rheol., 33, 97 (99). Yang, H. H., Han, C. D., and Kim, J-K., Polymer, 35, 503 (994). Yang, H., Shibayama, H., Stein, R. S., Shimuzu, N., and Hashimoto, T., Macromolecules, 19, 667 (986). Ye, Y.-C., La Mantia, F. P., Valenza, A., Citta, V., Pedretti, U., and Roggero, A., Europ. Polym. J., 27, 723 (99). Zaldua, A., Munoz, E., Pena, J. J. and Santamaria, A., Polym. Eng. Sci., 32, 43 (992). Zeichner, G. R., and Patel, P. D., 2nd World Congress Chem. Eng., Montreal, QC, Canada, 6, 333 (98).

H (ω) HG (ω) Hmax hc kB L M Mn, Mw, Mz N = σ − σ22 No,N

7.8

Notation and Abbreviations

No, N+

7.8.

Notations (Roman letters)

NT

a, b, k, K, n, u, β - equation constants av - interfacial area density BAB - reduced binary thermodynamic interaction parameter, BAB = χΑΒRT/V B - droplet width bi - segment length ci - Lee and Park relaxation parameters D - deformation d - droplet diameter d* - equilibrium droplet diameter D, DM, Ds - diffusion, inter diffusion and self diffusion coefficient, respectively Dp - particle diffusion coefficient E+ - threshold energy of coagulation EDK - macroscopic bulk breaking energy F - intrinsic thermodynamic function f - frequency G*, G’, G” - complex, storage and loss shear modulus, respectively H (τ) - relaxation time spectrum

n ni P Pe pr q R Re R(q) r 〈r2Θ〉/2 rad rN S(q), S0 So, S

s

535

- reduced frequency relaxation spectrum - Gross’ frequency relaxation spectrum - maximum of the relaxation spectrum - critical separation distance - Boltzmann constant - droplet length - Onsager-type mobility factor - number, weight and z-average molecular weight, respectively - first normal stress difference - initial and final number of particles respectively - number of coagulating drops, initially and at t = tc - total number of collisions per unit time - number of particles - number of moles unit volume - pressure - Peclet number - probability that two particles that have collided result in coalescence - wave vector, or sinusoidal distortion - ideal gas constant - Reynolds number or real part of a complex function - fluctuation function - reduced drop radius - unperturbed, average radius of gyration - radian - radius of the critical nucleus - virtual structure function - interfacial area per unit volume of the blend for monodispersed spherical particles before and after coalescence, respectively - spinodal

536

L. A. Utracki and M. R. Kamal

T Tg t tb

-

tb* tc UCST

-

V Vx/V

-

z

-

7.8.2

absolute temperature glass transition temperature time necessary time for breakup of droplets dimensionless breakup time coalescence time upper critical solubility temperature volume volume fraction of emulsion undergoing uniform shear reduced frequency, f • τ

Notation (Greek letters)

α αo β12 χΑΒ

-

∆

-

∆Ε

-

∆Gm ∆Hm, ∆Sm

-

ε, e

-

φ, φ2

-

φc

-

φi

-

φm

-

φperc γ, y

-

η ηo

-

orientation angle the distortion at t = 0 interlayer slip factor binary thermodynamic interaction parameter between polymers A and B thermodynamic distance from the spinodal; ∆ ≡ 2((χΑΒN)s - (χΑΒN) activation energy, e.g., of flow: ∆Εη Gibbs free energy of mixing enthalpy and entropy of mixing, respectively Hencky strain and Hencky strain rate in extension, respectively volume fraction of dispersed and matrix phase, respectively volume fraction of the crosslinked monomer units volume fraction of phase i at phase inversion maximum packing volume fraction percolation threshold shear strain and rate of shearing, respectively viscosity zero-shear viscosity

ηr [η] η’ η* η, η2

-

κ = σijd/ν2 κcrit Λ λ = η/η2 λs νο2

-

ν2

-

ρ ρd σ σ σ− σ22 = N σ2 sm σy σyo τ t* Ω(Λ, λ)

-

ω ωmax

-

ωx ψ, ψ2

-

relative viscosity intrinsic viscosity dynamic viscosity complex viscosity viscosity of dispersed and matrix phase, respectively capillarity number critical capillary number distortion wavelength viscosity ratio wavelength interfacial tension in a quiescent blends interfacial tension coefficient between phase  and 2 density droplet density stress extensional stress first normal stress difference shear stress stress in the matrix phase yield stress permanent yield stress relaxation time mean relaxation time tabulated function for capillary instability angular frequency frequency at which H(ω) is maximum crossover frequency first and second normal stress difference coefficient, respectively

537

The Rheology of Polymer Alloys and Blends

7.9

Appendices

7.9.

Appendix 7.1. Use of Rheology for Analysis of Polymer Blends

No. Study

Blends

Description

Reference

. Phase inversion Co-continuity Percolation

Empirical description Theory HDPE/PS

viscosity ratio dependence viscosity ratio viscosity ratio dependence

Phase inversion Co-continuity 2. Interphase theories

Theory Theory Immiscible blends

viscosity ratio dependence fibrillar morphology density profile

Paul & Barlow, 980 Metelkin & Blekht, 984 Lyngaae-Jørgensen & Utracki, 99 Utracki, 99 Luciani, 993 Helfand & Tagami, 97, 972; Helfand & Sapse, 975 Noolandi, 984 Utracki, 99; 992

PS/PB

3. Theories

PMMA/PS PMMA/SAN Low molecular weight solutions homologous polymer blends Thermodynamics

Several blends Miscible blends

3. Detection of miscibility

Theory of miscible blends Miscible blends

PPE/PS PMMA/PEG PS/PVME

compatibilized compatibilized - semi experimental, η vs. φ. compatibilized with styrene- butadiene block copolymer. interface thickness, T-dependence MW - dependence MW - dependence

Anastasiadis et al., 988, 989; Anastasiadis & Koberstein, 988 Kressler et al., 993 Glasstone et al., 94; McAllister, 960 Friedman & Porter, 975

PVT - dependence

Simha & Somcynsky, 969; Simha & Jain, 984 η vs. PVT and concentration Utracki, 983 free volume theory; Steller & Zuchowska, η vs. φ dependence. 990 viscoelasticity Tsenoglou, 988 dynamic shear; relaxation spectra; linear viscoelasticity dynamic flow, MW ratio =  to 00 dynamic shear behavior thermodynamic influence, dynamic shear flow

Utracki & Schlund, 987

Prest & Porter, 972; Araujo & Stadler, 988 Colby, 989; Booij & Palmen, 992 Schneider & Brekner, 985; Brekner et al., 985; Yang et al., 986

538

L. A. Utracki and M. R. Kamal

Appendix 7.1. Continued No. Study

Blends PS/PVME

4. Miscible blends PP/SEBS PE/PE Theory Theory PS/PVME

SAN/PMMA PS/PB

PS/PVME

5. Theory of immiscible blend

PS/PIB Theories

Theory

PS/LLDPE

Theory

Description phase separation, time temperature superposition viscosity dependence on composition dynamic shear phase separation, interactions stress/composition, encapsulation molecular weight ratio dependence

stress induced changes in miscibility strain compatibilization with block copolymers low steady state stress theory of strain demixing shear induced effects on miscibility, phase separation shear induced mixing viscosity - concentration correlation

Reference Ajji et al., 989; Ajji & Choplin, 99; Mani et al., 992 Steller & Zuchowska, 990 Utracki & Schlund, 987 Rangel-Nafaile et al., 984; Larson, 992 Metzner et al., 979; Doi & Onuki, 992 Caville et al., 987; Stadler et al., 988; Ajji et al., 989; Ajji & Choplin, 99; Mani et al., 992 Silberberg & Kuhn, 952; Wolf, 980, 984 Lyngaae-Jørgensen, 985 Hashimoto et al., 990 Horst & Wolf, 99, 992, 994 Mazich & Carr, 983; Katsaros et al., 986; Nakatami et al., 990 Wu et al., 992 Helfand & Fredrickson, 989; Doi & Onuki, 992; Utracki, 99 Doi & Ohta, 992; Takahashi et al., 994

constitutive equation: : mixture, Newtonian fluids, same viscosity and density; orientation and interfacial tension effects at interface. constitutive equation: orientation Lee & Park, 994 & ν ; interface effects, morphology in shear. constitutive equation generalizing Grmela & Aït-Kadi, 994 blends behavior on the basis of the thermodynamic and dissipative potentials

539

The Rheology of Polymer Alloys and Blends

Appendix 7.1. Continued No. Study

6. Microrheology deformation and breakup.

Blends

Description

PS/PE, Comparison of Palierne, Lee & PETG/EVAc, Park, and Grmela & Aït-Kadi PP/EVAc, theories with experimental data PP/EVAc+PEMA

Lacroix et al., 997, 998

Immiscible blends viscoelastic drops in Newtonian medium.

Gauthier et al., 97; Prabodh & Stroete, 99; De Bruijn, 989 Flumerfelt, 972

Newtonian drops in viscoelastic fluid. both medium viscoelastic

PMMA/PS

7. Flow imposed morphology, FIM. TSE TSE

PP/EPDM

PS/PE PS/PE Polyolefins

8. FIM Fiber formation. PP/PA-6 POM/CPA POM/PEVAc HDPE/PA-6 9. FIM Interlayer slip

Reference

viscoelastic liquids, extensional and non-uniform flow Newtonian and viscoelastic filaments. break up and coalescence shear induced coalescence molecular diffusion, shear induced flow induced collision probabilities, coalescence.

Wu, 987; Van Oene, 978 Han & Funatsu, 978; Chin & Han, 979 Bousfield et al., 986 Lyngaae-Jørgensen et al., 993 Roland & Bohm, 984 Ratke & Thieringer, 985 Tokita, 977; Fortelny et al., 988, 990

steady state shearing Huneault et al., 995 blending in a twin screw extruder. Shi & Utracki, 992 Huneault et al., 995 Plochocki, 978, 983 dispersion Utracki et al., 986 fibrillation Krasnikova et al., 984 fibrillation with a compatibilizer capillary flow - fibrillation viscosity ratio vs. fibrillation shear- induced coalescence. theoretical prediction

HDPE/PA-6 capillary PP/LCP concentration dependence of η Immiscible blends shear-induced interlayer slip

Sondergaard et al., 992 Tsebrenko et al., 976 Tsebrenko et al., 982 Utracki et al., 986 Lin, 979 Utracki et al., 986 Ye et al., 99 Utracki, 987, 988, 989; Bousmina et al., 988

540

L. A. Utracki and M. R. Kamal

Appendix 7.1. Continued No. Study

Blends

0. Shear flows η vs. φ

Description

Immiscible blends PS/PMMA PC/LLDPE PET/PA-6,6 PE/PMMA PMMA/PS

. Extensional flow mixing

7.9.2

PS/LDPE Blends with wide range of viscosity ratios

Theory: emulsion effects & interlayer slip steady state shear viscosity dynamic shear viscosity zero shear viscosity dynamic and steady state shear viscosity zero shear viscosity, microrheology zero shear viscosity Flow through convergence & divergence improves mixing

Reference Utracki, 989, 99 Lyngaae-Jørgensen, 983 Utracki & Sammut, 990 Utracki et al., 982 Matinez & Williams, 980 Valenza et al., 99 Austin et al., 987 Nguyen & Utracki, 995; Luciani & Utracki, 996; Utracki & Luciani, 996

Appendix 7.2. Rheology of Commodity Polymer Blends

Polymer 1

Polymer 2

Type

Method

References

Polyethylene, PE

PS

homologous

dynamic

PE

PP

immiscible

PE

Isotactic polypropylene, PP PP PET Ethylene maleic anhydride-g-PA PA-6, with maleated SEBS, SEBS-MA PS Polyamide, PA Polybutene  (PB)

compatibilized with 5 wt% EPR

Akovali, 967; Masuda et al., 970; Marin, 977; Frank & Meissner, 984; Montfort, 984 Plochocki, 978; Deanin & d’Isidoro, 980 Ho & Salovey, 982; Bartlett et al., 982

PE PE PE

PE

LDPE Polyolefin LDPE

impact strength

Dumoulin, 988 Gerasimchuk et al., 986 Serpe et al., 990

immiscible immiscible immiscible

capillary capillary

immiscible

capillary/ dynamic Lim & White, 994

immiscible compatibilized immiscible/ viscoelastic

Utracki & Sammut, 988 permeability Willis & Favis, 988 capillary/ dynamic Deri & Behar, 992

54

The Rheology of Polymer Alloys and Blends

Appendix 7.2. Continued Polymer 1

Polymer 2

Type

Method

References

LDPE LDPE

LLDPE PS or PS/SEB; PS-bpolyisoprene (diblock) HDPE/ UHMWPE HDPE, LDPE

immiscible immiscible

capillary dynamic, elongational

David et al., 993 Utracki & Sammut, 990

immiscible

dynamic

miscible

capillary dynamic

Dumoulin et al., 984; Vadhar & Kyu, 987 Dobrescu, 980; Chuang & Han, 984 Dumoulin et al., 984

HDPE HDPE LDPE HDPE HDPE

LDPE

miscible/ immiscible immiscible

HDPE HDPE

LDPE LDPE

immiscible immiscible

HDPE

PP

immiscible

HDPE

PA

immiscible/ lamellar

HDPE

immiscible immiscible

torque rheometry

Chen & Lai, 994

UHMWPE UHMWPE

PET with EGMA as a compatibilizer PET with EVAc as a compatibilizer HDPE LLDPE

Dobrescu, 980; Curto et al., 983; La Mantia et al., 984 rotational Kammer & Socher, 982 capillary Acierno et al., 988; La Mantia et al., 986; Bhijsels et al., 988 capillary Alle & Lyngaae-Jørgensen, 980; Alle et al., 98 dynamic/capillary/ Kamal et al., 984; permeability Dumoulin et al., 985; Utracki et al., 986 capillary Dagli & Kamdar, 994

dynamic capillary

UHMWPE HDPE

HDPE PS or PA-6 or PC

immiscible miscible/ immiscible miscible immiscible

HDPE or PP HDPE LLDPE LLDPE

LLDPE

Dumoulin et al., 984 Vadhar & Kyu, 987; Ree, 987 Dumoulin et al., 984 Min, 984; Utracki et al., 986 Tree & McHugh, 988

LLDPE

LLDPE

LLDPE

LDPE

HDPE

UHMWPE

UHMWPE LLDPE PC

immiscible/ fibrillation immiscible miscible immiscible miscible/ immiscible immiscible

capillary

dynamic steady state shear capillary capillary dynamic dynamic dynamic capillary

Bretas & Granado, 993 Utracki & Schlund, 987 Utracki et al., 987; Utracki & Sammut, 987 Utracki, 985c; Utracki & Schlund, 987b Speed, 982; Nancekivell, 985

542

L. A. Utracki and M. R. Kamal

Appendix 7.2. Continued Polymer 1

Polymer 2

Type

Method

References

LLDPE

PP

EPR, di-block

dynamic/ capillary

LLDPE

LDPE

immiscible

uniaxial elongation

LLDPE LLDPE LLDPE LLDPE

UHMWPE LDPE PP LDPE

immiscible immiscible immiscible immiscible

LLDPE PP PP PP

PC PP PP PE

immiscible homologous miscible immiscible

uniaxial elongation capillary capillary capillary, elongational dynamic dynamic dynamic dynamic

Dumoulin et al., 984, 986; Dumoulin, 988; Tree & McAugh, 988 La Mantia et al., 986, 988; Acierno et al., 986; Schlund & Utracki, 987c; Ghijsels et al., 988 Dumoulin et al., 984 Muller et al., 994 Bains et al., 994 Tremblay, 992

PP PP

LDPE PA-6

immiscible immiscible

capillary capillary

PP PP PP PP PP PP PP PP (PP-g-AA) PP PS

PA-6 Elastomer LLDPE SEBS EVAc ULDPE EPDM PET LCP PVME

immiscible immiscible immiscible immiscible immiscible immiscible immiscible immiscible immiscible miscible

capillary elongation elongation capillary dynamic dynamic dynamic capillary capillary dynamic

PS

Polyethylacrylate immiscible

dynamic

PS

PE

immiscible

dynamic

PS

Polycarbonate, PC TMPC

immiscible

dynamic

Utracki & Sammut, 990 Deopura & Kadam, 986 Zeichner & Patel, 98 Dumoulin et al., 986; Dumoulin, 988; Utracki & Schlund, 987 Santamaria et al., 985 Hayashida & Yoshida, 979; Utracki et al., 986; Han, 98; Liang et al., 983 Yakovlev et al., 984 Sebastian & Chen, 983 Dumoulin et al., 984 Gupta & Srinivasan, 993 Cassagnau et al., 992 Lee et al., 99 Van Gisbergen et al., 99 Xanthos et al., 990 Ye et al., 99 Nishi et al., 98; Ajji et al., 988 Smith et al., 987; Bazuin & Eisenberg, 986 Federic et al., 98; Bolotnikova et al., 982; Edel, 984; Sammut & Utracki, 986; Utracki & Sammut, 988 Bye & Miles, 986

miscible

dynamic

Wisniewsky et al., 985

PS

543

The Rheology of Polymer Alloys and Blends

Appendix 7.2. Continued Polymer 1

Polymer 2

Type

Method

References

PS PS PS PS PS PS PS PS PS PDMS PS PE, CPE PS PDMS PS ABS

PMMA PVME SBS LDPE & SEB PPE XPS PA PP PMMA PO PVME PBA EVAc PMMA; PEMA; polyethylenediol PVME ASA

immiscible immiscible immiscible immiscible immiscible crosslinked immiscible

capillary uniaxial elongation elongation dynamic/ elongation dynamic dynamic dynamic

Wang & Lee, 987 Katsaros et al., 986 Dreval et al., 983 Sammut & Utracki, 987 Kim et al., 992 Nemirovski et al., 993 Sundararaj et al., 992

immiscible/ viscoelastic immiscible

dynamic

Graebling et al., 993

immiscible

dynamic

Graebling et al., 994

dynamic shear / elongation

Ajji et al., 99 Munstedt, 98

ABS

SAN

elongation

Saito, 982

PVC

TPU

dynamic

Baranov et al., 992

PMMA PMMA PMMA PMMA PMMA PMMA

PMMA PVDF SAN ACS PVDF PE

immiscible immiscible/ yield stress immiscible/ yield stress immiscible/ viscoelastic homologous miscible miscible immiscible immiscible immiscible

dynamic dynamic dynamic dynamic dynamic dynamic

PMMA PMMA

PVDF and PS SAN; PVDF

immiscible miscible

dynamic shear flow

Onogi et al., 970 Wu, 987 Wu, 987 Han & Yang, 987 Wu, 987 Martinez & Williams, 980 Chuang & Han, 984 Yang et al., 994

Fernandez et al., 995

544

7.9.3

L. A. Utracki and M. R. Kamal

Appendix 7.3. Rheology of Engineering Polymer Blends

Polymer 1

Polymer 2

Type

Method

Reference

Polyamide-6, PA-6 PA PA-6 PA-6

PA

homologous

dynamic

Hinrichen & Green, 98

LLDPE NBR; EVAc EVAc

steady state shear dynamic steady shear

Utracki et al., 986 Wu et al., 990 Okoroafor et al., 992

PA-6 PA PA

PP, SEBS-MA ABS PP

immiscible immiscible immiscible/ interphase immiscible immiscible immiscible

capillary capillary dynamic/ capillary

PA-6 PA-6 PA PC PC

PE LCP LCP PS LDPE

immiscible immiscible immiscible immiscible immiscible

dynamic dynamic

PC PC

PMMA ABS

immiscible immiscible

steady state shear steady state shear

PC PC POM PC PC PC or PP or PET or PA PC PC PC

LDPE ABS TPU ABS 0-90 wt% PO 5-5 wt%

immiscible immiscible

steady state shear capillary/ model dynamic capillary

Ikkala & Reima, 994 Aoki & Watanabe, 992 Samurkas & Rogers, 992 Padwa, 992 La Mantia et al., 990 Jang & Kim, 994 Lipatov et al., 98 Kwok, 987; Sammut & Utracki, 986b Kasajima et al., 98 Dobrescu & Cobzaru, 978 Utracki & Sammut, 987 Kumar et al., 993

LCP LCP PETG; PETG/ LCP PBT/ LCP PA-6/ LCP LCP LCP PET PETG LCP (PET/PHB60) LCP LCP PBT & MBS PS PS High impact polystyrene, HIPS

immiscible immiscible

PC PSF PC PC PC PC PC PC PPE PPE PPE

immiscible immiscible

capillary capillary

Babbar & Mathur, 994 Huang et al., 994

immiscible

capillary, elongational capillary capillary capillary

La Mantia et al., 994 Beery et al., 992 Chapleau et al., 992 Beery et al., 99

immiscible

dynamic/capillary

Skovby et al., 99

immiscible miscible immiscible immiscible immiscible immiscible miscible miscible immiscible

dynamic

Miller et al., 990

dynamic/ capillary dynamic/ capillary dynamic/ capillary dynamic steady shear dynamic dynamic

Nobile et al., 989 Kohli et al., 989 Malik et al., 989 Memon, 994 Takahashi et al., 990 Prest & Porter, 972 Schmidt, 979 Marvin, 960

545

The Rheology of Polymer Alloys and Blends

Appendix 7.3. Continued Polymer 1

Polymer 2

Type

Method

Reference

Polyoxymethylene, POM POM

EVAc

immiscible

capillary

PS

capillary

POM

EVAc

Polyethyleneterephthalate PET PET

PBT

immiscible/ fibrillation immiscible/ fibrillation immiscible

Rezanova & Tsebrenko, 98 Romankevich et al., 982, 983 Tsebrenko et al., 980

Utracki et al., 98, 982; Utracki & Bata, 982; Guenther and Baird, 996 Dimov & Savov, 980 Kim, 988

dynamic dynamic

Lin & Winter, 992 Jin & Tao, 99

capillary

La Mantia et al., 989 Wakita, 993

immiscible immiscible

capillary/ dynamic dynamic dynamic

Sukhadia et al., 990 Ajji & Gignac, 992

immiscible immiscible immiscible

capillary dynamic dynamic

La Mantia et al., 993 Schneider & Wirbser, 990 Lee et al., 994

immiscible

PET PET PA PET PBT

PA-6 PAr EVAl LCP PET

PBT

Liquid crystal polymer, LCP PA-6

immiscible miscible/ immiscible immiscible immiscible; viscoelastic immiscible

PET PBT, PHMT, PC PET PVME SAN

LCP LCP (PHB/PET) LCP PS Phenoxy

Mishra & Despura, 984

steady state shear dynamic and steady state shearing capillary dynamic

Polyamide-66, PA-66

PBT

capillary

immiscible

546

7.9.4

L. A. Utracki and M. R. Kamal

Appendix 7.4. Rheology of Specialty Polymer Blends

Polymer 1

Polymer 2

Type

Method

References

Polyarylate, PAr PAr PAr PAr PAr PEEK PEEK PSF PES Polyetherimide, PEI Thermoplastic polyimide, TPI

PBT PS LCP LCP LCP PTFE LCP LCP Polyimide, PI PEEK/LCP LCP

miscible immiscible immiscible immiscible immiscible immiscible immiscible immiscible miscible immiscible immiscible

capillary dynamic dynamic dynamic slit die steady shear capillary capillary dynamic dynamic capillary

Ausin et al., 987 Ahn et al., 994 Hong & Kim, 994 Hong et al., 993 Zaldua et al., 992 Chan et al., 992 Mehta & Isayev, 99 Golovoy et al., 992 Liang et al., 992 Bretas et al., 994 Blizard & Haghighat, 993

CHAPTER 8

MORPHOLOGY OF POLYMER BLENDS

T. Inoue

Yamagata University, Yonezawa, Japan

8.1

Abstract

In the first part of the chapter several methods used to observe morphology of polymer blends are presented. Various optical microscopic methods are reviewed, including such modern techniques as photon tunneling microscopy (PTM), scanning near-field optical microscopy (SNOM), phase measurement interference microscopy (PMIM), surface plasmon microscopy (SPM) and optical waveguide microscopy (OWM). Many of these methods have been developed to study surfaces and thin films. However, they can also be applied to polymer blend morphology. For transmission electron microscopy (TEM), many staining techniques are described. The double staining method, e.g., using OsO4 and RuO4, was shown to be a powerful tool for enhancing contrast in three and four component systems showing complex morphology. Importance of the etching procedures for scanning electron microscopy (SEM) has been demonstrated to reveal real morphology. A SEM observation by back-scattered electrons from the microtomed and stained specimen was introduced as a novel method for obtaining high quality images. There are serious problems with the digital image analysis (DIA) of micrographs, using commercially available DIA systems. Recently developed, sophisticated DIA methods are introduced in the chapter. The methods provide means for quantification of the polymer blends morphology, and correlation of the extracted parameters with measured properties. In addition to DIA, a light scattering method has been also used for the quantification of morphology. For the better understanding of blend morphologies, the fundamental mechanisms of morphology development are discussed, viz. the liquid-solid phase transition (crystallization), the liquid-liquid phase separation (e.g., spinodal decomposition under non-isoquench depth), as well as the complex mechanism of the morphology generation that results from the competition between these two transitions. The effects of chemical reactions and flow fields on morphology development have also been discussed. Finally, several evidences of a local “structure” in single-phase polymer-polymer mixtures are presented.

L.A. Utracki (Ed.), Polymer Blends Handbook, 547-576. © 2003 Kluwer Academic Publishers. Printed in the Netherlands.

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8.2

T. Inoue

Introduction

By blending dissimilar polymers one can design a variety of morphologies. Morphology control is of prime importance for polymer blends. In the first part of this chapter, a review is given on the microscopic methods to observe polymer blend morphology. This will provide a guide for selection of the most suitable methods of morphology observation. The advantages as well as the limitations of microscopic methods are discussed. In the second part, digital image processing and light scattering methods are presented for morphology quantification. The quantification of morphology is required to discuss the morphologyproperties relationships. In the following part of this chapter, a number of mechanisms of morphology development are discussed. Emphasis is put on the structure development during phase transitions; viz. crystallization and liquid-liquid phase separation, as well as competition between these two transitions. Finally, the effects of chemical reaction and flow fields on morphology development are summarized.

8.3

Microscopic Methods

There is a broad spectrum of microscopic techniques available to observe polymer blend morphologies over several orders of magnitudes of length scale. It starts in the 10 µm range using optical microscopy (OM) and it goes down to the nm range using the high resolution transmission electron microscopy. 8.3.1 Optical Microscopy (OM) First information on the morphology of polymer blends is simply obtained by visual inspection. Blending two transparent, colorless amorphous polymers that have different refractive indexes usually leads to an opaque material, in which the size of phases exceeds the wavelength of visible light (>500 nm). It is usually assumed that to have opacity the difference of the refractive indexes should be larger than 0.003 [Paul, 1978]. In the case when crystalline polymers are involved, the situation becomes more complicated. These

polymers are frequently opaque because there are refractive index differences between the amorphous and crystalline regions and the supermolecular structures (e.g., spherulites) are also large (>500 nm) [Wunderlich and Sullivan, 1962]. Thus in the case of blends with semicrystalline polymers, the distinction between homogeneous and heterogeneous structures of the amorphous phase by visual inspection is limited to studies above the melting point of the crystalline component. A more detailed insight into the phase morphology of polymer blends can be obtained by using optical microscopy (OM). The classical light microscopy is limited by diffraction to domains not smaller than 500 nm. In OM, the necessary contrast for detecting different phases might arise from a number of different sources such as color, opacity, refractive index, orientation, absorption or dichroic differences [Shaw, 1985]. There are numerous imaging methods of OM, e.g., transmission-reflection, bright field-dark field, phase contrast, interference microscopy, polarized light, etc. [Hemsley, 1989; Vesely, 1996]. Prior to the use of more sophisticated methods, light microscopy should be the first experimental method to observe blend morphology in the micro meter range. The method is particularly useful when the light microscope is equipped with a hot stage. The phase behavior is usually best observed in the phase contrast transmission mode. When crystalline polymers are involved, valuable information is obtained using polarized light [Tanaka and Nishi, 1985; Keith and Padden, 1964]. Confocal microscopy was recently developed [Vesely, 1996]. This new OM method makes it possible to obtain images at various depths, then to combine these into a three-dimensional construct. The method has been mainly used for film specimens, but it can also be applied to polymer blends. Although the resolution is similar to the classic OM, this is the only OM method that records three-dimensional morphology. There are several developments in OM to overcome the classical limits. So far, they are mostly applied for the analysis of surface and thin film of polymers. However, they could be equally well used for studies of polymer blend morphology.

Morphology of Polymer Blends

The photon tunneling microscopy (PTM) is an analog of OM to the scanning tunneling microscopy. The method uses light tunneling through one surface to another. This is then observed by scanning near-field optical microscopy (SNOM) [Reddick et al., 1989]. It has been shown that SNOM is able to achieve lateral resolutions of 20 nm [Pohl et al., 1988]. Thus this technique was employed to study, e.g., the surface structure of styrene-methylmethacrylate diblock copolymers [Srinivasarao et al., l992]. The phase measurement interference microscopy (PMIM) is mainly used for surface analysis. A resolution depth of 0.6 nm can be achieved, whereas the lateral resolution is only in the range of about 1 µm [White et al., 1990]. A laser beam in the microscope is reflected from the sample surface and simultaneously from a semitransparent smooth reference surface [Stamm, 1992]. The interference pattern is recorded which can be converted into a contour map of the sample surface, e.g., to determine the root-mean-square roughness of materials. There are two additional methods that are especially useful for the characterization of thin films or monolayers, the surface plasmon microscopy (SPM) and the optical waveguide microscopy (OWM) [Knoll, 1991]. 8.3.2 Transmission Electron Microscopy (TEM) There has been a tremendous development in transmission electron microscopy (TEM) since Table 8.1.

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its discovery by Ruska. To study the polymer morphology by TEM, a sample must be thin, usually thinner than about 200 nm. However, using field emission guns (FEG), thicker specimen (≤ 2 µm) can be used. There are basically three methods to prepare thin specimens: • Casting of dilute polymer solution yields the thin film, which is then taken up with the grid and placed in the TEM. The film can also be stained as discussed below. The morphology of the solution cast films is sensitive to the used solvent and the preparation method. Frequently, these effects cannot be erased by thermal annealing. • The second method employs replication. Direct replica or two-stage replica of the free surface can be used. The free surface may be obtained, e.g., by cryofracture, chemical etching, ion bombardment, etc. Metal shadowing of the replicas with C/Pt coating is frequently used for contrast. • The most successful method is the preparation of ultrathin sections by microtoming. Cryosectioning is required for the samples that may be deformed during sectioning at room temperature. Since polymers are mainly composed of C, H, N and O atoms, the electron density difference between polymers is not large enough to achieve sufficient contrast in heterogeneous materials. Thus staining techniques are usually necessary. Table 8.1 gives an overview of the most common staining methods for polymers.

Fequently used staining methods for TEM [Heckmann, 1990]

Polymer

Staining agent

Reference

Polydienes

OsO4, Br2, RuO4

Polyesters, Copolyetheresters Polyamides Polyolefins

Allylamine/OsO4, RuO4 PTA, RuO4, Formalin/OsO4 Chlorosulfonic acid/ Uranylacetate, RuO4 Hydrazine/OsO4 OsO4 Uranylacetate, lead acetate RuO4 HgO in trifluoroacetic acid, RuO4, Br2

Schafer et al., 1963; Kato, 1965; Vitali and Motani, 1980; Trent et al., 1981 Wegner et al., 1981; Chen et al., 1987; Lutz, 1989 Spit, 1967; Weber et al., 1982; Lutz, 1989 Kanig, 1973; Montezions et al., 1985; Sano et al., 1986

Polyacrylates Polyvinylalcohol Polymers with COOH groups Polystyrene, styrene copolymers Poly(phenylene ether)

Kanig and Neff, 1975 Kanig, 1979; Kast, 1985 John, 1974; Morel and Grubb, 1984 Morel and Grubb, 1984; Trent, 1981 Hobbs et al., 1980; Trent, 1981

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For most cases staining with osmium or rhutenium tetroxide (OsO4 or RuO4, respectively) yields sufficient results. OsO4 reacts with isolated double bonds (C=O, C=C) such as the double bonds in polyisoprene and polybutadiene, but it does not react with conjugated double bonds [Parker and Vesely, 1990]. The OsO4 staining technique was first applied to ABS (acrylonitrile-butadiene-styrene terpolymer) resin [Kato, 1966], then to styrene-isoprene block copolymer [Inoue et al., 1968]. Kato proposed the term “OsO4 staining and fixation technique”. The term implies that the method not only improves the contrast for electron transmissions, but also it hardens the rubbery phase, preventing deformation of rubber particles during ultramicrotoming. Figure 8.1 shows TEM micrographs of ABS resin (a) and ABS latex (b). The small poly(styrene-coacrylonitrile) (SAN) particles are occluded within the rubber particles (so-called salami structure) and rubber particles are partially covered by graftedSAN (viz. “as-polymerized” latex). The OsO4 technique is also effective to enhance observations of the craze formation in ABS, thus providing a better understanding of the rubber toughening [Kato, 1967]. Furthermore, it was also applied to understand the metal plating process of ABS [Kato, 1968]. The ABS sample in Figure 8.2

was microtomed perpendicularly to the surface and stained after the surface was chemically etched and painted with ink. The ink coats the surface and also penetrates into the holes previously (before the etching) occupied by the rubber particles. As shown in Figure 8.2, the method can be used to demonstrate the anchoring of the metal-plating process. Recently, this well-known process started to attract new interests. This is caused by the increasing demand for computer housings that would have sufficient shielding against the electromagnetic noise from and to the environment. The mechanism of RuO4 staining seems to be different from OsO4. It does not react directly with chemical species, but rather it forms clusters. It stains most polymer but to a different degree dependent mostly on the diffusion rate. The diffusion of RuO4 into the polymer is most important [Trent et al., 1983]. For crystalline polymers, RuO4 diffuses preferably into the amorphous regions and stains them, whereas the lamellar crystals remain unstained. One example for successful RuO4 staining is for polyolefins. Figure 8.3 shows TEM micrographs of linear low density polyethylene (LLDPE) and polypropylene (PP) [Sano et al., 1986]. PE crystal lamellae are seen as stripes in Figure 8.3a. Cross-hatched lamellar crystals of PP are seen in Figure 8.3b. For polyolefins,

Figure 8.1. TEM micrographs of (a) ABS resin and (b) as-polymerized ABS latex by the “OsO4 staining and fixation method” [Kato, 1967]. In micrograph (a), polybutadiene phase is stained. SAN particles are occluded in the rubber particles. In micrograph (b), a latex particle is embedded in agar-agar/polybutadiene latex mixture and stained by OsO4. The occluded SAN particles are clearly seen. Surface grafted SAN partly covers the “salami” particle.

Morphology of Polymer Blends

another staining mechanism has been proposed. The process involves oxidation of polyolefin chains by RuO4 that yields: COOH, COORu and RuO2. The oxidation takes place only in amorphous, not in crystalline, regions [Sano et al., 1986]. Figure 8.4 is a TEM micrograph of so-called “block PP” stained by RuO4. The block PP

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is prepared by successive polymerization of ethylene and propylene. It is believed to be a ternary mixture of PP, PE, and poly(ethyleneco-propylene) (EPR). PP and PE phases can be distinguished by the different appearance of the lamellae. The third phase, stained deeply, can be assigned to EPR phase.

Figure 8.2. ABS resin painted with ink, showing the anchoring effect of metal-plating [Kato, 1967].

Figure 8.3. TEM micrographs of (a) LLDPE and (b) PP, both stained by RuO4 [Sano et al., 1986].

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Figure 8.4. TEM micrograph of “block PP” stained by RuO4 [Sano et al., 1990].

Figure 8.5 is a TEM picture of high-impact polystyrene (HIPS)/ABS blend [Hosoda et al., 1989]. The HIPS/ABS blend was stained with OsO4, microtomed, and then stained by RuO4. HIPS consists of PS matrix and polybutadiene particle with salami structure. In contrast, the matrix of ABS is the SAN (25 wt% acrylonitrile) which is less stainable by OsO4 than PS. Then, dark particles can be assigned to rubber particles stained by OsO4, gray region to PS stained by RuO4, and bright region to SAN less stained by RuO4. That is, salami particles with bigger occlusions are dispersed in PS matrix and the salami particles with finer occlusions reside in SAN matrix, suggesting a phase separation between HIPS and ABS. The double staining technique facilitates a clear contrast for more complicated systems, ABS/ polycarbonate (PC)/poly(butylene terephthalate)

(PBT) blend in Figure 8.6 [Watanabe, 1993]. PC particles (very dark) are dispersed in PBT (dark) and the elongated PC/PBT hybrid entities are dispersed in the ABS matrix in which salami particles (gray) are dispersed in SAN matrix (bright). Note here that PC particles are not dispersed in SAN matrix but just in PBT. This may be caused by a balance between the interfacial tensions of the component polymers, that will be discussed later. A double staining using RuO4 and phosphotungstic acid (PTA) was successfully applied for a ternary system, poly(phenylene ether) (PPE)/polyamide (PA)/hydrogenated styrene -butadiene-styrene block copolymer (SEBS) blend. RuO4 and PTA stained SEBS and PA, respectively [Heckmann, 1990]. The electron beam damage is often utilized for TEM visualization. In SAN/poly(methyl methacrylate) (PMMA) blends, TEM contrast was

Morphology of Polymer Blends

553

Figure 8.5. TEM micrograph of HIPS/ABS blend, double-stained with RuO4 and OsO4 [Hosoda et al., 1989].

Figure 8.6. TEM micrograph of ABS/PC/PBT blend, double-stained with RuO4 and OsO4 [by courtesy of Dr Watanabe from the Mitsui Chemicals Inc.].

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enhanced by the degradation of PMMA by electron beam and evaporation of degraded products [McMaster, 1974]. The electron beam radiation can also modify the polymers to be more susceptible for staining, e.g., they turn to be stainable with OsO4 by the formation of isolated double bonds [Parker and Vesely, 1993]. A similar effect was obtained by irradiating SAN/chlorinated-PE (CPE) blends with X-rays followed by OsO4 staining [Michler, 1984]. TEM was also used to measure the interfacial thickness, ∆l, of two-layer specimen [Kressler et al., 1993]. Figure 8.7a shows a TEM micrograph of a two-layer specimen of PMMA and SAN (38.7 wt% acrylonitrile) annealed at 140°C. This system is immiscible but it has small, positive value of the binary interaction parameter, χ12. Figure 8.7b shows the densitogram from the TEM micrograph. The interfacial thickness between PMMA and SAN is ∆l = 32 nm, i.e., much thicker than that observed for PMMA/PS interphase. TEM was used to study the toughening mechanism in rubber modified plastics. A crack tip of a damaged sample of an acrylic rubber modified epoxy resin was observed [Sue et al., 1993].

A double staining technique using OsO4 and styrene was used to achieve good contrast and to prevent damaging the sample during microtoming of ultrathin sections [Sue et al., 1991]. It was shown that near the crack tip the rubber particles are deformed up to 60%, thus suggesting that the highly crosslinked epoxy resin is able to undergo extensive yielding. The cavitation mechanism was also observed. Scanning electron transmission microscopy (STEM) provides the additional possibility to carry out an energy dispersive X-ray analysis and thus providing the possibility of mapping different polymers in a blend under the condition that they contain different elements such as N, O, S, Cl, etc. Bromine mapping was successfully applied for PS/poly(p-bromostyrene) [Voigt-Martin, 1984, 1987] and sulfur mapping for poly(ether sulphone)/ poly(ethyl oxazoline) [Nakamura et al., 1990]. A modified version is the elemental mapping after double staining by OsO4 and RuO4 [Hosoda et al., 1989]. The example is shown in Figure 8.8, for the same sample as shown in Figure 8.5. Thus, the assignment in Figure 8.5 can be confirmed by the mapping method.

Figure 8.7. (a) TEM micrograph of a bilayer specimen of PMMA/SAN system stained by RuO4 and (b) the densitogram across the interface [Kressler et al., 1993].

Morphology of Polymer Blends

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Figure 8.8. STEM analysis after a double staining by RuO4 and OsO4 of a HIPS/ABS blend (the same sample as in Figure 8.5): (a) STEM image, (b) elemental mapping based on Os, and (c) elemental mapping based on Ru [Hosoda et al., 1989].

8.3.3 Scanning Electron Microscopy (SEM) Scanning electron microscopy (SEM) is normally used to observe surfaces. The resolution is not so good as TEM. Usually, SEM micrographs are obtained by collecting secondary electrons emitted upon bombarding the samples with high energy electrons. This secondary electron image (SEI) gives information about the topography of the sample surface. To obtain an information of morphology in the bulk of the material, it is necessary to remove the surface layer. Only when adhesion between the phases is poor, “new surface”, that reflects the bulk morphology can be created by fracturing the sample. Usually, to prevent plastic deformation, the sample is first annealed in liquid nitrogen, then fractured. Another method of removing the surface layer is by etching. This process may be carried out by: • Chemical etching, where one polymer is degraded using a chemical reaction and the reaction products can be removed from surface. • Solvent etching, to selectively dissolve one of the polymers. • Ion beam etching, to preferentially degrade one of the polymers. The low molecular weight byproducts evaporate under high vacuum. An example of chemical etching is seen for polyimide/silica composite prepared by sol-gel method [Morikawa et al., 1992]. After chemical etching of the polyimide matrix with hydrazine,

Figure 8.9. SEM micrograph of “block PP” after Izod impact test. Contrast was obtained by ion etching. The interfacial debonding and rubber cavitation are seen [Nakagawa, 1976].

the finely dispersed silica particles were observed by SEM. An example of ion beam etching is shown in Figure 8.9 [Nakagawa, 1976]. Micrographs of high quality, similar to TEM, can be obtained by SEM observation of microtomed and stained surfaces. Figure 8.10 shows SEM micrographs of 70/30 PP/EPDM blends. Samples for SEM observation were prepared using four different methods [Goizueta et al., 1992]. The sample for Figure 8.10a was fractured after being placed in ethanol-dry ice bath for 30 min. The sample for Figure 8.10b was first fractured as described above, then it was etched

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Figure 8.10. SEM micrographs of 70/30 PP/EPDM blend; (a) SEI of the fractured surface, (b) SEI of the fractured and etched surface, (c) SEI of the microtomed and etched surface, and (d) CI of the microtomed and stained surface [Goizueta et al., 1992].

with cyclohexane to preferentially remove EPDM. The sample for Figure 8.10c was microtomed at -90°C (to provide a new flat surface), then etched with cyclohexane. The micrographs in Figure 8.10a-c were obtained using SEI. The sample for Figure 8.10d was microtomed, then its flat surface was stained by OsO4. Here, the composition image (CI) was obtained by collecting back-scattered electrons from the stained surface. Figure 8.10a shows holes and mounds where the crack has propagated. The phase morphology becomes clearer by solvent etching (micrographs c and d). However, the best contrast was obtained for the microtomed and OsO4-stained sample (micrograph d). Similar results were reported for HIPS and block PP [Goizueta et al., 1993]. It has been observed that SEI under low acceleration voltage provides better resolution for the microtomed and stained samples.

8.4

Quantification of Morphology

A microscopic image itself is important and useful. The quantification of morphology is often required, e.g., to establish processing conditions-morphology and morphology-properties correspondences. Among the methods of morphology quantification, most powerful for polymer blends are digital image analysis (DIA) of micrographs and light scattering.

8.4.1 Image Processing Simple image processing devices are commercially available and are frequently used to obtain statistical parameters from the micrographs, such as number or volume average particle size, histogram for size distribution, and volume (area) fraction. Recently, an international round robin study on DIA of blend morphology was carried out as a VAMAS project. Eleven groups from six countries participated. A wide variation was found to exist in both DIA systems and in the methodologies employed. The results were very different from each other. A complete lack of standardization in the procedure involved is the main culprit cause [Nishi and Ikehara, 1997]. A more sophisticated DIA was proposed for the studies of spherulite formation and phase separation in polymer blends [Tanaka, 1986; Tanaka and Nishi, 1987]. With a different computer vision, shape features of phase-separated structure was obtained [Gur et al., 1989]. The digitalization gives the possibility of two-dimensional Fourier transformation. A power spectrum of the twodimensional Fourier transformation was given for the structure developed by spinodal decomposition [Tanaka et al., 1986]. In the real space, one cannot see the order in the image clearly, whereas the characteristic wavelength and the distribution can be seen in the reciprocal space representation.

Morphology of Polymer Blends

Further, a point-pattern analysis was applied. This is especially suitable for the quantification of particle distribution patterns [Tanaka et al., 1989]. The nearest-neighbor center-to-center distance and effective coordination number (number of cell sides of corresponding Voronoi polygons) of dispersed particles were measured. For a ternary system, PPE/PA/rubber, TEM micrographs were analyzed to estimate the randomness and the degree of clustering [Hayashi et al., 1992]. Recently, a new DIA program was developed [Sigalov et al., 1997]. The particles in micrographs are considered as ellipses rather than circles (Figure 8.11), hence their asphericity (ζ = a/b - 1) and orientation (α) can be dealt with. Moreover, replacement of the particle patterns of complex shape by an effective ellipse, significantly simplifies the calculation of the morphological parameters such as center-to-center and surface-to-surface interparticle distances, particle section perimeters, and coordination numbers. As it will be discussed further in the text, the new DTA was successfully applied to reactive blending (see example of the calculation results in Figure 8.30).

ζ

Figure 8.11. Binarized particle section pattern showing pixel quantitization (exaggerated) and its effective ellipse for DIA [Sigalov et al., 1997].

8.4.2 Light Scattering Light scattering provides several morphological parameters for polymer blends. While the number

557

of parameters generated by light scattering parameters is limited (as compared to DIA), the method is faster and less expensive than electron microscopy followed by DIA. Light scattering provides a convenient and powerful way to generate routine data, e.g., during reactive processing of immiscible polymers. A comprehensive review on light scattering theory was given by Stein [1978]. Since then, there was little theoretical progress, but a very significant one in instrumentation. A light scattering photometer with a 46-photodiode-array system facilitates the time-resolved measurements of the angular distribution of scattered light in a time slice of 50 ms [Okada et al., 1992]. By a highly sensitive charge-coupled device (CCD) camera system, one can carry out the time-resolved measurements of a two-dimensional scattering pattern [Lee et al., 1995]. The light scattering methods have been used for the studies of morphology changes associated with crystallization, spinodal decomposition, reactive processing, and particle coalescence in polymer blends. In polymer blends, one most frequently encounters the two-phase morphology with irregular particle shape and size. For this reason, just the light scattering analysis of a system with irregular particles will be outlined. For irregular two-phase materials, the scattering theory goes back to the fundamental works of Debye and Bueche [1949]. The theory gives a relation between the intensity distribution of scattered light and the statistical parameters that describe the two-phase structure. The intensity of scattered light, I, from the irregular two-phase material usually decreases monotonously with increasing scattering angle θ. The decreasing function is described by: I(q)-½ = (1 + ξ2q2) (8π < η2 > ξ3)-½

(8.1)

where <η2> is the mean-square fluctuation of the refractive index and q is the magnitude of the scattering vector given by: q=(4 π/λ)sin(θ/2), λ being the wavelength of light in the specimen. The correlation distance ξ in Eq 8.1 can be obtained from the slope and the intercept in the Debye-

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Bueche plot of I(q)-1/2 vs. q2 plot. Once the value of ξ is known, the other morphological parameters, such as the specific interfacial area, Ssp, and the mean radius of dispersed particles R, are obtained from: Ssp = 4φ (1 - φ) ξ−1

(8.2)

-1

R = 3φSsp

(8.3)

where φ is the volume fraction of the dispersed phase. The results of R by light scattering were compared with TEM data for ABS and blends of block copolymer with homopolymer - good agreement was reported [Moritani et al., 1970]. Figure 8.12 shows a Debye-Bueche plot for a series of test specimens differing in the diameter of the dispersed phase, but having the volume fraction of the dispersed particles identical. From the linear part of this plot (large q values), the correlation distance ξ can be obtained. In this Figure, there is a strong deviation from linearity for small q-values. From the deviation, one can obtain the interparticle correlation distance a2. Combining ξ and a2, other parameters may be derived, viz. the “distance of heterogeneity” and “volume of heterogeneity” [Moritani et al., 1970].

Figure 8.12. The Debye-Bueche plots for ternary blends of polystyrene, polyisoprene, and styrene-isoprene block copolymers [Moritani et al., 1970].

8.5

Crystallization from a Single-Phase Mixture

8.5.1 Formation of Lamellar Crystals A miscible blend of amorphous and crystalline polymers usually means a single phase in the melt and a neat crystalline phase with a mixed amorphous region in the solid. Because of chain folding during crystallization, the crystal lamellae are formed. Their radical growth usually lead to the formation of spherulites [Nadkarni and Jog, 1991]. When the miscible blend is subjected to crystallization, the non-crystalline impurity is excluded from the crystalline region. The degree of exclusion may depend upon the relative rates of the diffusion rate of non-crystalline component (D) and the growth rate of the crystalline component (G), thus D/G. The exclusion is believed to occur in three ways [Keith and Padden, 1964]. If the value of D/G is comparable with the interlamellar distances, the non-crystalline component may reside within the interlamellar region. If D/G > 1, the non-crystalline component resides between fibrils consisting of lamellar bundle. For very large D/G the non-crystalline component may be excluded from the spherulite. However, in polymer/polymer system, D is usually small. Thus, the non-crystalline polymer is expected to be excluded only on the length scale of the lamellar bundle. The physical meaning of (D/G) itself has not been well understood, but it is a convenient parameter for discussing the morphology development in polymer blends. It was shown that amorphous polymers certainly reside between crystal lamellae as a single phase mixture in poly(ε-caprolactone) (PCL)/poly(vinyl chloride) (PVC) blends [Stein, 1978]. When the content of the amorphous polymer is high, its excess is excluded from interlamellar region to the outside of lamellar stacks in poly(etheretherketone)/ poly(etherimide) blends [Lee et al., 1997]. The light scattering in both the parallel (Vv) and cross polarization (Hv) modes, was measured for PP/partially hydrogenated oligo(styrene-co-indene) blends. The results indicated that the primary lamellae (thick line in Figure 8.13b) sufficiently grew in the

Morphology of Polymer Blends

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Figure 8.13. Schematic representation of lamellar structures in (a) neat PP (see Figure 8.3b), (b) 30/70 PP/hydrogenated oligo(styreneco-indene) (Arkon) blend crystallized at 145°C, and (c) 30/70 PP/Arkon blend crystallized at 85°C [Lee et al., 1995].

radial spherulite direction, while the crosshatched lamellae (thin line in Figure 8.13b) grew less. When the blend was crystallized at 105°C (a representative temperature above the maximum crystallization rate temperature, Tmax = 85°C), this produced spherulites with a positive birefringence. On the other hand, below Tmax = 85°C, non-crosshatched and fragmented lamellae developed and the impurity was segregated out from lamellar stacks to form amorphous domain of ca. 0.1 µm (Figure 8.13c). Thus, various lamellar structures are available by blending miscible impurities and controlling the crystallization temperature [Lee et al., 1995]. The crystal lamella of PE grows rather straight and the edge view of the lamella under TEM is usually a long stripe (see Figure 8.3a). By blending a miscible impurity, the lamella turned into a zigzag [Takagi et al., 1998]. The miscible

polymer was a hydrogenated butadiene block copolymer: ethylene-(ethylene-co-butene)-ethylene block copolymer. The zigzag lamellar structure was characterized by two parameters: the persistent lamellar length (PLL), defined by the length of a straight lamellar segment, and an angle (β) between the neighboring persistent lamellae (Figure 8.14). Incorporation of PE block chains in PE lamellae seemed to force the zigzag growth of lamellae. When the PLL was shorter and the β was wider, the spherulite was less ordered in radial direction. In a 50/50 PP/hydrogenated poly(styrene-cobutadiene) (hSBR) blend, crystal lamellae were extremely fragmented to form the fringed-micelle type crystallites, as schematically shown in Figure 8.15. Closed circles in the Figure are the hSBR domains of 20 nm diameter, which were formed by the spinodal decomposition (the upper critical

Figure 8.14. Persistence length of crystal lamellae: (∆) PE (quenched), (●) 50/50 PE/CEBC blend (Tc = 110°C), (▲) 50/50 PE/CEBC blend (Tc = 75°C), (O) 50/50 PE/CEBC blend (quenched). CEBC is ethylene-(ethylene-co-butene)-ethylene block copolymer [Takagi et al., 1998.

Figure 8.15 Fringed micelle-type PP crystallites in a 50/50 blend of PP with hydrogenated poly(styrene-co-butadiene) [Otsuka et al., 1998].

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solution temperature is below the melting point of PP, UCST < Tm) [Otsuka et al., 1998]. PCL/SAN (27.5 wt% acrylonitrile) system has the lower critical solution temperature type phase diagram, with LCST = 120°C). When the singlephase mixture prepared below LCST was subjected to crystallization, bent lamellae (with short PLL and wide β) were formed. In contrast, when the single-phase mixture was subjected to prolonged spinodal decomposition above LCST, and then the partially demixed system was crystallized, extremely long and straight lamellae were formed (see Figure 8.16). The lamellae length was more than 1 µm [Svoboda et al., 1996]. When the system was demixed above LCST for longer times, then crystallized, much shorter lamellae were formed. At present, one cannot explain why such giant lamellae develop from partially demixed blends.

lamellar crystals of PCL (Figure 8.18a). In contrast, in the blend crystallized below LCST, PCL-rich domains (dark region of 0.1 µm diameter) are dispersed in a PCL-rich matrix, in which the crystallization is highly suppressed (Figure 8.18b). Another type of morphology development was reported for PP/EPR blends [Inaba et al., 1986]. High molecular weight blends of PP and EPR are immiscible. One-phase blends were prepared by precipitation from a solution. This mixture underwent phase separation via spinodal decomposition (SD), followed by crystallization. However, in this case the spherulites were formed seemingly ignoring the bi-continuous SD-structure. Additional interesting information was reported for PCL/SAN blends [Schulze et al.] and PCL/ oligomeric-PS blend [Tanaka et al., 1985; Li and Jungnickel, 1993; Nojima et al., 1991]. 8.5.3 Blend Induced Crystallization

8.5.2 Crystallization via Competition with Spinodal Decomposition In amorphous/crystalline polymer blends with UCST < Tm, the phase separation would interfere with crystallization. That is, during the quenching from a single-phase melt, the system would cross a UCST phase boundary and the liquid/liquid phase separation would compete with the crystallization process. For poly(vinylidene fluoride) (PVDF)/ PMMA blends (see Figure 8.17), normal spherulite growth is observed when the single-phase melt is crystallized at shallow depth below Tm in the one-phase region above UCST (Figure 8.17a). At deeper quench depth below UCST, the liquid-liquid phase separation occurs to yield a bi-continuous morphology, typical for the spinodal decomposition. In the late stages, one phase in the bi-continuous structure begins to brighten under a polarizing microscope, indicating the formation of small spherulites in the PVDF-rich region (Figure 8.17b) [Saito et al., 1987]. Also in PCL/SAN (12.4 wt% acrylonitrile) blends, the interference of the liquid-liquid phase separation during crystallization is seen Figure 8.18 [Svoboda et al., 1994]. In the blend crystallized above LCST, SAN is segregated out of stacked

It is well-known that the rate of crystallization of a crystalline polymer is often reduced by blending with a miscible amorphous polymer. Some typical examples are blends of PVDF/PMMA [Tanaka et al., 1985], PVDF/PEA [Alfonso and Russel, 1986], PEO/PMMA [Briber and Khoury, 1987] and PCL/SAN (19.2 wt% acrylonitrile) [Kressler et al., 1991]. However, the opposite effects are also known, as for example, in blends of polyethersulfone (PES) with polyaramid (PARA) prepared by polycondensation of bis (4-aminophenyl)ether with isopthaloyl chloride [Nakata et al., 1990]. The PARA had been thought to be an amorphous polymer, whereas it did crystallize in the blend. The results suggest that the crystallization of PARA was induced by the presence of PES, i.e., this is an example of the blend-induced crystallization. The aramid seems to be a dormant crystalline polymer. In other words, the crystallizability is latent but crystallization in neat PARA was not possible even when annealing - it did take place in the presence of PES. This may be due to an increase of chain mobility by decreasing the glass transition temperature (Tg) on blending PARA with PES that has lower Tg.

Morphology of Polymer Blends

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Figure 8.16. TEM micrograph of PCL/SAN blend crystallized after spinodal decomposition (stained with RuO4) [Svoboda et al., 1996].

Figure 8.17. Phase diagram of PVDF/PMMA system and the structure formation. The virtual UCST phase boundary is drawn, basing on the kinetic analyses, liquid-liquid phase transition, and crystallization as detected by light scattering [Tomura et al., 1992]. Pictures (a) and (b) are for 60/40 PVDF/PMMA blends crystallized above UCST (150°C) and below UCST (110°C), respectively [Saito et al., 1987].

Figure 8.18. TEM micrographs of 60/40 PCL/SAN blends crystallized at (a) 46°C (above UCST) and (b) 24°C (below UCST) [Svoboda et al., 1994].

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However, since the phase separation also takes place during annealing, it is not certain whether the simple Tg (or mobility) argument is valid. That is, when the liquid-liquid phase separation precedes and crystallization follows, PARA should be dormant, as in the case of a neat resin. Only the crystallization precedes, the Tg argument is applicable. The phase separation by itself may induce crystallization. That is, by the “up-hill diffusion” associated with the SD, the crystallizable aramid chains are transported to the crystal growth front in aramid-rich region, thus increasing the crystallization rate [Matsuura et al., 1992]. Notably, the crystallized blends showed high temperature resistance at temperatures above 400°C. Polycarbonate of bisphenol-A (PC) is recognized as an engineering, amorphous polymer. However, there are many reports that PC does crystallize in blends with, e.g., PCL [Jonza and Porter, 1986], and other polymers [Utracki, 1998].

8.6

Spinodal Decomposition under Non-iso-quench Depth

8.6.1 Solution Casting Figure 8.19 shows a schematic phase diagram of the solution of immiscible polymer pair. A single phase is formed above the binodal (full line). During evaporation of a solvent, the system crosses the phase boundary and it is thrusted into the two-phase

region in which it should phase-separate into A-rich and B-rich phases. During further evaporation, owing to elevated viscosity, the separation process ceases. Then the bi-continuous structure could be frozen in the cast film, as shown in Figure 8.19 [Inoue et al., 1985]. Three factors are responsible for the structure formation during solvent evaporation: the rate of evaporation, the polymer-polymer interaction parameter χ12, and the blend ratio. The faster the solvent evaporates, the smaller the periodic distance in bi-continuous structure. The lager the χ12, the faster casting is required for the formation of bi-continuous structure. The blend ratio is important because it is closely related to the quench depth [Reichelt et al, 1992]. The bi-continuous morphology was not formed when the ratio strongly deviated from unity [Kammer et al., 1993]. When the χ12 value is small and the molecular weight is high, the system fails to phase separate in a limited time of casting, resulting in a homogeneous blend [Inoue et al., 1985]. These systems were named the “forced mixtures”, to indicate that the single-phase system was achieved by fast solvent evaporation, while at equilibrium the polymers are immiscible [Kammer and Kummerloewe, 1990]. When the forced mixtures are heated above their glass transition temperatures, they usually phase separate by spinodal decomposition. A forced mixture of PET with a liquid-crystalline copolyester (LCP) of ethylene terephthalate and

Figure 8.19. Phase diagram of PVC/NBR/THF and optical micrograph of the solution cast film. Spinodal and binodal curves are shown by the broken and solid lines, respectively. The arrow indicates the casting (evaporation) process [Inoue et al., 1985].

Morphology of Polymer Blends

p-oxybenzoate units was prepared [Nakai et al., 1986]. Forced mixtures were also prepared for the blends where one component is not fusible, e.g., poly(p-phenylene terepthalamide) (PPTA) [Kammer and Kummerloewe, 1990] and poly(p-phenylene 1,3,4-oxadiazole) (PODZ) [Kummerloewe et al., 1991]. In these cases, the phase separation had some different features, since immediately after phase separation, PPTA and PODZ behave as solids. A qualitative analysis of the percolation limits was attempted [Kammer et al., 1993]. It has been shown that the highly interconnected two-phase morphology obtained during solution casting via SD is suitable for preparation of charge-mosaic membranes [Kawatoh et al., 1988]. The membrane facilitates negative osmosis with a selective permeability, i.e., it is permeable to salts but not to nonelectrolytes. The high salt-permeability seems to be caused by the bi-continuous nature of two-phase morphology in the solution cast films. 8.6.2 Reaction-induced Spinodal Decomposition There is an increasing number of examples where simultaneous chemical reactions and phase separation are used to achieve desired material properties. A classical example is the production of HIPS where styrene is polymerized in the presence of polybutadiene under intensive stirring.

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In this case, the spinodal character is destroyed by intensive stirring. In contrast, the spinodal morphology can be preserved during the curing of multicomponent thermoses systems. This is typically demonstrated in a blend of epoxy with PES [Yamanaka and Inoue, 1989]. As shown in Figure 8.20, epoxy/ PES has a LCST-type phase diagram. As the molecular weight of the epoxy increases with curing, the LCST is expected to decrease and a two-phase structure should ultimately be formed. However, when the molecular weight increases, the Tg of the mixture also increases. These tendencies are schematically illustrated by arrows in Figure 8.20. At Tcure, a mixture of composition φ is initially a single-phase system. However, as the curing reaction proceeds, SD takes place, and the system is forced into a two-phase regime. In this particular system, the progress of SD will be suppressed by vitrification, as shown by an approaching Tg-line and by the gelation in the epoxy-rich region. Figure 8.20 also shows the phase morphology of the cured material. The SEM micrograph was taken from a fractured and etched surface so that remaining material is cured epoxy. The connectedglobule structure can be explained as a two-phase morphology of interconnected spherical domains of the epoxy-rich phase dispersed in a PES matrix.

Figure 8.20. Schematic representation of the variation of LCST type phase boundary and Tg with the curing reaction in an epoxy/PES mixture (left). SEM micrograph of cured system (epoxy/PES/DDM=100/50/26) (right) [Yamanaka and Inoue, 1989].

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To achieve this morphology the mechanism might be as follows. At first, the system phase separates via SD. At the late stage, the periodic distance of the structure will increase self-similarly and finally the interconnectivity will be maintained by the network formation in the epoxy-rich region. However, an increase in the interfacial tension will make the structure more spherical. The size of globule can be controlled in a range of 0.2 to 10 µm, by judicious selection of the curing temperature and the curing agent. By employing reactive PES, the size can be reduced to 20 nm [Kim et al., 1995]. Cured epoxy/PES blends are tough thermosets used as matrix resins in carbon fiber reinforced plastics (CFRP). A potential for injection molding of the reaction-induced phase separation system is also used in PPE/triallyl isocyanurate system to obtain high-temperature resin with excellent flexural strength and high chemical resistance [Fujiwara et al., 1996]. Relevant systems are rubber-toughened epoxy resins, high-impact acrylic (PMMA particles in rubber matrix, obtained by radical polymerization of 80/20 MMA/EVAc mixture), polyimide/silica hybrid materials (obtained using the sol-gel method), and a very high strength (~60 MPa) rubbers (obtained by peroxide cure of a hydrogenated nitrile rubber/ zinc dimethacrylate system) [Inoue, 1995]. Spinodal decomposition (SD) driven by the chemical reaction proceeds isothermally, but the quench depth ∆T (expressing the temperature difference between LCST and the reaction temperature), increases with time. This situation is quite different from the familiar SD under isoquench depth, where after a temperature jump (or drop) SD proceeds isothermally, and the ∆T is constant. However, the regular morphology is also obtained in the kinetically driven SD, as in the iso-quench SD. This observation was confirmed by the computer simulation using the Cahn-Hilliard non-linear diffusion equation [Ohnaga et al., l994]. This should also be the case for the solution casting, described in preceding section and the shear-dependent decomposition in next section.

8.6.3 Shear Dependent Dissolution and Decomposition The influence of shear fields on the phase transition of polymer mixtures has recently been the subject of a number of investigations. For example, in poly(ethylene-co-vinyl acetate)/solution-chlorinated polyethylene (EVAc/CPE) blend, both a depression and an elevation of LCST by shear were observed [Hindawi et al., 1992]. At low shear rates, the LCST shifted to lower temperatures, thus indicating the shear-induced demixing. This behavior reversed at a critical shear rate, above which the cloud point started to increase, indicating the shear-induced mixing. Noteworthy, these results refer to the shear rates below 10 s-1. The shear rates in melt extrusion or injection molding are significantly higher. PC/ABS blend is an injection moldable material widely used for automotive parts. These blends, developed in 1960, can be prepared by simple meltmixing of PC with ABS, without any compatibilizer. In this early stage of polymer blends technology, the phase diagram of polymer/polymer mixture, the spinodal decomposition, and the effect of the flow fields were not seriously considered. Nowadays, using these concepts, one can describe the processes that are behind the “simple melt-mixing”. Figure 8.21a shows the phase diagram of a blend consisting of a PC copolymer (cPC) and SAN [Okamoto et al., 1995]. Considering this phase diagram, the process of melt extrusion of a cPC/SAN = 50/50 blend may be described as follows. Cold pellets of both polymers are fed from a hopper to the extruder where they are gradually heated up. After the polymer temperature exceeds the glass transition temperatures of both polymers, the resins will melt and the phase dissolution will commence. The dissolution may continue even after attaining the spinodal temperature, Ts, since Ts can increase under shear. Under the high shear rate in the extruder, Ts might be raised above the barrel temperature (e.g., 260°C), thus homogenization of the mixture can be performed in a wide range of temperatures. After extrusion, the homogeneous state can be frozen by quenching. Then, the quenched-homogeneous blend is subjected to injection molding at 290°C (barrel temperature).

Morphology of Polymer Blends

Figure 21b shows the change of resin temperature in mold with time after the injection [Okamoto et al., 1995]. Here, the resin temperature was measured by embedding a thermocouple in the mold set at 95°C. When the melt is out of the shear field in the mold, the spinodal temperature will drop to the static one (Ts = 223°C) and, for about 4 s after injection, the blends is in the two-phase region. The phase separation with large diffusion coefficient, D, will proceed until the melt cools down to Ts. Then, at T < Ts, the phase dissolution takes place. The dissolution slows down near Tg

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of cPC-rich region (here Tg is in between Tg of neat cPC and that of a 50/50 blend), then vitrification takes place. Because the value of D is small, the phase dissolution within the limited, available timespan (about 2 s) will be negligible. Thus, the SD phase-separated structure can be frozen by approaching Tg of the demixed, cPC-rich, system. One could expect a similar scenario for the morphology formation in PC/ABS system. Rubber particles in ABS are good markers for the electron microscopy. It is expected that, after SAN (the matrix resin of ABS) dissolves with PC in an extruder, and

Figure 8.21. (a) LCST type phase diagram showing the processing temperatures for cPC/SAN blends. (b) Time-dependent temperature of the blend within a mold cavity after injection [Okamoto et al., 1995].

Figure 8.22. (a) TEM micrograph (stained by OsO4) of an injection-molded 50/50 PC/ABS blend [Kurauchi and Ohta, 1984]. (b) In the same TEM micrograph as shown in (a), the SAN-rich regions are marked by the presence of rubber particles (dark domains) so that a borderline between ABS-rich and PC-rich regions (the latter shown as bright) can be drawn, to yield the bi-continuous pattern characteristic of the spinodal decomposition.

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the blend is injected, the phase separation proceeds, leading to agglomeration of rubber particles. Figure 8.22a shows the transmission electron micrograph of injection molded PC/ABS = 50/50 blend [Kurauchi and Ohta, 1984]. SAN-rich region should be marked by the presence of agglomerated rubber particles (dark domains, strained by OsO4) so that a borderline between ABS-rich region and PC-rich region (bright) can be drawn to show the characteristic pattern of the spinodal decomposition (Figure 8.22b). PC/PBT blends are also widely used as engineering plastics. There are many experimental studies on the two-phase morphology in these systems. Among them, it is an interesting report that a highly interconnected two-phase structure was observed in melt-mixed blends [Hobbs et al., 1988]. The structure development could be explained considering the shear-dependent mixing and demixing. Both blends, PC/PBT and cPC/PBT, exhibit LCST type phase behavior. In Figure 8.23 are shown the Ts, Tg, and Tm of the constituent polymers, as well as the extruder barrel temperature [Okamoto and Inoue, 1994]. On the basis of the information presented in Figure 8.23, a scenario similar to that of Figure 8.21 may be given for the melt extrusion that yielded a regular, two-phase morphology. Thus, cold pellets of both polymers are fed into extruder where they are gradually heated up.

After the polymer temperature exceeds Tg of PC (145°C) and Tm of PBT (225°C), the phase dissolution starts. Since Ts increases under the shear, even after attaining Ts = 198°C (for PBT/PC), the dissolution will continue toward a homogeneous mixture. Then, once the melt is extruded out from nozzle, the shear rate drops to zero and Ts decreases to its static value. Thus, SD can proceed until the system is cooled down to Ts = 198°C. In the decomposed PBT-rich region, crystallization will also take place. TEM micrographs of an extruded PBT/cPC = 50/50 blend (see Figure 8.23) show a high level of connectivity of both phases and their regular spacing. At larger magnification, the crystal lamellae of PBT are seen in dark PBT-rich region. In short, the morphology developed in PC/PBT blends can be skillfully programmed by a combination of the shear-dependent phase behavior, the liquid-liquid phase separation (SD), and the liquid-solid phase separation (crystallization). Another example of shear-induced mixing comes from the rubber field [Izumitani and Hashimoto, 1985]. A solution cast film of SBR/BR blend with two-phase structure was folded and pressed by hand. After the folding and pressing process was repeated several times, a homogeneous film was obtained. Note that this process is similar

Figure 8.23. TEM micrograph (stained by RuO4) of an extruded 50/50 cPC/PBT blend (left) and listing of the phase transitions and processing temperatures (right) [Okamoto and Inoue, 1994].

Morphology of Polymer Blends

to roll-milling. When the homogeneous blend was quiescently rested at 40°C, the system phase separated by SD. Figure 8.24 shows a TEM micrograph of an injection-molded PP/HDPE blend stained by RuO4. PE and PP phases can be distinguished by the different appearance of lamellar crystals. For PP, cross hatched lamellae are seen. PE lamellae are thicker and longer. It is interesting that PE lamellae “penetrate” into PP region to about 100 nm depth. However, the crystal growth into PP phase is impossible in the immiscible system. That is, there is no way to transport PE chains to the growth front. If so, then PE chains had to exist in the PP phase to start with, i.e., under high shear rate in injection molding machine PP and PE could be locally miscible. Upon removal of the shear field the PE chains crystallized in PP-rich matrix. Also note that in the PE phase, there are the highly disordered regions in which normal PE lamellar growth may be prevented by the presence of PP chains as impurities. This implies that shear-induced mixing took place in the PE-rich phase as well.

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shows the PMMA/PC/PBT system, and Figure 8.25b the PS/PBT/PMMA system, both melt mixed in a twin-screw extruder [Hobbs et al., 1988]. In the first case, PMMA particles (bright) are encapsulated by PC (dark) and the hybrid particles are dispersed in the PBT matrix (gray). In the second case, the particles of PS (dark) and PBT (gray) are isolated and dispersed in a PMMA matrix (bright). The results can be explained by taking into account the interfacial tension coefficients between the components.

Figure 8.25. TEM micrographs of ternary blends: (a) PMMA/PC/PBT = 1/3/7, (b) PS/PBT/PMMA = 1/3/7 [by courtesy of Dr Hobbs from the GE Co.].

According to Harkins, the spreading coefficient Λij of liquid (component i) on solid (component j) is given by [Harkins, 1952]: Λij = νj - νi - νij

Figure 8.24. TEM micrograph of an injection molded HDPE/PP blend [by courtesy of Dr Sano from the Mitsubishi Chem. Co.].

8.7

Melt Blending of Highly Immiscible Polymers

8.7.1 Ternary Blends Figure 8.25 shows TEM micrographs of two different ternary polymer blends. Figure 8.25a

(8.4)

where νj, νi and νij are the surface and interfacial tension coefficients of the solid, liquid and solid/ liquid interfaces. The spreading can occur only when Λij > 0. This Eq 8.4 can be extended to the ternary systems [Hobbs et al., 1988]. The condition for the minor component 1 to spread over the component 3 in the matrix of component 2 can be expressed as: Λ31 = ν12 - ν32 - ν13 > 0

(8.5)

Eq 8.5 provides agood explanation for the morphology evolution in three-component systems. A more precise discussion can be based on the concept of Neumann’s triangle, shown in Figure 8.26 [Nakamura and Inoue, 1990].

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The cosΘ2 is function of the interfacial tension coefficients in ternary systems: cosΘ2 = (ν231 - ν212 - ν223) / (2ν12 ν23)

(8.6)

By the Neumann’s triangle arguments, not only the encapsulation, but also another morphology in three-component systems, that of stuck hybrid particles (Figure 8.26c), can be explained as well.

Figure 8.26. Schematic illustration of the interfacial tensions balance at the interface between three liquids. The Neumann’s triangle, and the morphologies to appear in the melt-blends of three polymers are also shown. Four morphologies are predicted for dispersions of liquids 1 and 2 in liquid 3: (a) encapsulated hybrid particles (1 in 2), (a’) encapsulated hybrid particles (2 in 1), (b) stuck hybrid particles, (c) isolated particles [Nakamura and Inoue, 1990].

8.7.2 Early Stage of Melt Blending The knowledge about the early stage of melt mixing of immiscible polymers, e.g., in an extruder, is limited. When pellets of two different polymers are fed into an extruder, two fundamentally different processes may take place. • When the major polymer melts before the minor one, the pellets are deformed into sheets, holes are formed in these sheets (their expansion is controlled by the interfacial tension forces), the hole coalesce, the sheets break up into ribbons and irregular fragments, and eventually the wellknown droplet morphology is obtained [Scott and Macosko, 1991]. • In systems where at first the minor polymer melts and then the major one, the mechanism is different. At first, the major pellets are coated and then (under the flow conditions) a phase inversion takes place [Sundararaj et al., 1992]. A review on the morphology development in polymer blends during compounding in a twinscrew extruder was published [Utracki and Shi, 1992]. The theoretical predictions of drop deformation and breakup are limited to infinitely diluted, monodispersed Newtonian systems. However, it is possible to obtain valid relationships between processing parameters and morphology. Thus it was found that in the system PS/HDPE the viscosity ratio, blend composition, screw configuration, temperature, and screw speed significantly influence the blend morphology [Bordereau et al., 1992]. For more detail on the topic see Chapter 9, Compounding Polymer Blends, in this Handbook. 8.7.3 Reactive Blending and the Interface

The angles in the Neumann’s triangle can be directly measured by a ‘droplet-sandwich’ experiment. For this purpose, small pieces of C polymer are placed between two layers of polymers A and B, then annealed at a melt temperature. After quenching, the cross-section of the ‘sandwich’ is observed under a microscope. From the value of cosΘ, in Neumann’s triangle, the results in Figure 8.25 can be justified. Also note that values of two νij coefficients can be obtained, if one of them is known [Nakamura and Inoue, 1990].

A typical example of reactive blending is a system containing polyamide (PA) and polyolefins (PO) with a small amount of maleic anhydride (MA). This type of blending has been used to produce “Super-tough Nylon”. The particle size during reactive blending of PA-6 and MA-functionalized polystyrene (PS) decreased by two orders of magnitude, in comparison to the non-reactive PA-6/PS = 80/20 system, in which particle size was about 3.2 µm [Park et al., 1992].

Morphology of Polymer Blends

Poly(styrene-co-maleic anhydride) (SMA) is frequently mixed with SAN before the reactive blending with PA [Takeda and Paul, 1992]. Much attention has been paid to morphology control during the reactive processing [Serpe et al., 1990; Campbell et al., 1990; Willis and Favis, 1990]. Frequently, a third polymer is added as a compatibilizer for binary systems, e.g., MA-grafted SEBS to compatibilize (and impact-modify) blends of PE with PET [Carte and Moet, 1993]. The chemical reaction at the interface during processing influences the morphology and thus the material properties. During the reaction, block or graft copolymers are formed. These copolymers are expected to reduce the interfacial tension coefficient, and to prevent coalescence of the dispersed particles. Furthermore, the chemical reaction influences the interfacial thickness. It was shown by ellipsometry that as a result of the reactive compatibilization, the interfacial thickness in the ternary system, PA/SMA/ SAN, increased up to ∆l = 50 nm [Yukioka and Inoue, 1994]. Figure 8.27 shows the change of morphological parameters in PCL/liquid nitrile rubber with α,ωcarboxyl end-groups (CTBN) obtained by light scattering and ellipsometry, as a function of mixing time and the added amount of γ-aminopropyl triethoxy silane (APS), which was used as a coupling agent [Okamoto and Inoue, 1993]. The size reduction process is shown in terms of time variation of the correlation length ξ, the specific interfacial area Ssp and the mean radius of dispersed particles R. The size reduction is accelerated by addition of APS.

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Another morphology parameter is the volume fraction of the interface, Vλ, expressed as a product of Ssp and the interfacial thickness: Vλ = ∆l.Ssp. It can be shown that the APS loaded system has a significantly larger volume fraction of the interface. Figure 8.28 presents the results of the morpholoogy stability studies for the reactive blend of PA/PP compatibilized with maleic anhydridegrafted PP (PP-MA) [Li et al., 1997]. After mixing PA, PP, and PP-MA, the blend was annealed under quiescent conditions at 240°C. The change of particle size, D, with time was evaluated using light scattering. In the non-reactive system (PP-MA = 0), D increased with the annealing time ta. This suggests that coalescence took place during the annealing. By contrast, in the reactive systems containing a large amount of the reactive compatibilizer (PP-MA = 35 or 70 phr) D reminded constant, suggesting that the formed copolymers prevent coalescence. Note that such a morphology stabilization was achieved already after a short reaction time (tr = 1 min). The slope, β, that is a measure of the coalescence rate, is plotted as a function of the residence time in Figure 8.28. One can see that, even at low PP-MA content (PP-MA = 1 phr), a good stabilization mechanism was established after the reaction time, tr = 10 min. This suggests sufficient coverage of the interface by the in situ formed graft copolymers, which provide the entropic repulsion between neighboring particles. When PP-MA content is higher, the mechanism is set up in shorter tr.

Figure 8.27. Time dependence of the morphological and interfacial parameters during the reactive processing - PCL/CTBN = 80/20 blends at 120°C [Okamoto and Inoue, 1993].

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Figure 8.28. Time dependence of the average particle size,D, during static annealing at 240°C, after melt mixing for 1 min. The slope β may be taken as a measure of the coalescence rate (see left). The coalescence rate vs. the residence time in the mixer tr (right). Numbers indicate the blend composition as: (PP + PP-MA)/aPA [Li et al., 1997].

The interfacial thickness of PP/PA/PP-MA system was estimated by ellipsometry to be thick, ∆l ≅ 40 nm, much larger than the coil size of the component polymers, viz. the root-mean-square radius of gyration <s2>1/2 = 3.7 10-2M1/2 ~10 nm for PP. The extremely thick interface may be caused by an undulation of interface. When the in situ formed block or graft copolymer chains are accumulated in too large a volume at the interface, the copolymer chains are forced to elongate perpendicularly to the interface and destabilize the interface. In other words, the excess accumulation may lead to a negative interfacial tension coefficient, so that the interfacial area will tend to increase by the undulation. As an extreme case of undulation, the copolymer will escape from the interface to form micelles. Such an example was present in the PS-COOH/PMMA-epoxy system

Figure 8.29. Undulation of the interface and micelles formation in reactive blend.

[Nakayama et al., 1993]. The undulated interface and micelle formation are schematically presented in Figure 8.29. The thick interface determined by ellipsometry may correspond to the amplitude of undulation, and sometimes may be including micelles near the undulation tips. During the melt mixing of immiscible polymers (e.g., in an extruder) the dispersed particles are alternatively subjected to shear stress and left to relax. The larger a particle, the more it becomes deformed during shearing and the less it recovers the spherical shape during the “rest” period. The relaxation to the spherical shape is driven by the interfacial tension and hindered by the viscosity of the particle and the matrix. The interface effects dominate over bulk effects when the particle size is small enough, and vice versa. Once the blend is quenched, the shape relaxation is virtually frozen, and the particles remain more or less aspherical - the larger the particle, the greater would be its residual asphericity ζ. Figure 8.30 results from a DIA analysis of the melt mixed PA/polysulfone (PSU) (non-reactive) and PA/PSU-MA (reactive) systems [Sigalov et al., 1997]. The data points are for the non-reactive system. As expected, the data points are scattered above a declined lower boundary (shown by solid line). The upper boundary may appear most likely due to the breakup of very elongated particles. A perfectly spherical shape may only be reached asymptotically, after an infinitely long time. Since

Morphology of Polymer Blends

the specimens are quenched immediately after sampling, the asphericity parameter should not be less than a certain value corresponding to the lower boundary. In the reactive system, the in situ formed graft copolymer would reduce the interfacial tension and lead to a more aspherical shape of the particles. Then, the lower boundary would be shifted upwards, compared with the non-reactive system. This is shown by the broken line in Figure 8.30.

Figure 8.30. Particle asphericity-size correlation for nonreactive 80/20 PA/PSU blend. The solid line shows the lower boundary above which 97% of data points lie. The lower boundary for reactive system (PA/PSU-MA) is shown by the broken line [Sigalov et al., 1997].

8.8

571

“Structure” in Single Phase Mixtures

The miscibility of amorphous blends is usually judged on the basis of the occurrence of a single glass transition temperature. But Tg of miscible system is broader than these of the component polymers [MacKnight et al., 1978]. Thus it has to be considered that such methods as DSC or DMA, measure only a macroscopically averaged quantity. In these systems local fluctuations of composition might exist, in addition to normal density fluctuations. It was shown by solid state 13 NMR [LeMenestrel et al., 1992], as well as by electron spin resonance [Muller et al., 1992] that some local ordering exists at a length scale smaller than 5 nm. The TEM measurements have also shown some evidence of a microstructure in miscible polymer blends [Vesely and Finch, 1988]. The miscibility of dissimilar polymers arises from the specific interactions, such as hydrogen bonding or weak charge-transfer complexing. The specific interactions are expected to induce some conformational changes of the individual chains. Alignments of dissimilar chains would be parallel to each other and the chains would be locally stretched, as compared to the random coil chains (see Figure 8.31). The local nematic ordering was suggested by the depolarized light scattering studies [Saito et al., 1991]. Monte Carlo simulations also showed that even for athermal mixtures, the number of heterocontacts is larger than estimated from random mixing rules [Cifra et al., 1988]. The non-randomness is not only caused by the enthalpic interactions, but also by the free volume contributions [Nies and Xie, 1993].

Figure 8.31. Chain conformations in a single-phase mixture of dissimilar polymers: (a) locally ordered chains and (b) random coil chains. The local nematic ordering is suggested by the depolarized light scattering [Saito et al., 1991].

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8.9

T. Inoue

Concluding Remark

To observe morphology in polymer blends, one should start with a careful observation by the naked eye and then by OM. This will provide information on a length scale of 10 µm, thus about rather macroscopic inhomogeneities, often created by inadequate control of extrusion, solution casting, etc. The macroscopic inhomogeneities can be critical for the materials performance. To observe finer details of morphology, such methods as SEM and TEM are to be used. There are many publications and reports based on SEM observation of carelessly prepared specimens, e.g., by simple fracturing that may provide information on fractography, but not morpholoogy. Proper specimen preparation is essential if the applied method is to reveal the true morpholoogy. Fine morphology can be determined using TEM on properly stained specimens. To get fine morphological details on the atomic scale, atomic force microscopy (AFM) is a promising new tool. However, special pretreatment of the sample, similar to those used for SEM, should be applied to obtain the real blends’ morphology. Once the morphology has been observed, it is important to consider the mechanisms responsible for its development. For example, consider ABS/PC morphology as discussed in Part 8.6.3. By recognizing the mechanisms, one can not only reach better understanding of the morphology, but also get ideas for designing better materials. For the DIA of micrographs, one should pay attention to commercial computer programs. The computations should be calibrated, using a series of model images with simple geometries. At present, the real-time observation of morphology development can be done only by OM. Real-time observation of finer morphology on sub-micron level is expected to be soon available. In the meantime, light scattering methods (LSc) can compensate for the present “hole”. LSc are quite powerful for the studies of the morphology changes during: SD, crystallization, reactive processing, coalescence, etc. Morphology control on the crystal lamellar level is a new, interesting research domain. A suc-

cess story is the commercial “Toyota Super Olefin Polymer (TSOP)” [Sano et al., 1996]. TSOP is reported to be a blend of low molecular weight PP and olefinic rubber, in which fragmented crystal lamellae in the core of injection-molded part provide the toughness and the crystallites in skin part, developed perpendicularly to the surface, lead to high surface hardness. Clearly, morphology control is a key-technology in polymer blends. The data should aim to define the processing-morphology-performance relationships.

8.10

Acknowledgement

The author expresses his sincere thanks to Dr. Jorg Kressler for helping in the literature survey for Section 8.3 (Microscopic Methods).

8.11

Nomenclature

8.11.1 Abbreviations ABS AFM APS BR CCD CEBC CI CPE CTBN DIA DMA DSC EPR EVAc FEG HDPE HIPS hSBR LCP

acrylonitrile-butadiene-styrene atomic force microscopy α-aminopropyltriethoxysilane polybutadiene, butadiene rubber charge coupled detector ethylene-(ethylene-co-butylene)-ethylene block copolymer composition image chlorinated polyethylene carboxyl terminated nitrile rubber digital image analysis dynamic mechanical analysis differential scanning calorimetry poly(ethylene-co-propylene); ethylene-propylene rubber copolymer from ethylene and vinyl acetate; poly(ethylene-co-vinyl acetate) field emission gun high density polyethylene high impact polystyrene hydrogenated poly(styrene-co-butadiene) liquid crystalline polymer

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Morphology of Polymer Blends

LCST LLDPE LSc MA MMA NBR OM OWM PA PA-6 PARA PB PBT PC PCL PE PEAc PEG PES PET PLL PMIM PMMA PO PODZ PP PP-MA PPE PPTA PS PSU PSU-MA PTA PTM PVAc PVC PVDF SAN (SB)ñ SBR SD

lower critical solution temperature linear low density polyethylene light scattering methods maleic anhydride methyl methacrylate nitrile rubber; poly(butadiene-coacrylonitrile) optical microscopy optical waveguide microscopy polyamide poly(ε−caprolactam) aromatic (mainly amorphous) polyamide polybutadiene poly(butylene terephthalate) polycarbonate of bis-phenol-A poly-ε-caprolactone polyethylene poly(ethyl acrylate) polyethyleneglycol poly(ether sulfone) poly(ethylene terephthalate) persistent lamellar length phase measurement interference microscopy polymethylmethacrylate polyolefin poly(p-phenylene 1,3,4-oxadiaxole) isotactic polypropylene (aPP — atactic; sPP — syndiotactic) maleated polypropylene polyphenyleneether poly(p-phenylene terephtalamide) polystyrene polysulfone maleated polysulfone phosphotungstic acid photon tunneling microscopy poly(vinyl acetate) poly(vinyl chloride) poly(vinylidene fluoride) styrene-acrylonitrile; poly(styrene-coacrylonitrile) multi-block styrene-butadiene copolymer styrene-butadiene elastomer spinodal decomposition

SEBS

styrene-ethylene/butene-styrene three block copolymer SEBS-MA maleated SEBS SEI secondary electron image SEM scanning electron microscopy SMA styrene-maleic anhydride; poly(styrene-co-maleic anhydride) SNOM scanning near-field optical microscopy SPM surface plasmon microscopy STEM scanning transmission electron microscopy TEM transmission electron microscopy THF tetrahydrofuran UCST upper critical solution temperature 8.11.2 Notation D D G I q R Ssp ta Tcure Tg Tm Tmax tr Ts V

α β ν12 ζ θ Θ λ ∆l ξ φ1 χ12 Λij

mean diameter of a particle diffusion coefficient crystal growth rate intensity of scattered light magnitude of the scattering vector mean radius of a particle specific interfacial area annealing time curing temperature glass transition temperature melting point temperature maximum crystallization rate temperature residence time in mixer spinodal temperature volume fraction of interface angle between neighboring persistent lamellae coalescence rate interfacial tension coefficient asphericity scattering angle angle in Neumann’s triangle wavelength of light interfacial thickness correlation distance volume fraction of dispersed phase binary interaction parameter spreading coefficient

574

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T. Inoue

References

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CHAPTER 9

COMPOUNDING POLYMER BLENDS

L. A. Utracki and G. Z.-H. Shi

National Research Council Canada, Industrial Materials Institute, Boucherville, QC, Canada

Introduction Cost of polymer alloys is mainly determined by the composition. By contrast, the profit that is based on the alloys’ performance is controlled by the way the material is put together, i.e., by the morphology and its stability. The compounding process must ascertain that the alloy has the desired spectrum of the performance characteristics. Undoubtedly, compounding is the most critical and difficult step in the polymer blends’ technology. It must combine the fundamental knowledge of thermodynamic and rheological material behavior, with engineering aspects of flow inside the mixing machines. Furthermore, since nowadays about 90% of blends are reactively compatibilized, good knowledge of polymer chemistry is essential. This Chapter has four parts: Fundamentals of polymer mixing, Blending methods and equipment, Non-mechanical methods of polymer blending, and Reactive processing. The information is presented in a concise form, with tables of data and references to the source literature. Owing to the complex nature of the topic, interested reader is encouraged to consult other chapters in this book dedicated to specific topics that impact on the Compounding of Polymer Blends viz. compatibilization (Chapters 4 and 5) and rheology (Chapter 7).

9.1

Fundamentals of Polymer Mixing

According to the English dictionaries mixing, blending and compounding are synonyms, indicating “an action to combine ingredients into one mass, so that the constituent parts are indistinguishable”. However, in the plastics processing, these terms have different meanings: mixing indicates the physical act of homogenization (e.g., mixing of fractions), blending usually indicates preparation of polymer blends or alloys, while compounding that of a compound, i.e., incorporation of additives into a polymeric matrix, viz. antioxidants, lubricants, pigments, fillers, or reinforcements [Utracki, 1994, 1998].

L.A. Utracki (Ed.), Polymer Blends Handbook, 577-651. © 2003 Kluwer Academic Publishers. Printed in the Netherlands.

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9.1.1

The Reasons for Mixing

Mixing is the most important operation in polymer processing. Uniformity of the molecular weight, degree of entanglement, temperature, and composition, is the key to good plastics’ performance. For example, fusion of rigid PVC formulation by heating, but without mechanical mixing, results in abominable mechanical properties, poor thermal stability and weatherability. Similarly, poor performance is obtained from compression molding of a dry blend of LLDPE reactor powder — low level of the intermolecular entanglements results in low strength of molded parts. Hydrodynamics recognizes two kinds of a fluid motion: (1) laminar, in which the streamlines are smooth, and (2) turbulent, in which the motion is irregular or disorganized. These two types are separated by the critical value of the Reynolds number (Recritical ≅ 2000). The Reynolds number is defined as: –

Re ≡ L V ρ / η

(9.1)

where L is the channel size, W is the flow velocity, ρ is the fluid density, and η is its viscosity. For a typical molten polymer with η = 1 kPas, ρ = 1,000 kg/m3, flowing through a one meter long channel, L = 1 m, the turbulence occurs when the flow velocity exceeds the critical value of W = 2 km/sec. Evidently, the turbulent flow is not to be expected for liquids with such a high viscosity, hence the polymers flow by laminar mechanism. [Note that the term “laminar” is used in two different connotations: to characterize the flow (as in laminar vs. turbulent), and that of mixing, viz. “laminar mixing model” — see below.] It is convenient to distinguish two types of flow: dispersive and distributive, schematically shown in Figure 9.1. The dispersive (or intensive) mixing involves application of stresses that break domains of the dispersed phase to the desired size. The dispersed phase may be composed of liquid drops, gel particles of the matrix material, aggregates of filler particles, etc.

The distributive (or extensive) mixing involves homogenization of a fluid, accomplished by the application of strain. Homogenization may involve a single-phase fluid (e.g., homogenization of temperature), a miscible system (homogenization of concentration), or a multiphase system, blend or composite (homogenization of the dispersion).

Figure 9.1. The schematic representation of mixing (top row, left to right): bad dispersion and distribution; bad dispersion, but good distribution; (bottom row, left to right): good dispersion, but bad distribution, and good dispersion and distribution.

Two other terms have also been used to describe mixing: the laminar and chaotic. In both types the system is considered “passive”, i.e., the rheological properties are identical and the interface is “invisible”. Thus, the flow-induced morphological changes do not affect the flow behavior — the effects of flow and morphology are “decoupled”. In real systems the rheological and thermodynamic properties of two phases are different, the changes of morphology affect the flow, which in turn affect the morphology, and so on, until a steady state is reached — the effects are “coupled”. Thus, the laminar and chaotic mixing models are excellent to start with, but excepting special cases, they do not represent reality. As indicated in Chapter 7, Rheology of Polymer Alloys and Blends, the effects of flow on morphology have been described by the microrheology. For highly diluted systems, good agreement between the experimental data and theoretical predictions has been reported. Furthermore, during the last few years, the efforts have been made to incorporate the kinematics of the interface generation into the rheological description of the system [Doi and Ohta, 1991; Grmela and Aït-Kadi, 1994, 1998; Lee and Park, 1994; Grmela et al., 1998].

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9.1.2.

Laminar Mixing

The following derivations for the laminar mixing are valid for the flow of passive mixtures. Neither the effects imposed by the interface, nor those originating from the differences in the rheological properties of the fluid components are taken into account. Thus, the laminar mixing (as presented below) provides the first approximation for the distributive (or homogenizing) mixing. The dispersive mixing has not been considered — there is no need for it when there is no interface and the rheological properties of both fluids are the same [Edwards, 1985]. 9.1.2.1 Growth of the Interfacial Area, R Laminar mixing depends on the strain tensor, which can be visualized as an ellipsoid formed upon straining a sphere. The strain magnitude is proportional to the relative size of the ellipsoidal axes while their relative positions to the orientational effects of the imposed flow. It can be shown that the interfacial area, Ao, will grow with the imposition of strain according to the relation [Erwin, 1991]:

Depending on the type of deformation, the interfacial ratio (R) takes on values listed in Table 9.1. For example, the principal elongation ratios in simple shear can be expressed as:

[ = [1 + γ

λ x = 1 + γ 2 / 2 + (γ / 2)( 4 + γ 2 )1/ 2 λy

2

/ 2 − (γ / 2)( 4 + γ 2 )1/ 2

1/ 2

] ]

1/ 2

λz = 1 9.1.2.2 Energy of Mixing, E It is important to be able to calculate how much energy is required to generate the same R in different flow fields, in particular in extension and in shear. The specific energy (per unit volume) required to generate the aerial strain R in a passive liquid mixture with viscosity, η, for time to, is given by the following expressions [Erwin, 1991]: Uniaxial extension: E = (12η / to )[ln(5 R / 4)]2 Biaxial extension: E = (3η / to )[ln(5 R / 4)]2 Planar extension: E = ( 4η / to )[ln(2 R)]2 Simple shear: E = ( 4η / to ) R 2

{

R ≡ A / Ao = (λ x λ y )2 + λ2y (λ2z − λ2x )cos 2 α + + λ2x (λ2z − λ2y )cos 2 β

1/ 2

}

(9.2) where λi (i = x, y, z) are the principal elongation ratios, and α and β are the orientation angles.

Table 9.1.

(9.3)

(9.4) These dependencies are shown in Figure 9.2. Evidently, the simple shear is inefficient for generation of a large interfacial area (for smaller deformations, R ” 10, the inefficiency is not as dramatic). The inefficiency can be reduced if a

Values of R for different types of deformation

Deformation

Elongation ratios

Plane elongation Plane elongation Pure elongation Pure elongation Simple shear

λx = λo ; λy = 1/ λo ; λz = 1 Averaging isotropic mixture λx = λo ; λy = λz = λo-1/2 λx = λo ; λy = λz = λo-1/2 expressed by shear strain, γ

R λo λo / 2 λo1/2 1 / λo γ cos β

Comment regular, flow between rolls random input orientation uniaxial stretching biaxial stretching Couette or Poiseuille flow

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large deformation is obtained not at once, but in a series of smaller deformation steps. The best formula is: stretch-twist-fold, the second best is stretch-&-fold, the third is stretch-&-randomize. Thus, to improve the mixing efficiency of the processing units that operate in the shear mode (viz. a single-screw extruder) the flow should be interrupted and the flow lines randomized. For uniform shear strains of randomly oriented elements, the interfacial area increases linearly with strain — when the process is repeated, the total increase of the interfacial area becomes proportional to square of strain, γ2, etc.:

ally adopted by the designers of mixing equipment. The process would be even more efficient if the elements could be reoriented at 90° (biaxial stretching), or even 180°(folding and stretching). Several attempts have been made to incorporate these ideas into extruders’ design. 9.1.2.3 Kinematics of Mixing The strain between two material points M and M’ can be expressed as:

(9.6) (9.5) The same approach may also be used to improve mixing in the extensional flow field. However, since the surface generation in extension is efficient to start with, the improvement is not as dramatic as that expected in shear.

The kinematics of the deformation is described by the following set of differential equations [Poitou, 1988]:

(9.7) where X is the coordinate of M, u is the velocity field, [grad u] is the velocity gradient tensor, [e], is their strain rate tensor, and m is the unity vector of MM’. Figure 9.3 presents the kinematics of strain evolution in simple shear and uniaxial elongation.

Figure 9.2. Energy consumption as a function of change of the interface area in extensional (uniaxial, biaxial and plane strain) and simple shear flows. The most efficient is the biaxial stretch, the worst (by a factor of 500,000!) is shear (after Erwin, 1991).

The validity of Eq 9.5 was experimentally confirmed using a specially designed mixer [Ng and Erwin, 1981]. Even if this device did not lend itself to be incorporated into a processing unit, the idea of the interrupted flow has been gener-

Figure 9.3. Kinematics of separation of two material points during laminar flow in the uniaxial elongation and in the simple shear.

Compounding Polymer Blends

9.1.2.4 Laminar Mixing — a Summary The elongational deformation provides more efficient mixing than shear. On all three accounts: the magnitude of the interface increase, the energy required for mixing, and the rate of spatial separation of two material points, significantly better mixing (frequently by several orders of magnitude) is expected for the extensional than for the shear flow. The fundamental assumptions in all these derivations have been that the system is rheologically decoupled, that there is neither an “active” interface, nor a significant difference between viscosity of the mixture components. In consequence, the coalescence does not exist — the flowing materials always are stretched, deformed, and the degree of mixedness continuously improves. By contrast with immiscible polymer blends, the model does not take into account either the shear coalescence or the thermodynamic coarsening. Thus, the laminar mixing provides good guidance for the distributive homogenization of idealized polymer mixtures. The conclusions that come from this model may be directly applied for mixing of homologous polymer blends, temperature homogenization of polymer melt, or to mixing different batches of the same resin. 9.1.3

Chaotic Mixing

The term chaotic mixing” was introduced by Ottino [1989] to describe laminar, distributive mixing

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with continuous or periodic translation of cavity walls. The phenomenon is universal, observed in many diverse branches of science and technology — from aerodynamics to geology [Ottino, 1989; Rauwendaal, 1994]. This type of mixing can be observed even in a simple geometry as a result of interplay of at least two velocity fields, which can be generated by applying either two type of forces (e.g., heating and shearing, imposition of two types of motion), or by imposing a specific type of flow geometry. To illustrate the principle, a “two-dimensional” cavity was constructed. The box-like geometry had a transparent front wall and movable top and bottom neoprene walls. An insertion of rod-shaped, fluorescent tracer drop (with only a crossection visible) made it possible to follow the laminar mixing as a function of time for diverse sequences of the top and bottom wall translations. As shown in Figure 9.4, periodic translations resulted in stretching, folding and transporting the drop much more effectively than the regular, steady state translation. At the same strain, the mixedness index, expressed as a ratio of the generated interface to the initial one, was significantly higher for the periodic translations than for the steady state. The experiments also demonstrated presence of the “mixing islands”, where very little mixing took place. In consequence, the chaotic mixing could be schematically represented by the streamline diagram comprising the elliptic points located in the center of the “blinking vortex” and a hyperbolic

Figure 9.4. Mixing in a 2D cavity with upper and lower neoprene walls. Left — effect of continuous motion of both walls, right — that of discontinuous motion (1/2 top, then 1/2 bottom, etc.). In both tests the total walls’ displacement was the same [Leong and Ottino, 1989].

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point. Little mixing takes place around the elliptic points surrounded by circulating streamlines, while extensive mixing takes place at the hyperbolic points where the streamlines (leading in opposite directions) crosslink. Thus, the principle of chaotic mixing involves a non-steady-state translation of the material by stretching and folding that creates horseshoe-shaped domains of the dispersed phases [Smale, 1967; Leong and Ottino, 1989; Ottino, 1990, 1991]. Formally, it was described as originating from a “blinking” vortex that is periodically displaced by imposed strain on different cavity walls [Aref, 1983]. As in the case of laminar mixing, here also it is advantageous to start consideration with a passive system. Imposition of a periodic motion causes the domains to stretch and fold. Several 2D flows were explored and mathematically treated. Of these, the cavity flow (with two moving walls) and the journal bearing flow were the investigated. For example, in the first geometry good stretch-&-fold mixing was observed for periodic, interrupted shearing of the upper, then lower wall (see Figure 9.4). More recently, three-dimensional chaotic mixing was studied by Liu and Zambrunnen [1995]. A cylindrical cavity of the journal bearing type was constructed. The cylinder diameter was 79.0 mm, the height of the cavity was 23.7 mm, and the offset of the shearing wall centers was 27.6 mm. The disks were alternatively rotated, each for half period. The cavity was used to generate fine fibrillar structures (aspect ration of about 1000) of PE in either EVAc or PS. The authors remarked that owing to extensional flow the mixing process is insensitive to the viscosity ratio. The principles of chaotic mixing are well established, but the development of mixing devices based on them still awaits commercialization. 9.1.4

Dispersive Mixing — Microrheology

This topic has been already presented in Chapter 7, The Rheology of Polymer Alloys and Blends. In the following text only the aspects pertinent to mixing will be summarized [Utracki, 1991, 1995]. It is important to recognize the advantages and the drawbacks of the microrheological approach.

By contrast with the discussed above laminar mixing, the microrheology focuses on the physical differences between the two blends’ components, the dispersed and the matrix fluids. Both, the thermodynamic (expressed by the interfacial tension coefficient) and the rheological properties (expressed by the strain, stress, and the viscosity ratio), are important for the deformation and breakup of the dispersed phase. The drawbacks are several, viz. assumptions of high dilution, low deformation rates, and that the systems are Newtonian. In spite of these, the microrheology provides invaluable guidance for mixing of commercially interesting systems. For example, the use of these principles led to successful predictions of the blends’ morphology evolution during flow in a twin-screw extruder [Shi and Utracki, 1992; Huneault et al., 1995a]. 9.1.4.1 Drop Deformability When a neutrally buoyant, initially spherical droplet is suspended in another liquid and subjected to stress, it deforms into an ellipsoid, and then breaks up into smaller droplets. Taylor [1932, 1934] described the process (at low stress in a steady state, uniform flow), using the following three dimensionless parameters: the viscosity ratio, the capillarity number, and the reduced time, respectively, defined as: .

λ ≡ ηd / ηm ; κ ≡ σd / ν12 ; t* = t γ / κ = γ / κ (9.8) where σ = yηm is the local stress, ηd and ηm are the dispersed phase and matrix viscosity, respectively, y is the deformation rate, d is the droplet diameter and ν12 is the interfacial tension coefficient. During flow the initially spherical drop deforms into a prolate ellipsoid with the long axis, a1, and two orthogonal short axes, a2. The drop deformability parameter, D, is a complex function. At low shear stress it can be expressed as [Taylor, 1934]: Dshear ≡ ( a1 − a2 ) / ( a1 + a2 ) = (κ / 2)[(19λ + 16) / (16λ + 16)]

(9.9)

Compounding Polymer Blends

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Since, for the full range of λ values from zero to infinity, the quantity in the square bracket in Eq 9.9 ranges from 1.00 to 1.18, the drop deformability Dshear ≅ 0.55 κ. For highly viscous dispersed phase, i.e., λ >> 1, Taylor calculated that D = 5/4λ, but since the condition for drop break is: D • Dcrit = 0.5, this means that shear flow can lead to drop breakage only when the viscosity ratio λ < 3.8. For the planar and the axisymmetrical extensional flow the theory predicts that the deformability is larger than in shear, viz. respectively: Dplanar = 2 Dshear ; Daxisymmatric = (3 / 2) Dshear (9.10) Cox [1969] derived the kinetic equation for the deformability of a drop in the planar flow: Dplanar = κ [(19λ + 16) / (16λ + 16)]

{

[1 − exp −(10κ / 19λ ) t *planar .

where:

t *planar ≡ (t γ

planar )λ

}

/ κ = γ planar λ / κ (9.11)

where γplanar is the strain in planar flow and t*planar is the reduced time for planar deformation. It is worth noting that in the hyperbolic flow the reduced time scale depends on the viscosity ratio. Unfortunately, Eq 9.11 is valid only for small values of λ and κ. The exponential master curve is shown in Figure 9.5. The effects of λ and κ on D are illustrated in Figure 9.6.

Figure 9.5. Drop deformability in a planar hyperbolic flow as a function of total strain. The dependence was calculated from Cox equation, Eq 9.11.

Figure 9.6. Drop deformability in a planar hyperbolic flow at three values of λ and t κ. The dependencies were computed from the Cox Eq 9.11.

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The above relations are valid for Newtonian systems undergoing small, linear deformations, smaller than that, which would lead to a breakup. As Figure 9.5 indicates, in planar hyperbolic flow about 10 units of strain are required to get into an equilibrium. In simple shear, t*d ≅ 25 of the reduced time scale is required [Elemans, 1989]. 9.1.4.2 Drop Breakup From the point of view of the drop deformation and breakup there are four regions of the reduced capillarity numbers κ∗ ≡ κ / κcr, both in shear and elongation: for κ∗ < 0.1 droplets do not deform, for 0.1 < κ∗ < 1 droplets deform, but do not break, for 1 < κ∗ < 2 droplets deform then split into two primary droplets, for κ∗ > 2 droplets deform into stable filaments. The critical capillary number, κcr, is defined as the minimum capillarity number sufficient to cause breakup of a deformed drop. It is customary to plot κcr as a function of λ. For simple shear and uniaxial extension flow, De Bruijn [1989] found the following dependence (see Table 9.2 and Figure 9.7):

breakup depends on the viscosity ratio, λ. In shear, four regions have been identified: for λ << 0.1 small droplets are shed from sigmoidal drops — tip spinning, for 0.1 < λ < 1 drop breaks into two principal and odd number of satellite droplets, for 1 < λ < 3.8 drop deforms into fiber, then disintegrates into small drop lets, for λ > 3.8 drops may deform, but do not break.

Figure 9.7. Critical capillarity number for drop breakup in shear and extensional flow.

log(κ cr / 2) = c1 + c2 log λ + c3 (log λ )2 + c4 / (log λ + c5 )

(9.12)

Similarly as κcr, also the critical time for drop breakup, t*b, varies with λ. When values of the capillarity number and the reduced time are within the region of the drop breakup, the mechanism of

Table 9.2. Flow Shear Elongational

9.1.4.3 Drop Fibrillation and Breakup For κ* > 2 drops deform affinely with the matrix into long fibers. When subsequently the deforming stress decreases, causing the reduced capillarity number to fall below two, κ* < 2, the fibers disintegrate under the influence of the interfacial

Constants for drop breakup in Newtonian systems 1000c1 -506.0 -648.5

1000c2

1000c3

1000c4

1000c5

-99.4 -24.42

124.0 22.21

-115.0 -0.56

-611.0 -6.45

Compounding Polymer Blends

585

tension [Tomotika, 1935; 1936]. The distortion amplitude α grows exponentially with time, t:

α = α o exp {qt}

(9.13)

where αo is the distortion at t = 0 and the growth rate parameter q = Ω(Λ, λ)(t*/t), with Ω(Λ, λ) being a tabulated function of the distortion wavelength and the viscosity ratio. The thread breakup occurs when α = R ≅ 0.81Ro. The most serious obstacle for the quantitative use of Timotika’s theory is the estimation of the initial distortion, αo. 9.1.4.4 Microrheology of Viscoelastic Systems The mechanisms governing deformation and breakup of drops in Newtonian liquid systems are relatively well understood. However, within the range of compounding and processing conditions the molten polymers are viscoelastic liquids. In these systems the shape of a droplet is determined not only by the dissipative (viscous) forces, but also by the pressure distribution around the droplet that originates from the elastic part of the stress tensor. Therefore, the characteristics of drop deformation and breakup in viscoelastic systems may be quite different from those in Newtonian ones. Some of the pertinent papers on the topic are listed in Table 9.3.

Figure 9.8. Computed radius at the neck for disintegrating jet stream of Newtonian (solid line) and Maxwell fluid (dashed line) [Bousfield et al., 1986].

Figure 9.9. Deformability of PS drops in PMMA during steady-state shear flow at 180°C. The points are experimental, the line is to guide an eye [Lyngaae-Jørgensen et al., 1993].

In summary, there is no theory capable to describe the deformability and break-up of viscoelastic drops in viscoelastic media, or even one that provides explanation for the divergent observations. The complexity of the problem involves three-dimensionality, free surface, and non-stationary flow leading to complex constitutive equations. Furthermore, the role of fluid elasticity in the breakup process is not well understood. In most experiments, the elasticity stabilized the deformed drops, thus making the dispersing process more difficult. However, there are also reports of the opposite effects. The microrheological behavior of two-phase liquid systems may provide but a general guidance for the drop deformation and breakup in polymer blends. For blends of industrial interest the phenomenon is further complicated by higher concentration of the dispersed phase as well as by the presence of numerous additives that affect the interfacial phenomena. One of the sources of confusion may originate from different types of liquids studied. As reported recently, large differences in deformability in convergent flow were observed for drop prepared from viscoelastic solutions or from Boger fluid [Bourry et al., 1999]. Spiegelberg et al. [1996] have shown that a Boger fluid is strongly strain hardening. At Hencky strains larger than two, the elongational viscosity is larger than shear viscosity by three orders of magnitude. These high values of strain hardening may explain the

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Table 9.3.

L. A. Utracki and G. Z.-H. Shi

Deformability of viscoelastic drop

Description

Reference

Higher κcrit for viscoelastic system than for Newtonian drops.

Gauthier et al., 1971

For Newtonian drops in viscoelastic fluid, elasticity of the medium stabilizes the drops.

Flumerfelt, 1972

ν12 = νo12 + (do/12)[(σ11 - σ22)d - (σ11 - σ22)m] (9.14) where: νo12: is the interfacial tension in a quiescent polymer blend; do is the initial diameter of the dispersed drop; and the bracketed expressions are the first normal stress differences of the dispersed (i = d) and of the matrix (i = m) phase. For (σ11 - σ22)d > (σ11 - σ22)m, Eq 9.14 predicts that higher elasticity of the dispersed phase than that of the matrix causes the interfacial tension coefficient to increase, what leads to more stable drops. For (σ11 - σ22)d < (σ11 - σ22)m, it predicts that: 0 < ν12 < νo12 thus the flow tends to enhance the dispersing process (flow compatibilization). However, as the shear stress increases, Eq 9.14 predicts physically untenable condition: ν12 < 0. To avoid the physical absurdity, the system morphology changes from dispersed to lamellar.

Van Oene, 1978

Viscoelastic drops in either Newtonian or viscoelastic media were reported more stable (less deformable) than the Newtonian ones. The critical shear rate for droplet breakup was lower for λ < 1 than for λ > 1.

Han and Funatsu, 1978; Chin and Han, 1979, 1980

Initially, disturbances grow more rapidly in viscoelastic filaments than in Newtonian ones, but at a later stage there is stabilization, resulting from large extensional stresses (see Figure 9.8). As a result, the time required to break the viscoelastic drop is longer than that for Newtonian. Formation of the satellite drops was found to be retarded by elasticity.

Bousfield et al., 1986

Both the drop and the medium were viscoelastic. Drops did break during extrusion, even when λ > 4, but the mechanism is not clear.

Wu, 1987

In viscoelastic systems κcrit is always higher than in Newtonian — drop elasticity hinders the drop breakup, whatever the λ value.

De Bruijn, 1989

When λ < 0.5, the drop elasticity has stabilitizing effect.

Prabodh and Stroeve, 1991

Predictive model for the morphology variation during simple shear flow under steady state uniform shear field was developed. The model considers the balance between the rate of breakup and the rate of drop coalescence. The theory makes it possible to compute the drop aspect ratio, p =a1/a2, a parameter that was directly measured for PS/PMMA = 1:9 blends. Theoretically and experimentally, p vs. shear stress show a sharp peak at the stresses corresponding to a transition from the Newtonian plateau to the power-law flow, i.e., to the onset of the elastic behavior (see Figure 9.9).

Lyngaae-Jørgensen et al., 1993; 1999

Deformation of a single drop in a medium subjected to convergent flow was observed. Both liquids were of the Boger fluid type. For a given matrix, the drop deformability decreased with elasticity of the dispersed phase. For a given drop liquid, the drop deformability increased with the matrix elasticity. The following empirical equation was proposed:

Mighri et al., 1997

(9.15) ld,E = ld,N - ∆P(lm - 1) where ld is the aspect ratio of the elastic (subscript E) or Newtonian (subscript N) drop, ∆P is a function of the elasticity ratio, and lm is the matrix deformation expressed in Hencky terms. The deformability of the viscoelastic drops in Newtonian matrix was studied in the convergent slit flow. Both, the experimental observations and the boundary element method computations were carried out. It was reported that deformation of the Boger fluid drop, was quite low —about 1/3 of that recorded for the deformability of a strongly shear-thinning, viscoelastic solution. The latter drops showed deformability similar to these observed for Newtonian drops of similar viscosities.

Bourry et al., 1999

587

Compounding Polymer Blends

large differences of deformability of two types of viscoelastic drops. 9.1.5

Newtonian fluids [Ottino, 1983]. An example of comparative evaluation of 11 commercial SM is presented in Table 9.4.

Distributive Mixing 9.1.6

The static mixers (SM) operate on the principle of repetitive dividing of a flow channel into at least two new channels, reorienting them by 90°, and dividing again. The flow is a pressure driven, laminar shear. Mixing by SM is related to the numbers of striations (Ns ) generated by a number of SM elements (ne) and the number of divisions (new channels) engendered by each element (nc): N s = ncne

(9.16)

Over 30 different SM designs have been commercialized [Rauwendaal, 1986]. Their efficiency is determined comparing: (1) the length-to-diameter ratio (L/D proportional to ne) required to produce the same degree of homogeneity; (2) the associated pressure drop, ∆Prel; (3) the holdup volume, ∆Vrel; and (4) the relative dimensions of the device, Drel and Lrel. The efficiency of SM also depends on the type of liquid. For example, Ross mixer was found better for the pseudoplastic polymer melts, whereas Kenics mixer for the

Table 9.4.

Mixing in Extensional Flow Field

The elongational flow field exists anywhere where the streamlines are not parallel. This type of deformation is quite common during the processing, viz. in fiber spinning, film blowing, blow molding, biaxial and uniaxial stretching and the similar operations carried out down-stream from the die. However, these operations have little to do with mixing. It is extremely rare to find a mixer where the energy dissipated in extension is nearly as large as that in shear — mixing is overwhelmingly carried out in the shear field. Since the easiest method for the generation of extensional flow is by means of convergence and divergence, the attention will be focused on these types of deformations. 9.1.6.1 Orientation of Solid Particles The extensional flow field has been used to generate orientation either in homopolymers or in filled or reinforced systems. The convergent flow results in high fiber alignment in the flow direction,

Comparison of static mixers performance

Static Mixer

L/D

∆Prel

∆Vrel

Drel

Lrel

Koch SMX Koch SMXL Koch SMV Kenics Etoflo HV Komax Lightnin PMR Toray N-Form Ross ISG

9 26 18 29 32 38 100 320 13 29 10

1.0 0.6 2.3 0.6 0.6 2.1 2.6 14.5 1.3 1.4 8.6

1.0 1.8 4.6 1.9 2.0 8.9 29.0 511.0 1.9 4.5 9.6

1.0 0.8 1.3 0.8 0.8 1.3 1.4 2.4 1.1 1.1 2.1

1.0 2.4 2.7 2.7 2.7 5.4 15.3 86.0 1.6 3.6 2.3

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whereas the diverging flow causes the fibers to align 90° to the major flow direction. Shearing reduces the alignment. The early theoretical studies of the elongational flows were carried out by Takserman-Krozer and Ziabicki [1963]. Molden [1969] developed a geometrical theory that showed the extensional strain to be more effective than shear for fiber alignment. Several other theoretical descriptions of the chopped fibers orientation in flow have been proposed [Lockett, 1972; Johnson, 1972; Vincent and Agassant, 1985]. For example, Goettler used the convergent-divergent die geometry to orient fibers in the transverse direction, greatly improving pipes’ hoop strength [Goettler, 1970; Goettler and Lambright, 1977; Lee and George, 1978; Crowson et al., 1980; Goettler et al., 1981; Goettler and Shen, 1983].

Table 9.5.

9.1.6.2 Convergent Flow of Viscoelastic Fluids The fundamental studies of drop deformation in convergent-divergent geometry were carried out using either axisymmetric [Mighri et al., 1997] or slit geometry [Bourry et al., 1999] (see Table 9.3). The analysis of converging flow of viscoelastic fluids using nonlinear viscoelastic models was carried out by several authors. The work is summarized in Table 9.5. 9.1.6.3 Blending in Extensional Flow Field The use of extensional flow field for mixing is relatively unexplored, while a growing number of reports show that mixing in extensional flow field is more efficient than in shear, especially for

Convergent flow of viscoelastic systems

Description

Reference

Kaloni used Oldroyd model, Schümmer a fourth order fluid model, while Wissler a nonlinear Maxwell model. Employing the perturbation method, the authors observed that the inclusion of second-order perturbation terms (which bring in the non-Newtonian effects) predicted velocity profiles with superimposed secondary circulation patterns. Normal stresses of HDPE and PP flowing through a converging channel were measured. Near the die exit, the wall normal stress rapidly increased. Relationships were derived for the wall shear stress, normal stress difference at the channel wall, and the wall normal stress gradient in the flow through a converging channel and through a conical duct. Contrary to viscometric flows, in the converging flow field the normal stress alone does not permit to determine the pressure gradient of viscoelastic fluids. Flow into either sharp-edged or tapered die was studied using stress-optical measurements. Distribution of shear stress and normal stress difference was obtained for PS melt — neither stresses nor velocities showed secondary motion. Flow through converging-diverging tubes has been computed (velocity fields, pressure distribution, and the wall stresses). The velocity profiles were measured using laser-Doppler anemometry. Flow through an axisymmetric pipe, whose diameter was slowly varying in the axial direction, was theoretically described. A converging flow rheometer was developed. Converging and diverging dies produced flat, biaxially oriented extrudates. Relationships between the die geometry and product orientation were established. Flow of multilayer viscoelastic fluids in converging channel slit die was analyzed. The neutral stability contours were altered in the channel flow. The effect of convergence on interfacial stability was similar to that of the elasticity. Channel convergence can be used to stabilize the interface at low depth ratios provided that shear thinning effects are more dominant in more viscous layers.

Kaloni, 1965; Schümmer, 1967; Wissler, 1971

Han, 1973, 1974, 1975

Han & Drexler, 1973

Theodorou et al., 1984 Williams & Javadpour, 1980; Williams & Williams, 1985 Mascia & Zhao, 1991; Lohfink & Kamal, 1993 Su & Khomami, 1992

Compounding Polymer Blends

blends with higher viscosity ratio, λ • 3.8, where the shear field is unable to cause drop breakup [Grace, 1971]. Most of the works on mixing in the extensional flow field consider convergent flow of Newtonian liquids, flowing from a reservoir to a capillary [Tsebrenko et al., 1974, 1976; Ablazova et al., 1975; Krasnikova et al., 1978; Vinogradov et al., 1982; Han and Funatsu, 1978; Chin and Han, 1979, 1980; Han and Yu, 1981; Han, 1981]. The following factors affecting the quality of blends were identified: (i) Diameter of the convergence, dc, (ii) The ratio of the reservoir-to-convergence diameters, usually expressed as the convergence ratio, C ≡ dr/dc, (iii) The capillary length-to-diameter ratio, R ≡ L/d, (iv) The initial drop size, d, (v) Extensional viscosity ratio, λε ≡ ηεd/ηεm [extensional viscosities of the dispersed phase and that of the matrix, ηεd and ηεm respectively, to be taken at the conditions existing during the mixing, viz. extensional stress, σ11, temperature, T, pressure, P, etc.], (vi) Absolute value of the elongational stress, σ11, and (vii) The number of passages through the convergence. These factors affect both the dispersive and distributive mixings. Suzaka [1982] patented an extensional flow mixer, composed of a series of plates placed across the flow channel. A polymer blend was forced to pass through a series of convergent and divergent orifices (or c-d for short), that elongated the drops and dispersed them. All the circular orifices had the same size — the convergence ratio was constant. The mixer provided good mixing for some systems but not for others. Furthermore, the device could be optimized only by trial and error. Any change in the blend composition required repetitive testing, the process had to be interrupted, and the mixing plates changed. In consequence, the mixer has not been commercially explored.

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9.1.6.4 Extensional Flow Mixer During the last few years there has been significant progress in the fundamental understanding of microrheology and its role for polymer blending. Flow visualizations of the Newtonian and nonNewtonian systems have been carried out [Mighri et al. 1996; Picot, 1997; Bourry et al., 1999]. The boundary element method (BEM) provided important information on the evolution of drop shape during the flow through c-d channel [Luciani et al., 1997; Khayat et al., 1996]. This led to designing of a new extensional flow mixer (EFM) in which the blend is repeatedly exposed to extensional flow fields and semi-quiescent zones. The c-d channels are of progressively increasing intensity, with flow in the radial not axial direction. To reduce the pressure drop, and to prevent blockage, slit restrictions are used [Nguyen and Utracki, 1995; Utracki and Luciani, 1997]. The EFM is a fully adjustable, general-purpose, motionless mixer. The mixing action is provided by the development of extensional flow field through c-d geometry. It is primarily a dispersive mixer that should be attached to a pressure-generating device. It efficiently homogenizes different liquid systems, even these where the components’ viscosity ratios are large. The mixer attached to a single-screw extruder provided comparable or better mixing than a twin-screw extruder with the so-call mixing screw geometry [Bourry et al., 1995; Utracki, and Luciani, 1996; Luciani, and Utracki, 1996]. EFM has been used for: polymer blending, incorporation of elastomers into resins, and dispersion of high viscosity resins or “gel particles”. Figure 9.10 shows a cross-section of the EFM. The molten polymer blend enters EFM from an extruder through adapter plate #1. The melt is distributed by the distributing plate #3 to six slots, located between the cone and the mounting ring of the part #3. Next, the melt enters the annular space inside the EFM body #5 where it is directed to the space limited by the upper (part #4) and lower (part #6) c-d plates. There it flows from the rim toward the center, undergoing the convergent and divergent deformations before sorting out

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L. A. Utracki and G. Z.-H. Shi

Figure 9.10. Cross-section of the extensional flow mixer, EFM, assembly. The numbered parts are described in the text.

through the central passage in the lower plate #6 and the central bore in the plate holder #7. The magnitude of the stresses is controlled by the pressure (generated by the extruder) and by the gap between the two c-d plates (parts #4 and 6). The latter is adjusted by turning the adjusting plate #8. The pressure drop, ∆P, across the EFM was calculated using an expression, proposed by Binding [1988]. The pressure loss across the c-d plates, ∆P, was computed for a series of blends with either PE, PP, or PS as matrix. The result can be expressed either by the proportionality or given by a linear fit: (9.17)

Figure 9.11. The total pressure drop across the EFM c-d plates and its extensional and shear components vs. throughput of polypropylene, using gap of one millimeter.

The computations made it possible to separate the pressure losses into the extension and shear contributions, Pe and Ps, respectively. As evident from Figure 9.11, in EFM the pressure drop caused by elongation is larger than that by shear. The performance of EFM attached to a singescrew extruder (SSE) was examined in reference to the results obtained using a co-rotating, intermesh-

ing twin-screw extruder (TSE) equipped with the recommended by the manufacturer high dispersion screws. Four types of mixing were studied: dispersion of a viscous polymeric blend component in at least four times less viscous polymer matrix, impact-modification of engineering resins, elimination of gel particles from reactor powders, and homogenization of new metallocene polyolefin

∆Pexp = (1.026 ± 0.195)∆Pcalc ; or ∆Pexp = 0.0512 + 1.044 ∆Pcalc , R = 0.9772

591

Compounding Polymer Blends

mixtures having very large ratios of melt flow index. The results demonstrated that compounding in SSE+EFM resulted in finer dispersion of HDPE in PS and vice-versa, than in a TSE. Similarly, the impact properties of EPR/PP blends prepared in SSE+EFM were better than of the same compositions mixed in TSE. The SSE+EFM mixing unit was also superior to TSE as far as dissolution of very high molecular weight fractions (polymeric gel particles) was concerned [Utracki and Luciani, 1996; Luciani and Utracki, 1996]. EFM was commercialized in 1999. 9.1.7

Interphasial Properties

9.1.7.1 Importance of the Interfacial Tension Coefficient At high dilution the morphology of an immiscible blend is controlled by the viscosity ratio, λ, the capillarity number, κ, and the reduced time, t*, as defined in Eq 9.8. The interfacial and rheological properties enter into κ, and t*. As the concentration increases, the coalescence becomes increasingly important. This process is also controlled by the interphasial properties. Blending two immiscible polymers always creates the third phase — the interphase. In binary blends, thickness of this third phase, ∆l, is inversely proportional to the interfacial tension coefficient, ν12. When the blend approaches miscibility, ν12 approaches zero and ∆l goes to infinity. Thus the interphase, with its own set of characteristic parameters (e.g., viscoelasticity) may dominate the behavior of nearly miscible systems, as well as that of compatibilized blends. For further details on this topic see Chapter 4. Interphase and Compatibilization of Polymer Blends. To modify the interphase properties, three strategies of compatibilization have been developed: (i) Addition of a small quantity of a third component that either is miscible with both phases (co-solvent), or it is a copolymer whose one part is miscible with one phase and another with the other phase (0.5 to 2 wt%, usually block-type, less frequently a graft one);

(ii) Addition of a large quantity, ” 35 wt%, of a core-shell copolymer that behaves like a multi-purpose compatibilizer-cum-impact modifier; and (iii) Reactive compatibilization designed to enhance domain interactions. The reactive blending may be employed to generate in-situ the desired quantities of either block and/or graft copolymer(s) (see also in this book Chapter 5. Reactive Compatibilization). Different strategies lead to different alloys, having different sets of properties, e.g., addition of a small amount of block copolymer mainly affects the interfacial tension coefficient, thus the size of dispersion, but under normal circumstances it only slightly affects the shear sensitivity of the blends’ morphology, or the solid-state behavior. The reactive compatibilization was found to engender thick interphase, which resulted in excellent stability during abusive processing and good mechanical performance. 9.1.7.2 Theoretical Aspects of the Interface Helfand and Tagami [1971, 1972] model is based on self-consistent field that determines the configurational statistics of the macromolecules in the interfacial region. The interactions between the statistic segments of polymers A and B are determined by the thermodynamic binary interaction parameter, χ12. The isothermal segmental density profile shown in Figure 9.12, ρi (i = A or B), was calculated for infinitely long macromolecules, Mw→ ∞. The interfacial thickness, ∆l∞, and the interfacial tension coefficient, ν∞, were expressed as: ∆l∞ = 2 b / (6 χ12 )1/ 2 ; ν ∞ = bρk B T ( χ12 / 6)1/ 2 (9.18) where b is a lattice parameter, kB is Boltzmann constant and T is the absolute temperature. The Helfand-Tagami lattice theory predicted that: (1) the product, ν∆l∞, is independent of the thermodynamic binary interaction parameter, (2) the chain-ends of both polymers concentrate at

592

L. A. Utracki and G. Z.-H. Shi

the interface, and (3) the low molecular weight components are repulsed to the interface. Addition of a compatibilizer reduces ν12 to the level corresponding to the critical micelles concentration, CMC [Utracki and Shi, 1992; Tang and Huang, 1994]. The radius of the dispersed drop, R, follows the same “titration curve” as the interfacial tension coefficient, ν12. The amount of the interfacial agent required to saturate the interface, wcr, per volume fraction of the dispersed phase, φ, can be expressed as [Paul, 1978; Matos, 1993]: wcr / φ = 3M / RN Avζ

(9.19)

where M is molecular weight, R is radius of the dispersed drop, NAv is Avogadro number and ζ is the area occupied by a copolymer molecule, taken either as a crossection of a copolymer chain (ca. 5 nm2), or proportional to the square of the end-to-end distance of copolymer chain, ζ = 〈r2〉/9.

Figure 9.12. Computed segmental density profile across the interface. The figure defines the interphasial thickness, ∆l [after Helfand and Tagami, 1971].

9.1.7.3 Calculations of the Interfacial Tension Coefficient The interfacial tension coefficient can be calculated from the solubility parameters, δ, that comprises contributions from the dispersive, polar and hydrogen bonding interactions. The following dependence was proposed [Luciani et al., 1996, 1997]:

[

ν12 = k ( ρRT ) n −1 Θ(δ1d − δ 2 d )2 + (δ1 p − δ 2 p )2 + (δ1h − δ 2 h )2

Figure 9.13. Interfacial tension coefficient computed from the chain structure of the polymeric chains for 46 polymer blends vs. the computed solubility parameter difference [Luciani et al., 1996].

n

]

(9.20) where k is a constant, Θ = 1 - 4 is a parameter, and exponent n (depending on the adopted theoretical assumptions) n = 1/2 to 3/2. The expression: (ρRT)n-1, is insensitive to temperature. All available experimental values of ν12 were plotted as functions of the computed values for polymer blends, assuming that all factors have the same value, Θi = 1 (see Figure 9.13).

9.1.8

Coalescence

During mixing, the dispersed phase progressively breaks down until a minimum drop diameter, d, is reached. For Newtonian systems the size of the smallest drop that can be broken can be calculated from the Taylor theory (viz. Eq 9.9). However, many experimental studies [Tokita, 1977; Roland and Bohm, 1984; Plochocki et al., 1990; Willis et al., 1990, 1991; Favis et al., 1991] have shown

593

Compounding Polymer Blends

that the final drop size is usually larger than predicted. For example, only when φd < φo § 0.005, the observed drop sizes approached the theoretical values [Elmendorp and Van der Vegt, 1986]. The source of the discrepancy is coalescence. There are two types of coalescence mechanisms, the first being determined by the equilibrium thermodynamics, and the second caused by flow. Thus, coalescence occurs in flowing as well as quiescent systems. To the latter type belong the “Ostwald ripening”, characterized by the linear increase of the drop volume with time, i.e., d3 ∼ t. The process involves diffusion from smaller drops (high interfacial energy) to the larger ones. Flow modifies the rate of this process — here rate depends on the drop dynamic crossection, i.e., dn ∼ t, where 3/2 < n < 3 [Ratke and Thieringer, 1985]. Tokita [1977] suggested that the drop diameter in polymer blends originates from the two competitive processes: continuous breakup and coalescence of the dispersed particles. The equilibrium drop diameter should increase with concentration, number of drops, and the interfacial tension coefficient, but decrease with shear stress, σ12. The dependence qualitatively agrees with experiments [Liang et al., 1983; White and Min, 1985; Willis et al., 1991]. The coagulation time, tc, of PVC lattices under steady state shear flow, followed the theoretical relation [Utracki, 1973]: . −2

tc ∝ d 3φ d−8 / 3 γ

On the other hand, the microrheology predicts that the rate of diameter change due to break is:

(d d / d t )break ∝

.

− γ d / κ cr tb*

(9.23)

These two equations indicate which factors can be used to enhance either dispersion or coalescence. Clearly, the shear rate is expected to similarly affect coalescence and breakup. However, the flow-induced coalescence is a strong function of concentration whereas the break is not, thus concentration may be used to discriminate between these two processes. Furthermore the rate of break is proportional to d, whereas the coalescence is proportional to 1/d. Thus, coalescence is not expected to play a major role in the beginning of the dispersion process. Equating Eqs 9.22 and 9.23 leads to [Huneault et al., 1995a]:

(

o + 6Cκ cr tb*φ 8 / 3 deq = deq

1/ 2

)

(9.24)

were κcr is the critical capillary number for breakup, t*b, is the time to break, and φ is the volume fraction of the dispersed phase, is the equilibrium drop diameter and doeq is its value extrapolated to zero concentration (see Figure 9.14).

(9.21)

The coagulation rate was related to the projected area of the drop d* ∝ φd2/3. The relation separates the effects of the particle diameter, concentration and the rate of shear. When the other variables are constant, the coagulation time increases with particle diameter following the Ostwald ripening kinematics. It decreases with nearly a cube of concentration and with a square of the shear rate. From Eq 9.21 the rate of the diameter change due to coalescence can be expressed as: .

(d d / d t )coalescence ∝ γ φ 8 / 3 / d

(9.22)

Figure 9.14. Average equilibrium diameter of polyethylene dispersed in polystyrene matrix as a function of PE volume fraction. The data were obtained blending the resins in an internal mixer at 200°C until an equilibrium drop size was obtained, e.g., for 300 s.

594

L. A. Utracki and G. Z.-H. Shi

9.1.9

Measures of Mixedness

As defined above, mixing is an operation that provides enhanced spatial homogeneity of a system. However, the homogeneity is not an absolute — depending on the scale of observation one can always find domains belonging to a single component of the mixture. By definition, mixtures that show heterogeneity of composition on a scale not larger than the radius of gyration of a given polymeric chain, are considered thermodynamically miscible. They may be considered as having the ultimate degree of mixedness. Assessment of the quality of a mixture, or mixedness, is difficult and time consuming. Since the performance of blends is controlled by the optimum degree of dispersion, as well as by other factors (e.g., compatibilization, adhesion in solid state, the level of degradation engendered during the blending, etc.), the mixedness must be considered a separated quantity, to be determined independently of the blends’ performance. In a binary mixture of two polymers, A and B, their local concentration (in weight fraction) can be expressed as: a(x) + b(x) = 1 (where x denotes vectorial location). Similarly, the sum of the average compositions: aav + bav = 1. For such a system, the mixedness can be described by the binary frequency function. For fine dispersions, variance of the composition and the intensity of segregation can be defined as [Hold, 1991; Tucker, 1991]: 2 a

Variance: σ =

[a( x ) − aav ]

2 a

where Av,o is the initial interfacial area at strain γ = 0. It was estimated that for adequate laminar mixing 18,000 strain units are required, what gives: Av / Av,o → ∞ [Hold, 1991]. For layered or laminar mixtures, one can also define mixedness by the striation thickness, s (defined as 1/2 of the layer thickness). It can be shown that in these systems there is a simple relation: AV = 1/s. Since in real blends there is a variety of striation thicknesses, s should be expressed by a volume distribution function. Another excellent description of mixedness is the power spectrum, derived by Fourier transformation of the correlation function: ∞

∫ R(r) exp{−2πinr}dr ≅

2

[a( x ) − aav ] [a( x + r ) − aav ]

∫ R(r)cos(2πnr) dr

(9.28)

−∞

where n is the wave vector. The power spectrum can be used to calculate the variance and the linear scale of segregation:

(9.26)

0

where: R(r ) ≡

(9.27)

(9.25)

∞

∫ {R(r) / σ }dr ;

r→o

−∞ ∞

Since segregation is not uniform, it is important to know the scale of segregation in the mixture. The volume scale of segregation is defined as [Danckwerts, 1953]: Sv = 2π

Av = 4 lim ( dR(r ) / dr ) = ( Av,o / 2) exp {γ / 2}

P( n ) =

2

Segregation intensity: Is = σ a2 / aav bav

The segregation intensity varies between 0 and 1. For the mechanical mixtures of immiscible polymers the parameter IS is close to 1, thus other measures of mixedness must be found. An excellent measure of mixedness is the specific interfacial area, or the interfacial area per unit volume of the mixture, AV. For completely separated phases, AV = 1 m2/1 m3 = 1 m-1. As the mixing progresses, the value of AV increases — larger is its value, better is the blend’s mixedness. In laminar shear mixing AV of randomly oriented elements depends on strain, γ:

∞

σ a2 =

∫

−∞

∞

P(n)dn ;

∫

SL = R(r )dr / 0

∞

∫ P(n)dn

−∞

(9.29)

595

Compounding Polymer Blends

9.1.10

Morphology Evolution During Processing

Blending is an economically viable, versatile method of manufacturing new materials with a wide range of properties [Utracki, 1987, 1993; Rauwendaal, 1986]. The parameters most frequently targeted for the improvement by blending are: the mechanical properties, impact strength, processability, heat deflection temperature, and the cost-to-performance ratio. When designing a blend, first the polymers and their approximate concentrations must be selected, then the most appropriate blend morphology for the envisaged application. For example, the need to improve impact properties implies that the minor phase ought to be dispersed as spherical droplets with micron or submicron diameter. On the other hand, if the material is to be used in vapor on solvent barrier applications, the minor phase should be dispersed in the form of relatively large, thin lamellas. The desired morphology is obtained by selecting the appropriate compatibilization and compounding/processing methods [Utracki, 1993, 1999].

Table 9.6. Mixer

9.1.10.1 Sample Collection for the Morphology Characterization Several methods have been proposed to follow the morphology evolution in an extruder. Ideally, the method should make it possible to instantaneously freeze the structure, without imparting changes due to either stress, stress relaxation, or coalescence. The frozen structure may then be analyzed for the degree of fusion, crystallinity, composition, degree and type of dispersion, etc. Table 9.6 lists some of the more popular methods used for the characterization of morphology evolution during blending in an internal mixer, an extruder or a special mixing device. Blends’ morphology has been studied by several research teams, viz. Elemans et al. [1988, 1990], Lindt et al. [1992], Shi and Utracki [1992], Huneault et al., [1995b], Delamare and Vergnes [1996], Cho and White [1996], etc. There are numerous small single- or twinscrew extruders developed for the preparation of small quantity of blends and evaluation of their morphology. These are particularly useful for compatibilizer optimization. The SSE type recirculating devices are well represented by

Experimental methods for morphology characterization Description

SSE, The oldest, and the most primitive method is the, so call “carcass method”, viz. stopping now used the motor, cooling the compound, and removal of either screw(s) or a barrel to access also for TSE the solidified resin. Due to the thermal inertia of the system, the minimum time for the solidification is about 15 to 20 min. During this period extensive changes of blends’ morphology may take place. The observed structure can hardly reflect the structure existing during the blending. TSE Not much better is the dead-stop method and sliding out the barrel (Clextral BC 45 mm extruder). Three methods were described: 1. Cooling the barrel with water requires at least 5 min. 2. Sliding out the barrel and then cooling by sprinkling water on the melt still requires several minutes. 3. The best choice is to slide out the barrel and then to take molten samples from the hot screws, and quench them in cold water. The methods 2 and 3 may affect the morphology. Significant differences of the average drop size were reported between samples collected using the three methods.

Reference Maddock, 1959

de Loor et al., 1994, 1996; Delamare and Vergnes, 1996

596

Table 9.6.

L. A. Utracki and G. Z.-H. Shi

Continued.

Mixer

Description

Reference

TSE

Similar to the sliding barrel is the method that uses a clam-shell barrel. The morphology is quenched by either pouring a cooling liquid (ice water or liquid nitrogen), or quenching the specimens taken from various locations in a cooling medium. Opening the barrel causes deformation of the molten blend. Here the time lag is 2-5 min. Prior cooling of the extruder and opening it after re-heating the surface requires a comparable time-lag to the carcass method, 15 to 20 min. A Baker-Perkins 50.8 mm and more recently Werner & Pfleiderer up to 70 mm are popular TSE’s that use this method. A quenching double barrel clam-shell section was designed and manufactured for a ZSK-30, to replace two barrel segments. The section played the same role as the regular barrel elements, but it had extra quenching channels for chilled water (4°C). Once blending reached a steady state, the screw rotation and heating were stopped and the quenching unit was switched on. It took 4-6 s to quench the polymer near the barrel wall and about one minute to quench the position near the screw root. It was calculated that in the kneading discs quenching took about 10 sec. The advantage of this method is the ability to quench a blend at any location along the screw, within seconds without deforming the material. TEX extruders with sampling nozzles along the barrel were introduced. The nozzles have been designed to follow the chemical reactions during the blending, but they can be used as well to follow the morphological changes. The disadvantage of the system is that the sampling barrel elements can be placed only at the locations where sufficient pressure is generated to force the melt through the nozzle. Furthermore, depending on the morphology and stresses in the nozzle, the blend morphology may be modified. The quenching time was estimated as 3-5 sec. The most rapid method of quenching involves scooping the specimens directly from an operating extruder and quenching these in either ice water or liquid nitrogen. It takes ca. 1 s to quench the sample. Furthermore, scooping does not impose extensive changes of morphology. Its disadvantage rests with the limited number of sampling ports. Large differences of morphology were observed within the first few seconds after TSE screws were stopped. A “dispersion tester” consists of rotating drum with a milled grove, in which a stationary spreader was placed. The gap, clearance, drum speed, temperature, pressure on and the shape off the spreader could be controlled. The specimen could be subjected to repeated passages, with the morphology analysis after each passage. The device was used to analyze the morphology generation of PS/LDPE blends with and without compatibilizing SEB di-block copolymer. The observed morphological changes are summarized in Figure 9.15. The mixer is capable of mixing ” 1 g of the resin. Recently, the device was reported to produce similar dispersions of reactive blends as those prepared in industrial mixers, TSE or internal mixer. It was found that Mini-Max makes it possible to pre-evaluate blends, especially those of polymers having low viscosity and elasticity.

Kalyon et al., 1988; Sundararaj et al., 1992

TSE

TSE

TSE

Special

Mini-Max

Bordereau et al., 1992; Shi et al., 1992

Nishio et al., 1990-2; Sanada et al., 1991; Sakai, 1993

Huneault et al., 1995b

Tadmor, 1988; Kozlowski, 1994

Maxwell, 1972; Sundararaj et al., 1995

Compounding Polymer Blends

597

mini-extruders, Microtruder™, having throughput • 10 g/hr [Luker, 1996]. Recirculating TSE with a sampling port in the bypass line was originally developed at DSM. Later, the unit was commercialized. Owing to the small size, its specimens can be rapidly quenched. Following on the success of the Extensional Flow Mixer (EFM — see Part 9.1.6.4) an extensional mini-mixer was designed. The device consists of two reciprocating pistons with a convergent-divergent, c-d, restriction between them. The material to be mixed is introduced to the mixing chamber, heated under vacuum until the desired temperature is reached. Next, it is forced repeatedly through the c-d restriction, by the action of the reciprocating pistons [Utracki, 1996]. 9.1.10.2 Modeling of the Morphology Evolution During the last few years several attempts were made to develop mathematical models capable of describing the morphology evolution during blending in different model flows and compounding machines. These efforts are summarized in Table 9.7, and briefly discussed below under separated headings. 9.1.10.3 Morphology Evolution in Internal Mixers

Figure 9.15. Morphology evolution during flow in Tadmor’s dispersion tester [Kozlowski, 1994].

A laboratory internal batch mixer was used by Shih et al. [1991, 1992] to simulate the progression of mixing in a compounding process of multicomponent polymer systems. A mixing bowl was modified with a glass window built into the front heating plate. A video picture taken through the glass window was recorded continuously. By using HDPE, PBT, and PAR, the authors distinguished four sequential characteristic states: I. Elastic solid pellets II. Deformable solid pellets III. Transition material A. Fluid with suspended solid particles B. Fractured or semi-fluid material C. Dough-like material IV. Viscoelastic fluid.

598

Table 9.7.

L. A. Utracki and G. Z.-H. Shi

Modeling of morphology evolution

Compounder

Modeling

References

Internal Mixer Internal Mixer

Flow visualization of blending elastomers and plastics. Experimental studies of mixing in an internal mixer, after installing glass window in a mixing bowl. Experimental and theoretical study of drop elongation in a steady-state flow, at constant shear stresses. Theoretical analysis of flow through a SSE using the “three-layer” model. Theoretical and experimental studies of the early stages of morphology development. Melting of two resins generated lamellas that during the flow through SSE were thinned out, then disintegrated into filament, which in turn broke into droplets. Pressure flow model for Newtonian liquid in the pumping section. Flow of Newtonian fluids in kneading disc elements and (right and left handed) screw elements. A simplified model was developed. It is based on the assumption that throughput is the same along the extruder as within the zone where the flights are full. Calculations of blend morphology evolution along an extruder and comparison with experimental data were published. The groove model to describe the flow in axially open screw channels by introducing two parameters, dimensionless pressure gradient and dimensionless output. Dimensionless calculations of throughput and pressure gradients for flow of non-Newtonian fluids in screw and kneading disk elements of a modular intermeshing self-wiping corotating twin-screw extruder, as well as the temperature rise. Experimental studies using Baker-Perkins corotating TSE. During the melting stage, first sheets, than sheets with holes labeled “laces,” and finally dispersed drops were observed. Theoretical and experimental studies of morphology evolution during compounding of PS/PE blends. A quenching barrel section was designed for a corotating intermeshing TSE (W&P). A predictive model for the morphology development was proposed. Morphology development along a corotating or counter-rotating intermeshing TSE was observed. Different screw configurations were used. Increasing the number of kneading blocks enhanced the degree of dispersion. Extension of the model developed by Shi and Utracki. Better fluid mechanics computations and microrheological drop fracture mechanisms were used. The model takes into account the coalescence. Experimental work. Samples were scooped directly from the venting ports located between the melting and kneading block sections, as well as after the latter section. After melting the dispersed phase consisted off fibers and droplets, both with diameters already below 10 mm.

Min and White, 1985 Shih et al., 1991, 1992

Rheometer Single-screw extruder

Twin-screw extruder Twin-screw extruder Twin-screw extruder

Twin-screw extruder

Twin-screw extruder

Twin-screw extruder

Twin-screw extruder

Twin-screw extruder

Twin-screw extruder

Corotating twin-screw extruder

Lyngaae-Jørgensen et al., 1991, 1993. Lindt, 1981; Ghosh et al., 1991; Lindt and Ghosh, 1992

Vergnes et al., 1983, 1986 Szydlowski et al., 1987 Elemans et al., 1988, 1989, 1990

Potente et al., 1989

Chen and White, 1992, 1993

Sundararaj et al., 1992

Bordereau et al. 1992; Utracki and Shi, 1992; Shi and Utracki, 1992, 1993 Lim and White, 1993

Huneault et al., 1995a

Huneault et al., 1995b, 1996

599

Compounding Polymer Blends

Table 9.7.

Continued

Compounder

Modeling

References

The kneading section (independently of its construction), further reduced the drop diameter by a factor of ten. Twin-screw extruder Experimental and theoretical. Non-isothermal model with drop diameter computations based on microrheology and coalescence. Changes in polydispersity of sizes can also be computed. Corotating twin-screw A 3D modeling of flow was carried out by finite elements method. extruder (CORI) The flow profile, the back-flow volume, pressure distribution, shear and elongation rates and adiabatic T-gradient were computed for the conveying and kneading sections. Co- or counter-rotating A 3D modeling of flow in full conveying screw elements of either a CORI twin-screw extruder or ICRR was reported. The flow velocity and stress fields were computed as well as the residence time distribution. CORI was found to have better distributive mixing, whereas ICRR performed better as a dispersive mixer. Twin-screw extruder Experimental work using CORI. Blends of 10 & 30 wt% PS in HDPE were prepared either by feeding dry blended pellets, or introducing molten PS to HDPE melt at different screw position. The morphology was insensitive to the feeding method. The authors concluded that the blend morphology is determined by the melt mixing within the first 25 sec.

The type of the transition material (No. III A, B, or C) mainly depends on the difference in the solid-liquid transition temperature of polymers. Furthermore, if the second polymer was crystalline, after melting a phase inversion has been frequently observed. For blends having widely different softening point (either melting point or glass transition temperature) the four stages were “shifted”. Thus, the polymer having the lower softening temperature melted first and became matrix of a “highly filled composite”. If the softening point and the concentration of the second polymer were high enough, the torque increased causing breakage of the shear pin in the mixer. In the extrusions, the consequences could be more serious. It was concluded that blends of polymers having widely different softening points should be fed consecutively, trying to match viscosities by judicious selection of the local temperature. Extensive studies of mixing in an internal mixer were carried out by Min and White [1985].

de Loor et al., 1994, 1996; Delamare and Vergnes, 1996 Goffart et al. 1996; van der Wal et al. 1996

Kajiwara et al. 1996

Bourry and Favis, 1997

The work combined the flow visualization with the modification of rotors and measurements of the pressure distribution. Detailed description of the blend morphology in an internal mixer is provided in a review chapter [Min, 1994]. 9.1.10.4 Blends’ Morphology in a Single-screw Extruder (SSE) Lindt and Ghosh [1991, 1992] dealt with the early stages of blend morphology development in a SSE. According to the model, melting of two types of pellets having similar solid-liquid transition temperature, generates lamellas that during the flow through SSE progressively thin out, then disintegrate into filaments, which in turn break into droplets. Comparison between experimental data and theoretically simulated striation thickness of polymers in a SSE channel showed good agreement. Unfortunately, owing to theoretical difficulties, the model did not go beyond the lamellas thinning. The authors argued that the

L. A. Utracki and G. Z.-H. Shi

melting zone plays an important role during continuous blending of polymers. Within the residence time in the melting zone the blend undergoes rapid morphological changes. Within a fraction of a second, the scale of mixing may decrease by several orders of magnitude.

PE-DROP DIAMETER vs. SCREW LENGTH 4

10

Q5, N150 Q5, N200 Q5, N250

d ( mm)

600

10

2

10

0

9.1.10.5 Morphology Evolution in a Twin-screw Extruder (TSE)

0

400

800

Position along the screw, L (mm)

PS-DROP DIAMETER vs. SCREW LENGTH

10

3

10

1

d ( mm)

Twin-screw extruders are becoming increasingly important tools for the plastics industry. Their principal advantages (over SSE) are: better feeding and more positive conveying, self-wiping of the screws, short residence times, a narrow residence time distribution spectrum, better mixing, larger heat transfer area, an easier scale-up, and large output (in 1999 up to 59 ton/h). The predictive models of the morphology evolution during polymer blending in a TSE start with a description of the flow mechanics. The fundamental assumption is that there are “strong” and “weak” zones. Within the former the screw elements are fully filled, there is a positive pressure, large deforming stresses, and usually material recirculation that offers a potential for modifications of the blends morphology. The “weak” zones are partially filled, made of mainly transport elements. Here the material is conveyed at either ambient pressure or under vacuum. The stresses are minimal. Within these regions two mechanisms are expected (provided that the reduced time, t*, is sufficiently long): breakup of filaments by the Raleigh instability mechanism, and coalescence. Table 9.7 summarizes the fluid mechanics models of flow through TSE. For more detailed discussion on the early modeling of flow through co- or counter-rotating TSE see White [1990]. Shi and Utracki [1992] proposed the first model of the morphology changes in a corotating, intermeshing TSE, CORI. The model incorporated the microrheological dispersion mechanism, but coalescence was neglected. Its validity was evaluated comparing the predictions with the experimentally measured drop diameter at different axial positions in the twin screw extruder. It was estimated that using the newly designed quenching barrel,

Q5, N150 Q5, N200 Q5, N250 10

-1

0

400

800

Position along the screw, L (mm)

Figure 9.16. Dynamic drop diameter as a function of position along the screws for φ = 0.05 of PE dispersed in PS (top) and PS dispersed in PE (bottom). The blend was extruded with a throughput of Q = 5 kg/h, at the screw speed N = 150, 200, and 250 rpm, at 200°C — lines are theoretical, the points are experimental [Huneault et al., 1995a].

the PS/PE blends were quenched within 7-10 sec. The model well predicted the morphology evolution of these non-compatibilized blends, but only to the last mixing block. In the assumed absence of coalescence, the model predicted continuous decrease of drop diameter, whereas the experiment indicated its stabilization. The second model refined these computations and introduced coalescence between drops [Huneault et al., 1995a]. The model is fully predictive. Good agreement with experimental data was found (see Figure 9.16).

601

Compounding Polymer Blends

More recently, these experiments were repeated using faster method for the samples’ collection. The specimens were periodically “scooped” from open ports in the extruder and quenched in ice water. It was estimated that the quenching time was t ” 1 sec [Huneault et al., 1995b]. Immediately after melting the dispersed phase consisted of fibers and droplets, both with diameters below 10 µm. The fibers were observed for φ • 0.02 — they were broken in the kneading section. The kneading blocks reduced the drop diameter to d ≅ 1 µm. Screw configuration had little influence on the resulting drop diameter. As the concentration increased so did the final diameter of droplets, as well as the fiber content. Evidently, in spite of the good agreement between the previous models and observations obtained after 7-10 seconds of the quenching time, the “incipient” blends’ morphology is more complex than the models predicted. The morphology changes rapidly within the first few seconds after blending is stopped — this rapid variation of morphology would be particularly important for the reactive blending. Furthermore, the morphology evolution during the melting stage may be more important than initially assumed. During this stage, the viscosity ratio may vary by several orders of magnitude, both locally (stress and temperature gradients) and along the screw length. Furthermore, the coalescence affects not only the drops, but fibrils as well. Other types of morphological changes during blending in TSE were also observed [Sundararaj et al., 1992, 1995]. The authors reported that both reactive and non-reactive blends in an extruder, internal mixer, or a miniature cup-and-rotor mixer, show similar morphological features. Initially, during the melting, the polymers stretch into sheets and ribbons that first broke into fibers then into drops.

9.2.

Blending Methods and Equipment

The term “processability” refers to the relative ease with which neat or compounded resin can be handled in production operation and equipment.

In most cases it is a synonym of a high throughput, low pressure at the die, thus low viscosity. For the blending operations another parameter is important — the energy consumed per mass or volume of the material, or the specific energy, Esp. The polymer compounding or blending involves: 1. Preparation of ingredients (drying, sizing, heating, etc.), 2. Premixing (dry blending, homogenization, breakage of agglomerates, fluxing, etc.), 3. Melt-mixing (usually with degassing), and 4. Forming, e.g., granulation, pelletization, or dicing. The most frequently applied operations are mixing, milling, extrusion, molding, and curing. The resin processability and the processability window are of main concern. Since they affect the material performance, their stability and control are of paramount importance. 9.2.1

Historical Evolution

Mixing is the oldest process. The first annular container with a spiked rotor (for rubber compounding) was developed in 1820 by Thomas Hancock. A more efficient, counter rotating twin shaft internal mixer with elliptical rotating discs was patented by Freyburger in 1876, while its sigma-blade versions four years later by Pfleiderer, and forty years later by Banbury. These machines provided adequate mixing, but (owing to the batch mode) the production was slow and the performance erratic. An alternative was extrusion mixing. Paul Troester started manufacturing single-screw extruders (SSE) in 1892, but only in 1935 the first SSE for thermoplastics was introduced. Four years later, Paul Leistritz built electrically heated, air cooled SSE, having L/D = 10, automatic temperature control, variable screw speed, and nitrided barrel. The machine is considered prototype of the modern extruders. The SSE offered continuous processing capability, but it was notoriously poor as a mixer. Furthermore, under the standard processing conditions, a significant non uniformity in the shear history had resulted in large temperature differences of the melt, ∆T ≅ 60°C. To alleviate the

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problem, numerous types of mixing screws, mixing sections, or “add-ons” (inserted between the extruder barrel and the die) have been developed, for example: 1. Mixing screws — S-shaped kneader, Eagle mixing, Turbine mixing, Cohen double-wave, or Maillefer; screws with pin-, blister ring-, fluted-, barrier-screw, Dynamic Extentional Flow Mixer (DEFM), etc. 2. Mixing sections — Dulmage, Saxton, Pineapple, Strata-Blend, and many others. 3. Add-on’s — Barmag add-on torpedo, Staromix™, Maddox, Cavity Transfer Mixer (CTM), Twente Mixing Ring (TMR), the Extensional Flow Mixer (EFM), Dynamic Melt Mixer (DMX), static or motionless mixers (SM) from Koch, Ross, or Kenics, etc. Many of these devices were developed in parallel with the evolution of twin-screw extruders (TSE) — in part to reduce the competitive advantage of TSE as far as their recognized capability for good mixing was concerned [Rauwendaal, 1986; White, 1990]. In 1937 Roberto Colombo and his associates developed an intermeshing, co-rotating TSE that during the early 1940’s was used by I. G. Farbenindustrie for the continuous reactive extrusion of poly-ε-caprolactam. During the early 1940’s, recognizing the need for improved compounding capability, Meskat and Erdmenger designed and manufactured intermeshing, co-rotating TSE’s equipped with self-wiping corotating disks. After the 2nd World War, the development of the TSE technology has been carried mainly in Germany, then in Japan. In 1959 Werner & Pfleiderer introduced ZSK – Zwei Schnecker Kneter, a CORI with segmented screws and barrel. The screws have been assembled from at least six different types of conveying and mixing elements. Japan Steel Works (JSW) started production of TSE in 1951. Thus, by the late 1950’s, TSE technology has reached end of the pioneering period — directions of the screw rotation, their intermeshing, principal left- and right- handed screw elements, the kneading and mixing discs, were designed, evaluated, and patented.

Recent years brought significant improvement in the throughput capabilities by providing screws with higher “free volume”, higher screw speeds (experimental machines operate at screw speeds exceeding 1500 rpm), and higher torque. Several new mixing elements as well as restrictors have also been introduced. Probably the most significant is the progress in the construction of the intermeshing, counter-rotating machines, ICRR. Several manufacturers offer TSE’s in changeable configuration: CORI or ICRR, both capable of operating within similar ranges of the barrel lengths and screw speeds. While developing either SSE or TSE machines for the mixing applications, the employed strategy was simple: when distributive mixing was required — the split-and-recombine the flow stream was introduced, when dispersive mixing was needed — the high shear stresses zones were incorporated. In spite of the recognized advantages of the extensional flow mixing, there has not been much effort to incorporate this concept into the processing equipment. The extensional flow mixer (EFM) is the only device that consciously utilizes the elongation flow for mixing. 9.2.2

Melt Mixers

The melt mixers are either batch or continuous type. The formers require lower investment cost, but are more labor-intensive, have low output and poor batch-to-batch reproducibility. Recent developments in process control and automation eliminated some of these disadvantages [Utracki, 1991]. The continuous melt mixers comprise: extruders, continuous shaft mixers and specialty machines — these will be discussed in the following part of this chapter. A brief overview of the melt mixing devices is given in Table 9.8. 9.2.2.1 Batch Mixers Before introduction of the intensive internal mixers, the mixing was accomplished on open mills — a slow, dirty process requiring skilled operators. An efficient, counter rotating twin shaft internal mixer with elliptical rotating discs was

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Compounding Polymer Blends

Table 9.8.

Melt mixing machines

No. Machine

1.

2. 3.

4.

Internal mixer batch continuous SSE TSE — CORI TSE — ICRR TSE — CRNI Special Pin-barrel Kokneader Planetary gear Disk

N

P

RTD

Feeding

Dispersive

Powder

Filler

Sticky

as such

+ 0 + ++ + +

+ 0 + -

+ + -

+ + 0 + + +

+ + + 0 +

+ + 0 + +

+ + + + -

+ 0 +

0 +

0 + + +

0 + 0 0

+ 0

+ + -

+ + ++ 0

Distributive

with additions

as such

with additions

0 + + 0

+ + 0 -

0 + 0 0

+

+ + 0 0

+

Symbols: N - high screw speed, P - pressure generation, RTD - residence time distribution. Feeding: Capability to be fed with difficult material, distributive and dispersive mixing without and with extra mixing elements or devices. CORI co-rotating intermeshing, ICRR, intermeshing counter-rotating, CRNI - counter-rotating non-intermeshing. Evaluating symbols: + means good, 0 means acceptable, and - means poor.

patented by Freyburger, while its sigma-blade versions by Pfleiderer and by Banbury. The latter two inventors started the manufacturing companies, Werner & Pfleiderer (W&P), and Farrel Corporation, respectively. An important feature of these machines is that they are enclosed and pressurized, so that fine powders and additives would mix into the compound, not drifting away. Batch mixing is not efficient for handling large capacities, but it is well suited for short-run operations, e.g., for the manufacture of colorconcentrate masterbatches and products that have to have tailored identity. Because of the high stresses, an internal mixer can complete a cycle in minutes. The residence time, shear, and temperature can be controlled — all of these may be of critical importance in the compounding of heat-sensitive materials or in the alloying of two materials of varying melt indices. Batch mixing

also encourages close monitoring of formulations that combine expensive ingredients produced in small quantities. Several basic designs of the internal batchtype mixers are available. The most popular are the laboratory mixers manufactured by, e.g., Brabender or Haake, and their homologues on the larger scale manufactured by, e.g., Banbury or Moriyama dispersion mixers. The laboratory machines are often used for the evaluation of compounds on a small scale, prior to production. The test requires a small amount of materials, short time, little efforts and operational expense. The scale-up of the test results to production size is usually done through the “unit work” concept. The mixer makes it possible to assess the effects of changes in temperature, torque, and shear characteristics.

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Most serious drawback of these internal mixers is the heat conduction by the mixing shafts. The mixer chamber is only heated from the outside — the shafts conduct the heat away from the molten polymers. During mixing at relatively low temperatures, a dynamic equilibrium may be reached —heat from the heaters is transferred to the melt, then the shafts conduct it to the internal drive system of the mixer. The dynamic equilibrium means that higher temperature is at the chamber wall and lower on the shafts’ surface. However, for blending the engineering or specialty polymers with high melting point, the temperature drop across the chamber can be as large as 100°C. In the Readco High-Intensity Mixers, the agitators ascertain that the material is continuously moved in a circular pattern from one side of the chamber to the other, undergoing compression and expansion. At the same time, due to the helical angle of the agitators and the offset between the outer and inner agitators, the material is moved from one end of the mixing chamber to the other. The mixer is recognized for providing good dispersive mixing. 9.2.2.2 Continuous Mixers Continuous mixing involves the continuous loading and unloading of components. When properly performed, mixing decreases the compositional variations to the desired level. Continuous operations have the advantage of providing a stable process. The power consumption is usually lower than in batch operations. The stresses are imposed systematically, either in the shear or in the shearand-elongation mode of deformation. In spite of high capital costs and complex mixing the continuous mixing is easy to justify on the basis of the production volume (mixers with throughputs up to 80 ton/h are available) and quality. The continuous mixers make it possible to control: the feed rate, screw speed, temperature, as well as the discharge orifice setting and temperature. The new machines that operate at high screw or shaft speeds require a short residence time. However, care must be taken that high

speeds will not lead either to excessive specific energy consumption, or thermal decomposition of the resin. The continuous mixers require high capital investment, but are easy to automate and robotize, have high output, and can be run with a statistical quality loop control [Kearney, 1991; Canedo and Valsamis, 1994]. Extruders Extrusion is one of the most important forming methods in polymer processing. Virtually all polymers go through an extruder at least once, viz. compounding and pelletizing reactor powders. Furthermore, most forming operations involve extrusion, viz. extrusion of profiles, films, sheets, fibers, wire or paper coating, injection or blow molding, thermoforming, etc. More thermoplastics are converted into useful products by extrusion than by any other method. The word “extrude” originates in the Latin words “ex” (out) and “trudere” (to thrust), and thus it closely described the process itself as shaping by forcing through an exit opening called the “die”. In the 19th century, extruders were slow, thus rightly called “plodders”. The extrusion is accomplished by a screw rotating in a cylindrical barrel. The feed to the extruder may be either solid or liquid. In the plastics industry the solid-fed machines are called “plasticating extruders”, while that fed by melt are known as “melt extruders”. There are three layers of the basic functions that extruder must perform: • Primary: melting, pumping and forming. • Secondary: devolatilize and mix. • Tertiary: conduct chemical reactions. The extruders are classified according to the principal element of their construction, as: • Single-, Twin-, & Multi-screw extruders. • Single-, Twin-, & Multi-shaft compounders. • Gear or Disk extruders, e.g., Maxwell Meltelasticity Extruder, Tadmor’s Disk Extruder,... • Special extruders, viz. Gelimat, Patfoort, etc. In a plasticating extruder the homopolymers, copolymers, or their blends are introduced at one end. As they advance along the extruder length,

Compounding Polymer Blends

they are melted, homogenized, reacted, devolatilized, and transported either to a die (e.g., for extrusion or blow molding) or into the mold cavity (i.e., for injection molding). The extrusion that aims to produce pelletized material is conducted at higher screw speeds, using more intensive mixing than that aimed to manufacture the finished product. Hence, it is appropriate to distinguish: the compounding extrusion and the forming extrusion. The extruder operation is supported by an army of ancillary equipment. Their type and quality depend on the specific machine, type of operation, and material to be extruded. It is advantageous to divide this equipment into pre-extrusion, and post-extrusion type. To the first category belong: dryers, dry-mixers, pre-compounders, solid and liquid feeders, etc. To the second group of postextrusion equipment belong: gear pumps, add-on mixers, dies, pelletizers, coolers or chillers, dewatering systems, sizing equipment, and many others. Operation of an extruder is conducted considering the throughput, Q, vs. the pressure drop at the die, P, relation, schematically illustrated in Figure 9.17.

Large die Deep screw

Q

Mixedness

limit

Degradation

limit

Small die

Shallow screw

Economic

limit

0 0

P

Figure 9.17. The throughput, Q, vs. pressure drop at the die, P, dependence for a SSE. The operational triangle is defined by three limiting lines: the economy determined by the lowest acceptable value of Q, the mixing capability determined by the stresses (thus, by the screw speed, N, and P), and by the degradability limit, also determined by N and P.

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Older SSE’s have been equipped with minimum instrumentation: a pressure gauge at one point along the barrel (usually at the head) and a thermocouple in the hot melt region. For non-critical operations the operator would record pressure, temperature, screw speed, and mass flow rates (typically determined by the sample weight-time method). In operations requiring close tolerances of the extrudates, a greater degree of instrumentation and automation is employed. Usually, several pressure transducers and thermocouples along the barrel are used to ensure uniform extrusion and to control barrel and stock temperatures. Some designs may include thermocouple on the screw to monitor and control conveying flights. For specific applications a close-loop control may involve monitoring of composition, density, morphology, rheological performance, and other variables pertinent for the product. Standard methods of monitoring these variables may be used, viz. radio-isotopes, ultrasonics, rheological detectors, near and far infrared spectroscopy, mechanical testing, clarity scans, etc. 1. Single-screw extruders (SSE) The single screw extruder is a relatively inexpensive machine for small or medium size production lines. For lines with the throughput exceeding 10 tons per hour, the capital costs of SSE and TSE are comparable. SSE that can produce 40 tons/h can be custom-built by only few manufacturers. For example, in 1996 Berstorff demonstrated 600 mm machine for homogenization and degassing of LDPE at a throughput of 23 tons/hr. SSE’s are characterized by a simple design, raggedness, reliability, are easy to operate and maintain, have favorable performance-to-cost ratio, and theoretical description of their operation is well documented. On the other hand, they are difficult to scale-up, are notoriously poor mixers with broad residence time distribution and relatively long residence time [Hold, 1982]. Over the years, SSE’s have been made more versatile by introduction of special mixing screws, by using add-on mixing devices, by utilization of two or more extruders operating in tandem or as coextruders, etc. Refinements have been

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developed in feed preparation and feeding methods, including improved drying and premixing techniques, as well as downstream feeding of some formulation components (such as glass fibers). The principal components of a single-screw extruder are illustrated in Figure 9.18. The machine has a motor drive, a gear train, and a screw that is keyed into the gear reducing train. The fluid layer between the screw flights and the barrel wall maintains the screw balanced and centered. Modern units are equipped with continuously variable speeds and electrically heated barrels. The barrel can be “zoned” according to the number of controllers or the heater bands. Depending on the application and type of service, the screw may be cored for heating or cooling. The standard barrel (inside) diameter ranges from 12 to 600 mm. Larger extruders are available by special order. Standard screw has length-todiameter ratio: L/D = 16-36, with 20 ” L/D ” 24 being most common (feed length 4-8D; transition length 14-16D; and metering zone length 6-10D). It has a single parallel flight, with pitch = 1D and pitch angle φ = 17.66°. The flight width is usually about 0.1D, the channel depth in feed section is 0.1-0.15D. The channel depth ratio varies from less than 2 to 4. The compression ratio, CR, has been defined as a ratio of the screw flight volume

at the entrance to that at the exit. The CR can be expressed as a function of the internal barrel diameter and external diameters of the screw, viz.

[( ) ( / [(πD / 4) − (πD

)]

CR = πDb2 / 4 − πDS2, feed / 4 / 2 b

2 S , die

∴ CR =  DS, feed + H1 

2

(

[( D

S , die

+ H2

)

2

)

−

/4

)]

(9.30)

DS2, feed  

− DS2,die



/

]

For the standard screw, CR varies from 1.5 (for rigid PVC) to 6 (for PA). The pressure generated at the die ranges from P = 70 to 200 MPa, or 10 to 30 kpsi. The rotational screw speed ranges from N = 20 (for rigid PVC) to 360 rpm (for PE). There are three zones in the SSE: the solid conveying, the melting, and the melt conveying. In addition, there is the up-stream (feeding), and down-stream (cooling) zone. Owing to the importance of feeding to the extrusion process, it is appropriate to start with that zone. Since SSE operates at the fully flooded conditions, the quality and throughput depend directly on the feed. The feed zone consists of a hopper

Figure 9.18. Schematic representation of a single-screw extruder, SSE.

Compounding Polymer Blends

or feed arrangement and a solid stock conveying region. Its purpose is to transfer a polymer from the feed hopper into the barrel, where it is initially compressed. This compression forces the air out from between the interstices of resin pellets or rubber chunks (air is expelled back through the hopper). It also breaks up lumps and polymer agglomerates, creating a more homogeneous feedstock that can be readily melted. Flow in hoppers depends on their type (gravitational or forced), bulk density of the material, the shape and location of the feed inlet (direct or tangential feed, straight or with a chamfer on the down-going side), as well as on the type of screw and the barrel. For good feeding the flight volume in the feed zone must be as large as possible. Since the solids’ conveying ability of an extruder depends on the ratio of the friction coefficients on the barrel and on the screw root, frequently grooved barrel (cooled by circulating cold water and thermally insulated from the rest of the barrel) is used in the feed zone. Its advantages are: increased throughput, greater extrusion stability, and smaller effects of pressure variation at the die on the flow. For this one pays by higher torque and pressure at the die, high wear and higher energy losses caused by the extra cooling. The solid conveying zone begins at the feed inlet and extends to a point where the solid particles just begin to melt. The screw section is characterized by a deep flight between the root of the screw and the barrel wall. In most cases the pitch (i.e., the screw length for a single turn of the screw flight) is 1D, that means that the screw pitch angle is 17.66° [Darnell and Mol, 1956]. Here the solid bed is compressed and it moves forward as a plug. The flow causes the material to be compressed and to melt under the influence of frictional forces between the barrel surface and the solid plug. There are two important factors controlling these operations: the need to remove the entrapped air, and to ascertain high friction coefficient on the barrel. Removal of air is underrated activity that may seriously affect the performance of the extruder. Its importance depends on the type of feed as well as on the

607

shape of the feed inlet. For powder feeding a vacuum-assisted air removal may be required. The melting zone follows the solids conveying zone. The purpose of this zone is to compress the solid bed and to provide intense friction between it and a barrel. Within this zone the channel depth decreases, and the solid polymer coexists with its melt. In a SSE, this zone extends over a major part of the extruder, e.g., L(melt) = 10-14D. Melting is caused by the two principal forms of energy: the thermal, applied to the polymer by conduction from the external heaters through the barrel surface, and the mechanical converted to heat through friction and viscous dissipation. The friction coefficient of a polymer on the metal surface (f) depends of pressure (increasing P by a factor of 10 increases f by a factor of 3 to 4), temperature (e.g., for PE: f = 0.6 - 0.002T°C), and the sliding velocity. The effect of the friction coefficient ratio: f(barrel)/f(screw) is illustrated in Figure 9.19. The conveying is controlled by the difference in the friction coefficient between the solid bed and either the barrel or the screw. It is important to maximize this difference. Good balance of heating and shearing within this zone is crucial. For the heat-sensitive polymers such as PVC, too much shearing may cause degradation. For this reason the frictional heat generation should be kept low, e.g., by means of low screw rotational speed (N = 20 to 30 rpm). Good screw design should efficiently eliminate unmelted particles. Shallow screws and high pressure give better melting, but their throughput is relatively low and there are problems with scaling. Deep screws at low N provide better melt homogeneity and throughput stability at the cost of inefficient mixing. The best compromise seems to be high N and shallow screw. The length of the melting zone should also take into account the type of feed — powder feeding requires longer melting zone than pellets. However, vacuum-assisted removal of entrapped air significantly improves the melting rate and the melt homogeneity. The melting zone is particularly important for the generation of blends’ morphology. The variation of temperature and high stresses encountered within this zone may profoundly affect the extrudate quality.

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6

f / f = 3.0 Calculated pressure increase as a function of relative friction coefficients and # turns

P (kPa)

4

B

800 f B / fS = 2.0

10

f / f = 1.5

2

B

10

f / f = 3.0

S

S

Q (kg/h)

10

Calculated throughput for: H = 0.01, fB = 0.5

B

S

f / f = 2.0 B

400

S

f B / fS = 1.5

f B/ fS = 1.0 0

10

0

Number of 1screw turns

2

200 20

40

60 N (rpm)

80

Figure 9.19. Effects of the friction coefficient ratio on the pressure evolution within the feed zone (left) and on the throughput vs. screw speed (N).

The melt conveying zone begins at the point where all solid particles are melted and it extends to the outlet of the extruder, thus it encompasses the screen-pack, breaker plate and the die. Once the resin is in the molten state, the extruder acts on it as a pump, transferring, homogenizing the molten polymers, and building up the pressure to the level required to force the material through a discharge nozzle or a die. SSE is capable of delivering up to 70 MPa or 10 kpsi of pressure. In the classical approach to the mathematical modeling of flow through SSE, the screw and the barrel are unwound, the screw is assumed to be stationary and the barrel moves over it at the correct gap height and the pitch angle, φ. Near the tip of the screw there is the melt metering or the pumping zone, where the pressure builds up. In this region the polymer melt is essentially homogenized and raised to the proper temperature for the formation of products. Many types of screws have been developed. The screw configuration should be selected considering properties of the material being processed. A constant-pitch metering screw is usually employed in applications not requiring intensive mixing. Where mixing is important, for example for color dispersion, a two-stage screw equipped with a let-down zone in the center of the screw is used. Mixing promoters can also be incorporated in the metering zone or near the tip of the screw [e.g., see Gagliani, 1991]. Figure 9.20 shows two screws developed

for improved dispersive mixing. All the material must pass through the narrow slit between the restriction and the barrel. To enter the slit, the material is exposed to an elongational flow field, while within the slit it is subjected to high shear stress. Many other screw modifications have been developed to enhance mixing. In some applications with two-stage screws, venting at the let-down section may be needed [Rauwendaal, 1991]. The Buss Kneader is one of specialty SSE-type mixing extruders. The machine has relatively narrow residence time distribution and relatively low residence time. Here, the constant channel depth screw consists of interrupted flights and three rows of stationary pins in the barrel. During extrusion the screw rotates and axially oscillates providing good self-wiping of the blades. This periodic reciprocation minimizes the material flow in one direction only. The intensive shearing strain is periodic and the gaps tend to reduce the effective shear stress. The large surface areas of exposed materials are a result of shear stresses generated by high torque. There is a good deal of distributive mixing during the high-stress dispersive process. The clam-shell barrel design makes it easy to clean, service, analyze, and optimize the process. The machine is used for polymer blending, compounding of PVC formulations, preparation of engineering resins (with and without reinforcement), extrusion of foodstuffs, carbon electrode pastes, etc. [Jakopin and Franz, 1989; Anonymous, 1989].

Compounding Polymer Blends

Figure 9.20. Two SSE screws with restrictions designed to improve dispersive mixing.

By contrast with SSE’s, the twin-screw extruders operate with partially filled screws. The polymers from a feeder are transported toward the pressure zone, usually created by a flow restrictor, e.g., a left-handed or reverse screw element. Under the pressure the resins are compressed into a solid plug. The plug fractures and melts mainly by friction between the fragments of the initially compressed solid bed. The process is highly efficient. Its length can be as short as less than one diameter, L ” D. There is evidence that melting in a TSE differently affects the generation of morphology than that in a SSE. 2.

Twin-screw extruders (TSE)

2.1. Introduction to TSE The first patent on the precursor of the TSE dates from 1869. However, the first commercial CORI for handling polymers was marketed 70 years later. Several intermeshing counter-rotating twin screw extruders (ICRR) were developed in the 1920-30’s. After the 2nd World War, the corotating, fully intermeshing TSE (CORI) started

609

to dominate the field. Originally, CORI was developed for dewatering, devolatilization and conduction of chemical reactions. The names of Meskat and Erdmenger are prominent on the early patent. The latter author is credited with designing the kneading blocks for the dispersive mixing, sometimes called “Erdmenger eggs”. The prototype modular CORI was built by Werner & Pfleiderer (W&P) in 1955. The non-intermeshing counter-rotating TSE (CRNI) also dates from the 1920’s when it was first designed by Ahnhudt for mixing and pumping rubbers and plastics. However, already in 1930 the Welding Engineers started to develop CRNI compounding extruders. In 1950’s the company introduced a modular design. An excellent record of the TSE history is part of the monograph on the twin-screw extrusion [White, 1990]. The TSE development work was completed by the mid-1950. The modular design of screws and barrel was universally accepted. The main types of screw elements (conveying forward, conveying backward, mixing, and neutral) were developed. Looking today at the screw elements from dozens of TSE manufacturers one must but note how similar they are. For example, the bi-lobal (eggshaped) Erdmenger’s mixing elements are universally popular. They provide high shear stress in the overflight (between the mixing element and the barrel) as well as in the direct intermeshing (between two mixing elements) regions. Many TSE manufacturers still insist that to control the dispersive-to-distributive mixing ratio simply the width of the mixing block should be used — a set of narrower blocks resulted in more distributive, and less dispersive mixing. TSE’s are classified according to the following three categories:

Screw rotation

Intermeshing

Confinement

1. Co-rotating 2. Counter-rotating

A. Separated B. Tangential C. Intermeshed

a Lengthwise closed or open (L-closed or open) b Crosswise closed or open (C-closed or open)

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L. A. Utracki and G. Z.-H. Shi

In total, seven types of twin screw extruders are on the market: I. Counter-rotating, not intermeshing (mostly tangential), L & C-open (CRNI) II. Counter-rotating: partially intermeshing, L & C-open (rare) III. Counter-rotating: partially intermeshing, L-open & C-closed (rare) IV. Counter-rotating, fully intermeshing, L & C-closed (ICRR) V. Co-rotating: not intermeshing, L & C-open (rare) VI. Co-rotating: partially intermeshing, L & C-open VII. Co-rotating, fully intermeshing, screw: L-open & C-closed, discs: L& C-open (CORI — the most popular). The three most popular types: I, IV, and VII are indicated in bold. In Table 9.9 a partial list of TSE manufacturers is provided. Whether the screws are open or closed lengthwise or crosswise, have a direct effect on the conveying capability, mixing, and the pressure buildup capacity of the system. For example, the non-intermeshing systems are open lengthwise and crosswise. Fully intermeshing, counterrotating systems can be closed lengthwise and crosswise — here the material is locked in closed

Table 9.9.

Selected manufacturers of TSE

No. TSE-type 1. 2.

3. 4. 5. 6.

chambers (this is also the case for the screw pumps). The theoretically impossible systems are: (i) lengthwise and crosswise closed CORI, and (ii) lengthwise open and crosswise closed ICRR. Thus, CORI’s are open lengthwise (normal screw flights are closed crosswise, and staggered screw discs are open crosswise), while ICRR are lengthwise and crosswise closed. A screw system lengthwise open has a passage from the inlet to the outlet of the apparatus. This means that material exchange can take place lengthwise along the channel. In a closed arrangement the screw flights in the longitudinal direction are closed at intervals. It is important that the cross section of the screw channel be open in order for the material exchange to take place from one flight to the other in a direction normal to the screw channel. There is usually some leakage over the screw crests and through the areas required for the mechanical clearances. The great advantage of TSE’s stems from the modular design of barrel and screws. The manufacturers provide a diversity of the barrel elements (usually their length, L = 4D, sometimes 3D), that can be assembled in the configuration required for specific applications, viz.: • gland (or sealing) barrel, • feeding barrel,

CORI - low-to-medium speed CORI - high speed

Manufacturer

Bandera, Colombo, Creusot-Loire, Ikegai, JSW APV, B&P Process Equipment, Betol, Berstorff, Clextral, Davis-Standard, Delaware Extruder, Egan, Ermefa, Farrel, ICMA San Giorgio, Ikegai, IKG, JSW, Kobe Steel, Leistritz, Mapré, Maris, Mitsubishi, OMC, Pomini, Reifenhäuser, Rockstedt-Farrel, Teledyne-Readco, Theyson, Toshiba, Werner & Pfleiderer ICRR - parallel screws; high speed Amut SPA, Bandera, Battenfeld, Bausano, Clextral, Ermefa, Ikegai, JSW, Krauss-Maffei, and long L/D Leistritz, Maplan, Mapré, Reinhäuser, Toshiba, Weber ICRR - conical (mainly for AGM, Anger, Cincinnati-Milacron, Haake, JSW, Maplan, Mitsubishi, Toshiba, Weber PVC profiles) CRNI - equal length screws; Bandera, Bausano, Bolling, Farrel, Davis-Standard, JSW, Kobe Steel, Leistritz, Pomini high speed and short L/D CRNI - non-equal screws JSW, Welding Engineers

611

Compounding Polymer Blends

• • • • • •

standard-closed barrel, venting barrel, side feeder unit barrel, injection-sampling barrel, melt pressure barrel, mixing barrel, e.g., with hexagonal channel, or profiled grooves, • main gate barrel (a closing unit), and others.

• • • •

polygon shifting elements, continuous mixing block, Erdmenger’s bi- or tri-lobal kneading blocks, low energy distributive vane mixers (usually the last elements before the screw tip, for homogenization of the melt within the highest pressure zone), • and many others.

Similarly, the screws are assembled on shafts from selected elements. These are usually classified according to their functions into three categories: conveying, neutral (e.g., end or centering disks, screw tip), and mixing. There is again diversity of elements within each of these categories — the greatest being for the mixing type. The conveying elements may differ in the: 1. number of treads, 2. the helix angle or pitch, 3. the length, and 4. the thread direction (forward moving the material, or right handed, and back moving, or left handed). The elements are usually provided with a pitch angle of 14, 20, 28, 42° and reversed -14 and -20°. As the pitch increases, the material is pumped faster, but at the same time the dwell time and the conveying capability decreases. The conveying elements exercise very little shearing and mixing action on the transported material. The mixing elements are available as individual elements or as their assemblies. There is a great diversity of these elements, for example: • perforated or slotted restrictor elements, • compression elements, • shearing elements, • mixing turbines, • multi-lobal elements (bi-lobal for CORI, tri- and hexa-lobal for ICRR)

The multi-element kneading blocks are assembled of discs of different thickness (to adjust the ratio of dispersive to distributive mixing) and displacement angle (see Table 9.10). Blocks constructed from the same number of elements, but of different thickness also behave differently. Wide discs provide mach greater sharing action than the narrow ones, i.e., the former provides dispersive mixing while the latter the distributive one. As the disc thickness in the block decreases its conveying capability increases. 2.2. Main types of TSE Out of the seven types of twin screw extruders, the following three are of principal interest (in order of increasing commercial importance): intermeshing counter-rotating (ICRR), counter-rotating non-intermeshing (CRNI), and co-rotating intermeshing (CORI). Table 9.11 provides examples of TSE specifications. Some manufacturers (e.g., Leistritz) offer four types of extruders to be mounted on the same base, viz. CORI and ICRR with D = 34 mm, L/D = 42, and centerline distance of CL = 30 mm, CRNI with D = CL = 34 mm, L/D = 42, and SSE with D = 36 mm, L/D = 40. The screw speed is ” 390 rpm and the torque (for TSE) = 2x123 Nm

Table 9.10. Effect of the twist angle on the performance of a block made of six Erdmenger’s kneading discs Angle (°)

Mixing

Shearing

Conveying

+30 +60 +90 -30

low medium high highest

low medium high highest

high forward medium forward nil medium backward

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Table 9.11. Example of TSE specifications No. Manufacturer

Type

D (mm)

L/D

N (rpm)

Torque (kNm)

1. 2. 3. 4. 5.

Bandera (PVC) Bandera W&P Theyson JSWa**

ICRR CORI CORI CORI

46 - 175 57 - 170 25 - 315 30 - 140

18 - 22 24 - 46 42 - 48 28 +

3 - 45 150 - 350 220 - 1200 300 - 500

0.75 - 25 0.95 - 19.5 0.16 - 340 ” 40

6. 7. 8. 9. 10. 11. 12. 13.

Leistritza Welding Engineers Clextral Berstorff Kobe Steel Kobe Steel Farrel - FCM Farrel - FTX

CORI/ICRR CORI/ICRR CRNI CORI CORI CORI CRNI CRNI CORI

32 - 460 18 - 135 20 - 50 25 - 200 25 - 305 32 - 315

” 65 20 - 60 37 - 50c ” 60 ” 63 36 5-10

185 -450b 400 - 500b 400 - 525 ” 400

0.024 - 313 0.021 - 11

22.5-125

24-36

200-526 350-1050 250-450 400-500

0.19-338 37-7,000* 150-6,000* 3.7-670*

Notes: a The same specifications are listed for co- and for counter-rotating fully intermeshing TSE’s; b On a request machines can be furnished with N ” 1500 rpm. c On a request L/D = 120 can be provided. * Motor power in kW. ** JSW stands for Japan Steel Works, also known (esp. in Japanese patents) as Nippon Steel Works.

and (for SSE) = 320 Nm. For the standard barrel, the maximum pressure and temperature was, respectively, P = 30 MPa and T = 400°C. For many years, CORI has been the universal compounder of choice. It operated at higher screw speeds with longer barrel than ICRR, advantageous especially for the reactive processing. However, during the last ten years or so, several manufacturers of these two types of extruders started to demonstrate that unbiased test results frequently show better performance of ICRR as far as the compound quality and the throughputs are concerned. Even at lower screw speed ICRR frequently outperformed CORI. This resulted in renewed interest in redesigning ICRR. Tables 9.12 and 9.13 list specifications from JSW and American Leistritz, respectively, for their standard machines that can be operated either as coor as counter-rotating fully intermeshing TSE. Since these manufacturers provide a diversity of processing equipment, they neither have a bias toward one of these types of machines, nor they pretend that one type can do all operations equally

well — for certain applications CORI may be better than ICRR and vice-versa. However, it is important to recognize the enlarged selection of the compounding equipment available on today’s market. Counter-Rotating Non-Intermeshing (CRNI) Originally developed for natural rubbers, foodstuffs, and elastomers, at present these machines are used as a stirred tank reactors or a twin-rotor continuous mixers. Their bearings are exposed mainly to radial, not axial, forces. By using one screw longer than the other, one can separate the mixing and pumping functions. In CRNI the material flow is based on drag, not positive pumping. There is a low shear stress field, responsible for the absence of dispersive mixing. However, the interchange of material between the screws provides good distributive mixing. The chemical reaction proceeds on the continuously renewed surfaces, related to reorientation of the laminar flow patterns and the total strain. CRNI is well suited for the polymerization

613

Compounding Polymer Blends

Table 9.12. Specifications of co- and counter-rotation TEX; L/D=34 (Japan Steel Works, JSW, 1996) Model

D(mm)

N (rpm)

Torque (Nm)

Drive (kW)

32 47 58 69 83 97 113 130 152 174 196 214 246 268 300 327

524 564 528 523 556 548 433 432 300 297 298 220 210 200 190 185

402 1520 2864 4803 8235 13159 21170 31817 42914 64376 92012 88984 133280 171500 245980 306740

22 90 158 264 480 756 960 1440 1350 2000 2875 2050 2930 3600 4900 5950

SuperTEX- 30 αII SuperTEX- 44 αII SuperTEX- 54 αII SuperTEX- 65 αII SuperTEX- 77 αII SuperTEX- 90 αII SuperTEX-105 αII SuperTEX-120 αII SuperTEX-140 αII SuperTEX-160 αII SuperTEX-180 αII TEX-200 TEX-230 TEX-250 TEX-280 TEX-305

Table 9.13. Specifications of co- and counter-rotation ZSE with L/D = 20 to 60 (American Leistritz, 1996) Model Micro-18 ZSE-27 ZSE-40 ZSE-50 ZSE-60 ZSE-75 ZSE-100 ZSE-135

D (mm)

N (rpm)

18 27 40 50 60 75 100 135

500 500 500 500 500 500 400 400

of miscible, low viscosity systems, viz. polymerization of monomers, grafting, halogenation, and PP-visbreaking (degradation by the addition of organic peroxides). The main advantage of CRNI is long enough residence time, sufficient to complete slow reactions. The molecular weight of the generated polymer was found to be virtually independent of Q and N. Mild mixing during

Torque (Nm) 21 100 320 600 1100 2200 5000 11000

Vol (mL/D)

Drive (kW)

2.8 8.4 29 51 100 196 436 1125

2 11 34 63 116 242 373 746

compounding prevents the mechanical degradation of elastomers as well as of crushing glass microbaloons. CRNI can offer larger outputs, and more interchange of material between the two screws than the other type TSE’s. These machines are frequently used for the preparation of composites with fibrous fillers. In general, the closer the clearances and

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intermeshing the more rapid the build-up of pressure. The narrower the lands and the larger the clearances between the screws, the greater the longitudinal mixing [Wood, 1979, 1980]. Intermeshing Counter-Rotating (ICRR) ICRR was developed as a positive displacement screw pump for viscous, difficult fluids (e.g., coal-oil, ceramic or rubber compounds). This is the only TSE type that is fully, axially and radially, closed. The extrusion speed depends on the intermeshing geometry and the screw speed. The fully intermeshing ICRR has narrower distribution of residence times and better precision in controlling rapid reactions between liquid reagent and molten polymer than a CORI. The low speed ICRR’s have been used for PVC compounding and forming. At higher screw separations and speeds, the machines can be used for the incorporation of high viscosity toughening elastomer. For example, ICRR’s have been used for the reactive extrusion of ABS. During the 1990’s there has been a significant progress in making ICRR fully competitive with the more popular CORI. This has been done by opening the calendering gap and moving the shearing action toward the barrel. While the older ICRR operated at the maximum screw speed of 150 rpm, the new ones can turn at least at 500 rpm without excessive gap pressure. For some applications, ICRR offers unique advantage of strong extensional flow field, able to disperse high viscosity ingredients in a low viscosity matrix. Co-rotating Intermeshing (CORI) These are the most popular and commercially important twin screw extruders. The machines were originally developed in the 1860’s for pumping stone paste. Their advantage arises from the movement of the intermeshing surfaces in opposite directions, thus the melt free surface is continuously renewed and the screws clean each other. In addition, since at the intermeshing the material passes from one screw to the other (change of the drag direction) there is low probability that the material would go through the gap. Thus, there is no calendering pressure that

may cause the screws to bend (a possibility for the ICRR machines). This in turn permits to use higher screw speeds and longer barrels. In 1938 CORI was used for the polymerization of butadiene and PA-6. The latter resin was commercialized by I. G. Farbenindustrie in 1939. During the 1950’s and 1960’s the design criteria for these extruders were: 1. Incorporation without degradation. 2. High shear stress for dispersing. 3. Homogenization of materials with different viscosity. 4. Uniform stress and heat history 5. Precise control of the mixing process. The present tendency in CORI development is to: (1) maximize throughput from the given size machine by maximization of the extruder free volume, or the OD/ID screw ratio, (2) provide capability to efficiently transmit the required amount of power into the material, and (3) increase the extruder speeds. The modern machines from Werner & Pfleiderer are supplied with OD/ID = 1.55, those from Berstorff with 1.74 (for reactive compounding and degassing). The later company also developed several mixing elements (turbo or gear mixers) well suited for mixing resins with different viscosities. Smaller industrial CORI machines are available from JSW with the screw speed of N ” 1500 rpm. Similar speeds are also offered by Werner & Pfleiderer. There is also talk in the industry of super-fast extruders operating at 3000 rpm. Maris increased free volume of their 58 mm extruder from 84 mL/D to 103 mL/D — an increase by 23%. To make use of this extra throughput capability, the machines must have high torque capabilities. Screws are available with 1, 2, or 3 screw flights; the first has the highest solid conveying capacity (used mainly for feeding powders). Twoflight screws have large free volume per unit length and low average shear rate — they are mainly used for processing the temperature-sensitive materials. The three-flight screws are used when high shear stress is required. CORI’s are efficient in alternating the direction of applied stresses, providing distributive mixing by lamellae folding, as well

Compounding Polymer Blends

as controllable shearing (dispersing mixing). They are operated in starve-fed mode that demands high performance feeders. The material transport mainly depends on drag flow, with only local contribution from the screw pumping in the so-called pressure zones. The maximum conveying is achieved when the screw pitch is equal to screw diameter. CORI’s are used for compounding, polymerization, and devolatilization. CORI’s are the prime machines for polymer blending and reactive extrusion [Brown, 1992]. They have been used as reactors for the addition polymerization (polyacrylates, SAN, S-MMA, PA-6, POM, or TPU) and for the polycondensation (PA-66, polyarylates, PEST, PEI). Polymer grafting (polyolefin + silane, maleic anhydride, acetic anhydride, etc.) as well as mechanical and chemical degradation of polypropylene have also been carried out. The reaction is usually conducted within the pressure zone, e.g., in a kneading block section. This section is followed by left-handed (reversed) elements that control the residence time within the pressurized section. Devolatilization (removal of the reaction byproducts) is usually carried out immediately following the left handed screws section. The sliced screw mixing elements provide gentler mixing than the kneading ones (now two-lobed not three-lobed). There is a trend towards development of an intelligent compounding plant, where the polymerization, compounding and shaping will be carried out in sequence. Comparison of CORI and ICRR TSE performance There are several publications dealing with the comparison of efficiency of compounding, mixing, or reactive extruding in different types of extruders. Most of these studies suffer from the same aspect — the evaluated machines were not operated at a comparable level of performance efficiency. A summary of the reported observations is given in Table 9.14. Historically, the co-rotating geometry has been preferred by the plastics industry, as having more even distribution of stresses, providing easier control of compounding, and operating at higher

615

screw speeds and throughputs. The counter-rotating TSE’s were known as the “calendering” extruders, with high stresses existing between the two screws and low stresses outside. In ICRR, owing to the high calendering pressures, the screws could rub against the barrel causing premature wear. Thus, shorter barrels, slower speeds, and large intermeshing gaps have been recommended. One of the advantages of these machines has been the presence of the elongational flow field within the calendering zone. The machines have been successfully used in numerous applications requiring high dispersive stresses. During the last few years there has been renewed interest in upgrading the TSE performance. This was spurred by the advances in the theory and methods of analysis of elemental contributions. The analysis resulted in modification of the screw profiles (higher free volume of the process), separation of screws that made it possible for CORI and ICRR to operate at the same screw speeds, as well as in development of new mixing (or kneading) elements [American Leistritz, 1996]. The increased free volume (thus slender screw profile) resulted in lowering the average shear rate, thus the screw speeds needed to be increased, hence higher power/torque motors, which in turn improved the throughput. New kneading and mixing elements have been designed to improve either the distributive or dispersive mixing. To the first category belong the turbineand-gear mixing elements (viz. ZME, SME, or TME from Werner & Pfleiderer). The intermeshing TSE’s have self-cleaning screws. The material that adheres to the screw root may degrade and eventually fall off and be carried out with the product, showing up as contaminants [Wood, 1979, 1980; Salden, 1978]. Self-cleaning action is achieved in both counter-rotating and co-rotating screws, through an opposite roll-off or the wiping motion. The calendering flow does not take place in CORI, and a more efficient and uniform self-cleaning action is achieved. In CORI the material is transferred from one screw to another in a tangential path [Salden, 1978]. Here, edges of one screw crest wipe the flanks of the other screw with a tangentially oriented,

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constant relative velocity. There is a high relative velocity, and hence sufficiently high shear velocity to wipe the boundary layers. In ICRR the roll-off process between the screw crest and screw root and between the screw flanks simulates the action of a calender. The necessary shear velocity required to wipe the boundary layers is proportionately lower because of the low relative velocity. Counter-rotating screws require greater clearances between them since their mode of action is rather like a two-roll mill, passing material through the nip between them [Schoengood, 1973]. The material is drawn into the roller gap and is squeezed onto the

surface. Consequently, ICRR has less efficient self-cleaning action. The shear stresses and their distributions are higher in ICRR than in CORI. In the former it only depends on the screw pitch, while in the latter the shear stress distribution changes with the pitch angles, the throughput and the screw speed. In addition, the flow pattern in these two types is different, with a higher radial flow in ICRR. CORI’s are used for the reactive processing, compounding of alloys and blends, color masterbatches, and fiber-reinforced composition. The quality of the product critically depends on the selection of screws and their configuration.

Table 9.14. Advantages and disadvantages of CORI and ICRR type TSE Extruder

Advantages

Disadvantages

Co-rotating, intermeshing, CORI

* Self cleaning * Dispersive and distributive mixing in transverse and longitudinal direction * Lower than in ICRR, uniform and controllable stress * Good control of resin degradability * Local pressure control * High screw speeds and throughputs * Local high vacuum possibilities * Double-walled barrel for rapid temperature changes * High specific energy input possible * Residence time up to one hour

* Capital cost * Only residence times below about 10 min are economic * Limited ranges of pressure and vacuum are available

Counter-rotating, intermeshing, ICRR

• Self cleaning • High dispersive mixing in the shear and extension • It can accommodate multi-lobal mixing elements of the gear type with excellent dispersive mixing capabilities • It has better specific energy than CORI • It has narrower residence time distribution than CORI • High local pressures possible • High local vacuum possible • Better devolatilization capability than CORI • High specific energy input possible • Shorter screws than in co-rotating • Easy separation of desired functions, viz. mixing, devolatilization, chemical reaction, compounding, etc.

• Transverse mixing is moderate • Longitudinal mixing is poor • Only residence times below 10 min are economic • Limited screw speed • Capital cost

617

Compounding Polymer Blends

Table 9.15. Comments on CORI vs. ICRR performance CORI

ICRR

Reference

Better control of shear and of mixing than in ICRR. Broader distribution of residence times, slower melting and cleaning, but better for dispersing glass fibers. For the same screw

Less efficient control of shear and of mixing than in CORI. Narrower distribution of residence times, better dispersing agglomerates of small particles, faster cleaning time (of TiO2 trace). Owing to periodic

Murray, 1978

speed and throughput higher melt temperature was obtained (an indication of not-optimized screw configuration). Material transfer from one screw to another, axially open extruders. Melt is conveyed by frictional forces, allowing for partially filled melting zone. Screws suspended in melt allow for high rotational speeds. In axially open system the pressure is controlled by purposely placed restrictions: either left hand side screw, valves or die. Material is compressed by the first screw and conveyed to the other. The melting process is similar to ICRR, but particle to particle rubbing makes it more efficient. Broader distribution of residence times, better distributive mixing, better for blending. Higher screw speed and output than the ICRR. Co-rotating extruders: are especially suitable for reactive processing. Polyethylene/polyamide-11 blend was compatibilized by radical copolymerization using diverse organic peroxides. At N = 400 rpm comparative blending produced lower throughput and poorer quality. Produced the same maximum lubricant concentration at different output. More aggregates than in ICRR.

Sakai, 1978; Sakai et al., 1988, 1992

compression-depression effects better devolatilization can be obtained. There is no transfer of material from one screw Herrmann and Burkhardt, to another, closed C-shape chambers are 1981 developed. Melt flow is plug type. Materials see a wide range of pressures. Pressure within the calendar gap forces the screws Herrmann and Burkhardt, to rub on the barrel. Low screw speeds must 1981 be used. Calendering gap problems are exacerbated in the feed region, leading to severe wear. The melting process is similar to that in single screw extruders. Unmelted particles are observed in the discharge.

Better pumping capability, dispersion of small Rauwendaal, 1981 particles, and devolatilization. Calendering leakage results in high lubrication White, 1990 pressure which in turn forces the screws apart, causing barrel abrasion. High calendering pressure causes mechanical degradation of polymers. The best performing material was obtained using Lambla and Seadan, 1993 co-rotating fully intermeshing machine. At N = 400 rpm comparative blending produced higher throughput at lower T and better quality. Higher maximum lubricant concentration achieved at higher power consumption.

However, optimization of screw configuration and operating conditions remains an art. TSE’s are indispensable in several applications, viz. reactive processing, high devolatilization, high additive loading, and/or intensive dispersive

Thiele, 1995

melt-mixing. Furthermore, they are the preferred extruders for high throughput lines with the extrusion rates exceeding 10 tons/h. Table 9.14 summarizes the relative advantages and disadvantages of the intermeshing co- and

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counter-rotating TSE (respectively, CORI and ICRR). A summary of published comments about the advantages and disadvantages of ICRR in comparison to CORI is provided in Table 9.15. Comparison between SSE and TSE performance Within TSE the shear and extensional stresses are generated by virtue of interactions between two screws. The degree to which this occurs depends on the relative direction of the screw rotation (co- or counter-rotating), shapes of the screw elements, the degree to which they intermesh, and the rotational speed of the screws [Fisher, 1958; Matthews, 1962]. The conventional TSE’s are claimed to be superior to SSE as compounders because they can provide [Murray, 1978; Wood, 1979, 1980; Adams, 1974; White, 1990]:

(1) better feeding and more positive conveying characteristics, (2) shorter residence times and narrower residence time distribution, (3) improved kinetics and melt temperature control, (4) high and controlled deformational stresses, (5) positive pumping action, (6) reduced melt slippage, (7) self-wiping action, (8) lower power consumption, (9) generation of high extrusion pressures with a short backup length. Table 9.16 summarizes these differences. The extrusion characteristics are given in Table 9.17. The TSE’s have also drawbacks, the main being their cost. Owing to complex design of the screws and the thrust bearings these machines are

Table 9.16. Fundamental Differences Between the Single (SSE) and Twin Screw Extruders (TSE) Function

SSE

TSE

Degree of channel fill Throughput, Q, is determined by: Total strain Shear strain Material transport Material flow path Heat transport Liquid and powder additives High viscosity additives Mixing capability Extruder length for melting Extruder length for mixing Advantages

Fully filled Screw speed, N Independent of Q and N Low By friction (drag flow) Smooth, regular thinning Ineffective Present problems Impossible to disperse Poor Long Long Easy to make, inexpensive, theoretical description of its function well developed, lot of experience Lack of flexibility, poor feeding characteristics, ineffective devolatilizer and mixer, not suitable for reactive processing

Partially filled Feed rate Dependent on Q High, controllable Positive conveying Tortuous, diverse Effective Do not present problems Can be dispersed Good, adjustable Very short Short High flexibility (modular design), effective feeding, devolatilization, mixing, reactive processing Difficult fabrication, cost up to 10x that of SSE, early stage of the theoretical analysis, poor back mixing (narrow residence time distribution), require accurate feeders

Disadvantages

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Compounding Polymer Blends

more expensive than SSE [Prause, 1967, 1968; Gras, 1972]. Furthermore, they are susceptible to quick overloading. The low bulk densities of preblends influence the throughput rate and the die pressure. Mechanically, a major difference between SSE and TSE is the type of transport that takes place within the extruder. The material transport in the former is drag-induced — the frictional drag in the solid conveying zone and viscous drag in the melt conveying zone. There are many materials with unfavorable frictional properties that cannot be fed into a SSE without getting into severe feeding problems. On the other hand, the transport in an intermeshing TSE is by positive displacement. Its degree depends on how well the flight of one screw closes the opposing channel of the other screw. The most positive displacement is obtained in a closely intermeshing, counterrotating geometry. Another difference is the velocity profiles in these machines. In SSE these are well defined and

fairly easy to describe. The situation in TSE is considerably more complex. The complex flow patterns have several advantages, such as good mixing and heat transfer, large melting capacity, good devolatilization capacity and control over the stock temperatures. The theory of TSE is not nearly as well developed as that of SSE. The extruders are assumed to be nominally “once through equipment”. This implies that preblending, or bulk mixing, is essential for good mixing. In SSE that operates under the flooded conditions, feeding is often difficult and preblending may be necessary. In counter-rotating tangential TSE the degree of fill is low and feeding is relatively easy (provided that feeding devices have sufficient accuracy). Thus, uniform product quality is obtained by feeding separate streams either at the same or sequentially into various axial locations. Most TSE’s are of a modular design, having removable screw and barrel elements. The screw design can be altered by changing the sequence

Table 9.17. Summary of extrusion characteristics for SSE and TSE Function

Single Screw

Flow mechanism

Continuous shear

Stress field

Shear in the overflight and in the channel

Conveying Pumping efficiency Die pressure Die restriction L/D ratios Compression Air entrapment Screw speeds Heating mode Residence time

Drag flow Variable High Often severe >20 Decrease channel Possible 20-100 rpm High proportion by shear Large spread wide distribution Low

Pressure in the gap

Twin Screw Counter-rotating

Co-rotating

Discrete c-sections, lengthand cross-wise closed High shear and elongation in the calendering gap, low in channels Positive conveying Good, positive Low Smaller effects 16-20 Various designs Possible, thus larger clearances 500 rpm Controllable low shear Narrow distribution, often easy to control High, thus possible wear

Figure-of-eight; more uniform shear history for all flow elements Uniform stress field, mainly shear; convergent-divergent flow in mixing blocks Less positive conveying Good Lowest Smaller effects Various Various designs No Up to 1500 rpm Near-adiabatic Wider distribution at high die pressures Low, effective self-wiping

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and type of the screw elements along the shaft. In this way, almost an infinite number of screw configurations can be put together. The modular design, therefore, creates great flexibility and allows careful optimization of screw and barrel geometry for each particular application. This flexibility is not available for SSE. The modular design of TSE makes it possible to adjust the relative magnitude of the distributive and dispersive mixing. The shear stress in normal melt pumping SSE channels is not sufficient to provide adequate dispersive mixing. Similarly, the laminar flow through SSE channel is not sufficient to distribute the flow elements having different performance characteristics. Often, high shear stress mixing devices, such as modified screw, screw torpedoes or external high energy planetary roller mixers are required. For the distributive mixing a static mixer may be inserted between the end of the extruder screw and the extrusion die. The modern technology often requires that TSE performs a series of different functions. Owing to limitation of L/D, even these universal machines may profit from additional, external, either dispersive or distributive mixers. These devices are discussed in the following parts. There are twelve criteria for the economic, high-quality compounding: 1) sound fundamental design principles of the compounding machine; 2) flexibility of operation; 3) high screw torque; 4) appropriate selection of the screw elements; 5) a balanced combination of the high-shear dispersive and low-shear distributive screw elements; 6) low pressure at the die; 7) reliability of components and parts; 8) long lifetime of parts; 9) reliable feeding/dosing equipment; 10) reliable granulation method; 11) automation of the process; 12) on-line quality inspection and closed-loop control [Hermann, 1988]. Flexibility and user friendliness are the key words for many of today’s extruders (easy-to-use control systems and displays, quick access for repair, and quick changeover). Several of the new machines feature: segmented barrel, the ability to change screw-rotation direction, and the ability to be modified in place, e.g., to compound, to extrude a product or to recycle plant scrap.

3.

Other mixing devices

1. Planetary roller extruder In these extruders, six or more evenly spaced planetary screws revolve around the circumference of the central or so-called “sun” screw. The planetary screws intermesh with the sun screw and the barrel. The planetary barrel section has helical grooves corresponding to the helical flights on the planetary screws. This section is usually a separate barrel with a flange-type connection to the feed barrel section. In the first part of the machine, before the planetary screws, the material is processed as in a SSE. As the plasticated composition reaches the planetary section it is exposed to intensive mixing by the rolling action between the planetary screws, the sun screw and the barrel. The helical design of the barrel, sun screw and planetary screws lead to a large surface area relative to the barrel length. The clearance between the planetary screws and the mating surfaces, is about 1/4 mm. This allows thin layers of the compound to be exposed to large surface areas, resulting in effective devolatilization, heat exchange (thus good temperature control) and efficient mixing. The planetary roller extruders are mostly used for processing heat-sensitive compounds, viz. rigid or plasticized PVC formulations [Collins, 1987; Anders, 1979]. The planetary roller section can also be used as an add-on mixer to SSE to improve the mixing performance [Rust, 1983; Huszman, 1983]. Degassing and addition of fillers or reinforcements into melt has also been carried out. 2. Four-screw extruder These machines are used primarily for removal of solvent [Gras, 1972; Gras and Eise, 1975]. Flash devolatilization occurs in a flash dome attached to a barrel. The polymer solution is delivered under pressure and at temperature above the boiling point of the solvent. The solution is then expanded through a nozzle into the flash dome. The foamy material resulting from the flash devolatilization is then transported away by the four screws. In many cases, downstream vent sections are also used to further reduce solvent content.

Compounding Polymer Blends

3. Disk extruders There are several extruders that do not use an Archimedean screw for transportation of materials, but still belong to the category of continuous extruders. These machines are sometimes referred to as screw-less extruders. Usually they employ a disk or a drum. Most disk extruders are based on viscous drag transport principle. To this category belong: stepped disk extruders [Westover, 1962; Raleigh, 1879], drum extruders, spiral extruders [Ingen Housz, 1975], diskpack extruder [Tadmor, 1979, 1980; Tadmor et al., 1979, 1983; Hold et al., 1979; Valsamis, 1983], and many others. Dispack [Tadmor and Gogos, 1979] has the inherent capability of performing the elementary steps of plastics processing by combination of differently shaped rotating disks in a drum-like housing. The wiping action is provided by stationary channel blocks that cause the material to transfer from one disk-gap to another. Melting, laminar mixing, venting, and the pumping functions are all separated. Diskpack has been used for reactive processing, blending, compounding, mixing, and devolatilizing [Tadmor et al., 1983]. The devolatilization capabilities are limited in comparison to TSE. The mixing capability and die pressure are influenced by: 1. Method of feeding 2. Disk diameter 3. Gap between disks 4. Shape and size of the feeding zone 5. Diameter and length of discharge 6. Hydraulic resistance of forming device 7. Viscoelastic properties of the melt 8. Shape of material particles 4. Elastic melt extruder (EME) In EME the polymer is sheared between two plates, one stationary and one rotating. The extruder makes use of the viscoelastic properties of polymer melts. When a viscoelastic fluid is sheared, the normal stresses develop in the fluid, trying to push the shearing plates apart. Thus, leaving a hole in the center of the stationary plate makes it possible for the melt to flow continuously from the rim toward the center then out [Maxwell and Scalora, 1959; Blyler, 1966; Fritz, 1968, 1971;

621

Macosko and Starita, 1971; Kocherov et al., 1973; Good et al., 1974]. The EME mixers have been used for mixing and extruding plastics formulations. Despite the relatively low pressures (0.3 to 0.7 MPa) in many applications (e.g., melt blending and incorporation of carbon fibers) EME was found superior to the screw-type extruder. The machines provide good mixing and melt homogenization at short residence time (e.g., for a unit with 150 mm disk diameter and a capacity of 25 kg/hr, the average residence time is about 15 seconds). EME is two to three times smaller than screw-type extruders of similar capacities. The flow is laminar dispersive, preferentially generating co-continuous blend morphology that lead to good performance [Thornton et al., 1980]. Unfortunately, neither the original EME nor its modifications found a general acceptance in the plastics industry [Westover, 1962; Frederix, 1978; Michaux, 1979]. 5. Ram extruders The ram extruders are divided into: (1) single ram extruders that operate discontinuously [Berzen and Braun, 1979; Zachariades et al., 1979; Sperati, 1983], and (2) multi-ram extruders that offer a continuous flow of materials [Westover, 1963; Yi and Fenner, 1975]. 6. FN-Plastifier This is a short (L/D = 5) single-screw extruder, developed for polymer blending and recycling, with a three-start screw extended from the feed zone 2/3 over the screw length. The frontal part of the screw is smooth ending with a flat disk. The material is transported and partially fluxed by the grooved part of the screw and barrel. The pumping is assured by the normal stresses between the flat part of the screw-end and the die (Maxwell-Scalora’s EME principle). The short screw assures short (and narrow distribution) residence time. Large thrust bearing makes it possible to control the die gap, thus the magnitude of the normal stresses, hence morphology and performance [Frederix, 1978; Anonymous, 1981, Michaux, 1979].

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The extruder was developed by Patfoort, then licensed to Fabrique Nationale Herstal, S. A., (FN) [Patfoort, 1976; Fabrique Nationale Herstal, 1977]. The machines were found to be particularly suitable for recycling commingled polymers from the municipal waste streams. Under high stresses developed between the smooth part of the screw and the barrel, the free radicals were generated compatibilizing in-situ the polymeric mixture. The usual product was of the “plastic wood” type. It had excellent mechanical performance characteristics. The FN plastificator was used for recycling by the city of Liege and other parts of Europe, then exported to China under the UNIDO program. However, FN discontinued manufacturing these machines. Several other machines of this type have been developed since. They all involve intensive mechanical shearing that produces extensive chain scission. Recombination of the free radicals in situ generates sufficient concentration of copolymer, to compatibilize the system. The generated under high stress, non-equilibrium morphology is then locked by quenching. Best performance has been observed for systems with co-continuous morphology. 7. Gelimat, K-mixer and Homomicronizer The residence time in the discussed above Patfoort’s extruder is 7 < τ < 25 s [Patfoort, 1976]. The machine was originally designed for the extrusion of the thermally unstable resins, in particular PVC. For similar applications Carlew Chemicals developed its Gelimat, latter re-named K-mixer (K- for the kinetic energy) [Crocker, 1981]. The mixer uses high-speed rotor with staggered blades mounted on a horizontal shaft at different angles. The number and position of the blades vary with the mixer size. The mixing is carried out at the blade tip velocity of 30-45 m/s. There is no external heating — the kinetic energy generated by the particles impacting on each other and the mixer elements produces sufficient heat to flux the material within 8 to 150 s. Once the material reaches the desired temperature, the infrared sensor activates the bottom doors of the chamber. The discharged dough can either be fed

to a short screw extruder or passes between rolls and diced. The mixing is batch-type. However, with two Gelimat units each providing a load of plastified resin every 10 s or so, the down-stream operation (e.g., calendering) is continuous. When preparing polymer blends, the starting resins should have similar particle size and dimensions. There is an obvious advantage of using fine powders. The high stresses developed during the plastification are often sufficient to generate compatibilizing copolymers. Similarly like in the Patfoort extruder, here also good performance of antagonistically immiscible polymer blends was obtained. More recently, Newplast developed recycling process line that is based on a Homomicronizer — the machine that very much resembles Gelimat or K-mixer. The unit has short, stubby radial blades attached to a horizontal shaft that rotates at high speeds, melting and homogenizing the post-consumer waste within 35-120 sec. The blades are of two types, pushing the material away from the side walls toward the center of the cylindrical mixing chamber. The clearance between the blades and the chamber wall is 0.5 to 1.0 mm [Dubrulle D’Orhcel, 1996; La Mantia et al., 1996]. 8. Mixing or calendering rolls The mixing rolls subject pastes and deformable solids to intense shear by passing them between smooth or corrugated metal rolls that revolve at different speeds. The principal design consists of two horizontal rolls or cylinders, arranged side by side and rotating toward each other at different speeds. The ratio of the peripheral roll speeds (known as the function ratio) ranges from 1 to 2, but it is usually about 1.2. The rolls’ temperature is controlled by the circulating oil. The higher friction ratio leads to greater heat generation. Friction, speed, sizes of the rolls and gap between them influence the material temperature and the intensity of mixing. Usually, the results obtained from one unit are specific to this machine and cannot be applied to a different one. The material enters the mixing rolls in a form of lumps, powder or friable laminates. As a result

Compounding Polymer Blends

of rotation, adhesion, and friction, the material is entrained into the gap between the rolls, and upon discharge it sticks to one of them. In batch mixing the mass after loading passes through the gap between the rolls several times. In continuous mixing the mass enters continuously from one side of the machine and passes between the rolls in rotational and forward motion along the unit’s axis. The mixed material is continuously discharged in form of a narrow strip. Both the shearing action and entertainment of material into the gap are important to the mixing process and transporting the material through the unit. 9. Ribbon blenders and planetary mixers There are a number of design variations of this type of machines, with differences primarily in the agitator configuration. Ribbon blenders are essentially self-contained mixers best suited for batch or semi-continuous mixing. They are mainly used for preblending, e.g., dry mixing a composition. The continuous ribbon agitator is standard, with inner and outer-ribbons that may be arranged for center or end discharge. Planetary mixers are well suited for a wide range of liquid and solid applications. Double planetary mixers consist of two regularly shaped stirrer blades revolving around the tank on a central axis. Each blade also revolves on its own axis at approximately the speed of the central rotation. With each revolution on its own axis the blade advances along the tank wall. This movement ensures good homogeneity of the material within short time. Double planetary mixers have no packing glands or bearings within the product zone, thus cleaning between batches is relatively simple.

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mixing accompanied by either heating or cooling. The operation involves compressing the fluid mass, folding it over, and then compressing it again. The material is usually torn apart by high shear stresses engendered between the moving and stationary elements. The mixing is performed using two Z-shaped by blades rotating in opposite directions on parallel horizontal shafts. There is great variety of the commercially available blade designs, ranging from lightweight to heavyweight constructions. Selection of the specific design depends on the consistency of the mix. The size of the mixer is limited by the power input, batch weight, speed of mixing, materials of construction, and methods of mix discharge. Many machines are equipped with tilting facilities to discharge the batch while the blades are turning. The mixers require heavy drive mechanisms and large motors. The agitator blades in the sigma-blade mixer may be mounted so that their paths are either tangential or overlapping (intermeshing). In the tangential arrangement, the two blades rotate side by side with their circular paths of rotation not quite touching. The blades can have any relative speed because their paths of rotation do not overlap. In the overlapping arrangement, the paths of rotation overlap. Consequently here the blades must be designed and the speed of rotation adjusted so that the blades always clear each other. Many manufacturers provide double-arm mixers with interchangeable blades. Examples of materials handled by double-arm kneaders of the sigmablade type are resins, putty, adhesives, baker’s dough, and cellulose additives. Sigma-type agitator blades are best for general-purpose use. 12. Dispersion-type screw mixers

10. Emulsifiers and blenders Mixer-emulsifiers are sometimes used as an alternative to slow-speed impeller mixing or highpressure homogenization for a wide range of processing requirements. These mixers are normally used in dished or conical-bottom vessels. 11. Internal (blade) mixers The internal mixers are used for kneading and

Dispersion-type screw mixers are in an entirely separate class from the extruder designs described above. They are applicable to processing of either dry materials or moderately viscous pastes, creams, or lotions. These machines normally consist of a conical vessel equipped with a conical or inclined screw. There are either single- or twin-screw models that provide a gentle mixing action, and thus

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are used for handling materials sensitive to attrition or fiber disintegration (one example is in dry cell battery manufacturing, where acetylene black has lace-like structure that easily breaks down). The twin-screw design, in addition to the axial, orbital, and gravitational action offered by the single-screw design, provides intermixing between the two screws. Depending on the material properties, the mixing times can be 90% faster than the single-screw system. The Combimixer employs two intersecting conical tanks with two orbiting spiral agitators. The axial and orbiting motions of the two agitators, along with the additional intermixing currants occurring through the juncture of the two tanks, provide for greater blending action than the single screw mixer. Two motors and two gear boxes drive the smaller Combimixer. One set drives both screw flights axially and the other drives the two orbiting arms. For larger Combimixer units, separate motors are used for each screw flight and a single motor drives both orbiting arms. 4. Add-ons A need for better mixing in a SSE was already evident in the 1950’s. Large effort was spent to re-design the screw by either changing the flow profile (mixing screws, viz. Eagle, Chief, Meltstar, ODM, Pulsar, Strata-blend, Toss-andRock, Stat-Dyn, etc.), adding mixing elements (pins, rectangles, or trapezoids), or modifying the screw tip. In the latter category quite a long series of designs has been proposed, e.g., Axon, Dulmage, Egan, Maddox, pineapple, turbine, etc. Today, these devices are parts of the extruder technology and have been well described [Manas-Zloczower and Tadmor, 1994; Rauwendaal, 1994]. By definition, an add-on device is an optional part to be attached to a standard SSE. While some of these use the extruder torque, the others can be used either as an attachment or stand-alone units with their own power (e.g., CTM or DMX mixers). Barmag 3DD torpedo mixer is a forerunner of this category of mixers. The unit was designed to improve the temperature homogeneity of the film and fiber spinning extrusion lines. The device is

made of a torpedo (to be attached to the tip of the SSE screw) and a sleeve (to be attached to the extruder barrel). Both the torpedo and the sleeve have slot-shaped grooves, arranged in rows (usually 12). Significant improvement of the melt homogeneity was reported by the manufacturer. Cavity transfer mixer (or CTM) was designed at RAPRA as an add-on distributive mixer [Gale et al. 1982]. The device is made of a torpedo (to be attached to the tip of the SSE screw) and a sleeve (to be attached to the extruder barrel). Both the torpedo and the sleeve have semi-spherical indentation. The number of rows of indentations must be selected to satisfy the mixing needs of the production line. The polymer flowing between them is sheared and split into different indentations. Furthermore, the extensional flows generated at the entrance and exit from the indentations introduce additional dispersive action. The flow through a simplified CTM was recently analyzed experimentally and by means of a fluid dynamics analysis package — FIDAP [Manas-Zloczower and Tadmor, 1994; Wang and Manas-Zloczower, 1996]. More recent is the University of Twente mixing ring (or TMR). This device is made off three parts: smooth wall barrel extension, a torpedo with semispherical indentations (similar to these in CTM) and a perforated sleeve that goes between them. The sleeve provides shallow indentations. It slowly rotates between the torpedo and the barrel by the virtue of a drag flow. The flow in TMR is both dispersive and distributive. The advantage of TMR is its suitability to improve mixing in injection or blow molding machines [Housz, 1989]. The newest in this series of add-on mixing torpedoes is the Dynamic Melt Mixer, DMX. The device consists of a series of mixing modules inside a tubular housing. Two types of modules are available. Each module is made of two components: a stationary ring fixed to the sleeve, and veined rotor keyed to an axial shaft either attached to the SSE screw or independently powered. Part of the rotor that does not enter the sleeve, provides the cutting action across the flow lines that induces the distributive action. The primarily distributive one consist of a stator and a rotor

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Compounding Polymer Blends

(keyed to the central shaft) without the rolling semi-cylindrical lugs. As the rotating modules turn, they provide dispersive and distributive mixing [Miller, 1996; Petren, 1997]. An interesting distributive add-on mixer was proposed by Jurkowski et al. [1992]. The mixer is made off a short barrel attached to the front of SSE. Within the barrel there are several plates with channels directing the melt to flow from the rim toward the center and vice-versa. Other mixers, viz. the “static (or motionless) mixers” were discussed in Part 9.1.5 “Distributive mixing” while the Extensional Flow Mixer (EFM) and the dynamic extentional flow Mixer (DEFM) were described in Part 9.1.6.4 “Extensional Flow Mixer”.

9.3

Non-mechanical Methods of Polymer Blending

9.3.1

Latex Blending

Historically, the preferred method of mixing was the latex blending. The early emulsion polymerization of rubbers (natural lattices or synthetic ones) and thermoplastic acrylates, methacrylates and vinyls provided raw ingredients for the latex blending. A wide range of blends was described in the early I. G. Farbenindustrie patents. Later on, owing to the development of loop technology [Lanthier, 1966] the method was revived by the Shawinigan Chemicals. The latex blends were used either directly, for example, as paints, adhesives or sealants, or they were spray-dried or pelletized. Spray-drying has been the most frequently used method. Since emulsion polymerization was able to produce either a fine dispersion (drop diameter d = 100 to 200 nm) of homo- or co-polymer, or alternatively large drop aggregates, comparable to those generated in suspension polymerization (d < 15 µm), the latex blending offered not only a wide range of compositions but also diverse morphologies. The principal disadvantage of the latex blending was high content of contaminants: emulsifiers, residuals of the initiators, chain transfers, stabilizers, etc.

9.3.2

Solvent and Spin Casting

There are several methods for the production of polymeric films (see Table 9.18). The oldest is calendering, but the most common methods are the film blowing and sheet extrusion. Others can be made by post-treating the extruded sheet, e.g., polyimide films [Sweeting, 1971]. Polymer films can also be polymerized onto a substrate from monomers in the vapor state, or formed by the coalescence of polymer dispersions. However, several polymers cannot be formed without dissolving them. Solvent casting is an important commercial technique utilized to fabricate thin layered films for diverse applications [Sweeting, 1971]. Most familiar is the solution casting of cellulose acetate (CA) for photographic films having good dimensional stability, clarity, flexibility and fracture resistance. The following advantages of solutioncasting method (as compared to the melt process) have been given [Ricklin, 1983]: • higher quality (uniformity) and thinner film • freedom from pinholes and gel marks • purity and clarity • lack of residual stresses • possible to produce patterns or dull finishes The solution-cast process consists of dissolution of the film ingredients in a suitable carrier that conveys the solution through a drier where the solvent is evaporated. The resulting film is removed from the substrate and wound into rolls. Solvents and scrap materials are reprocessed. As illustrated by examples in Table 9.19, the drying process can be used to manipulate the final film properties. Commercial solution-casting is continuous (see Figure 9.21) [Roberts, 1967; Sittig, 1967; Park, 1969; Chow et al., 1976; Heffelfiger, 1978; Ricklin, 1983, Powers and Collier, 1990]. The mathematical modeling of the process is complicated by the diffusivity dependence on composition. Low boiling solvents exhibit high diffusivity and usually are preferred, but for polymers with high Tg, high-boiling solvents may be more appropriate. The drying step is rate-

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Table 9.18. Methods for production of polymeric films Process

Method

Film blowing or sheet extrusion PEST, PA, etc.) Calendering

Continuous melt extrusion through a flat or circular die (PE, ionomers, PP, PS, PMMA,

Batch PVC Plastisol or Organosol casting Continuous solvent casting Batch in situ polymerization Continuous in situ polymerization

Casting onto a substrate

Usually it follows the sheet extrusion. However, there are commercial lines where plastified polymer (e.g., semi-rigid PVC formulations) are directly fed to the calendar mil Plastisol and organosol casting process involves melting the polymer or polymeric mixture by an action of the temperature and presence of either a plasticizer or solvent Extrusion of a solution through a flat die (CA, cellulose xanthate, PMMA, PVF, blends) Monomer casting, followed by controlled (co)polymerization (e.g., PMMA, epoxies) Extrusion of Polymer-1 containing monomer of Polymer-2, followed by post reaction to Polymer 2 (polyimides, electrically conductive polymer blends, reactive blends, e.g., with crosslinked elastomers), etc. Electro-deposition, Plasma-deposition, or Dispersive deposition (e.g., vapor deposition of polyxylylene)

limiting. An early mathematical model predicts the temperature and concentration profiles in the film during the drying [Collier, 1981, Roehner, 1982; Roehner and Collier, 1983]. As indicated in Table 9.19, casting may induce orientation of the polymeric chains. The process is caused by the polymer movement induced by the solvent removal. Solvent removal from a solution constrained in two dimensions by a substrate results in one-dimensional thickness reduction. The orientation is caused by effectively stretching the polymer in the planar directions during thickness reduction. The resulting orientation depends on the evaporation rate. The competing rates for molecular relaxation and solvent removal are best considered using the Deborah number, defined as: Db = τ/t1/2

(9.31)

where τ is the polymer relaxation time and t1/2 is the time required for the removal of one half of the solvent. For constrained processes with Db >> 1 orientation occurs; for Db << 1 solvent removal results in only thickness reduction [Collier, 1981]. Birefringence measurements confirmed that dry-

ing conditions can affect the level of orientation in the films, but the techniques could not detect an orientation gradient developed across the thickness [Prest and Luca, 1980]. Solvent casting of polymer mixtures (usually with additives) may affect the density, porosity, mechanical properties, mode of failure, etc. These effects are particularly pronounced in films of immiscible polymers. Before discussing these effects, it is appropriate to consider the intermediate step — solvent casting of copolymers; Table 9.20 provides few examples. For multicomponent systems, solution casting provides means for generating the desired degree and type of morphology. Profiting from the different (and controllable) rates of the phase separation and solvent evaporation one may generate either co-continuity of phases (spinodal decomposition, SD) or dispersed structure (nucleation and growth or late stage of SD). The solvent casting offers additional flexibility in producing precisely what is needed. It is not necessary to use a single component solution — one frequently uses a mixture of solvents having different evaporation rates and different miscibil-

Compounding Polymer Blends

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Table 9.19. Examples of solution casting of polymers Polymer

Solvent

Remarks, Methods & References

Poly(g-methylD-glutamate)

Chloroform, DCA, MC, FA, TCE

Poly-2,5(6) benzimidazole

Methane sulfonic acid, HCOOH/cresol

Solvent controls the degree of a, ß-configuration, or random coil contents in the film, its mechanical behavior and morphology. Rheovibron, Stress-strain measurements, WAXS, SAXS, Polarizing photomicrography, IR [Mohadger and Wilkes, 1976] Solvent nature affects the morphologies of polymer films thus the extensibility and mechanical properties of films. DSC, Rheovibron viscoelastometer, WAXS [Wereta et al., 1978]

PS, PC, PVK

CH3Cl Benzene, THF,

Polyvinyl acetate

Methanol, Acetone, Chloroform

PE, PC, PET, PS, i-PS, PVP, PMMA, PTFE, PVAc, PVC

Toluene, CCl4, CHCl3, Xylene, Decalin, EtOH, H2O, MeOH

PVC

MeOH, CHX, Toluene, THF, Decane, CS2

Ether polysulfone, PVDF

A high water affinity solvent

Figure 9.21 The solution-casting process.

Casting aligns polymeric chains preferentially in the plane direction. DSC, Optical anisotropy [Prest and Luca, 1980] Mechanical properties of the films depended on the solvent and the time of exposure of film at each temperature; CHCl3 was recommended. Mechanical tests [Olayemi and Adeyeye, 1982] The casting solvent has a pronounced influence upon the friction coefficient — it may confer either ductile or brittle failure. “Good” solvent generally tends to promote brittle mode of failure with little temperature dependence. Friction measurements [Briscoe and Smith, 1983] The slowest evaporating solvent has the greatest effect on the surface properties; annealing above Tg reduces the surface porosity. Inverse Gas Chromatography, DSC [Wildman and Hsu, 1984] Swelling and the viscosity of the casting solution were controlled by addition of small amounts of partially miscible polymers. SEM, Contact angle measurements [Stengaard, 1988]

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Table 9.20. Examples of solution casting of copolymers Material

Solvent

SBR

CHX, Toluene, Et-acetate, Tetralin, Cyclohexanone

SEBS, and SEPS

n-Heptane, Cyclohexane, Toluene

SEBS

THF, CHCl3, Toluene, CHX, Bromobutane, C6H14, C7H16

SBS block Toluene, CCl4, EtAc, MEK copolymer Silicone-PC C6H14, methylene chloride block copolymers

Poly(styrene-bbutadiene-b4-vinylpyridine)

THF, Benzene, MEK, EtOH, MMA, Nitromethane, Hydroxyethyl methacrylate

PDMS-b-PS copolymers Ter-copolymer of acrylonitrile methylmethacrylate

CH3Cl, Toluene, Cyclohexane Glycerol, Water

Remarks, Methods & References Film properties depended on the evaporation rate, but there was no evaporation rate dependence when the casting solvent was CHX. Density measurements, Rheovibron dynamic viscoelastometer, TMA, SEM [Beamish and Hourston, 1976; Beamish et al. 1977] Spherical, cylindrical and lamellar PS structures were obtained in the order of increasing solvent affinity to PS. Density, DSC, tensile test, dynamic viscoelastometer, SAXS [Séguéla and Prud’homme, 1978]. The complex modulus varied from 4.4 MPa (cast from low δ -solvents) to 205 MPa (cast from high d -solvents); the area under the damping peak also depended on solvent. DMA, DSC [Cowie and McEwen, 1979, 1980; Cowie et al., 1979] The solvent nature influenced film morphology. Position annihilation technique [Djermouni and Ache, 1980] Casting from a homogeneous solution that favors the hard segments yielded better results than from mixed solvents from which PC blocks preferentially separated to form a discontinuous dispersion or network. Capillary viscometer, GPC, TMA, DMA [Maung et al., 1983] A ball-in-a-box structure morphology was formed when the solvent for poly(4-vinylpyridine) is lost; a lamellar structure morphology was formed when solvents were equally distributed into the three phases; a random structure was formed when the solvent for one or two of the three block segments vaporizes faster than other solvent. SEM, Mechanical test [Arai et al., 1984] The interactions between the casting solvent and polymer affected the phase separation and the final morphology. DSC [Feng et al., 1988] An increase of glycol in the casting solution resulted in the formation of membranes with high pore density and small pore size, thus increasing the selectivity with respect to low molecular weight components. Selectivity of membrane, SEM, Porosimetry [Petrov et al., 1991]

ity with the polymeric components. Furthermore, the films can be annealed, selected polymeric components extracted to form the desired size and shape of domains. For example, it is possible to generate semipermeable membranes with well-controlled morphology, pore size, flux, and mechanical performance. The method is preferred for industrial preparation of semi-permeable membranes from polymer blends (see Table 9.21).

Another method of morphology control in multicomponent polymeric systems is by the use of mixed solvents having different affinity to the polymeric components. When the non-solvent is less volatile than the good solvent, evaporation results in a two-phase structure. When solvents are removed solely by evaporation, the membrane formation is known as a dry phase inversion process [Kesting, 1985]. When the phase separation and structure forma-

Compounding Polymer Blends

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Table 9.21. Solution casting of polymer blends Material

Casting Solvent

Remarks, Methods and References

Grafted block copolymer blends

Divers

The soluble polymer usually was the matrix with the other resins forming isolated domains. Physical properties and morphology [Merret, 1957; Beecher et al., 1969; Inoue et al., 1960; Pedemonte et al., 1975; Kraus and Rollman, 1976; Kawai et al., 1971]

PMMA/PS/SA

n-Butanol, Toluene, CHCl3,

Good solvent produced a tightly packed films, while poorer solvent yielded more porous ones. IGC, Surface tension [Schreiber and Croucher, 1980] The Tg of polymers varied because of residual solvent, possible due to the plasticizing effect, and thermal history. DSC [Brostrom et al., 1980] Blend properties were found to be a combination of these of the constituent homopolymers. Segmental motions of the PS chains in the copolymer were enhanced even though those chains were attached to more rigid PC chains. SEM, Rheovibron viscoelastometer, Instron [Eastmond and Haraguchi, 1983] Regions of miscibility and immiscibility as functions of a casting solvent and molecular weight for both components were established based on Tgs consideration. DSC [Saeki et al., 1983] Faster rates of solvent evaporation yielded a smaller periodic distance in the modulated structure. The rate of phase separation was lower and the periodic distance of the modulated structure smaller as the blend ratio deviated from 50/50. For high MW polymers, the system failed to phase-separate during fast casting, resulting in a homogeneous blend without appreciable composition fluctuations. Light-scattering [Inoue et al., 1985] The spherulitic crystalline morphology of PVDF showed strong dependence on the casting temperature and compositions. Mixed a- and g-form crystals were obtained. DSC, SAXS, Microscope, Dielectric thermal analyzer [Kyu and Yang, 1990] Casting induced distinct molecular order in the solid films; birefringence was influenced by the solvent type, the casting temperature, the surface energy of the substance, and the timing of the peel-off step; non-birefringence films could be produced by setting proper operating conditions. Birefringence [Machell et al., 1990]

terpolymer blends

Ethylene dichloride, Ethanol

Poly(alkyl methacrylates), PS/BR

Methanol, Toluene, DMF

PC, PS, PC/PS blends, PC-PS copolymer

Dichloromethane, Acetone

PS/poly(α-methyl styrene)

Toluene, Cyclohexane, Propylene glycol

PVC/NBR or EVAc; CR/NBR or EVAc; BR/SBR or IR; ABS/NBR or PMMA; PS/PMMA

THF

Perfluorinated Nafion ionomer/PVDF blends

Dimethylformamide, Dimethylacetamide, 1-Methyl-2-pyrrolidinone

PC, PS, PPE, t-methyl PC

Dichloromethane

tion are achieved by immersion of a cast membrane in a quench medium, the process is known as a wet phase inversion process [Heffelfinger, 1978]. The latter process is used to prepare asymmetric membranes for either: microfiltra-

tion [Roesink, 1989], ultrafiltration [Michaels, 1971], reverse osmosis [Strathmann et al., 1971], or gas separation [Henis and Tripodi, 1980; Van’t Hoff, 1988]. Post-treatment may be necessary to prepare ultra-thin asymmetric membranes for gas

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separation [Henis and Tripodi, 1980; Zampini, 1984; Zampini and Malon, 1987; Murphy et al., 1989; Ekiner et al., 1989]. Membrane structures can also be formed by a microphase separation process in which the outermost region of the cast membrane undergoes phase separation induced by solvent evaporation, while the bulk of the structure is formed by solvent/non-solvent exchange during a quench step. This type of structure formation is defined as a dry/wet phase inversion process [Pinneau et al., 1990].

tion was poured slowly into ice water undergoing high-speed stirring in a blender. After filtration on a sintered-glass filter the naphthalene powder “solutions” were dried in an air stream for 24 h. The powders were then transferred to round-bottom flasks, and naphthalene was sublimed from the polymer blends at room temperature on a highvacuum rack for 10 days. The first DSC scans showed single Tg for these blends, intermediate between those of the two polymers, while the second and the following ones showed two Tg ‘s, corresponding to those of the two polymers.

9.3.3

9.3.3.2 Spray Drying

Special Methods

9.3.3.1 Freeze Drying The freeze drying of co-solutions of thermodynamically immiscible polymer pairs is a method of capturing kinetically the unlike molecules in molecular or “segmental” mixtures. PMMA/PVAc blends freeze-dried from co-solutions in benzene were reported homogeneous by dynamic mechanical and dilatometric observations. They showed a single Tg intermediate between those of the pure polymers [Miyata and Hata, 1970]. Further dilatometric and differential calorimetric data corroborated the findings of homogeneity [Ichihara et al., 1971]. Berghmans and Overbergh [1977] found that freeze-drying of dilute solutions containing two miscible polymers produced a system with three glass transition temperatures, corresponding to the Tg‘s of the two polymeric species and an intermediate Tg, characterizing their blend. A comparison of the density of freeze-dried and bulk PS has also led to the inference that freezedrying from dilute solution leads to a system in which individual polymer chains are collapsed to globular particles rather than being intertwined [Miyamoto et al., 1970]. Schultz and Young [1980] prepared homogeneous blends of immiscible PMMA/PS blends by freezing and sublimation. The solutions containing 5 wt% of polymer or polymer mixture were dissolved in naphthalene at 100°C with a magnetic stirrer, under N2 blanket. The hot solu-

Spray drying is by definition the transformation of feed from a fluid state into a dried particulate form by spraying the feed into a hot drying medium [Masters, 1985]. This is an ideal process where the end-product must comply with precise quality standards regarding particle size distribution, residual moisture content, bulk density, and particle shape. It involves the atomization of a liquid feedstock into a spray of droplets and contacting the droplets with hot air in a drying chamber. The sprays are produced by either rotary (wheel) or nozzle atomizers. Evaporation of solvent of matrix liquid from the droplets and formation of dry particles proceeds under controlled temperature and air flow conditions. The initial contact between spray droplets and drying air controls the evaporation rate and the product temperatures. The drying-chamber and air-disperser design must create a flow pattern that prevents deposition of partially dried product to the walls and the atomizer. There are three modes of contact: 1. Co-current: Drying air and particles move through the drying chamber in the same direction. On discharge the product temperatures are lower than the exhaust air temperature — this is an ideal mode for drying heat sensitive products. 2. Counter-current: Drying air and particles move through the drying chamber in opposite directions. This mode is suitable for products that require a high degree of heating.

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Compounding Polymer Blends

3. Mixed flow: Particle movement through the drying chamber experiences both co-current and counter-current phases. Spray drying has been used to prepare polymer blends from mixed lattices or co-solvent solutions.

Between 1966 and 1987 over 1000 patents on reactive extrusion was granted to 200 companies. In order of decreasing number of patents they were: Asahi Chem., Bayer, Kabel Metal Gatehoff, BASF, Showa-Electric, Mitsubishi Petrochem., Hitachi Cable, DuPont, Exxon, and many others [Utracki, 1998]. 9.4.1

9.4

Reactive Processing (compatibilization)

Details of the reactive compatibilization are discussed in Chapter 5 of this book. For this reason, the following text is only concerned with the mechanical aspects of the process. The term “reactive processing” is used to describe a polymer processing that involves chemical reactions. In principle, any processing operation can be conducted as a reactive process, viz. reactive injection molding (RIM). However, most often the term refers to reactive extrusion, and in particular, to the reactive compatibilization of immiscible polymer blends, usually conducted in a TSE. During the last 50 years, the latter machines have been used as chemical reactors for the polymerization, depolymerization (chemical recycling), polymer modification and compatibilization [Brown, 1992, Xanthos, 1992; Utracki, 1989, 1991, 1994, 1997]. Four features make TSE particularly well suited as a reactor for highly viscous polymeric reactants [Thiele, 1996]: (1) An ability to compartmentalize the reactive mass into small melt pools, bounded by the screw flights and barrel walls, with short mass and heat transfer distances. (2) The possibility of conducting a series of operations in a logical, well controlled sequence along the length of the extruder, viz. melting, mixing, reacting, devolatilization, homogenization and forming. (3) The ability to provide the desired extend of the interface renewal that controls the overall reaction rate. (4) The continuity of the process.

The Use of Twin Screw Extruders in Reactive Polymer Processing

TSE’s have been used to prepare new families of engineering materials of high performance, polymers and their blends. Polycondensation, free radical, anionic and cationic polymerizations were conducted to obtain, e.g., PA, PEST, POM, styrenic or acrylic resins. When the reaction is conducted in a low molecular weight liquid (e.g., solution or emulsion polymerization), usually devolatilization and compounding are carried out in a cascade second extruder. Functionalization and chemical modifications have been performed in TSE on virtually all polymers. 9.4.1.1 Advantages and Disadvantages of Reactive Processing Reactive processing is a continuous, flexible process offering the technical and economic advantages vis-a-vis other reactors [Kamal and Ryan, 1984; Snell, 1985; Tzoganakis, 1989; Kowalski, 1990]: • Low production volume, hence small losses for equipment breakdown. • Rapid and precise control of pressure and temperature; good heat exchange. • Ability to introduce reactants and catalysts at the desired locations. • Working with 100% of reactant content, using rapidly renewed interface, what leads to higher reaction rates then those possible in batch processes. • Self cleaning system, hence narrow and adjustable distribution of residence time. • Good control of the residence time and residence time distribution — the residence time distribution is narrowest for ICRR, CORI, and

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then NIRT. • Control of the reaction environment (airtight configuration, or inert gasses). • Ability to ascertain sequential high reaction rates. • Capability of rapid removal of a large volume of volatiles. • Controllable amount of dispersive and distributive mixing. • Ability to disperse high viscosity materials without a solvent. • Ability to combine polymerization (or copolymerization), degassing, compounding, and forming, what results in overall more economic industrial process • Possibility of automation via the close-loop control of product quality. • Discharge through a forming die. The disadvantages of the reactive extrusion are: • High initial capital investment. • High cost of the reactions that require long residence times (reactions requiring residence time greater than one hour, are not acceptable). • All the contaminants, inhibitors, and possible sources of side reactions must be removed or deactivated — this includes the moisture, residual monomers, catalysts, etc. • Difficulties in rapid removal of locally generated large heats of reaction. • A need for high accuracy of ingredient addition and precise control of the process parameters. • The reaction may need to be carried out under a blanket of dry, inert gas. • Product quality is influenced by the residence time distribution. • Run out reaction usually leads to high viscosity product, thus for safety reason high torque machines are needed with capability for rapid feed reduction, and increase of temperature in the final barrel sections. • The process must be close-loop controlled, using either rheological (MFM), infrared, or ultrasonic sensors. • The mathematical models of the reactive extrusion are in early stage of development.

The reactive processing is controlled on two levels: by selection of equipment, screw configuration, and ingredients, as well as by the control of process variables. Depending on the type of reaction, the most appropriate TSE type must be selected, and the appropriate steel alloy for the inner barrel and screw elements selected. Most manufacturers offer a selection of alloys with different resistance to abrasion and corrosion (to acids or halogens). Similarly, there is a variety of the control systems that must be considered knowing the number of sensor to be used and a need for protocolling. It is important that the metering system is selected to deliver the materials with the required precision and continuity. For liquid injections, the check valves that prevent freeze-off during the startup or shut down should be used. Reactive compatibilization requires good dispersive and distributive mixing of reactants and sufficient residence time for the reaction to take place. The introduced components should be immediately uniformly mixed. Thus, appropriate mixing screw elements (kneading, turbine, etc.) should be placed in the reaction zone of the extruder, where to the reactants are introduced. Furthermore, adequate heat removal capabilities must be provided for the zone where the exothermic reaction takes place. It has been found that barrel cooling is not always adequate and screw (shaft) cooling has been used. Usually, the venting zone immediately follows the reaction one. Since in the venting zone the screw is only partially filled, between the two zones there is a restriction of reversed screw elements, kneading element, check valve, or a throttling ring. Within the venting zone multiple flight elements should be introduced to provide rapid surface renewal. Owing to milder operating conditions, corotating TSE (CORI) have been preferred. However, the counter-rotating TSE’s (ICRR) are used with increasing frequency, e.g., for polymerization of either polymethacrylates or polyurethanes. In the latter case, owing to higher pressures and stresses than those in CORI, the reaction must be controlled with greater precision. For example, during PBT modification with epoxy,

Compounding Polymer Blends

the reaction started before the liquid was properly distributed in the melt, what resulted in poor quality products. In conventional TSE the residence time of up to 10 min can be achieved [Dreiblatt, 1989]. Special TSE systems for slow reactions have also been developed. For example, a TSE with a prereactor vessel may offer the residence time up to 45 min. However, the modern tendency goes in the opposite direction. TSE’s are offered with the screw speed of up to 1200 rpm (laboratory experiments are run at speeds of up to 3,000 rpm) that results in the residence time of the order of seconds. Hence, there is a rush to accelerate the reaction rates by catalysis, intensive mixing, as well as temperature and pressure. The residence time distribution depends on the screw configuration, throughput (or the degree of fill) and the screw speed. All three factors need to be optimized considering the rate of reaction, degradability of materials and productivity. 9.4.1.2 Development of Reactive Compatibilization Process Reactive processing combines fine polymer chemistry with polymer processing. Thus, development of the reactive compatibilization process involves: 1. Studies of chemical kinetics in glassware and internal batch mixers. 2. Determination of the rheological and thermal parameters of the reaction ingredients and products. 3. Mathematical modeling of the reactive processing. 4. Reactive extrusion. 5. Detailed evaluation of the properties of reaction products. 6. Scale-up to industrial scale. 9.4.1.3 Laboratory Equipment To develop reactive compatibilization, the laboratory extruder should match the production unit. Thus, a TSE, either CORI or ICRR, with segmented barrel and screw should be used.

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The machine should have the screw diameter D • 30 mm, with length to diameter ratio L/D • 40. The extruder must provide an easy access to the reactive medium as well as the quenching capability with either screw or barrel removal. It must be fully instrumented with temperature and pressure sensors all along the barrel. It should offer the possibility of working under dry inert atmosphere, and have the capability of closed-loop control operation with an appropriate set of sensors. The machine should be equipped with a minimum of three feeders for solid and two for liquid ingredients. At least one vacuum system with cold trap is necessary. The laboratory should have drying facility for all components that are to be used, e.g., a drying tower with closed resin transfer facility, as well as pre-blending facilities that can be operated under controlled temperature and atmosphere. The extrudate is usually granulated (thus, a need for the appropriate dies, cooling facility and either granulator or pelletizer), and subsequently must be formed into suitable specimens and tested. Forming is usually done by injection molding into a standard, multi-cavity test mold. The reactive extrusion is a process demanding precise operation by well-trained personnel. The knowledge of the extruder, compounding and mixing principles, design of screw and temperature profile, local residence time distribution, etc., is required. For optimization of the reactive extrusion in- or on-line process monitoring, control, and data logging are recommended — this imposes further demands on the extruder operator. 9.4.1.4 Reactive Compatibilization Strategy The experience indicates that the best results are obtained if: 1. The chemical reaction has low activation energy. 2. The required residence time is short (t0.5 ” 75 s), what implies high reaction rates, and frequently catalysis. 3. The process can be simpler if the reaction is insensitive to moisture or oxygen.

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The kinetics of the process should be studied on a model system in glassware. The kinetic parameters should be determined within the full range of the expected process parameters. Sensitivity of the reaction to moisture and oxygen should be examined. Next, the kinetics should be verified using polymeric ingredients. It is essential to know the life time of the activated sites — this determines the sequence and location of the reactant feeds. For example, if the radical is unstable, one may be better off feeding the peroxide and grafting monomer at the same location. Factors affecting group reactivity are listed in Table 9.22. Since reactive compatibilization is carried out across the interface, the screws of the selected TSE must provide a suitable geometry for the mixing, mass and heat transfer. The length of the reaction zone and the mixing intensity must ascertain sufficient mixing for the reactive groups to meet, but without the accompanying mechanochemical degradation. It is essential that there is adequate devolatilization. The reaction byproducts should be identified, removed or rendered harmless (catalyst deactivation may be needed) [Dagli et al., 1994; Maier and Lambla et al., 1995; Sundararaj et al., 1995; Hu et al., 1996a]. In principle, first high melting point polymer should be fed and melted, then the second polymer, and liquid(s) should be added. After the reaction zone the vacuum line should remove the volatile reactants (e.g., unreacted monomers) and byproducts. Finally, just before the die, a filler (e.g., glass fiber) and unstable ingredients should be added. The reaction must be terminated when the product leaves the extruder. It is important that the process is well monitored and controlled [Broadhead et al., 1996; Gendron et al., 1996] During the development stage, it is essential that samples are frequently taken from the extruder and analyzed for the extent of reaction as well as for the morphology and the key properties. (The sampling methods are discussed in Part 9.2.1.2 Morphology evolution in an extruder). It may also be profitable to verify the whole reactive compatibilization process by the carcass method.

The initial strategy formulated on the basis of chemical kinetics can serve as a starting point for the mathematical modeling of reactive extrusion. However, to ascertain successful strategy mathematical “experiments” must be carried out, using realistic models of the process. To optimize the process parameters and the screw design, the model requires information on kinetics, rheology and thermal properties [Menges and Bartilla, 1987; Polance and Jayaraman, 1995]. The type of required rheological and thermal information depends on complexity of the model. The minimum rheological information is the shear viscosity as a function of the deformation rate, temperature and the degree of conversion, η = η(y, T, k). The minimum thermal information is the heat of reaction, transition temperatures of the ingredients and their enthalpies, as well as heat capacities and heat conductivity coefficients. The model must provide information on the sequence of extrusion steps, e.g., melting, homogenization of ingredients, chemical reaction, devolatilization, morphology, performance. The reactive extrusion should be optimized using a computer modeling procedure. The mathematical model of the reactive extrusion should combine three subroutines: 1. Pressure and temperature profile as a function of the degree of fill and screw length [Michaeli et al., 1991, 1993, 1995; Michaeli and Grefenstein, 1996]. 2. Rheology and chemical reaction (heat balance!) as a function of screw length [Chen et al., 1996; Kye and White, 1996; Hu et al., 1996b]. 3. Variation of morphology along the screw length [De Loor et al., 1996b]. An analytical model based on the energy balance in CORI was proposed. Anionic polymerization of ε-caprolactam in CORI was modeled, considering continuity of the process from one screw element to another. The model took into account the heat of reaction, viscous dissipation and heat transfer across the barrel. The calculations were performed for three different screw configurations comprising the transport and kneading elements. The predicted total conversion agreed with the experimental values within ± 3% [Kye and White, 1996].

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Compounding Polymer Blends

Table 9.22. Factors affecting group reactivity during reactive extrusion Factor Molecular weight Neighboring groups Molecular configuration Electrostatic effect Accessibility Supra molecular effect Mechanical shearing

Influence No influence of chain length onto which a functional group is attached was observed. Exceptions occur in stiff chain polymers. May increase or decrease group reactivity. Depending on chain configuration, the functional groups may co-interact what alters their reactivity. Polarity of the reactive groups leads to repulsion or attraction effects. Stereo-barriers between functional groups. Aggregation, associations, and clustering slow down the reaction, but extent of these effects is difficult to predict. Straining the macromolecules result in generation of stresses that may activate some bonds. Mechanically induced chain scission has been explored for grafting polymers and rubbers.

Effects of the reactive processing are usually assessed either by analytical means or by their effects on morphology and performance. The analytical methods usually involve selective either extraction or destruction of one polymeric ingredient, followed by GPC, rheology, DSC, 13 CNMR, FTIR, and other analyses [Polance and Jayaraman, 1995; Maier and Lambla, 1995; Miyoshi et al., 1996; Kye and White, 1996]. The morphological studies should be conducted on several levels of magnification, from the optical to the transmission electron microscopy, TEM [Vesely, 1996]. For detailed discussion of the methods used to characterize the morphology see Chapter 8 Morphology of Polymer Blends. Table 9.23 lists examples of the compatibilization studies conducted in laboratory TSE’s, whereas Table 9.24 provides examples from the commercial patent literature. It is noteworthy that about 90% of patents on polymer blends published during the last few years, specify a TSE as the preferred compounder. Exceptions are blends formulated for oriented fibers and films (e.g., with LCP) that require high die pressure and thus are usually prepared in a SSE. Similarly, elastomeric blends of either PO or PVC are preferably prepared using one of the older methods, viz. roll mill or Banbury mixer.

9.4.2

Scale-up to Industrial Size

Scale-up is one of the most difficult engineering task. The main source of difficulty originates in the multifunctional capabilities of TSE. In addition to the primary functions of an extruder (solid conveying, melting and pumping) TSE is used to mix, devolatilize, and more and more often to conduct the chemical reactions. For these reasons, several schemes for the scale-up have been developed, each focusing on one particular process variable, viz. constant specific energy (CSE), constant thermal conduction (CTC), constant residence time (CRT), etc. Usually, the CTC forms the lower, whereas the CRT the upper boundary. The scale-up starts with optimization of the compounding conditions in a laboratory-scale TSE, having the screw diameter D ” 35 mm. Owing to the difference in thickness of the molten polymer between this and commercial size extruders (D • 100 mm), it is advisable that the scale-up involves an intermediate stage, a mid-size TSE with D • 50 mm. The experience also showed that it is more difficult to scale up from 30 to 50 mm than from 50 to 200 mm TSE. Thus, the intermediary step may save a large amount of time, material, and the plant production capability. In principle, good mathematical model that combines melting, fluid mechanics, morphology evolution, devolatilization and chemical reactions

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Table 9.23. Examples of reactions conducted in laboratory TSE’s Process

TSE

Comment

Reference

PP peroxide degradation

CORI or SSE

PA-6 polymerization Polyarylate synthesis TPU polymerization Radical polymerization of styrenics Rubber polymerization PA toughening MMA grafted onto EPR PO-grafting PET/PA-6 + p-toluenosulfonic acid Radical polymerization of MMA POM synthesis from trioxane EVAc grafted with MA & styrene Continuous polymerization of PEI Crosslink: PE + CH2=CH-Si(OR)3

CORI or CRNI CRNI + SSE CORI CORI CORI CORI CRNI CORI, CRNI CORI ICRR CORI CRNI + SSE modified CORI

Melt flow monitor (MFM) is needed Xanthos, 1992; Fritz & Stöhrer, 1986; others Nitrogen blanket viz., Michaeli et al., 1991 Devolatilization using 3 vents T = 90 to 260°C 3 CORI used Illing, 1970 T = -5 to 100°C Sutter & Peuker, 1977 T = 310°C Epstein, 1979 Staas, 1981 N = 100 rpm Binsack et al., 1981 Compatibilization of PA-6/PET Pillion & Utracki, 1984, 1987 Kinetic studies Stuber & Tirrel, 1985 Conversion • 90% Todd, 1988 Study of the grafting parameters Sakai et al., 1988 T = 25 to 400°C Schmidt et al., 1989

CORI CORI

Peroxide added Controllable crosslink density

Ultsch & Fritz, 1990 Curry & Andersen, 1990

CORI CORI CORI CORI

Moisture crosslinking MFM monitoring Plug-flow reactor Incorporation of NBR by free radical grafting (DCPO) MW depends on initiator concentration not N Compatibilization and dynamic vulcanization T = 180 to 240°C, 3 screw configurations examined T = 195 to 230°C ICRR better for visbreaking, CORI for maleation To enhance paintability Reactive blends compounded with fillers Compatibilization of HDPE/PA-6 Toughened PP had crosslinked EVAc/EMAc phase Accelerated reactivity caused by devolatilization Peroxide modification

Ultsch & Fritz, 1990 Curry & Andersen, 1990 Hagberg, 1992 Sakai et al., 1992

Reacting acrylic acid grafted PE with oxazoline grafted PS Vinyltrimethoxysilane grafted PE PS-PO blending Imidization of MA-polymers Toughening of unsaturated polyesters PS and S-I block copolymer anionic polymerization initiated with Bu-Li PA-6 with brominated isobutylenemethylstyrene elastomer Peroxide-induced degradation of PP

CORI with nitrogen blanket CORI

Maleation of PP Peroxide-induced degradation and maleation of PP Polyetheramine modification of PP Maleation of post-consumer plastics (HDPE, PET, PP, PVC) PE grafting with GMA Reactive compatibilization of EVAc/EMAc/PP blends Mono-esterification of SMA using 1-octanol Reactive extrusion of PO blends

CORI ICRR and CORI

ICRR

CORI CORI CORI CORI’s CRNI CORI

Michaeli et al., 1993 Kuwamoto, 1994 Ebner & White, 1994 Martinez et al., 1994 Kim & White, 1995 Clark, 1995 Xanthos et al., 1995 Wang et al., 1995 De Loor et al. 1996 Chen et al., 1996 Kim, 1996

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Compounding Polymer Blends

Table 9.24. Examples of compatibilization in TSE’s Process

TSE

Comment

Reference

Free radical degradation of PO by addition of peroxides Reactive toughening of PC with functionalized TPO’s PPS modification by partial crosslinking with di-nitro compounds Preparation of PC-PBT copolymers

CORI, ICRR, CRNI CORI

Graftable monomer may also be added The blends may also contain PPE PPS, PEST or acrylic Fillers can also be added

McCullough & Bradford, 1996

Reactive blending of PPE/SBS with dioctylamine Reactive blending of epoxy

CORI CORI CRNI CORI

triazine capped PPE with PBT/SBS Epoxidized-PPE blended with TSE grafted-PO and SBS PE/PB non-reactive blends Leistritz TSE Reactive toughening of PA CORI

Compatibilizer for PBE/PBT/SEBS blends To improve resistance to loss of toughness upon thermal cycling SEBS, (SB)n with epoxidized

Khouri & Stoddard, 1995 Kohler & Sarabi, 1995 Chambers et al., 1995 Richards & Pickett, 1995

Yates, 1995 elastomer can be added PE or PP grafted with t-alkyl carbamate Campbell & Presley, 1995 Sato, 1995 Nakajima, 1994

CRNI

Impact modification of PE EEA or CHR containing glycidyl ether moiety Incorporation of PP and styrene-grafted PP Blends can also be prepared in solution or using internal mixer Other PEST or PC may also be added

TSE

Melt blending in TSE at 285°C

Smits et al., 1994

CORI

Ethylene-butylmethacrylateglycidylmethacrylate added Impact modifier can be added

Flexman, 1994

PPE/PS/PO/SEBS blends

ICRR

PA toughened by incorporation of solution-maleated SEBS Polyglutarimide/Liquid Crystal Polyesters PPE/PA blends improved by incorporation of a polysiloxane POM with TPU, PA, PO, PEST

CORI

PBT poly(styrene-co-methacrylic acid) blends PPE/HIPS/EPDM-SAN

TSE

Poly(dimethylsiloxybiphenylen epoxide) PBT blends Impact modification of PA PPE/PA compatibilized by addition of trimellitic anhydride acid chloride HDPE/LDPE non-reactive blends Syndiotactic P, maleated PPE, and elastomer blends Semicrystalline, semi-aromatic PA blended with polyesteramide

CORI

TSE

DeNicola & Guyer, 1994 Gelles et al., 1994 Hallden-Abberton et al., 1994

Yang & McCready, 1994

CORI CORI

Fillers (carbon black, mica, ...) can also be added Improved flame retardancy and impact strength Acidified core-shell copolymers T = 232-287°C, N = 290 rpm

Fujii & Ishikawa, 1994

Liu & Liwak, 1994 Aycock & Ting, 1994

TSE TSE

Improved low-T properties Compounding at 300-350°C

Bohm et al., 1994 Okada & Masuyama, 1994

CORI

Extruded at T = 288°C, N = 170-215 rpm

Goldwasser & Chen, 1994

Jordan & Webb, 1994

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Table 9.24. Continued Process

TSE

Comment

Reference

PAI/PPS/organic di-isocyanate Impact modification of POM by blending with TPU, pentaerythritol and di-phenylmethatane-di-isocyanate PPE/PA/SEBS blends compatibilized by addition of maleic anhydride Reactive blends of PPS with epoxy-functionalized polysiloxane PC/PEST/epoxidized-PO/ABS

TSE CORI

T = 360°C T = 210°C, residence time t = 1-2 minutes

Kawaki et al., 1994 Nagai et al., 1994

CORI

Lee, 1994

Grafted PMP with glass fiber sized with carboxylic styrene-butadiene PP/PEST/PP-PEST graft copolymer Recycled domestic or industrial plastics, reactively stabilized PPS/PPE/PA/ SEBS and acidified-PO PPE/PEST/polystyrene-polycarbonate copolymer PC/acrylic multipolymer/grafted, unsaturated rubber PEST/PA/EEA-GMA copolymer Polyphthalamide/PP/carboxylated-PO

CORI

Blends compounded at T = 232-293°C, N = 280 rpm A dry blend extruded at T = 130-290°C, vacuum devolatilization T = 270°C, N = 250 rpm, vacuum devolatilization T = 260-290°C, N = 250 rpm,

Hagenson et al., 1994

TSE TSE

D = 45 mm, T = 280°C, N = 200 rpm T = 190°C;

Fujita et al., 1994 Pauquet et al., 1994

TSE CRNI

D = 50 mm, T = 300°C D = 20 mm, T = 120-266°C, N = 400

Ishida & Kabaya, 1994 Brown & Fewkes, 1994

TSE

D = 30 mm, T = 232°C, N = 100 rpm

Zimmerman et al., 1994

TSE CORI

T = 270°C D = 30, T = 285 - 320°C, N2

PPE/maleated-EPR/EVAc grafted with GMA/PA/PEST/4-functional epoxy resin/oleic acid amide

High speed TSE

Blends of semicrystalline PA (matrix) and amorphous PA (dispersed phase) PP/EPR/SEBS PC/PEST/amine functionalized elastomer with maleated EPR Blends of aliphatic and aromatic PA’s with aliphatic polyester (PCL) Semi-aromatic PA blends with PC POM with core-shell copolymer PET or PBT with 1 wt% sodium sulfonate groups, PE-ionomer, PBT Polycycloolefin (PCO)/graftedPCO/grafted elastomer/PA PC/ABS containing hydroxyethyl methacrylate/flame retardant

CORI

Maleation: D = 30 mm, L/D = 28, T = 230°C, N = 60 rpm; compounding: D = 47 mm, T = 310°C, N = 500 rpm D = 30 mm, maleated EPR added

CORI CORI

Han, 1994 Laughner, 1994

Natarajan et al., 1994 Paschke et al., 1994 Nishio et al., 1994

Schmid & Thullen, 1994

CORI (tapered) T = 246°C, N = 350 rpm CORI D = 34 mm, L/D = 40, T = 250-270°C, N = 200-250 rpm TSE D = 30 mm, T = 285°C

deNicola & Conboy, 1994 Akkapeddi & vanBuskirk, 1994 Hamada et al., 1994

CORI TSE TSE

D = 30 mm, T = 270°C D = 30 mm, T = 200°C D = 30 mm, T = 245°C, N = 80 rpm

Gambale et al., 1994 Sasaki et al., 1994 Tajima et al., 1994

TSE

D = 30 mm, T = 260°C

Moriya et al., 1994

CORI

D = 28 mm, T = 249°C

Vilasagar & Rawlings, 1994

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Compounding Polymer Blends

Table 9.24. Continued Process

TSE

Comment

Reference

Modified PPE reacted with terminal amino groups of PA + SEBS PET/PC/PEC/elastomer PPE/PA-6, PO or PBT/maleic anhydride/SEBS PA/PO/SEBS compatibilized by acidified SMA, SEBS-MA, ... PPE/PO compatibilized by EAA PP/ionomer/EGMA/EPR PA/maleated-EPR PC/PP/acryloid PEI/PP Reactive grafting and polymerization of vinyl acetate in PP/polyester blend

TSE

Vacuum devolatilization, T = 290°C

Kodaira et al., 1994

TSE TSE

T = 260-300°C T = 230°C, N = 250 rpm

Freitag et al., 1994 Takagi et al., 1994

Leistritz ICRR

L/D = 28, T = 249°C

Chundury, 1994

Leistritz TSE CORI TSE TSE CORI TSE

T = 270°C D = 28 mm, L/D = 28, T = 210°C D = 44 mm, T = 240-270°C T = 255-295°C T = 304-322°C T = 105-210°C

Cottis & Natarajan, 1994 Dawson, 1993 Ohmae et al., 1991 Liu, 1985 Giles & Schlich, 1984 Waniczek et al., 1983

should provide solid guidance for the scale-up. However, so far such a complex model has not been developed. In consequence, the methods are pragmatic and less formal. First, it must be decided which function is the most important and wherever possible, an appropriate mathematical model should be selected. At least three experiments must be conducted — both computer-type and extruder-type — using three sets of values, e.g., composition, throughput, and screw speed. For each composition the nine specimens should be examined for the desired performance characteristics, viz. mechanical performance, permeability, solvent resistance, etc. Once the model is found to well describe the experiments, the exercise must be repeated for the larger diameter extruder, the model parameters re-adjusted, and the optimum screw configuration and the processing conditions computed, then experimentally verified. Less formal method uses the power-law exponents. Here, the ratio of the screw diameters, δ = D2/D1 > 1, is the principal variable. When scaling, the screw channel angle is to be kept constant, and the screw section length increased by δ. The channel depth is also increased as:

H2 = H1δ0.92. If the ratio of heat transfer to throughput is assumed constant, then: Q2 = Q1δ2, and Ν2 = Ν1/δ. On the other hand, if ESP is taken as constant: Q2 = Q1δα, and Ν2 = Ν1/δβ, where the exponents α = 2.84 - 0.08/n, and β = -0.08(1 + 1/n), with n being the power-law exponent of the shear viscosity vs. rate of shear. Assuming CRT: N2 = N1, and Q2 = Q1δ2.92, the experimental data were found to closely follow the ESP = const. approximation [Potente, 1985; Rauwendaal, 1986]. For the applications where degassing is crucial, the degassing constant, K”, for specific TSE geometry must be established: K”=(Q/LN1/2) ln (∆c0/∆cf)

(9.32)

where: Q - throughput, L - vent length, N - screw speed, ∆co = co - ce,o, ∆cf = co - ce,f, co being the initial concentration of volatiles, ce the equilibrium value (initially with subscript o and finally with subscript f) [Hashimoto et al., 1993]. From the scale-up strategy Q and N are computed, then Eq 9.32 is used to determine the required vent length and/or the pressure of the vacuum line (that controls the equilibrium values, ce,o, and ce,f).

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In applications where mixing efficiency is most critical, the scaling should be based on the total induced strain, γ — its value (optimized in the small TSE) should be preserved in the larger machine [Manas-Zloczower, 1991].

9.5

Summary

This chapter was divided into four parts: Fundamentals of mixing, Blending methods and equipment, Non-mechanical blending, and Reactive blending. The theory of mixing distinguishes the dispersive and distributive mixing, the former dependent on stress, while the latter on strain. Both types of mixing are more efficient in the extensional than in shear flow field. The extensional flows more economically generate large dispersing fields and induce larger strains. Independently of the flow field, the best mixing strategy is generation of a series of deformation and folding steps. A majority of commercial polymer blends is melt-blended, mostly in TSE, less frequently in SSE, and rarely in batch mixer and other

specialty mixing devices. During the last five years 94% of patents on polymer blends describe preparation of blends in a TSE. This is particularly significant, since to make the invention more attractive, many of the texts claim that the blend “may also be prepared by other methods, viz. solution blending, in a batch and SSE mixer”. The high percentage of TSE being used to develop blends corresponds to the recent estimates of these machines use in the engineering alloys’ production — some analysts put the number as high as 90%. This dominance of TSE most likely stems from the growing use of the reactive blending. Since SSE is not well suited for conducting chemical reactions, the demands gravitate toward the twin-screw machines (an abbreviated comparison of the SSE and TSE performance was given in Tables 9.16 and 9.17). The TSE selection is based on the production requirements. A summary of the performance characteristics is given in Table 9.25 [Rauwendaal, 1993]. The single-screw extruders are being used for the preparation of blends that either do not require

Table 9.25. Functional characteristics of TSE Function

Throughput Distributive mixing Dispersive mixing Devolatilization Melting Conveying Self-wiping High screw speed Residence time distribution Pressure development Screw separation Feed capacity

CORI

ICRR

CRNI

low N

high N

low N

high N

+ + o o + + + + + +

++ ++ + + + o ++ ++ o o + o

+ + o + + ++ + o ++ ++ o ++

+ + ++ ++ ++ + o + + + o ++

++ ++ + + ++ o o + +

641

Compounding Polymer Blends

compatibilization, or are compatibilized by addition of non-reactive compatibilizers. For both applications, SSE is frequently supplemented by additional mixing devices, viz. mixing screws, add-ons, extensional flow mixer, etc. To the first category of blends also belong these systems that are produced as filaments or fibers. It has been found that creation of a large interface (by stretching) reduces the need for compatibilization (when the viscosity ratio is large, it may be difficult to transfer stress to the non-compatibilized dispersed phase). To the second category of blends belong these systems that are compatibilized and at the same time impact modified by the addition of multi-component core and shell additives. To the third group of blends belong systems based on either PO or PVC with a large amount of elastomer. These highly viscous systems are prepared on roll mills, in a Banbury-type mixer, or non-intermeshing twin-shaft continuous mixers. When selecting the mixing strategy one must consider the thermo-mechanical degradability of the blends ingredients (including the compatibilizer!), the time scale of the relaxation processes, as well as the stress-induced coalescence. Some polymers, e.g., PP, are sensitive to degradation. For this reason better performance may be obtained from PP/EPR blends using low dispersing energy, just sufficient to adequately disperse the elastomer without inducing too extensive degradation of the matrix. One of the most efficient mechanisms of mixing is development of fibrils that upon reduction of stress disintegrate by the capillarity mechanism into thousands of micro-droplets. This mechanism, well known in the Newtonian systems, seems to be less efficient in systems comprising highly elastic dispersed phase and low interfacial tension coefficient, e.g., in R-TPO systems. The loss of efficiency is related to, on the one hand small interfacial energy, and on the other long relaxation times. However, it must be stressed that dispersion of high viscosity polymer into a less viscous matrix is greatly facilitated by the extensional flow field. When selecting the mixing devices it is important to remember that the final degree of dispersion

is related to dynamic equilibrium between the dispersive and coalescing mechanisms. Excepting the concentration of the dispersed phase, they both depend on the same parameters. Thus, the rate of drop diameter decrease during dispersing is: ∂d/∂t = -k1yd, whereas that of coalescing: ∂d/∂t = k2y φ8/3/d, hence, the rate of shear similarly affects the dispersion as the coalescence process. For this reason, when all other parameters are constant, the drop diameter is often found to be insensitive to the screw speed. It has been even reported that, owing to generation of very large interface and/or shear degradation of a compatibilizer, an increase of the shear rate may cause the drop size to grow.

9.6

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9.7

Nomenclature and abbreviations

9.7.1

Nomenclature

a b B C, ci CL D d d* do

area occupied by a copolymer molecule channel width in a rectangular flow channel droplet width, or the smaller diameter of prolate ellipsoid constants screw center-to-center distance deformability parameter droplet diameter. equilibrium droplet diameter. initial diameter of droplet

Dp dP/dz E e’ E+ EDK Fd, Fp h H hc K” kB L L L/D m M or MW MWD n N Nav nc Ns p P pr q Q, Qc R Re Ri T t tb tb* tc u V v vx Vx/V

particle diffusion coefficient. pressure gradient energy of mixing relative flight width = e/(πDssinϕ) threshold energy of coagulation. macroscopic bulk breaking energy. correction shape factors mean channel depth Hamaker constant of the system. critical separation distance degassing constant Boltzmann constant. length of a drop screw length, or vent length length to diameter ratio (e.g., for TSE barrels) coefficient of the power law molecular weight MW distribution power flow index, or number of particles screw rotation speed Avogadro’s number number of divisions in a static mixer, SM number of striations aspect ratio pressure probability that two particles that have collided result in coalescence. sinusoidal distortion. throughout, throughput capacity radius of a dispersed drop Reynolds number interfacial area temperature. time, or pitch necessary time for breakup of droplets. dimensionless breakup time. coalescence time velocity field free volume of the extruder velocity velocity in the x direction volume fraction of emulsion undergoing uniform shear.

Compounding Polymer Blends

~ V x α αi αo β χ12 δ ∆l∞ ∆P [ε˙] φc φd φm η ηd, ηm

γ˙ κ κcrit Λ λ ν12 ν°12 ρ, ρd σ11 − σ22 σ12 τ Ω(Λ, λ) 9.7.2 CHO CHX CMC CORI CRNI CRT CSE CTC CTM DCA

average flow velocity linear distance orientation or tip angle. break-up parameter the distortion at t = 0 orientation angle thermodynamic binary interaction parameter solubility parameter interphase thickness for ∆M∞ pressure drop velocity gradient tensor percolation threshold volume fraction volume fraction of dispersed phase maximum packing volume fraction. viscosity. viscosity of dispersed and matrix phase respectively. shear rate capillarity (or Taylor) number. critical capillary number distortion wavelength viscosity ratio interfacial tension coefficient between phase 1 and 2 interfacial tension in a quiescent polyblend. density, droplet density the first normal stress difference. shear stress relaxation time Tomotika’s tabulated function. Abbreviations cyclohexanone cyclohexane critical micelles concentration co-rotating intermeshing TSE counter-rotating non-intermeshing TSE constant residence time constant specific energy constant thermal conduction cavity transfer mixer dichloroacetic acid

DCPO DMA DMF DMX DSC EFM EME ESP EtAc EtOH FA FCM FTIR GPC

ICRR JSW MA MC MEK MeOH MFM MMA MW MWD NMR SD SEM SM SSE TCE TEM THF TMR TSE W&P ZSK

651

dicumyl peroxide dynamic-mechanical analysis dimethyl formamide dynamic melt mixer differential scanning calorimetry extensional flow mixer elastic melt extruder specific energy (kWh/kg) ethyl acetate ethanol formic acid Farrel continuous mixer Fourier transform infra-red spectroscopy gel permeation chromatography, or size exclusion chromatography, SEC intermeshing counter-rotating TSE Japan (Nippon) Steel Works methyl acrylate methyl chloride methylethylketone methanol melt flow monitor (measuring shear viscosity vs. shear rate) methyl methacrylate molecular weight MW distribution nuclear magnetic resonance spectroscopy spinodal decomposition scanning electron microscopy static (motionless) mixer single-screw extruder tetrachloroethane transmission electron microscopy tetrahydrofuran Twente mixing ring twin screw extruder Werner and Pfleiderer Zwei Schnecken Kneter (two-screw kneader)

CHAPTER 10

POLYMER BLENDS FORMING

Michel M. Dumoulin

National Research Council Canada, Industrial Materials Institute Boucherville, QC, Canada

10.1

Abstract

The key to optimizing polymer blends performance is control of morphology. Distinct types of morphology are required to improve specific properties. The matrix/dispersed phase type is necessary to obtain the maximum impact strength, while dual phase co-continuity will yield better stiffness/ductility balance and dispersion of lamellar domains allows to modify the barrier performance. Optimized properties therefore result from strict control over processing conditions and equipment, and from methods of generating and stabilizing the morphology, such as reactive compatibilization. The reviewed information shows that in addition to improving final properties, blending is increasingly carried out to improve processability. Easier processing can often be achieved through a negative deviation from the viscosity additivity rule. This route is particularly critical for high performance resins difficult, if not impossible, to process, such as PPE, PEI and aromatic polyamides. Easier processing can also be achieved indirectly, i.e., by adding a toughener to reduce a resin’s MW to lower viscosity while maintaining the system’s toughness. Other less dramatic changes include less temperature dependent rheology or increased melt elasticity. The chapter attempts to review the behavior of blends during forming. The processes are discussed first from the point of view of the general principles and the effect of material parameters. The general behavior of blends is discussed in terms of processability, flow-induced morphology and the effect of forming on blend performance. The various important forming methods are examined: extrusion, injection molding, blow molding, calendering, thermoforming, film blowing, fiber spinning, as well as solid state forming and compression molding. In each case, a description of the process and of the main process parameters is followed by a discussion on the behavior of blends of commodity, engineering and specialty resins. Processing conditions of commercially available blends are tabulated whenever possible. Forming of reinforced blends, quality and process control and finally reprocessing are also discussed.

L.A. Utracki (Ed.), Polymer Blends Handbook, 653-755. © 2003 Kluwer Academic Publishers. Printed in the Netherlands.

654

10.2

M. M. Dumoulin

Introduction to Polymer Forming

McKelvey defined polymer processing as an engineering specialty concerned with the operation carried out on polymeric materials or systems to increase their utility [McKelvey, 1962]. He further suggested to subdivide the field into three categories: (i) modification (including mixing), (ii) forming and (iii) bonding (i.e., finishing). The present chapter focuses on forming, i.e., on the processes that fix the shape of the products. Shaping is usually preceded by preparation (e.g., drying, mixing) and followed by finishing (e.g., painting, joining) steps. Some of these will be briefly discussed here whilst operations leading to preparation of polymer blends, e.g., mixing, compounding, are discussed in Chapter 6. Compounding. Similarly, processes involving chemical changes are beyond the scope of this chapter. A brief introduction to thermoplastics forming is given first, followed by sections discussing the principal processes: extrusion, injection molding, blow molding, calendering, thermoforming, film blowing, fiber spinning, solid-state processing, compression molding, and forming of filled or reinforced blends. The choice of forming method is dictated by the product geometry, final part weight, performance requirements and cost. Extrusion is a continuous operation that fixes two dimensions of the resulting products while molding generally describes discontinuous operations that fix the three dimensions of a part. Molding may involve physical changes (injection molding), chemical changes (monomer casting), or simultaneous physical and chemical changes (reaction injection molding or transfer molding of thermosets). For all polymers, significant changes in microstructure take place during and after forming, affecting molecular orientation, crystallinity and crystalline morphology. In addition, in the case of polymer blends, blend morphology is usually affected. All these changes may profoundly influence the product properties. Polymer forming involves a number of physical and sometimes chemical changes that need to be understood. Because deformation and flow are involved in all shaping operations, the concepts

of rheology are essential. Similarly, heat transfer, and pressure-volume-temperature, PVT, relations constitute critical components of the forming behavior and will be examined in this chapter. Polymer forming began with the chemical modification of natural polymers such as natural rubber vulcanization and cellulose acetylation. The first efforts to shape natural polymers and early synthetic ones into useful products such as textile fibers and films for packaging date from the middle of the 19th century. The first mixing devices, in the form of rotors in chambers for rubber compounding, started to appear in the 1820’s. The two-roll mill, the precursor of the modern calendering process, was developed in the 1830’s. Early ram extruders used for forming rubber were developed in the 1850’s. Screw extruders were first developed in the late 18th century to produce miscellaneous products like lead pipes, soap and food. Screw extruders for wire and cable coating with rubber were introduced in the 1890’s. The development of the first extruders designed for thermoplastics followed closely the development of the thermoplastics themselves and were first introduced in the mid 1930’s. The historical development of polymer processing, including forming processes, is detailed in Chapter 1 Introduction to Polymer Blends. 10.2.1

General Description of Polymer Forming

In forming operations, the material must flow to be shaped, and solidify while retaining the desired shape. Analysis of plastics forming involves defining a number of elementary steps, viz., solids handling, melting, mixing, forming, etc. In order to analyze the material’s behavior, three distinct stages may be considered: Plastication, during which the polymer changes state from solid to liquid. Forming, during which the polymer is formed into a specified shape by flowing under pressure, e.g., into molds or through dies. Solidification, during which the material solidifies, and possibly undergoes other post-forming operation such as demolding, trimming, cutting, polishing, etc.

655

Polymer Blends Forming

Polymers have unique properties important for processing, viz., low thermal diffusivity, high viscosity and viscoelasticity. Because of the low thermal conductivity, efficient plastication cannot be based on thermal diffusivity alone. Polymer melting requires generating heat dissipation through intensive deformation of the highly viscous melts, leading to thin films. As discussed in Chapter 7 Rheology of Polymer Alloys and Blends, molten polymers show constant viscosity (a Newtonian behavior) only in the low strain rate region. At higher shear rates, the viscosity generally decreases with increasing rate of deformation, what is know as a pseudoplastic behavior. The viscosity of synthetic polymers is generally high, resulting in large pressure buildups in processing equipment. The viscoelastic nature of polymers manifests itself during processing in several ways, e.g., (i) resulting in extrudate swelling at the die exit, (ii) stabilizing stretching operations, (iii) causing flow instability leading to extrudate distortions. Modern processing equipment usually employs an Archimedes-type screw that conveys the solid polymer (powder or granules) and causes it to melt progressively through heat transfer from the barrel and heat generation by viscoelastic dissipation of the energy introduced by the screw rotation. The processability window of a polymer system is defined as the processing conditions that will yield acceptable products. The window is not defined by a single set of conditions, but rather it may be represented schematically by an area on a processing diagram. For example, a processing window for the extrusion process is shown in Figure 10.1. The curves illustrate limits of the process. Line 1 indicates the maximum screw rotation speed of the extruder, Curve 2 represents the torque limitation of the drive system (region to the right and below of curve exceeds the allowable torque on the drive system). Curve 3 represents the maximum temperature allowable to avoid degradation. Line 4 indicates the minimum screw rotation speed to achieve the minimum throughput dictated by the economics of the process. In the present chapter, processability of the different polymer blends will often be discussed

in terms of degree of difficulty. Product shapes fabricated in commercially used plastic forming processes are listed in Table 10.1. The abbreviations used in this Chapter for the various processing operations and the degree of difficulty or processability are presented in Table 10.2.

Figure 10.1. Schematic representation of the processing window for the extrusion process. The four curves identify limits to the process: line 1 indicates the maximum screw rotation speed, curve 2 represents the torque limitation of the drive system (region to the right and below of curve exceeds the allowable torque on the drive system), curve 3 represents the maximum temperature allowable to avoid the thermal degradation, and line 4 shows the minimum screw rotational speed to provide the minimum throughput, acceptable by the process economics.

10.2.2

Influence of Material Characteristics

Three types of material properties have significant impact on processing: bulk, rheological, and thermal properties. 10.2.2.1 Bulk Properties Bulk density, as well as the physical shape of the polymer, influences the solid material flow in hoppers and in the feed section of screws. Low bulk density results in low mass flow rate, possibly causing solids conveying rates too low to feed enough solids to the plasticating section. In this situation, specially designed force-feed

656

Table 10.1.

M. M. Dumoulin

Product shapes obtained in forming processes

Product shape

EXT

Flat - films Flat - sheets, panels Linear - tubes, profiles Linear - fibers, yarns Hollow - bottles, containers Hollow - drums, tanks Open - Cups, trays Complex parts

Table 10.2.

IM

N N

BM

TH

COM

CAL

P

P N

FS

P

SSF

CAST

ROT

D P P

P P

D

FB N

N P D N N

N N

P N

P N P P

P

Abbreviations of the processability terms used in the chapter Processing method

Difficulty level

Code

Process

EXT IM BM THER COM CAL FS SSF CAST ROT FB

Extrusion Injection molding Blow molding Thermoforming Compression or transfer molding Calendering Fiber spinning Solid state forming Casting Rotomolding Film blowing

systems may have to be used. The particle size and shape can also affect the solids conveying characteristics in both the hopper and the barrel. For example, material in the form of spheroidal pellets will tend to flow better and entrap less air than granules or powder. In some operations, polymeric materials in tape or fiber form may have to be fed - these will require special feeding methods and careful adjustment of the processing parameters.

Code N P D

Processability Normal Possible Difficult

10.2.2.2 Rheological Behavior Polymer melt flow properties determine to a large extent the characteristics of the forming process. The detailed discussion on rheology is provided in Chapter 7 Rheology of Polymer Alloys and Blends. Rheological behavior will only be summarized here in order to provide sufficient information to understand and discuss polymer forming.

Polymer Blends Forming

657

Rheology relates the stress components to the respective rate of strain of the fluid. For example, the rheological equation for a Newtonian fluid shows a linear relation between stress, sxy, and shear rate, y: sxy = ηo(dvx/dy) = ηoy

(10.1)

where the zero-subscript indicates the zero-shear, or Newtonian viscosity, ηo. However, most synthetic polymers are non-Newtonian and their viscosity depends on the deformation rate. Figure 10.2 shows a typical viscosity curve for a high polymer melt. The rheological functions of a polymer depend on the temperature, pressure, deformation rate and strain, as well as on the macromolecular features: molecular weight, MW, its distribution, MWD, and the chain structure (e.g., tacticity, type and level of branching, etc.). Figure 10.3 demonstrates the effect of several factors, viz., MW, temperature, pressure, as well as additives and fillers concentration. A thorough analysis of forming processes relies on mathematical descriptions of the physical steps. These are independent of the material and are formulated with the equations of continuity, momentum and energy. These equations applied to the particular geometry must be combined with the rheological equation of state of the material. Furthermore, the relations expressing the temperature and pressure dependencies of such properties such as density, thermal conductivity, etc., are required. The combination of these relations constitutes the mathematical formulation of the flow process. These formulations are beyond the scope of the Chapter. The range of shear rates encountered in polymer processing varies approximately from y = 1 to 106 sec-1. For proper analysis of polymer processing operations, it is necessary to represent the behavior in each region with an appropriate model. The one most frequently used for non-Newtonian fluids is the power-law model: σ = m y n or

η = m y n-1

(10.2)

which approximates the behavior in the high shear rate region. The parameters m and n are referred

Figure 10.2. Shear rate dependence of viscosity for a high molecular weight PP at 180°C.

Figure 10.3. Effect of several variables, viz., pressure, temperature, resin molecular weight, as well as filler and additive concentration on a flow curve - shear viscosity vs. shear rate.

to as the consistency (or reference viscosity) and the power law index respectively. When n = 1, the power law reduces to Newton’s law. Hence, the extent of deviation from unity indicates the extend of non-Newtonian behavior. Most commodity polymers such as polyethylenes, PE, polypropylene, PP, polyvinylchloride, PVC, and

658

M. M. Dumoulin

polystyrene, PS, are strongly shear-thinning (or in other words, pseudoplastic) and therefore have small values of n ” 0.5. Engineering polymers such as polyamides, PA, polycarbonates, PC, polyethyleneterephthalate, PET, have lower molecular weight, and thus generally show higher values of n. Typical values of n for several common polymers are listed in Table 10.3. Table 10.3.

Values of the power-law index for selected polymers [Rauwendaal, 1986]

Polymer

n

HDPE LDPE LLDPE PP PS PVC PA-6 PA-66 PET PBT SAN ABS PC PMMA

0.50 0.35 0.60 0.35 0.30 0.30 0.70 0.75 0.60 0.60 0.30 0.25 0.70 0.25

improve its ability to describe the flow behavior of polydispersed polymers [Utracki, 1987a]: η = ηo [1+(y τ)m1]-m2

Temperature and pressure also strongly affect the viscosity of polymer melts and thus are important for processing. Increasing the temperature lowers the viscosity. The dependence may be expressed by an Arrhenius-type expression that involves the concept of an activation energy of flow, Ea: η = A exp{Ea/RTK}

(10.5)

where R is the ideal gas constant and TK the absolute temperature. For most common thermoplastics, the effect of temperature on flow can be approximated by Eq 10.5 [Westover, 1959]. Another commonly used relation describing the temperature dependence of thermoplastics above the glass transition temperature, Tg, is the WLF equation [Williams et al., 1955], discussed in Chapter 7 Rheology of Polymer Alloys and Blends. 10.2.2.3 Thermal behavior

The utility of mathematical models incorporating power-law Eq 10.2 is limited to either a small range of shear rate, or to its high values. To improve the description of the viscosity dependence on shear rate, several alternative relations were proposed, e.g., by Ellis, Bueche, Eyring, Carreau [Bird et al., 1960]. The latter relation: η = ηo [1+(yτ)m]-2

(10.4)

(10.3)

was derived for monodispersed systems [Carreau, 1969]. It is simple and it has the advantage of being able to describe the shear rate dependency of the viscosity from the Newtonian plateau to the power-law region. Eq 10.3 was modified to

The thermal behavior of polymers is a complex subject that encompasses several topics relevant to polymer forming: thermal transitions, thermal transport and PVT behavior. Table 10.4 lists a number of thermal properties of selected materials. Thermal transitions control the processing temperature, Tp. Forming normally takes place at Tp > Tx + 40°C, where Tx is the melting point, Tm, for semi-crystalline polymers or the glass transition temperature, Tg, for amorphous resins. Increasing the temperature lowers the viscosity, thus facilitating processing, but Tp must be kept low to minimize degradation. The thermal transport properties of polymers impose severe limitations on the forming processes. Heating and cooling stages are part of all the polymer forming processes. Important considerations are the total energy change and the rate at which the change takes place. The temperature distribution in the polymer system

659

Polymer Blends Forming

Table 10.4. Material

Thermal properties of selected materials Tm (°C)

PMMA PS ABS PVC HDPE PP PET PA-66 PC PTFE Aluminum Copper Steel

95 100a 105 199a 134 168 250 260 266a 330 600 980 1500

Tg (°C)

100 90 102 90 -110 5 70 50 150 -115

Thermal conductivity (103 W/m.K) 2.53 1.27 1.27 2.11 5.06 1.18 1.52 2.45 1.98 2.53 1270 1900 370

Heat capacity (10-3 J/kg°C) 2.34 2.09 2.09 2.51 3.77 3.77 1.88 3.13 2.09 1.26 0.96 0.38 0.46

Thermal diffusivity (10-8 m2/s) 8.9 5.7 3.8 6.2 13.9 3.5 5.9 6.8 7.8 9.1 4 900 5 700 1 000

Thermal expansion (10-6 m/m°C) 50 50 60 50 59 81 65 80 68 70 19 18 11

a: Low levels of crystallinity

during processing depends on heat transfer within the system and between the system and the outside as well as on heat capacity. Because of the low thermal conductivity, efficient plastication cannot be based on thermal diffusivity alone. Melting plastics requires generating high heat dissipation through intensive mechanical deformation of the highly viscous polymers, leading to thin films with uniform (on a macroscopic scale) properties. Melting therefore relies on heat generated from two mechanisms: conduction from the metal surfaces through thin films and viscous dissipation. Limitations due to the low conductivity also manifest themselves during cooling. Slow heat transfer causes longer cycle times during molding operations. It is also the source of variations in crystalline morphology in semi-crystalline polymers - longer cooling times allow for growth of spherulites before freezing. For thick parts, this phenomenon may cause core imperfections: on the one hand the longer cooling results in higher crystallinity and larger crystals (in comparison to

the skin region), and on the other in the presence of voids. The amount of heat required for melting varies considerably between resins, with the semicrystalline resins requiring more heat to melt their crystalline phase. For example, acetal or polyoxymethylene, POM, has a fairly high heat of fusion (∆Hm = 160-200 kJ/kg). Examples of total heat energy required for forming are: POM = 435-475, high density polyethylene, HDPE, = 720, PA-66 = 756-786 kJ/kg. As a first approximation, the PVT behavior is often neglected. Contrarily to gases, liquids and solids are relatively incompressible. Engineering approximations frequently consider that density does not change with pressure and only slightly with temperature. However, these assumptions are rarely valid for most aspects of polymer processing. For example, density changes during injection molding are significant. Upon an increase of temperature the change of specific volume, ∆v, is the sum of instantaneous changes, ∆vi, and a delayed expansion that occurs

660

M. M. Dumoulin

at a rate proportional to the difference between the instantaneous, v, and the equilibrium, ve, specific volumes. The first expansion corresponds to changes in intermolecular spacing due to increased thermal vibration and is proportional to the temperature jump. The second mechanism is delayed because the increase in free volume depends on diffusion and viscous flow. Computing the PVT behavior requires a thermodynamic equation of state determined from statistical mechanics, and verified by experimental data of volume changes as functions of pressure and temperature.

10.3

Forming Polymer Blends

10.3.1

Introduction

The history of development of polymer blends, detailed in Chapter 1, Introduction to Polymer Blends, provides insight into the history of the polymer forming processes. It demonstrates the speed at which blending has followed the introduction of new polymers and the speed at which processing methods were adapted and optimized for the new materials. Some of the oldest references to processing found in the literature go back to the late 18th century where calendering was described. Textile calenders made of wood and driven by horsepower were patented in the early 1800’s. Calendering was the preferred method for production of rubber sheets after the discovery of vulcanization by Goodyear in 1838. As soon as the first blends were prepared, in the 1840’s, when natural rubber, NR, and gutta-percha, GP, were mixed together, attempts at molding the newly developed materials were made [Parkes, 1846]. Obviously, the early blends were formed by methods adopted from the rubber industry, e.g., milling, calendering, extrusion. Wire coating with blends was first described in the 1930’s for PVC formulations [Fikentscher and Schmidt, 1936] and NBR/PVC [Badum, 1942]. Over the years, the patent literature on polymer blends has claimed mostly improvements of a number of properties, particularly impact resistance, paintability or heat resistance. However,

recent patents indicate that blending is increasingly carried out to improve processability. The last decades saw the introduction of a number of high performance resins that proved to be difficult, if not impossible, to process. Blending often proved the only way to process these materials, as was the case for PPE/PS (Noryl™) blends. Blending can also bring about less dramatic, but still useful, changes such as less temperature dependent rheology or increased melt elasticity. These will be discussed for the different systems below. Processing must be carefully examined in all cases since blend properties depend on properties of ingredients, composition and morphology. The forming stage often has a dramatic influence on the final morphology. During processing, the material undergoes complex deformation affecting the morphology and thus the final properties of the material. For example, a strong elongational flow component is present during convergent flow through a die. Depending on the rate of cooling, extrudates may show uniaxial orientation of elongated drops or enhanced dispersion. Therefore, tailoring polymer blends for specific applications requires control of the morphology through processing. The presence of a second phase often makes it necessary to modify process parameters. The most important aspects to be kept in mind are: (i) the generation and preservation of the optimum morphology, and (ii) ascertaining the uniform distribution of the components to avoid flow segregation. A spatial redistribution of the ingredients, caused by the tendency of the flowing systems to minimize energy, can occur during processing. This phenomenon was observed for HDPE/PA-6 mixtures during extrusion through capillaries causing the low viscosity component to migrate toward the wall and the high viscosity toward the center of the capillary [Dumoulin et al., 1986]. The present chapter aims at discussing the behavior of blends during forming. Tables 10.5 to 10.11 show the applicability of the usual forming methods to all commercial polymer blends on which information was gathered. The different processes will be described in separate sections of this chapter along with more in-depth discussions

661

Polymer Blends Forming

on the behavior of specific blend materials. Blend behavior will first be discussed from a general point of view, looking at the impact of blending on processability, flow-induced morphology and morphology-properties relationships. Then the peculiarities of specific systems will be examined, discussing in more details blends based on commodity, engineering and specialty polymers in the subsequent sections.

Table 10.5.

10.3.2

General Behavior of Blends During Forming

This section aims at discussing the behavior of blends from a general point of view, looking at the impact of blending on processability, flowinduced morphology and the effect of forming on properties.

Forming methods for commodity resin blends: polyolefins-based blends

Blend

LDPE/PIB PE/PIS HDPE/LDPE PE-co-MA, or PE-co-AA/PE PE, PP/α-olefin random copolymer PP/elastomer PP/EPDM PP/EPDM PP/EPDM PP/EPDM PP/EPDM PP/EPDM PP/EPDM +filler PP/EPDM +filler PP/PA-66 or PA-mXD,6 PP/PA PP/elastomer PP/olefin elastomer EPR/PP EPDM/silicone

Trade name

Processes Ext

IM

BM

Ther

Com

Cal

FS

SSF

Cast

Rot

FB

Lupolen 250H Pax-Plus Hostalen GC

N P N

N P N

N D D

P P D

D P D

P P D

D D D

D D D

D D D

D D D

N D

Formion

N

N

D

D

D

D

D

D

D

D

D

Oleflex Moplen SP Hifax Hostalen Vestolen EM Santoprene Keltan TP Kelburon Modylen IPCL Eref Gapex Ferrolene Ferro Flex Polytrope Royaltherm

N N N N N N P P N N N P P N N N

N N N N N N N N N N N N N N N N

N N N N N N D D

N N P D D D D

D N N N N D D

N D P P D N D D

D D D D D D D N

D D D D D D D

D D D D D D D

D D D D D D D

D D D

D D

P D D

N D D

P D D

D D

D D D

D D D

D D

N D

N N D

D D

P D D

D D

P D

D D

D D

Note: Process codes are listed in Table 10.1.

D D

662

Table 10.6.

M. M. Dumoulin

Forming methods for commodity resin blends: styrenics-based blends

Blend

HIPS Styrenics/rubber PS/butadiene rubber PS/butadiene rubber PS-co-PMMA/elastomer HIPS/PPE ABS/PS brominated ABS/PVC ABS/PVC ABS/PA ABS/PBT ABS or AES/PC ABS/PC ABS/TPU ABS/PBT ABS/BA or PB ABS/TPU ABS/elastomer SAN/olefinic elastomer SAN/elastomer SAN/elastomer SAN/PBD SMA/ABS SMA/rubber ASA/PVC ASA/PC ASA/AES/ elastomer ASA/PC ASA/AES

Trade name

Polystyrol Limera Styroblend Styroplus Zylar Vestoran Starflam ABS Lustran Enplex Techniace TA Techniace TB Techniace TC Cycoloy Cycolin Terluran Prevail Lustran Rovel Luran S Lustran Magnum (ABS) Cadon Stapron S Geloy Geloy Centrex Terblend S

Processes Ext

IM

BM

Ther

Com

Cal

FS

SSF

Cast

Rot

N N P N N N

N N N P N N N N N N N N N N N N N N N N N N N N N N N N N

N N P P D D

N D P P D D

P D D D D D

N D P P D D

D D D D D D

D D D D D D

D D D D D D

D D D D D D

D

N

D

D

D

D

D

D

D D D N N D N N D P N D N D P D D

D D N D N D N N N D N N N N D P P N N N

D D D D P D D P D D P D P D P D D

D D D D P D P P D D P D P D D D D

D D D D

D D D D P D D P D D D D P D D D D

D D D D

D D D D

D D D D D D D D D D D D

D D D D D D D D D D D D

D D

D D

D

D D

D D

D D

N D D N P N P N N N N N N N N P N N N N N

10.3.2.1 Processability Blending is increasingly carried out to improve processability. In general, low viscosity is considered beneficial to processing. When immiscible polymers are blended, the blend viscosity rarely follows the additivity rule. A negative deviation

D D

D D D D D D D D D D D D

often opens the way for easier processing. For example, blending ABS with PC improves significantly processability of the latter. Similarly, the viscosity of PVC/TPU blends, particularly those in the middle of the composition range, is significantly lower than that of the pure components and it is also less shear-sensitive, which

663

Polymer Blends Forming

Table 10.7.

Forming methods for blends of commodity resins - Others

Blend

PVC/acrylate elastomer PVC/PMMA PVC/elastomer Acrylic based

Trade name

Processes Ext

IM

BM

Ther

Com

Cal

FS

SSF

Cast

Rot

Vinidur

N

N

N

N

D

N

D

D

D

D

Kydex Oxyblend OxyClear Acrylite, Cyrolite, XT

P N

P N

D N

N N

D N

N P

D D

D D

D D

D D

N

N

D

N

D

D

D

D

D

D

facilitates processing [Das et al., 1992]. Styrenics are added to several engineering resins such as PC and PPE primarily to lower the cost and improve the processability. Addition of a small amount of polyolefin improved the processability and impact properties of several engineering resins [Utracki, 1987b]. The optimum performance has usually been obtained at 2 to 4 wt% level, although up to 20 wt% is frequently claimed in the patent literature [Rosenquist, 1982]. PA-6 or PA-66/PP blends have flow properties similar to that of PP, which constitutes a significant improvement over the neat PA. The PA/PP blends are easier to extrude than PA is, because the melt viscosity varies less with temperature. Another way to reduce viscosity and improve polymer processability is to lower molecular weight. This, however, has detrimental effects on mechanical properties, e.g., modulus or toughness. Fortunately, in many instances, the loss in toughness can be minimized by the addition of a second component at little or no cost in the ease of processing. For example, if optical clarity is not essential, the molecular weight of PC can be reduced, enabling easier processing and the toughness can be maintained by addition of an elastomeric impact modifier [Weihrauch, 1985]. This however requires a good understanding of the structure/property relationships.

Using blending technology to improve processability has been particularly true for recent high performance resins that proved difficult, if not impossible, to process, such as for PPE/PS (Noryl™) blends. PPE was developed in 1959 but because processing temperatures were in the thermal degradation range, its successful commercialization was only possible after blending with PS. Because the two polymers are miscible, melt rheology and processability depend primarily on composition. Commercial PPE compositions, introduced in the 1960’s as Noryl ™ by GE Plastics, could then be processed on conventional equipment. Similarly, processing proved to be difficult for polyetherimide, PEI, and aromatic polyamides, PARA. PEI introduced by GE Plastics in 1982 is not moldable by itself: its processability also relies on blending. PEI moldability was improved by blending it with either polyarylate, PAr, [Holub and Mellinger, 1981; Holub, 1990; Holub and Schlicht, 1990], or with PA-6, using PEI-PA-6 block copolymer as compatibilizer [Robeson and Matzner, 1984] and with PC (Ultem LTX™). PARA’s, developed in the 1960’s, are also difficult to process. For example, spinning of Kevlar™ fibers is carried out from concentrated sulfuric acid solution. Blending PARA with PPS (Ryton™) or PAr [Bapat et al., 1992; Dean, 1992] results in improved processability [Shue and Scoggins, 1981].

664

Table 10.8.

M. M. Dumoulin

Forming methods for blends of engineering resins: PA based

Blend

PA-66-co- 6/ elastomer PA-6/elastomer PA-66/elastomer PA-6/elastomer PA 6,6T/elastomer PA-6/BA copolymers PA-6, PA-66/ elastomer PA-66, PA-6/66/elastomer PA-6, PA-66/elastomer PA-6, PA-66 or 6,12/ionomer or elastomer PA-6 or PA-66/elastomer PA-6, PA-66/elastomer PA-6 or PA-66/PP PA-66 or PA-mXD,6/PP PA-6 or 66/PS brominated PA-6, PA-66/ABS PA/ABS PA-6 or 66/ABS PA-6/aromatic aliphatic polyamide PA-12/aromatic aliphatic polyamide PA-6/PPE PA-666, 610, 12/PTFE

Trade name

Processes Ext

IM

BM

Ther

Com

Cal

FS

SSF

Cast

Rot

Ultramid C

P

N

D

P

D

D

D

D

D

D

Grilon A Ultramid A Ultramid B Ultramid T Durethan

N P P P N

N N N N N

D D D D N

D P P P D

D D D D P

D D D P D

D D D D D

D D D D D

D D D D D

D D D D D

Fiberfil TN

P

N

P

D

P

D

D

D

D

D

Technyl

P

N

D

D

D

D

D

D

D

D

Akuloy

P

N

P

D

P

D

D

D

D

Zytel

N

N

N

D

N

D

D

D

D

Star

N

N

Nybex

N

N

P

D

D

D

D

D

Orgalloy Eref

N N

N N

N D

P D

P D

P D

D

D D

D D

D D

Starflam PA

D

N

Triax 1000 N5 Novalloy-A Grilon

N N P N

N N N N

D D D D

N D P D

D D P D

D D D D

D D D D

D D D D

D D D D

D D D D

Grilamid

N

N

D

D

D

D

D

D

D

D

Dimension

N

N

N

D

D

D

D

D

D

D

RTP B, H, 200

D

N

D

D

D

D

D

D

D

D

665

Polymer Blends Forming

Table 10.9.

Forming methods for blends of engineering resins: Polyester, POM and PPE based blends

Blend

PET/elastomer PET/PTFE PBT/elastomer PBT/elastomer PBT/PS brominated PBT/ABS PBT/ASA PBT/PTFE PCTG/PC or SMA PPE/HIPS PPE/PS PPE/PS PPE/PA POM/elastomer POM/TPU, PTFE POM/TPU

Trade name

Rynite RTP 1100 Ultradur Pibiter Hi Starflam PBT Novalloy-B, Cevian Ultrablend S RTP 1000 Ektar MB Luranyl Noryl Prevex Noryl GTX Delrin Hostaform Ultraform

Processes Ext

IM

BM

Ther

Com

Cal

FS

SSF

Cast

Rot

N D P N

N D D D

D D P D

N D D D

D D D D

D D D D

D D D D

D D D D

D D D D

P

N N N N N N

D

D

P

D

D

D

D

D

P D N N N N N N P P

N N N N N N N N N N

D D D N N N N N D P

P D D N D D D D D P

D D D D D D D N P D

D D D N D D D D D P

D D D D D D D D D D

D D D D D D D D D D

D D D D D D D D D D

D D D D D D D D D D

Table 10.10. Forming methods for blends of engineering resins: PC based Blend

PC/BA-co-MA PC/ABS PC/ABS PC/ABS PC/ABS PC/ABS/SAN PC/SMA PC/PET or PBT PC/PET or PBT PC/PET/HDPE PC/PET/ABS PC/PET, PBT + PB PC/TPU PC/PTFE

Trade name

Makrolon Triax 2000 NovalloyPulse Bayblend T Bayblend Arloy Sabre Xenoy Makroblend Makroblend Stapron E Texin 3000 RTP 300

Processes Ext

IM

BM

Ther

Com

Cal

FS

SSF

Cast

Rot

N N P N N N N P P N N P N D

N N N N N N N N N N N N N N

D D D N N N D P N N N P D D

D N D N D D N P D D D D N D

D D P P P P D P D P P P N D

D D D P D D D P P D D D N D

D D D

D D D P P P D P D D D D D D

D D D

D D D

D D D D D D D D D D

D D D D D D D D D D

D D D D D D D D

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Table 10.11. Forming methods for blends of specialty resins Blend

PES/PEEK PPS/PPE PEI/PC PC/PEI PSF/ABS PSF/PET PSF, PPS, PEI, PEEK/PTFE Nitrile rubber/ PVC, cPE, TPU, PA Crosslinked PE copolymers/ PVC NBR/PVC NBR/EPDM NBR/PVC TPU/PC

Trade name

Processes Ext

IM

P P N P N N D

N N N N N N N

Chemigum TPE

N

Alcryn

Nipol Geolast Krynac NV Texin 4000

Sumiploy DIC PPS Ultem LTX Mindel A Mindel B RTP

BM

Ther

Com

Cal

FS

SSF

Cast

Rot

P P D

D D D

D D

D D D

D D D

D D D

D D D

D D D N D D D

D D D

D D D

D D D

D D D

D D D

D D D

N

D

D

D

N

D

D

D

D

N

N

N

D

D

N

D

D

D

D

N N N N

P N P N

D N D D

D D D N

D D P N

N D N N

D D

D D D D

D D D D

D D D D

D D

In general, blends require tighter control over the process parameters. Processing temperature, for example, has been shown to influence strongly the performance of blends. Using too high processing temperatures was shown to affect the structure and thus to be detrimental [Menges, 1989]. Increasing the melt temperature during injection molding of a thermoplastic/elastomer system resulted in an increase of the relaxation peak of the elastomer, thus a decrease of impact strength. Similar effects were observed when temperature was increased beyond the recommended range, when the dwell time was increased, or when the material was repeatedly reprocessed. Examination of the morphology pointed toward a number of different mechanisms, responsible for these effects, viz., aggregation of dispersed particles in regions of high shear, or destruction

D

of grafting around dispersed particles. In the case of HIPS, it was reported that five successive injection cycles caused reduction of the impact strength of the sample by a factor of two. The microscopic examination indicated that the morphology had been significantly altered: the dispersed particles had shrunk and the elastomer phase around dispersed PS islands seemed to have buckled. 10.3.2.2 Flow-induced Morphology The theoretical treatment of droplet deformation is very complex and most of the work reported is on simpler Newtonian systems. Consequently, rules guiding efforts on morphology control during forming are often qualitative and based on experimental observations.

Polymer Blends Forming

A flow field can induce a number of changes in multiphase systems. These changes will affect either the drops or the whole system. On the one hand, drops can: i) undergo deformation from a spherical shape to ellipsoids or fibrils, ii) break up into smaller ones or iii) coalesce when they collide. Coalescence of PA drops was demonstrated in HDPE/PA-6 blends during capillary flow [Dumoulin at al., 1986]. The few micrometers large PA dispersed drops produced by the compounding stage gave, after capillary extrusion, very long fibrils. The flow-induced deformation of drops into fibers. However, the diameter of the fibrils was of the same order of magnitude as that of the initial drops, indicating that coalescence of the drops in the shear flow had to take place to create the much larger particles. On the other hand, if one looks at the whole system, a number of phenomena can take place: segregation, encapsulation, interlayer slip or variations of miscibility. Components can segregate during the flow, leading to spatial redistribution, encapsulation of one phase by the other being a special case of flow segregation. A redistribution of the phases was observed for HDPE/PA-6 blends during capillary flow [Dumoulin et al., 1986]. At 150°C, the solid PA-6 particles flowed toward the center of the capillary, in the region of lower shear stress. At 250°C, both phases were molten and the lower viscosity PA-6 was found more concentrated at the rim of the extrudate. In both cases, minimization of the total energy of the system is the driving force behind the flow of the lower viscosity component toward the higher stress region. The extrudates showed “tree ring” structures typical of telescopic flow in the capillary, caused by interlayer slip. Figure 10.4 shows a micrograph of an extrudate that was freeze-fractured. The circular patterns can be seen on the surface. In systems with weak interactions between the phases, flow can take place in the form of layers flowing past each other, giving rise to a telescope effect. This phenomenon induces a decrease in viscosity and was described by Lin [1979]. Segregation was also observed for PP/PVAl and HDPE/PA-6 blends [Kamal et al., 1993]. More detailed discussion on the flow-

667

induced morphology can be found in this Handbook, viz. Chapter 7 The Rheology of Polymer Alloys and Blends, and in Chapter 8 Morphology of Polymer Blends.

Figure 10.4. Micrograph of a freeze-fractured surface of an HDPE/PA-6 extruded blend, showing “tree ring” structures typical of telescopic flow in the capillary, caused by interlayer slip [Dumoulin et al., 1986].

10.3.2.3 Effect of Forming on Properties The present section will discuss the main types of morphology observed in polymer blends, the morphology-properties relationships and how, in turn, the forming process affects the performance. Dispersed, lamellar and co-continuous are the three main classes of morphology found in polymer blends. These different blend structures are required for different applications. For example, a lamellar morphology is effective for reduction of permeability, while impact properties can be improved by engendering either a co-continuous structure (useful for large impact energy absorption) or small drops (useful for dissipation of crazes and cracks). Dispersed morphology Toughness is probably the most examined property for polymer blends. Impact resistance depends

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strongly on morphological features. Depending on conditions (i.e. temperature, velocity), three fracture mechanisms can be observed in polymers. Brittle fracture with pure elastic deformation occurs at low temperature or high impact velocity. The two other mechanisms involve plastic deformation of some form: either shear yielding or cavitation. At higher temperature or lower velocity, plastic deformation occurs, causing extensive crazing, accompanied by volume increase. The required fracture energy in this mode depends on the number and size of crazes and is higher than for brittle fracture. At still higher temperature or lower velocity, some polymers exhibit larger deformation between extremely small flaws through shear, either diffuse shear yielding or formation of localized shear bands. The fracture energy is still higher for this type of mechanism. In general, crazing seems to be the principal toughening mechanism but shear yielding also contributes [Bucknall, 1977]. The shear contribution is particularly important for more ductile plastics. Crazes and shear bands are both initiated at rubber particles. In fact the two mechanisms could be synergistic as shear bands act as obstacles to the crazes and keep the latter’s size smaller. To obtain the maximum impact strength from a matrix/dispersed phase blend, the latter phase must be in the form of particles and these should be as round and as uniformly distributed as possible. There is an optimum particle size for toughening. The size dependence might be explained by the fact that small particles appear less efficient in terminating crazes initiated by neighboring particles [Bucknall, 1977]. On the other hand, shear bands can terminate crazes, providing an alternate mechanism, which explains why the critical particle size is smaller in ductile polymers. In ductile polymers, the rubber particles must be large enough to initiate crazes and shear bands but do not need to act as craze terminators since shear bands can play that role. Particle size is controlled by: i) control of reaction conditions in cases where the modifier is solution or suspension-grafted on the resin, such as HIPS or ABS, ii) microrheology during mixing.

The optimum size differs from system to system but is usually in the range 0.5 to 2 µm [Bragaw, 1971]. Increasing rubber particle size from 0.1 to 0.3 µm produced a significant increase in impact resistance in ABS [Parsons and Suck, 1971]. PS generally requires rubber particle size over 1 µm in diameter while it should be less than 0.1 µm for PVC [Bucknall, 1977]. Some PP/EPDM blends showed optimum impact strength when the elastomer particle size was d = 0.5 µm [Speri and Patrick, 1975]. However, this drop size was difficult to achieve, because of the need to match viscosities of the two phases. Thus the high viscosity of EPDM required that high viscosity PP is to be used, making the system unsuitable for injection molding (which is the principal forming method for the blends). Because the elastomer phase is neither grafted to the matrix nor crosslinked (as it is the case for HIPS or ABS), the morphology of these blends is complex and dependent on flow conditions. Consequently, these blends are sensitive to the presence of weld lines [Thamm, 1977]. Strong matrix-rubber adhesion is needed since weakly bonded rubber particles are mechanically equivalent to voids. Voids are much less effective in toughening than rubber particles. Consequently, a poorly bonded particle will not be able to dissipate crazes but could convert the craze into a crack, accelerating failure. Glass beads were observed to induce yielding, stress whitening and higher elongation at break in SAN [Lavengood, 1973]. Mixing NBR with carboxylic acid functionality with PS with and without oxazoline functionality allowed to examine the effects of processing parameters and adhesion on the morphology and impact properties of rubber toughened PS [Liu and Baker, 1991]. Higher impact energy was observed for reactive blends because of increased adhesion. Large stress whitening was observed in reactive blends: crazing occurs in PS phase because of good stress transfer, in areas of high stress concentrations. Thus strong interfacial adhesion is beneficial for effective crazing. In commercial blends, the morphology is normally stabilized by the addition of a compatibilizer or by inducing its formation during reactive processing. Particles grafted to the matrix are less

Polymer Blends Forming

likely to agglomerate in a shear field than when no grafting is present. This is the case for toughened engineering polymers that contain relatively small concentration (10% or less) of a dispersed phase, which exists in the form of small droplets. Co-continuous morphology A different type of morphological structure, cocontinuity of the phases, can prove beneficial by drawing on properties of both constituents: stiffness from a rigid polymer and high strain at break from a ductile one. This is illustrated by Kyu et al. (1991) who studied the toughness enhancement of PC/PMMA blends. The blends were prepared using three different procedures: i) solvent casting from a THF solution and compression molding, resulting in a single phase, ii) melt mixing of the polymer pellets, which yielded a disperse morphology, iii) melt mixing of the single phase blend in the two phase temperature region, causing spinodal decomposition and co-continuous structure. The optimum tensile properties were obtained for the latter structure, which showed ultimate strain equal to the two phase system, significantly superior to the single phase one, and enhanced toughness compared to the other two. A co-continuous morphology was also observed to improve toughness: formation of a network, or honeycomb, structure in which the modifier is actually the continuous phase. Such a structure was observed when chlorinated PE, cPE, or EVA were added to rigid PVC [Menges, 1989]. It was found that extrusion temperature significantly affects the impact properties of these mixtures. This effect originates in the characteristic structure of the suspension-type PVC grains. These particles, with diameter d = 100 µm, were created during polymerization by agglomeration of the “primary particles” (0.1-0.5 µm). At low extrusion temperatures, PVC grains do not disintegrate into primary particles. At higher temperatures, the latter particles merge into a homogeneous material. It was shown that for extrusion temperatures T = 185°C, honeycomb structures were formed, in which the modifier phase surrounds the PVC primary particles and, at a composition of only

669

10%, formed the continuous phase. The impact strength of the compound extruded at the optimum temperature of 185°C was almost 5 times that of the same composition extruded 5°C lower, as shown in Figure 10.5. The best conditions are those that lead to break-up of the large grains, but preserve the primary particles, which enables formation of the honeycomb structure. When the optimum temperature is exceeded, the PVC particles merge, a phase inversion occurs and the honeycomb structure disappears. The modifier is then in the form of droplets but in insufficient number to significantly affect the impact resistance. Producing the right morphology for these ultra high impact strength compositions is process-dependent and therefore can be difficult.

Figure 10.5. Effect of extrusion temperature on structure and impact strength of CPE modified rigid PVC [Menges, 1989].

This also explains why the trend in PVC impact modification has been toward the use of grafted acrylate elastomers. Because of the high shear stability of the elastomer particles (that are responsible for impact modification), these PVC blends are relatively easy to process, i.e., the machine settings can be varied in a broader range.

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Lamellar morphology As mentioned above, a lamellar morphology is effective for reduction of permeability. The effect of blending up to 30 wt% of PA-6, PA-66 or elastomer-modified PA-66 on the oxygen and water vapor permeability of LDPE, LLDPE and HDPE was examined [Kamal et al., 1984]. The data tended towards the series law only above PA loadings of 20 to 25 wt%, indicating the generation of lamellar morphologies at these compositions. To minimize both concentration of the barrier polymer and permeability of the blend, the barrier resin must be distributed as large thin oriented sheets parallel to the container wall or film surface. When this type of morphology was imposed on PE/PA-6 blends, permeability similar to that obtained by coextrusion was observed [Subramanian, 1985; Subramanian, 1987]. Figure 10.6 shows hydrocarbon permeability for HDPE containers [Subramanian, 1990]. The lamellar morphology shows significant improvement in permeability over the pure HDPE or the blend with droplet dispersion. The permeability is strongly reduced by the presence of PA lamellae. The platelet-type dispersion, responsible for the high efficiency of the barrier phase, is obtained under specific processing conditions, in particular under low shear conditions. Intensive melt compounding and dispersive mixing destroy the platelets and reduce the effectiveness of the PA. Figure 10.7 shows an example of the morphology of lamellar blends [Champagne et al., 1995]. The lamellar morphology was here generated through blow molding. Similarly, a common method for generating these structures is film blowing, as it will be discussed in later sections. Barger et al. [1993] demonstrated the production of lamellar morphologies by co-injection, using a special three layer feed-block and a layer generation stage. Layer multipliers subdivided the melt stream into four sub-streams, which then were reoriented, stacked and recombined. The process produced net shape parts with micron size layers. Lamellar morphologies can also be obtained through extrusion by using special die geometries [Kamal et al., 1993]. Holsti-Miettinen et al. [1995] examined the permeability and mechanical properties of PP/PA-6 blends com-

Figure 10.6. Hydrocarbon permeability of HDPE containers in terms of hydrocarbon weight loss as a function of time. a: HDPE, b: HDPE/PA = 80:20, the PA being in the form of dispersed domains, c: HDPE/PA = 80:20, the PA being in the form of lamellae [Subramanian, 1990].

Figure 10.7. SEM micrograph of a cross-section of an HDPE/PA-6 blow molded bottle [Champagne et al., 1995].

pounded in a twin screw extruder, then injection molded. They found that, under some conditions, injection molding could induce the formation of lamellar morphology, in particular in the skin region. These parts yielded the largest reduction in oxygen permeability and also showed significant increases in tensile modulus, strength and impact strength.

Polymer Blends Forming

10.3.3

Forming Blends of Commodity Polymers

This section aims at providing a more in-depth look at forming for some specific commercial blend systems on which information could be gathered. Forming requirements and/or advantages are given for each type of blend. 10.3.3.1 Polyolefins PE blends HDPE/LDPE blends marketed under the trade name Hostalen™ GC by Hoechst, are used in injection molding (grade GC 7260), for bottle crates in particular, and in extrusion (grade GC 7260 F) for paper coating, fibers and films for typewriter ribbons. In injection molding, the blends show low warping tendency, high density, hardness and stiffness, good impact and stress cracking resistance. In extrusion coating, the blend contributes to very high drying rates of emulsion coating due to its barrier properties. The reverse side coating prevents curling effect of the face coating. Elastomer modified PP Because of its glass transition temperature around 0°C, PP shows brittle fracture in the sub-zero range. Therefore, most PP grades currently produced include an elastomeric phase in order to improve impact resistance, particularly at low temperature. The elastomer is usually of the polyolefin type. In recent years, the elastomer is being introduced at the reactor stage, by adding comonomers during polymerization. Some PP/ EPR systems showed optimum impact strength for an elastomer particle size of approximately 0.5 µm [Speri and Patrick, 1975]. Such a degree of dispersion is difficult to achieve when the elastomer phase is compounded in. One method is to match viscosities of the two phases as closely as possible. The high viscosity of EPR requires the use of a high viscosity PP, making the blend less suitable for injection molding, which is the principal forming method for these systems. Because the elastomer phase is neither grafted to the matrix

671

nor crosslinked (as is the case for HIPS or ABS), the blends’ morphology is complex and sensitive to flow conditions. Consequently, during injection molding, these blends are prone to show weak weld lines [Thamm, 1977]. Elastomer modified PP (Hostalen™) shows very high toughness, particularly at low temperature, and excellent acoustic properties. The material is injection molded into a wide variety of automotive exterior (bumpers, trim, spoilers), interior (steering wheels) and engine compartment applications (battery casings, covers, ducts). It is also extruded into pipes and ducts for domestic and automotive engine compartment uses. In PP/EPDM systems (Vestolen™ EM), advantage has been taken from the combination of crystalline and amorphous polymer properties. The blends are formulated by selecting molecular weights of the components and concentration of the elastic segments in the EPDM. Some of these blends are prepared by dynamic vulcanization, during which molten EPDM phase is dispersed then progressively vulcanized. When the degree of vulcanization is high the EPDM particles may be broken into small elastomeric filler particles, physically bound to the matrix. Therefore, since the concentration of EPDM as well as the extent of its vulcanization can vary, a great variety of morphologies, thus properties, can be obtained. A co-continuous morphology is observed for some PP/EPDM systems (e.g., Santoprene™) and the melt is highly viscous. Montell markets two families of elastomer modified PP: Hifax is the trade name for PP/EPDM formulations, while the Moplen SP series also includes elastomer-modified, reinforced and nonreinforced grades. Blends from both series show a dispersed morphology. Their properties include good paint adhesion, low temperature ductility, wide viscosity range and good thermal oxidative stability. They have been developed for automotive applications and are formed by injection (grilles, body panels, trim, spoilers, dashboards,...) and by extrusion/blow molding primarily into ducts. Non-automotive applications include roofing.

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EPDM/silicone The advantages of EPDM/silicone (Royaltherm™) blends are better heat resistance than EPDM and better mechanical properties than silicone, in particular tensile strength and flex life, in addition to flame retardancy. The blend is being formed either by injection molding or by extrusion for the use in diverse applications: radiator hose, wire, cable, sponge, rubber grips, vibration damping, gaskets, seals. Royaltherm™ may be compounded with fillers such as silica, calcium carbonate, talc, clay and barytes. 1-5 PHR of carbon black will improve weathering, but large amounts should be avoided because, unlike silica, no particular advantage is obtained. PE-MA or PE-AA /PE Blends of ethylene-methacrylic acid or ethyleneacrylic acid copolymers with PE (Formion™) show either dispersed or co-continuous morphologies. They have outstanding fuel and oil resistance, adhesion to glass and metals, and ease of foaming. The blends are injection molded into bumper end caps, bumper rub strips, fender stone guard, extruded into products such as carpet heel pad and window trim. Moldings can also be foamed for bumper guards. 10.3.3.2 Styrenics Polymers derived from styrene are amorphous with Tg ranging from 100 to 130°C and processing temperatures from 160 to 200°C. The unmodified polymers are characterized by ease of processing but also by high brittleness. Consequently, they are often modified by addition of an elastomer phase (viz., HIPS or ABS). The polybutadiene phase, in levels of at least 5 wt%, is added during polymerization to result in a dispersion of small (0.1-5 µm) particles. The graft copolymers with sub-micron rubber particles are transparent, but more brittle than those with larger rubber particles (e.g., about 1 µm in diameter). HIPS and ABS, which can be considered as blends in their own right, are frequent component of blends with other resins.

HIPS High impact PS, HIPS, a dispersion of poly(butadiene-g-styrene) in polystyrene matrix, has been commercialized under a number a trade names, e.g., Edistir™, Lacqrene™, Lastirol™, Polystyrol™, Styron™, Styroplus™, Styroblend™ or Vestyron™. During processing, PS can degrade and emit styrene in the atmosphere. The amount of degradation depends on processing conditions. Since styrene affects the nervous system when its concentration exceeds TLV = 20 ppm, good ventilation is required. Among the critical factors affecting product performance are processing (i.e., mold and barrel) temperatures and the injection rate. The chemical resistance, especially environmental stress cracking resistance (ESCR), is more affected by processing than by chemical composition. As a general rule, the materials that have greater orientation exhibit higher resistance to environmental stress cracking. This illustrates how important is the temperature control and the molding rate. Limera™ is a family of styrenic thermoplastic/ rubber compositions. The compositions are suitable for compounding with a variety of additives, fillers and polymers such as PS, PVC, PMMA, ABS and PPE. They are processable by injection molding, extrusion and blow molding. PS-PMMA/elastomer Elastomer modified styrene-acrylic copolymer, PS-PMMA, (Zylar™) display high clarity and toughness superior to clear ABS and acrylics. The material is formed by injection molding or extrusion and used in parts for blenders, food processors, diskette boxes, face shields, glasses and tool handles. HIPS/PPE Composition of commercial HIPS/PPE blends range from 25/75 to 60/40. Because of the miscibility of PS/PPE system, the blends’ properties linearly depend on composition. The melt rheology, thus processability is also dominated by composition. In most cases standard processing equipment can be used. Historically, HIPS/PPE

Polymer Blends Forming

blends were first formed by injection molding. This forming method still remains most widely used, with blow molding and extrusion becoming increasingly popular. Reinforced grades of HIPS/ PPE blends are available, e.g., Noryl GFN, as well as formulations for structural foam (Noryl™FN, Vestoran™). ABS ABS itself is an acronym for a wide variety of blends comprising styrene-acrylonitrile copolymer, SAN, toughened by incorporation of a grafted elastomer. PB has generally been used as the elastomer. The PB phase is finely dispersed and chemically grafted to SAN during the copolymerization of styrene and acrylonitrile. Emulsion polymerization allows production of grades containing high levels of rubber (25 to 50 wt%) while bulk polymerization produces low contents (up to 10%). ABS (e.g., Magnum™ from Dow) exhibits good processability: good thermal stability, mold release characteristics and easy flow. The high HDT of ABS allows parts to be used in higher temperature applications than for HIPS. High HDT also contributes to shorter molding cycles, because parts set up faster, making it possible for the mold to be opened and the part to be ejected without distortion sooner. The excellent thermal stability of ABS allows the parts to hold their original color and physical properties over a wide range of processing conditions (218 to 274°C). Since ABS is amorphous, it can be processed with good dimensional tolerances and low warpage. Numerous modifications of ABS are known. For example, replacing part of acrylonitrile with methylacrylate led to a modified version, MABS, advantageous when blended with PVC [Wilsher, 1979]. A new and technologically important class of acrylonitrile-styrene-acrylate, or ASA (e.g., Centrex™ from Monsanto, Geloy™ from General Electric, Luran™S from BASF), was discovered when PB was replaced with saturated acrylate elastomer. Replacing PB by chlorinated polyethylene led to acrylonitrile-chlorinated polyethylene-styrene, or ACS, terpolymer. Replacing PB by EPDM led to acrylonitrile-ethylene rubber-

673

styrene, or AES copolymer (e.g., Hostyren™XS, Novodur™ AES, Rovel™). The terpolymers ASA, ACS, and AES have better oxidative degradation resistance and weatherability. A new ABS-type copolymer suitable for reactive blending with polyesters and polyamines was introduced in 1982 by Daicel Chemical Industries. The material, comprising di-olefins, vinyl aromatics, unsaturated carboxylic acid and acrylonitriles, ABSMA, is prepared by dissolving SBR in the monomer mixture, then polymerizing it. ABS/PC During processing, ABS/PC blends (Techniace™ TC from Sumitomo Dow) behave similarly to ABS. These blends fill a need for a combination of high flow and high modulus required for molding thinner walls. For example, in the office product industry, the move from installed or laptop computers to notebook types demands significant downsizing and weight reduction without loss of strength. Thus, the laptop wall thickness is typically 2.2-2.5 mm while that of a notebook, is 1.0-1.8 mm. ABS/TPU ABS/TPU blends are used mostly for automotive parts, e.g., paintable, soft bumper fascias. They are mostly formed by injection molding or extrusion. ASA/PVC Acrylic-styrene-acrylonitrile terpolymers, ASA, are produced by simultaneous polymerization of styrene and acrylonitrile monomers in the presence of an acrylic rubber (polybutylacrylate). Thus, the material has a two-phase structure similar to ABS. However the elastomer phase in ASA is saturated and thus significantly more resistant to oxidative degradation. ASA is used in applications requiring good weatherability, mostly with PVC. The blends are primarily extruded for exterior trims and window profiles applications. SAN/elastomer The processing temperature should be well controlled for elastomer modified SAN (Luran™ S). Long exposure to 260°C may cause degradation.

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Since Luran™ S is immiscible with most thermoplastics, mixing it with other plastics should be avoided. However, good performance can be obtained by blending it with PVC, provided that a high dispersion is obtained. SMA (Cadon™, DYGL™, Dylark,™ Stapron,™) Random copolymers of styrene with 10 to 30 wt% maleic anhydride are amorphous, with Tg higher than that of PS by 20 to 60°C. Like PS, their brittleness makes it necessary to incorporate an elastomer for many applications. Toughening is usually accomplished by blending in styrenegrafted PB (like HIPS) or by melt blending with SBR. SMA is a component of many high temperature resistant blends, viz., PC/SMA/HIPS, Arloy™ from Amoco. The terpolymers of styrene-maleic anhydride-methylmethacrylate, SMA-MMA, also show interesting properties, e.g., in blends with styrenics such as the commercial SMA-MMA/ ABS blends (Cadon™), or with PVC. 10.3.3.3 Vinyls Owing to its versatility of application and low cost, PVC is one of the most widely used thermoplastics, and by far the most important member of the vinyl family. However, PVC is thermally unstable and difficult to process. To improve processability, the resin must be plasticized, stabilized and processed under strictly controlled conditions. The necessity to improve processability by incorporation of a “permanent plasticizer” was the main reason for blending PVC, in addition to the need to increase its HDT and impact strength. Toughness of PVC has been improved by plasticizing it with low molecular weight plasticizers, such as octyl phthalates, adipates or sebacates, but this method not only decreases the modulus, but also leads to materials that are prone to “sweat out” the plasticizer and progressively becoming brittle. By contrast, the use of elastomers as impact modifiers leads to blends characterized by stable, improved toughness, and good rigidity. Several elastomeric modifiers have been specifically designed for PVC. Some of these, for applications such as packaging films or sheets and

clear bottles, toughen the resin without affecting its optical clarity. These modifiers are either designed to match the refractive index of PVC or have small particle size (0.1-2.0 mm). A number of PVC blends (viz., PVC/EVA, PVC/NBR, PVC/ABS) have been used for extrusion of window profiles, wires and cables coating, or for water hoses. PVC/elastomer Oxychem markets two lines of impact modified PVC: Oxyblend™ and OxyClear™, in which an elastomeric phase is dispersed in the PVC matrix. OxyClear™ is a series of clear, high impact, UV stabilized compositions developed particularly for blow molding bottles for cosmetics, chemicals, foodstuff. The OxyBlend™ series is designed for injection molding and extrusion of housing applications such as gutters and window profiles. PVC/ABS 10 to 40 wt% ABS is usually added to PVC to improve its processability and impact resistance. Because of thermal degradation, the blends have to be prepared under well controlled conditions. To avoid repeated heating and cooling cycles, the blends are frequently compounded and formed by the processor in a single step operation. Sheets of PVC/ABS blends are often used for thermoforming large parts, e.g., for automotive or airplane interior panels. The blends have also been molded and extruded as replacements for higher cost flame-retarded ABS. PVC/PMMA PVC/PMMA blends (e.g., Kydex™, Royalite™) in which PMMA is dispersed in the PVC matrix, show deep formability, rigidity, breakage resistance, chemical and abrasion resistance. They are easily thermoformed, thus used for manufacture of thermoformed chairs, seats and aircraft interior panels. They are also formed by calendering and, to a lesser extent, by injection molding and extrusion. Applications: wall covering, ceiling tiles, shelving and counter laminates, telephone areas, ...

Polymer Blends Forming

10.3.3.4 Acrylics Acrylic resins include a wide range of polymers and copolymers, PMMA being the most common one (viz., Acrifix™, Acrylite ™, Acryrex™, Altuglas™, Bonoplex™, Deglas™, Lucite™, Perspex™, Plexiglas™, Resartglas™, Shinkolite™, Sumipex™, Vedril™, etc.). PMMA is amorphous, optically clear and UV resistant, thus widely used for glazing. For this reason, the products from PMMA are often formed in two steps: extrusion followed by thermoforming. Impactmodified grades, usually obtained by blending with acrylic elastomers, are also available. The transparency is preserved by keeping the elastomer particle size not larger than 0.1 mm. These blends are mainly formed by molding, for such applications as outdoor signs or automotive light housings. 10.3.4

Forming Blends of Engineering Polymers

10.3.4.1 Polyamides PA’s are semi-crystalline polymers with strong amide links, responsible for high interchain hydrogen bonding. The resins have high melting point and excellent mechanical strength, but are sensitive to moisture. The amide linkage is produced by condensation of amine and acid groups, with water as a by-product. The moisture sensitivity of PA originates from the reduction of strength of the intermolecular hydrogen bonding caused by the presence of water, leading to a decrease of the resin mechanical properties, viz., strength and modulus. Furthermore, water may also displace the chemical equilibrium of the condensation reaction, thus induce a drop in molecular weight. The two most widely used PA’s are PA-66, a condensation product of hexamethylene diamine with adipic acid, and PA-6, obtained by ring opening polymerization of ε-caprolactam. Both are primarily used for fiber spinning, with a portion used as films or moldings. Others, more recently developed PA’s include: PA-6,12, PA-11, PA-12, PA-12,12, PA-4, and aromatic polymers

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obtained from condensation of aliphatic or aromatic amines with phthalic or naphthoic acids. Blends of PA’s with most thermoplastics show coarse morphology indicative of large interfacial tension. To obtain reasonable performance from the PA blends, these must be compatibilized. The most popular method of PA compatibilization is by two step reactive extrusion. In the first step, the compatibilizer is prepared by condensing a polymer bearing functionalities reactive with the amine terminal groups of a low molecular weight polyamide, and the second by blending the principal two components of the blends with the compatibilizer. For example, PA/PP blends (Orgalloy™) are compatibilized by incorporation of a low molecular weight PA adduct with maleated PP. PA/ABS blends (Triax™ 1000) are compatibilized by incorporation of acrylamide and maleic anhydride comonomers during the preparation of ABS. Introduction of acrylamide increases polarity and hydrogen bonding while the maleic anhydride reacts with the terminal amine groups of PA. To lower PA’s sensitivity to crack propagation under impact conditions, the resins are blended with 10-25% of an elastomer or a polyolefin. The impact-modified PA’s exhibit good processability. They are mostly formed by injection molding. Special compositions have been formulated for extrusion of tubings. Several PA/PA blends (viz., Bexloy™ C, Eref™, Grilon™ BT, Nyref™, Ultramid™) have been developed, particularly those comprising high-Tg PA with PA-6 or PA-66. Two uses have been envisaged: (i) preparation of yarns to reduce flatspotting in tires, and (ii) manufacture of films having improved barrier properties for packaging applications. PA-6, PA-66, or PA-612 with ionomer, elastomer, or fluorinated polymer (Zytel™) The type of polyamide determines the blend melting and softening temperatures, and therefore forming range. Table 10.12 lists melting temperatures for a number of polyamides. Addition of a toughening or impact modifying resin has little effect on the melting temperature.

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Table 10.12. Melting temperatures of selected polyamides and blends PA type

Trade name

Tm (°C)

PA 6 PA 66 PA 612 PA-12/aromatic aliphatic polyamide PA-6/BR or EPR elastomer PA-6/PPE

Zytel Zytel Vestamid D Grilamid Grilon A Dimension

210 255 212 178 220 220

On the other hand, the amount of heat required for processing depends on the crystalline structure and content of material, thus it depends on the resin grade and blend composition. Heat of fusion of crystalline PA-66 is 130 kJ/kg (equal to LDPE but only half of that reported for HDPE). It should be noted that the heat of fusion almost doubles the heat required for processing of crystalline resins in comparison with that necessary to process amorphous ones. Copolyamide-66 and 6/elastomer Ultramid™C blends consist of rubber particles finely dispersed in the PA matrix. The blends are marketed in 15% and 30% mineral filled grades. They are used particularly in injection molded sports and leisure goods such as helmets, roller skate chassis and automotive parts such as car wheel covers. D&S International markets a family of polyamide-based engineering blends: Dexpro™ (PA/PP) and Dexlon™ (PA/PP) in which the PP is maleated, and Dexcarb (PA/PC) in which a polyetheramide and a maleated PP or EPR are used as compatibilizer/impact modifier additives.

PA/ABS The main reason for incorporating semi-crystalline engineering resins, such as PA or PBT, is to improve the chemical resistance of ABS, without

loosing its excellent processability and impact strength. The PA/ABS blends (e.g., Alphaloy ™MPA, Elemid,™ Diaaloy ™N, Kane ace MUH, Malecca™ S, Maxloy ™A, Monkalloy ™N, N5, Novaloy-™A, Stapron™ N, SX, Techniace™ TA, Toyolac™Alloy, Triax™ 1000, Ultramid™) must be compatibilized. Here, polyamide, PA, is either a matrix, or part of a co-continuous structure. The blends have been formulated for injection molding, extrusion, or thermoforming, but blow and compression molding can also be used. They show good processability and flow, high heat and chemical resistance, high resistance to oil, wear, and abrasion, dimensional stability, low temperature impact strength, reduced moisture sensitivity and economy. The main applications are in automotive, chemical, electrical, consumer and sport industries. Some grades are glass fiber reinforced. The material has also been used in vibration damping structures. 10.3.4.2 Polyesters PET and PBT are the most widely used thermoplastic polyesters. Two other well known polyesters are: (i) polyethyleterephthalateglycol, PETG, a copolymer of terephthalic acid with 66 mol% of ethylene glycol and 34 mol% of cyclohexylene dimethanol, and (ii) poly(1,4-cyclohexanedimethylene terephthalate), PCT. Both PET and PBT are semi-crystalline with high melting point, Tm = 265 and 225°C respectively. Most PET is used for fiber spinning, with only a small portion dedicated to thermoplastic applications, viz., blow molding of soft drink bottles, injection molding of engineering parts and thermoforming of the so-called “dual-ovenable” trays. PET’s rather low crystallization rates (much lower than those observed for polyamides) has hampered its use in injection molding applications, but development of fast crystallizing grades, containing nucleating agents and/or blends, has opened new opportunities. PBT, on the other hand, exhibits significantly faster crystallization rate, which results in shorter molding cycles. This combined with lower viscosity has resulted in better acceptance of PBT than PET for molding applications.

Polymer Blends Forming

Molten polyesters show low viscosity and small extrudate swell. For these reasons, they have been blended with amorphous polymers to improve the latter’s processability and chemical resistance. Elastomers have also been added to polyesters to improve impact resistance. Blends of polyester (either PET or PBT) with polycarbonate, PC, are the most popular (viz., Bayfol™, BCT 4201, Calibre™, Dialoy ™P, Ektar™ MB (with PCTG), Idemitsu™ SC, Makroblend™, MB4300, Novadol™, Pocan™, R2-9000, Sabre™, SC 600, Stapron™ E, Ultrablend™ KR, Valox™, Xenoy™ 1000, etc.). Presence of PC in PET/PC increases crystallization rate of PET, which translates into faster injection molding cycle and lower part distortion upon demolding than those observed for neat PET. PET/PBT (e.g., C, CN, Celanex™, EMC, Valox™) Advantages of PET/PBT blends are fast crystallization of PBT combined with lower cost of PET. The amorphous phases are miscible, yielding acceptable mechanical properties without any compatibilizer. In fact, an ester-ester exchange reaction, or transesterification, takes place during compounding and processing between PET and PBT, creating PET/PBT copolymers, thus compatibilizing the system. However, since the copolymer is amorphous the extent of transesterification must be carefully adjusted. PET/PBT blends give good surface appearance and gloss and are used largely for injection molded appliances and automotive components. PBT/elastomers (e.g., Arloy™ 2000, Celanex™, Rynite™, SC-150, Bexloy™ J, Pibiter™ Hi, Makroblend™, Techster™ T) PBT shows low Izod impact strength and consequently must be impact modified — core-shell type elastomers are most frequently added. The blends have been used for injection molding exterior automotive parts such as bumpers and mirror housings. These blends show low melt viscosity, thus are easy to process. However, owing to the semi-crystalline nature of PBT, the moldings have high mold shrinkage and tendency to warp — process optimization and part design must take these into account.

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PBT has also been blended with copolyether-ester thermoplastic elastomer in which the hard segment is composed of the polybutyleneterephthalate chain. The thermoplastic elastomer component brings about the typical elastomeric behavior while addition of PBT reduces the cost. Elastomerrich blends found application for injection molding flexible exterior automotive components. On the other hand, PBT-rich blends (Celanex™) are used for molding parts requiring high impact strength in engineering resin applications. PET/elastomer blends have been used commercially either filled or reinforced. Owing to slow crystallization of PET in these blends, unfilled grades would be very difficult to mold. In warm molds, the cycle would be too long, the parts would tend to stick and distort when removed, making the process expensive. In cold molds, amorphous parts with low heat deflection temperatures, would be produced, reducing the usefulness of the blends. PCTG/PC or SMA (Ektar™ MB) Blends of PCTG with PC or SMA show good clarity, good impact and chemical resistance and excellent forming characteristics. The material is formed by injection molding and extrusion and is used for business machines and appliance parts, vertical body panels for cars, trays and tote bins, lawn equipment housings. 10.3.4.3 Polycarbonate PC/PA (e.g., Dexcarb™, Iupilon™, SC 720) PC and PA are immiscible and require compatibilization. The blends show good processability, high HDT (HDT > 200°C) and impact strength, low mold shrinkage, solvent and moisture sensitivity. These relatively new materials have not being widely used. PC/PET or PBT (BCT 4201, Dialoy™ P, Ektar™ MB (with PCTG), Idemitsu SC, Makroblend™, MB4300, Novadol™, Pocan™, R2-9000, Sabre™, SC 600.) The blends have been formulated mainly for injection molding, extrusion and blow molding.

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However, they can also be transfer or compression molded, solid state formed or thermoformed. PC is the matrix phase in most of these blends. The impact modifier can be PE, ABS or acrylic copolymer, e.g., MBS. The blends show good processability, heat resistance, ductility, HDT, high modulus, impact, tensile and flexural strength over a wide temperature range, good adhesion, solvent, chemical, and UV resistance. They can be painted, hot stamped, metallized and plated. Some blends (containing PB as impact modifier) may have poor weatherability. The blends are mainly used in the automotive industry. PC/ABS (Alphaloy™, Bayblend™ T, Cycoloy™, Denka™ HS, Dialoy™ C, Exceloy™, Iupilon™, Lynex™ B, Malecca™ P, MB2200, Multilon™, Novaloy™ S, Proloy™/Lexan™, Pulse™, PX, Ryulex™, SC 250, Techniace™ TC, Toplex™, Triax™ 2000, ABS/PC 7901) Blending improves the flowability of PC, and results in low mold shrinkage and warpage. The blends have been formulated mainly for injection molding, extrusion and thermoforming or vacuum forming, but there are grades also suitable for compression or transfer molding. PC/ABS blends show good processability, excellent plating, high HDT and impact strength, high stiffness and strength, dimensional stability, toughness, delamination resistance, low temperature impact strength, good weldline strength, and solvent and chemical resistance. These are three-phase blends, with 30-65 wt% PC usually being the matrix. The blends are used for computer housings, business machines and electrical appliances, connectors, carriages, switches, fans, power and agricultural, garden and lawn tools, cameras, optical instruments, exterior automotive components, electrical, electronic or telecommunication parts, sporting goods, etc. Molding of PC/ABS blends takes place at temperatures close to those recommended for ABS. Grades with glass fiber reinforcement are also available.

10.3.4.4 Polyoxymethylene (acetal) POM/elastomer (Celcon™, Delrin™, Duraloy™, Formaldafil™, Fulton™ KL, Hostaform™ S, RTP 800, TC, Thermocomp™, Ultraform™) Due to its high crystallinity, polyacetal is more rigid and stronger than other thermoplastics especially at elevated T (50-120°C). Because of its semi-crystalline character, POM requires a greater heat input for melting than amorphous resins. It has a fairly high heat of fusion (160-200 kJ/kg) but because of lower melt temperature (210°C), it requires considerably lower total heat energy than other crystalline polymers (viz., POM = 435-475, HDPE = 720, PA-66 = 756-786 kJ/kg). As other semi-crystalline resins, POM benefits from toughening. The most common elastomer used for this task is TPU. Most blends contain 10-30 wt% TPU. For example, Ultraform™ blends consist of TPU particles dispersed in POM obtained by mechanical blending. For sliding applications there are grades lubricated by addition of PTFE. These materials are formulated for extrusion, injection, compression, blow and transfer molding. They show excellent processability, rigidity, high impact strength, high fatigue, flexural and tensile strength, high toughness, low water absorption, gloss, resistance to chemicals. Co-continuous morphology gives especially good performance. These blends find use in sports equipment, plumbing, electronic/ mechanical parts, automotive, appliances, hinges, etc. A family of acetal based blends (Hostaform™), used in injection molded engineering parts, include grades based on POM/TPU elastomer (Hostaform™ S) and on POM/PTFE (Hostaform ™C 9021 TF). The material shows high tensile and impact strength, high toughness, very good fatigue resistance, glossy surface and good solvent resistance. The blends are also processed by extrusion and compression molding. 10.3.4.5 Polyphenylene Ether PPE is an amorphous polymer with a high glass transition temperature, Tg = 210°C. Since the processing temperature ought to be at least 60°C

Polymer Blends Forming

above Tg, the resin is notoriously difficult to process since PPE degrades readily above 230°C. The PPE homopolymer was introduced by GE in 1964. A year later Westlake Plastics Co. started production of Alphalux™ 400, extruded rods, sheets and tubing of PPE. The same year Richardson Co. began manufacturing miscible blends of PPE with crystal PS, but the single-phase blends were not a commercial success. PPE/HIPS (Gepax™, Luranyl™, Noryl™, Prevex™, Verton™, Vestoran™, Xyron™) In mid-1966 GE introduced Noryl™, a family of PPE blends with HIPS [Boldebuck, 1962; Cizek, 1968]. A year later, glass-reinforced and flameretardant grades were introduced. After expiration of the original GE patent on PPE, several manufacturers (e.g., BASF, Hüls, Borg-Warner, Asahi, Engineering Plastics Ltd., Mitsubishi) began selling their own blends based on PPE or its copolymers. In 1979 Asahi-Dow introduced styrenegrafted PPE blend, Xyron™ [Izawa et al., 1973; Izawa, 1992]. In 1984 Borg-Warner started production of Prevex™, a copolymer obtained by oxidative coupling of 2,4-dimethyl phenol with 5 to 10% of 2,4,6-trimethyl phenol [Grant et al., 1990; Bartges, 1990]. Blends of PPE with most styrenics, viz., HIPS, SBS, SEBS, ABS, SB, etc. have been commercialized. The PPE/HIPS blends, with 25-60 wt% PPE, have been designed for injection or blow molding, calendering, thermoforming and extrusion. Some grades are reinforced with < 30 wt% glass fibers. They show good processability, heat resistance (HDT = 90-150°C), toughness, good dimensional stability, resistance to hot water, flame retardancy, low density, cost/performance ratio and low moisture absorption. PPE/PA (Artley™, Dimension™, Lynex™ A, Noryl™ GTX, Remaloy™, Ultranyl™) PPE has been blended with semi-crystalline polymers, PA in particular, in order to improve chemical resistance. The success of PPE/PA blends depends on adequate compatibilization. Because the PPE/PA blends show low value of notched impact strength, toughening by addition of an

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elastomer is required. The blends have been designed for extrusion, thermoforming or injection molding. The compatibilized blends usually contain 40-60 wt% PA with the PPE dispersed phase acting as an organic filler. These materials show moderate processability and impact strength, good high temperature creep and tensile and flexural strength, as well as solvent and chemical resistance and low moisture absorption. They are dimensionally stable, paintable and platable. They are being used in a number of automotive components and are formed mostly by injection molding. PPE/PBT (Dialoy™ X, BE, Gemax™, Iupi-ace™) Polyesters such as PBT or PET offer to PPE blends similar advantages as PA (chemical resistance because of semi-crystalline nature) but with lower moisture sensitivity. Commercial compatibilized PPE/polyester blends normally contain an elastomer and PC for improved shear stability. The blends show excellent processability, high solvent and temperature resistance as well as dimensional stability. 10.3.5

Forming Blends of Specialty Polymers

Specialty polymers are generally high price resins with either some specific properties or performance superior to engineering resins. They are high Tg amorphous resins or high Tm semi-crystalline polymers. Their common feature is an aromatic backbone chain. Amorphous specialty resins include polysulfones (PSF), polyimides (PI), polyetherimides (PEI), polyamideimides (PAI) and polyarylates (PAr), while semi-crystalline ones include liquid crystal polymers (LCP), polyetheretherketone (PEEK) and polyphenylenesulfide (PPS). Blends based on specialty resins that are easily formed by conventional thermoplastic techniques include PEI/PC (Ultem LTX), PPS/PPE (DIC™ PPS), PEEK/LCP (Sumiploy™ EK) and PEEK/PES (Sumiploy™ SK). PEEK PEEK-based blends include grades with polyethersulfone, PES (SumiployTM SK1800 and SK1660) where SK1800 comprises PES dispersed in PEEK

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and SK1660 is glass fiber reinforced SK1800. Advantages of Sumiploy™ SK compared to PEEK or glass fiber filled PEEK are low temperature impact strength and good processability. The blends are processed mainly by injection molding. Extrusion, compression or transfer molding are also possible but rarely used. PSF/ABS (Mindel™ A) PSF/ABS blends show a dispersed morphology. The advantages of the blend are high HDT (it can withstand temperature about 50°C higher than ABS), high toughness, dimensional stability and hot water resistance. It is formed by injection molding and extrusion and used in plumbing, food service and fiber optics systems. PSF/PET (Mindel™ B) PSF/PET blends show a dispersed morphology. The combination of crystalline PET and amorphous polysulfone provides chemical resistance and warp-free properties. The amount of crystalline polymer is varied to meet requirements in thermal properties and the level of reinforcement is varied to tailor the modulus. The blends have electrical and mechanical properties similar to PET but only a third of its shrinkage and warpage. Also, stress crack resistance to common solvents is improved. Similarly, the service T is upgraded compared to PET. The blends are formed by injection molding and extrusion. Lubricated compositions (Lubricomp™, RTP) In plastics, wear depends on adhesion and deformation of soft material. Frictional forces are proportional to velocity rather than load as in the case of metals. A number of additives can improve wear and slip properties. PTFE has lower friction coefficient than any other material. Additionally, its particles form a film on shearing surfaces thus decreasing friction. Its addition is very effective in high-pressure applications. Optimum additions are 20% in crystalline polymers and 15% in amorphous ones. Molybdenum disulfide is primarily used in PA compounds. It works as a nucleating agent, promoting increased crystallinity in PA, thus providing harder, more wear

resistant surfaces. It also has high affinity to metal surfaces, filling pores and rendering the metal surfaces more slippery. In the case of graphite, its lattice structure allows its molecules to slide past one another thus decreasing friction. This property is emphasized in aqueous environment, making graphite extremely useful in underwater applications. Silicon fluid migrates to the surface because of difference in free energy and forms a slippery film on the surface. Glass fibers, if well bound to matrix, improves wear resistance because of limiting creep and heat deformation and increasing thermal conductivity. Carbon fibers act similarly but additionally are softer and thus unable to score the surface of iron. Aramid is much less abrasive than glass or carbon fibers, essential if the mating surface is sensitive to abrasion. RTP produces series of engineering thermoplastics lubricated with PTFE, molybdenum disulfide, graphite powder, silicone and glass, carbon and aramid fibers. Grades AR, TFE and SI include respectively aramid fibers, PTFE (10-20%) and silicone fluid (0.5%). Matrix materials used include PA-6, PA-66, PA-6,10, PA-12, PC, POM, PET, PBT, PSF, PPS, PEI, PEEK and polyphthalamide. The blends have been developed for injection molding of various parts for engineering applications. LNP developed a similar family (Lubricomp™) in which PTFE has custom-tailored MW to serve the purpose of formation of lubricity film over the mating surface under shear. Matrix materials include PE, Poly(ethylene-co-chloroethylene), PP, PS, ABS, SAN, PA-6, PA-66, PA-6,10, PA-6,12, PA-12, amorphous PA, PBT, polyester elastomer, PC, POM, PPE, PEI, PEEK, PSF, PES, PPS, TPU, PVDF, polyfluoroalkoxy and ethylene tetrafluoroethyene copolymer. 10.3.6

Forming elastomeric Blends

Elastomers are often blended with thermoplastics in order to combine in the final product the thermoplastic resin processability and elastomeric behavior in the solid state.

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Polymer Blends Forming

PVC/NBR PVC/NBR blends have been developed in the 1940’s as an alternative to natural rubber. In addition, the blend provided improved processability and high resistance to oxidation, a notorious drawback for the natural rubber wire coatings. NBR/PVC (Geon™, JSR NV, Vynite™, Nipol™, Chemigum™ TPE, Paracril™ OZO) The different formulations of NBR/PVC blends marketed by Zeon under the trade name Nipol™ show either a co-continuous or dispersed structure. The NBR/PVC formulations vary from 50:50 to 70:30 wt%. The blends are processed mostly by extrusion into products such as cable jackets and hose covers. Injection molding is also being carried out. The blends show good ozone and abrasion resistance. Grades with higher AN content in the NBR display better ozone and fuel resistance while lower AN content results in improved low temperature flexibility. Goodyear markets blends of nitrile rubber with PVC, chlorinated PE, TPU or PA, or their combination, under the trade name Chemigum™ TPE. The blends show either a disperse or cocontinuous morphology. Properties include excellent oil and fuel resistance, high wear resistance and strength, a brittleness temperature below -45°C and a service temperature up to 120°C. Applications include gaskets, seals, joints, geomembranes, seals and liners. NBR/PVC blends marketed by Uniroyal (Paracril™ OZO) cover the concentration range NBR/PVC = 70:30 to 60:40 and show a dispersed morphology. They are typically formed by injection molding, extrusion and calendering and show the good fuel and oil resistance and low compression set. Dynamically vulcanized blends Dynamically vulcanized blends are produced by melt mixing an elastomer phase with a thermoplastic (the elastomer being the major component) and then curing the elastomer during melt mixing. The resulting blend comprises the thermoplastic resin filled with high concentration of rubber particles well bonded to the matrix. Examples include EPDM/PP (Santo-

prene™), PB/PP (Vyram™, Trefsin™), NBR/PP (Geolast™), PVC/crosslinked ethylene copolymers (Alcryn™). The cured elastomeric phase provides high elastic character while the thermoplastic enables easy melt processing. Curing the elastomer phase provides improved melt strength, better morphology stability during processing while maintaining recyclability. Because of the ease of processing, dynamically vulcanized blends have displaced fully cured rubbers in many applications. Curing takes place during melt mixing, for example in a twin screw extruder, and the blend pellets can then be used in typical thermoplastic forming processes. The blends can be extruded into window glazing, weather stripping or wire and cable jacketing. They can also be injection or blow molded into parts such as sleeves, air ducts, and protective covers, in particular for automotive applications. Crosslinked ethylene copolymers/PVC (Alcryn™) In these dynamically crosslinked elastomers, high volume fractions of ethylene terpolymers, such as ethylene-n-butylacrylate-carbon monoxide and ethylene-vinylacetate-carbon monoxide terpolymers, are dispersed into the PVC matrix. Due to the carbonyl, the ethylene terpolymers can be crosslinked selectively through the use of peroxide/bis maleimides or diamine curatives, or radiation techniques. Crosslinking takes place during melt mixing. The blend pellets can then be used in typical thermoplastic processes. The advantages are that the material performs and looks like a rubber but is processable. All grades have excellent oil and weather resistance. Main uses for the material include seals, gaskets, weather-stripping, conveyor belts, tubing and hose, and wire and cable. 10.3.7

Drying

Moisture in the materials during forming can cause problems such as bubbles, voids or streaks. In addition, in condensation polymers such as PA’s or PEST’s, water can react with the polymers causing molecular weight loss. Moisture level is then particularly critical and, in general, blends

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containing condensation polymers require careful drying. Improper drying would affect a number of properties, e.g., tensile strength, maximum strain at break, low temperature impact strength. Moisture absorption varies with the nature of the polymer, particularly the polarity of its chemical bonds. Because most engineering polymers contain polar groups, they are hygroscopic and must be dried prior to processing. Maximum allowable moisture content, and consequently the drying procedure, varies, based on the hygroscopic nature of the components and the composition. For engineering polymers, the maximum moisture content generally recommended is 0.02 wt%. The water content is usually determined using the Karl-Fischer titration method. The blends that do not contain hygroscopic components, such as PPE/HIPS in principle do not have to be dried. However the material, when exposed to high humidity, adsorbs surface moisture and small voids and bubbles can appear after processing. These flaws can be avoided and better product appearance and performance can be obtained if the blend is thoroughly dried before or during (e.g., using vented barrel) processing. Hot air, desiccant-bed type dehumidifying hopper-dryers are generally recommended. These remove moisture from the process air to deliver hot air with a dew point usually between -30°C and -40°C at the hopper inlet. An air flow rate of one liter per second for each kg/hour of resin flow is recommended. The drying temperature depends on the material and is usually set at temperature lower by 10 to 15°C than the heat deflection temperature. A drying time of more than 2 hours is typical. During the first hour of drying the blend normally looses 90% of the original moisture. Excessive drying should be avoided since it may result in surface oxidation. Tables 10.13 to 10.16 show typical drying conditions for selected commercial neat polymers and blends. The drying times shown in Tables 10.13 to 10.16 are averages of the recommended drying times once the resin has reached the desired temperature. Larger size pellets require longer drying. The following sections give additional information for the different types of blends.

10.3.7.1 Commodity Blends Polyolefins PP/EPDM (Hifax™) Drying is normally not required. If the material was exposed to excessive moisture, drying can be carried out for 1-2 hours at 90-120°C. Polyvinyls PVC/PMMA (Kydex™) Kydex™ is primarily formed into sheets, to be thermoformed later. The sheets pickup moisture when exposed to humid atmosphere. If sheets contain 0.2% moisture or more, bubbles will result when the material is heated to the thermoforming temperature (205°C). Sheets can be dried in a circulating air oven at 70°C in 16 h (for thickness of 3 mm) or in 24 h (thickness of 6 mm). Drying can be avoided if the sheets are stored in a low humidity environment or thermoformed at 180°C or less. Styrenics HIPS In HIPS (Polystyrol™) or PS/butadiene rubber (Styroplus™, Styroblend™ and Styroblend™ WS), the moisture absorption is dominated by PS, which as a non-polar material has little tendency to pick up moisture. These materials normally do not require drying, unless stored outside for prolonged periods. If necessary or for parts for which appearance is critical, 2-3 hours at 80°C are sufficient. In the case of HIPS/PPE blends (VestoranTM), neither component is hygroscopic. In the majority of cases, pre-drying is not needed. Vented barrel processing is sufficient. ABS Drying requirements of ABS-based blends vary widely, depending on the other components and their concentration. In general, the presence of moisture will produce surface defects such as silver streaks or blisters. With non hygroscopic

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Polymer Blends Forming

Table 10.13. Drying conditions for commodity polymers-based blends* Blend

Trade name

PP/EPDM PP/EPDM PP/EPDM PP/EPDM PP/PA-66 or PA-mXD6 PS/PB Styrenic/rubber PS-co-PMMA ABS ABS ABS/elastomer ABS/BA or PB ABS/PVC ABS/PVC ABS/PVC ABS/PC ABS/PBT ABS/PBT ABS/PA ABS/PA ABS/PA-6 ABS /TPU SAN/elastomer SAN/elastomer SAN/saturated olefinic elastomer SMA/ ABS SMA/ rubber ASA/PC ASA/AES terpolymer /elastomer Acrylic based

IPCL JSA Keltan TP Kelburon Santoprene Eref Styroplus, Styroblend Limera Zylar Novodur Magnum Lustran Terluran Triax CBE Enplex Lustran Techniace TC Techniace TB Cevian Techniace TA Triax 1000 Stapron N Prevail Lustran Luran S Rovel Cadon Stapron S Terblend S Centrex Acrylite, Cyrolite, XT

Moisture** (wt%)

0.3

0.2 0.1

0.3 0.02

0.1 (0.03) 0.1 (0.03)

Temperature (°C)

Time (hours)

95 100 100 65-75 120 80 90-100 82 80 82-88 75-80 80 75 70-80 75-80 100-110 90-100 80-120 90-100 90 80 75-80 75-80 80 85 90 (82-93) 80 10-20 below Vicat 80-90 (83-93) 80

0.75 2 2 2-3 4-5 2-3 3-4 2 1-2 2 4 2-4 2 2-3 4 2 2-3 3-5 2-3 2-4 3 4 4 2-4 3 2 (3-4) 2 2-4 2 (3) 3

* Conditions recommended by manufacturer for injection molding. Values for extrusion indicated in parentheses if different. Values for other processes indicated in text when different. ** Recommended final moisture content

components, such as in blends with PB or acrylic rubber (Terluran™), the systems show moisture absorption similar to pure ABS. On the other hand, requirements for blends with engineering polymers are more stringent. ABS/PA-6 (Stapron™ N) alloys absorb 1.5 wt% of moisture in normal

atmosphere or 4 wt% when immersed in water. PA is highly hygroscopic - insufficient drying of ABS/PA blends (Techniace™ TA) brings silver streaks. Drying of ABS or AES blends with PC (e.g., Techniace™ TC) is necessary to reduce risk of hydrolysis of PC.

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up to 8 wt% moisture absorption. Consequently, their blends need proper drying, regardless of the other component. Molding of PA-6 or PA-66 blends with ABS (Novalloy™-A) without predrying can not only impair physical properties but also may cause adverse changes in appearance. Glass fiber reinforced grades of semi-aromatic co-polyamide (from caprolactam, hexamethylenediamine and terephthalic acid) blends PA6IT6/elastomer (Ultramid™ T) are particularly sensitive to moisture. The ester linkage is particularly susceptible to hydrolysis. Thus blends containing PEST must be dried to prevent splaying: PC/PB/PET or PBT blends (Stapron™), PCTG with PC or SMA (Ektar™ MB), PBT/ABS (Novalloy™-B or Cevian), PC/PET or PBT/PB (Stapron™ E), PC/ABS/ SAN (Bayblend™) and PC/ABS (Bayblend™ T). In PET/elastomer systems (Rynite™), if moisture exceeds 0.2 wt% the processability and final

A number of styrenics-based blends do not undergo significant changes when exposed to moisture but surface quality of molded parts is affected: SAN/elastomer (Luran™ S), ASA/PC (Terblend™ S), PS-co-PMMA/elastomer (Zylar™) and acrylic-based multipolymer compounds (Acrylite™, Cyrolite™, XT). For all these systems, it is therefore essential to dry pellets before processing. Drying is not important from the point of view of mechanical performance that remains fairly unaffected, but the presence of moisture may give rise to streaks, blisters, haze, low surface gloss or, in extreme cases, even bubbling. The use of vented extruder is not a substitute to good drying. 10.3.7.2 Engineering Blends Engineering resins are in general hygroscopic and require thorough drying. For example, PA’s show

Table 10.14. Drying conditions for polyamides-based blends* Blend

Trade name

PA-6, PA-66 or PA-612/ ionomer or elastomer PA-6/elastomer PA-66/elastomer PA-6/elastomer PA 6/6T+ elastomer PA-66-co-6/elastomer PA-6/BA copolymers PA-6/methacrylate and n-acrylate copolymer PA-6 or PA-66/ elastomer PA-66 or PA-66/PA-6/ elastomer PA/PP PA-6 or PA-66/ABS PA-6 or PA-66 / ABS PA-12/aromatic aliphatic polyamide PA-6/PPE

Zytel Grilon A Ultramid A Ultramid B Ultramid T Ultramid C Durethan Durethan RM Fiberfil TN Technyl Orgalloy R Novalloy-A Triax 1000 Grilamid Dimension

Moisture (wt%)

Temperature (°C)

Time (hours)

0.25

80

3-6

0.2

100

2

0.15 (0.10) 0.15 (0.10) 0.1 (0.05) 0.15 (0.10) 0.1 0.1

80-100 80-100 80-110 80-100 80 80

10-24 4-6

0.1 0.2 0.15 0.02 0.02 0.2 0.2 (0.1)

100-110 70-80 80-90 80-100 90 80 110 (80)

2 5 4 4-6 4 6-10 as needed

* Conditions recommended by manufacturer for injection molding. Values for extrusion indicated in parentheses if different. Values for other processes indicated in text when different.

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Table 10.15. Drying conditions for engineering polymers and their blends: others* Blend

Trade name

PET/elastomer PBT PBT/elastomer PBT/ABS PBT/ASA PCTG/PC or SMA PC PC/Copolymer PC/BA-co-BMA elastomer PC/ABS PC/ABS PC/ABS/SAN PC/ABS PC/ABS PC/SMA PC/PET PC/PET/ABS or PC/PET/HDPE PC/PET or PBT PC/TPU or TPU/PC POM/TPU, PTFE POM/TPU PPE/ HIPS PPE/ PA PPE/HIPS

Rynite Pocan Ultradur Novalloy-B or Cevian Ultrablend S Ektar MB Makroblend EC Makrolon Novalloy-S Pulse Bayblend Bayblend T Triax 2000 Arloy Makroblend DP-4-1368 Makroblend Sabre Texin Hostaform Ultraform Noryl Noryl GTX Luranyl

Moisture** (wt%)

Temperature (°C)

Time (hours)

0.2 0.04 0.05 0.02 0.02

120-135 120 100-140 80-120 100-140 77-107 120 120 121 80-120 90-95 100-110 100-110 90-100 105-110 100-102 100-110 110-120 80-110 105-120 110 110-120 110 80

2-6 2-3 3-4 3-5 3-4

0.02 0.02 0.02 0.02 0.03 0.02 0.02 (0.1) 0.02 0.02 0.02 0.02 0.02 0.02

0.02-0.05

3-6 2-3 3-4 3-5 3-4 1-4 1-4 4 2.5-3 4-6 4-6 3-4 1-3 3-6 2 2 3 3

* Conditions recommended by manufacturer for injection molding. Values for extrusion indicated in parentheses if different. Values for other processes indicated in text when different. ** Maximum moisture content recommended by manufacturer for injection molding. Values for extrusion indicated in parentheses if different. Indicated in text when different for other processes.

properties are affected by hydrolysis of PET, e.g., the melt viscosity is significantly reduced, so is the tensile and impact strength. However, it is worth noting that molded parts do not show the surface defects, e.g., splay, reported for PA resins and their blends. Therefore polyester blend parts may show excellent surface appearance and still have poor mechanical properties. Since the decreased molecular weight is not easy to restore,

the hydrolyzed regrind should be discarded. For PC/ABS/SAN (Bayblend™) and PC/ABS (Bayblend™ T), the recommended moisture content is less than 0.02 wt% for injection molding and 0.1 wt% for extrusion. However, drying before processing might not be necessary for the latter blends if extrusion or injection equipment provides for efficient devolatilization.

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PC/TPU (Texin™ 3000) or TPU/PC (Texin™ 4000) blends are hygroscopic. Since the TPU in the blends is a polyester-urethane, the material is susceptible to hydrolysis. Equilibrium moisture at 50 %RH is below 0.4% and it is reached in less than 40 hours, while that at 95 %RH, is 1.2% and reached in 100 hours. Pigments or additives to be added to the blend must be dried as well. The POM/TPU and POM/PTFE (Hostaform™) compositions contain very little moisture and usually pre-drying is not required unless the product was exposed to excessive moisture. It should be noted that 0.2-0.3% moisture can reduce the material’s thermal stability by 20-30°C. 10.3.7.3 Specialty Blends Drying is recommended for PSF/PET (Mindel™ B) blends. Dehumidified hopper drier or circulating air oven with air dew point of -29°C, are preferred.

10.4

Extrusion

10.4.1

General Description of Extrusion

Extrusion is a continuous process in which an Archimedes screw, rotating within a heated barrel, melts and pumps the polymer through dies to form continuous shapes such as pipes, profiles or sheets. In addition, extrusion is the basic element in a number of other polymer processing operations. The plastication and pumping functions

are performed by screw extruders in most forming operations such as calendering, film blowing, wirecoating, as well as injection, or blow molding. Although extrusion is used commercially for rubbers, and for solid state forming of thermoplastic polymers, in the chapter only melt extrusion of thermoplastic blends will be covered. Interested readers will find in depth discussion on extrusion in recent monographs [Rauwendaal, 1986; White, 1992]. The screw geometry, temperature profile and screw rotation speed are the most important operation parameters of the process. The screw is the heart of the process and its design is the process’ most important factor. The main variables are geometry (i.e., one or two stages, barrier screw, etc.), screw diameter and length, usually expressed as length-to-diameter ratio, L/D. These latter two variables influence extruder throughput, polymer residence time and rate of heat transfer. Typical L/D ratios for thermoplastics are around 24 for simple operations and higher for more complex ones, where venting or mixing, for example, are important. Extruders can have different designs and modes of operation, e.g., single or multi-screw, continuous or discontinuous. The basic operation of the most widely used extruder for forming thermoplastics into finished products, the single screw extruder, will be reviewed in the next section. This provides the concepts that can be extended to other, more complex machines, viz. the multiscrew and specialty extruders discussed briefly in

Table 10.16. Drying conditions for blends of specialty polymers Blend

Trade name

PSF/ABS PSF/PET TPU/ABS NBR/EPDM

Mindel A Mindel B Prevail Geolast

Moisture * (wt%)

0.02

Temperature (°C)

Time (hours)

121 163 75 75-90

3 3-4 4 4-5

* Maximum moisture content recommended by manufacturer for injection molding. Values for extrusion indicated in parentheses if different. Indicated in text when different for other processes.

Polymer Blends Forming

the subsequent sections. Review of the historical developments of extruders is given in Chapter 1 while a more detailed discussion of the operation of multi-screw extruders, particularly useful for compounding, is given in Chapter 9, Compounding Polymer Blends. 10.4.1.1 Single Screw Extruders (SSE) Early extruders were ram presses used for forming rubbers, for wire and cable coatings. Although some ram presses are still used for extrusion (e.g., for extrusion of propellant grains), the screw extruders dominate the market. The screw extruders were first developed in the late 19th century to produce miscellaneous products like lead pipes, soap and foodstuffs. The early extruders had small diameter, short (L/D = 4) single screw and steamheated barrel. However, even these early beltdriven units showed all the characteristic features of the modern extruders, viz., speed controlling gears, hopper, screw, dies of different sizes and shapes. Screw extruders for wire and cable rubber coating were introduced in the 1890’s. The development of the first extruders designed for thermoplastics followed closely the development of the thermoplastics themselves. These machines were first introduced in the early 1930’s. By the late 1930’s, electrically heated, air cooled extruders with L/D = 10 were available. Breaker plates, screen packs, crosshead dies, coextrusion, monofilament extrusion, film blowing and biaxial sheet orientation were introduced in the 1940’s. By the late 1950’s, extruder venting and twostage screws became available and the extruder length-to-diameter ratio increased to L/D = 15 and then to L/D = 30 by the mid-60’s. The latest developments in extrusion systems, introduced in the 1980’s, feature helical grooved feed barrel, high pressure gear pumps allowing 40 to 80 MPa pressure on the die, air lubricated dies, biaxial orientation of blown film and strong emphasis on computer control of the process. Today, automated extruder lines, equipped with closed loop control, generate output in excess of 20 tons/hr. In the modern extruder, the screw is usually divided in three zones that correspond to different

687

physical states of the polymer: Feed zone The polymer in the solid state is fed from a hopper and falls into the channels of the screw. It is compacted and moves forward as a result of the frictional forces. The diameter of the screw root is constant in this zone. The material is initially compressed, which forces the air out from between the resin pellets and back through the hopper. This creates a more homogeneous feedstock that can be melted more easily. Compression zone In this second stage of the screw, the depth of the channel is progressively reduced. Here, both solid and liquid polymer is present. The narrowing of flow channel induces a pressure build-up and helps the melting process. The compression zone usually coincides with the plastication zone (the latter being defined as the region where solid and melt coexist), but the coincidence is not perfect since the boundaries of the plastication region vary with the polymeric systems and operating conditions. Metering zone In the third section, characterized by a constant diameter of the screw shaft, the molten polymer is mixed and progressively compressed. The flow brings the molten polymer to the pressure required for the melt to flow through the die. In addition to the three regions of the screw, the extruder includes the die zone. As the polymer flows through the die, it takes the shape of its flow channel. A certain die-head pressure is required to force the flow of polymer through the die. This pressure is provided by the action of the screw and depends on the polymer rheology and the shape of the die. The single screw extruder has several advantages: relatively low cost resulting in good performance-to-cost ratio, straightforward design, raggedness, reliability, easy operation and maintenance. However, it is a poor mixer with relatively long residence time and broad residence time distribution. Exceptions are the Buss Kneader and Barmag-type extruders, providing good mixing

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action, but at increased complexity and cost. However, the versatility of single screw machines has been upgraded by utilization of special mixing screw designs, improvements of feeding methods, such as drying and premixing or downstream feeding, as well as diverse mixing screws and add-on homogenizing elements. More complex screw geometries are also available. These are used for specific purposes. For example, twostage screws provide a decompression zone for devolatilization and barrier screws provide a more efficient melting. In the first part of the extruder, the solids are conveyed in the down-channel direction through particle/barrel frictional drag. It is accompanied by the compaction of the deformable solid bed and increase in pressure. Friction generates heat mainly at the barrel/solid bed interface. Heat generation is proportional to pressure and the friction coefficient between the barrel and the solid particles. This friction-generated heat is added to heat conducted through the barrel to raise the temperature of the polymer up to the desired melt processing level. Once a film of molten polymer has formed at the interface with the barrel, the transport mechanism changes from friction to viscous drag. The polymer is conveyed through shear stresses generated in the molten film. A number of models with diverse sets of assumptions exist to describe the melting process [Halasz, 1993]. Only an abbreviated description of the most widely accepted views will be presented here. Figure 10.8 shows a schematic view of the screw channel cross-section in the melting region. The melt starts accumulating at the pushing flight (on the right hand side) and the size of the solid bed continuously decreases. A circulating flow develops in the melt phase and eventually around the solid bed. At the same time, the depth of the channel decreases, resulting in increased pressure, helping the melting process. The melting process depends on material properties, particularly the friction coefficient. A significantly different melting mechanism was proposed for rigid PVC, because of the lack of adhesion to the barrel wall [Gale, 1970, Chung, 1976].

Figure 10.8. Cross-section view of extruder screw, showing the melt pool and the solid bed against the screw’s pushing flight.

10.4.1.2 Twin Screw Extruders (TSE) Although single screw extruders dominate the world of extrusion, there are many applications where a twin screw extruder is a better choice despite higher purchasing and maintaining costs. More complex extruders with several planetary screws have also been developed. These find a commercially marginal use in specific applications. The development of multi-screw systems began in the 19th century and saw rapid progress. The first twin screw extruder was introduced in 1869 and the first modular counter-rotating machine was patented in 1882. The direct predecessor of the modern twin screw extruders was an intermeshing, co-rotating twin screw extruder designed by Colombo in 1938. The machine was used since 1939 by BASF for continuous reactive extrusion. After the second world war significant progress in equipment design took place in West Germany. As a result, in 1959 Werner & Pfleiderer introduced the ZSK twin screw extruder family, which was to become the industry’s standard. These machines are vented, intermeshing, corotating, with segmented screw and barrel. In 1972 Cincinnati Milacron developed conical fully intermeshing twin screw extruder for extrusion of PVC profile. Twin screw extruders are generally classified by the direction of rotation: co-rotating and counter-rotating, and by the degree of intermeshing: non-intermeshing (tangential), completely or partially intermeshing. Currently, a wide variety of twin screw machines are being offered. Some

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Polymer Blends Forming

find niches in specific applications, particularly when extensive devolatilization, high additive loadings, and/or intensive dispersive melt mixing are important. Twin screw machines are superior to single screw machines in these applications because they can provide [Murray, 1978; Wood, 1980]: • better feeding and more positive conveying characteristics, • shorter residence times and narrower residence time distribution, • improved kinetics and interphase renewal, • high and controlled shear rates, • self-wiping action, • high level of devolatilization. In intermeshing, co-rotating extruders, the material is conveyed in a figure-eight pattern, alternating between moderate shear stress in the overflight zone (between screw and the barrel) and high shear stress in the apex zone. Owing to segmented design both the screw and the barrel can be custom assembled to optimize for a specific production. The screw is usually made of several types of conveying and kneadingmixing elements, both being able to convey the material either forward or backward (for enhanced mixing). These extruders are easier to scale-up, provide good and adjustable balance between the dispersive and distributive mixing and they can be operated at relatively high output rates. In general, TSE subjects the material to more intense and more uniform shear stress than SEE does. The level of stress depends on the shape of the screw flights, the degree to which they intermesh and the relative directions of the screw rotation. Because of this, the conventional twin screw extruders are more frequently employed as compounding machines than forming units. There are several major differences in operation of the single and twin screw machines, some of which are summarized in Table 10.17. The processes in the twin screw extruder depend primarily on the mode of operation (co- or counterrotation) and degree of fill. The length of the feeding and melting zones can vary considerably, depending on the screw configuration and rotation

speed, but in general these are much shorter in twin screw than for single screw machines. While in the single screw extruder, a continuous solid bed is formed, this does not happen in a twin screw machine since the screw chambers are normally not totally filled and the solids conveying zone is continuously divided by intermeshing with the other screw. Table 10.17. Differences between single and twin screw extruders

Melting Material transport Mixing

Single screw

Twin screw

Long melting zone Drag flow

Short rapid melting Positive displacement

Poor

Excellent

The melting process in TSE is similar but not identical to that in a SSE and follows the general stages: i) melt film formation engendered by friction between the polymer granules and between the granules and the barrel surface and the screw flights, ii) material from the melt film moves to form a melt pool, iii) the size of the melt pool increases while that of the solid bed gradually diminishes to zero. However, melting in a twin screw extruder is much faster because of: i) the more complex flow pattern of the polymer pellets which leads to intensive friction not only with the extruder surfaces, but also between the pellets, and ii) the rapid increase in pressure in the regions of the extruder subjected to backward pumping engendered by downstream presence of the reverse (so called left-handed) screw elements or kneading blocks. The melting process is strongly influenced by the screw geometry. The material is compacted rapidly in the calendering gaps. In counter-rotating machines, the polymer is forced through the calendering gap between the screws, helping the melting process. On the other hand, in co-rotating or counter-rotating extruders, the solid bed is rapidly broken, enhancing the friction between polymer granules. The break-up

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of the solid bed does not cause disruption in the flow, as is the case for single screw extruders, because of the absence of a continuous bed. The length of the melting zone is nearly independent of the screw speed and for well designed screw configuration is approximately equal to the screw diameter. The material transport mechanism in a twin screw extruder also significantly differs from that in a single screw machine. Material transport in a single screw extruder takes place mainly through drag flow while in an intermeshing twin screw extruder, it relies more on positive displacement. This can be a significant advantage for processing materials with poor frictional properties that cause problems in single screw extrusion. The degree of positive displacement depends on how well the flight of one screw closes the opposite channel of the other screw. The most positive displacement is obtained in a closely intermeshing, counterrotating geometry. The third major difference is the velocity patterns in the machine. The velocity profiles in single screw extruders are well defined and fairly easy to describe. The situation in twin screw extruders is considerably more complex. However, the complex flow pattern in a twin screw extruder gives several advantages, such as good mixing, good heat transfer, large melting capacity, good devolatilization capacity, and good control over the melt temperatures. Further discussion on the relative merits of SSE and TSE can be found in Chapter 9, Compounding Polymer Blends. 10.4.1.3 Specialty Extruders Extruders with special configuration are mainly used for compounding with limited applications for forming. The Buss Kneader is one of these specialty mixing single screw extruders that is relatively easily to scale up, has good mixing capability resulting in narrow residence time distribution and relatively low residence time. It is characterized by a constant channel depth screw with interrupted flights and three rows of stationary pins in the barrel. During operation

the screw rotates and axially oscillates providing self-wiping action of the blades. There is a good deal of distributive mixing during the high-stress dispersive process. The clam-shell barrel design makes it easy to clean, service, analyze, and optimize. The machine is used for polymer blending, for compounding PVC formulations or engineering resins with reinforcements, for foodstuffs, carbon electrode pastes, etc. [Jakopin, 1989]. Planetary extruders with several parallel screws have also been developed for polymer compounding. In the planetary roller-extruder, six or more planetary screws, evenly spaced, revolve around the circumference of the central (or sun) screw. The planetary screws intermesh with the sun screw and the barrel. The planetary barrel section, therefore, must have helical grooves corresponding to the helical flights on the planetary screws. This planetary barrel section is generally separated by a flange connection from the feed barrel section. In the first part of the machine, before the planetary screws, the material moves forward as in a single screw extruder. As the plasticated material reaches the planetary section it is exposed to intensive mixing by the rolling action between the planetary screws, the sun screw and the barrel. The helical design of the barrel, sun screw and planetary screws result in development of large interfacial area that leads to good mixing. The planetary extruders are used for extrusion/compounding of PVC formulations, both rigid and plasticized [Collins, 1987; Anders, 1979]. Planetary roller sections are also used as add-ons to regular extruders to improve the mixing performance [Rust, 1983; Huszman, 1983]. There are a number of machines that do not utilize Archimedean screw for transport of the material, but still fall within the category of continuous extruders. Sometimes these machines are referred to as screwless extruders or elastic melt extruder. These machines employ disks or drums to melt, mix, pump, and extrude the material and are based on viscous drag transport principle. They are not widely used in industry as forming machines, and will not be discussed here.

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Polymer Blends Forming

10.4.2

Processing Parameters

10.4.2.1 Screw and Barrel The geometry of the screw is the most important design parameter controlling the performance of an extruder. Screw configuration depends to a large extent on the properties of the material being processed. A constant-pitch metering screw is usually employed in applications not requiring intensive mixing. Where mixing is important (e.g., for color dispersion) a two-stage screw equipped with a let-down zone in the center of the screw is appropriate. Turbulence-promoters can also be included at or near the tip of the screw. In some applications with two-stage screws, venting at the let-down section may be needed. Theoretically, to optimize the performance, a different screw is needed for each type of material, temperature profile and extruder die. This requirement is not realistic in practice. Thus, rules ensuring the most satisfying results without major problems have been developed. For example, it is important that the depth of the metering zone allows satisfying pumping while avoiding overheating, especially at high screw speeds and with highly viscous materials. On the other hand, a limited depth ensures a sufficient high pressure to maintain regular throughput. Typical geometries are shown in Table 10.18. The change in flight depth along the screw (i.e., the ratio of the solids metering channel depth to the melt pump channel depth) is known as the

compression ratio. The purpose of the screw flights is to enable the screw to transport polymer down the barrel. The pitch angle of the flights depends on the type of material. Helix angles are usually chosen to optimize feeding characteristics. Many applications employ a general-purpose screw with a constant pitch (i.e., flight equals the diameter). The pitch angle of this single-flighted screw is usually 17.61°, but can be varied from 12 to 20°. The screw outside diameter (OD) is generally just a fraction of millimeter smaller than the barrel internal diameter (ID). The reason for such a tight clearance is to limit the amount of material flowing back along the barrel, resulting in a loss of pumping capacity. Typical single screw extruder specifications are as follows: • Compression ratio = 2:1 to 6:1 (for materials ranging from LDPE to PA) • Pitch angle = 12 to 20° • Length-to-diameter ratio, L/D = 16:1 to 36:1 (Ratios of 20/1 and 24/1 are common for most thermoplastics). • Extrusion pressures = 70 to 200 MPa • The lead on a screw (defined as the distance between the flights) approximately equals the inner diameter of the barrel. • The radial clearance between the flight tip and the barrel is tight and usually constant. Extruder screws and cylinders are made of standard tempered and high strength steels. For increasing the wear resistance, nitrating or heat hardening is used. Cooling of the entrance zone is

Table 10.18 - Typical geometries of a single screw extruder (lengths expressed as L/D ratio) Thermoplastic type

Standard Wide softening range (amorphous, ex: PS) Wide softening range and low heat stability (PVC) Narrow softening range (semi-crystalline)

Screw zone length

Compression ratio

Feeding (L/D)

Melting (L/D)

Metering (L/D)

4-6 4-6 4-6 14-16

7-14 7-13 12-18 1.5-2.0

4-8 3-8 0-3 4-7

2.5-3.0 1.5-2.0 3.0-4.5

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M. M. Dumoulin

recommended for screws with diameter larger than 90 mm. Cooling the feed hopper also improves feeding. Furthermore, when using grooved barrel entrance zone, the temperature of these should be separately controlled to improve the feeding and throughput. The grooves should not be larger than 2x2 mm and only 3D long.

that, on the one hand, elongate the dispersed drops into fibers, which may disintegrate into fine droplets downstream from the die. On the other hand, the convergence induces coalescence. The net effect of these two competing influences depends very much on the nature of the blend and how well its morphology has been stabilized.

10.4.2.2 Die

10.4.3.1 Extrusion of Blends of Commodity Resins

The function of the die is to form the melt into a desired shape and to deliver this shape to the take-off equipment. This has to be done as fast as possible with minimum error in dimensions. It is important that the die uniformly distributes the melt across the channel in order to ensure that the extrudate exits the die at even velocities. The flow inside the die depends on the rheology of the melt, therefore a die is usually designed for a specific resin. For example, dies designed for film blowing of LDPE will not be appropriate for processing LLDPE which, because of its higher viscosity and different extensional viscosity behavior, generates higher back pressure and has propensity for flow instability. To accommodate more than one resin, dies are often built with temperature and pressure adjustment capabilities and with adjustable lips or blades. Changing the temperature profile across the die land or the gap might allow for different resins to be extruded using the same die. Die design must take into account the viscoelastic nature of polymers, e.g., the extrudate swell depends on the rheology of the resin, blend morphology, and geometry of the die — longer die lands tend to minimize swell.

This section aims at discussing the conditions and recommendations for successful extrusion of blends containing commodity resins as the major component. The conditions recommended by manufacturers are listed in Table 10.19, while the recommended procedures and accepted practices are discussed below for each type of blend.

10.4.3

Extrusion of Polymer Blends

The whole spectrum of the different extrusion processes is used to form blends, i.e., production of profiles, sheets, tubes, wire and cable coating. The products find use in diverse industries: electrical, appliances, automotive, etc. The die region is extremely important for forming polymer blends. The convergent flow into the die engenders strong elongational forces

Blends with Polyolefins Blending polyolefins affects viscosity and elasticity whilst the morphology obtained also affects rheology. For example, addition of HDPE to LLDPE caused increases in viscosity and higher melt elasticity, indicated by large extrudate swell during extrusion, with higher molecular weight HDPE causing higher elasticity [Karbashewski et al., 1993]. A higher elasticity may have a negative effect on processability since it is often associated with melt defects. In these blends, surface roughness increased with addition of HDPE, and the larger effect was seen for the higher molecular weight. Haze in extruded films appeared at lower extrusion rates in the presence of HDPE. It should be noted that molecular weight effects on properties might be difficult to isolate, since changes in MW will also affect mixing. In these blends, an increase in the fraction of linear molecules resulted in decrease of processability, both in terms of elasticity and viscosity. As indicated in Chapter 7, Rheology of Polymer Alloys and Blends, an increase in viscosity is associated with an emulsion-like blend behavior of the dispersion. On the other hand, extrudate swell results from the recovery of drops strained into filaments in the convergent flow to the die. Polyolefin blends may be extruded using either the pre-compounded blends or by directly blend-

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Polymer Blends Forming

Table 10.19. Extrusion conditions for commodity resin blends Blend

LDPE/PIB PP/EPDM PP/EPDM(Fibers) PP/EPDM(Sheets) EPR/PP TPO PS/BR ABS/BA or PB ABS/elastomer ABS/PVC SAN/elastomer SAN/elastomer SAN/saturated olefinic elastomer ASA/AES/ elastomer ASA/PVC ASA/PC SMA/ABS PVC/elastomer Acrylic based

Trade name

Lupolen 250H Santoprene Modylen Modylen Polytrope Oleflex Styroplus Terluran Lustran Lustran Lustran Luran S Rovel

Temperature (°C) Barrel

Die

182-215 160-180 150-230 190-210

192-215 240-270 230-250

Screw Melt

L/D

CR*

180-240 190-210

24

2.5-3.5

195-215 200-220 180-280 230-240

24

2.5-3

30 25-30

2-2.5 3

230

25-30

2-4

Speed (rpm)

185-250 185-250 185-250 175-215

Centrex Geloy Terblend S Cadon Oxyblend OxyClear Acrylite, Cyrolite, XT

40-42 210-240

230-250

215-240 180-195 250-280 230-250 180 140-240

2.8-3.8

* CR: compression ratio

ing the resins in the extruder during the forming operation. The pre-compounded material is normally extruded using a single screw extruder, but if the compounding is to take place simultaneously with forming, a twin screw extruder with a screw configuration providing good mixing is recommended. For example, it is recommended to extrude the PP/EPDM blend (Modylen™) in a twin screw extruder with L/D of 25-30, equipped with several zones of kneading blocks, at the temperatures ranging from 160°C at the entrance to 270°C at the die. The pre-compounded pellets can also be fed to a single screw extruder with a two stage screw with L/D of 24-30 at temperatures of 150 to 250°C [Utracki, 1994].

Kamal and co-workers obtained lamellar morphologies in extrusion of PP/PVAl and HDPE/ PA-6 blends by using a special die and adapter design [Kamal et al., 1993]. They obtained the best permeability results at low die temperature and under conditions of low shear (i.e. metering screw at low speed). Also, they modified the viscosity ratio by changing temperature and found that a viscosity ratio close to unity was more conducive to laminar morphology and, thus, to lower permeabilities. Maleation of the polyolefin phase was introduced in order to compatibilize the blend and improve the impact toughness. Increasing the maleation level had a detrimental effect on the platelet structure and permeability.

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However the impact properties of the lamellar blends were lower than that obtained for dispersed morphology. Fluoropolymers have been used as processing aids because small quantities can reduce significantly the overall viscosity and thus facilitate extrusion. Feng et al. [1996] examined the mechanism of viscosity reduction in the capillary flow of HDPE/fluoroelastomer blends. X-ray photoelectron spectroscopy, used to characterize the composition of the extrudates surface, indicated only very small traces of the fluoroelastomer on the extrudate, pointing to the fact that the viscosity reduction is due to adhesive failure between the fluoropolymer layer and HDPE. HDPE/LDPE (Hostalen™ GC) HDPE/LDPE blends are marketed use in extrusion, in particular for paper coating, fibers and films for typewriter ribbons. In photographic paper coating, the material is used in 2 layers: a transparent coating for the back and a white colored layer on the front, on which gelatin and photosensitive layers are coated. The coating is very thin (0.012-0.030 mm) and takes place at high speeds (250 m/min). It requires very high drawability and homogeneity from melt as well as excellent adhesion to paper. The high melt temperature used (300°C) must not result in fish eyes which lead to distortion of photographic images. In the LDPE/PIB blend (Lupolen™ 250H), a LDPE with relatively low branching density (higher crystallinity) is used. Thus both LDPE and PIB have similar densities. The extruded parts are less transparent and glossy than those made of neat LDPE. All-purpose 3-zone screws with L/D = 18 to 22 was found suitable for processing. Standard practice favors single screw extruders with barrels of 25 to 200 mm. PP/EPDM blends (Santoprene™, Sarlink™) show a co-continuous morphology. The melt has high viscosity, but all the grades are easily extruded. The recommended extruder should have L/D • 24 and the screw a moderate compression ratio, e.g., 2.5 to 3.5 (similar to those used for PVC or PE extrusion). Temperature settings

should be increased with the blend hardness. A 20 to 60 mesh screen pack is recommended. Owing to the two-phase nature of the blend, the material has relatively high extrudate swell which increases with shear stress, i.e., with resin hardness at constant shear rate. For Oleflex™ TPO (PP and PE with α-olefin random copolymer), reactive processing (crosslinking) is used to get a finely dispersed morphology of copolymer dispersed in the polyolefin matrix. Recommended processing conditions are similar to those for neat PP or PE. For calendering, the melt temperature should be in the range 165-175°C. For Modylen™ polyolefin blend (PP/EPDM/ filler/interfacial additive), fiber filled grades should be formed using a twin screw extruder with L/D = 24 to 30, having the plasticizing zone equipped with kneading elements. However, sheet grades may be extruded using a single screw extruder with L/D = 24 to 30, preferably with a two-stage screw. Polytrope™ EPR/PP blends have either dispersed or co-continuous morphologies. For extrusion, approximately 60 mm diameter screws are recommended, with screen packs typically 40-40 to 60-80 mesh. Blends with PVC Extrusion temperature can have a significant impact on the final morphology of a PVC/ elastomer mixture fed in the form of a dry blend, thereby affecting performance. It was found that extrusion temperature significantly affects the impact properties of PVC/CPE or EVA mixtures. This effect originates in the characteristic structure of the suspension-type PVC grains. These particles, with diameter d = 100 µm, were created during polymerization by agglomeration of the “primary particles” (0.1-0.5 µm). At low extrusion temperatures, PVC grains do not disintegrate into primary particles. At higher temperatures the latter particles merge into a homogeneous material. It was shown that for extrusion temperatures of 185°C, honeycomb structures were formed, in which the modifier phase surrounds the PVC primary particles and, at a composition of only

Polymer Blends Forming

10%, formed the continuous phase [Menges, 1989]. The impact strength of the compound extruded at the optimum temperature of 185°C was almost 5 times that of the same composition extruded 5°C lower. The best conditions are those that lead to break-up of the large grains, but preserve the primary particles, which enables formation of the honeycomb structure. When the optimum temperature is exceeded, the PVC particles merge, a phase inversion occurs and the honeycomb structure disappears. The modifier is then in the form of droplets but in insufficient number to significantly affect the impact resistance. Producing the right morphology for these ultra high impact strength compositions is process-dependent and therefore can be difficult. Consequently, the trend in PVC impact modification has been toward the use of grafted acrylate elastomers. In Vinidur™ blends, the elastomer forms a reinforcing network in the PVC matrix, which gives a high impact strength. Extrusion of PVC/acrylate elastomer (Vinidur™) is easier because of the high shear stability of the graftpolymerized elastomer particles that are responsible for impact modification. The machine settings can thus be varied in a broad range. This feature distinguishes Vinidur™ from PVC modified by chlorinated PE or EVA. Both twin and single screw extruders can be used. Twin screw extruders can be fed with the dry blends whereas single screw extruders should preferentially be fed with pre-compounded pellets. Styrenics Blends In general, extrusion of styrenics blends is easy and single or two-stage screws can be used. However, when a single screw machine is selected, a two-stage screw with venting should be used to remove volatiles (moisture, styrene from degradation and entrapped air). HIPS (Polystyrol™) and PS/PB (Styroplus™) extrusion grades have high viscosity, but the use of external lubricants is not recommended since it may cause throughput fluctuations. The desired balance of rigidity to toughness can be tailored by mixing the high impact and general purpose grades. In such cases, products with similar flow characteristics should

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be used. High outputs are obtained using threezone screws arranged in tandem, having L/D = 30 and compression ratio 1:2 to 1:2.5. The vent, if necessary, is typically located between the two screws. If no vent is used, then the material should be dried before. For HIPS/PPE blends (Vestoran™), the die temperature must be exactly controlled and jet cooling nozzles should be installed adjacent to the die. Since ABS degrades by the simultaneous action of high temperature and atmospheric oxygen, the only products that may be used for extrusion and thermoforming are those with exceptional heat stability. Both ABS/elastomer blend series (Terluran™), those with butylacrylate elastomer, BA, and with polybutadiene, PB, satisfy this requirement. The ABS/BA series is especially suitable for critical pale shades. Vented extruders are preferred with L/D = 25-30 and compression ratio of 3 in both sections. Luran™ S SAN/elastomer blend is used for sheet or pipe extrusion. The resin withstands high temperature and shear forces without yellowing and can be considered as an easy material for processing by extrusion. Vented extruders are used with L/D = 25-30 and compression ratio of 1:2 to 1:4. In sheet extrusion, the stock temperature should be 230°C. For ASA/PVC blends (Geloy™), single screw vented extruders are recommended with L/D = 24, and compression ratio = 2.1-2.5. PVC degrades rapidly, thus residence time at 200°C should never exceed 1 min. For normal processing at temperatures well below 200°C, residence time should not exceed 5 min. ASA/PC blends (Terblend™ S) are prepared by grafting acrylic ester elastomer on the SAN matrix, resulting in elastomer particles very finely distributed within SAN which is then homogeneously mixed with PC in the melt. The blend should be extruded on the same machines as SAN and ABS resins. Sheets and films with a good, somewhat dull, finish can be produced at stock temperatures of 250-280°C. Chill roll temperature settings of 70-110°C are standard practice and under no circumstances should the temperature be higher than 5°C below the Vicat softening

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temperature. In general, melt temperature for extrusion of profiles is set at lower temperature than for sheets. For extrusion of Centrex™ ASA/AES terpolymer rubber modified blends and SMA/ABS blends (Cadon™), standard practices recommended for other styrenics systems should be used. Thus, when a single screw machine is selected, a two-stage screw is preferred, but a single-stage screw may also be used. When using two-stage screws, the recommended compression ratios are 2.5-2.7 for the first stage and 1.5-2.0 for the second stage. Introduction of excessive shear into the polymer during extrusion should be avoided. PMMA, PS or ABS is the best purging agents.

diameter D = 80-150 mm, compression ratio of 2.8-3.8, be moderately long, with constant depth feed zone and a gradual transition to a long metering section.

Blends of Other Commodity Resins Acrylic-based blends (Acrylite™, Cyrolite™ and XT) are often used to extrude sheets for applications where transparency is important. Thus, to ensure clarity, it is important to avoid contamination. The combination of long residence time and high temperature will increase viscosity leading to gel-like skin, with loss of gloss and transparency. Orientation in the final product should be kept below 25%. Suitable screws should have

PA/PP PA-6 or PA-66 blends with PP (e.g., Akuloy™ RM, Dexpro™, Dexlon™, Eref™, Flexloy™, Gapex™, LAX 23, MCX-Q, NB, Novamid™ AC, Orgalloy™, Poliblend™ NH, Snialoy™, Systemer™ S, Thermocomp™, Ube Alloy CA, Ultramid™ KR, UTX) have flow properties similar to PP. This constitutes an improvement over the flow behavior of neat PA. The PA/PP blends are easier to extrude than PA is because the melt viscosity is

10.4.3.2 Extrusion of Blends of Engineering Resins Conditions recommended by the manufacturers for extrusion of engineering resin blends are listed in Table 10.20. The recommended procedures and accepted practices are discussed below for a number of different types of blends. Polyamide Blends

Table 10.20. Extrusion conditions for blends of engineering resins Blend

Trade name

Temperature (°C) Barrel

PA/PP PA-6, PA-66/ABS PA/PPE PA-6/PPE PC/ABS PC/ABS PC/SMA PC/TPU or TPU/PC PPE/PS PPE/PS 1. CR: compression ratio

Orgalloy Triax 1000 Noryl GTX Dimension Triax 2000 Bayblend T Arloy Texin Noryl Prevex

Die

240-265

210-240 260-290 200-230

Screw Melt

L/D

CR1

220-290 250-265 277-288 265-295 215-240 220-240

>20

2.5-3

20-24

3-4

>24

3

260 200-230 190-220 240-270

Polymer Blends Forming

less sensitive to temperature variation. Owing to high thermal stability, the blend can be processed over a wide temperature range, T = 220 to 290°C. The single screw extruder barrel should be L/D • 20, fitted with a standard three-zone screw, with compression ratio 2.5 to 3. Excessive shearing may cause destruction of morphology, and thus should be avoided. Purging with PP followed by LDPE is recommended. PA/ABS For PA/ABS blends, (e.g., Alphaloy™ MPA, Elemid™, Diaaloy™ N, Kane ace MUH, Malecca™ S, Maxloy™ A, Monkalloy™ N, N5, Novaloy™A, Stapron™ N, SX, Techniace™ TA, Toyolac™ Alloy, Triax™ 1000, Ultramid™), single or twin screw extruders are suitable. However, the latter are preferred. Typical ABS extrusion procedure should be followed. For two-stage screws, a first stage compression ratio of 2.5-2.7 and pump ratio 1.5-2 are recommended. PA/PPE blends (e.g., Artley™, Dimension™, Lynex™ A, Noryl™ GTX, Remaloy™, Ultranyl™) may be processed in a single screw extruder, using a single-stage metering screw with compression ratios = 2.2 to 3.5, and metering sections from 3 to 10 flights long. Two stage screws should have compression ratio of 3:1 with second stage metering approximately 1.6 to 1.7 times deeper than the first stage metering. A reverse temperature profile may increase output. PS or ground acrylic is best for purging. Polyester Blends PCTG/PC or SMA (Ektar™ MB) Blends of PCTG with PC or SMA are typically injection molded into business machine, appliance and automotive parts but can also be extruded. Films of thickness 0.025 mm can be extruded and sheet thickness up to 3.2 mm is possible when conventional extrusion equipment is used. Typical processing conditions vary depending on film or sheet thickness. PP, HDPE or an acrylic compound is used for purging.

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Polycarbonate Blends PC/ABS and PC/ABS/SAN blends (Alphaloy™, Bayblend™ T, Cycoloy™, Denka™ HS, Dialoy™ C, Exceloy™, Iupilon™, Lynex™ B, Malecca™ P, MB2200, Multilon™, Novaloy™-S, Proloy/ Lexan™, Pulse™, PX, Ryulex™, SC 250, Techniace™ TC, Toplex™, Triax™ 2000, ABS/PC 7901) can be extruded on typical ABS extrusion lines. Both single and two-stage screws are suitable but the latter are preferred. For two-stage screws, the recommended compression ratios are: for the first stage = 2.5 - 2.7 and for the second = 1.5 - 2. Three-stage extrusion screws normally used for PC or ABS are also suitable. The power requirements depend on the screw diameter. For diameters from 45 to 120 mm, the power consumption varies from 10 to 130 kW/kg. PC/SMA blends (Arloy™ 1000, Ektar™ MB, R4 9900) show outstanding processability. Low shear screws without mixing sections are preferred. On sheet extrusion lines, polishing rolls should be set at the highest temperature that does not cause web sticking. For PC/TPU or TPU/PC blends (Texin ™), the screw should have a long gradual transition zone and a long metering zone. Polyphenylene Ether Blends PPE-based blends can be formed into a wide variety of products: cable trays and protection pipes for electrical systems, car and truck rooftops, profiles and pipes for transport of various fluids. The materials can be processed on a large number of extrusion systems, including single screw machines without venting, and twin screw extruders of various types and sizes. Generally, PPE blends are extruded using single-stage screws with constant pitch, L/D = 20 to 25, compression ratios for densifying the melt = 2.2 to 3.5 and metering sections from 3 to 10 flights. The compression zone should be linear, to give a melting process with good stability. Typical zone lengths are given in Table 10.21 [Bussink et al., 1993].

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Table 10.21. Typical zone length for PPE blends extrusion screws Zones

Length (in D)

Feed Compression Metering Total screw length

6-8 5-6 9-6 20

7-9 7-8 11-8 25

Two-stage screws should have compression ratio of 3. The second stage metering should be 1.6 to 1.7 times deeper than the first stage metering. HDPE, PS and ground acrylic are best purging resins. Because of the hardness of the PPE granules before melting as well as high melting temperature and high melt viscosity, high power drive motors are needed. The recommended motor power for single screw extruders with different screw diameters is given in Table 10.22 [Bussink et al., 1993]. Table 10.22. Motor power as function of screw diameter Screw diameter (mm)

Motor power (kW)

45 60 90 120

15-25 25-35 60-70 100-120

PPE/HIPS blends (Gepax™, Luranyl™, Noryl™, Prevex™, Verton™, Vestoran™, Xyron™ 200). Easy flow and good heat stability make these blends suitable for extrusion of sheets, sections, hollow sections and tubes. 10.4.3.3 Extrusion of Blends of Specialty Resins Conditions recommended by the manufacturers for extrusion of specialty resin blends are listed in Table 10.23. For PEI/PC blends (Ultem™ LTX), drying is essential. Furthermore, for extrusion grades, adequate moisture reduction may take 3-4 times longer than for injection molding grades. Excessive die drool is one of several indications that the resin was not sufficiently dried, even though some moisture analyzers may indicate adequate dryness. LCP Blends The effects of extrusion conditions on PP/LCP blends were studied, in particular focusing on the effects the mixing history and post-extrusion drawing have on mechanical properties [Machiels et al., 1997]. Two methods of mixing were used: simultaneous extrusion of both PP and LCP in a SSE and separate feeding of both components in a static mixer. In both cases, post-extrusion drawing was necessary to obtain improved mechanical properties. The morphology of the blends ranged from a skin-core structure (at low draw ratios) to a high aspect ratio fiber/matrix morphology (at high draw ratios). The deformation of the LCP phase was less than affine, indicating slip between the phases. The tensile modulus went from 1 GPa for

Table 10.23. Extrusion conditions for blends of specialty and elastomeric resins Blend

PEI/PC NBR/EPDM

Trade name

Ultem LTX Geolast

1. CR: compression ratio

Temperature (°C)

Screw

Back P

Melt

L/D

CR1

(MPa)

300-340 205-220

24

2-3

13.8

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Polymer Blends Forming

the neat PP to ca. 13 GPa for the blend containing 30 vol% LCP at draw ratio of 15. 10.4.3.4 Extrusion of elastomeric Blends Conditions recommended by the manufacturers for extrusion of elastomer are listed blend in Table 10.23. Articles made of NBR/PVC blends (Carloy™, Geon/Hycar™, Denka™ LCS, JSR NV, Krynac™ NV, Nipol™, Oxyblend™, Paracril™ OZO, Vynite™) can be manufactured by dispersing powdered nitrile elastomers in a continuous mixing operation, which enables them to be used directly in extrusion. Easy and smooth extrusion and calendering are usual. Plasticizers and stabilizers should be pre-dispersed in PVC otherwise they may swell the elastomer and cause formation of lumps. Blends of nitrile rubber with PVC or chlorinated PE, TPU, PA or their combination (Chemigum™ TPE) show extremely high resistance to melt fracture and die swell. The material can be fed at high melt temperature into crosshead to promote strong melt-fusion with rigid plastics such as PVC/ABS blends during coextrusion. 10.4.4. Summary Extrusion is the basic polymer processing operation and is a relatively simple process. Nevertheless, a number of process parameters have a profound influence on final product properties: screw and die geometry, melt temperature and screw speed, to name the most important. Although some blends can be processed in the same fashion as singlephase materials, in general, blends are more sensitive to process variables such as temperature, stress and residence time. The key to performance is to generate in the final part the right morphology. Therefore, blends generally have to be compounded in a twin screw extruder in order to obtain the necessary amount of mixing. The use of single screw extruders is often limited to a secondary step, after compounding, which aims strictly at forming the product. It can however be used to both compound and form in some cases, in particular when the blend components have similar flow behavior.

The die region is extremely important for forming polymer blends. A number of flow induced phenomena can be observed in the die region, such as formation of lamellar morphology under the right conditions and geometry. The convergent flow into the die engenders strong elongational forces. On the one hand, the forces may elongate the dispersed drops into fibers which, downstream from the die, may disintegrate into fine droplets. On the other, crowding flow lines in the convergence may induce coalescence. The net effect of these two competing influences depends very much on the nature of blend — how well its morphology has been stabilized.

10.5

Injection Molding

10.5.1

General Description of Injection Molding

Injection molding is the most widely used forming operation. It is also the most complex in many ways since it is a discontinuous process: there is a large number of process variables and flow conditions rarely reach steady state, causing large transient effects. Optimization of the process for a given product requires adjustment of several parameters such as pressure and temperature, each taking different values at different stages of the process. The process cycle can be broken down into different stages. Plastication. First, the thermoplastic resin is fed from a hopper into the barrel of a single screw extruder in which the resin is softened (or plasticized) by the action of the rotating screw. Injection. The molten polymer is forced through a nozzle by the forward action of either a piston or the screw itself acting as a piston. The resin enters through the nozzle and a series of channels, the runners, a number of cavities in a colder mold. It should be noted that in the case of thermoset injection, the mold will be hotter than the injected resin. Cooling. Once the resin is in the mold, it solidifies under pressure. After a predetermined cooling time, the mold opens and the part is

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ejected. In the case of thermosets, the mold is opened after the curing period. Contrarily to the metals used in die casting, polymers are highly viscous, requiring large pressures to force them to flow into the mold cavity. Once the mold is filled and solidification starts, high pressure must be maintained to pack more melt in the cavity to offset polymer shrinkage and achieve an accurate reproduction of the mold. Moreover, owing to the notoriously low heat conductivity of polymers, injection molding requires significant time to heat up the material and cool it down. Injection equipment consists essentially of two parts: the injection and clamping units. The injection unit melts the polymer and injects it in the mold. The vast majority of modern injection presses use a reciprocating screw. As in an extruder, the screw rotates, melts the polymer and pumps it. Furthermore, in most injection molding units, the whole screw can also axially move forward, acting like a plunger to inject the resin into the mold cavity through runners and gates. The beginning of the injection cycle is marked by the forward motion of the screw to fill the mold with melt while a non-return valve at the end of the screw prevents the polymer from flowing back in the extruder barrel. Following the injection, during the holding time, the screw maintains the injected melt under pressure (hold pressure). This pressure pushes in additional melt, as the molten polymer cools down and shrinks in the mold, to offset contraction. The pressure must be sufficiently high to avoid sink marks, but low enough to ensure easy removal of the part from the mold. When the gates eventually freeze, isolating the mold from the injection unit, the screw starts to rotate and move backward. It melts new volume of resin and pumps the melt into the cylindrical space in front of the screw, until there is sufficient material for the next shot. The chronology of the development of the injection molding process was discussed in Chapter 1, Introduction to Polymer Blends. Injection molding of polymers began in the 1870’s. It was adapted from the die casting process to form diverse articles from nitrocellulose, NC. The

early machines already had the essential features (although in primitive forms) of the modern injection presses: heated (with steam) cylinder, plunger (not screw), nozzle and mold. Phenolics and cellulose acetate, CA, were injection molded in the 1910’s, as soon as they were developed. PS was already formed by this method in 1931. It is only in the 1940’s that screws began to systematically replace rams in injection presses. Developments that followed include increase of clamping pressure and shot size as well as evolution of electronic control systems. More recent developments focused on computer aided design of parts and molds, as well as on closed loop process control to increase reliability and precision. With the availability of robots, injection molding is the most automated forming process. Automated plants, producing hundreds of tons a day of injection molded parts, operate 24 hours a day with minimal human intervention. Modern injection molding has evolved into a wide and varied, and often complex, range of processes. Co-injection to produce multi-layer parts with inserts is now common. Thermosetting polymers, elastomers and hybrids (thermoplastic/ thermoset systems) have been injection molded. For example, the reaction injection molding, RIM, process is widely used to inject liquid reactive ingredients to produce polyurethanes, polyamides, or polyesters. The possibility of generating cellular structures and varying polymer rigidity makes the process flexible, capable of producing a wide variety of products. Reinforced RIM has been used to produce large panels for the automotive industry. Other recent modifications include lost-core injection molding and gas-assisted injection. The former involves inserting in the mold a low melting metal alloy insert over which the polymer is injected. After solidification and ejection of the part, the insert, which has a melting point lower than that of the polymer, is melted away to yield a hollow part. Gas injection also makes it possible to manufacture hollow parts. After injection of a small volume of polymer, but before its solidification, gas is injected in the mold through a different gate. The gas pushes the molten material

Polymer Blends Forming

to adhere uniformly to the mold, producing finished products with a hollow central part. The quantity of injected polymer and the rate of gas injection are the principal critical parameters of the process. 10.5.2

Processing Parameters

The properties of injection molded parts depend strongly on processing conditions. In addition, the dependence is usually more pronounced for polymer blends. The determining factors for moldability are injection pressure, injection speed, residence time of the material in the cylinder and plastication performance. Plastication of new material takes place during the cooling phase of the previous shot in the injection cycle. Therefore, plastication performance has a strong bearing on the cycle time. An unsatisfactory plastication performance means an increase in cycle time. On the other hand, if the plastication is too quick, a waiting time has to be introduced in the cycle, bringing about risks of thermal degradation of the material. The influence of the most important parameters is discussed below. 10.5.2.1 Pressures Pressure varies widely during the injection process. In addition, the pressure is different depending on where it is measured: at the nozzle, in the runners or in the cavities. The injection cycle is schematically shown in Figure 10.9, which shows the time dependency of pressure measured at two different locations: nozzle, Pn, and mold cavity, Pm. The cycle is usually described in terms of pressure-time stages. Stage 1 begins when the mold closes and the plunger/screw starts moving forward. This corresponds to a period known as dead time since Pn immediately increases, whereas Pm = 0. The filling stage (stage 2) starts when the material reaches the cavity and Pm starts to increase. Pn then reaches a predetermined plateau: the injection pressure. The next step (3) is the packing phase. It consists of a period of time during which very little material is being pumped into the mold cavity, to compensate

701

for thermal shrinkage of the cooling polymer. Thus Pm is maintained at the packing pressure, near the maximum pressure level. Once the gate has been sealed by the solidification, the mold is isolated from the injection system. Pm then drops rapidly. When the mold pressure reaches the established minimum value, the mold opens (5), thus Pm = 0, and the part is removed. At the same time, the screw pulls back. Pn shows a low but non-zero value: the back pressure.

Figure 10.9. Typical pressure profiles during injection molding. Pn and Pm stand respectively for nozzle and mold cavity pressure. Stages are: 1 - dead time, 2 - filling, 3 - packing, 4 - cooling, 5 - ejection.

The injection and packing pressures depend on material type, quality of tools, and configuration of the parts. For example, the PPS/HIPS blends with their low viscosity and heat resistance need lower pressures than the high heat resistant ones. This statement is valid for both the injection and packing pressures. Both must be chosen so that efficient forming pressure, sufficient for filling the mold and avoiding flaws, is maintained. Packing pressure affects the final dimensions and formation of sink marks and shrinkage. It should be applied and controlled until the gate area is completely frozen off. During the filling stage, it is generally recommended to set the machine pressure near its

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M. M. Dumoulin

maximum and control the speed of the ram with velocity controls. This helps to fill the cavity at a uniform rate (despite differences in viscosity due to temperature fluctuations or other factors), and as a result to provide consistent parts. It is recommended to switch the molding control from filling to packing on the basis of the ram position or cavity pressure. It is important to effect the change when the part is 95-98% full, otherwise the high pressure of the first stage will cause the part to flash and may cause damage to the mold. Switch-over based on hydraulic pressure and time is recommended if ram position or cavity pressure controls are not available. While maintaining a cushion, packing pressure should be set at the lowest level necessary to properly pack out the part to achieve the desired surface aesthetics. Under-packing the part will result in excessive sink marks, poor dimensional tolerances and poor surface aesthetics. Over-packing will result in excessive molded-in stress, increased part weight and poor part performance. In injection molding, the pressure at the tip of the screw is called back pressure. It is generated by mechanical obstacles (valves, gates, or dies) against which the material is forced to flow. Thus, it is the back pressure during extrusion that causes the material to mix. Its magnitude is particularly important while the screw is being pulled back after the injection shot. Back pressure is important to insure a homogeneous melt and maintain consistent shot size. Higher back pressure increases and improves melt mixing. The magnitude of back pressure varies greatly with design of the plasticating screw and the material being injected. It is best to use the minimum level of back pressure, just sufficient to provide a good quality, homogeneous melt. Too small back pressure may result in entrapment of air in the melt causing splay and short shots. Too large back pressure may lead to excessive shear stress, causing material degradation and splay. To optimize the conditions, it is best to start out with minimum back pressure and gradually increase it checking the quality of moldings. Decompression (negative back pressure, or suck-back) tends to pull air through the nozzle

into the plasticating extruder. The air, when injected into the cavity, may result in splay. It is recommended that no decompression be used. However, when it is required to reduce drool, minimum decompression should be used. If excessive amount of decompression is needed a different checking valve should be used. 10.5.2.2 Processing Temperatures In the injection molding process, setting the temperature involves optimization of the temperature profile of the plasticating unit (extruder barrel), temperatures of the runners and gates, (all these determine the molten polymer temperature) as well as the mold temperature. The temperature setpoints depend on the material type (viscosity profile, thermal and shear stability, thermal properties) as well as machine or process considerations (machine capacity to shot size ratio, screw design, mold and part design, cycle time, etc.). Temperatures of the two basic units, the injection system and the mold, should be discussed separately since their selection stems from very different considerations. The optimum melt temperature represents a compromise. It should be as low as possible, to ensure minimal thermal degradation of the material. On the other hand, it should be as high as possible, to allow the easiest possible flow, thus necessitating lower pressures. It will therefore depend directly on the part shape and geometry of the runner system and mold. The higher the melt temperature, the lower the polymer viscosity, thus the more easily the material flows. Melt temperatures in the upper end of the recommended range are necessary when processing parts with thin walls, small gates or long flow lengths. Thermal degradation of the material is a function of melt temperature and residence time. To reduce the risk of degradation, it is recommended that cycle times be optimized and that the shot size be kept between 40% and 75% of the total barrel capacity. This is particularly important when producing large parts that tend to have longer cycle time.

703

Polymer Blends Forming

Mold temperature is critical by its impact on final properties and plant economy. The cooling phase is the major contributor to the overall cycle time. Therefore, low mold temperature will ensure faster cooling and shorter cycles. Uniform heat removal is critical for securing the best properties and consistent dimensional tolerances. Mold temperatures in the upper end of the recommended range may be necessary for thin wall parts or large parts to avoid early freezing and ensure proper mold filling, and parts that require dimensional stability and a low level of moldedin stress. Mold temperatures in the lower end of the recommended range will result in shorter cooling times, less initial shrinkage and higher molded in stress. 10.5.2.3 Injection Speed The optimum injection speed depends on the runner system, the gate, as well as the size and complexity of the part. Generally, slow to moderate injection speeds give the best results. Too fast injection may result in excessive shear heating, degradation, delamination and splay. However, average to high injection speeds are sometimes required to avoid early freezing of the melt, e.g., in such blends as for PPE/HIPS and PPE/PA. For small parts, with relatively simple geometry, a constant fill speed is recommended. For large or complex parts, it is often helpful to profile the injection speed. The injection speed should

be optimized so as to form a symmetric flow front inside the mold cavity. Some polymers, e.g., those with high molecular weight and/or rigid macromolecules, show high susceptibility to stress (mechanical degradation). In these cases, to avoid compromising the mechanical properties of the finished product, the rate of injection must be carefully optimized and controlled. 10.5.2.4 Screw Geometry and Speed As in extrusion, the homogeneity of the melt and reproducibility of the screw plastication efficiency depend on the screw geometry and processing conditions. The melt temperature is affected by three factors: barrel temperature (which is controlled), screw design, and screw speed. Since polymers have low heat transfer coefficient, the screws must be designed to maximize the heat transfer. Screw designs vary considerably from one polymer to another depending on, for example, the heat stability of the resin. Shorter lengths are used for resins that tend to degrade. Screws for cyclic operations, such as in injection molding units, are generally 40-50% shorter than screws used for extrusion. Table 10.24 lists some typical screw geometries for common plastics. It should be noted that the geometry can vary depending on the rheology of a particular resin. Because glass reinforced grades often give wear problems, bimetallic liners (Xyloy™ 800) should be used. Abrasive wear occurs primarily

Table 10.24. Recommended screw length (expressed in L/D ratio) for injection molding Polymer

Total (L/d)

Feed zone (L/d)

Compression zone (L/d)

Metering zone (L/d) (L/d)

Rigid PVC HIPS LDPE HDPE PA-66

16-20 20 20 20 20

3-5 6 5 8 15

13-17 4 10 4 1

0 10 5 8 4

Compression Ratio 1.5-2.0

2.5-3.0

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on the lands and edges of the screw flights, and then with time, the root will wear out in the transition and metering zones. Use of treated and stress relieved alloy steel with a hard surface finished with an abrasion resistant coating (e.g., Colmonoy™ 56) is recommended. Similar precautions against wear should be included in the design of screw tips and check valves. Generally, slow screw speeds are recommended to provide homogeneous melt without excessive shear heating. Screw speeds should be set so that plastication is complete two to three seconds before the cooling step of the preceding shot is terminated and the mold begins to open. This allows for a minimum screw speed without causing an increase in the total cycle time. 10.5.3

Injection Molding of Polymer Blends

The majority of commercial blend grades has been developed specifically for injection molding and thus can be processed in conventional machines as long as the processing conditions recommended by the blend manufacturer are followed. Reciprocating screw machines are preferred over plunger-type ones. Phase morphology is strongly affected by the type of flow. Injection molding gives rise to a number of phenomena affecting morphology: high shear strain and rate in sprues and runners, fountain flow in mold, presence of weld lines, non-uniform heat history throughout the sample. Figure 10.10 shows schematically two of these phenomena: deformation of the advancing front during flow in an injection mold and the presence of weld lines. As shown on the schematic, fountain flow causes volume elements, upon reaching the flow front, to move towards the outer edge and to deform gradually into elongated shapes. This gives rise to the phenomenon of skin-core morphology whereas the skin of molded parts shows different crystalline and phase morphologies than the core region, as depicted in Figure 10.11. The morphology of multiphase systems will reflect these deformation patterns. Thamm examined the phase morphology in injected PP/EPDM blends [1977]. PP/EPDM injected

Figure 10.10. Schematics of the advancing front in an injection mold, showing material’s deformation and the presence of weld lines.

Figure 10.11. Schematics of skin-core effect on phase morphology in injection molded part.

samples showed a skin-core crystalline morphology. At the outer skin layer, the molecules were oriented and non-spherulitic whereas a spherulitic crystalline morphology was observed in the core region. The thickness of the skin layer was found inversely proportional to melt temperature and injection speed. Phase morphology also varied depending on position. EPDM

Polymer Blends Forming

domains in the skin were larger and oriented in the flow direction resulting in disk-shaped particles. EPDM in the core of the moldings consisted of spherical domains. Moldings of polymer blends have often showed higher susceptibility to weld lines. Weld lines are created in injection molding when the flow front is split and recombined or when multiple gate molds are used. The presence of weld lines in general strongly affects the mechanical behavior of the molded part. The weakness of knit lines has been attributed to the fact that the flow front is generally colder than the bulk, resulting in a less adequate mixing of the two highly viscous streams meeting at the weld line. In the case of blends, there is an additional difficulty. Because of the fountain flow, the dispersed phase reorients in a direction parallel to the surfaces. Thus the weld line constitutes a discontinuity in the morphology of the molded part, where failure might preferentially occurs. Thamm examined the phase morphology in the vicinity of weld lines in PP/EPDM blends [1977]. The presence of weld lines strongly influenced phase morphology. At weld lines, EPDM was in the form of thin sheets parallel to the knit line, thus perpendicular to flow, as illustrated in Figure 10.10. The sheets are assumed to prevent bonding across the interface. Under SEM and optical microscopy, there appeared to be no particles at all at the center of the weld line, caused either by flow segregation or exceedingly high dispersion at the outer edge of the flow front. The generation of lamellar morphologies through a modified co-injection process was also shown [Barger et al., 1993]. The process included a feedblock producing from the two melt streams three layers that fed a series of layer multipliers. These multipliers subdivided the melt stream into four sub-streams, which then were reoriented, stacked and recombined. The process produced net shape parts with micron size layers. A reasonable viscosity match of the two polymers, and low interfacial tension coefficient between them, ensured good compositional uniformity throughout the part.

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10.5.3.1 Injection Molding of Commodity Resin Blends Injection molding conditions of blends of polyolefin resins and other commodity resins are listed respectively in Tables 10.25 and 10.26. The conditions represent an average for each type of blend. They usually vary slightly from one grade to another. Recommended practices and specific comments are discussed below for each type of blend. Polyolefin Blends HDPE/LDPE (Hostalen™ GC) blends are injection molded, in particular to produce bottle crates. Various factors affect shrinkage of the blends. Shrinkage is very dependent on wall thickness. Increasing demolding temperature from 20 to 60°C increases shrinkage by up to 0.4%. Wall thickness increase from 2 to 6 mm may cause a shrinkage increase of up to 0.5%. When molding objects with a wall thickness in the range of 4-6 mm, shrinkage decreased by 0.5% when hold pressure time increased to 20 s or when injection pressure increased from 600 to 1400 bars. Increasing melt temperature from 220 to 280°C increases shrinkage by 0.3%. LLDPE/PP Injection molding involves high pressures and high shear stresses, strongly affecting the blend morphology. High dispersion in LLDPE/PP blends, resulting in domains smaller than the resolution of the scanning electron microscope, SEM (0.1 µm), was observed under the high pressure and temperature conditions of injection molding [Dumoulin, 1988]. PP/elastomer Most polyolefin blends to be injection molded are based on elastomer-modified PP (the elastomer is mainly EPDM). The elastomer can be mechanically incorporated, such as in Oleflex™ TPO. However, there is a growing number of PP/EPDM blends prepared directly in the reactor: the rubber modifier is incorporated during PP polymerization (i.e., Kelburon™). The blends are often used in

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Table 10.25a. Injection molding conditions for polyolefin-based blends — Part A Blend

Trade name

Temperatures (°C) Melt

HDPE/LDPE LDPE / PIB PE-MA, PE-AA/PE PP/EPDM PP/EPDM PP/EPDM PP/EPDM PP/EPDM PP/EPDM PP/EPDM/filler PP/elastomer PP/elastomer PP/PA-66 PP/PA-66 or PA-mXD6 TPO EPR/PP

Hostalen GC Lupolen 250H Formion Santoprene Hifax IPCL Vestolen EM Keltan TP Kelburon Modylen Ferrolene Ferro Flex Systemer S Eref Oleflex Polytrope

Barrel

Shrinkage (%)

Nozzle

Mold 1.5-3.0

180-240 230 180-200

177-210 190-215 185-205

188-220

200 200

240-260 240-260

20-60 15 or less 10-80 20-40

0.8-1.4

200

250-260

220-230 200-240 215 250 260-280 180-210 220-225

0.2-0.6

20-40 <40 <40 50-80 20-60

250-245 240-280

0.6-1.6 1.65-1.8 1.0-1.3 1.3

80-120 80 20-40 25-50

0.2 0.9

Table 10.25b. Injection molding conditions for polyolefin-based blends — Part B Blend

Trade name

Injection rate

Screw L/D

LDPE/PIB PE-MA, PE-AA/PE PP/EPDM PP/EPDM PP/EPDM PP/EPDM PP/EPDM PP/EPDM/filler PP/PA-66, PP/PA-mXD6 EPR/PP

Lupolen 250H Formion

Speed (rpm)

Injection

Back

600-700

10

Cycle time Hold

(sec)

18-22 mod (5-8 s)

Hifax Keltan TP Kelburon IPCL Santoprene Modylen Eref

med-fast

Polytrope

med-fast

1. CR: Compression ratio

CR1

Pressure (MPa)

20-24 >20 >20

2-3 2.3-2.8 2.3-2.8

50 100-200

20 max

2. IP: Injection pressure

60-100

25-80 <120 <120 70-80

10-15 10-15 48-65 0.1-1

40-60% of IP2 40-60% of IP2

30-150

0-1

95

65

20

2-3

5-7.5

0.3

30-150

707

Polymer Blends Forming

Table 10.26a. Injection molding conditions of blends of other commodity resins — Part A Blend

PVC/NBR PVC/TPU HIPS PS/BR PS-PMMA/elastomer Styrenic/rubber ABS ABS/elastomer ABS/BA ABS/PB ABS/PVC ABS/PVC ABS/PVC ABS/PA ABS/PA ABS/PA-6 ABS/PBT ABS/PBT ABS/PBT ABS or AES /PC ABS or AES /PC ABS/PC ABS/PC ABS/TPU ABS/PTFE SAN/elastomer SAN/olefinic elastomer SAN/elastomer ASA/AES ASA/PVC ASA/PC ASA/PC SMA/rubber SMA/ABS SMA/PBT Acrylic/ABS Acrylic based

1. MC: machine capacity

Trade name

Vynite Vythene Polystyrol Styroblend Zylar Limera Magnum Lustran Terluran Terluran Enplex Lustran Triax CBE Techniace TA Triax 1000 Stapron N Cycolin Techniace TB Cevian Techniace TC Techniace TC-IR Cycoloy Bayblend Prevail RTP Lustran Rovel Luran S Centrex Geloy Geloy Terblend S Stapron S Cadon Dylark Terlux Acrylite, Cyrolite and XT

Temperatures (°C) Melt

Barrel

Nozzle

160-170 177 180-280 180-280 215-220

150-170 150-170

170 170

218-274 200-270 240-280 210-270 205-210 188-199 200-215 252-275 240 245-270

193-204 240-260

150-170

170

177-200 210-260 232-265

193 230-260 250-266

210-260

230-260

210-260 180-230 240-293 221-277

230-260 210-230 232-300 249-266

240-260

240-290 221-282 210

193-271 200-220 230-250 240-280 250-290 190-200 260-290 260-290 245-275 250-265 245-260 230-260 200-245

238-274 165-195 204-250

252-288 188-204 227-288

230-250 235-257

240-260

Shrinkage (%)

Shot Size (% MC1)

Mold

27-32 10-70 10-70 54 70 27-60 40-80 50-80 50-80 40-60 40-80 50-70 40-60 21-93 40-55 50-65 40-60 60-80 40-60 40-60 50-100 50-100 25-30 65-93 40-80 70 40-80 45-70 25-60 55-75 80 60-80 50-80 66-82 50-74 30-65

0.3-0.7 0.4-0.8 0.4 0.4-0.7 0.6-0.8

0.3-0.5

50-65

0.4-0.6 0.6-0.8 0.8-1.2 0.2-1.6 0.4-0.9 0.6-0.8 0.2-0.9 0.4-0.7 0.4-0.7 0.4-0.7 0.3-0.7 0.5-0.7 0.7 0.3-0.4 0.5-0.7

35-65

0.5-0.7 0.3-0.5 0.5-0.7 0.5-0.8 0.4-0.9 0.52-0.56 0.4-0.7

30-80 40-75

30-80

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Table 10.26b. Injection molding conditions of blends of other commodity resins — Part B Blend

Trade name

Injection rate

Screw L/D

PVC/NBR PVC/TPU HIPS PS/BR HIPS/PPE ABS/BA or PB ABS/elastomer ABS/PVC ABS/PVC ABS/PA ABS/PBT ABS/PBT ABS/PBT ABS/PC ABS or AES/PC ABS/PTFE SAN/elastomer SAN/SOE* SAN/elastomer ASA/AES ASA/PVC ASA/PC ASA/PC SMA/rubber SMA/ABS SMA/PBT Acrylic/ABS Acrylic based

Vynite Vythene Polystyrol Styroblend Vestoran Lustran mod-fast Triax CBE Enplex low-med Techniace TA mid-low Techniace TB mid-low Cycolin Cevian Cycoloy Techniace TC mid-low RTP Luran S Rovel fast Lustran mod-fast Centrex Geloy slow-med Geloy slow-med Terblend S moderate Stapron S Cadon mod-fast Dylark Terlux Acrylite, Cyrolite and XT

CR1

Pressure (MPa) RPM

Injection

Back

Cycle time Hold

(sec)

0.35 0.7 16-20 16-20 20 16-20

18-22

2-3

2.0-2.6

30-55 30-80 30-60

25-75

80-110

60-80 0-0.3

8-10 78-98 78-98 55-95

20-75

0.35-0.7 0.3-0.7 0.5-1.5 0.35-0.7

78-98 0.35 16-20

25-60 30-100

7-10 80-110 mod-high 8-12 8-12

30-55

80-110

30-55

16-22

0.3-1

14-20 60-80

0.2-0.4 0.35-1

2-3

< 3.5

20-100

60-80 0.35-0.7 10-38 70-140

1. CR: Compression ratio * SOE: saturated olefinic elastomer

automotive applications as bumpers (IPCL™), dashboard and glove boxes (FerroFlex™, Ferrolene™). The PP/EPDM blend morphology can be of the matrix/dispersed phase type (Oleflex™) or co-continuous (Santoprene™, Kelburon™), depending on the composition and manufacturing method. Showa Denko manufactures the Ole-

flex™ TPO by means of compounding and reactive processing (crosslinking). The method results in a fine dispersion of α-olefin random copolymer in PP or PE matrix. A wide range of hardness can be obtained by adjusting composition. In general, the blend is processed at conditions similar to polyolefins.

Polymer Blends Forming

PP/EPDM blends (Santoprene™, Kelburon™, Sarlink™) with co-continuous morphology have a high and shear-dependent viscosity. Because of the viscosity, large nozzles, sprues, runners and gates are required to injection mold these resins. The injection and packing pressures also must be high. Due to the excellent thermal stability of these materials, short interruptions are permitted without purging. When purging is necessary, PE or PP can be used. Because of the low tendency of polyolefins to stick, no mold release agents are needed. The IPCL™ PP/EPDM blends comprise 5 to 25 wt% EPDM, 2 to 10 wt% of mica filler, and small amounts (0.01-0.1 wt%) of sulfur and initiator to induce some crosslinking in order to stabilize the morphology. For Hifax™ PP/EPDM blends, a general purpose screw is recommended. The clamping pressure increases with part size. The nozzle should be as short as possible with orifice larger than usual. All tools should be vented: trapped air may cause localized burning, surface sinks and voids. Keeping the moving half of the mold cooler than the stationary part aids demolding. For injection of PP/EPDM blends (Keltan™, Kelburon™), both core progressive and 3 zone screws have been used. The plasticating extruder temperatures should be set to increase progressively from hopper to nozzle, but T • 300°C are to be avoided, since the material will thermally degrade. Screws with long metering zones may also cause thermal degradation whereas short metering ones may lead to inhomogeneous melts. To enhance the surface quality and gloss of filled grades, mold temperature might be increased to 50-60°C. If mold temperature is kept higher, toughness and cycle time will be affected. To avoid shrinkage differences and stress in the product it is important to fill the mold quickly. If direct coloring is intended, then a mixing device should be fitted to the end of the plasticating screw. Injected samples of PP/EPDM blends are known to have the skin-core crystalline morphology. The macromolecules in the skin zone are oriented parallel to the flow direction and non-spherulitic whereas a spherulitic crystallinity is observed

709

within the core [Thamm, 1977]. Skin thickness is inversely proportional to melt temperature and injection speed. Differences in the blend morphology have also been reported. Usually, the EPDM domains in the skin zone are larger, in the form of disk-shaped particles oriented in the flow direction, while in the core zone the EPDM particles are globular. The shape and orientation of the domains are explained by the typical flow patterns of the flow front [Tadmor, 1974]. Contrarily to HIPS or ABS, the elastomer phase in polyolefin blends is not grafted to the matrix and generally not or only slightly crosslinked. The morphology of PO blends is thus complex and more dependent on flow conditions. Consequently, these blends are sensitive to the presence of weld lines, occurring when multiplegate injection molds are used. As mentioned above, at weld lines, the dispersed EPDM phase can be in the form of thin sheets parallel to the weld line (perpendicular to flow). The sheets prevent bonding across the interface at the weld line and may also act as stress concentrators. The orientation gradually curves outward to a position parallel to the mold surface away from the weld line. The spherical shape of EPDM particles in the core region is gradually transformed, as the observation point moves closer to the weld line, into vanishingly smaller, highly elongated streaks. The presence of sheet-like domains confirms that the orientation and shape of the phases originate from the flow patterns at the melt front. The tendency to form sheets increases with decreasing viscosity of the elastomer phase, which correlates with the observation that lower viscosity elastomers yield poorer weld line strength. Another consequence of flow fronts meeting in the mold is that air (that did not escape through mold vents) is entrapped within. This may result in microscopic cracks on the surface of the weld lines caused by the presence of compressed air [Hagerman, 1973]. In three-component blends (PP/HDPE/EPDM), HDPE was reported to be within the EPDM phase, regardless of order of addition. It was assumed that HDPE acts as a viscosity modifier and improves EPDM dispersability [Thamm, 1977].

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A slight crosslinking of EPDM helped to stabilize the morphology against the deforming action of the high shear field imposed on the material during injection molding [Dao, 1982, 1984]. In LDPE/PIB blends (e.g., Lupolen™ 250H) the LDPE has relatively low branching (higher crystallinity). The material is soft and crack resistant. Injected or extruded parts are less transparent and glossy than neat LDPE. Since the molded parts are softer than neat LDPE, low mold temperature and good care had to be used while demolding. PP/PA-66 blends (Akuloy™ RM, Dexpro™, Dexlon™, Eref™ Flexloy™ Gapex™, LAX 23, MCX-Q, NB, Novamid™ AC, Orgalloy™ R-6, Poliblend™ NH, Snialoy™, Systemer™ S, Thermocomp™, Ube Alloy CA, Ultramid™ KR, UTX™) The processing conditions for these blends are comparable to those normally used to process PA-66. However, it should be noted that drying of these blends is much slower than that for the neat PA. Gloss can be improved by raising the mold temperature. PP/PA-6 blends Holsti-Miettinen et al. [1995] examined the permeability and mechanical properties of PP/PA-6 blends, compounded in a TSE then injection molded. The effects of different compositions and blend preparation methods were examined. It was concluded that injection molding could induce lamellar morphology, in particular in the skin region of the parts. The best results were obtained for the compatibilized, dry-blended system. In this case, first, a PP-b-PA block copolymer was prepared by reacting maleated PP (PP-MA) with large excess of PA. Next, the PA containing the compatibilizer was dry blended with PP, and processed in an injection molding press. The resulting parts showed clear lamellar structure and yielded the largest reduction of oxygen permeability. By contrast, the blends compounded in a TSE before injection molding, showed a fine, dispersed morphology that had only a minor effect on permeability. The lamellar morphology also resulted in significant increases in tensile modulus, strength and impact strength.

Ethylene-methacrylic acid or ethylene-acrylic acid copolymer/PE When injecting blends of ethylene-methacrylic acid or ethylene-acrylic acid copolymer with PE (Formion™), it is recommended to keep full water cooling of hopper and throat and ensure that the rear zone temperature does not exceed 190°C to prevent bridging. Styrenics Blends HIPS (Styroblend™, Polystyrol™) Screw type injection molding machines are used for molding impact modified PS (HIPS). Special screws or plungers might be employed only in special cases, e.g., to generate marble-like effects. Settings can be chosen with some flexibility due to the amorphous nature of the resin and its low shrinkage. The mold temperature control system is of great importance because moldings surface properties (gloss, brilliance, absence of flow lines), strength of weld lines, resistance to warpage, and shrinkage all depend on the mold temperature. It is typically set in the range 10-70°C depending on product type and requirements. For example, for rapid molding of thin wall containers the mold may even require temperature below -10°C. In some cases, warpage may be controlled by having both halves of mold at different temperatures. Injection molding temperature is usually 180-280°C with exception of flame retardant grades which require 230-250°C. Melt temperature affects toughness, especially in modified grades. High injection rates yield glossy moldings but at the same time the impact modified grades may loose some of their resistance to high temperature deformation. A designed experiment is a good way to initially determine the temperature settings that provide the most robust process for a particular mold. The recommended melt temperature of 220-275°C with an optimum of 245°C can be determined from an air shot. Keeping a uniform melt temperature is essential to ensure good color matching of mating component parts. A higher melt temperature generally produces parts with

Polymer Blends Forming

better surface and higher gloss, while a lower one reduces cycle time and the risk of thermal degradation. However, excessively low processing temperatures may result in high residual molded-in stresses. Uniform heat removal from the mold is critical for securing consistent dimensional tolerances and the best set of properties. Higher mold temperatures are required for thin walls, large parts, and parts that require dimensional stability and low level of molded-in stresses. Mold temperature has minimal effect on gloss level. The recommended screw back pressures range from 0.3 to 1 MPa assuming a nominal hydraulic intensifying ratio of 10:1. The amount of back pressure needed depends on the design of the plasticating screw. It is best to start with minimal back pressure and increase it gradually to the level that provides a good quality, homogeneous melt. Too low back pressure results in trapped air in the melt, causing splay and short shots, while too high may cause excessive shearing, resulting in material degradation and splay. It is particularly important to have sufficient back pressure when using color concentrates. Operating at the upper end of the recommended range may be necessary to achieve good distributive mixing and uniform color of the moldings. The injection speed should be adjusted, depending on the size and complexity of the part, as well as of the gate and runner system. Generally a moderate-to-fast speed gives the best results. High injection speeds create high shear as the melt passes through the runners, gates, and along the mold cavity surface, reducing the polymer viscosity, thus allowing the material to flow more easily. However, too fast speeds lead to excessive shear heating, degradation, flash or burning. Medium to slow injection speeds usually provide more uniform and lower gloss. It is recommended that during retraction of the plasticating screw no decompression be used. If it is necessary to reduce drool, it may be used at the lowest possible level to avoid sucking back air and inducing splay. During the filling stage, the machine pressure should be set near its maximum, controlling the

711

ram speed with velocity controls. The packing pressure should be set at the lowest level, just necessary to properly pack the part to achieve the desired surface aesthetics. Purging of the machine should be done with neat PS. The impact strength of HIPS was reduced by a factor of two after the material was injected five times [Menges, 1989]. Microscopic examination showed that the morphology had been significantly altered - the dispersed particles had shrunk and the elastomer phase around dispersed PS islands seemed to have buckled. PS-co-PMMA/elastomer (Zylar™) Chrome plated molds are recommended when injecting elastomer modified styrene-acrylic copolymer. Runners, gates and sprues should be smooth to minimize polymer turbulence during injection. Good mold temperature is needed for optimum clarity and surface gloss. Contamination with other materials will cause streaking and haze. Styrenic thermoplastic/rubber (Limera™) Conventional free-flow type nozzles with an orifice diameter of 3 mm or larger are commonly used for Limera™ R. Although screw design is not critical for Limera™ R, small L/D and compression ratio are preferable. Forcibly plasticizing screws such as the Dalmage screw are not recommended. Melt temperature should not exceed 280°C. ABS (Magnum™) The effect of temperature on injection molding of ABS was discussed by Menges [1989]. Increasing the melt temperature beyond the recommended range, as well as increasing the dwell time or repeatedly reprocessing the material, all induced an increase in the position of the relaxation peak of the elastomer phase (enhanced crosslinking), and decreased the impact strength. For different systems, analysis of morphology also pointed toward different mechanisms, viz., aggregation of dispersed particles for material subjected to high shear, destruction of grafting around dispersed particles, etc. High injection rates caused the elastomeric particles within the surface layer

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M. M. Dumoulin

to be highly deformed, resulting in lower impact strength. The optimum temperature profile should be selected according to the ratio of shot size to machine capacity, screw design, mold and part design, as well as cycle time. Barrel temperature profile should be set so the material melts gradually — lower near the feed, hotter near the die zone. However, for vented barrel machines, a relatively flat temperature profile is recommended to ensure that the polymer is melted by the time it reaches the vent zone. Reverse temperature profiles are also sometimes used, primarily to compensate either for improper screw design or with short L/D ratios, and to reduce machine amperage or torque requirements. Elastomer-modified SAN Elastomer-modified SAN (Luran™ S) is highly viscous. All-purpose screws with L/D = 16-20 with a constant pitch of 0.8-1 D over the entire length are suitable. Non-return valves are usually installed. Because of resin viscosity, open nozzles are used, typically equipped with shutoff devices, especially when high back pressures are required. Mold temperature control system is of great importance because surface properties of extruded product (gloss, brilliance, absence of flow lines), strength of weld lines, resistance to warpage and shrinkage depend on the mold temperature control. Mold temperatures of 40-80°C are typically used. In some cases, warpage might be controlled by having both halves of the mold set at different temperatures. The melt temperature is normally 240-280°C. At the upper end of that range, care should be taken to ensure that the residence time is as short as possible to avoid degradation. Increasing the melt temperature increases the impact strength. A relatively high rate of mold filling is advantageous because it entails that the temperature does not drop too much during molding. This result in brilliant finish, strong weld lines and good strength of seams. However, air must be able to escape from the mold cavities during the injection phase, otherwise auto ignition may occur and cause burn marks.

ABS/PS Starflam™ ABS is a flame retarded unreinforced ABS containing brominated PS as the minor phase. The main application is for electrical parts. Reported shrinkage is 0.2-0.35 vol%. The recommended melt and mold temperatures are respectively 220-240°C and 40-50°C. ABS/BA, ABS/PB For ABS blends with acrylic elastomer (polybutylacrylate, BA) or with polybutadiene, PB, (Terluran™), injection molding is similar to neat ABS. Settings are not critical, owing to thermal stability and low shrinkage of the material. When using inserts, these should be heated to 80-120°C before placing them in the mold to avoid moldedin stresses. As for ABS, injection temperature affects the impact strength. A high injection rate yields glossy moldings but at the same time the impact-modified grades may loose some of their resistance, because of modification of morphology imposed by high stress and high temperature deformation. ABS/PVC ABS/PVC blends (e.g., Triax™ CBE, Enplex ™) are formulated to provide a broad window of molding conditions. Because of the presence of PVC, shear or thermal degradation must be minimized. Shot size should be about 50% of machine capacity — less than 35% will result in a long residence time, whereas in excess of 65% will increase plasticating time, both of which create conditions for resin degradation. High compression screws should be avoided since they generate excessive shear heat — a linear taper screw with a compression ratio between 1.5 and 2.5 is desirable. Successful operation depends on the quality of temperature control. Purging with PS, acrylic and SAN is recommended. In Enplex ™ ABS/PVC blends, ABS contains some α-methyl styrene and a methylmethacrylate, MMA, copolymer. An in-line screw injection machine for molding general-purpose ABS, PS and similar resins can be used. The machine should have multi-stage injection speed control unit and a cooling blower unit to control shearing

Polymer Blends Forming

heat in the barrel. To produce a quality product, the shot should be 1/3 to 2/3 of the machine’s maximum injection shot capacity. To prevent sudden compression, the nozzle should be short, open bore, without a shut-off valve. The nozzle should be 20% shorter than a general purpose nozzle, with the nozzle orifice size determined by considering the relationship between size of the molding machine and the shape/weight of the molding product, viz., Table 10.27.

Table 10.27. Nozzle orifice size for ABS blends Molding machine clamping pressure (tons)

Nozzle orifice diameter (mm)

150-250 300-350 400-450 500-550 600-650 800-850 1000-1300

3 3.5-4 5 5-6 6-7 7-8 8-10

A low shear heat screw is recommended, with L/D = 18 to 22 and a compression ratio from 2.0 to 2.6. The screw should be protected against corrosion and wear, by either plating or nitrating. If any scratch is observed on the surface of the barrel, screw, or nozzle, it must be repaired immediately, otherwise owing to resin retention, it may cause burn marks or black spots on the moldings. The resin temperature may be checked with a few air shots. Due to thermal degradation, temperatures above 210°C may cause burn marks, black spots or mold-rust. It is recommended to start the operation at a low injection speed and gradually increase it. If the injection speed is too high, the shearing heat and the adiabatic compression heating (intense heat caused by insufficient venting of the mold, generated when air or the exhaust gases are compressed at the final stage of mold filling) may lead to high temperature, causing burn marks or

713

mold-rust. ABS or PS may be used to clean the barrel. ABS/PA, ABS/PBT When engineering resins are added to ABS, such as in ABS/PA (Techniace™ TA) or ABS/PBT (Techniace™ TB, Cycolin™) blends, mold shrinkage increases significantly from about 0.4-0.6% to above 1%, as for ABS/PA-6 Stapron™N blends. (These values should be compared with values below 0.01% reported for shrinkage of ABS/elastomer blends). Thus, the mold must be carefully designed taking this high shrinkage into account. Stapron™N blends use proprietary chemical compatibilization method, providing the system with stable morphology. This makes it possible to mold these materials within a wide range of processing variables, facilitating optimization of the process. ABS/PC For ABS/PC blends (Cycoloy™, Techniace™ TC), L/D = 20 screws, having a compression ratio = 2.5, with short feed zone (5 flights), long compression zone (11 flights) and a gradual taper leading to a short metering zone (5 flights) are recommended. Screws that increase the melt temperature above that of the cylinder should be avoided, since such a situation may lead to PC decomposition. High back pressures may result in an increase of the melt temperature and improvement of melt mixing. Mold shrinkage depends on flow characteristics, molding conditions and mold design. ASA/PVC In the case of ASA/PVC blends (Geloy™), the melt temperature should not exceed 200°C for more than 1 minute, to avoid PVC degradation. Furthermore, the total time at the recommended processing conditions should never exceed 5 minutes. PS is the best purging agent. ASA/AES blends (Centrex™) have outstanding weatherability. Their excellent stability suggests their use as a protective layer for other polymers. Smooth uniform fill at moderate to fast rate is recommended. To control sink marks, the ram

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Table 10.28a. Injection molding conditions for polyamide-based blends — Part A Blend

Trade name

Temperature (°C) Melt

PA/elastomer PA-6/elastomer PA-6/elastomer PA-6/elastomer PA-6, PA-66, PA-612/ elastomer PA-66 /elastomer PA-66/elastomer PA-66 or PA-66/6 + elastomer PA-6/6T+ elastomer PA-66-co-6/elastomer PA/PP PA-6/PO PA-66/PO PA-6/ABS or PA-66/ABS PA-6, PA-66/ABS PA-6/BA copolymers PA-6 or PA-66/Acrylic PA-6/aromatic- aliphatic PA PA-12/aromatic- aliphatic PA PA-66 or PA-612/ionomer PA-6/PPE PA/PPE

Star Grilon A Ultramid B Fiberfil TN Zytel ST

260-290 230-265 270-290 270-280

Fiberfil TN Ultramid A Technyl Ultramid T Ultramid C Orgalloy Akuloy Akuloy Novalloy-A Triax 1000 Durethan Zytel FN Grilon Grilamid Zytel 71G, 77G Dimension Noryl GTX

Barrel

Nozzle

200-224

50

Shrinkage (%) Mold 60-90 70-90 38-90 50-85

0.4-0.9 1-2.5 0.3-1.2 1.4-2.2 0.8-1.5

280-305 270-300 230-290

55-95 40-60 40-80

2-2.8 1.4

340 270-290 250-300 235-250 260-280 240-260 250-270 240-280 225 250-300 290-305 255-310 280-305 285-315

60-90 80-90 20-30 50-70 50-70 60-80 20-90 70-120

0.4-0.8 1-1.1 0.7-2 0.3 0.3 0.2-0.7 0.8-0.1 1.6

80 38 70 80-100 55-90

0.7-1.0

232-254

243-254

Shot size % MC1

40-70

30-70

0.2-0.6

10-70

0.4-1.6

40-80

1. MC: machine capacity

forward time should be low. Feed cushion should be set at minimum. ASA/PC (Terblend™S) In ASA, an acrylic ester elastomer, which forms very finely distributed domains, is grafted onto the SAN matrix. When blending ASA with PC, the latter component is homogeneously mixed with the elastomer, thus SAN remains as a matrix with PC/acrylic elastomer forming the dispersed phase.

The recommended melt temperature range is 260-290°C, with the optimum impact resistance reported for parts molded in the upper range (280 to 290°C). Moderate injection rates result in high gloss. Shrinkage for these blends is comparatively low and almost isotropic. After-molding shrinkage is usually negligible, and the highest value recorded was 0.5%. Flame-retardant compositions should be processed under milder conditions (low shear and short residence time).

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Polymer Blends Forming

Table 10.28b. Injection molding conditions for polyamide-based blends — Part B Blend

Trade name

Injection rate

Screw L/D

PA-6/elastomer PA-6/elastomer PA-66/elastomer PA-66, PA-666 /elastomer PA-66-co-6/ elastomer PA/PP PA-6, PA-66 /ABS PA-6/BA copolymers PA-12/aromatic aliphatic PA PA-6/aromatic aliphatic PA PA-66, PA-612 / ionomer PA/PPE PA-6/PPE

Grilon A Ultramid B Ultramid A Technyl Ultramid C Orgalloy Triax 1000 Durethan Grilamid Grilon Zytel 71G, 77G Noryl GTX Dimension

CR1

18-22 18-22 28-32 18-22 15

2.5-3

18-20

2-3

high high 16-20

Pressure (MPa) RPM

Injection

50-100

medium

Back

Cycle time (sec)

0.35-0.7 >150

2-3 30-80 50-100

100-120

0.3-0.6 0.35-1

41 35-140

45 0.35-0.7

fast

3.5-12

1. CR: Compression ratio

ABS/TPU Forming conditions for ABS/TPU blends (Prevail™) are not drastically different from those for ABS. The recommended melt temperature is 210°C. Keeping a uniform melt temperature is essential to ensure good color matching of mating component parts. The shot size should be between 40-75% of the machine capacity to avoid thermal degradation. This is particularly important when producing large parts that tend to have longer cycle times. As usual, the injection speed depends on the size and complexity of the part, as well as on the gate and runner system. Slow to moderate speeds give the best results — too fast speeds result in excessive shear heating, degradation and splay. For small parts or parts with relatively simple geometry, a constant filling speed is recommended, while for larger or more complex parts, it may be helpful to profile the speed. Mold

shrinkage of ABS/TPU blends is average between that of ABS and TPU. For Prevail™ the shrinkage is 0.5 - 0.7%, whereas for ABS and TPU it is respectively 0.4 - 0.6 and 0.9 - 1.3%. PS should be used for purging the machine. Blends of Other Commodity Resins Cyro produces a family of acrylic-based resins (Acrylite™, Cyrolite™ and XT). All types of injection molding machines can be used to mold these blends. If reciprocating screw machines are used, the screw compression ratio should not exceed 3.5 (to avoid material breakdown). Injection temperature should be the lowest at which the part will fill. Since the material has sparkling clarity, great care should be taken to avoid contamination.

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Table 10.29a. Injection molding conditions of PC-based blends — Part A Blend

Trade name

Temperature (°C) Melt

PC/BA-co-MA PC/ABS PC/ABS PC/ABS PC/ABS PC/ABS PC/ABS/SAN PC/ABS or AES PC/SMA PC/PET PC/PET or PBT PC/PET PC/PET/PB or PC/PBT/PB PC/PBT or PET PC/PET/ABS or PC/PET/HDPE PC/TPU or TPU/PC

Makrolon Pulse Cycoloy Triax 2000 Bayblend T Novalloy-S Bayblend Techniace TC Arloy Ektar MB Sabre Makroblend Stapron E

270-310 274-302 250-300 260-280 240-280 240-280 220-280

Xenoy Makroblend Texin

Barrel

Nozzle

220-290

235-300

Shrinkage (%) Mold 50-90 82-96 55-100 70-90 80 60-80 80-100 40-60 65-100 15-65 60-71 37-77

0.5-0.7 0.6-0.8 0.7-0.8 0.5-0.8 0.5-0.7 0.5-0.8 0.4-0.7 0.7 0.4-0.7 0.6-0.9 0.5-0.7 0.5-0.7

210-260

230-260

270-275

270-275

245-265

258-268

240-280 240-270

227-275

240-277

50-90 80-100

0.5-1.0 0.1-0.8

215-240

210-246

219-246

15-50

0.8

275-305 271-332 248-282 260-277

Shot size % MC1

40-75 40-80 50-70

50-80

1. MC: Machine capacity

Table 10.29b. Injection molding conditions for polycarbonate-based blends — Part B Blend

PC/BA-co-MA PC/ABS PC/ABS PC/ABS/SAN PC/ABS, PC/AES PC/SMA PC/PET PC/PET, PC/PBT PC/PET/ABS, PC/PET/HDPE PC/TPU, TPU/PC

Trade name

Makrolon Triax 2000 Bayblend T Bayblend Techniace TC Arloy Makroblend Xenoy Makroblend

Injection rate

mid mid low-mid

mid-high

2. IP: injection pressure

Pressure (MPa)

L/D

CR1

RPM

Injection

<15

2-3

50-75 moderate 40-70 40-70

70-105

18-20 18-20

slow-mid

Texin

1. CR: Compression ratio

Screw

70-105 70-105 78-98

Back

Hold

0.2-0.3 0.35-0.7 0.35-0.7

50 50-75

70-140

0.7-1 0.35-0.7 50-75% of IP2

20 18-20

2.2 2-3

25-60 40-70

60-80 70-140

0.15-0.35 60-80% of IP2 0.35-0.7

15

2-3

40-70

40-100

0.35-0.7

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Polymer Blends Forming

Table 10.30a. Injection molding conditions of blends for other engineering resins — Part A Blend

Trade name

Temperature (°C) Melt

PET/elastomer PBT/elastomer PBT/ABS PBT/ASA PBT/PET PCTG/PC or SMA PPE/PS PPE/HIPS or PS PPE/HIPS POM/elastomer POM/TPU POM/TPU POM/TPU

Rynite Ultradur Novalloy-B or Cevian Ultrablend S Valox Ektar MB Noryl Prevex Luranyl Delrin Hostaform S Hostaform C 9021 TF Ultraform

Barrel

Shrinkage (%)

Nozzle

Mold 0.2-1.0

243-260

260-271

227-295

260-300

90-120 40-70 60-80 60-100 66-93 23-93 65-100 65-95 60-100 65-90 65-90 60-90

1.0-2.0 1.8-2.0

260-300 250-275 240-260 250-280 260-271 265-304 230-330 220-290 260-300 175

190-220 170-200 170-200

Shot size (% MC1)

0.2-0.9 30-80 0.5-0.7 0.5-0.7

30-70 40-60 max 50

180-220

1. MC: Machine capacity

Table 10.30b. Injection molding conditions of blends for other engineering resins — Part B Blend

Trade name

Injection rate

Screw L/D

PET/elastomer PBT/ASA PBT/elastomer PBT/PET PPE/PS PPE/PS PPE/HIPS POM/TPU, PTFE POM/elastomer POM/TPU

Rynite Ultrablend S Ultradur Valox Noryl Prevex Luranyl Hostaform Delrin Ultraform

1. CR: Compression ratio

fast med-fast high

IP: Injection pressure

Pressure (MPa) Speed (rpm)

18 16-20 16-20 20

Injection

Back

Hold (% of IP)

55-70

0.2-0.7 0.2-0.7 0.35-0.7

60-80

16-20 50-100 20 16-20

50-100

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M. M. Dumoulin

10.5.3.2 Injection Molding of Blends of Engineering Resins Injection molding conditions for selected blends of engineering resins are listed in Tables 10.28 to 10.30. The conditions represent averages recommended for each type of blend. These should be judiciously varied for each particular grade. Recommended practices and specific comments are discussed below for each type of blend. Polyamide Blends PA is the oldest engineering resin. It offers outstanding mechanical performance, easy processability and good chemical resistance. It is being widely used in injection molded parts for automotive, appliances, sports and leisure applications. There is consequently a large number of blends based on PA, particularly the impact modified compositions. For example, the Zytel™ family of PA’s include PA-6, PA-66 and PA-612, all of which are available in elastomer-modified versions, blended with ionomers, elastomers or fluorinated polymers. Because of the high energy required for melting PA, a shallow metering screw gives the best results. Zytel™ FN is a recently introduced plasticizer-free, low modulus composition with good low temperature toughness, resistance to thermal aging and to solvents (particularly to fluorocarbon refrigerants). It is a flexible blend containing less than 50 wt% PA-6 or PA-66 and an acrylic elastomer. The blend is more temperature sensitive and requires low residence times, thus larger shot size than other grades. Mold shrinkage for compositions containing PA is generally high. Table 10.31 illustrates the effect of composition on mold shrinkage of PA-6 or PA-66 blends with elastomers (e.g., Fiberfil™TN). Addition of fillers reduces shrinkage significantly. PA/elastomer Elastomer-modified PA compositions (Zytel™, Ultramid™ B, Ultramid™ C, Technyl™, Grilon™ A) are easily processable on conventional machines with single flight, with three-section screws, and

L/D = 18 to 22. For PA-66-co-6/elastomer compositions (Ultramid™ C), warpage may be controlled in some cases by having both halves of the mold at different temperatures. Because of the high temperature required, at least one band heater, rated at 200-300 W, is required for the nozzle where severe heat losses to the mold may occur as a result of radiation and conductivity. The feed characteristics can often be improved by controlling temperature in the hopper area at about 80°C. Non-return valve is subject to excessive wear that can be reduced by wear-resistant coatings. In barrel design, it should be considered that PA requires twice as much heat than PS to bring it to the molding temperature (700 kJ/kg, compared to 350 kJ/kg). The nozzle design is essential, e.g., wrongly positioned thermocouple, too long, or poorly designed nozzle may lead to local overheating. Table 10.32 lists the general impact of some parameters on the properties of impact modified PA. Surface properties, degree of crystallinity, resistance to warpage and shrinkage, all depend on the mold temperature. In some cases, warpage may be controlled by having each half of the mold at a different temperature. Typical mold temperature range is 60-80°C, but if good finish, greater hardness and strength are required, it should be increased to 70-90°C. In exceptional cases, the mold can be as high as 120°C. Some properties of the moldings can be modified by post-treatment. Internal stress generated during molding, can be dissipated either during their use or by means of controlled conditioning. The stress relaxation process increases impact strength but also contributes to dimensional variability, although well controlled conditioning makes it possible to reduce the latter problem. The conditioning is usually carried out in mineral or water soluble oils at temperatures 10-20°C above the maximum service temperature, e.g. 175°C. Parts thicker than 5 mm are treated for 45 min, while those thinner than 5 mm for 20 min. Cooling must be gradual to prevent thermal shock. Freshly molded PA’s are dry and they will absorb moisture, increasing in dimensions and

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Polymer Blends Forming

Table 10.31. Mold shrinkage of PA-6 or PA-66/elastomer blends Fiberfil™ TN grade

Description

NY-7 J-7/13 J-7/33/IT NY-8 J-8/13 J-8/33/IT

Toughened PA-6 PA-6, fiberglass PA-6, fiberglass, and elastomer Toughened PA-66 PA-66, fiberglass PA-66, fiberglass, and elastomer

Melt temperature (°C)

Shrinkage (%)

270-280 270-280 270-280 280-305 280-305 280-305

1.4-2.2 0.3-0.5 0.1-0.3 2.0-2.8 0.4-0.7 0.1-0.3

Table 10.32. Effects of some process parameters on PA/elastomer blends Increase in or effect Crystallinity Deformation Appearance Internal stress Shrinkage Mechanical properties Molding rate Viscosity Cycle time

Melt temperature Holding pressure

Holding time

Injection speed

+ + -

+ -

+

+ + -

Mold temperature +

+ + -

decreasing in rigidity. These changes are slow. It takes typically 6 months for 2 mm thick objects to reach equilibrium. Soaking in water accelerates the process, especially at higher temperatures. PA/PO PA’s are also blended with polyolefins (Orgalloy™R, Eref™, Akuloy™). Orgalloy™ R 6000 and R 6600 consist of PP blended respectively with PA-6 and PA-66. An elastomer is also incorporated in some grades. The morphological structure of Orgalloy™ is different from that observed for mechanical blends of PA/PP. The flow properties are similar to PP, thus improved over PA. Lower injection pressures are needed than those for PA.

+ + +

+

Owing to its high fluidity, Orgalloy™ is well suited for processing in standard injection molding machines. Three-zone polyamide-type screws are preferred, with L/D = 15 and compression ratio of 2 to 3. The clearance between screw and barrel should be low to ensure the best mold cycle repeatability. Due to high thermal stability, depending on the grade, the blend can be processed over a wide range of temperatures: from 250 to 300°C. It is generally advisable to use high process temperature and low injection pressure to achieve good mechanical properties, in particular impact strength.

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Mold temperature strongly affects the product properties. Reduction of mold temperature increases orientation and internal stress, while increasing it augments the shrinkage, gloss, duration of dwell, cycle time and flow length. Usually, mold temperature should be 20-30°C for standard grades and 50-80°C for the glass fiber reinforced ones. Dwell (or packing) pressure periods always demand close attention when processing alloys of semi-crystalline polymers. An increase in degree of compaction causes increases in strength of weld lines, flashing, mass of molded part, risk of overfill and decreased shrinkage, sink marks and inclusions. It is interesting to note that shrinkage can be reduced by lowering the dwell pressure, mold temperature, material temperature and component thickness. In addition, it should be noted that in most moldings PA/PP blends show anisotropic shrinkage, lower in machine direction (0.7-1.5%) than in transverse direction (1.2-2%). Blends produced under the Eref ™ trade name include both alloys rich in PP and alloys rich in PA where PA is PA-66 or PA-mXD6 (a semicrystalline PA, made by polycondensation of m-xylylenediamine and adipic acid). A compatibilizer is added in order to provide optimal adhesion between PP and PA. Injection molding machines used for injecting PA are recommended, with peripheral screw speeds of 3-10 m/min. Cooling time can be estimated as 2.5×e seconds where e is the wall thickness in mm. Melt temperature for flame retarded grades must be well controlled and not exceed 270°C.

comprise mainly PA-66 whereas PA-6 is used for Dimension™ blends. Two developments led to this technology: (1) thermal stability of PPE was improved by controlling the molecular weight, and (2) adequate compatibilization methods were found. For injection molding of PA/PPE blends, presses with plasticating screws having L/D = 18-20, and a compression ratio = 2 to 2.5 (depending on screw diameter) are recommended. Injection pressure depends on several variables such as melt and mold temperature, part geometry, wall thickness, flow length. Injection pressure controls the filling of the part and should be applied for 90% of ram travel. PS or ground acrylic is best for purging.

PA/PS The Starflam™ family consists of blends of PA-6 or PA-66 with brominated PS as dispersed phase (Starflam™ PA66, Starflam™ PA6) used in flame-retarded compositions. The moldings find applications as automotive and electrical parts. The processing conditions are similar as those for the neat PA used as the matrix material.

PET/elastomer For PET/elastomer blends (Rynite™), general purpose, gradual compression screws, with L/D of 18 are suitable. Three zone heating control and separate temperature control for the nozzle are recommended. Proper drying provides a better solution than vented barrels for ensuring adequate absence of moisture. There are several important variables that may cause a reduction of toughness and strength of the moldings. These include the moisture content in virgin and rework resin, melt residence time (should be less than 10 min), too high or too

PA/PPE Blends (Noryl™ GTX, Dimension™) were introduced in 1985, twenty years after the introduction of PPE/PS blends. Noryl™ GTX blends

Polyester Blends PET’s low crystallization rate, much lower than polyamides, has hampered its use in injection molding applications, but development of fast crystallizing grades (containing nucleating agents) has opened new opportunities for PET. PBT, on the other hand, crystallizes significantly faster, making it possible to mold in short cycles. The crystallizability, combined with low viscosity of PBT melts, has resulted in better acceptance of this resin for molding applications. The thermoplastic polyesters have been often blended with amorphous polymers to improve chemical resistance and processability. Elastomers have also been added to improve their impact resistance.

Polymer Blends Forming

low melt temperature, improper mold venting, effective glass fiber length, contaminated rework or molding equipment, additives such as mold release, lubricants and pigments, part design with sharp corners or non-uniform wall thickness, etc. PBT/styrenics PBT blends with styrenics include those with brominated PS (Starflam™ PBT) as a flame retarded glass reinforced PBT composition for electrical parts, ABS (Novalloy™-B or Cevian™) and ASA (Ultrablend™ S). All are easy to process in standard screw-type machines, with three-zone screws. The maximum molding temperature is 270°C. PBT will decompose at higher temperatures. PBT/ABS blends show excellent moldability, good dimensional accuracy and reduced warpage resulting from substantial reduction of the thermal shrinkage (a typical disadvantage of crystalline resins such as polyesters). The viscosity of PBT/ASA blend rapidly decreases with temperature, therefore it is important to optimize the process parameters to obtain good moldings performance. PEST/PC The most important polyester blends are those with PC. The presence of PC in PET/PC increases the crystallization rate of PET and allows faster cycles. These materials are discussed below, under the heading of PC/PET blends. PCTG/PC or SMA (Ektar™ MB) Blends of PCTG with PC or SMA are typically formed by injection molding into business machines, appliance parts and automotive body panels. Properties are affected by injection speed, screw speed, back pressure, injection pressure, melt temperature, clamping pressure and cycle time. Polycarbonate Blends Numerous blends based on PC have been developed for injection molding applications. Amongst the blends with commodity resins, styrenics have found the widest applications, particularly the blends with ABS.

721

PC/ABS PC/ABS blends combine the good heat resistance and toughness of PC with the excellent processability of ABS. The blends are being offered by several manufacturers: Alphaloy™ from Kanegafuchi Chem., Bayblend™T from Bayer, Cycoloy™ from GE Plastics, Denka™ HS from Denki Kagaku, Dialoy™ C from Mitsubishi Rayon, Exceloy™ from JSR, Iupilon™ from Mitsubishi Gas, Lynex™ B from Asahi Chem., Malecca™ P from Denka, MB2200 from Mitsubishi Gas, Multilon™ from Teijin Chem. Ltd., Novalloy™S (or Cevian™) from Daicel Chem. Ind., Pulse™ from Dow Chem. Co., PX from Teijin Chem. Ltd., Ryulex™ from Dainippon, SC 250 from Idemitsu Petrochem., Techniace™TC from Sumitomo Dow, Toplex™ from Multibase, Inc., Triax™ 2000 from Bayer, ABS/PC 7901 from Diamond Polymers, etc. PC/ABS blends have better flow properties than PC, which means that the parts have fewer knit lines, sink marks as well as less molded-in stress. Molded parts offer excellent surface appearance and good paint adhesion. Up to 20% regrind may be used for all grades. Injection can be done on standard screw-type machines with three zone screws having L/D = 18-20 and compression ratio of 2. Screw back pressure should be kept as low as possible with 350-700 kPa usually sufficient. Shrinkage is low and virtually identical in all directions for nonreinforced grades: 0.5-0.8%, and even lower for the reinforced ones: 0.2-0.4%. Shot size should be between 50 and 70% of the machine capacity to avoid excessive residence time and reduce the risk of degradation. Values below 30% must be avoided. PS, SAN and acrylic are best for purging. The optimum temperature profile depends on many variables such as the ratio of shot size to machine capacity, screw design, mold and part design and cycle time. The upper range melt temperatures should be used while molding thin wall parts. Similarly, mold temperatures in the upper range are advisable for thin wall parts, large parts, and parts that require dimensional stability and a low level of molded in stresses. Mold temperature has minimal effect on gloss level of finished part. Uniform heat removal from the

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mold is critical to obtain consistent dimensional tolerances. It is recommended that no decompression be used. However, when it is necessary to reduce drool, as little decompression as possible should be used. PC/SMA blends (e.g., Arloy™, Ektar™ MB, R4 9900) offer outstanding processability. Machines designed for PS, ABS, PC are suitable. PC/PET The vast majority of blends of PC with engineering resins involve PET or PBT (BCT 4201, from Toray, Dialoy™ P from Mitsubishi Rayon, Ektar™ MB (with PCTG) from Eastman, Makroblend™ from Bayer, MB4300 from Mitsubishi Gas, Novadol™ from Mitsubishi Chem., Pocan™ from Bayer, R2-9000 from Thermofil, Sabre™ from Dow, SC 600 from Idemitsu Petrochemicals, Stapron™ E from DSM, Ultrablend™ KR, from BASF, Valox™ from GE Plastics, Xenoy™ 1000 from GE Plastics). By combining PC with PET or PBT, the best properties of each component can be incorporated. Thus, the blends combine the impact resistance and toughness of PC with the chemical resistance and thermal stability of PBT or PET. However, the morphology of these materials is shear sensitive. Therefore, during injection molding, sharp corners, shut-off nozzles and mixing screws should be avoided. Injection speed of 25 mm/sec is suggested. Screw compression ratio of about 2.2 is preferred, but depending on the screw diameter, it can vary from 2.0 to 2.5 — higher ratios than that may cause shear degradation. The screw length-to-diameter ratio of 20:1 is preferred. A short feed zone (5 flights) and a long compression zone (11 flights) with a gradual constant taper leading to a short metering zone (4 flights) has been suggested. It is important to minimize barrel residence time. HDPE and PS are the best purging agents. Since good heat removal is important, molds should be water or oil cooled. Mold cavity should be vented to prevent blistering or burning.

From an air shot, the recommended melt temperature is 250-285°C, with an optimum of 265°C. Keeping a uniform melt temperature is essential to ensure good color matching of mating component parts. To avoid thermal degradation, the cycle time must be optimized and shot size should be between 40-75% of machine capacity. This is particularly important when producing large parts that tend to have longer cycle times. Injection speed should vary depending on the size and complexity of the part, as well as the gate and runner system. Generally, slow to moderate speeds give the best results. Too fast speeds result in excessive shear heating, degradation and splay. For small parts or parts with relatively simple geometry, a constant filling speed is recommended. For larger or more complex parts, it is often helpful to profile the speed. Generally slow screw speeds are recommended to provide a homogeneous melt without an excessive amount of shear heating. Screw speed should be set so that plastication is complete 2 to 3 seconds before the cooling of the preceding shot is over and the mold begins to open. This will allow for a slow screw speed without causing an increase in the total cycle time. For impact-modified blends (Xenoy™ 1200 and 5720, Stapron™ E), screw speed should be set at minimum, as low as possible (less than 50 rpm), just sufficient to assure a reasonable cycle time. Melt temperature should be kept as low as possible, without binding the screw. When temperature control is critical, as is the case here, the actual melt temperature should be checked with a pyrometer or digital thermometer, rather than relying on barrel set points. The material runs best in thick parts with oversized gates. Injection speed of 25 mm/sec is suggested. When running a large shot (>70% of machine capacity), a flat or reverse barrel heat profile may help to convey the melt. PC/TPU PC is also blended with a polyester urethane TPU (Texin™ 3000: PC/TPU, Texin™ 4000: TPU/PC). For these systems, screw machines are preferable with L/D = 15 and compression ratio = 2-3.

Polymer Blends Forming

Polyphenylene Ether Blends Polyphenylene ether resins were developed in the early 60’s but could not be processed because the thermal degradation temperature was a few degrees below the melt temperature required for processing. Later discovery that PS can act as a plasticizer lead to the introduction of PPE/PS, one of the most successful blend (Noryl™, Prevex™, Luranyl™). PPE/PS All types of injection molding machines have been successfully used for molding PPE/PS blends. Inline reciprocating screw machines are preferred because of their rapid heating and reduced pressure loss through the cylinder. Depending on screw diameter, L/D = 18-20, compression ratio = 2 to 2.5, and screw pitch of 0.8-1×D, are suggested. The entrance zone can make about 50% of the compression zone and about 25% of the total screw length. Short screws, short compression screws, and screws with progressively changing channel depth are inappropriate. Injection pressure depends on several variables such as melt and mold temperature, part geometry, wall thickness and flow length. Injection temperature affects ductility, thus it must be optimized for each product. Back pressure is important to ensure a homogeneous melt and consistent shot size. Higher back pressure improves melt mixing as a result of higher melt temperatures. Surface properties, degree of crystallinity, resistance to warpage and shrinkage depend on mold temperature, usually set at T = 60-100°C. HDPE, PS and ground acrylic are best for purging. For the injection molding of foamable blends, either inert gas is introduced directly into the melt or the resin is preblended with a chemical blowing agent. The gas present or formed expands when the material is injected. Polyoxymethylene (acetal) Blends POM/Elastomer Elastomer modified POM compositions (Celcon™ from Celanese, Delrin™, from DuPont, Duraloy™

723

from Hoechst-Celanese, Formaldafil™ from Wilson-Fiberfil, Thermocomp™ and Fulton™ KL from LNP, Hostaform™S from Hoechst, RTP 800 from RTP, TC from Polyplastics, Ultraform™ from BASF) are molded in a variety of designs. The plasticating screw design should take into account the high heat requirement of POM caused by its crystalline nature. Screw designed for optimum processing has shallow flight depth in the metering section and a higher compression ratio than general purpose screw. Its length-todiameter ratio, L/D = 20, should be divided as follows: feed 45-50%, transition 20-30%, and metering 25-30%. In such design, part of the energy is provided from mechanical work — thus there is a need for an extra powerful motor. The heater band should extend as close as possible to the nozzle tip. The non-return valve should meet the following requirements: no holdup spots, no flow restrictions, good seal and control of wear. The injection molding machine should be capable to deliver a shot size of 50% or less of the injection stroke. Typically, the machine should be able to inject its rated capacity in less than 2 seconds, except for high viscosity resins or when warping or jetting must be prevented by slower injection rate. Depending on the grade, the cylinder temperature setting should be between 190 and 220°C. The clamping force depends on the injection pressure and speed, dimension of parts and runners, precision and tolerances. Typical clamping force varies from 48 to 69 MPa. Thin parts with restricted flow may require pressures up to 100 MPa. POM/TPU For POM/TPU blends (Ultraform™, Hostaform™ S), it is important that cooling water is well sealed off from the mold cavity. Any leakage may cause formation of corrosive solutions. Like other semi-crystalline polymers, the mechanical properties of injection molded materials are governed by the degree of crystallinity. Raising mold temperature increases hardness, rigidity and strength. Typical mold temperature range is 60-90°C, but for precision parts, T = 90-120°C may be required. PE is rec-

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M. M. Dumoulin

ommended for purging. Screw rate should be decreased with increasing screw diameter.

10.5.3.3 Injection Molding of Specialty Resin Blends

POM/PTFE (Hostaform™ C 9021 TF) Screw rate should be decreased with increasing screw diameter. Back pressure should be kept as low as possible while allowing an homogeneous melt, in order to reduce shear. Higher mold temperature will result in higher strength, smoother surfaces, lower shrinkage and longer cycle time. PE is recommended for purging. When changing from one grade to the other, only temperature adjustments should be made.

Blends based on specialty resins used routinely in injection molding applications include PEI/PC (Ultem™ LTX), PPS/PPE (DIC™ PPS), PEEK/ LCP (Sumiploy™ EK) or PEEK/PES (Sumiploy™ SK). Injection molding conditions for these systems are listed in Table 10.33. PPS/PPE PPS/PPE blends (DIC™ PPS) are immiscible and prepared by reactive processing. Commercial

Table 10.33a. Injection molding conditions for blends of specialty and elastomeric resins — Part A Blend

Trade name

Temperature (°C) Melt

PEEK/PES PEEK/LCP PEI/PC PPS/PPE PSF/ABS PSF/PET Nitrile/PVC, cPE, TPU, PA NBR/EPDM

Sumiploy Sumiploy Ultem DIC PPS Mindel A Mindel B

340-380 280-330 280-310 270-290

Chemigum TPE Geolast

135-185 215-220

Barrel

Nozzle

350-390 350-380 310-400

338-412

Shrinkage (%)

Shot size (% MC1)

0.3-0.5 0.24-1.1 0.66 0.3-0.9

20-80

Mold 160-200 160-200 110-150 120-180 70-120 66-104 10-40 50-75

1. MC: Machine capacity

Table 10.33b. Injection molding conditions for blends of specialty and elastomeric resins — Part B Blend

Trade name

Injection rate

Screw CR1

PEI/PC PSF/ABS PSF/PET NBR/EPDM

Ultem Mindel A Mindel B Geolast

CR1: Compression ratio

slow-mod max

Pressure (MPa)

Speed (rpm)

Injection

Back 0.35-0.7 0.6-2.4

2 100-140 100-200

0.7-1.4

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Polymer Blends Forming

grades include those with 30 or 40 wt% of glass fibers. During injection molding, the blends show reduced flash in comparison to PPE. Minimum recommended wall thickness is 0.3 mm. PEEK/LCP, PEEK/PES Two series of blends based on PEEK are available from Sumitomo, PEEK/LCP (Sumiploy™EK) and PEEK/PES (Sumiploy™ SK). The blends are to be processed by injection molding. The critical factors affecting performance of both series of blends are crystallinity of the matrix resin (PEEK) and blend morphology. It is important to control the diameter-to-length ratio of LCP fibers or the particle size of PES dispersed in PEEK matrix to achieve the expected performance. Sumiploy™ K series have excellent processability compared with PEEK and glass fiber filled PEEK. Since PEEK, the matrix phase, has a high melting point, it is necessary to crystallize completely with the mold temperature as high as possible taking into consideration the deformation of molded parts and molding cycle. Standard mold temperature is 160-200°C. If it is lower than 160°C, the degree of crystallinity of PEEK is reduced. It may be subsequently increased by annealing the moldings for more than 1 hour at 200-240°C. The morphology of Sumiploy™ depends on MW of each polymer, additives and processing conditions. In PEEK/LCP blends (Sumiploy™ EK 1400), the oriented LCP macromolecules act as reinforcing fibers. In low molecular weight PEEK the LCP fibers are large, diameter d = 1-2 µm and length l = 30-50 µm, and well oriented, thus providing the necessary reinforcement. By contrast, in high molecular weight PEEK, the LCP fibers are smaller and disordered, thus the reinforcing effect is negligible. PEEK based blends also include grades with PES (Sumiploy™ SK1660 and SK1800) where SK1800 comprises PES dispersed in PEEK, while SK1660 is a glass fiber reinforced grade of SK1800. The advantages of Sumiploy™ SK compared to PEEK or glass fiber filled grades are primarily the low temperature strength and good processability.

In SK1800 and SK1660, PEEK is the matrix and PES is dispersed as nearly spherical particles with diameter, d = 3-5 µm. This has been accomplished by using relatively low MW PEEK and relatively high MW PES. On the other hand, by combination of high MW of PEEK and low MW of PES, blends in which PEEK and PES form co-continuous phases can be obtained. Therefore some of the excellent characteristics of PEEK such as fatigue and chemical resistance are retained. Presence of fillers in SK1660 requires increase of cylinder temperature to 360-400°C. PSF/ABS (Mindel™ A) PSF/ABS blends are more sensitive to shear than modified PPE or PC. When the shear rate increases, melt viscosity rapidly decreases, which provides an opportunity to manufacture parts having thin cross-section or which require long flow. Melt temperature should always be kept below 320°C. Upon completing molding, the material should be immediately purged and immersed in water. If this recommendation is not followed, spontaneous decomposition of hot mass may result producing toxic fumes containing HCN and acrylonitrile. 10.5.3.4 Injection Molding of elastomeric Blends NBR blends For injection molding of nitrile rubber/PVC or chlorinated PE or PU or PA or their combination (Chemigum™ TPE), processing temperature depends on the grade: the harder the material, the higher the temperature. For Chemigum™ TPE, recommended melt temperature is 135-185°C and mold temperature: 10-40°C. For NBR/EPDM blends (Geolast™), recommended melt T is 215-220°C, mold T: 50-75°C, screw speed: 100-200 rpm and back pressure: 0.7-1.4 MPa. 10.5.4

Summary

Blending affects significantly a number of properties that, in turn, impact on the injection molding process and resulting parts. For example, the

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presence of PC in PET/PC increases the crystallization rate of PET and allows faster cycles. The rheological behavior and, consequently, processing requirements can also be modified. For instance, addition of EPDM to PP yields materials with high and shear dependent viscosity. Because of the viscosity, large nozzles, sprues, runners and gates are required to injection mold these resins, as well as high injection and packing pressures. On the other hand, flow properties of PA/PP blends are similar to those of PP, in particular less dependent on temperature, making processing significantly easier. Injected parts usually show a skin/core morphology. The skin corresponds to the early frozen material where the macromolecular chains are oriented whereas no or little orientation is observed in the core region. In the case of semicrystalline polymers, fibrillar, non-spherulitic crystallinity is usually produced in the skin while a spherulitic crystalline morphology is normally observed in the core. The phase morphology of blends also depends on position. Because of higher strain and faster cooling in the skin, the morphology is often oriented in the flow direction. Dispersed particles in the core of the moldings are more likely to appear in the form of spherical domains. The relative thickness of these two regions and the extent of deformation of the dispersed phase are related to the molding conditions, particularly injection speed and temperatures. The thickness of the skin layer is normally found inversely proportional to melt temperature and injection speed. The problem of weld lines in injection molding can be more severe for blends than for single-phase, neat polymers. It is recognized that strength at the weld line can be as low as 10% of that away from it. Thus weld lines may significantly reduce part performance. The weakness of knit lines is attributed to poor mixing of the meeting, colder flow fronts. In the case of blends, there is also a reorientation of the dispersed phase in a direction parallel to the surfaces, creating a discontinuity in the morphology of the molded part.

As in the blend extrusion, the key to obtaining the expected properties in injection molding is control of blend morphology. One clear example is shown by PEEK/LCP blends in which the oriented LCP macromolecules act as reinforcing fibers. The choice of proper molecular weight and processing conditions allows to produce reinforcing LCP fibers with adequate shape (L/d = 25-30) and orientation. By contrast, the use of high MW PEEK causes the LCP fibers to be smaller and disordered and, thus, non-reinforcing. Similarly, in PEEK/PES blends, the right morphology (small dispersed spherical PES particles) is achieved by the proper choice of MW. A different combination of MW will yield co-continuous phases. In all cases, the morphology that was generated during the compounding stage can be dramatically affected by the molding conditions, particularly temperature and shear stress. Long dwell time can induce coalescence of the phases. Local shear stresses during the injection molding may be much higher than those generated during compounding. During processing, the material also encounters extensional strains, e.g., flowing through convergences such as a gate, or expanding at the moving front while flowing to fill the mold cavity. Dispersed elastomer particles in impact modified materials can be affected by the flow: it was observed that elastomer particles in HIPS had shrunk and the elastomer phase around dispersed PS islands had buckled, resulting in a significant decrease in impact strength. The injection molding process generates some of the harshest conditions encountered in polymer processing: high shear stress and strain, local temperature gradients, convergences and divergences, weld lines, etc. Generating the right morphology requires careful design of both material and process parameters, in addition to strict control.

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10.6

Blow Molding

10.6.1

General Description of Blow Molding

Blow molding is a process for manufacturing hollow articles such as bottles and containers. It involves first forming a parison or preform, which is a pre-shaped sleeve made by extrusion or injection. The parison is engaged between two mold halves and air is blown into it, causing the parison to take the shape of the mold. The plastic solidifies when it comes in contact with the cold mold. After sufficient cooling, the mold opens and the part is ejected. The surface of the parts can be smooth and bright or engraved, depending on the surface of the mold itself. Blow molding competes against other methods such as injection molding and rotomolding both from a quality and cost points of view. New developments, particularly in coextrusion blow molding and stretch blow molding, extend the range of properties achievable and consequently the range of products. Blow molding benefits from good economics, as well as processing advantages: low forming stress, varying wall thickness and the possibility to mold irregular and re-entrant shapes. Blow molding of plastics is the third most widely used polymer forming process. As described in Chapter 1, Introduction to Polymer Blends, the technology was introduced to mold cellulose acetate already in 1880. From 1930 to the mid40’s, the method was used mostly for cellulosics and styrenics, until LDPE containers were introduced and became the focus of large mass production efforts. Although the primary goods produced with blow molding are bottles of all shapes and sizes, modern blow molding plants currently manufacture a great variety of products, such as automotive fuel tanks, air ducts and manifolds, toys, sports equipment and many others. There are four different blow molding processes: extrusion blow molding, injection blow molding, stretch blow molding and the less frequent dip blow molding. Extrusion blow molding can be semi-continuous, with a continuously rotating screw, extrud-

ing parisons into one or more dies, or intermittent, either using reciprocating screw or ramaccumulator extrusion systems. Intermittent systems give more flexibility and control over the most important part of the process, the parison shape. The blowing step is fast and does not allow any control over the thickness of the finished part. Thickness control is accomplished by parison programming by varying the opening of the die. The shaping of the parison obviously also depends on rheological properties of the polymer. Good control over the parison shape can be obtained in injection blow molding, where the preform is injection molded first, then moved to the blowing mold for the final forming. Some orientation is also achieved during the injection process. Furthermore, additional orientation may be obtained by first mechanically stretching the preform (in the axial direction), then by blowing it, which induces hoop orientation. Since orientation improves the mechanical, optical, and barrier properties of the material, it makes it possible to reduce the weight of the part without paying any performance penalty. The process has been well established over the years for bottle production. The recent developments in both blow molding technology and materials are leading to new advanced applications of the process. Vacuumassisted blow molding ensures closer contact between the material and the mold and thus improves heat transfer. This reduces the molding cycle time by up to 10%. In the present section, the discussion will focus on the conventional, and most widely used, extrusion and injection blow molding processes. 10.6.2

Processing Parameters

Table 10.34 shows typical values for some processing parameters for some commodity and engineering resins. 10.6.2.1 Pressure Pressures needed for blow molding are different than those used in other competing processes used for production of similar parts. They are much

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Table 10.34. Blow molding of conditions for commodity and engineering resins Material

Air pressure (MPa)

HDPE LDPE PVC PP PS ABS POM PPE PA PC PMMA PBT

Swell (%)

0.4-0.7 0.15-0.4 0.5-0.7 0.5-0.7 0.3-0.7 0.35-1 0.7-1

15-65 30-65 30-35

0.5-1 0.35-0.55

5-10

higher than used in rotomolding but significantly lower than those required for injection molding. While pressures in the range 15 to 200 MPa are normal for injection molding, they vary in blow molding from 0.15 to 2 MPa. The lower pressures mean lower equipment costs, lower residual stresses in the parts and more even stress distribution, resulting in less warpage. 10.6.2.2 Wall Thickness The ability to blow mold containers with widely different wall thickness represents the most important advantage of the process. Since the wall is not constrained by two mold halves, it can be varied by simply changing processing conditions, allowing to change part weight at will. Much thinner walls than in injection molding can be produced. Obviously, the wall thickness depends on parison thickness, but parison programming makes it possible to control the variations in parison wall thickness: a hydraulic positioner in the die modulates the internal die diameter.

10-20

Temperature (°C) Material

Mold

160-220 130-180 180-205 200-220 280-300 180-230 150-280 240-250 240-270

15-30 15-30 15-30 30-60 40-65

50-70 40-60 245-260

10.6.2.3 Material Properties For successful blow molding (especially extrusion blow molding), several rheological functions of the molten resin are of critical importance: melt strength, extrudate swell, elongational viscosity, etc. Good melt strength is essential to avoid sagging of the parison. For this reason, extrusion blow molding requires resins with higher molecular weight than in injection molding. The low MW condensation polymers (PA’s, PEST’s) are processed by injection blow molding. The use of high MW resins, in turn, makes it possible to design parts for applications that require properties of higher MW materials. This is particularly beneficial for environmental stress cracking resistance, which benefits from higher MW, particularly when molding containers designed for fluids that promote environmental stress cracking. Furthermore, the melt may need to be used in large blow-up ratio applications, or in those with sharp corners and edges. It is important for the finished product to have wall thickness devoid of thin spots. The extensional viscosity must allow these extensions.

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The blowing phase of the process requires a large gap between the processing and the solidification temperature. In extrusion blow molding, swell of the material leaving the die must be taken into account. This may be complex since not only swell of the outside diameter and that of the wall thickness, but also sagging due to gravity affect geometry of the parison. Typical values of the extrudate swell for selected commodity and engineering resins are listed in Table 10.34. Shrinkage of the materials also has to be considered. Shrinkage of the material during the cooling stage in blow molding, as in injection molding, depends mainly on crystallizability of the material, molding temperature and pressure. Owing to the anisotropic distribution of the residual stresses, shrinkage in the axial direction is usually greater than that in the transverse (hoop) direction. These effects may lead to variation in wall thickness rather than parts dimensions. 10.6.2.4 Mold The mold quality as well closing speed and force strongly affect product quality. For example, increasing the mold closing speed reduces the gravitational sag effects, thus increases the wall thickness. To avoid excessive flash, mold closing pressure P = 7-14 MPa is normally used. In addition, mold surface quality will dictate that of the finished part.

10.6.3

Blow Molding of Polymer Blends

Information on conditions to be used for blow molding of commercial polymer blends is limited. The available data are listed in Table 10.35. 10.6.3.1 Blow Molding of Blends of Commodity Resins Polyolefins are the material of choice for blow molded parts, either containers, bottles or tanks. However they show high permeability to many organic solvents and gases. This problem can be overcome by a variety of means: co-extrusion of a thin layer of a barrier polymer, surface treatment or coatings. This has prompted the investigation of polymer blends in which one of the polymers has good barrier properties. The concentration of the immiscible second polymer has to be kept as low as possible to keep the good toughness of polyolefins and minimize the costs. To minimize both concentration of the barrier polymer and permeability of the blend, the barrier resin must be distributed as large thin oriented sheets parallel to the container wall or film surface. When this type of morphology was imposed on PE/PA-6 blends, permeability similar to that obtained by co-extrusion was observed [Subramanian, 1985, 1987]. Figure 10.6 shows hydrocarbon permeability for HDPE containers [Subramanian, 1990]. The lamellar morphology shows significant

Table 10.35. Blow molding conditions for blends Blend

Trade name

Temperature (°C) Melt

SAN/elastomer PPE/PS PA/PPE PC/PET or PC/PBT PA-6/PPE NBR/EPDM

Luran S Noryl Noryl GTX Xenoy Dimension Geolast

220-230 230-255 278-288 220-230 265-295 205-220

Shrinkage (%)

Accumulator

Mold

230-255 278-288 210-220

65-95 79 50-80 60

0.8-1 1.3-1.7

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M. M. Dumoulin

improvement in permeability over the pure HDPE or the blend with droplet dispersion. The platelet-type dispersion is obtained under specific processing conditions, in particular under low shear conditions. Intensive melt compounding and dispersive mixing destroy the platelets and reduce the effectiveness of the PA. Blending PA in PE, for example, can significantly reduce oxygen permeability albeit at the cost of increased water permeability [Kamal et al., 1984]. In this particular case, the generation of lamellar morphologies was indicated by the permeability data, which tended towards the series law only above PA loadings of 20 to 25 wt%. Champagne et al. [1995] used extrusion blow molding to generate PA lamellae in HDPE. Intensive dispersive mixing, such as that induced during twin screw mixing, hampered the phenomenon by creating too small PA droplets which could not be deformed into lamellae during blow molding. Around 50% of PA phase was in the form of lamellae with an aspect ratio > 10 under the best processing conditions. Figure 10.7 shows an example of the morphology of lamellar blends. A limited number of blends have been developed for the blow molding process. Two polyolefin blends are commercialized for this purpose, those of the PP/EPDM and PA/PO type. PP/EPDM The harder grades of PP/EPDM blends (Santoprene™) can be blow molded. The resins are characterized by a co-continuous morphology and high viscosity. The recommended parameters for blow molding are similar to profile extrusion. Moisture content is critical and should be kept below 0.1%. PA/HDPE Lamellar dispersion of polyamide domains within HDPE matrix has been achieved by proper choice of rheology of the resins and processing conditions [Subramanian, 1985]. The platelet shape of PA domains enhanced its barrier properties against permeation by gases or liquids. The technology was commercialized and marketed by DuPont (Selar™ RB), aiming at packaging applications

such as blow molded containers. The blends usually include PA or poly(ethylene-co-vinylalcohol). DuPont’s technology usually involves addition to a polyolefin of a custom tailored mixture of a polyamide with ionomeric compatibilizer or polyamide-maleated polyethylene graft copolymer, the latter copolymer reacting with the PA. Formation of the copolymer reduces the interfacial tension and stabilizes the morphology. The flow imposed morphology of these immiscible blends may be visualized as overlapping PAlamellae dispersed in PO matrix. Biaxial stretching of the polyamide drops dispersed in the polyolefin creates the overlapping lamellae. This structure creates surprisingly high barrier properties. The grafted PA/PE blend is offered as a concentrate to be blended into HDPE at the customer’s location. It should be noted that in the case of the Selar™ technology, the processing know-how and not the product is marketed. The actual mixture’s formulation and concentration depends on the customer’s resin, processing equipment and product requirements. The processing conditions have to be optimized for each specific application and strictly adhered to. Deviations from the optimum settings resulted in the loss of the desired morphology with ensuing deterioration of barrier properties. Co-extrusion blow molding represents a commercially viable technology that has hampered the wide commercial acceptance of the Selar™ technology. However, increasing needs to recycle production scrap and post-use waste might improve its market penetration. Styrenics blends used for blow molding include SAN/elastomer (Luran™ S) and ABS/BA or PB (Terluran™). The high viscosity grades are the most suitable. Parisons should be extruded at 220-230°C. In some cases, the temperature might be kept as low as 200-210°C to avoid sagging.

Polymer Blends Forming

10.6.3.2 Blow Molding of Blends of Engineering Resins Blends of engineering resins used in blow molding applications include systems based on PA, PC and PPE. When using impact modified PA-6, PA-66 or PA-6,12 (Zytel™), the moisture content must be significantly lower (0.05%) than that required for injection molding (0.25%). PA/PP Bottles and containers can be blow molded using PA/PP blends (Orgalloy™ R). The series of blends contain either PA-6 (Orgalloy™ R-6000) or PA-66 (Orgalloy™ R6600). Some grades are impactmodified through the addition of an elastomer. The morphological structure of Orgalloy™ R is different from that observed in PA/PP mechanical blends. The flow properties of these blends are similar to those of PP (i.e., improvement over those of PA). Due to its high thermal stability, the blend can be processed over a wide range of temperature, T = 220-290°C. Orgalloy™ blends are easier to process than PA since their melt viscosity is less sensitive to temperature. For extruding the parison, the extruder barrel should be at least 20D long, fitted with standard three-zone screw, with compression ratio of 2.5:1- 3:1. Excessive shearing should be prevented. Purging with PP followed by LDPE is recommended. The blow molding head may be either crosswise or in-line type. It was reported that the use of nitrogen instead of air injection resulted in improved weld strength. PA/PPE Blow molding grades of PA/PPE blends can be based either on PA-6 (Dimension™) or PA-66 (Noryl™ GTX). For PA-6/PPE blends, a melt temperature of 265-295°C (as for extrusion) is recommended. Mold temperature of 60°C and typical blow ratio of 2 to 3 are normally used. The PA-66 blends are compatibilized by addition of a PA-PPE copolymer. In the blends, PA constitutes the matrix phase in which particles of PPE are dispersed, playing a role of low density organic filler. A melt temperature of 278-288°C is recommended, with adapter and accumulator

731

temperature set at T = 278-288°C and mold temperature at T = 79°C. PS and ground acrylic are best for purging. PC/PET or PBT For PC/PET or PBT blends (Xenoy™), on startup, the barrel temperature can be set around 245-260°C, then it can be lowered as the production commences. Moderate to slow screw speed should be used, otherwise material will overheat. A reverse temperature profile may increase throughput. PPE/HIPS Impact-modified PPE/HIPS blends parts have been blow molded for automotive applications since the mid 1980’s. In the last few years, this processing method has gained wide acceptance, as it was expanded to produce parts for the office equipment industry using flame-retarded PPE/HIPS blends and PPE/PA blends. The low melt index grades of PPE blends with PS or HIPS (Noryl™, Luranyl™) are most suitable for blow molding. The optimal processing conditions depend on the article type and machine and thus should be experimentally optimized. The typical temperature range for parison extrusion is 250-260°C whereas it is 230-255°C for the adapter and accumulator. HDPE should be used for purging. PPE blends are essentially processed by extrusion-blow molding using standard blow molding arrangements. Normally, the extruder has a grooved cylinder entrance zone. This zone must be heated to 80 to 100°C for minimizing torque and achieving regular material flow. Separate control of the temperature of the cylinder, adapter and nozzle are needed. The clamping force must be high enough to build a back pressure of 0.6 to 1.0 x10-6 Pa above the blowing pressure. Melt temperature depends on the material type and complexity of the parts and varies between 220 and 270°C. To obtain a sufficient stability of the melt during blowing, melt temperature during the blowing stage is about 50 to 60°C lower than that used during injection molding, while the mold temperature is

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the same. The PPE/HIPS and PPE/PA blends show minimal extrudate swelling effects. Blow forming normally results in mat surfaces, especially with application of texture. These textures are easy to reproduce and don’t leave surface defects. Aluminum and zinc are the usual mold materials while epoxy can be used for short production series. 10.6.3.3 Blow Molding of Blends of Specialty Resins Information available on commercial grades of specialty blends is even scarcer than that for other blends. The recommended temperature range for NBR/EPDM blends (Geolast) is given in Table 10.35. Excessive melt temperature reduces the blow ratio and should be avoided. 10.6.4

Summary

Information on blow molding of commercial blends is limited. A number of commercial blends show both the rheological behavior and final physical properties that make them suitable for the process. Moreover, some blend systems have been developed specifically to take advantage of the good barrier properties of one specific component. An example of such systems is duPont’s technology for polyolefin blends with high barrier properties. The technology involves addition to a polyolefin of a custom tailored mixture of a polyamide with compatibilizer. Blends offered commercially as Selar™ RB involve maleation of polyethylene and subsequent grafting with polyamide. The flow imposes a morphology of overlapping PA-lamellae dispersed in PO matrix, resulting in superior barrier properties. Processing conditions in this case are crucial and are optimized for each application. The beneficial performance depends on the generation of specific lamellar morphology. Processing conditions have a great impact on the morphology of blends, resulting in potentially significant advantages for blow molded containers.

10.7

Other Forming Processes

10.7.1

Calendering

10.7.1.1 General Description of Calendering Calendering refers to a continuous process in which the forming takes place through a gap between cylindrical rolls with parallel axes. The technology is used to produce films, sheets or to coat a prefabricated sheet material. The calendering rolls are rotating and the forces needed to push the material through are generated by the surface of the rolls. Several polished rolls are typically used and the resulting films are usually 25 to 250 µm thick, whereas the sheets are 1 to 4 mm thick. Stretching often takes place between the different sets of rolls. The most common roll configuration is an inverted-L shape. The polymer mass is fed in the nip of the first pair of rolls, emerges as a sheet which passes through the remaining rolls. The first nip controls the feed rate and the others the total stretching and the final thickness. Typically, films or sheets of up to 3 m in width can be produced at rates of at least 120 m/min. Typical dimensions of the rolls are: 0.15 to 1.0 m in diameter and 0.3 to 3 m in length. Pressures of 7 to 70 MPa can be found between the rolls of large calenders. Compared to other forming methods, calendering has the advantages of high productivity, relatively low investment costs and high quality of product. The advantages of this process compared to film blowing are lower space requirements, no blocking, ability to make almost transparent films from semi-crystalline polymers, better thickness control and the possibility to apply treatments to both sides of the film. More complete descriptions of the process and its applications have been published [Kohlert et al., 1989]. 10.7.1.2 Processing Parameters Shear rates during calendering range typically from 10 to 103 sec-1. The pressure and forces encountered within the nip region are affected by the characteristics of the material (rheology,

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homogeneity), processing temperature, height of roll gap and speed of rolls. The description of flow in a calender assumes that all the material is eventually drawn inside the gap. As the material moves towards the nip region, it is sheared and elongated, the cross-sectional area of the gap between the rolls decreases and the pressure gradually increases to reach a maximum somewhere shortly before the minimum gap distance. Up to this maximum, the pressure is opposed to the direction of flow whereas after the pressure peak, the pressure gradient helps the flow rate to increase. 10.7.1.3 Calendering of Polymer Blends TPO’s A number of polyolefin blends are routinely calendered. TPO’s consisting of blends of PP or PE with α-olefin random copolymer (Oleflex™) are suitable. The copolymer is finely dispersed as a rubbery phase in the polyolefin matrix. The blend is processed at conditions similar to polyolefins. For calendering, typical melt temperature is 165-175°C. PP/EPDM PP/EPDM blends (Santoprene™), with their high melt viscosity and melt strength, are easily calendered. The material is normally premixed in a Banbury or other intensive mixer. The melt temperature should be about 190°C. Roll temperature should be about 70-120°C. Conventional fourroll calenders can produce quality sheet with thickness 0.4 to 1 mm. PVC/acrylics PVC, in either rigid or flexible formulation, is the most widely used thermoplastic for the calendering process. A number of other polymers can be calendered but frequently these can be more easily extruded with high throughputs. Acrylics are often added (1 to 5 wt%) to PVC as plasticizing processing aids, allowing calendering speed and film thickness to be increased. External lubricants may also have to be added to improve release of the film from the rolls, thus speed up the

processing line. Blends of PVC with 11% acrylate elastomer (Vinidur™) are used for the production of calendered film having high impact strength and suitable for outdoor exposure. 10.7.2

Thermoforming

10.7.2.1 General Description of Thermoforming Thermoforming consists of forming a preheated plastic sheet or film into a desired shape against a mold surface. Forming can take place pneumatically or mechanically. In the first case, a differential air pressure is used, either vacuum on the mold side, compressed air on the other, or both. In the second case, the preheated sheet of material is pressed into the mold using either a preform, a plug, or by a combination of both. Virtually all thermoplastics can be thermoformed. However, because of limitations in the material characteristics, some can only be used for simple shapes with small draw ratios. Historically, thermoforming was introduced in the 1930’s for cellulosics. During the following decades, as the new thermoplastics were commercialized the method was gradually extended. Commercial vacuum forming units were introduced in the 1950’s, mainly for the fast growing packaging industry. Relatively small parts such as drinking cups, throw-away food containers, bowls and trays have constituted the bulk of the thermoforming production. The process began to be used in the 1970’s to make parts for other industries, such as the automotive, recreational and furniture. For example, thermoformed ABS hulls are being produced for 4-5 m long boats. Some office machine housings are also being thermoformed. Among the advantage of the method is the ability to produce thin walls, often unattainable by other methods, and the low tooling cost, particularly advantageous for short runs. The complexity of the process varies widely. Forming can be accomplished in a simple single-stage operation, where a precut sheet is inserted into a clamping frame, heated and then formed. The method can also take place in multistage comput-

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er-controlled operations in which the preforms are mounted on a rotary table and the different phases take place simultaneously. The output can be increased in continuous operations where extruded sheets are fed directly into the forming station. Moving tracks clamp the moving sheet during forming with multi-cavity molds. 10.7.2.2 Processing Parameters Thermoforming is one of the simplest forming processes. Nevertheless, successful forming depends on a number of parameters, related to characteristics of the material and shape of the product: material’s rheology and its temperature dependence, minimum and maximum thickness, temperature gradients, extent of uniaxial and biaxial orientation, etc. The melt viscosity and strength will dictate the rate and depth of draw achievable, which in turn affect orientation. High orientation can be achieved using semi-crystalline materials and controlling the temperature during the stretch. This semi-solid state formation makes it possible to produce rigid, transparent containers with highly reduced wall thickness over the standard thermoformed product. In all cases, temperature control is critical. The processing range is generally narrow: the temperature must be high enough to ensure that the material is molten and easily deformable. However it should be low enough to ensure that it will not sag. Using infrared heating can be advantageous since it allows shorter heating times. 10.7.2.3 Thermoforming of Polymer Blends Styrenics HIPS Temperature of HIPS (Polystyrol™) or PS/PB blends (Styroplus™) during thermoforming is usually within the range 130-150°C. High draw ratios, e.g., 5:1, are common. It should be noted that PS, being an amorphous polymer with low heat capacity, requires less heat to reach the processing temperature than other commonly used plastics, thus it is a “thermally efficient” resin.

SAN/elastomer Thermoforming of SAN/elastomer blends (Luran™ S) is usually carried out at 140-170°C. The presence of moisture may cause blister formation. Sheets should be stored at lower than 30% relative humidity at 23°C. For example, a sheet stored at 50%, may absorb 0.4% moisture whereas it has been reported that as little as 0.05-0.1% of moisture may cause problems. Sheets are typically stored in PE film. ABS/BA or PB For ABS/BA or PB blends (Terluran™), the stock temperature is in the range 135-170°C. Panels of ABS blends should also be protected against moisture pick up by sealing them in PE film. Moisture concentrations above 0.05% may cause blistering. PVC Blends PVC/ABS Sheets of PVC/ABS blends are often used for thermoforming large parts such as automotive or aircraft interiors. These blends have also been formed by molding and extrusion processes as low cost alternatives to flame-retarded ABS. The low thermal stability of PVC requires strict control of the processing conditions. PVC/PMMA (Kydex™) Heating PVC/PMMA blends at 205°C for a prolonged period of time results in degradation, during which HCl is evolved. If thermoforming is done in the specified temperature range, less than 1 ppm of HCl is produced. If temperature is too low, the product will have internal stress due to difference in temperature throughout its cross-section. Heating time will depend on sheet thickness and varies between 2 and 18 min for sheets ranging from 0.7 to 6 mm. Shorter times are obtained with infrared heating. Optimum forming temperatures are 175-200°C. The cooling rate does not have any significant effect on product property as long as cooling is uniform. Non-uniform cooling, on the other hand, will cause residual stress in the material. In order to

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obtain repeatability, material should be left in the mold until the temperature falls below 40-50°C. All presses can be used but the best results are obtained when using automatic thermoforming or vacuum forming machines.

ture between 180-210°C. The forming mold is normally at 80-100°C. 10.7.3

Film Blowing

10.7.3.1 General Description of Film Blowing Acrylics Most types of equipment are satisfactory for thermoforming acrylic-based compounds (Acrylite™, Cyrolite™ and XT). Typically, the temperature is in the range 160-175°C, but 205°C might be needed for high speed forming. High heating temperature may affect gloss and clarity. Stock storage is important to avoid moisture pickup. It is preferable to keep orientation below 25% since a high orientation causes product distortion. Mold design should assume shrinkage of 0.4-0.8%. Mold temperature should be kept at 70°C, i.e., below the heat distortion temperature. Engineering Blends PA/ABS Owing to the presence of PA, the PA/ABS blends (Triax™ 1000) have narrow processing window. Forming temperature range from 235 to 250°C. Prior to thermoforming, the stored sheets must be dried at 85°C for 8 hours. PC/SMA Several PC blends are routinely thermoformed. For PC/SMA blends (Arloy™), the stock temperature should be between 135 and 150°C. PC/ABS (Bayblend™ T, Triax™ 2000) PC/ABS blends require more heat to form than ABS. This can be accomplished by means of longer cycles and/or increased stock temperature - the normal forming temperature for the blend is 205-230°C, i.e., 10°C higher than that for ABS. Stored sheets need to be dried at 82-93°C for up to 8 hours. PC/ABS/SAN (Bayblend™) PC/ABS/SAN blends are thermoformed using standard methods at sheet tempera-

The vast majority of plastic films, particularly thin films, with thickness less than 100 µm, are manufactured by the film blowing process. The polymer melt is extruded through an annular die to create a tube, which is then inflated by blowing air from inside. The blowing production lines can be horizontal or vertical (up or down). Besides obvious different building requirements, the line direction influences the process itself. The horizontal lines might show problems engendered by the gravity and uneven cooling. Vertical, upward lines, are the most widely used. The bubble created in the vertical direction is drawn by rollers at the top of the blowing tower. Cooling is generally achieved through forced air convection, by blowing air through an air ring above the die. The air blown inside the tube, via an air duct inside the die, is trapped inside since the pair of pull rolls at the top seals the bubble. The advantages of this process (compared to the extrusion and calendering process) are biaxial orientation of the films, larger widths (potentially up to 12 m with film blowing vs. 3 m in calendering) and no loss through edge trimming. Cellulose derivatives were the first thermoplastics formed in the film blowing process in the beginning of the 20th century. By the 1930’s, biaxially oriented films of cellulose acetate were commercially available. Forming of synthetic high polymers, especially LDPE, by the film blowing process began only in the 1940’s. 10.7.3.2 Processing Parameters Film blowing die designs generally involve a multiple threaded helical mandrel that discharges the melt into a smooth cylindrical land. Depending on the desired film thickness, the land length is 10 to 15 times the annular gap. Cooling is

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extremely important to maintain good product properties and uniformity. Gentle cold air is generally better than warmer air at high velocities. Proper cooling enhances bubble stability, increases the line speed, and augments the production rates. The film thickness and the degree of uniaxial and biaxial orientation of the material are controlled by the blow-up ratio and the haul-off rate. The haul-off rate controls the film orientation in the machine direction, while the blow-up ratio controls the orientation in the transverse direction. Common values for the blow-up ratio are in the range of 1.5-4, depending on the material and desired film thickness. Take-off speeds are usually around 10 to 50 m/min. Appearance of semi-crystalline polymer films changes during cooling from transparent to hazy at a point in the bubble called the frost line. Local unevenness in the film can be overcome by rotating the die or the haul-off and wind-up gear. The film has to be cooled down well below Tm or Tg (e.g., around 50°C for PE) before it reaches the pinch rolls. This prevents blocking and reduces shrinkage on the rolls. 10.7.3.3 Film Blowing of Blends The film blowing operation is characterized by a biaxial drawing of the melt as it leaves the die. The deformation of the dispersed phase during drawing can be described by the microrheology. For sufficiently large capillarity number κ (see Chapter 7, Rheology of Polymer Alloys and Blends), thus for large drop diameter and/or stress, the drop deformation follows essentially that of the matrix, leading to the formation of a lamellar morphology. For the production of a 25 µm film of Orgalloy™ R60ES, the final lamellar structure observed corresponded to a biaxial deformation of the initial droplets without coalescence [Glotin, 1992]. It should be noted that poorly compounded blends, i.e., having large diameter of the dispersed phase, are known to show large differences in the tear strength in the machine and the transverse directions.

Lee and Kim examined the effect of lamellar structure of the dispersed EVA phase on the gas permeability of LDPE/EVA blends produced by film blowing (Lee and Kim, 1995, 1996). The viscosity ratio and dispersed domains size had a predominant influence on the formation of lamellae. An addition of LDPE-g-MA as a compatibilizer increased the number of particles and reduced the thickness of the layers. Between 5 and 6 wt% of compatibilizer produced a dispersed phase size of around 5 µm and the best oxygen permeability reduction (by a factor of 1600). Above this optimum concentration, the particle size becomes too small, resulting in shorter and thinner, and less effective, lamellae. Dispersed phase stretchability increased when a compatibilizer is present and when the viscosity ratio decreases, but it was not affected by the initial particle size. Polyolefins Most of the LDPE and LLDPE production is used in film applications. These resins have been blended with different types of polyolefins to obtain specific film properties and improve processability. As a result, improved tear and tensile strength lead to downgauging and material savings. HDPE/PIB (Pax-Plus™) HDPE/PIB blends can be produced on almost any LDPE film blowing line. The equipment condition affects the production rate and the quality of the film. A throughput rate of 200 kg/hr per one meter of die circumference is achievable. The HDPE/PIB blends have high viscosity, therefore the extruder should have adequate power. Any screw can be used but for high throughput, high shear mixing screws should be avoided. The extruder should be equipped with a 60 mesh screen-pack, a die with a 0.01 mm gap and melt temperature and head pressure gages. Melt temperature should be in the range of 205-225°C. The air-ring may be either conventional or a dual-lip type. To ensure wrinkle-free finished film, care should be taken to eliminate bubble movement. The collapsing frame should

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be lowered to a short cylinder configuration and collapsing frame and bubble guides should be set to just touch the bubble.

10.7.4

PP/PA blends The mechanical and transport properties of PP/PA blends processed by film blowing were studied [Di Lorenzo et al., 1996]. As a compatibilizer, 4 wt% of a maleic anhydride grafted PP (PP-MA with 3 wt% MA) was used. It was reported that upon addition of 10 or 20 wt% of PA-6 to PP, the films showed reduced tensile modulus, tensile strength and maximum elongation (in the transverse direction). Addition of PP-MA somewhat improved the blends’ performance. For example, the ultimate strain went from 840% for PP, to 16% for PP/PA = 90:10, and to 120% for the same blend with 4 wt% of PP-MA. Permeability to CO2, O2 and N2 were also measured. Although PA has better barrier performance to these permeants, addition of 10 or 20 wt% of PA to PP worsened this property, indicating that the permeant could diffuse through voids between the phases. Addition of PP-MA resolved the problem. The composition PP/PA/PP-MA = 30/70/4 showed tensile and transport properties similar to neat PA, indicating that it is possible to partially substitute PA with PP with significant cost savings.

Synthetic fibers are manufactured by extruding a polymeric fluid through a die consisting of a large number of small holes, the spinneret, followed by drawing and wind-up. The fluid can be either a molten polymer (i.e., polyamides, polyesters, polyolefins) or a solution. The latter is used in these cases when the polymer degrades or does not readily melt, i.e., for fiber spinning of PAN (polyacrylonitrile, acrylics), PVAl, aramid fibers, gel spinning of PE, etc. In cases where a melt is used, the process involves melting the polymer in an extruder followed by pumping the melt through a distribution system. The melt or solution is filtered through screen packs and extruded through the spinnerets. Filtering is critical because any solid impurities may cause the thin fibers to break. Solution spinning is called “dry” if the solvent is evaporated in air or “wet” if the solvent is extracted in a bath containing another fluid, typically water, causing the polymer solution to coagulate. The cross-sectional shape of fibers is sometimes non-cylindrical but flat, “kidney-bean” or “clover-leaf”. Dog bone-shaped cross-section can be obtained in solution spinning when the initial porosity is collapsed by removing the solvent. Hollow fibers can also be spun.

Polyamides PA/BA Blends of polyamide copolymers with butyl acrylate copolymers (Durethan™ C, with PA being a copolymer of ε-caprolactam and other monomers) can be blown into films. The recommended processing conditions are: screws having L/D = 25-33 and compression ratio = 3.5-4.0, barrel feed zone at T = 225-260°C and melt temperature of T = 250-280°C. For multi-layer films, adhesion between layers can be achieved either by addition of an adhesion promoter to either of the two materials or (preferably) by co-extrusion with an adhesive layer.

Fiber spinning

10.7.4.1 General Description of Fiber Spinning

10.7.4.2 Processing Parameters In a typical melt spinning operation, filaments emerge from the hundreds of spinneret holes with diameter d = 150 to 300 µm at speeds of the order of 1000 m/min. The extruded fibers go through a drawing stage during which they are stretched by a factor of 3 to 10. For some fibers, e.g., PET and its blends, the drawing process is carried out at higher temperatures in a horizontal oven with zone temperature control. To fix the fibers’ orientation, they are cooled during stretching, usually in air, but also sometimes in water or some other fluid. Deformation and cooling take place simultaneously with crystallization. The extension

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is beneficial from two points of view: (i) it makes it possible to manufacture small diameter fibers (not accessible by direct extrusion), and (ii) it generates high molecular orientation, thus improved mechanical properties. It is important to note that chemical reactions may take place during compounding or plasticating and pressurizing the polymer blend prior to feeding it to a spinneret. For example, exchange reactions can take place between ester groups of polyesters or amine/amide linkages of polyamides. These reactions result in generation of random copolymers. Condensation polymers in the molten state will also see an evolution of their MW. The MW can, within the residence time in the extruder and transfer line, either increase or decrease, to reach an equilibrium dictated by the moisture level. 10.7.4.3 Fiber Spinning of Blends The matrix-fibril morphology is most common in melt spinning of blends containing low percentage of the dispersed phase, whereas phase co-continuity has been frequently observed for systems containing nearly equal concentration of the two components. The different morphologies can be roughly divided into [Paul, 1978]: (i) analogous to laminate side-by-side structures, (ii) sheath-core morphology, where one phase is encapsulated by the other, and (iii) dispersion of fibrils in the matrix. Clearly, there can be an almost infinite number of structures engendered by combinations of the above three types. The distinction between morphologies (ii) and (iii) is the size of the second phase - in the sheath-core morphology both phases have roughly the same size, while in the matrix-fibril structure the fibrils are significantly smaller in comparison with the fiber size. In all cases, because of the orientation phenomenon inherent to the fiber spinning process, the phases are all aligned in the longitudinal direction. In the fibril-matrix structure, the number and size of fibrils depend on composition, rheological properties of the constituents and the type of deformation during spinning. A “tree-ring structure”, as shown

in Figure 10.4, has been observed under certain flow conditions [Lin, 1979; Utracki et al., 1986]. In virtually all cases, the blends used for fiber spinning are heterogeneous and give rise to structured phase morphologies. On the other hand, fibers of miscible blends obviously yield a single-phase structure. This type of fiber is rare, firstly because the known miscible polymer blends do not lend themselves to fiberforming (usually they are amorphous), and secondly because the effects sought by blending often require properties provided by a two-phase structure. One property that could be modified by fiber spinning a miscible blend is dyeability, an important parameter of some commercial fibers, e.g., used for textile or carpet yarns. An addition of a miscible polymer could enhance the number of dye-active sites. Spinning a mixture of polymers can be done using a mixed-stream process or by spinning precompounded blends. In mixed-stream spinning processes, two different polymer streams are extruded side-by-side. This technique can be useful in special cases, for example to take advantage of different shrinkage behaviors of two polymers, thus creating particular bulking effects. Here, additional equipment is required for melt distribution. Special spinneret hole design is also needed to produce the side-by-side or sheath-core structures of the filament. Some precautions must also be taken to maintain the configuration of the fluids during transport to and through the spinneret hole. Rearrangements of the fluid can be caused by interfacial tension or rheological considerations. The rheological properties of the two polymeric liquids can alter the configuration in order to minimize the total energy of the system by ensuring that the lowest viscosity component is located in the regions of highest shear rate [Southern and Ballman, 1973, 1975]. Thus a sideby-side arrangement of two fluids of significantly different viscosities tends to be unstable. However these processes take time to occur while time in the spinneret hole is limited since the holes are short (L/D ≅ 1) while the throughput is high. There is therefore little concern for the rearrangement, unless the viscosity difference is large.

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The problem is best avoided by matching the two polymer melts’ viscosities at the particular shear stress occurring in the hole. Spinning a pre-compounded blend is simpler and more economical than mixed-stream spinning. The compounded blend can be fed directly to the spinning machine or pelletized to be subsequently re-melted for spinning. The polymers are typically pre-compounded in extruders or mixers to obtain a fine dispersion and then delivered by the extruder, and frequently a gear pump, to the spinneret via a transfer line and manifold. During the lowstress transportation, the blend morphology may be considerably affected by the thermodynamic coalescence processes with little particle breakup. However, stresses increase rapidly as the fluid reaches the spinneret hole. The elongational stresses created at the convergence to the holes strongly affect the structure, normally resulting in a fibrilmatrix structure. The elongational flow induces deformation of dispersed drops into elongated fibrils, which upon reduction of stress may disintegrate into a pearl-necklace structure of fine droplets (through the capillary instability process). Low viscosity and high interfacial tension favor this process. Rapid cooling and take-up force will restraint the phenomenon. The large extensional strains induced by drawing can cause large changes in the morphology of blends. For example, during fiber spinning of PP/PA-6 blends (Orgalloy™ R6000) spherical particles with mean diameter d ≅ 0.9 µm were found to elongate into 100 nm diameter fibrils [Glotin, 1992]. The extent of particle elongation depends on the microrheological parameters, thus the interfacial forces, rheological characteristics of the polymers, hole geometry and the spinning parameters. For example, low interfacial tension and lower elongational viscosity of the dispersed phase than that of the matrix (at the extensional stresses occurring in the convergence) lead to the largest L/D ratios for the fibrils. Frequently a gradation of the fibrillar aspect ratio is observed along the radial direction. This originates from inhomogeneity of the stress field during the fibril formation as well as the presence of a thermal gradient during solidification.

A number of phenomena, such as extrudate swell and melt fracture, depend on blend composition and, in turn, affect the process. Melt fracture obviously puts limitations to the production rate. Swelling of fibers at the exit from spinneret holes was reported to be a non-linear function of the blend composition for a number of systems [Han and Yu, 1972]. The multi-component, multiphase nature of polymer blends affects crystallization kinetics, crystalline morphology and level of crystallinity [Nadkarni and Jog, 1991]. In particular, the presence of one polymer may affect the crystallization of the other, the phase boundary enhancing nucleation of crystallization. If the two blend components have different crystallization temperatures, which is likely, the solid particles of the higher melting component will nucleate crystallization of the lower melting component. It was established that each polymer has a maximum draw ratio, λmax, also called natural draw ratio or maximum strain at break. Owing to the influences of one polymer on another, the relationship between λmax in the blends and that of the components is not straightforward. λmax values larger than that of either component have been reported in blends. Commercial PP/EPDM blends (Modylen™) have been routinely spun. The PP matrix contains EPDM as impact modifier, a filler and an interfacial additive. The compositions can be spun into fibers and yarns and extruded into monofilaments. 10.7.5

Solid State Forming

10.7.5.1 General Description of Solid State Forming A number of solid state forming methods have evolved from metal forming techniques. These methods (e.g., rolling, solid state extrusion, stamping, and forging) use limited amounts of heat. They received limited attention in the field of plastics. In some cases, they can be the best processes for forming particular polymers, such as been the case for forging of ultra high molecular

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weight polyethylene, UHMWPE, which cannot be processed by conventional techniques. Also, thermoplastic composites can be stamped to economically produce high performance parts. Solid state forming is normally understood as forming a semi-crystalline polymer above its glass transition temperature, Tg, but below its melting point Tm. Amorphous polymers have been used but with less success. They are too stiff and too brittle below Tg, whereas above Tg the relaxation mechanisms are too fast for maintaining the induced orientation and shape during the cooling step. Orientation in semicrystalline polymers induces dramatic changes in physical properties [Ward, 1975]. Fibers and films for many years have taken advantage of the improvements in strength brought about by alignment of the macromolecular chains. More recently, ultra high modulus polyethylene fibers, with moduli of the order of 120 GPa [Porter et al., 1994] have been produced by the gel spinning technique, which induces high degrees of orientation to the macromolecules indicated by the draw ratio, λ ≅ 200. However impressive are these achievements, commercial products showing high degrees of orientation have been limited to small dimension fibers and films. During the last decade, extensive research and development efforts have been spent on developing processes for manufacturing structural shapes of highly oriented polymers. These have been focused on solid state deformation by either tensile drawing, hydrostatic or ram extrusion, die drawing, or roll-drawing [Zachariades and Porter, 1988]. The efforts have been focused on commodity semi-crystalline materials, such as PE and PP, and more recently to semicrystalline engineering resins such as POM and PET. The high melting point and chemical resistance of a number of polymers such as PA’s, PES or PEEK make them attractive for structural applications. The principal solid state forming processes that have been used for plastics are ram extrusion, hydrostatic extrusion, rolling, roll-drawing and die-drawing. During solid state ram extrusion, the polymer, in the form of a billet, is placed in the

reservoir of a ram extruder and forced through a die at a temperature just below the melting point. Tension is often applied on the resulting profile. The process of hydrostatic extrusion is similar to ram extrusion, except that the extrusion pressure is applied through a fluid that also acts as a lubricant, helping the material slip on the die surface. The process of rolling has been used extensively for forming metal profiles and sheets and involves forming the material by the action of parallel rolls, much like calendering. In rolldrawing, the additional feature is that the sheet or profile is pulled as it exits the rolls, therefore resulting in a complex stress pattern combining tensile and compressive forces. Die drawing is also derived directly from metal forming technology. It involves pulling a rod through a die to form wires or thins rods. 10.7.5.2 Processing Parameters The most important parameter controlling the final properties is the draw ratio, λ. There is a direct relationship between λ and the mechanical properties, in particular the modulus. The maximum draw ratio is generally limited to about 5 for PET and most PA’s, about 10 for PP whereas values of 15-20 are common for HDPE. Crystallinity also has a strong effect on final properties. Strain-induced crystallization develops during the orientation process as the molecular order progressively increases upon drawing. Temperature has a dual impact on the process. Firstly, it affects the modulus of the material, thus affecting the forces needed to form it. The higher the temperature, the easier is the forming process. However, molecular relaxation also depends on temperature. Increasing the temperature facilitates molecular movement and thus faster relaxation, which in turn tends to lower the final molecular orientation in the material. It is customary to form the material 10 to 20°C below Tm. 10.7.5.3 Solid State Forming of Polymer Blends Solid state forming is a young area of polymer processing and relatively few reports can be found

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on development of the processes. This is even more so for forming of polymer blends. Blends of PEO/PMMA were oriented by solidstate coextrusion in a capillary rheometer [Kim and Porter, 1988]. Addition of PMMA to POM dramatically decreased its drawability from a draw ratio λmax = 36 for pure PEO to λmax = 5 for a blend of PEO/PMMA = 40/60 wt%. Drawing has a strong effect on crystallinity, χ, and melting point, Tm. In pure PEO, χ increases from 83% at λ = 1 to χ = 95% at λ = 36, while Tm increased from 62.4 to 70°C, indicating an increase in crystal size and perfection. A different behavior was found in the blends. The crystallinity remained at χ = 70%, while Tm decreased not only with increasing PMMA content but also with λ. This may be taken as an indication that in POM/ PMMA blends the latter polymer gets trapped between PEO lamellae and disrupts the crystalline morphology. Lee and Cakmak described the solid state extrusion of PVDF and its copolymers with vinylidene trifluoride. The authors described processability of different copolymer compositions as a function of the extrusion rate, draw ratio, and temperature [Lee and Cakmak, 1993]. Khait [1995, 1996] proposed using solid state shear extrusion to recycle commingled plastic waste. A twin screw extruder is used for pulverizing the plastic mixtures into a powder. The ground plastic flakes are comminuted by the combined effects of pressure - between 0.25 and 3 MPa - and shear, at temperatures between 15 and 60°C, to produce fine powder. Powder formation involves the breaking of bonds, which results in reactive sites. The mechanical properties of a HDPE/LDPE/PP = 60:30:10 blend formed after pulverization showed minimal changes when compared to the dry blended one: tensile strength went from 18 to 20 MPa while elongation at break, impact strength and HDT remained unchanged. No commercial activity has been found in this area. A major stumbling block to significant developments with polymer blends seems to be two apparent conflicting requirements for the process: i) having a miscible blend, to avoid creation of voids and flaws because separate

domains would deform differently and ii) having at least one semi-crystalline polymer. 10.7.6

Compression Molding

10.7.6.1 General Description of Compression Molding Similarly to injection molding, compression molding is a sequential process which involves forming a part in a cavity under pressure. However, the material is not injected in the mold while molten. The cold or slightly preheated thermoplastic or partially cured thermoset is transferred in a cavity kept at high temperatures. The cavity is closed and while the polymer melts, pressure is applied through a piston, forcing the material to flow into the cavity. The part is cooled and ejected after complete filling (and curing in the case of a thermoset). Compression molding is the most common forming method for thermosets. Thus compression molding was the preferred processing technique in the early part of the 20th century. Its relative importance diminished as thermoplastics were introduced, and other forming methods better suitable for the new materials were developed. Compression molding of thermoplastics has always been limited to special cases. The advantages of the process are low tooling costs, very little waste produced, possibility of producing very large parts and high flexibility for short runs. 10.7.6.2 Processing Parameters The important parameters are temperature, pressure (or piston force) and time. The process is usually described in terms of pressure cycle. Figure 10.12 shows schematically the different stages in compression molding of either thermoplastic or thermoset materials. The material is heated up during the first stage. The second stage begins when the material has reached its melting point and starts to flow. Once the cavity is filled, the material is packed during the compression cycle. The pressure is maintained during the cooling

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cycle of thermoplastic materials at least until the polymer has solidified. The thermoset material, partially cured at this point, can be ejected and curing can be completed outside the mold.

Figure 10.12. Typical cavity pressure profiles during compression molding. Stages are: 1 - polymer heating, 2 - flow, 3 - compression, 4- cooling. Full and dashed lines represent respectively thermoplastic and thermoset molding processes.

10.7.6.3 Compression Molding of Polymer Blends Compression molding involves limited flow and, consequently, less strain than injection molding. Consequently, shear sensitivity typical of polymer blends is less critical. On the other hand, the cycles are usually longer and the long soak times at high temperatures can affect blends’ morphology through coalescence. Compression molding is not widely utilized for thermoplastic materials and information on commercial blends is scarce. As mentioned above, PP/EPDM blends (Santoprene™, Sarlink™) exhibit high melt viscosity. For compression molding, pellets or preformed slabs should be heated for 20 min at 190-215°C to generate a melt that is then compression molded at 120-190°C under pressure of 1.5-3 MPa. For large parts, preformed slabs should be used to avoid air entrapment. In order to fill the mold cavity, a slight excess (5-10%) of melt is recom-

mended. Pressure should be released once the mold is closed which allows material to retain its shape.

10.8

Forming of Filled and/or Reinforced Blends

10.8.1

Filled and Reinforced Polymers

Reinforcements and fillers are widely used in plastics to extend the use of resins and modify their properties. The technology has begun to be used in the early 1940’s. Glass is the most widely used type of reinforcement: about 80% of composites comprise some form of glass fiber (chopped, continuous, mats, etc.) [Rosato, 1989]. In thermoplastics, the fillers and/or reinforcements can be in the form of particles, fibers or flakes. Fibers can be particularly effective since the flow patterns can create preferential orientation distributions and, therefore, directional properties. Table 10.36 shows typical properties of fibers, indicating the wide difference in strength and modulus compared to that of polymers. Glass fibers are typically 1-2 to 10-50 mm long, with an aspect ratio of 10-50. Table 10.37 lists a number of properties of selected polymers with and without glass fibers (GF) [Rosato, 1989]. Density, strength, stiffness and thermal conductivity systematically increase with addition of glass fibers, whereas the ultimate elongation and mold shrinkage are drastically reduced. Thermal and rheological behavior is also affected by addition of fillers, and thus these have impact on processing. Adding glass fibers increases the thermal conductivity, but only slightly affects the specific heat. Filled resins are easier to mold to tight tolerances, as illustrated by the lower mold shrinkage values. 10.8.2

Filled and Reinforced Polymer Blends

Introduction of fillers or reinforcements influences not only properties but also processing and, most frequently, the blend morphology. Several material parameters that influence the forming are

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Polymer Blends Forming

Table 10.36. Properties of some reinforcing fibers Fiber

Density (kg/m3)

Tensile strength (MPa)

Tensile modulus (GPa)

Glass Boron Carbon Aramid

2500-2550 2630 1750-1940 1440

3500-4590 4590 2100-2750 2750

72-85 400 185-380 125

Table 10.37. Effect of glass fiber addition on some polymers [Rosato, 1989] Polymer

PP PP PA-66 PA-66 PC PC PBT PBT

GF content (wt%)

Density

0 20 0 30 0 10 0 30

890 1040 1130 1480 1200 1260 1310 1520

(kg/m3)

Tensile strength (MPa)

34 45 79 159 66 83 57 131

Tensile modulus (GPa)

Tensile strain (%)

Impact strength * (J/m)

Thermal conduct. (W/mK)

Mold shrinkage (%)

0.7 3.7 2.8 8.3 2.3 5.2 1.9 8.3

200 3.0 60.0 1.9 110.0 9.0 50.0 4.0

50-1000 59 43 117 854 107 43 96

2.10 14.5 2.94 2.60 2.34 7.97 1.8-2.9 12.1

0.020 0.003 0.008 0.002 0.006 0.005 0.017 0.003

* Izod impact strength at 22°C

affected by the presence of fillers, viz., viscosity, mold shrinkage, heat deflection temperature, crystallization rate, thermal conductivity, etc. Most notably, the presence of fillers adds new phases to the system (the reinforcement and the sizing agent) and creates several new interfaces: between the new phases and each of the phases already present in the system. Each of these interfaces has to managed. Proper interactions have to be ensured between the different phases. Presence of fillers increases viscosity of the melt, particularly in the case of fibers. This brings about a large increase in torque and pressure requirements, often requiring higher processing

temperatures. For example, unfilled PPE/HIPS and PPE/PA blends are processed at temperature between 260 and 300°C, whereas 10 to 30°C higher temperatures are needed for filled materials. Fiberglass-filled blends are normally processed at temperatures up to 320°C [Bussink et al., 1993]. One of the most tangible impacts of reinforcements on the forming processes is indicated by the mold shrinkage values, shown in Table 10.38. Lower mold shrinkage translates into tighter molding tolerances. On the other hand, the table also indicates that mold temperature specifications for injection molding are higher for blends with

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Table 10.38. Effect of filler content on mold temperature and shrinkage Blend

Trade name

Grade

PA-66/elastomer PA-66/elastomer PA/PP PA/PP PA-6 or PA-66/ABS PA-6 or PA-66/ABS PBT/ABS PBT/ABS PC/ABS PC/ABS PPS/PPE

Ultramid™ A Ultramid™ A Orgalloy™ R Orgalloy™ R Novalloy™-A Novalloy™-A Novalloy™-B or Cevian Novalloy™-B or Cevian Novalloy™-S Novalloy™-S DIC™ PPS

Unfilled 30-40% GF Unfilled 20-30% GF Unfilled Filled Unfilled Filled Unfilled Filled SE-730 (30% GF)

PPS/PPE PA/ABS PA/ABS PA-6/aramid/PTFE PA-6/aramid/PTFE PA-12/aramid/PTFE PA-12/aramid/PTFE

DIC™ PPS N5 N5 RTP™ 200A TFE RTP™ 200A TFE RTP™ 200F TFE RTP™ 200F TFE

SE-740 (40% GF) 15% GF 30% GF Unfilled 30% GF Unfilled 30% GF

glass fibers. Higher mold temperatures slow the solidification process, providing more time for the higher viscosity melt to flow into the mold before solidification. GE Plastics markets a range of composites (Azloy™) based on glass fiber reinforced PC, PC/PBT (Xenoy™) or PPE alloys. The glass fiber reinforcement is in the form of mats or chopped fibers. Parts with these materials are formed by compression molding. The materials are commercially available in the form of precompounded sheets from which blanks are cut and stamped into the final shape following the standard rules for glass filled materials. For example, better flow is obtained with the reinforcement in the form of chopped fibers than with mats. The type of reinforcement determines the ability to fill complex parts, such as parts with ribs. Some blends contain graphite fibers. For example, Xenoy™ grade DX6302 contains 20% graphite fibers, which are electrically conductive. It is

Mold T (°C)

Shrinkage %

40-60 80-90 20-30 50-80

1.4 0.25-1.05

0.7-0.9 0.2-0.4 0.7 0.2-0.4 0.5-0.7 0.15-0.3 0.3-1.1 0.24-1 0.004% 0.002% 1.7 0.6 1.3 0.5

not recommended to regrind this material. Sensitive electrical equipment should be protected and explosion-proof motors are suggested. Addition of reinforcements can also reduce wear and adhesion. Reinforcement by glass, carbon and aramid fibers enhances wear resistance by increasing thermal conductivity and creep resistance. RTP produces series of engineering and specialty thermoplastics lubricated with PTFE, molybdenum disulfide, graphite powder, silicone and glass, carbon and aramid fibers. Molybdenum disulfide is primarily used in PA compounds. It works as a nucleating agent, promoting increased crystallinity in PA, thus providing harder, more wear resistant surface. It also has high affinity to metal surfaces, filling pores and rendering metal surface more slippery. In the case of graphite, its lattice structure allows its molecules to slide past one another thus decreasing friction. This property is emphasized in aqueous environment, making graphite extremely useful in underwater

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applications. Glass fiber, if well bound to the matrix, improves wear resistance by limiting creep and heat deformation and increasing thermal conductivity. Carbon fibers act similarly but additionally are softer and thus unable to score the surface of iron. Aramid is much less abrasive than glass or carbon fibers, essential if the mating surface is sensitive to abrasion.

10.9

Quality and Process Control

Constant pressure for increased productivity and improved product quality is forcing plastic product manufacturers to examine closely process control issues. The topic of control of polymer processing operations and quality control, QC, of the products has been the subject of increased attention [Halasz, 1993]. A number of conventional process parameters are traditionally controlled during processing such as temperature, pressure, throughput or screw rotation rate. In addition, finer and more extensive control is needed in order to achieve the desired blend properties. Variables such as rheology, composition, morphology, color, optical properties, electrical properties may need to be controlled. It is desirable to include the property measurements in closed loop control systems so that disturbances such as variations in feed or changes in equipment conditions can be compensated. QC of polymer products will first be discussed, followed by an examination of techniques available to measure on-line properties that can be particularly important for blends.

ability for a particular process whereas postprocessing QC enables to verify that the material did not degrade during processing. Usually, this is done by measuring either the melt flow index or intrinsic viscosity. The former (often formally referred to as melt index, MI) is normally used for polymer blends. MI corresponds to the weight of polymer extruded through a capillary during a ten minutes period under specified temperature and load conditions. Low MW materials have high MI and low viscosity and are “relatively easier to process”. Melt index test conditions are given in Table 10.39. This test, despite the fact that an absolute measure of viscosity cannot be obtained, is widely used and most polymers are specified on the basis of their MI value. International standards describe the test, viz., ASTM D1238 in the U.S.A., DIN 53735 in Germany, JIS K7210 in Japan, or the international ISO R1133/R292. It should be recognized that MI corresponds to one point on the flow curve, viscosity vs. shear stress. Differences in structure (MWD, branching, etc.) influence the shape of the curve. Therefore MI might sometimes lead to erroneous estimations of the flow behavior if the processing is to take place under conditions (shear rate and stress) that are significantly different from those of the test. Ranges of MI recommended for different processes are given in Table 10.40. The test is sometimes performed at two loads, making it possible to calculate the melt flow ratio (MFR), which is the ratio of the MI values at the two loads. MFR can been related to the molecular weight distribution, MWD. 10.9.2

10.9.1

On-line Monitoring

Quality Control

Quality control of thermoplastic materials has relied predominantly on viscosity measurements, either in the melt or in solution. This is also the case for polymer blends. For example, QC of PPS/PPE (DIC ™PPS) blends (prepared by reactive processing) relies on measurements of density and melt viscosity in a capillary rheometer. Viscosity control before processing allows to assess incoming materials variability and suit-

Advanced control systems are becoming a necessity because of the ever increasing pressures toward product quality and reliability. All material processing operations can benefit from improved process control. However, increased control on the more complex systems, such as composites, compounds and blends, and on the more critical applications will yield the highest benefits. In addition, process control schemes based on process variables alone are not sufficient anymore

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Table 10.39. Melt index (MI) testing conditions Material Blend PE PP PS ABS SAN PMMA PVA POM PA PET PC PPS PBT/PET PET/PC PC/ABS/SAN, PC/ABS

MI test conditions Trade name

Temperature (°C)

Plexiglas, Perspex

Valox Xenoy Bayblend, Bayblend T

Load (N)

190 230 200 200 220 230 150 190 235-275 250 300 315 250 250 250

212 21.2 49 49 98 11.8 21.2 21.2 3.2-49 21.2 11.8 49 50 50 50

Table 10.40. Melt index ranges corresponding to forming processes Process Injection molding Rotational molding Film extrusion Blow molding Profile extrusion

since these might not reveal all material property variations. For example, drifts in raw material characteristics might go unnoticed. It is therefore necessary to monitor material’s characteristics. The technical and commercial literature has witnessed, over the years, ever increasing efforts to improve process control systems, as exemplified by the numerous attempts to monitor polymer properties during processing, such as composi-

MI range (g/10 min) 5-100 5-20 0.5-6 0.1-1 0.1-1

tion, viscosity and dispersion of a minor phase. Typically, an efficient control scheme would rely on a real-time analysis of the polymer characteristics during its manufacturing, compounding, or processing, while the polymer is in the molten state. As opposed to the traditional offline control, where a discrete sample is taken from the process flow line and characterized in the laboratory, in-line process control requires that the

Polymer Blends Forming

characterization be performed directly in the process stream. Consequently, an in-line monitoring device would have to deal with the unfriendly processing conditions: elevated temperature and pressure, fluctuating conditions, chemically aggressive environment, electrical noise, dust and vibrations. On-line control can also be envisaged, where a sampling stream is drawn from the process flow line, usually by the use of a gear pump, and transferred to the measurement apparatus. Each approach has its pros and cons. Sampling delays are very short, or non-existent, for in-line devices, whereas they can be significant in on-line installations, due to the transit line and gear pumps. However, an in-line apparatus may interfere with the main process, and is also very dependent on the temperature and pressure of the melt, so compensation may be required. On the other hand, on-line devices are isolated from the main stream by the use of a gear pump, and temperature and pressure of the polymer sampling flow can thus be controlled. Their maintenance can therefore be done without a complete process shutdown. Development of process instrumentation is often accomplished through the transfer of well established laboratory technique toward the processing line, with modifications in order to enhance their robustness against the severe in-plant conditions. This is typically the case for rheometry or infrared spectroscopy (IR), for which recent developments have led to a number of commercial on-line instruments. Other techniques, such as ultrasonics, are still at the research level even for laboratory characterization, but because of their nonintrusive nature and their mechanical simplicity, they became rapidly very attractive for in-line monitoring applications. Rheometry is the most widely used on-line monitoring technique. Rheometers of different types have been developed for use in closed loop process control. Commercially available instruments all use steady state shear deformation. One commercial instrument attempted to operate in the dynamic mode, Rheometrics’ On-line Rheometer (ROR) but is not currently commercialized. In dynamic mode, the frequency dependence

747

of the storage and loss shear moduli, G’ and G”, respectively, was measured. The empirical observation that for PP, the moduli cross-over point is related to MW and MWD can and has been used for real time determination of MW data [Zeichner and Patel, 1981]. Steady state shear flow measurements provide shear viscosity as a function of shear rate or stress. Both capillary and slit geometries have been used. Great care has to be exercised when using rheometry, whether on-line or in the laboratory, with polymer blends since the rheological response of the material greatly depends on the morphology, which in turn can easily be modified by the flow. This will occur in real time rheometry, die flows or during transfer using gear pumps. In principle, the steady state mode is not well suited for multiphase systems, since the flow may induce morphological changes, migration of components and the measurements may reflect the properties of the skin layer rather than the bulk. Thus, the rheological properties measured might not represent the real blend behavior outside the measuring device. The dynamic mode is more adequate to probe the structure because the normally used smaller strains are less prone to affect the morphology. However, dynamic instruments are more complex and have not met with commercial success. On the other hand, steady state shear is very useful insofar as it is more representative of the processing conditions. Attempts have been made to adapt infrared spectroscopy, IR, for real time monitoring of polymer processing operations. Chemical bonds absorb IR radiation at characteristic frequencies. The resulting spectra are thus characteristic for each chemical species. IR has traditionally proved very powerful for laboratory analysis and compound identification. Early attempts using the mid IR range for on-line monitoring were plagued with several disadvantages such as: (i) very small possible path lengths (typically 15-30 µm) due to high absorptivity, (ii) lack of suitable fiber optics, resulting in the need to build complex delivery systems to bring the melt to the IR cell. Although off-line near-infrared (NIR) analysis is a well-known technique, only a few papers

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M. M. Dumoulin

dealing with the use of NIR for in-line analysis of polymers have been published. The NIR region (12000 to 4000 cm-1) provides the overtones and combination bands of the fundamental vibrational bands observed in the mid-infrared (mainly C–H, O–H and N–H stretching). The low absorptivity of the overtone vibrations (10 to 1000 times less than for the fundamental ones) allows the analysis of thicker samples. The broad and overlapping nature of the bands causes difficulties for quantitative analysis, but these can be overcome by the use of chemometric techniques like the multivariate statistical method known as PLS (Partial Least Squares) [Hansen and Khettry, 1995]. On the other hand, the NIR range lends itself to the use of fiber optics, allowing sensors to be mounted directly on the process stream. NIR measurements were used to monitor composition of PPE/PS during extrusion [Hansen et al., 1993; Thomas et al., 1995]. Fiber optics probes were mounted in a cell in which the polymer was flowing. The composition was varied between 0 and 75 wt% PPE and, after calibration via multivariate analysis, could be predicted with very low standard error of prediction, typically 0.7-0.8 wt%. Other systems tested include polyphthalamides blends and EVA samples with vinyl acetate content varying from 9 to 35 wt% [Hansen and Khettry, 1995; Khettry and Hansen, 1995]. NIR can also be used to monitor extent of reaction during reactive processing of polymer blends. Attempts have been made to determine morphology of blends in processing conditions. Online morphology determination is clearly in the early stages of development. Information in the literature is scarce, let alone any commercial development. Multiple light scattering patterns at the surface of model compounds was related to the morphology of the dispersed phase [Bélanger et al., 1992]. Ultrasonic velocity and attenuation through flowing polymer melts have been shown to be correlated with the size of filler particles and to the morphology of polymer blends [Gendron et al., 1993, Gendron et al., 1995]. Efforts have also been directed at on-line microscopy observations through a special quartz cell mounted on extrusion dies [Gogos, 1994; Yu et al., 1995].

10.10

Reprocessing

10.10.1 Introduction In general, trim or scrap from the fabrication of parts may be reground and reused when care has been taken to prevent contamination and thermal degradation. Certain amounts of uncontaminated regrind may be used with virgin material with minimal or no loss in physical properties and color shift. With contaminated regrind, or for applications where low temperature impact is critical, greater care must be used and the level of regrind should be evaluated for each particular application. Tables 10.41 to 10.43 show the amount of regrind recommended by the manufacturers that should allow processing without loss of physical properties of the finished part. Further details on reprocessing can be found in Chapter 16 of this Handbook, entitled: Role of Polymer Blends’ Technology in Polymer Recycling. 10.10.2 Reprocessing of Commodity Blends Styrenics HIPS (Polystyrol™) and PS/butadiene rubber blends (Styroplus™ and Styroblend™) can be recycled with PS, at the rate of up to 50% regrind. Processing of stock containing regrind does not affect properties (unless the regrind was previously degraded), even if different grades are mixed together. Degradation causes an increase in free monomer content, decrease in MW, increased crosslinking of rubber, changes in morphology and tendency to yellowing. When recycling ABS/BA or PB blends (Terluran™), it should be noted that ABS is not miscible with any thermoplastic except PVC. Addition of any other thermoplastic will reduce the strength. 10-20% regrind of ABS or AES/PC blends (Techniace™ TC) can be used with virgin material with minimum loss of physical properties. When regrind is contaminated or when the molded parts will be used where low temperature impact strength is critical, the use of regrind should be avoided or reduced. In the case of ABS/TPU blends (Prevail™), up to 20% painted

749

Polymer Blends Forming

Table 10.41. Acceptable regrind content for blends of commodity resins Blend

Trade name

LDPE/PIB PS/butadiene rubber PP/EPDM PS-co-PMMA/elastomer ABS/PC or AES/PC ABS/TPU SAN/elastomer SAN/PB Acrylic based

Lupolen 250 H Styroplus Hifax Zylar Techniace TC Prevail Luran S Magnum Acrylite, Cyrolite and XT

Regrind content (wt%) 30 50 25 15 20 20 30 10-20 100

Table 10.42. Acceptable regrind content for blends of engineering resins Blend

Trade name

PA-6, PA-66 or PA-612/ ionomer or elastomer PA-6/elastomer PA-66-co-6/elastomer PA-6/6T/elastomer PA-6 or PA-66/ABS PA-6 or PA-66/ABS PA-6/BA copolymers PET/elastomer PBT/ABS PBT/ASA PC/BA-co MMA PC/ABS PC/ABS PC/ABS/SAN PC/PBT or PC/PET PC/PET/ABS, PC/PET/HDPE PC/TPU or TPU/PC PPE/HIPS PPE/HIPS, PPE/PA POM/TPU

Zytel Grilon A Ultramid C Ultramid T Triax 1000 Novalloy-A Durethan Rynite Novalloy-B or Cevian Ultrablend S Makrolon Bayblend T Pulse Bayblend Sabre Makroblend Texin Luranyl Noryl, Noryl GTX Ultraform

Regrind content, wt% 25 25 10 10 25 20 10 25 20 20 25 20 10-20 20 10-20 20 20 30 30 20

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M. M. Dumoulin

Table 10.43. Accptable regrind content for blends of specialty and elastomeric resins Blend

Trade name

TPU/ABS PSF/ABS Nitrile rubber/PVC, cPE, TPU, PA NBR/EPDM

Mindel A Chemigum TPE Geolast

regrind may be used if care has been taken to prevent contamination or thermal degradation. Acrylics When recycling acrylic based compounds (Acrylite™, Cyrolite™ and XT) 100% regrind can be used without an effect on properties if properly dried and free of contamination. However, if regrind contains fines these may adhere to the screw and eventually cause discoloration. 10.10.3 Reprocessing of Engineering Blends In general, engineering resin blends can be reprocessed. PA-6/ABS or PA-66/ABS (Novalloy™-A) and PBT/ABS alloys (Novalloy™-B or Cevian™) are thermally stable. During injection molding, regrind from sprues and runners can be recycled. Up to 10% regrind can be used for PA-66/ elastomer blends (Ultramid™ A), PA-6/elastomer (Ultramid™ B) and PA-66-co-6/elastomer blends (Ultramid™ C). Reground glass-reinforced Ultramid™ is particularly sensitive to moisture. Mixing regrind with virgin pellets may alter the feed and flow characteristics, demolding, shrinkage and in particular the mechanical properties. As a rule, extrusion scrap is not recycled. Different grades are miscible to a certain extent. For example, small amounts of Ultramid™ B may facilitate processing of glass reinforced Ultramid™ A, but excessive differences in viscosity must be avoided. Ultramid™ is immiscible with most other thermoplastics and even small quantities of these materials may create laminate structures.

Regrind content, wt% 20 25 25 100

For PET/elastomer blends (Rynite™), some loss of properties results from the addition of regrind. About 10% loss for 25% regrind, and 15% for 50% regrind was reported. Blends of POM/TPU (Ultraform™) must be reprocessed with care. When the processing stock contains regrind of different pellet size, the change in the flow properties may affect the product characteristics. Furthermore, since these blends are immiscible with other thermoplastics, contaminated regrind may lead to delamination. Contamination with other thermoplastics, especially with small amount of PVC, may cause sudden, uncontrollable decomposition engendered by dehydrochlorination of PVC. Formed parts of PPE blends can be ground and reprocessed without problems as long as no significant thermal or mechanical degradation took place and no contaminants were introduced. For optimum processing, the recycled product should have a size similar to the final granules and careful drying is needed before processing. An addition of 30% recyclate is possible without the loss of the essential properties.

10.11

Conclusions

Over the years, developments in the area of polymer blends have primarily aimed at improving final properties, particularly impact resistance. To obtain the maximum impact strength, the morphology should be of the matrix/dispersed phase type, and the dispersed phase must be in the

Polymer Blends Forming

form of spherical particles uniformly distributed and well bound with the matrix. One of the methods of binding the particles to the matrix is by grafting these with the matrix polymer. Such particles were found to be less prone to agglomeration in shear field than those without grafting. In commercial blends, the morphology has been frequently stabilized by addition of a compatibilizer or by chemical reaction. This is the case for the toughened engineering polymers containing relatively small concentrations of an impact modifier. Other typically observed morphologies are co-continuity of phases and dispersion of lamellar domains. Co-continuity might allow to draw on properties of both constituents, yielding for example a stiffness/ductility balance not achievable with pure polymers. Lamellar morphologies will obviously yield improved barrier performance when the proper constituents are selected. On the other hand, blending is increasingly carried out in order to improve processability. In general, a lower viscosity is considered beneficial to processing. Blending immiscible polymers will often yield a negative deviation from the viscosity additivity rule, opening the way to easier processing. Blending technology can also enable easier processing indirectly: adding a toughener can allow to reduce a resin’s MW to lower viscosity while maintaining the system’s toughness. On the other hand, the last few decades saw the introduction of a number of high performance resins that proved to be difficult, if not impossible, to process. PPE, PEI and aromatic polyamides are good examples of such materials whose processing relied on blending. Blending can also bring about less dramatic, but nevertheless useful, changes such as less temperature dependent rheology, an advantage gained when adding PP to PA, or increased melt elasticity. For a large number of systems, processability is not the primary rationale for blending. Nevertheless, processing must be carefully examined since blend properties depend on properties of ingredients, composition and morphology. Therefore, tailoring polymer blends for specific applications requires control of the morphology. The right

751

morphology has to be generated and preserved. Theoretical rules guiding morphology control during forming are still sketchy and will require significant work. It is however already evident that in general, blends require tighter control over the process parameters, in particular processing temperature or residence time. Too severe processing conditions can not only cause thermal and/or mechanical degradation as in simpler systems, but in the case of blends, they can trigger various detrimental mechanisms such as damage to the grafting around dispersed particles or aggregation of particles in high shear regions. In injection molded parts, blends can be particularly sensitive to the presence of weld lines, as the flow patterns in welded areas might not yield the optimum dispersed particle size for effective toughening. Flow can impact on the blend’s morphology in a number of ways. It can affect either the drops, causing drop deformation to ellipsoids or fibrils, break up or coalescence, or the whole system, inducing then domain segregation, encapsulation, interlayer slip or variations of miscibility. Examples of each of these phenomena abound. The last decade saw tremendous growth in the use and understanding of polymer blends. A number of challenges still remain. The problem of preserving the right morphology is central to polymer blends technology. One of the challenges is to overcome the processing difficulties by developing new alloys and blends with robust morphologies, i.e., morphologies that do not unravel under commercial processing conditions. Scale-up of polymer mixing and forming is still very difficult. Classical scale-up rules do not necessarily allow preservation of the morphology. In addition, the need to preserve the morphology makes the blends more sensitive to process variables such as temperature, pressure and shear. This calls for efficient process monitoring and control. It also justifies the efforts spent toward the development of techniques for measuring material’s properties on-line. Techniques able to rapidly determine blend morphology, online preferably, are needed and are being investigated. When available, such techniques will help accelerate the already fast growing penetration rate

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of polymer blends in an ever expanding number of commercial applications.

10.12

References

Anders D., Kunststoffe, 69, 194 (1979). Anon., Rheometrics ROR Technical Brochure. Badum, E., U. S. Pat., 2,297,194, 29 Sep. 1942, Appl. 21 Dec 1938, to Bergisch-Gladbach. Bapat, P. M., Harris, J. E., and Chen, Y. T., Europ. Pat. Appl., 477,757, 1 Apr 1992, Appl. 1990, to Amoco Corp. Barger, M. A., Schrenk, W. J., Pawlowski, D. F., and Bremmer, J. N., Proc. Injection Molding Outlook Regional Technical Conference, Society Plast. Eng., Dallas, TX, A-1 (1993). Bartges, B. A., in Handbook of Plastics Materials and Technology, Rubin, I. I., Ed., J. Wiley & Sons, New York (1990). Bélanger C., Cielo P., Favis B. D., and Utracki, L. A., Polym. Eng. Sci., 32, 1886 (1992). Bird, R. B., Stewart, W. E., and Lightfoot, E. N., “Transport Phenomena”, Wiley, New York, 1960. Boldebuck, E. M., U.S. Pat. 3,063,872, Nov 13,1962; App. Feb 15, 1960, to General Electric Co. Bussink, J., Sederel, W. L., and Minderhout, W., “Polyphenylenether-Blends” in “Kunststoff Handbuch, 3/2 Technische Polymer-Blends”, Hanser Verlag, Munich (1993). Bragaw, C. G., “Multicomponent Polymer Systems”, R. F. Gould, Ed., ACS Adv. Chem. Ser., 99, 86 (1971). Bucknall, C. B., “Toughened Plastics”, Applied Science Publishers, Ltd., London (1977). Carreau P. J., Ph.D. Thesis, Univ. of Wisconsin (1969) Champagne, M. F., Huneault, M. A., Roberge, H., SPE Techn. Pap., 41, 1552 (1995). Chung, D. I., SPE Journal, 32, 48 (1976). Cizek, E. P., U.S. Pat. 3,383,435, 14 May 1968; Appl. 6 Jan 1965, to General Electric Co. Collins, S. H., Plastics Compounding, Nov./Dec., 29 (1987). Dao, K. C., J. Appl. Polym. Sci., 27, 4799 (1982). Dao, K. C., Polymer, 25, 1527 (1984). Das, A. P., Gattani, S. S., and Ramamurthy, K., J. Polym. Mater., 9, 1 (1992). Dean, B. D., U. S. Pat., 5,115,046, 19 May 1992, Appl. 6 May 1991, to Amoco Corp. Di Lorenzo, M. L., Di Maio, L., and Acierno, D., Polym.

Networks Blends., 6, 147 (1996). Dumoulin M. M., Ph.D. Thesis, Ecole Polytechnique, Montreal (1988). Dumoulin M. M., Utracki, L. A., and Toma, P., Polym. Eng. Sci., 26, 34 (1986). Feng, J., Chan, C.-M., and Kim, S. C., Proc. Polym. Process. Soc. Regional Meeting for Asia/Australia, Singapore, November 26-28 (1996). Fikentscher, H., and Schmidt, F., Ger. Pat., 636,014, 1936, to I. G. Farbenindustrie. Gale, G. M., Plast. Polym., 6, 183 (1970) Gendron, R., Dumoulin, M. M., Piché, L., and Tatibouët, J., Proc. Polym. Process. Soc. Ninth Annual Meeting, Manchester, UK, April (1993). Gendron, R., Dumoulin, M. M., Piché, L., Tatibouët, J., Polym. Eng. Sci., 35, 79 (1995). Glotin, M., Proc. Colloque Génie de la transformation des polymères, Groupe Français des Polymères, Lyon, France, 23-25 November 1992. Grant, T. S., Jalbert, R. L., Whalen, D., Bartges, B. A., and Thompson, T. L., U.S. Pat. 4,923,924, 8 May 1990, Appl. 6 Dec 1985, to Borg-Warner Chemicals, Inc. Gogos, C., SPE Techn. Pap., 40, 270 (1994) Hagerman, E. G., SPE Techn. Pap., 19, 278 (1973). Halasz, L., “Control Methods in Polymer Processing”, Studies in Polymer Science, vol 10, Elsevier, Amsterdam, 1993. Han, C. D. and Yu, T. C., Polym. Eng. Sci., 12, 81 (1972). Hansen, M. G., Khettry, A., Batra, J., and Steward, D. A., SPE Techn. Pap., 39, 1437 (1993). Hansen, M. G., and Khettry, A., SPE Techn. Pap., 39, 2220 (1993); ibid., 41, 2820 (1995). Holub, F. F., and Mellinger, G. A., U. S. Pat., 4,258,155, 24 Mar 1981, Appl. 3 Dec 1979, to General Electric Co. Holub, F. F., and Schlicht, W. R., U. S. Pat., 4,908,419, 13 Mar 1990, Appl. 1982, to General Electric Co. Holub, F. F., U. S. Pat., 4,908,418; 13 Mar 1990, Appl. 1982, to General Electric Co. Holsti-Miettinen, R. M., Perttilä, K. P., Seppälä, J. V., Heino, M. T., J. Appl. Polym. Sci., 58, 1551 (1995). Huszman, J., Kunststoffe, 73, 347-348 (1983). Izawa, S. I., Toyama, K., Harada, K., and Sugawara, Y., U. S. Pat., 3,781,382, 25 Dec 1973, Appl. 1 Sep 1971, to Asahi-Dow Ltd. Izawa, S. I., Japan Plast. Age, 1979, 1; J. Soc. Materials Sci. Japan, 41, 789 (1992). Jakopin, S., Franz, P., Plast. So. Africa, 32-46, April (1989). Kadokura, M., Momii, T., Hirata, H., Okuda, Y., and Kobayas-

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hi, H.; J. Vinyl Technol., 15, 51 (1993). Kamal, M. R., Jinnah, I. A., and Utracki, L. A., Polym. Eng. Sci., 24, 1337 (1984). Kamal, M. R., Luhfik, G., Arghyris, L., and Hozhabr-Ghelichi, S., US Pat., 5,188,784, Feb. 23, 1993. Kamal, M. R., Garmabi, H., Hozhabr, S., and Arghyris, L., Polym. Eng. Sci., 35, 41 (1995) Karbashewski, E., Kale, L., Rudin, A., and Tchir, W. J., J. Appl. Polym. Sci., 47, 1143 (1993). Khait, K., SPE Techn. Pap., 41, 2066 (1995); SPE Techn. Pap., 42, 3120 (1996). Khettry, A., and Hansen, M. G., SPE Techn. Pap., 41, 2824 (1995). Kim, B. S. and Porter, R. S., J. Polym. Sci. Polym. Phys., 26, 2499 (1988). Kohlert, C., Dietmar, B., and Kohlert, W., Chap. 6 “Calendering Operations” in “Handbook of Polymer Science and Technology, vol. 3, Applications and Processing Operations”, N. P. Cheremisinoff, Ed., Dekker, New York, 1989. Kyu, T., Saldanha, J. M., and Kiesel, M. J., in “Two-Phase Polymer Systems”, L. A. Utracki, Ed., Progress in Polym. Processing Series, Hanser, Munich (1991). Lavengood, R. E., Nicolais, L., and Narkis, M., J. Appl. Polym. Sci., 17, 1173 (1973). Lee, J. S., and Cakmak, M., Polym. Eng. Sci., 33, 1559 (1993). Lee, S. Y., and Kim, S. C., Proc. Polym. Process. Soc. Regional Meeting for the Americas, Akron, OH, November (1995); ibid., Singapore, November 26-28 (1996). Lin, C.-C., Polym. J., 11, 185 (1979). Liu, N. C., and Baker, W. E., Polym. Mater. Sci. Eng., 65, 357 (1991). Machiels, A. G. C., Denys, K. F. J., Van Dam, J., and Posthuma de Boer, A., Polym. Eng. Sci., 37, 59 (1997). McKelvey, J. M., “Polymer Processing”, John Wiley, New York, 1962. Menges, G., Makromol. Chem., Macromol. Symp., 23, 13 (1989). Murray, T. A., Plast. Tech., 24(11), 83; 24(12), 65 (1978). Nadkarni, V. M., and Jog, J. P., “Crystallization Behavior in Polymer Blends”, in “Two-Phase Polymer Systems”, L. A. Utracki, Ed., Progress in Polym. Processing Series, Hanser, Munich (1991). Parkes, A., Brit. Pat., 11,147, 25 March, 1846. Parsons, C. F., and Suck, E. L., ACS Adv. Chem. Ser., 99, 340 (1971).

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Paul, D. R., “Fibers from Polymer Blends”, in “Polymer Blends”, D. R. Paul and S. Newman, Eds., Academic Press, New York (1978). Porter, R. S., Kanamoto, T., and Zachariades, A. E., Polym., 35, 4979 (1994) Rauwendaal, C., “Polymer Extrusion”, Hanser Verlag, Munich (1986). Robeson, L. M., and Matzner, M., Europ. Pat. Appl., 104,659, 4 Apr 1984, Appl. 29 Sep 1982, to Union Carbide Corp. Rosato, D. V., and Rosato, D. V., “Plastics Processing Data Handbook”, Van Nostrand Reinhold, New York (1989). Rosenquist, N. R., US Patent 4,335,032 to General Electric, Jun 15, 1982. Rust, H., Kunststoffe, 73, 342 (1983). Shih, C.-K., Adv. Polym. Technol, 11, 223 (1992). Shue, R. S., and Scoggins, L. E, U. S. Pat., 4,292,416, 29 Sep 1981, Appl. 21 Feb 1980, to Phillips Petroleum Co. Southern, J. H., and Ballman, R. L., Appl. Polym. Sci., 20, 175 (1973); J. Polym. Sci., Phys. Ed., 13, 863 (1975). Speri, W. M., and Patrick, G. R., Polym. Eng. Sci., 15, 668 (1975). Subramanian, P. M., Polym. Eng. Sci., 25, 483 (1985). Subramanian, P. M., and Mehra, V., Polym. Eng. Sci., 27, 663-668 (1987). Subramanian, P. M., ACS Symp. Series, 423, 252 (1990). Tadmor, Z., J. Appl. Polym. Sci., 18, 1753 (1974). Thamm, R. C., Rubber Chem. Technol., 50, 24 (1977). Thomas, Y., Cole, K. C., Pellerin, E., and Dumoulin, M. M., ACS Spring Meeting, Anaheim, CA, April 1-4, 1995, Polym. Mater. Sci. Eng., 72, 19 (1995). Tomotika, S., Proc. Roy. Soc. London, A150, 322 (1935). Utracki, L. A., Dumoulin, M. M., and Toma, P., Polym. Eng. Sci., 26, 34 - 44 (1986). Utracki, L. A., and Schlund, B., Polym. Eng. Sci., 27, 367, 1512 (1987a). Utracki, L. A., in “Current Topics in Polymer Science,” Ottenbrite R. M., Utracki L. A., Inoue S., Eds., Hanser Verlag, Munich (1987b). Utracki, L. A., Ed., “Two-Phase Polymer Systems,” Vol. 2 in the Polymer Processing Society, “Progress in Polymer Processing” series, Hanser Verlag, Munich (1991). Utracki, L. A., “Encyclopaedic Dictionary of Commercial Polymer Blends,” ChemTec Publishing, Toronto (1994). Ward, I. M., “Structure and Properties of Oriented Polymers,” London, Applied Science (1975). Weihrauch, K., Kunststoffe, 75, 700 (1985). Westover, R. F., Section III of Processing of Thermoplastic

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Materials, E. C. Bernhardt, Ed., Reinhold Publishing, 1959. White, J. L., “Twin Screw Extrusion”, Hanser Pub., Munich, 1992. Williams, M. L., Landel, and R. F., Ferry, J. D., J. Am. Chem. Soc., 77, 3701 (1955). Wilsher, R. S., Brit. Pat., 1,192,204, 20 May 1979, Appl. 31 May 1967, to Princeton Chemical Research, Inc. Wood, R., Plast. Rubb. Intl, 5(1), 25 (1980). Yu, D.-W., Esseghir, M., and Gogos, C., SPE Techn. Pap., 41, 136 (1995). Zachariades, A. E., and R. S. Porter, “High Modulus Polymers”, Marcel Dekker, New York (1988). Zeichner, G. R., and Patel, P. D., Proc. 2nd World Congr. Chem. Eng., 6, 333 (1981).

10.13

Notation and Abbreviations

10.13.1 Notation (Roman letters) d D e Ea ∆Hm L m m1, m2 n P Pm R RH T Tg TK Tm Tp v

particle diameter screw diameter wall thickness activation energy of flow heat of fusion screw length power law factor (also called consistency index) (Eq. 10.2) modified Carreau equation factors (Eq. 10.4) power law exponent (Eq. 10.2) Pressure mold cavity pressure ideal gas constant relative humidity temperature glass transition temperature absolute temperature melting point processing temperature specific volume

10.13.2 Notation (Greek letters) y η

shear rate viscosity

ηo λ sxy κ τ χ

zero-shear or Newtonian viscosity draw ratio shear stress capillarity number relaxation time crystallinity

10.13.3

Abbreviations

AAc ABS ACS

acrylic acid acrylonitrile-butadiene-styrene acrylonitrile-chlorinated polyethylene-styrene terpolymer AES acrylonitrile-ethylene rubberstyrene AN acrylonitrile ASA acrylonitrile-styrene-acrylate BA polybutylacrylate BMA butylmethacrylate CA cellulose acetate cPE chlorinated PE CR screw compression ratio EPDM ethylene-propylene-diene terpolymer EPR ethylene-propylene rubber ESCR environmental stress cracking resistance EVA ethylene-vinyl acetate copolymer GF glass fibers GP gutta-percha HDPE high density polyethylene HDT heat deflection temperature HIPS high impact polystyrene ID internal diameter IP injection pressure LCP liquid crystal polymers LDPE low density polyethylene LLDPE linear low density polyethylene MA maleic anhydride MABS methacrylonitrile butadiene styrene copolymer MAc methacrylic acid MC machine capacity (for injection shot size) MI melt index MW, MWD molecular weight, molecular weight distribution

Polymer Blends Forming

NBR NC NIR NR OD PA PAI PA-6IT6

PA-mXD6 PAr PARA PB PBT PC PE PEEK PEI PE-MAc PE-AAc PEMA PES PEST PET PETG PHR PIB PLS PMMA PnBA POM PP PP-MA PPE PPS PS PSF PTFE PVAl PVC PVME PVT QC

nitrile-butadiene rubber nitrocellulose near-infrared spectroscopy natural rubber outside diameter polyamide polyamideimide co-polyamide of caprolactam, hexamethylenediamine and terephthalic acid PA of m-xylylenediamine and adipic acid polyarylate aromatic polyamides polybutadiene polybutyleneterephthalate polycarbonate polyethylene polyetheretherketone polyetherimide ethylene-methacrylic acid copolymer ethylene-acrylic acid copolymer poly(ethylmethacrylate) polyethersulfone polyester polyethyleneterephthalate polyethyleterephthalateglycol parts per hundred resin polyisobutylene partial least squares polymethylmethacrylate poly(n-butyl)acrylate polyoxymethylene polypropylene maleic anhydride grafted polypropylene polyphenyleneether polyphenylene sulfide polystyrene polysulfone polytetrafluoroethylene polyvinyl alcohol polyvinyl chloride polyvinylmethylether pressure-volume-temperature quality control

RIM SAN SBR SBS

755

reaction injection molding styrene-acrylonitrile copolymer styrene-butadiene rubber styrene-butadiene-styrene tri-block copolymer SEBS styrene-ethylene/butene-styrene triblock copolymer SMA styrene - maleic anhydride copolymer SSE single screw extruder THF tetrahydrofuran TLV threshold limit value TPU thermoplastic urethane TSE twin screw extruder UHMWPE ultra high molecular weight polyethylene

CHAPTER 11

USE OF HIGH-ENERGY RADIATION IN POLYMER BLENDS TECHNOLOGY

Ajit Singh1 and Kamarudin Bahari2

1

AECL, Whiteshell, Laboratories, Pinawa, Manitoba, Canada

2

Nuclear Energy Unit, Kompleks Puspati, Malaysia

Abstract Exposure of organic systems to high-energy radiation (radiolysis, radiation processing) at room temperature leads to the formation of ions and free radicals. These species, in turn, react to give stable products, often through other free radicals and metastable species. The course of their reactions is influenced by the conditions under which the system has been radiation processed; for example, in the presence of oxygen, oxidative degradation predominates. The course of the reactions can also be influenced by the presence of other additives. Ambient temperature processing and the ability to change the course of resulting chemical reactions by suitable additives have been important factors in the attraction of this technology for the polymer industry. The major industrial applications of this technology in the field of polymers are based on (i) polymerization, (ii) crosslinking, (iii) scission, and (iv) sterilization of medical disposables. Radiation processing of polymer blends can lead to crosslinking or scission of one or more components, or even to inter-phase (inter-polymer) crosslinking. These effects modify the blends’ properties. Many research groups have been working on radiation processing of polymer blends, to stabilize the phases and improve performance. The polymers that have been used in the radiolytic studies of blends include polyethylene (PE), polypropylene (PP), ethylene propylene rubber (EPR), polyvinyl chloride (PVC), polystyrene (PS), and polymethylmethacrylate (PMMA). The primary aim of this chapter is to summarize the literature on radiation processing of polymer blends, to highlight the achievements, and to discuss the problems encountered. The text is preceded by brief descriptions of relevant radiation chemistry of organic compounds and known effects of high-energy radiation on the leading polymers used in blends, with the aim of enabling future workers to draw upon this knowledge in the design of experiments, in industrial applications, or both. The effects of radiation processing, first of commodity resin blends, then on the engineering and specialty resin blends, are discussed. L.A. Utracki (Ed.), Polymer Blends Handbook, 757-859. © 2003 Kluwer Academic Publishers. Printed in the Netherlands.

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A. Singh and K. Bahari

Introduction

Radiation processing involves the use of natural and manufactured sources of high-energy radiation on an industrial scale, to give products that are safe, practical and beneficial [Silverman, 1981]. In comparison to conventional methods, radiation processing is an energy-conserving and environmentally benign technology. The number of electron accelerators in the industrial use has been steadily increasing despite the fact that radiation-based technologies face close scrutiny by society and regulators [Saunders, 1988; Machi, 1989; Leemhorst and Miller, 1990; Singh and Silverman, 1992]. The growth of this technology depends on market forces, which restrict the use of industrial electron and gamma processing to applications that either offer significant cost savings or produce unique and useful products [Silverman, 1981; Bradley, 1984; Fraser, 1988; Singh and Silverman, 1992; Kunstadt et al., 1993; Singh and Singh, 1994; Woods and Pikaev, 1994; Singh, 1995]. The major applications of the high-energy radiation processing technology result from (i) polymerization or curing, e.g., of coatings or rubbers; (ii) crosslinking, e.g., in wire and cable industry, or plastic films; (iii) scission, e.g., for degradation of polytetrafluoroethylene (PTFE); and (iv) sterilization of medical disposables [Silverman, 1981; Saunders, 1988; Fraser, 1988; Stannett et al., 1989; Cook, 1990; Schroeder, 1990; Tabata, 1990; Tenorth, 1990; Morrissey and Phillips, 1993]. Radiation processing of polymer blends can also lead to crosslinking or scission of one or more of the polymeric blend phases. These effects may result in the modification of the blends’ properties. Many research groups have been working on radiation processing of polymer blends to improve properties, particularly stabilization of morphology [viz., van Gisbergen and Overbergh, 1992]. Though effort has been made to include as many papers and patents on irradiation of blends, as possible, comprehensive coverage of all the publications on the topic is not claimed. A brief description of relevant radiation chemistry of organic

compounds and known effects of high-energy radiation on the leading polymers used in the blends is also given [Charlesby, 1960; Chapiro, 1962; Williams, 1968; O’Donnell and Sangster, 1970; Dole, 1972, 1973; Wilson, 1974; Farhataziz and Rodgers, 1987; Spinks and Woods, 1990]. To use high-energy radiation sources either in a laboratory or on an industrial scale, one must be familiar with their characteristics, the safety requirements and dosimetry. These topics are well covered in the literature and therefore will not be discussed [Bradley, 1984; Cleland, 1992; IAEA, 1994; Woods and Pikaev, 1994; Singh et al., 1996]. Usually, the suppliers of high-energy radiation equipment are also helpful in providing relevant information to their customers.

11.2

Radiation Chemistry of Organic Compounds

Exposure of organic compounds to high-energy radiation (irradiation, radiolysis) leads to chemical changes through the formation of short-lived ionic and free radical species [Ausloos, 1968; Gäumann and Hoigné, 1968; Burton and Magee, 1969-1976; O’Donnell and Sangster, 1970; Singh, 1972; Földiák, 1981; Baxendale and Busi, 1982; Farhataziz and Rodgers, 1987; Freeman, 1987; Spinks and Woods, 1990; Woods and Pikaev, 1994]. The extent of substrate decomposition and the nature of the radiolytic products formed depend on many factors, including (i) the functionality of the groups; (ii) the environment (for example, the presence or absence of oxygen); (iii) the presence of additives, e.g., olefinic or aromatic compounds; and (iv) temperature and pressure. 11.2.1

Saturated Hydrocarbons

By definition, organic compounds contain mainly carbon-carbon and carbon-hydrogen bonds. For this reason radiolytic products formed by their scission are common to various organic compounds. The main reactions that occur during irra-

Use of High-Energy Radiation in Polymer Blends Technology

diation of liquid saturated hydrocarbons, (RH2), (under inert atmosphere, vacuum or argon, at room temperature, RT, unless otherwise specified) are summarized below [Ho and Freeman, 1964; Gäumann and Hoigné, 1968; Földiák, 1981; Spinks and Woods, 1990]. On exposure to high-energy radiation, the energy deposition process engenders ionized, e.g., RH2+•, and excited, e.g., RH2*, species [Klots, 1968; Willard, 1968; Burton and Magee, 1969-1976; Mozumder, 1972; Freeman, 1987; Sauer et al., 1991]:

For example, in the case of ethyl radicals, the recombination and disproportionation reactions may produce n-butane, ethylene, and ethane: •CH2 - CH3 + •CH2 - CH3 → CH3 - CH2 - CH2 - CH3 (R-12) •CH2 - CH3 + •CH2 - CH3 → CH2 = CH2 + CH3 - CH3 (R-13) 11.2.1.1

RH2 - /\/\/ → RH2+• + RH2* + e-

759

Additives

(R-1)

This is followed by ionic recombination to produce free radicals and excited molecules:

Additives (A) can play several roles; for example, by charge and energy transfer, they can protect the substrate from radiolytic decomposition:

RH2+• + e- → •RH + •H

RH2+• + A → RH2 + A+•

(R-14)

RH2* + A → RH2 + A*

(R-15)

(R-2)

RH2+• + e- → RH2* (R-3) and ion-molecule reactions to produce molecular products, or free radicals: RH2+• + RH2 → RH3+ + •RH

(R-4)

RH3+ + e- → •RH + H2

(R-5)

RH3+ + e- → RH2 + •H

(R-6)

RH2+• → R+• + H2

(R-7)

The excited states formed in reactions (R-1) and (R-3) can react in several ways, e.g., to produce free radicals, molecular products, luminescence and regeneration of the substrate [Calvert and Pitts, 1967; Ausloos, 1968; Singh, 1972; Walter et al., 1976; Katsumura et al., 1983]. Reactions of the free radicals include hydrogen abstraction, recombination, and disproportionation: •H + RH2 → H2 + •RH

(R-8)

•RH + •RH → RH-RH

(R-9)

•RH + •H → RH2

(R-10)

•RH + •RH → RH2 + R

(R-11)

Alternatively, they can react with cation, electron, or free radicals, thus to change the course of the subsequent reaction chain. Examples of such effects of additives are given below. In the case of a substrate with additives, the initial deposition of energy follows the “law of mixtures” — the total energy absorbed by each component of the mixture is proportional to the relative amounts of orbital electrons of each component in unit volume, or in practical terms, it is very nearly proportional to the fraction of the component by weight [Klots, 1968; Spinks and Woods, 1990]. However, it is not unusual to see disproportionately large radiolytic changes in the smaller component of a mixture. This is usually a result of the participation of the smaller component in energy and charge transfer reactions, as discussed below. 11.2.1.1.1 Mixtures of Hydrocarbons The ionization potentials of homologous hydrocarbons decrease with increasing molecular weight [Spinks and Woods, 1990]. Thus in mixtures of two hydrocarbons, say cyclopentane and cyclohexane, positive charge transfer, i.e., (R-14), would be expected to occur from cyclopentane cations to

760

A. Singh and K. Bahari

cyclohexane. This effect has been demonstrated in several studies, particularly by work on mixtures of perdeuterated cyclohexane in cyclopentane [Dyne, 1965; Stone et al., 1966; Gäumann and Hoigné, 1968]. In these systems the radiolytic yield of D2 and HD was much greater than would be expected on the basis of the law of mixtures.

11.2.1.1.3 Halogenated Compounds Halogenated compounds are also good electron scavengers [Gäumann and Hoigné, 1968]. For example, the following reaction occurs in the case of halogenated methane (X = Cl, Br or I):

11.2.1.1.2 Nitrous Oxide Nitrous oxide is known to react efficiently with electrons [Gordon et al., 1963]:

Thus in the presence of halogenated compounds, the radiolytically produced electron may preferentially react with the halogenated molecules. Also, the free radical moiety produced from the halogenated molecule, e.g., the free radicals formed in reaction (R-20), may participate in the reactions leading to the formation of several new products. For example, radiolysis of cyclohexane in the presence of carbon tetrachloride was reported to lead to a chain reaction [Henglein et al., 1963; Stone and Dyne, 1964].

N2O + e- → N2 + O-•

(R-16)

Radiolysis of cyclohexane in the presence of 20 mmol/L of N2O resulted in the reduction of the H2 yield by about one third [Scholes and Simic, 1964]. Later, Asmus et al. [1970] estimated the contribution of ionic mechanisms to the radiolytic production of hydrogen from hydrocarbons, to be about one third. These observations suggest that the recombination of electrons with the cations contributes to the formation of H2 in reactions such as (R-5) and (R-7). In the presence of N2O, the reactions (R-5) and (R-7) could be modified resulting in reduced yield of hydrogen because of the occurrence of (R-16) and the following types of reactions [Földiák, 1981]: O-• + C6H12+• → •OH + •C6H11

(R-17)

C6H12* + N2O → C6H12 + N2O

(R-18)

C6H12 + •OH → H2O + •C6H11

(R-19)

In addition to being an electron scavenger, nitrous oxide can also deactivate excited species by energy transfer, such as reaction (R-18) [Singh, 1972; Klassen, 1987]. Subsequent to the formation of the hydroxyl radicals through reaction (R-17), water formation through reaction (R-19) would be expected. Water formation has been observed during radiolysis of cyclohexane in the presence of N2O [Freeman, 1996].

CH3X + e- → •CH3 + X-

(R-20)

11.2.1.1.4 Oxygen During radiolysis of organic compounds, oxygen can also act as an electron scavenger: O2 + e- → O2-•

(R-21)

However, what is more important, oxygen also takes part in the formation of peroxy-radicals [Ingold, 1969; Singh, 1989], which leads to the radiation-induced oxidative degradation [Spinks and Woods, 1990; Kashiwabara and Seguchi, 1992; Williams, 1992, Soebiantio et al., 1996], viz.: HR• + O2 → HRO2•

(R-22)

HRO2• + RH2 → HROOH + •RH

(R-23)

Reaction (R-23) can be followed by reaction (R-22) leading to chain oxidation of the substrate. On exposure to light or heat, the hydroperoxide can decompose, producing more free radicals. In the case of cyclohexane, as shown by the data in Table 11.1, the yields of free-radical-based products (such as cyclohexene and dicyclohexyl) are reduced in the presence of oxygen [Ho and Freeman, 1964], at least partly because of reac-

761

Use of High-Energy Radiation in Polymer Blends Technology

Table 11.1.

Effect of oxygen on radiolytic yieldsa (G(µmol.J-1)) of cyclohexene and dyclohexene and dicyclohexyl from cyclohexyl from cyclohexane [Ho and Freeman, 1964]

Substrate Cyclohexane Cyclohexane a

Additive

Cyclohexene

Dicyclohexyl

— Oxygen (20 kPa)

0.33 0.15

0.18 0.03

Traditionally, G value has been defined as the number of molecules of product (or of the modified substrate) per 100 eV of absorbed energy. In the SI units, the designation G remains, but the units are µmol.J-1. The old G value divided by 9.63 gives the G in µmol.J-1. In this chapter, the G values are given in the SI units.

tions (R-21) and (R-22). The radiolysis is also affected by the reactions of oxygen with hydrogen atoms, e.g., reactions (R-24) and (R-25): H• + O2 → HO2•

(R-24)

11.2.1.1.5 Unsaturated Compounds In the presence of unsaturated compounds, formed either as a result of radiolysis or added externally to the system, the free radicals can react by addition, to give the secondary free radicals:

HO2• + RH2 → H2O2 + •RH

(R-25)

•RH + CH2 = CH2 → •CH2 - CH2 - RH

(R-30)

•H + CH2 = CH2 → •CH2 - CH3

(R-31)

The organic peroxy radicals (R-22) can also undergo a variety of unimolecular reactions [Ingold, 1969; Dole, 1973b]; the most relevant of these is the chain scission leading to the formation of carbonyl and hydroxy compounds: •O2  CH3CHCH2CH3 → CH3CHO + CH3CH2O• (R-26) CH3CH2O• + RH2 → CH3CH2OH + •RH (R-27) The mutual reactions of the peroxy radicals also lead to the formation of carbonyl compounds and alcohols, along with the formation of peroxides [Ingold, 1969]: R2CHO2• + R2CHO2• → R2C = O + R2CHOH + O2 (R-28) R2CHO2• + R2CHO2• → R2HCO - OCHR2 + O2 (R-29)

A result of reaction (R-31) is that hydrogen formation by reaction (R-8) is reduced. Since, in saturated substrates, unsaturated products are formed by radiolysis, e.g., by (R-7) and (R-13), the yield of hydrogen decreases with increasing dose, along with a decrease in the yield of the unsaturated products. For example, in the case of cyclohexane, the variation of the yields of hydrogen and cyclohexene as a function of dose is shown in Table 11.2. The reduction of the hydrogen and cyclohexene yields is due to the consumption of cyclohexene in carbon-centered free radical reactions, such as reaction (R-30) and (R-31). As shown by the data in Figure 11.1, the yield of hydrogen is also reduced by changing the solvent composition, for example by replacing part of the cyclohexane by either cyclohexene or benzene [Manion and Burton, 1952; Freeman, 1960]. This reduction is attributable to the charge and excitation transfers, as well as to hydrogen atom scavenging (R-14, R-15, R-30 and R-31).

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Table 11.2.

Variation of hydrogen and cyclohexene yields as a function of dose, on radiolysis of cyclohexane [Ho and Freeman, 1964]

Dosea (kGy)

G(H2), (µmol.J-1)

G(Cyclohexene), (µmol.J-1)

0b 20 40

0.58 0.55 0.53

0.32 0.24 0.22

a b

Gamma irradiation. The initial yield is obtained by extrapolation to zero dose.

11.2.2

Aromatic Hydrocarbons

The main difference between the radiolysis of aromatic and corresponding saturated hydrocarbons is the stability of the former towards irradiation; for example, the radiolytic decomposition of benzene (C6H6) is only about one seventh of that of cyclohexane (C6H12) [Földiák, 1981]. As in the case of saturated hydrocarbons, upon irradiation both ionization and excitation of the aromatic substrate would occur: C6H6 - /\/\/ → C6H6+• + C6H6* + e-

(R-32)

The benzene cation is relatively stable, though it could undergo addition to a benzene molecule and recombination with the electron to produce excited states,

Figure 11.1.

Reduction in the yield of hydrogen on irradiation of cyclohexane, with increasing cyclohexene (top) or benzene (bottom) concentration [Manion and Burton, 1952; Freeman, 1960].

C6H6+• + C6H6 → (C6H6)2+•

(R-33)

C6H6+• + e- → C6H6*

(R-34)

(C6H6)2+• + e- → (C6H6)2*

(R-35)

which are relatively stable [Calvert and Pitts, 1967; Holroyd, 1968; Singh, 1972; Födiák, 1981]. The yields of the major radiolytic products formed on γ-radiolysis of benzene (under inert atmosphere) are shown in Table 11.3 [Gäumann and Hoigné, 1968; Földiák, 1981; Spinks and Woods, 1990]. It may be recalled that benzene also acts as a protective agent when present in a saturated hydrocarbon (Figure 11.1), as discussed earlier.

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Use of High-Energy Radiation in Polymer Blends Technology

Table 11.3.

γ-radiolysis of benzene: yields of major productsa G (µmol.J-1)

Product Hydrogen Acetylene C6 to C12, Cyclohexadienes and Bicyclohexadienes Biphenyl Polymer a

0.004 0.002 0.011 0.005 0.114

At RT based on the data given in Földiák [1981].

11.2.3

Halogenated Organic Compounds

The carbon-halogen bonds in iodo-, bromo-, and chloro-organics are generally weaker than the carbon-carbon or carbon-hydrogen bonds; hence they are susceptible to breakage during reactions of excited molecules, as well as to dissociative capture of electrons during radiolysis [Calvert and Pitts, 1967; Spinks and Woods, 1990]. The main distinguishing reaction of the halogenated compounds is the dissociative reaction with electrons: XCH2-CH2R + e- → •CH2-CH2R + X-

(R-36)

In this reaction X denotes either Cl, Br or I atoms. Reactions of the hydrocarbon moiety would

Table 11.4.

Radiolytic yields (G(µmol.J-1)) of major products from butyl chloridesa. Dose = 240 to 770 MGy.

Product Hydrogen Hydrogen Chloride Methane n-Butane iso-Butane sec-Butyl Chloride tert-Butyl Chloride 1,2-Dichloro-2-methylpropane 1,3-Dichlorobutane a

result in products similar to those obtained from hydrocarbons. For example, the yields of the major products formed on gamma-irradiation of n- and isobutyl chlorides are given in Table 11.4 [Dismukes and Wilcox, 1959]. The high yields of isomerization (sec-butyl chloride and tert-butyl chloride, Table 11.4) suggest the occurrence of chain reactions. This has been further confirmed by work on other halogenated substrates in which it was also found that the isomerization reaction is catalyzed by the acids formed (HX) [Takehisa et al., 1966; Martin and Williams, 1970]. In the case of halogenated aromatics, the overall decomposition is lower, and there seems to be no isomerization. The products formed on irradiation of chlorobenzene are given in Table 11.5 [Choi and Willard, 1962].

n-Butyl chloride

iso-Butyl chloride

0.15 0.10 0.01 0.72 0 6.22 0 0.03 0.08

From the data reported by Dismukes and Wilcox [1959], quoted by Spinks and Woods [1990].

0.09 Not determined 0.02 0 0.72 0 2.49 0.23 0

764

Table 11.5.

A. Singh and K. Bahari

Radiolysis of chlorobenzene: yields of major productsa. Dose = 208 kGy.

Product Hydrogen Hydrochloric acid Benzene Biphenyl Dichlorobenzene Chloro-biphenyl a

G (µmol.J-1), 25°C

G (µmol.J-1), 100°C

0.001 0.026 0.301 0.012 0.037 0.093

0.075 0.373 0.016 0.051 0.187

From the data of Choi and Willard [1962], quoted by Spinks and Woods [1990].

During radiolysis of the fluorinated organic compounds, the relative yield of carbon-carbon bond breakage is higher, and that of carbon-halogen bond breakage is lower than that of chlorinated organic compounds [Spinks and Woods, 1990]. 11.2.4

Oxygenated Compounds

A general feature of the radiolysis of oxygenated organic compounds is that the oxygenated site becomes a preferential, but by no means exclusive, center of the radiation-induced reactions. [Farhataziz and Rodgers, 1987; Spinks and Woods, 1990; and Tabata et al. 1991]. The products formed on irradiation of alcohols include hydrogen, hydrocarbons, water, carbon monoxide, glycols and carbonyl compounds. In the absence of oxygen, glycols are among the most important products, following the loss of an α-hydrogen atom and dimerization of the resulting free radicals. The presence of oxygen (in the hydroxyl group) significantly modifies the radiolytic reactions, as compared to the corresponding hydrocarbons. A major difference is the trapping of the radiolytically produced electron by liquid alcohols [Freeman, 1970; Tabata et al., 1991]. Solid hydrocarbons and ethers can trap electrons at low temperatures (196 K). However, with the increasing polarity of the substrate, the trapped or solvated electrons are also observable at higher temperature. The yield of hydrogen decreases with increased branching, whereas the yield of

hydrocarbons increases with increased branching, as shown by the data on the various alcohols in Table 11.6. In the case of ethers, the C-O and the C-C (adjacent to the oxygen) bond scission is favored over the C-H bond scission. The products formed in the case of the radiolysis of diethyl ether are shown in Table 11.7. In the case of the carbonyl compounds, the reactive center of the radiolytic reactions is the carbonyl group. The radiolytic reactions in carbonyls include formation in ketones of carbon monoxide and breakage of the C-H bond adjacent to the carbonyl, in carboxylic acids of carbon dioxide, and in esters of carbon monoxide and carbon dioxide. The products formed during radiolysis of acetone and methyl acetate are listed in Table 11.7.

11.3

Radiation Effects on Polymers

The chemical changes that occur when polymers are irradiated in inert atmosphere (e.g., under vacuum) include crosslinking: inter- or intramolecular bond formation, viz. (R-9); bond scission: chain breaks, viz. (R-26); changes in the level of unsaturation: formation of carbon-carbon double bonds, viz. (R-11) and (R-13); and formation of hydrogen and methane: viz. (R-5) to (R-8); [Charlesby, 1960; Chapiro, 1962; Williams, 1968; Dole, 1972, 1973; Wilson 1974; Wilski, 1987;

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Use of High-Energy Radiation in Polymer Blends Technology

Table 11.6.

γ-radiolysis of liquid alcohols: major products [Spinks and Woods, 1990] G(Product) in µmol.J-1

Alcohol

Methanol Ethanol 1-Propanol 2-Propanol n-Butanol t-Butanol a

H2

>COa

0.56 0.52 0.46 0.38 0.47 0.10

0.21 0.29 0.50 0.19 0.52 0.33

Glycol 0.378 0.230 0.160 0.020 0.161 ~0

Bond cleavage (%) Hydrocarbons

C-H

C-C

C-O

0.07 0.09 0.21 0.16 0.19 0.44

88 84 62 69 65 ~11

0 0 32 27 30 ~77

12 9 6 4 5 ~12

Carbonyl compounds

Table 11.7.

γ-radiolysis of liquid acetone, diethyl ether and methyl acetate — T ~ 22-28°C [Spinks and Woods, 1990] G(Product) (µmol.J-1)

Product

Hydrogen Carbon Monoxide Carbon Dioxide Methane Ethane Ethylene Formaldehyde Acetaldehyde Ethanol Dimethyl Ether Acetic Acid Biacetyl Methyl Ethyl Ketone 2-Propanol CH2=CHOC2H5 (CH3COCH2) 2 CH3COCH2COCH3 CH3CH(CH3)OC2H5 C2H5 CH(CH3)OC2H5 (CH3CHOC2H5) 2

Acetone

Diethyl ether

Methyl acetate

0.052 0.075 0.180 0.025 0.009 0.032 0.029 0.019 0.009 0.028 0.011 -

0.35 0.041 0.14 0.06 0 0.02 0.22 0.11 0.01 0.08 0.27

0.079 0.170 0.099 0.210 0.035 0.016 -

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A. Singh and K. Bahari

Gillen and Clough, 1992]. The chemical reactions occurring in polymers are at least as complex as those observed in their simpler analogues discussed in Section 11.2. However, the polymers do not lend themselves to detailed analyses, as compared to the simpler systems (Tables 11.1-11.7). This is partly because of the large molecular weights of the polymers and partly because their purity is generally lower than that of the simpler compounds (such as cyclohexane). Although many workers have studied details of some of the chemical changes occurring in polymers on irradiation, there has been a greater emphasis on the observed changes in the physical and mechanical properties of the polymers. In the presence of oxygen, oxidative degradation dominates, leading to the introduction of hydroxyl and carbonyl groups, as well as chain scission via (R-21 to R-23) and (R-26) [Williams, 1968; Dole, 1973b; Wilson, 1974; Makhlis, 1975; Shalaby, 1979; Randall et al., 1983; Carlsson et al., 1991; Gillen and Clough, 1992; Kashiwabara and Seguchi, 1992; Williams, 1992; Soebiantio et al., 1996]. The oxidation may cause deterioration of mechanical properties [Clough and Gillen, 1981; Gillen and Clough, 1989a and b; Bradley, 1984; Clough et al., 1991; Kashiwabara and Seguchi, 1992; Kubo and Otsuhata, 1992]. Oxidative degradation occurs not only during radiolysis in the presence of air or oxygen, but, if oxygen is available, it can take place for months during the post-irradiation period. In polymers, the oxidative degradation is caused by the free radical reactions that lead to the formation of the peroxy radicals, viz. (R-22). The role of the peroxy radicals has been elucidated by several groups, e.g., by the luminescence and electron spin resonance methods, ESR [Partridge, 1972; Dole, 1973b; Kadir et al., 1989; Kashiwabara and Seguchi, 1992; Williams, 1992; Jipa et al., 1993]. The presence of antioxidants may reduce the extent of the degradation, but progressively the antioxidants get consumed [Bourges et al., 1992]. Several factors affect the extent of oxidative degradation, e.g., solubility of oxygen, dose rate and thickness of the sample (oxygen diffusion) [Charlesby, 1960; Dole, 1972, 1973; Clough

and Gillen, 1981; Gillen and Clough, 1989 a,b; Clough et al., 1991; Gillen and Clough, 1992; Kashiwabara and Seguchi, 1992; Williams, 1992]. The additives usually are excluded from the crystalline regions, e.g., in PE or PP [Winslow et al., 1966], and since the diffusion of oxygen into the crystalline regions and its solubility there are negligible, long-lived free radicals are formed in the crystalline regions [Dole, 1972, 1973; Clough, 1987; Kadir et al., 1989]. These slowly migrate to the crystalline surfaces and then react with the oxygen present in the amorphous regions, during the post-irradiation period. The role of crystallinity is evidenced by the post-irradiation oxidation of PE; it was found that the postirradiation oxidation was higher in HDPE than in LDPE [Dole, 1965]. Consistent with the above observations is the recently proposed order of the intermolecular reaction rates in polyethylene oxide: amorphous > interfacial > crystalline [Zhang et al., 1992]. The role of oxygen diffusion during the radiationinduced degradation was also brought out by the work of Burillo et al. [1992], who found that, because of the reduced diffusion of oxygen, the extent of oxidative reactions was lower in compressed PVC samples (P ≤ 880 MPa) than in uncompressed ones. Since oxygen is consumed during the radiationinduced oxidation, continuous oxidation during irradiation depends on the diffusion of oxygen into the sample. Thus one would expect that the extent of the degradation is inversely proportional to (i) the sample thickness, (ii) G-value for oxygen consumption, and (iii) the dose rate. These expectations are borne out by the work of several workers [Dole 1972, 1973; Feng and Ma, 1992; Kashiwabara and Seguchi, 1992; Williams, 1992]. In general, at the dose rates obtained from laboratory-scale gamma-irradiation equipment (mostly Co-60, at about 5 Gy·s-1), the oxidative degradation predominates. On the other hand, on irradiation with electron accelerators, where the dose rates are high (104 - 105 Gy·s-1 [Singh, 1991]), the dissolved oxygen is rapidly consumed, and any further oxidation is limited to the surface. However, in most cases where thick samples

Use of High-Energy Radiation in Polymer Blends Technology

were electron-irradiated, the effects of the surface oxidation were found to be negligible, as for example, during radiation-induced crosslinking of wire and cable insulation [Feng and Ma, 1992]. In the presence of additives, both crosslinking and scission are usually affected; e.g., antioxidants reduce both of these, whereas crosslinking agents predominantly enhance crosslinking [Makhlis, 1975; Sasaki et al., 1979; Charlesby, 1987; Carlsson et al., 1991; Feng and Ma, 1992; Kashiwabara and Seguchi, 1992; Lyons, 1992; Williams, 1992]. Lyons [1992] has pointed out that in the presence of antioxidants, chain scission, which is usually a unimolecular reaction, may be affected to a smaller extent than crosslinking, which is usually a bimolecular reaction. Most commercial polymers contain low amounts of antioxidants. Thus the results obtained with them may vary from batch to batch and supplier to supplier. Crosslinking converts linear molecules into three-dimensional structures with increased molecular weight, improvement of most mechanical properties, heat tolerance, and reduced solubility. On the other hand, degradation reduces the molecular weight, increases solubility, and lowers the mechanical performance. Usually, the degradation and crosslinking reactions occur simultaneously, though in some polymers crosslinking may dominate, e.g., in PE; but in others the degradation plays the main role, e.g., in PTFE [Charlesby, 1960; Chapiro, 1962; Wilson, 1974; Bradley, 1984; Woods and Pikaev, 1994]. In polymers having carbon atom main chain, the crosslinking was observed to decrease (while the degradation increased) when the hydrogen atoms on the carbon main chain were substituted with bulkier groups. Data on the efficiencies of crosslinking and degradation of various polymers have been reported [Woods and Pikaev, 1994]. Most polymers tend to decompose as they are heated [Grassie, 1972; Allcock and Lampe, 1981]. Many of the heat-induced reactions are due to the formation of carbon-centered free radicals, followed by peroxidation, such as reaction (R-22). Thus some of the radiolytic chemical reactions being described in this section may also take place during conventional processing

767

of polymers, particularly at elevated temperatures (>100°C for systems containing peroxides, and >150°C for polymers without any reactive ingredients), though the extent of such reactions may be low. 11.3.1

Polyethylene, PE

Irradiation of polyethylenes dominates the commercial activities in the polymer irradiation domain [Bradley, 1984; Charlesby, 1987; Feng and Ma, 1992; Silverman, 1992; Woods and Pikaev, 1994]. The main radiolytic chemical reactions in PE are (i) inter- and intra-chain crosslinking, viz. (R-9); (ii) formation of trans-vinylene groups (-HC=CH-), viz. (R-13); (iii) free-radical-addition reactions, viz. (R-30); (iv) formation of alkyl and allyl radicals.; (v) reduction of crystallinity (at doses of about 1 MGy); (vi) formation of H2 and low molecular weight hydrocarbons, viz. (R-5) and (R-8); and (vii) in the presence of oxygen, oxidation and chain scission, viz. (R-22), and (R-26 to R-29) [Crook and Lyons, 1963; Cross and Lyons, 1963; Lyons and Vaughan, 1967; Waterman and Dole, 1970; Dole, 1972, 1973; Silverman, 1981; Keller and Ungar, 1983; Lyons, 1983; Randall et al., 1983; Bradley, 1984; Clough et al., 1991; Lyons, 1992]. The crosslinking yield for the various PEs, G = 0.03-0.26 µmol.J-1, depends on the type of PE and the conditions of irradiation [Woods and Pikaev, 1994]. Radiation-induced crosslinking, and oxidative degradation of PE have been widely studied [Charlesby, 1960; Chapiro, 1962; Wilson, 1974; Keller and Ungar, 1983; Lyons, 1983; Randall et al., 1983; Bradley, 1984; Clough et al., 1991; Feng and Ma, 1992; Lyons, 1992; Silverman, 1992; Woods and Pikaev, 1994]. The changes observed in the physical properties of PE on irradiation include (i) decrease of elongation with increasing crosslinking, and on oxidative degradation; (ii) increase of modulus and its stability; (iii) increased tensile strength on crosslinking but its decrease on oxidative degradation; and (iv) decreasing solubility of PE on increasing crosslinking.

768

A. Singh and K. Bahari

One of the widely used effects of irradiation of PE is the so-called memory effect [Charlesby, 1960, 1987; Silverman, 1992]. PE is irradiated at RT in an inert atmosphere, and heated above the melting point, then stressed and given a particular shape. During the subsequent cooling, the crystallization fixes the imposed shape. However, upon reheating, the crystals melt and the crosslinked PE returns to its original unstressed shape. This memory effect is the basis of the production of many heat-shrink products, e.g., food-packaging film and electrical connectors [Bradley, 1984; Silverman, 1992]. It is widely accepted that free radical reactions, e.g., (R-37), which is an analogue of (R-9), contribute to PE crosslinking, though opinions on the details and the involved types of the free radicals may vary (for example crosslinking through the H-link involves combination of two sec-alkyl radicals and crosslinking through the Y-link involves addition of a free radical at a double bond in the neighboring molecule) [Charlesby, 1960; Chapiro, 1962; Dole, 1972, 1973; Lyons, 1983; Randall et al., 1983; Tabata et al., 1990; Singh, 1995]: PE• + PE• → PE-PE

(R-37)

where PE. represents the free radical formed by the loss of a hydrogen atom from PE. It is noteworthy that whereas in free radicals from small molecules both disproportionation and recombination reactions occur (reactions (R-12) and (R-13)), in the case of the free radicals from larger species (e.g., polymers), only recombination seems to take place [Silverman, personal communication, 1996]. On irradiation in air at low dose rates, the peroxy radical formation (reaction R-22) dominates over recombination (reaction R-37). However, with increasing dose rates, there would be increasing competition between these two reactions, and the reaction of PE. with the peroxy radical that may lead to PE• + PEO2• → PEOOPE

(R-38)

At the high dose rates typically obtained with electron accelerators [Singh, 1991], the recombi-

nation reaction, (R-37), dominates, mainly because the dissolved O2 gets rapidly consumed and the rate of the peroxy radical formation becomes dependent on the rather slow rate of diffusion of O2 into the PE [Dole, 1973; Seguchi and Yamamoto, 1986; Clough et al., 1991; Gillen and Clough, 1992]. Besides the expected crosslinking of PE on irradiation with electron accelerators, there are reports of crosslinking of PE at the low dose rates of gamma irradiation in air [Torikai et al., 1986; Minkova and Mihailov, 1990]. There is also a report that dicumyl peroxide can act as a sensitizer for radiation-induced crosslinking of PE [Kurian et al., 1993]. The implications of these observations and the mechanisms of crosslinking of PE have been discussed by Singh [1995]. The various radiation-induced free radicals formed in polyethylene, as shown by the reactions given in this Section, are also key species that can participate in the process of grafting of monomers to polyethylene [Chapiro, 1962; Dworjanyn and Garnett, 1992]. The products formed on radiolysis of PE include hydrogen, low molecular weight hydrocarbons, and unsaturation (C=C) [Chapiro, 1962; Dole, 1972, 1973]. The extent of radiationinduced crosslinking and the effects of irradiation on the mechanical properties of PE have been widely studied [Charlesby, 1960; Chapiro, 1962; Silverman, 1981; Lyons, 1983, 1992; Singh et al., 1995; also see the data given in Section 11.4]. 11.3.2

Polypropylene, PP

Commercial PP is predominantly isotactic and contains 50-60% crystallinity [Bradley, 1984]. PP normally has some unsaturation (e.g., vinylidene groups), a few hundred parts per million of combined oxygen (as carbonyl, hydroxy and hydroperoxy groups) and a catalyst residue (~100 ppm). It usually contains antioxidants, stabilizers, and other additives [Geymer, 1973]. The combination of high crystallinity, low density, high melting point, Tm = 166°C, and HDT > 90°C, make it ideal for molding plastic medical devices and disposables, which have to be autoclaved for sterilization [Bradley, 1984]. As the use of radiation sterilization increased, the effects of

Use of High-Energy Radiation in Polymer Blends Technology

high-energy radiation on PP became a major area of investigation [Williams, 1992, Kashiwabara and Seguchi, 1992, Morrissey and Phillips, 1993]. On irradiation, hydrogen and methane are amongst the products formed; the yields for PP are G(H2) ~ 0.28 and G(CH4) ~ 0.01 µmol.J-1 (for atactic PP these values are slightly higher) [Geymer, 1973]. The yields can be reduced by ca. 20% by the addition of nitrous oxide (electron scavenger) or nitric oxide (free radical scavenger) [Geymer, 1973]. Unsaturations (vinylidene and vinyl groups) as well as chain scissions, in roughly equal yields, result on irradiation of PP. These unsaturations would contribute to increased crosslinking by free radical and cationic reactions (e.g., R-11 and R-30; also see Singh [1995]). A comparison of the yields of scission vs. crosslinking for atactic and isotactic PP suggests that crosslinking affects mainly the amorphous regions [Geymer, 1973]. At high doses, the crystallinity of PP declines (viz. to 73% at 6 MGy and to 15% at 18 MGy) and crosslinking increases [Geymer, 1973]. The chain scission of PP is favored by higher irradiation temperatures [Bradley, 1984]. The intrinsic viscosity of PP decreases on irradiation under inert atmosphere or in air [Bradley, 1984]. Harnischfeger et al. [1990] found that the α-relaxation temperature of iPP decreases from 90°C to 84°C, on electron irradiation (at room temperature under vacuum) to a dose of ~200 kGy. This was attributed to the introduction of defects into the iPP crystals that, in turn, lower the overall melting point of the iPP. The reduction in the intrinsic viscosity is due to reductions in molecular weight brought about by chain scissions. On irradiation under inert atmosphere, the free radicals formed most likely undergo intramolecular scission [Lazár et al., 1989]: PP• → ppCH = CH2 + pp•

(R-39)

Here pp denotes fractional chains of polypropylene. Crosslinking of PP can also occur by reactions analogous to (R-30), and (R-37), and by the analogues of the cationic reactions in PE [Singh, 1995].

769

Irradiation of PP in air leads to oxidative degradation, evidenced by discoloration and embrittlement. The extent of the degradation depends on crystallinity, MW, MWD, and chain mobility [Kadir et al., 1989; Kashiwabara and Seguchi, 1992; Williams, 1992]. Neat PP does not discolor on irradiation up to 100 kGy [Williams, 1992]. The antioxidants should be selected so as not to cause the discoloration. However, most commercial preparations containing phenolic antioxidants turn yellow on irradiation. Phenolic antioxidants produce stable phenoxyl radicals that convert into colored quinonoids. Other stabilizers and antioxidants are compounds that contain either phosphorous [Bentrude, 1965; de Paolo and Smith, 1968], sulfur [Jirackova and Pospisil, 1979], or hindered piperidine derivatives [Carlsson, et al., 1980; Felder et al., 1980; Allen et al., 1981]. A comprehensive list of stabilizers and their mode of action was given by Dexter [1992]. It is noteworthy that antioxidants and stabilizers are excluded from the crystalline regions [Winslow et al., 1966]; thus they would provide protection only within the amorphous domains. In the presence of oxygen, oxidative degradation proceeds by chain scission, lowering the MW, and forming carbonyl, hydroxy and hydroperoxy groups [Geymer, 1973]. The oxygen uptake values are quite high: G(-O2) = 16.6 µmol.J-1 for PP powder and 3.8 µmol.J-1 for PP fibers [Kimura and Sakurada, 1967]. These values suggest that peroxidation proceeds via reactions (R-22), (R-23), and (R-26) to (R-29), similar to that in saturated hydrocarbons or PE [for detailed mechanisms, see Williams, 1992; and Kashiwabara and Seguchi, 1992]. The addition of agents that enhance mobility of the polymer chains (e.g., mineral oil) results in enhanced crosslinking and decreased oxidative degradation [Williams et al., 1982]. The protective effect of the mobilizing agents is greater if the PP has narrow MWD [Williams, 1992]. Grafting other polymers onto PP chain has been widely studied [Geymer, 1973]. The longlived free radicals formed when PP is irradiated in inert atmosphere, and the hydroperoxides formed when PP is irradiated in air, enable one to graft

770

A. Singh and K. Bahari

other monomers or polymers to pre-irradiated PP, e.g., glycidyl methacrylate [Iwakura et al., 1965], or acrylonitrile [Chapiro, 1962]. Irradiation of PP in inert atmosphere results in similar yields for chain scission and crosslinking [Geymer, 1973]. Because of scission, the intrinsic viscosity of PP initially decreases. At higher doses, >100 kGy, gelation caused by crosslinking has been observed [Black and Lyons, 1959; Kondo and Dole, 1966]. Two factors contribute to this effect: (i) the presence of trace amounts of hydroperoxides in the PP, which get consumed in various radiolytic reactions, e.g.: PPOOH + e- → PPO• + OH-

(R-40)

and (ii) the participation of unsaturations in the crosslinking reactions (through analogues of (R-30) and cationic reactions [Singh, 1995]). Since unsaturation increases as the radiolysis progresses, crosslinking dominates at higher doses. Trapped free radicals may also contribute to PP crosslinking. For example, when PP is irradiated at liquid nitrogen temperature (T = -196°C), the free radicals are trapped in the polymer; on warming to RT, crosslinks are formed. However, if the irradiated sample is treated with methyl mercaptan (a free radical scavenger) while it is still cold, the free radicals react with it and no crosslinking takes place: PP• + CH3SH → PP + CH3S•

(R-41)

Many other additives have also been found to act as crosslinking inhibitors, including iodine, 2-propanol, dilauryl thiodipropionate and various phenols [Geymer, 1973]. By contrast, other additives increase the crosslinking yield by modifying the free radical reactions. Addition of nitrous oxide or chlorobenzene increases the crosslinking yield by electron scavenging. Polyfunctional monomers, such as allyl acrylate and allyl methacrylate, can bring about crosslinking between somewhat distant chains, or reaction sites. They also increase the yield of crosslinks, through additional reactions [Geymer, 1973] (e.g., analogues of reaction (R-30)).

In general, the mechanical properties of PP degrade on irradiation, more so when irradiated in air [Bradley, 1984; Kashiwabara and Seguchi, 1992; Williams, 1992; also see data in Section 11.4]. 11.3.3

Polyvinylchloride, PVC

PVC, a polymer that is widely used because of its low gas permeability and high fire retardancy, is sensitive to light and heat. The degradation leads to dehydrochlorination and color formation [Chapiro, 1962; Bradley, 1984; Thomas et al., 1986]. To process the resin it is necessary to add heat stabilizers. The most common additives include metal organic compounds, acid acceptors and alkalis, dioctyl phthalate, and dioctyl sebacate [Bradley, 1984]. PVC is only slightly crystalline. However, it has been used to produce heat-shrink products [Bradley, 1984]. Dehydrochlorination, brought about by free radical intermediates, is characteristic of both thermal and radiolytic decomposition of PVC [Salovey, 1973]. HCl is the major gaseous product of irradiation. Its formation depends on temperature; at T = -196°C, H2 is the major product but HCl is also formed [Ohnishi et al., 1962]. As the temperature increases, the yield of HCl increases, first linearly up to T ~ -60°C, and then at higher rates from T ~ -40°C to ~ +30°C. At T = 30 to 70°C the rate of HCl formation again increases [Ohnishi et al., 1962]. PVC does undergo similar decomposition when heated to T > 150°C. However, irradiation seems to sensitize it to thermal degradation at much lower temperatures. The G-values (in µmol.J-1) of HCl formation at T = -145 to -90°C are G = 5.6; at T = 30°C, G = 1.3; and at 70°C, G = 2.4 [Salovey, 1973]. The color formation is attributed to the formation of polyene chains with long sequences of conjugated double bonds [Ohnishi et al., 1962; Salovey, 1973]. The free radical origin of the color formation is confirmed by its inhibition by oxygen. Lindberg et al. [1989] have compared the degradation of PVC with that of PVDC, Cl-PE, Hypalon and chloroendic ester polyester, on electron irradiation. The mass loss in the case of PVC appears to be the highest.

Use of High-Energy Radiation in Polymer Blends Technology

Hegazy et al. [1981] reported the formation of H2, CH4, CO and CO2 on irradiation of PVC in air. Thomas et al. [1986] studied the gammairradiation effects in the presence of air on PVC (suspension polymerized containing 60 phr of DOP, 4 phr barium cadmium stabilizer, and 2 phr barium stearate lubricant) and its blends with a thermoplastic, elastomeric polyether-ester block copolymer, TPE. PVC samples turned pale yellow at a dose of 10 kGy, reddish brown at 100 kGy and dark brown at 500 kGy. The authors also studied changes in the mechanical properties. The tensile strength of PVC decreased marginally at doses up to 100 kGy, with larger decreases at higher doses, up to 500 kGy. The modulus of PVC decreased up to a dose of 100 kGy after which it was unchanged up to 500 kGy. The energy at rupture and elongation at break for PVC remained steady up to 10 kGy, but decreased at higher doses. SEM data showed closely packed well-defined surface patterns. The corrugations seen on the surface were associated with the degradation of PVC. The tensile fracture surface of non-irradiated PVC showed discontinuous fracture paths, attributed to the shearing action between the amorphous and semi-crystalline chains. The surface of the irradiated PVC showed brittle failure. These results are consistent with the more recent observations by Kashiwabara and Seguchi [1992]. Electron spin resonance, ESR, measurements of the irradiated samples at T = -196°C showed the formation of alkyl radicals, resulting from carbon-chlorine and carbon-hydrogen bond scission [Ohnishi et al., 1962; Salovey, 1973; Burillo et al., 1992]. On warming to 25°C in vacuum, the ESR spectra changed, suggesting the formation of polyenyl radicals. Ultra-violet, UV, spectroscopy shows the formation of allyl and polyenyl radicals. In the presence of oxygen, the allyl radicals decay rapidly. Detailed reaction mechanism has been advanced by Ohnishi et al. [1962]. In air, the alkyl and the allyl radicals formed on irradiation of PVC most likely react with O2. Although O2 inhibits dehydrochlorination, the formation of the polyenyl radicals, and color development, it does promote oxidative degradation through reactions

771

(R-22 to R-23). Sodium stearate appears to be quite efficient in preventing discoloration of airirradiated PVC whereas the samples containing organo-tin stabilizers turn brown at rather low doses [Chapiro, 1962]. Naimian et al. [1994] report the formation of hydroxyl and carbonyl groups on gamma irradiation of plasticized PVC in air. IR spectrometry and rheological testing of the PVC samples showed that chain scission was not significant at the sterilizing doses used (25 and 50 kGy). Post-irradiation degradation of the mechanical properties of the PVC samples (tensile strength and elongation) was also insignificant, on storage for 6 months. Crosslinking of PVC is known to occur on irradiation, with a low yield [Ohnishi et al., 1962]. More efficient crosslinking may be induced by irradiating PVC in the presence of multifunctional crosslinking agents [Salovey, 1973]. For example, plasticized PVC has been used as a wire and cable insulation. To enhance the crosslinking of PVC during electron irradiation, triallyl cyanurate (TAC), trimethylolpropane trimethacrylate (TMPTMA), or trimethylolpropane triacrylate (TMPTA) have been used as the additives [Bradley, 1984]. 11.3.4

Polystyrene, PS

PS is one of the most stable polymers that requires large doses of irradiation for any effects to be observed [Chapiro, 1962; Parkinson and Keyser, 1973; Shultz and Mankin, 1976; Bradley, 1984]. The G-values for crosslinking are somewhat higher than those for the formation of H2, but both are low, G < 0.004 µmol.J-1. For PS chain scission under vacuum, G ~ 0.0016 µmol.J-1 [Parkinson and Keyser, 1973]. When PS was irradiated at 25°C under vacuum and then heated to 95°C, additional 50% of crosslinks were formed. This was attributed to the presence of trapped free radicals. Irradiation of PS in air resulted in slow increase of -OH and >C=O groups. Most of the work on irradiation has been done using atactic PS, with limited studies on isotactic PS [Parkinson and Keyser, 1973].

772

A. Singh and K. Bahari

ESR studies have established the formation of two different free radicals on irradiation of PS: a di-substituted benzyl radical, formed by the loss of a hydrogen atom from the α-carbon atom, and a cyclohexadienyl type of radical formed by the addition of a hydrogen atom to the benzene ring. In addition to the reactions of these two free radicals, the ionic species formed during the irradiation may also take part in the radiationinduced crosslinking of PS [Heusinger and Rosenberg, 1969]. Radiation-induced crosslinking of PS is accelerated by the presence of nitrous oxide [Amemiya et al., 1963], which supports the participation of cationic reactions (e.g., reaction R-33) in its crosslinking. The incorporation of multifunctional crosslinking agents (e.g., allyl methacrylate, allyl maleate and TAC) markedly increases the crosslinking yields [Chapiro, 1962; Parkinson and Keyser, 1973; Bradley, 1984]. Radiation-induced free radical formation in PS is known to facilitate grafting reactions [Chapiro, 1962; Dworjanyn and Garnett, 1992]. The effect of temperature on the crosslinking was also reported [Parkinson and Keyser, 1973]. The G-value for crosslinking remains essentially constant at G ~ 0.005 µmol.J-1 at T = -196 to 65°C, and then it decreases to zero at T ~ 100°C. At T = 130°C, only scission, but no crosslinking was observed. In contrast, the G-value for hydrogen formation increased linearly from G = 0.0024 µmol.J-1 at T = -196°C to G = 0.0042 µmol.J-1 at T = 65°C; at T = 130°C it decreased to G = 0.0032 µmol.J-1 [Burlant et al., 1962]. The mechanical properties of PS are affected to a much smaller extent, than those of PE or PP, on irradiation, though the elongation-atbreak appears to be more sensitive than the other properties; it decreases by about 10% at a dose of 50 MGy [Parkinson and Keyser, 1973]. 11.3.5

Polymethylacrylate, PMMA

PMMA is a widely used transparent polymer. A noticeable yellowing develops in PMMA at about 30 kGy, which at higher doses deepens to a reddish brown color. The color is due to trapped free

radicals or charged species — it slowly fades away as the oxygen diffuses into the irradiated sample [Chapiro, 1962]. The color also disappears on heating the irradiated sample to T ~ 80°C. Irradiation of isotactic PMMA leads to its racemization, with the formation of syndiotactic and heterotactic sequences, attributed to a recombination of chain scission fragments [Dole, 1973a]. Irradiation of PMMA in vacuum results in the formation of several gaseous products including CO, CO2, CH4, H2, and methyl formate, a similar effect to that seen earlier for methyl acetate [Dole, 1973a]. The combined G-value (in µmol.J-1) for the gaseous products was G = 0.16 at a dose of 560 kGy, decreasing to G = 0.11 at 2800 kGy [Chapiro, 1962]. Crosslinking of PMMA has not been observed [Chapiro, 1962; Dole, 1973a]. The G-values for the main chain breaks and the loss of the -COOCH3 groups are G ~ 0.12 and 0.08 µmol.J-1, respectively [Dole, 1973a]. Irradiation of PMMA leads to the formation of carbon-centered free radicals [Dole, 1973a]. In PMMA polymerized by irradiation, ESR identified the free radical on the carbon between the two -CH2- groups as the precursor of the one with the free radical site on the carbon next to the ester group. Another radical, which could result from main chain scission, has never been seen. It probably undergoes recombination fairly quickly, but lives long enough to lead to racemization [Dole, 1973a]. For the radical(s) formed during gamma irradiation of PMMA under vacuum, G = 0.25 µmol.J-1 [Kircher et al., 1965]. For non-irradiated PMMA the elongation at break at 325°C was ~ 300%. On irradiation, the temperature at which this value of elongation at break was obtained decreased with increasing dose: for 31 kGy, it was 210°C [Chapiro, 1962]. This decrease in temperature was greater for low dose rates than for the high ones; thus significant differences in the values for elongation at break for electron- and gamma-irradiated PMMA were found. Contrary to the case of most polymers, the radiation-induced degradation of PMMA is inhibited by O2, though the mechanism of this inhibition has not been established [Chapiro, 1962].

Use of High-Energy Radiation in Polymer Blends Technology

Peroxy radicals are known to be formed on exposure of irradiated PMMA to O2. Thus in the presence of O2 some oxidative degradation would take place but its nature is most likely different from that of the degradation in the absence of O2. The addition of benzene and many free radical scavengers (e.g., allyl thiourea) also protects PMMA against radiation-induced degradation. PMMA undergoes grafting by direct irradiation, post-irradiation by trapped free radicals or on irradiation in a blend [Chapiro, 1962]. 11.3.6

Elastromers

Ethylene-propylene copolymers, EPR, and terpolymers containing diene monomers, EPDM, crosslink on irradiation [Bradley, 1984; Aoshima et al., 1992]. The crosslinking can be enhanced by the use of crosslinking agents; for example, EPDM containing 5-ethylidene-2-norborene, ENB, may be efficiently radiation crosslinked. Analogues of reaction (R-30) contribute to such crosslinking. 11.3.6.1

Ethylene Propylene Rubber, EPR

The radiation-induced crosslinking efficiency of EPR is higher than that of PP, but lower than that of PE. The crosslinking is not an additive function of the PE and PP content [Geymer, 1973]. The products obtained on irradiation of EPR are similar to those obtained from PE and PP, and the effects of the free radical scavengers and crosslinking agents are also similar [Geymer, 1973]. Chen and Lee’s [1986] results show that on gamma irradiation under nitrogen, the gel fraction levels off at about 80 kGy (gel fraction ~20% at 50 kGy, ~80% at 80 kGy, and ~86% at 150 kGy). Eldred [1974] electron-irradiated EPR and EPDM in air, in the presence or absence of the following crosslinking agents: ethylene glycol dimethacrylate, EDMA, 1,4-hexadiene, HD, 1,2-poly(butadiene), TAC, triallyl isocyanurate, TAIC, triallyl phosphate, TAP, ENB, and TMPTMA. In the presence of unsaturated crosslinking agents, the dose required to cure these rubbers

773

was drastically reduced. For example, typically EPDM requires over 1500 kGy dose for radiation curing. On incorporation of 15 phr of TMPTMA the curing dose was reduced by 90%. EDMA was found to be more efficient in reducing the curing dose than the allylic compounds, and ENB was more efficient than HD. Eldred [1974] reported that the physical properties of the radiation-cured EPDM compared favorably with those of the conventional sulfur-cured EPDM. Moreover, it showed much better compression set and oil resistance. Kammel and Wiedenmann [1976] have also reported on crosslinking of EPDM, on electron irradiation in air. 11.3.6.2

Thermoplastic Elastomer, TPE

Thomas and co-workers [1986] studied effects of gamma-irradiation on a thermoplastic elastomer, TPE, a polyether-ester block copolymer from 1,4-butanediol, polybutylene glycol and terephthalic acid (Hytrel™ D40), TPE, and its blends with PVC, in air, at doses of 1, 10, 100 and 500 kGy. Visual inspection of the samples showed no color change up to 500 kGy, but the samples irradiated at 500 kGy showed surface cracks. TPE crosslinked on irradiation. Its tensile strength increased by about 30% at 10 kGy, but at 500 kGy it dropped to about one third of the value for the non-irradiated sample. The modulus was not much affected by irradiation. The elongation at break, tear strength, tensile set after failure, and energy at rupture were similarly affected: up to 10 kGy they were unchanged, but at higher doses their values drastically declined, most likely because of radiation-induced oxidation. Surface of TPE samples irradiated to 500 kGy showed two types of cracks — a number of broad cracks interconnected by a network of microcracks. Thomas et al. [1986] attributed these to embrittlement of the bulk as a result of crosslinking, and to radiation-induced surface oxidation. The tear failure surface of non-irradiated TPE showed a rough surface with sinusoidal foldings and closely spaced vertical striations. The development of these patterns was attributed

774

A. Singh and K. Bahari

to the presence of crystalline segments in TPE, which arrested the propagation of the growing cracks. In contrast, samples irradiated at low doses showed a smooth tear failure surface, with a few tear lines. This was attributed to the loss of crystallinity, leading to loss of resistance to the crack propagation. Some TPE samples have also been irradiated with electrons from a 1.5-MeV accelerator [Du Pont, 1995]. As shown in Table 11.8, on irradiation to 150 kGy at 23°C in air, the samples showed excellent retention of physical properties.

11.4

Irradiation of Polymer Blends

Irradiation of polymer blends has been the focus for several teams. The polymers most often used

Table 11.8.

Non-irradiated Tensile strength (MPa) Elongation at break (%) 100% Modulus (MPa) Hardness, Durometer D Irradiatedb, 50 kGy Tensile strength (MPa) Elongation at break (%) 100% Modulus (MPa) Hardness, Durometer D Irradiatedb, 100 kGy Tensile strength (MPa) Elongation at break (%) 100% Modulus (MPa) Hardness, Durometer D Irradiatedb, 150 kGy Tensile strength (MPa) Elongation at break (%) 100% Modulus (MPa) Hardness, Durometer D

b

11.4.1 11.4.1.1

Polyolefins, PO Normal and Ultra-High Molecular Weight Polyethylenes, PE/UHMWPE

Minkova [1987], and Minkova and Mihailov [1990] studied gamma-irradiation of blends of these two polymers, in air at RT, to doses up to 60 kGy. The blends of the two were prepared by passing their mixtures through a melt-flow indexer at 190°C, four times. The pure components were similarly treated. The melting temperatures for the first and

Effect of electron irradiation of Hytrel™ polyester elastomers, TPEa ASTM Method

a

in the blends are PE, PP, EPDM, and PS. Radiation effects on immiscible blends have been recently reviewed [van Gisbergen and Overbergh, 1992]. A list of the systems discussed in this chapter is given in Table 11.9.

Hytrel 4056

Hytrel 5556

Hytrel 7246

D 638 D 638 D 638 D2240

24.1 550 6.8 40

27.2 390 14.4 55

35.7 430 22.0 72

D 638 D 638 D 638 D2240

22.8 510 7.3 40

28.3 470 14.5 55

36.6 410 23.6 72

D 638 D 638 D 638 D 638

22.8 500 6.2 40

28.9 470 14.5 55

37.4 370 23.9 72

D 638 D 638 D 638 D2240

22.1 490 6.1 40

30.3 490 14.2 55

38.6 390 24.6 72

Data from Du Pont [1995]. Irradiation with 1.5-MeV electrons at 23°C, 70% relative humidity.

Use of High-Energy Radiation in Polymer Blends Technology

Table 11.9.

775

Polymer blends: summary of irradiations

Blend

Sample size (mm)

T (irrad.)

Cover gas

PE/UHMWPEa

Source dose, kGy

Additives

6 x 4a

RT

air

γ, 60

-

HDPE/LLDPEb

-

RT

air

γ, 100

DCP

LDPE/PPc

-

RT

N2

e-, 100

P-1c

LDPE/PPd

5

RT

vac

γ, 300-1500

-

LDPE/PPd LDPE/PPe

0.5 0.5 x 4 x 5

RT RT

vac air

γ, 1500 γ, 100

-

EVAc/PPf LDPE/PPf

100-µm film

RT?

air

e-, 50-160

f

PP/cPPg

0.5

RT

air

e-, γ

g

LDPE/PAh LDPE/PSh LDPE/PA-6i

-

-

-

γ, 50

-

10 x 100 x 1

RT

air

γ, 100

-

LLDPE/PA-6i

10 x 100 x 1

RT?

vac, air

γ, 50-400

-

PE/PA-6i

h

RT

air

γ, 0-50

-

PE/EVAcj PE/CPEj

-

RT

air

e-, 100

j

LDPE/EVAck

2

RT

Argon

γ, 100-160

k

LDPE/EVAcl

-

RT?

air, N2, O2

γ, e-, 35

-

Properties

Ref.

Enthalpies and rates of melting and crystallization Crosslinking, crystallinity, tensile and yield strength, abrasion, modulus, DSC measurements Viscosity, flow behavior, chemical and radiation crosslinking Crosslinking, scission, crystallization, thermal analysis, density Elongation, modulus, crystallinity Elongation, tensile modulus and strength, crystallinity Melting temperatures, enthalpy of melting, Elastic modulus, Tensile strength, Elongation at break Oxidative degradation, elongation at break Gel fraction

1, 2 3

4 5 6 7

8 9 10

Gel fraction, carbonyl formation 11-14 torque, viscosity, tensile strength and modulus, elongation impact strength, SEM, DSC Gel fraction, melt viscosity, tensile 12, 15, strength and modulus, elongation, 16 impact strength, carbonyl and hydroxy group formation, SEM, DSC Viscosity, SEM, IR, temperature and 17 heat of crystallization Hardness, tensile strength, elongation, 18, 19 heat aging, chemical resistance, dielectric properties, resistivity, dissipation factor, softening point, contraction Gel fraction, volatile products, tensile 20 strength, elongation at break, Shore ED hardness, Young’s modulus, aging, crystallinity Gel fraction, carbonyl groups, 21 density, foam characteristics, tensile strength, elongation at break.

776

Table 11.9.

A. Singh and K. Bahari

Continued.

Blend

Sample Size (mm)

T (irrad.)

Cover gas

Source dose, kGy

Additives

1.9

RT

air

e-, 300

m

Crosslinking

22

46 x 4 x 0.3

RT?

N2

γ, 50-150

-

23

EPDM/PE°

20 x 5 tape

RT?

air?

e-, 0-150

o

LDPE/EPDMp

1.6; 100µm

RT?

air

e-, 0-200

p

EPDM/LDPEq

2

RT

air

γ, 500

-

LDPE/EPDMr

1 x 4 x 20

RT?

air?

γ

r

LDPE/EPDM/CBs 1 LDPE/EEAs PP/EPDMt 240 x 240 x 1

RT

air?

s

RT

air

γ, e-(1MeV) e , 100

PP/EPDMu PP-PE/EPDMu PBT/EPRv

52u

RT

vac

Storage and loss modulus, tan δ, Tg, fracture surfaces Heat aging, crosslinking, oxidative degradation, Gel fraction, tensile strength, elongation, thermo-oxidative stability Tensile strength, elongation, modulus, permanent set Fracture stress, yield stress, elastic modulus Positive temperature coefficient, crystallinity Viscosity, DSC, rheology, impact strength, SEM, field strength Gel dose, tensile strength, elongation

-

RT?

air?, N2

EPDM/BRw PE/CDBx

-

RT? RT?

air? air?

e-, 0-160 e-, 0-150

x

PB/CDBy

150 x 150 x 1.5 RT?

air?

e-, 0-200

PP/CDBy

-

RT

air?

e-, 0-40

Irganox 1010 -

PP/PBz

-

RT?

air?

e-, γ, 100

z

PE/Craa

3-3.5

RT

air

e-, 150

aa

HDPE/NRab

2

RT

air, N2

γ, 1000

DCP

PE/CSRac

-

20°C

N2

EPR/LDPEm EPDM/LDPE EPDM/PPm EPR/PPn

Properties

Ref.

m

-

e-, 1 MeV, 400 γ, 0-150 DCP, GF

e-, 50-200

TGM-3ac

19, 24 25

26 27 28, 29 30, 31 32

Impact strength, tensile strength, 33, 34 modulus, delamination Energy at break, morphology 35 Crosslinking, heat sealability, 36 heat-aging Modulus, elongation, tensile strength, 37 melt index Modulus, elongation, tensile strength, 37 melt index Radiation resistance, tensile strength, 38 elongation, heat seal, tear, puncture resistance, grafting Tensile strength, true stress, 39 crosslinking, crosslinking promoters, fillers Crosslinking, crystallinity, tensile 40 strength, stress, elongation, SEM Tensile strength, elongation at break, 41 equilibrium modulus, rigidity, density, gel fraction, degree of swelling, flexing resistance, aging

Table 11.9.

Use of High-Energy Radiation in Polymer Blends Technology

777

T (irrad.)

Cover gas

Source dose, kGy

Additives

Ref.

Continued.

Blend

Sample Size (mm)

PVC/Hytrelad

-

RT

air

γ, 0-500

-

TPU/POMae

-

RT?

air

γ, 25

-

PP/SKTV-1af PP/SKEP-60af EVAc/PPag

0.1

RT

air, H2O

γ, 50-1000

-

2

RT

air

γ, 0-1000

DCP

NR/EVAcah cPI/tPIai BR/EPDMaj PS/PMMAak

50 x 6.5 x 3

RT? RT? RT? RT?

air? air? air? vac

γ, 0-500 γ e-, 0-140 γ

ah ai aj -

PS/PMMAal

-

RT?

N2

e-, 0-1660

-

PS/PMMAam PS/PMMAan PS/EPDMa°

80 x 80 x 4

RT RT? RT

air air? air

γ e-, γ e-, 0-200

SP, SEP

PS/HDPEap PMMA/SANaq

60 x 10 x 3

RT? RT?

air vac

e-, γ, 250 γ

-

PS/LDPEar

-

RT

air

e-, 200

-

PS/BRas PS/EPDMat

at

RT? 50-55

vac air

PS/PVMEau PSD/PVMEav

75 µm 0.1, 0.5

RT 45°C

N2 vac

60 61

LCP/PVDFaw LCP/HDPEaw PSF/PVPax

-

RT

air?

γ, 1-4 e-, 3 MeV, Kratonat 0-75 Cariflexat e ,175 keV Gel fraction, Tg γ, 0-1250 Spinodal temperature, phase diagram, small angle neutron scattering, crosslinking e-, γ, 500 aw Reduced shrinkage, linear expansion

-

RT?

air?

γ, 10-50

63

-

Properties

SEM, surface cracks, tensile strength, modulus, elongation, tear strength, tensile set Heat deflection temperature, tensile strength, yield stress, elongation, modulus Electrical strength, crystallinity, spherulite formation Color, rupture, brittle behavior, elongation, tensile strength, crosslinking Modulus, tensile strength Crosslinking, X-ray diffraction Domain size, ozone resistance Modulus, elongation, flexural strength, ESR Crosslinking and scission, molecular weights Viscosity, oxidation Microhardness Impact strength, SEM, compatibilization, grafting, Tg, IR, NMR Fixation, crosslinking Flexural strength, crosslinking, degradation, grafting Crosslinking, microrheology, thread break-up and coalescence, modulus, elasticity Gel fraction Impact strength, delamination

Chromatography with hollow fibers

42

43

44 8, 45

46 47 48 49 50 51 52 53, 54

55 56 57

58 59

62

778

Table 11.9.

A. Singh and K. Bahari

Continued.

Blend

Sample size (mm)

PP/EVAcay

T (irrad.)

Cover gas

Source dose, kGy

Additives

Properties

Ref.

-

RT?

N2

e-, 50

-

Peel adhesion, Young’s modulus

64

-

RT?

air?

γ, 0-50

-

Yellowness index, elongation, tensile strength, flexural modulus, Izod impact strength

65

-

RT?

N2

e-, 30-100

Heat shrinkage, haze, shrinkage temperature, shrinkage tension, tensile strength, elongation, impact strength Mass loss in electron microscope

66

ay

PP/PB/EVAc PP/PB/EVAcay PP/EAAay PP/PB/EAAay CO1/CO2ay CO1/PPay ZYLARaz St-94-560 to 562 ST-94-568 Filmaaa

Peroxide

PVC/SANaab 200 nm PVC/PMMAaab PVC/PCLaab POM/PVCaab POM/PMMAaab ETFE/TPEaac 0.64

RT

vac

e-

RT?

air?

aac, 0-100

EPDM/PAcaad PC/Polyesteraae

76 x 32 x 64

RT? RT?

air? air?

γ γ, 30

-

PMMA/COaaf PMMA/PVDFaag PMMA/PCTFEaag PVC/PMMAaah PVC/PCLaah PVC/SANaah HDPE/PP/PSaai

0.8-1.3 µm RT? 10 x 10 x 0.5 RT? 10 x 10 RT? 100 to 200 nm RT

vac air? air? vac

X-ray γ, 10-1000 e-, 10-1000 e-

-

4 x 10 x 80

air

γ, 0-600

RT

-

aac

TAC

67

Tensile strength, elongation, secant modulus, solvent resistance, heat aging, cold impact brittleness temperature, blocking, volume resistivity, flex durability Electrical conductivity, crosslinking Yellowness index, chemical resistance, mold release energy Degradation Hardness Hardness Chlorine and mass loss

68

71 72 73 74

Flexural strength

75

69 70

Use of High-Energy Radiation in Polymer Blends Technology

Table 11.9.

779

Continued.

air? Presumably in air RT? Presumably at room temperature a

b

c

d e

f

g h i

j

k

l

m n

° p

q r s

t

u

v

UHMWPE, melt index 0g/10 min (190°C, 21.6 kg), Mw = 4 x 106; PE, melt index 4.8g/10 min (190°C, 5 kg). Cylindrical sample, 6 mm diameter. HDPE, GA 7260 from PIL, Bombay (density, 0.957 g/mL, MFI = 5.2 g/10 min); LLDPE, Ladene 218 W from IPCL, Baroda (density = 0.92 g/mL, MFI = 2 g/10 min). Unstabilized PP (Profax 6501, Himont) and LDPE (DYNH-1, Union Carbide). P-1: 2,5-dimethyl-2,5-bis-(t-butylperoxy)hexane (Lupersol 101, Atochem); phenolic antioxidant (Cyanox 2246, American Cyanamid). LDPE, FERTENE AF5 and iPP, MOPLEN D60P, both from Montedison. LDPE, Petrothene (mol. wt. = 44,000; density = 0.916 g/mL); iPP, Anic Kastilene M261 (mol. wt. = 136,000; density = 0.908 g/mL). EVAc, Lupolen V5510SX from BASF (melt index = 18.4g/10 min; density = 954 kg/m3 at 23°C); LDPE, Ropoten FV 03223, made in Bulgaria (melt index = 0.907 g/10 min., density = 920 kg/m3); iPP Buplen 7623, made in Bulgaria (melt index = 4g/10 min; density = 903 kg/m3). All samples contained 0.15% of an unspecified stabilizer. PP and cPP from Toyoshoda, Japan; 0.2% 1,3,2,4-di(p-ethylbenzylidine) sorbital as a nucleating agent. No details given by the authors. LDPE, Riblene A22-SAK (Enichem); PA-6 (ADS 40, SNIA); LLDPE Riblene Lx BF 2211 (Enichem). LDPE, BB2700 and HDPE, ZB5015 all from Enichem; PA6, ADS40 from SNIA; sample size most likely similar to that used in Spadaro et al. [1992]. HDPE, PEI12-1 MI:2, from Qian Jin Chemical Plant, Beijing; EVAc, from Shanghai Research Institute of Chemical Industry; CPE 035, from Anhui Province Wuhu Chemical Plant; zinc stearate, from Jilin City Plant; and antioxidant 1010, from Beijing Chemical Industry Plant no. 3; details on the materials used by Brooks [1983] are not available. LDPE (36 wt%) melt index = 1.8g/10 min, mp = 1200C (PX 20020X, Petróleos Mexicanos); EVAc (35 wt%) melt index = 2.0g/10 min, mp = 1060C (vinyl content 18 %, DuPont); TMPTMA, 4.5 wt% (Sartomer); flame retardant, 24.25% (Sb2O3 or decabromodiphenyloxide, Great Lakes Chemical Co); and antioxidant, 0.25 wt% (pentaerythrityl-tetrakis3-(3,5-di-tert.butyl-4hydroxyphenyl)-propionate (PTTBHP) or octadecyl3-(3,5-di-tert.butyl-4-hydroxyphenyl)-propionate (OTBHP)). LPDE (MFI = 2 g/10min, from No.1 Factory of Yanshan Petrochem. Corp.), EVAc (MFI= 2g/10 min, VAC= 18%, from Beijing Organic Chem. Factory), Blends: LDPE/EVAc with 90-50 wt%, LDPE. Blending rolling at 114-120°C. Electron dose = 360 KGy/4, γ-dose = 16.2. Sheet expansion at T = 200°C. Vistalon 404 and 707 (EPR) and 2504, 3708, 4608 and 6505 (EPDM); TMPTMA; carbon black (TEDX-15). EPR (EPCAR 306, B.F. Goodrich); PP (PPF-101, Taiwan Propylene Corp.). Sources of materials were not given (see information given in Table 12.30). DCP and an antioxidant (polymerized dihydroxytrimethyl quinoline) added in some of the samples. Details were also not given by Brooks [1983]. LDPE, density = 0.918 g/mL, Mw ~ 75,000; EPDM, density = 0.86 g/mL containing 45% propylene and 8% ENB; for thermo-oxidative work, 100-mm-thick film contained 0.5% polymeric 2,2,4-trimethyl-1,2-dihydroquinoline as a stabilizer. Vistalon 2555 from Exxon, Belgium, and LDPE from Dow, Spain. LDPE/EPDM, 1/1, containing 30 phr HAF carbon black. LDPE (melt index = 2.0), EPDM, EEA (ethylene-ethyl acrylate, melt index = 6.0) and CSF carbon black (CB; average size ca. 70 nm, surface area = 230 m2/g, dibutylphthalate value 280mL/100 g, pH = 7-9). PP112, PP 13E, and Keltan 514 and 578 (DMC); PP, 13E10 from DSM; EP-H, Keltan 512 from DSM (ethylene content 55%); EP-L, Keltan 740 from DSM (ethylene content 60%); and EP Vistalon 808 from Exxon (ethylene content 77%). iPP, Vestolen P 7000, Hüls AG; EPDM (70 wt% PE, 30 wt% PP, < 1 wt% ENB), Brina AP 437, Hüls AG; iPP-PE (6% PE), Vestolen P 8502, Hüls AG; 52 test bars, according to DIN 53455. PBT, polymerization in laboratory (DMT from China Petrochemical Co., and BDO from Merck), and (Celanex 2001, Celanese Co.); EPR (EPCAR 306, Mooney viscosity = 30, B.F. Goodrich); DCP (Merck) detailed experimental conditions were not

780

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aa ab

ac

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af

ag

ah

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°

ap

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au av

aw

A. Singh and K. Bahari

given; some PBT samples were irradiated under nitrogen; GF, glass fiber. EPDM, Mn = 106,000; Mw = 329,000; Tg = -60°C; BR; Mn = 97,000; Mw = 165,000; vinyl content = 98%; Tg = -5°C. CDB, Exxon; Butyl 268; HDPE, density ~ 0.96; Irganox 1010, Ciba-Geigy; TMPTA, Microcel E (diatomaceous earth silica filler). PB, polybutylene, Shell grade 1600A, 20-70%; CDB, Exxon, 30-80%, MW = 350-400; Irganox 1010; carbon black; iPP, Norchem. PP (Cosden 8771, melt index = 9); PB (Shell 0400, melt index 20); Exxelor or PA-23 (Exxon); several vinyl monomers also used for grafting. Crystalline polyethyelene (Marlex EHM 6006) and Neoprene W; crosslinking agents PEGDMA, TAC, TAP and TMPTMA. NR, crumb rubber from The Rubber Institute of India, Kottayam; PE (HDPE), Hostalen GA 7260 from Polyolefin Industries, Bombay. PE, grade P-2020T; no other details were given for any of the polymers used; six elastomers were used to make blends with PE (see Section 11.4.2.3.3); TGM-3, an unsaturated oligesteracrylate. HytrelTM, Du Pont; source of PVC was not given. TPU, Thermoplastic polyurethane, Estane 58111, Goodrich; POM polyoxymethylene, Celcon M 140, Hoechst Celanese. Test specimens were prepared according to ASTM D638 and D648. iPP, grade 04P000-XI; Polymethylvinylsiloxane, SKTV-1; ethylene-propylene elastomer, SKEP-60; some of the samples were irradiated under liquid water. EVAc (Bayer, vinyl content = 45%; density = 0.975 g/mL, Mooney viscosity (ML(4)100°C, 30); iPP (Indian Petrochemical Corporation; density = 0.910 g/mL; Mw = 530,000). NR, ISNR-5; EVAc, Exxon 218 (vinyl content = 18%; melt-flow index = 1.7 g/10 min.; Vicat softening point = 64°C); additives include zinc oxide, stearic acid, styrenated phenol, dibenzothiazyl disulfide, DCP and sulfur. Guayule rubber (cPI); tPI (Trans-PIP from Polysar); experimental details were not given. BR (Diene 35 from Firestone Tire and Rubber Company); EPDM (Royalene 301 from Uniroyal); DBPC (1%); see the text of the patent for another example, names of other elastomers and details about the use of conventional additives. PS, Shell, high purity; PMMA, prepared in laboratory. PS from Pressure Chemical Co. (5a, M = 498,000; Mw/Mn = 1.20; 2b, M = 20,400; Mw/Mn = 1.06); PMMA, L1489 (Mn = 42,000; Mw = 79,000; Mz = 134,000) and 1490 (Mn = 2.74 x 105; Mw = 1.20 x 106; Mz = 2.84 x 106) from Rohm and Haas Co.; also used irradiation degraded PMMA. PMMA, PS and w-PS-PMMA were prepared by radiation polymerization, after removal of inhibitors and vacuum purification. PS and PMMA from Chemical Agencies, Bombay. PS, Styron 634, Dow; EPDM, Keltan 514, DSM; 5B, Finaprene-410, Fina; SEP, Kraton G-1701X; Shell; and Kraton D-1118X, Shell. Materials are probably similar to those used in other publications from this group (see refs. 10, 31, 53, 54, and 57 in this Table). SAN, Luran 368R from BASF, acrylonitrile content = 36%; PMMA, 7H from Röhm AG, copolymer made from 95% methyl methacrylate and 5% ethyl methacrylate. LDPE (Stamylan 1808 AN, MFI = 7.5 dg/min) from DSM and PS (Styron 638, MFI = 25 dg/min) from Dow Chemical. PA and PB (polybutadiene) were prepared by emulsion polymerization. PS, Styron 686 E (RTM) from Dow; EPDM, Keltan 312, 378, 512, 578, 712 and 812 (RTM) from DSM; compatibilizers, Kraton G1650 and G1652 (RTM), and Cariflex TR 1108 (RTM), both from Shell; the authors claim that irradiations can be done under air or under inert atmosphere (vacuum or nitrogen); however, they do not specify the atmosphere for the data given. Sample size was not given. PS, Mw = 300,000, Mw/Mn = 2.5 (Dow Chemical) and PVME, Mw = 60,000, Mw/Mn = 2.1 (Dajac Labs). PVME was polymerized cationically with boron-trifluoride-ethyl ether complex as a catalyst; the fraction with Mw = 1,690,000 was used for preparing the blends. PSD, with Mw > 4,000,000, was prepared by emulsion polymerization. PVDF, KynarTM and other related fluorocarbons; HDPE, Petrothene LB 832 and other commercial grades; LCP, Vectra A950

Use of High-Energy Radiation in Polymer Blends Technology

781

and other liquid crystalline polymers; crosslinking agents; anti-oxidants, and thermal degradation inhibitors. See the reference for details of the various recommended materials. ax

Polysulphone (P-1700 and P-3500 from AMOCO); PVP (K-30 from BASF); and poly(ether sulphone) (VICTREX 4800P from ICI). ay CO1, PP/PE copolymer (#7371 from Fina); CO2, PP/PE copolymer (#6571 from Fina); EAA, ethylene/acrylic acid copolymer (PRIMACORTM, #3340 from Dow); EVAc, ethylene/vinyl acetate copolymer (#UE656 from Quantum); PB, polybutylene (#PB400 from Shell); PP, polypropylene (#PP3576 from Fina, and #5A95 from Shell). az ZYLAR series of polymer alloys from Novacor Chemicals, Inc., contain styrene/methyl metharcrylate copolymer and styrene/butadiene block copolymers. aaa See the patent for detailed composition of the blends and the results. aab Details were not given. aac Details about the source of ETFE and irradiation conditions were not given; TPE (Dai-el T-530 from Daikin Industries); many other components for many possible blends were mentioned; possible uses of many additives and crosslinking agents were mentioned. aad EPDM, Epcar 346 and 585 from Polysar; PAc, polymerization of acetylene with Ziegler-Natta catalyst contained in EPDM solution. aae Many experimental and materials details, not given. aaf PMMA (A-100 from Rohm and Haas); CO (Hydrin 100); for other details see the reference. aag PMMA (BDH, MW = 500 000, Tg = 105°C, MP = 225°C, specific gravity = 1.19); PVDF (Aldrich, MW = 350 000, Tg = 13°C, MP = 171°C, specific gravity = 1.8); PCTFE (Chemical Agencies, Bombay; Tg = 65-100°C, MP = 220°C, specific gravity = 2.08-2.20). aah PVC, Pevikon S 656 from Kema Nobel; PMMA, Diakon M0900 from ICI; PCL 700 from Union Carbide; and SAN, Luran 368 from BASF. aai HDPE (LH400, Asahi Chemical); PP (J640P, a propylene-ethylene copolymer, the ethylene content not given, Mitsui Petrochemical); PS 9 general purpose, PSGP, TOPOREX525-51, Mitsui Toatsu Chemicals), and PS (high impact grade, PSHI, TOPOREX825-21, Mitsui Toatsu Chemicals); and TAC as the multifunctional crosslinking agent. Ref: (1) Minkova, 1987; (2) Minkova and Mihailov, 1990; (3) Kurian et al., 1993; (4) Yu et al., 1992, 1994; (5) Rizzo et al., 1983; (6) Spadaro et al., 1984; (7) Kostoski et al., 1986; (8) Minkova and Nikolova, 1989a,b; (9) Bahari et al; 1993; (10) Meijer et al., 1988; (11) Spadaro et al., 1992; (12) Valenza et al., 1994; (13) Valenza et al., 1992, 1993a; (14) Valenza et al., 1993b; (15) Spadaro et al., 1993; (16) Acierno et al., 1994; (17) Spadaro et al., 1996; (18) Zhang and Xu, 1993; (19) Brooks, 1983; (20) Martínez-Pardo and Vera-Graziano, 1995; (21) Siquin and Chen, 1995; 1996; (22) Spenadel, 1979; (23) Lee and Chen, 1983; (24) Grossman, 1977; (25) Kammel and Wiedenmann, 1976; (26) Abdel-Aziz et al., 1992; (27) Hamza and Elwy, 1995; (28) Jia and Chen, 1994; (29) Jia et al., 1996; (30) van Gisbergen et al., 1989a,b; (31) van Gisbergen et al., 1991a; (32) Harnischfeger et. al., 1990; (33) Lee et al., 1982; (34) Chen and Lee, 1986; (35) Hamed, 1982; (36) Berejka, 1981; (37) Berejka and Bradley, 1981; (38) Rolando and Krueger, 1992; (39) Coran and Patel, 1982; (40) Akhtar et al., 1986; (41) Ivchenko et al., 1981; (42) Thomas et al., 1986; (43) Kumar et al., 1993; (44) Magerramov et al., 1981; (45) Thomas et al., 1987; (46) Koshy et al., 1992; (47) Bhowmick et al., 1986; (48) Bohm and Nelson, 1981; (49) Garrett et al., 1979, 1985; (50) Schultz and Mankin, 1976; (51) Torikai et al., 1994; (52) Katare et al., 1994; (53) van Gisbergen et al., 1990; (54) van Gisbergen et al., 1991b; (55) Elemans et al., 1988; (56) Nguyen and Kausch, 1984; (57) van Gisbergen and Meijer, 1991; (58) Witt, 1959; (59) Schrijver et al., 1988; (60) McHerron and Wilkes, 1993; (61) Briber and Bauer, 1988; (62) Toy et al., 1994; (63) Kobayashi and Tanaka, 1992; (64) Wilfong and Rolando, 1993; (65) Hauser et al., 1993; (66) Yoshimura et al., 1989; (67) Vesely and Finch, 1988; (68) Cheng et al., 1990; (69) Lee and Jopson, 1983; (70) Mason, 1995; (71) Jubinsky et al., 1986; (72) Mishra et al., 1994; (73) Bajpai et al., 1994; (74) Lindberg et al.; 1989; (75) Numata and Fujii, 1995.

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the second melting, were Tm(UHMWPE) = 142 and 133°C; and Tm(PE) = 129 and 126°C, respectively. The enthalpies of crystallization of individual PEs were higher after irradiation, whereas those of the blends hardly changed. Higher enthalpies of melting or crystallization of the irradiated components were attributed to chain scission within the amorphous regions and on the crystal surfaces, which for irradiation in air dominates over the crosslinking [Minkova and Mihailov, 1990]. In the presence of air, the chain scission would be primarily due to oxidative degradation through reactions such as (R-22) and (R-26). Crystallization enthalpies of the irradiated blends were lower than predicted by the additivity rule. The crosslinking of the amorphous domains lowered the melting enthalpies. Irradiated blends showed separated crystallizations, at their characteristic temperatures of crystallization, independent of composition. On irradiation, the crystallization temperature of PE did not change, whereas that of UHMWPE slightly decreased. Separate irradiation of UHMWPE and PE favored molecular scission over the crosslinking, whereas irradiation of the blends favored crosslinking. The melting temperature, Tm, increased linearly with

the concentration of UHMWPE. The value of the rate of melting, Vm, for PE was slightly higher, and for the irradiated UHMWPE it was 2 times higher than that of the non-irradiated UHMWPE. The increase of Vm of PE on irradiation was also attributed to the molecular scission. The temperature intervals for the melting of UHMWPE and PE on irradiation remained constant, while the melting enthalpies increased, leading to the higher rates of melting. Some of their data are given in Table 11.10. 11.4.1.2

High-Density and Linear Low-Density Polyethylenes, HDPE/LLDPE

Kurian et al. [1993] examined gamma-irradiation of the HDPE/LLDPE blends, at a dose rate of 10 kGy/h (Table 11.9). The blends of HDPE and LLDPE were prepared by melt mixing in an extruder attached to an internal mixer (L/D ratio, 20; screw compression ratio, 3:1; screw speed, 30 rpm; extruder heater zones, 140°C; die temperature, 150°C). For irradiation, the samples were compression-molded at 150°C in a steamheated laboratory hydraulic press for 5 min. For the dicumyl peroxide (DCP) crosslinkable samples, 1% DCP was added during compound-

Table 11.10. Melting and crystallization enthalpies of ultra-high molecular weight polyethylene (UHMWPE) and normal molecular weight polyethylene (NMWPE) blends [Minkova and Mihailov, 1990] Composition (wt%) NMWPE 100 100 80 80 60 60 30 30 0 0

Dose (kGy)

Melting enthalpy ∆Hm, J/g

Crystallization enthalpy ∆Hcr, J/g

0 60 0 60 0 60 0 60 0 60

183.5 189.4 170.1 154.6 214.6 154.6 180.6 155.8 139.4 176.4

157.1 183.5 143.6 183.5 170.1 179.8 157.1 178.5 127.3 201.6

UHMWPE 0 0 20 20 40 40 70 70 100 100

Use of High-Energy Radiation in Polymer Blends Technology

ing at higher temperatures (extruder at 190°C, die at 200°C) to trigger the crosslinking. Crosslinking by gamma-irradiation of HDPE, LLDPE, or their blends is more efficient when DCP (1-2%) is present. By contrast with the chemical crosslinking using DCP that significantly lowers the crystallinity, the radiation crosslinking does not seriously affect the crystallinity; thus it improves mechanical properties, such as tensile strength, yield strength, and abrasion resistance. The gel content data in HDPE irradiated in the presence of different DCP concentrations show that the crosslinking is ineffective without the peroxide. The gel content increased with DCP concentration and irradiation doses. For the same DCP concentration and irradiation dose, HDPE had higher gel content than LLDPE. On gamma irradiation of these polymers and blends in air, oxidative degradation would be expected, particularly in thin samples. The overall crosslinking of the polymers and blends reported by Kurian et al. [1993] suggests that the samples were most likely thick and that the crosslinking reactions caused by the presence of DCP (such as reactions (R-37), (R-38), and their analogues, also see [Singh, 1995]) outweighed the oxidative degradation reactions (such as reactions (R-22), (R-26) and related reactions). Although the authors report overall crosslinking in their samples, there would be oxidative degradation on the surface of the samples; the depth of oxidative degradation would depend on the dose rate [Kashiwabara and Seguchi, 1992]. The melting endotherms of HDPE, LLDPE and 50/50 HDPE/LLDPE blends showed only a single peak. The DSC data indicated that HDPE was more crystalline than LLDPE and that the crystallinity in the 50/50 blend was in between that of HDPE and LLDPE. The DSC thermograms of DCP crosslinked HDPE, LLDPE and HDPE/ LLDPE blends showed that crystallinity was sharply reduced during chemical crosslinking. However, since the radiation crosslinking takes place mainly in the amorphous regions, crosslinking by gamma irradiation did not severely affect the crystallinity.

783

The tension stress-strain curves of irradiated (100 kGy) HDPE, LLDPE, and their 50/50 blends with different amount of DCP, showed that the yield strength and tensile strength increased with DCP concentration, whereas the modulus and elongation at break decreased marginally. Whereas the abrasion loss of HDPE and LLDPE increased with irradiation and DCP content, the abrasion resistance of the blends increased with increase of the irradiation dose and the DCP content. 11.4.1.3. Low-Density Polyethylene and Polypropylene, LDPE/PP(1) Yu et al. [1992, 1994] have investigated electron irradiation of LDPE/PP blends in the presence and absence of peroxides (Table 11.9). The blends were prepared either in an internal mixer at 30 rpm (chamber temperature, 180°C), or in a 30-mm co-rotating, intermeshing twin screw extruder (100 rpm, 180°C, with vacuum devolatilization). The polymers were fed by two volumetric feeders, and either 0.1 or 1 wt% of an initiator was injected (as a 10% solution in cyclohexane). For the mechanical tests, the specimens were prepared using a 40-ton, 3.5-ounce capacity injection molding machine (ASTM D638). For injection molding, the injection pressure was 55.2 MPa, barrel temperature 205°C and mold temperature 38°C. To measure the uniaxial elongation viscosity of polymers and their blends, cylindrical samples were prepared by transfer molding under vacuum at 180°C, annealed and affixed with epoxy to aluminum ties. The electron irradiations were done with the I-10/1 electron accelerator (AECL). The samples were placed in plastic bags under nitrogen and were irradiated to doses of 30 or 100 kGy. These radiation doses were considered to be equivalent to 0.1 and 1 wt% peroxide, respectively, based on the gel fraction obtained in PE. If PE alone was irradiated, the gel fraction would be about 30% and 70%, for 30 kGy and 100 kGy, respectively. With 1 wt% peroxide in PE, the gel fraction obtained was 67%. The LDPE/PP blends that were unmodified by peroxide or irradiation showed non-Newtonian

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flow behavior and high viscosity. On irradiation, because of degradation of the PP, the complex viscosity of the blends decreased with the increase of irradiation dose and peroxide concentrations. The viscosities of the irradiated blends showed higher values than those of the peroxide-reacted blends. The complex viscosity of the peroxidetreated (1 wt%) and irradiated at 100 kGy blends, showed a crosslinked LDPE and highly degraded, low viscosity PP. The crosslinking in the LDPE component, and the degradation because of chain scission in the PP component are consistent with the effects expected for these two polymers. Since the work was done under nitrogen, no oxidative degradation of either of the components would be expected, in this case. The highly modified blends (100 kGy and 0.1 wt% peroxide) showed two distinctly different regions of the rheological behavior. At higher frequencies, the storage modulus, G’, of the peroxide-modified blends was lower than that of the irradiated ones. At lower frequencies the crosslinked LDPE phase, while at higher frequencies, the degraded PP matrix dominated the blends’ behavior. Thus in LDPE/PP blends, as the peroxide reaction takes place, LDPE becomes crosslinked and PP degrades. The size of the dispersed LDPE domains remained constant, the same as in unmodified blends, d ~ 0.5-1 µm. The yielding behavior of the irradiated, PP-rich blends slightly improved, whereas that of the peroxide-modified blends decreased with increasing peroxide concentration. For LDPE-rich blends the yield behavior showed improvement with increasing peroxide concentration and dose. In this case the peroxide would act as a mild crosslinking agent because of the formation of free radicals from it [Singh, 1995]. As a result, the increased formation of LDPE free radicals (e.g., reaction (R-37)), as well as the increased time available for the cationic reactions would lead to increased efficiency of crosslinking. Overall, for the PP-rich blends the irradiation process gives better mechanical properties, than the peroxide treatment. Some of the data reported by Yu et al. [1994] are shown in Tables 11.11 and 11.12.

An interesting application of irradiation has been used for medical devices [Lenzi, 1982, 1983]. Polypropylene is widely used for making medical disposables. However, it degrades during radiation-sterilization, and the degradation continues in the post-irradiation period [Williams, 1992; Kashiwabara and Seguchi, 1992; Wilfong and Rolando, 1993]. Lenzi [1982,1983] claimed that medical products (particularly hypodermic syringes) made from a PP/LDPE blend (LDPE between 2-8%, preferably 4-5%; materials and experimental details not given) could be sterilized using gamma irradiation, without losing the mechanical characteristics, particularly their impact strength. The LDPE should be carefully and intimately mixed with the PP, so that a homogeneous mixture of the two is obtained. When the LDPE content increases to greater than 8%, there is a tendency of the two polymers to segregate on cooling. On radiation sterilization, PP degrades, but the resulting loss of the mechanical properties is made up by the simultaneous crosslinking of the PE in the blend. The presence of the LDPE also inhibits the post-irradiation degradation of PP. Recent work by Burlinska et al. [1996] provides ESR evidence of the protection offered by PE against radiation-induced degradation of PP. 11.4.1.4

Low-Density Polyethylene and Polypropylene, LDPE/PP(2)

Rizzo et al. [1983] investigated changes in the physico-chemical properties caused by gamma irradiation of LDPE/PP blends (100:0, 75:25, 50:50, 25:75, and 0:100) (Table 11.9). On irradiation of the individual polymers with 1500 kGy dose, the gel fraction was obtained as 95% in LDPE and 65% in PP. On irradiation of the blends the gel fraction increased with the LDPEcontent and dose. On irradiation under vacuum, the crosslinking reactions predominated in LDPE, as discussed in Section 11.3.1. In the case of PP, accumulation of unsaturation with increasing dose contributes to extensive crosslinking at the high dose used here (1500 kGy), as discussed in Section 11.3.2.

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Use of High-Energy Radiation in Polymer Blends Technology

Table 11.11.

Complex viscosity and storage modulus of LDPE/PP blends at 180°C [Yu et al., 1994]

LDPE (wt%)

PP (wt%)

Peroxide (wt%)

Dose kGy

0 0 100 100 25 25 25 25 25 25 25 25 25 25 25 25 25 25 25 25 25 25 25 25

100 100 0 0 75 75 75 75 75 75 75 75 75 75 75 75 75 75 75 75 75 75 75 75

0 0.1 1.0 0 0.1 1.0 0 0.1 1.0 0 0.1 1.0 -

30 100 30 100 30 100 30 100

Frequency (ω), Complex Storage rad/s viscosity (kPas) modulus (kPa) 0.1 100 0.1 100 0.1 0.1 0.1 0.1 0.1 1 1 1 1 1 10 10 10 10 10 100 100 100 100 100

26 0.9 11 0.5 14 4.0 6.0 6.9 6.9 7.3 1.0 1.0 2.0 1.3 2.9 0.43 0.22 1.0 0.54 0.72 0.19 0.009 0.32 0.19

1.0 79 0.4 41 0.59 0.23 0.60 0.22 0.42 4.1 0.80 0.90 1.2 1.0 20 2.2 1.3 5.4 3.3 61 11 5.0 23 12

Table 11.12. Tensile and elongation properties of LDPE/PP blends [Yu et al., 1994] LDPE (wt%)

PP(wt%) (wt%)

Yield stress (MPa)

Tensile stress at break (MPa)

Elongation at yield (%)

Elongation at break (%)

Peroxide (%)

Dose (kGy)

25 25 25 25 25 75 75 75 75 75

75 75 75 75 75 25 25 25 25 25

27.6 23.2 20.9 28.7 28.5 13.3 14.3 13.0 14.0 14.0

18.8 23.5 20.8 17.5 23.8 11.5 14.2 12.9 9.5 14.0

7.4 7.7 4.1 7.6 7.6 38.0 55.0 71.5 30.0 51.5

100 9 4 25 16 80 76 77 78 62

0 0.1 1 0 0.1 1 -

30 100 30 100

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A. Singh and K. Bahari

An internal mixer was used to blend the samples (180°C, 20 rpm, 15 min). The blends were compression-molded into 5-mm-thick sheets in a laboratory press (200°C, 25 MPa). The specimens were evacuated for 24 h and sealed in glass under vacuum. The irradiations, using a Co-60 source, were performed at a rate of 5.2 kGy/h to doses of 300, 900 and 1500 kGy. To remove the trapped free radicals, and thus avoid their oxidative degradation on exposure to air, the irradiated samples were kept at 80°C for 5 days before breaking the seal. Thermal analysis of the blends showed that as a result of irradiation with progressively increasing dose, the melting temperatures of LDPE and PP decreased, by up to 5 and 30°C, respectively. The lowering of the melting temperature in the second run (remelting of the cooled melted samples) was attributed to the crosslinking of LDPE, which inhibited the structural reorganization of the polymeric chains. In the blends

with 75 wt% LDPE, crosslinking increased as did the difficulty of PP crystallization (Table 11.13). In blends comprising 25 wt% LDPE, the difficulty of LDPE crystallization was due to the crosslinks and the presence of PP. The density measurements of the irradiated blends were consistent with these results. The elongation at break was reduced by the blends’ irradiation. Whereas the LDPE and the blends remained ductile, PP became brittle. With increasing irradiation dose, the Young’s modulus of PP and the blends decreased, whereas that of LDPE was not significantly affected. The irradiation slightly decreased the degree of crystallinity — most likely the effect was located in the boundary regions between crystalline and amorphous phases. However, since modulus of a semicrystalline polymer is strongly affected by the properties of the boundary regions, a slight decrease of crystallinity could cause the observed effects.

Table 11.13. Fractional crystallinity of LDPE and iPP, in LDPE/iPP blends, evaluated from melting enthalpies [based on the data of Rizzo et al., 1983] LDPE (wt%)

0 25

Homopolymer

iPP LDPE iPP

50

LDPE iPP

75

LDPE iPP

100

LDPE

Heating run

1st 2nd 1st 2nd 1st 2nd 1st 2nd 1st 2nd 1st 2nd 1st 2nd 1st 2nd

Dose (kGy) 0

300

0.58 0.56 0.51 0.35 0.48 0.37 0.48 0.37 0.48 0.40 0.51 0.40 0.55 0.40 0.48 0.42

0.58 0.50 0.43 0.35 0.48 0.36 0.48 0.33 0.50 0.35 0.49 0.30 0.50 0.27 0.48 0.40

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The ultimate elongation of both homopolymers decreased with dose, but for the blends the effect was less pronounced. The ultimate strength can be related to the effect of crosslinking, which tends to increase the stress and reduce the ultimate elongation. Later, these workers investigated the effect of irradiation on the tensile properties of the LDPE/PP blends. For the non-irradiated blends, the mechanical behavior gradually changed with composition. The elongation at yield gradually increased from PP to LDPE, whereas the elongation at break showed a minimum at 75 wt% PP [Spadaro et al., 1984]. The results are shown in Table 11.14. Kostoski et al. [1986] also investigated gamma-irradiation of LDPE/PP blends, but in air

(Table 11.9). Melt blending was done in an internal mixer (200°C, at 30, 60 and 120 rpm, for 8 min). For comparison, the neat polymers were similarly treated. The samples were compressionmolded into 0.5-mm-thick sheets in a laboratory press (200°C, 1.75 MPa). The molded sheets were either air-cooled or quenched in ice water. Dogbone specimens (0.5 x 4 x 5 mm) were irradiated (Co-60, dose rate 14 kGy/h), to doses of 0, 25, 75, and 100 kGy. Tensile testing was done at RT, at 50 mm/min. Average values (four samples each) were calculated from the stressstrain curves: yield, tensile strength, elongation at break and tensile modulus at 5% elongation. In a later brief study, Kostoski and Kacarevic-Popovic [1988] report on the accelerated aging behavior of iPP/LDPE blends.

Table 11.14. Tensile properties of LDPE/PP blends [Spadaro et al., 1984] Dose (kGy)

E (GPa·102)

Ey (%)

Eb (%)

97 74 50 31 15

17 21 27 53 -

900 750 600 450 900

LDPE (wt%)

PP (wt%)

0 25 50 75 100

100 75 50 25 0

0 0 0 0 0

0

100 75 50 25 0 0

300

93

-

10

25 50 75 100

300 300 300 300

67 48 25 15

20 27 -

30 500 400 450

0 25 50 75 100

100 75 50 25 0

900 900 900 900 900

82 54 40 27 16

27 41 -

5 65 200 200 170

0 25 50 75 100

100 75 50 25 0

1500 1500 1500 1500 1500

80 55 39 26 16

30 43 -

35 75 90 160 95

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A. Singh and K. Bahari

For the blends, the elongation at break was found to be more sensitive to composition and morphology than the tensile modulus or yield strength. For the air-cooled samples (which would show normal crystallinity) the elongation at break decreased with dose up to 25 kGy, showing a minimum close to the 50:50 composition. Doses greater than 25 kGy had negligible effect on the reduction of the elongation at break. This reduction, as well as the inability of the cold drawing of the blends, was attributed to the two-phase character of the blends. On the other hand, for the ice-water-quenched samples (which would show very low crystallinity), the yielding, necking, and cold drawing were observed at all compositions. In this case, the elongation at break had a shallow minimum at LDPE/PP = 75:25. The tensile modulus and the yielding tensile strength did not show any significant dependence on the blend compositions. The melting and crystallization behavior was studied using the differential scanning calorimetry (7 mg samples, 10°C/min, up to 470 K, in N2).

The melting temperature of LDPE was unaffected by irradiation. The decrease in the melting temperature of PP indicated that its crystal structure was more affected by gamma irradiation than that of the LDPE, which is consistent with the radiation-induced chain scission expected in PP (Section 11.3.2). A summary of the results is presented in Table 11.15. Crosslinking of LDPE and degradation of PP during irradiation were used to explain the results, but the radiationinduced oxidative degradation was not considered. Although the dose rate used by Kostoski et al. [1986] is relatively high, some oxidative degradation of the samples at the surface would occur, since the irradiations are being done in air, as discussed for PE (Section 11.3.1) and PP (Section 11.3.2). On the basis of the changes of the elongation at break as a function of composition, the authors postulated undefined interactions between the amorphous phases of the two resins.

Table 11.15. Elongation at break of the LDPE/PP blends, irradiated in air [Kostoski et al., 1986] LDPE(wt%)

PP(wt%)

Dose (kGy)

Elongation at break (%)

0 0 0 0 10 10 10 10 30 30 30 30 50 50 50 50

100 100 100 100 90 90 90 90 70 70 70 70 50 50 50 50

0 25 56 100 0 25 56 100 0 25 56 100 0 25 56 100

730 19 14 12 650 20 13 11 70 22 14 11 40 20 18 12

Use of High-Energy Radiation in Polymer Blends Technology

11.4.1.5

Low-Density Polyethylene and Polypropylene, LDPE/PP(3)

Minkova and Nikolova [1989a,b] investigated the effect of electron irradiation on the thermodynamic and mechanical properties of thin films (100 µm) of LDPE/iPP blends (Table 11.9). Melt blending of the polymers was done at 180-225°C. The tubular films were extruded (145-190°C) and cooled slowly to room temperature (this would keep the crystallinity level high). The samples were irradiated (50-120 kGy) with electrons from an E-250 accelerator (energy not specified; probably 250 keV). In both types of blends, they observed segregational melting of the blend components. The peak melting temperature of either LDPE or iPP, in the blends, did not depend on the blend blend composition. The temperature of melting of the polymers, whether pure or in the blends, decreased with increasing dose (Table 11.16). The enthalpy of melting of the pure LDPE film decreases slowly with dose (Table 11.16). In the blends containing 70-80 wt% LDPE, the enthalpies of melting first increase with dose to 100 kGy, and then decrease to 160 kGy. However, in the case of the blend containing 55 wt% LDPE, the enthalpy of melting decreases with increasing dose, more rapidly than in pure LDPE. Minkova and Nikolova [1989a,b] attributed the observed effects to crosslinking in the amorphous regions and crystallite surfaces of LDPE as well as to scission in iPP. They also suggest that low levels of iPP in the blends (10 wt%) sensitize scission in LDPE. If this suggestion is correct, it could be that the lower molecular weight peroxy radicals formed by oxidative degradation of PP can migrate to the LDPE causing scission of its chains. It should be realized that even though electron irradiation is being used, based on the half-value layer data for oxidation given by Kashiwabara and Seguchi [1992], 10% or more of the film would undergo oxidative degradation, as discussed in Sections 11.3.1 and 11.3.2. The effect of the presence of the unspecified stabilizer is also difficult to predict. It is possible that the change in the trends of the effects at 100 kGy dose is related to possible depletion

789

of the 0.15% stabilizer. In the case of iPP, the authors suggested that the stabilizer affected the radiation effects up to a dose of 30 kGy [Minkova and Nikolova, 1989b]. The gel content of the LDPE/iPP blends (30-70%/70-30%) remained in the narrow range of 50-63% at doses of 50-120 kGy [Minkova and Nikolova, 1989a,b]; the reasons for this narrow range are not clear, though increasing crosslinking in the LDPE will be counterbalanced by degradation of iPP. Minkova and Nikolova, [1989a,b] found that the irradiated and the nonirradiated films were two phases, exhibiting two melting points. The temperatures of melting did not depend on the blend composition. The elastic modulus was found to increase with increasing dose. However, the tensile strength did not seem to be affected by the blend composition or dose (Table 11.17), the reasons for that are also not clear. The elongation at break decreased with increasing dose (Table 11.17). In a recent brief report, Tokuda and Kemmotsu [1995] suggest that electron-irradiated LLDPE/ PP-PE foams offer good mechanical properties. 11.4.1.6

Polypropylene and Co-Polypropylene, PP/cPP

Bahari et al. [1993] studied the radiation stability of PP/cPP (polypropylene/co-polypropylene, supplied by Toyoshoda, Japan) blends, by gamma and electron (3 MeV) irradiation (Table 11.9). The samples (40 g beads) were blended at a speed of 50 rpm for 5 min at 190°C. Polymer sheets (0.5 mm thick) were made from the blended material under a pressure of 15 MPa for 3 min at 190°C. The sheets were then immediately cooled between two plates of a cold press at 25°C under a pressure of 10 MPa for 5 min. The data for elongation at break for the samples is shown in Table 11.18. The values improve with the addition of cPP to PP, and they are better in the absence of the nucleating agent as compared to those in its presence. The decrease in the elongation at break values with increasing dose is because of oxidative degradation. A comparison of the values obtained on gamma irradiation with those

790

A. Singh and K. Bahari

Table 11.16. Effect of electron irradiation in air on melting parameters of 100-µm thick film made from EVAc/iPP and LDPE/iPP blends [Minkova and Nikolova, 1989a,b] Compositiona (wt%)

EVAc

iPP

99.85

0

89.85

10

69.85

30

49.85

50

39.85

60

0

99.85

LDPE

iPP

99.85

0

84.85

15

79.85

20

Dose (kGy)

0 53.3 100.0 163.2 0 53.3 100.0 163.2 0 53.3 100.0 163.2 0 53.3 100.0 163.2 0 53.3 100.0 163.2 0 53.3 100.0 163.2

0 53.3 100.0 163.2 0 53.3 100.0 163.2 0 53.3 100.0 163.2

Peak melting temperature (°C)

Enthalpy of melting (J/g)

EVAc

iPP

EVAc

iPP

124 118 116 114 124 119 116 115 124 117 115 113 123 117 117 116 123 119 117 116 -

164 157 149 166 157 150 149 164 158 152 150 165 157 152 149 162 159 154 151

144 146 146 126 146 153 169 133 142 147 155 120 139 149 167 113 143 142 142 123 -

154 152 150 116 151 150 132 113 b b b b 150 113 114 105 150 130 120 107

LDPE

iPP

LDPE

iPP

110 107 106 103 110 107 107 100 110 108 106 101

163 157 152 148 164 157 151 149

150 151 150 142 152 174 174 126 150 166 166 115

151 153 150 117 c c c c

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Use of High-Energy Radiation in Polymer Blends Technology

Table 11.16. Continued Compositiona (wt%)

EVAc

iPP

69.85

30

54.85

45

Dose (kGy)

0 53.3 100.0 163.2 0 53.3 100.0 163.2

Peak melting temperature (°C)

Enthalpy of melting (J/g)

EVAc

iPP

EVAc

iPP

109 106 108 100 110 107 106 100

164 156 152 146 164 158 151 150

140 162 162 111 150 143 140 117

c c c c 148 139 139 114

a All samples contained 0.15% of an unspecified stabilizer. b The data are similar to that for 30% iPP. c The data are similar to that for 15% iPP.

obtained on electron irradiation show that the oxidative degradation in the latter case is lower, as expected. 11.4.1.7

Low-Density Polyethylene and Polyamide-6, LDPE/PA-6

In a preliminary investigation, Meijer et al., [1988] reported that at a dose of 50 kGy, 42% of the LDPE was crosslinked in a blend with 3 parts of PA-6 (no details given) (Table 11.9). In comparison, the crosslinking of the LDPE was less efficient in blends with PS. Extensive work has been reported by Spadaro and co-workers (Table 11.9), to examine the effects of gamma-radiation-induced functionalization of PE (e.g., reactions (R-22), (R-26 to R-29), and (R-38)). It is known that polar groups on polyolefins help their compatibilization with polymers containing polar groups [Gaylord, 1989]. Use of chemically functionalized PE for preparing blends with PA-6 has also been examined [Chen et al., 1988; Xanthos, 1988; Raval et al., 1991]. Valenza et al. [1992, 1993a] examined the processing characteristics and mechanical properties

of the blends of PA-6 and irradiated LDPE (Table 11.9). Their experimental details are the same as those given in their other papers (described in this section). The addition of the non-irradiated LDPE to PA-6 leads to lower values of tensile modulus, tensile strength and elongation at break, as compared to the values for PA-6 alone (Table 11.19). Irradiation of the LDPE at the lower dose rate (0.108 kGy/h) slightly improves these values, whereas irradiation at the higher dose rate (3.6 kGy/h) leads to further reductions. The Izod impact strength values of the blends containing irradiated LDPE are higher than those containing non-irradiated LDPE (Table 11.19). Again, the values at the lower dose rates are higher. In another paper Valenza et al., [1993b] made a preliminary examination of the structure and properties of blends of PA-6 with gamma-irradiated LLDPE (10 wt%), using the experimental and processing techniques described before (Table 11.9). In the irradiated LLDPE, they observed the formation of both hydroxyl groups and carbonyl groups, as a result of radiation-induced oxidative degradation (through reactions such as (R-22), and (R-26 to R-29)). Both the elongation at break, and the Izod impact strength of the blends

792

A. Singh and K. Bahari

Table 11.17.

Effect of Electron Irradiation in air on 100-µm thick film made from EVAc/iPP and LDPE/iPP blends [Minkova and Nikolova, 1989 a,b]

Compositiona EVAc

iPP

89.85

10

69.85

30

49.85

50

39.85

60

LDPE

iPP

84.85

15

79.85

20

69.85

30

54.85

45

Dose (kGy)

Elastic modulus E(MPa)

Tensile strength (MPa)

0 53.3 100.0 113.2 0 53.3 100.0 113.2 0 53.3 100.0 113.2 0 53.3 100.0 113.2

nr 2.7 10.0b

29.0 27.5 21.7 24.0 nr 21.5 26.1 22.8 29.3 23.1 21.5 22.8 34.5 24.7 23.5 23.8

753 599 357 344 714 446 360 336 681 469 337 326 666 503 353 327

0 53.3 100.0 113.2 0 53.3 100.0 113.2 0 53.3 100.0 113.2 0 53.3 100.0 113.2

nr 2.0 4.0 5.1 nr 1.7 4.1 5.0 nr 2.0 2.9 3.4 nr 1.7 2.4 2.4

22.2 20.5 22.1 20.4 24.7 21.4 21.8 20.9 25.9 25.5 22.2 24.0 27.5 23.5 25.3 24.5

429 394 352 334 431 438 334 327 521 472 328 324 527 446 329 324

nr 1.7 4.1b nr 0.9 1.8b nr 0.7 1.3b

nr Not reported a All samples contained 0.15% of an unspecified stabilizer b Data for 113.2 kGy are the same, within experimental error (see their Figure 3)

Elongation at break (%)

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Use of High-Energy Radiation in Polymer Blends Technology

Table 11.18.

Effect of electron and gamma irradiation on the elongation at break of PP/cPP blends [Bahari et al., 1993] Eb, (%), e-

Eb,(%), γ PP (wt%) 100 100 100 100 100 70 70 70 70 70 50 50 50 50 50 30 30 30 30 30 0 0 0 0 0 cPP NA Eb,%,γ Eb,%,e-

cPP (wt%) 0 0 0 0 0 30 30 30 30 30 50 50 50 50 50 70 70 70 70 70 100 100 100 100 100

Dose, (kGy)

no NA

NA

no NA

0 10 25 50 80 0 10 25 50 80 0 10 25 50 80 0 10 25 50 80 0 10 25 50 80

86 50 6 471 297 204 74 0 500 471 307 173 0 564 524 410 213 0 682 695 556 300 0

51 11 0 200 21 3 0 0 310 153 10 0 0 363 285 100 12 0 495 508 351 300 280

85 21 4 0 0 470 422 247 196 81 500 463 340 276 82 551 521 458 450 420 674 700 688 596 450

NA 10 0 0 200 21 0 310 153 0 363 250 112 21 0 500 450 347 307 278

co-polypropylene nucleating agent, 1,3,2,4-di(p-ethylbenzylidine) sorbitol, 0.2% elongation at break (%) on gamma irradiation elongation at break (%) on electron irradiation

increased up to a dose of 50 kGy but decreased when the dose was increased to 135 kGy. For the three doses chosen (10, 50 and 135 kGy), there was no clear trend for the values of tensile modulus and strength. The SEM data showed that the compatibilization of the irradiated LLDPE was better in the blends with the PA-6. Spadaro and co-workers have studied the effect of gamma irradiation on the blends of

LDPE [Spadaro et al., 1992; Valenza et al., 1994] (Table 11.9). PE/PA-6 blends are antagonistically immiscible, showing poor ultimate properties [La Mantia and Valenza, 1988, 1989]. In an earlier study [Spadaro et al., 1992], gamma irradiation was used to functionalize LDPE, so that it may interact with PA-6. The LDPE sheets were compression-molded in a laboratory press (25 MPa at 240°C for 4 min) followed by cool-

794

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Table 11.19. Mechanical properties of LDPE/PA-6 blends [Valenza et al., 1992, 1993a] Blend (wt%)

Tensile modulus (GPa)

Tensile strength (MPa) Irrad.

Unirrad.

Elongation at break (%) Irrad.

Unirrad.

Izod impact strength (J/m)

PA-6

LDPE

Irrad.

Unirrad.

Irrad.

Unirrad.

100 90

0 10

1.33a

1.71

-

70

-

230

-

90

90 75 75

10 25 25

0.79b 1.09a 0.84b

1.19 1.19 0.99 0.99

44a 23b 40a 22b

43 43 33 33

20a 3b 31a 3b

260c 260c 310c 310c

130a 115b 150a 88b

110 110 73 73

a Dose rate, 0.108 kGy/h. b Dose rate, 3.6 kGy/h. c Interpolated value.

ing the platens with cold water. The sheets were irradiated using a Co-60 source at dose rates of 3.6 or 0.108 kGy/h, to a total dose of 100 kGy. The LDPE/PA-6 blends were prepared in an internal mixer (at 240°C at 15 rpm for 10 min). The LDPE samples irradiated at the higher dose rate (3.6 kGy/h) showed the formation of carbonyl groups and gel fraction of 30%, and those irradiated with the dose rate of 0.108 kGy/h showed much higher carbonyl content but absence of crosslinking. These observations are consistent with the expected radiation-induced oxidation in the system, as has been discussed earlier in Section 11.3.1. The authors attributed interaction between these carbonyl groups and PA-6 functionalities for the observed amelioration of the blends’ properties. However, the samples were not analyzed for the presence of hydroperoxides or peroxides that should also be formed by reactions (R-22), (R-23) and (R-29). The latter groups may also chemically interact with PA-6 through the formation of alkoxy free radicals [Ingold, 1969; Singh, 1995]. The trend of the property modification was similar in the blends containing 10 or 25 wt% LDPE. The torque was higher for the blends made with the irradiated LDPE. For example, when LDPE was irradiated at the lower dose rate, at the beginning of blending the torque was higher by

about 60% and about 700% at 10 min. Spadaro et al. [1992] attributed the high torque to two factors: crosslinking of LDPE (only for the high dose rate sample), and interaction of the carbonyl groups with PA-6. They also cited the development of turbidity in the Molau tests (blends containing irradiated LDPE) as evidence for the LDPE-carbonyl group-PA-6 interactions. Under the scanning electron microscope, SEM, the blends containing non-irradiated LDPE, fractured at liquid nitrogen temperature, showed phase separation, and those containing irradiated LDPE were homogenous. The latter observation was attributed to interactions between the carbonyl group and PA-6. The DSC data were also consistent with this hypothesis. In later works [Spadaro et al., 1993; Acierno et al., 1994; Valenza et al., 1994] the authors described the effect of irradiation on blends consisting 0, 10, 25, 75, 90 and 100 wt% of LLDPE and PA-6 (Table 11.9). Blending was done at 235°C at 20 rpm for 8 min. The samples were dried under vacuum at 100°C for 12 h. The irradiations, using a panoramic Co-60 source, were conducted in sealed glass tubes under vacuum, at a dose rate of 10 kGy/h up to doses of 50, 200 and 400 kGy. For the mechanical tests, the samples were compression-molded at 15 MPa.

795

Use of High-Energy Radiation in Polymer Blends Technology

The gel fraction and melt viscosities of the irradiated samples indicated that in the blends LLDPE was crosslinked (Table 11.20). The tensile modulus of the LLDPE-rich blends was found to be improved by irradiation. In contrast, the PA-6 underwent predominantly chain branching. As a result, impact strength of the PA-6-rich blends was found to be ameliorated [Valenza et al., 1994] (Table 11.20). For the neat LLDPE, the gel fraction varied from 84% at 50 kGy to 95% at 400 kGy. The gel fraction of the LLDPE decreased only marginally as the LLDPE component was reduced in the blends, e.g., from 95% at 100% LLDPE to 92% at 25% LLDPE at 400 kGy. However, at 10% LLDPE the gel fraction dropped sharply to 19%, indicating that at very low concentrations the radiation-induced crosslinking may be affected by partial solubilization of the LLDPE in PA-6 [Spadaro et al., 1993]. In the case of the irradiated neat PA-6, no gel fraction was observed. A comparison of the dynamic melt viscosity versus frequency for the non-irradiated and irradiated (200 and 400 kGy) PA-6 showed that the behavior of the irradiated samples was non-Newtonian, which was

attributed to an increase in the molecular weight through chain branching [Spadaro et al., 1993]. In later work, the melt flow of PA-6 indicated a viscosity increase without gel formation, attributable to chain branching [Valenza et al., 1994]. The Izod impact strength of the LLDPE/PA-6 blends showed an increase on irradiation to 200 kGy. However, (i) the impact strength of the non-irradiated blend containing LLDPE was higher than that of the non-irradiated PA-6; and (ii) irradiation increased the impact strength of the blend containing 25 wt% LLDPE, suggesting that the crosslinking of LLDPE contributed to the increase. In a recent paper, Spadaro et al. [1996] discused physical and structural characterization of the blends containing gamma-irradiated LLDPE, LDPE and HDPE (in air, at room temperature), with PA-6 (PE/PA-6, 25/75 wt%) (Table 11.9). The experimental details are similar to their previous work described in this section, except that the dose rates used were 0.1 and 0.04 kGy/h. In the case of the LLDPE/PA-6 blends, the SEM micrographs showed that the morphology of the samples containing irradiated LLDPE was

Table 11.20. Effect of gamma irradiation under vacuum on LLDPE/PA-6 blends [Spadaro et al., 1993; Valenza et al., 1994] Blend (wt%)

Dose

LLDPE

PA-6

100 100 75 75 25 25 25 10 10 10 -

25 25 75 75 75 90 90 90 100 100 100

(kGy)

Gel (%) LLDPE

Sol (%) PA-6

50 400 50 400 0 50 400 0 50 400 0 50 400

84 95 81 95 83 92 19 24 -

113 129 130 110 109 109 100 100 100

Impact strength (J/m) 14.9 24.5 13.4 13.5 8.0 11.1

796

A. Singh and K. Bahari

uniform, in comparison to the immiscibility clearly shown in the blends from the non-irradiated LLDPE. The SEM of the etched surfaces showed that the particle size of LLDPE decreases, and its distribution becomes more uniform, with increasing dose to the LLDPE. In general, the results obtained with the blends made with LDPE and HDPE also gave similar results, except that the distribution of the irradiated HDPE is somewhat less uniform than that of the LLDPE and LDPE. The melt viscosity of the blends made with irradiated LLDPE is higher than that of the LLDPE. This increase is attributed to increased interaction of the functional groups formed on the LLDPE, during radiation-induced oxidation, with the PA-6. The calorimetric data on these blends are also consistent with the interaction of the functional groups formed on LLDPE; the Tc of the LLDPE decreases with increasing concentration of the functional groups (e.g., with increasing dose or decreasing dose rate) because of increasing interaction with the PA-6 which hinders the crystallization process. Spadaro et al. [1996] also report that the complex viscosity of the LLDPE/PA-6 blends increases for the samples containing non-irradiated LLDPE but decreases for the samples containing irradiated LLDPE, as the mixing time for blending increases from 5 to 30 min. The blends were prepared at 240°C. It is known that polymers degrade at such high temperatures, through free radical reactions [Grassie, 1972]. The results

reported by Spadaro et al. [1996] suggest that in the case of the blends containing irradiated LLDPE, degradation takes place as a function of time (5-20 min, 240°C), but in the case of the blends containing non-irradiated LLDPE, crosslinking takes place during mixing, under the same conditions. 11.4.1.8

Blends of Polyethylene with Ethylene Vinyl Acetate and Chloro-Polyethylene, PE/EVAc and PE/CPE

Zhang and Xu [1993] blended EVAc and CPE (chlorinated-polyethylene) with PE (Table 11.9), because crosslinked PE has low flexibility for use as cable insulation. They mixed 30-70% HDPE with 70-30% EVAc or CPE, along with 0.5-1.0% of the antioxidant 1010 and 0.5-1.0% of zinc stearate in a SK-160B double roller mixer at 130-140°C. The blends so made were either hotpressed into sheets or extruded into tubes, and irradiated with 1.5-MeV electrons to a dose of 100 kGy. Some properties of the irradiated blends and PE are compared in Table 11.21. As the data in this Table show, the flexibility and the elongation of the HDPE/EVAc and HDPE/CPE blends are better than those of the PE but their hardness and softening points are lower. The two blends also show the memory effect better than that shown by the PE. Moreover, the heat-shrunk tubes of the PE/EVAc and PE/CPE blends were more flexible than those of the heat-shrunk PE.

Table 11.21. Some properties of HDPE/EVAc and HDPE/CPE blends [Zhang and Xu, 1993] Property Hardness, Shore A Tensile strengths (MPa) Elongation at break (%) Dielectric strength (MV/Mb) Softening Point (°C) a b

Presumably, irradiated to 100 kGy It is most likely MV/meter.

HDPEa

HDPE/CPEa

98 21.3 396 51.0 99

95 13.0 438 29.0 87

HDPE/EVAca 94 19.6 712 34.4 82

797

Use of High-Energy Radiation in Polymer Blends Technology

The heat-aging properties of the blends were also better than those of the PE. Blending of EVAc with PE results in reduction of the tensile strength [Brooks, 1983]; however, irradiation (with a 3-MeV electron accelerator in air [Berejka, 1996]) of the blend results in increase of the tensile strength, and elongation, as shown by the data in Table 11.22. 11.4.1.9

Low Density Polyethylene and Ethylene Vinyl Acetate Blends, LDPE/EVAc

Martínez-Pardo and Vera-Graziano [1995] studied gamma irradiation of LDPE/EVAc blends with the aim of improving their mechanical properties. LDPE, 36 wt%, EVAc, 35 wt%, TMPTMA, 4.5 wt%, flame retardant, 24.25 wt% and antioxidant, 0.25 wt% (Table 11.9) were first drum mixed and then fed into an extruder (L/D = 20/1) operated in the temperature range of 170 to 210°C. The extrudate was pelletized, the pellets masticated in a two roll mill (160°C 5 min) and then 2 mm thick plates were prepared in a hydraulic press (11.6 MPa, 160°C 5 min, then cooled with water for 10 min). Two samples of the LDPE/EVAc blends (C-1 and C-2) were made, one with each of the two antioxidants (Table 11.9). The LDPE samples were processed, similarly (it is not clear whether the LDPE sam-

ples also contained some of the additives). The samples were gamma-irradiated under argon, at room temperature, in sealed aluminum containers, to doses in the range of 100 to 160 kGy. Irradiated and non-irradiated samples were also aged for 168 h at 94°C (non-irradiated) and 136°C (irradiated) (125 changes of air/h), to see the effect of aging on their mechanical properties. According to the data shown in Table 11.23, the sample C-1 containing PTTBHP crosslinked more efficiently than the sample C-2 containing OTBHP; the samples that did not contain TMPTMA did not crosslink in the dose range 10 to 160 kGy. The authors claim that according to previous work, LDPE and EVAc do not crosslink below a dose of 250 kGy. This contradicts the data given by Chapiro [1962], Kammel and Wiedenmann [1976], Siquin and Chen [1995, 1996], Yu et al. [1992, 1994], and Singh et al. [1995], which showed that LDPE gets crosslinked at doses well below 150 kGy. It is possible that the LDPE used by Martínez-Pardo and Vera-Graziano [1995], contained high concentrations of antioxidant and other additives. The authors estimated that due to the presence of TMPTMA, the crosslinking dose was about 50-60% lower. As expected, the authors did find hydrogen and low molecular weight hydrocarbons as the volatile products; however, the

Table 11.22. Effect of irradiation on HDPE (1), HDPE/EVAc (2) and HDPE/EPDM/EVAc (3) blends at 150°C (based on the data given by Brooks [1983] Dose (kGy)

0 25 50 75 100 150

Tensile strength at break (kPa)

Elongation (%)

1

2

3

1

2

3

no break >97 no break >345 690 1028 -

138 379 393 641 655

131 -

>850 >850 725 680 -

>850 >800 >800 770 660

770 -

798

A. Singh and K. Bahari

Table 11.23. Gamma irradiation of LDPE/EVAc blends [Martínez-Pardo and Vera-Graziano, 1995] Samples

Dose (kGy)

Gel fraction (%)

Tensile strength (MPa)

Elongation at break (%)

Shore “D” hardness

Aged

Unaged

Aged

Unaged

C-1

0 10 25 100 125 150

1 23 68 70 70

8.2 10.3 10.8 10.3

525 202 207 195

552 182 189 130

40.90 44.40 44.85 45.90

34.90 48.00 47.00 47.00

C-2

0 10 25 100 125 150

1 19 59 52 67

7.6 8.7 8.9 9.1

510 157 141 124

569 131 121 98

41.00 47.90 48.00 47.50

39.30 48.00 48.30 47.70

absence of any oxygenated products suggested negligible decomposition of the EVAc or the TMPTMA. The tensile strength of the irradiated samples increases with increasing dose (Table 11.23); however, only the tensile strength of the sample C-1 reaches above the limit set by the UL 1581 standard (10.3 MPa). The elongation at break values of both C-1 and C-2 decrease with increasing dose (Table 11.23); however, again the values for C-1 at 100 to 125 kGy meet the requirements of the UL 1581 standard. The Shore D hardness of the two samples also increases with increasing dose, but the values are rather independent of the dose in the region 100 to 160 kGy (Table 11.23). The increase is again greater for the sample C-1, as compared to that for the sample C-2. The elastic limit increases by about 10% and the Young’s modulus by about 35% on irradiation of the samples. On aging of the non-irradiated samples, the tensile strength decreased by about 10%. However, aging of the irradiated samples led to

increases in their tensile strengths; 10% for C-1 and 15% for C-2 (Table 11.23). The crystallinity of the samples decreased by about 5% and their Shore D hardness decreased by about 2%, on aging. Siquin and Chen [1995, 1996] investigated irradiation of LDPE/EVAc blends for the production of crosslinked foam (Table 11.9). On gamma irradiation in air the gel fraction increased up to 30% with increasing EVAc content. (see Tables 11.24 and 11.25 for some of their data). EVAc by itself did not crosslink up to a dose of 35 kGy. For the blend containing 30% EVAc, the gel fraction on irradiation decreased in the order N2 > air > O2. The gel fraction at 35 kGy was much higher in the electron irradiated sample (55%) as compared to the gamma irradiated sample (35%). As one would expect, the formation of carbonyl groups on irradiation of the LDPE was higher in the sample in air than that in nitrogen. The apparent density of the irradiated samples increased with increasing gel fraction, whereas

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Use of High-Energy Radiation in Polymer Blends Technology

Table 11.24. Effect of irradiation on crosslinking of LDPE/EVAc blends [Siqin and Chen, 1995] LDPE (wt%)

EVAc (wt%)

Dose (kGy)

Gel fraction (%)

100 100 100 80 80 80 70 70 70 70 70 70 70 70 70 70 70 70 70

0 0 0 20 20 20 30 30 30 30 30 30 30 30 30 30 30 30 30

15 25 35 15 25 35 15 25 35 15 15 15 25 25 25 35 35 35 25

10 24 30 12 26 32 17 30 37 20 17 13 34 31 27 41 37 32 48(e-)

70 70 70

30 30 30

25 35 35

32 57(e-) 39

Cover gas air air air air air air air air air nitrogen air oxygen nitrogen air oxygen nitrogen air oxygen air air air air

e- denotes electron irradiation; all other data is for gamma irradiation

the expansion ratio to produce the foam, and the cell diameter of the foam, decreased with increasing gel fraction. The cell diameter on irradiation with electrons was larger than the cell diameter obtained on gamma irradiation. The tensile strength of the foam increased with increasing gel fraction and followed the order N2 > air > O2 for gamma irradiation. The tensile strength obtained by electron irradiation in air was greater than that obtained by gamma irradiation. The elongation at break of the gamma irradiated samples first increased and then decreased, as the gel fraction increased, following the order N2 > air > O2. The elongation at break values of

the samples irradiated with electrons were lower than those of the samples irradiated with gamma rays. 11.4.2 11.4.2.1

Elastromeric Blends Blends with Ethylene Propylene Rubber

11.4.2.1.1 PO blends with either EPR or EPDM Spenadel [1979], using an electron accelerator (3 MeV, 25 mA) at RT, crosslinked blends of either EPR or EPDM with PE or PP (Table 11.9). In EPR, crosslinking was easier when ethylene

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A. Singh and K. Bahari

Table 11.25. Effect of irradiation on tensile strength and elongation of LDPE/EVAc (70/30) blend samples [Siquin and Chen, 1995] Gel Fraction, (%) 20 20 20 30 30 30 40 40 40 50 a b c d e-

Cover gas

Tensile strength (Relative)a

Elongation at break (%)

3.0 2.9 2.4 3.8 3.6 2.8 4.5 7.4

283 270 263 347 333 253b 293c -d 263 257d

nitrogen air oxygen nitrogen air oxygen nitrogen air air (e-) air (e-)

Presumably, relative to the unirradiated sample. Peak value 300 at 27% gel fraction Peak value 367 at 35% gel fraction Peak value 337 at 32% gel fraction denotes electron irradiation; all other data is for gamma irradiation

content was higher and the viscosity was intermediate — for low ethylene content and low viscosity, the crosslinking response was poor. The specimens were prepared by injection molding (nozzle and barrel T = 205-221°C; mold T = 25°C; injection time t = 12 s; cooling time t = 20 s). The elastomeric specimens were prepared by pressing between PET sheets in a mold designed to prepare tensile pads at T = 110-120°C. The author argued that since LDPE requires 150-300 kGy doses for adequate crosslinking, whereas EPR and EPDM require 100-150 kGy, blends of EPDM with LDPE may respond well to radiation crosslinking. As the data in Table 11.26 show, LDPE blended with 30% VISTALON 3708, after a dose of 100 kGy had 80% gel fraction, whereas at that dose LDPE alone did not crosslink (this is unexpected; though the samples were ~ 2-mm-thick and electron-irradiated in air, at this dose the gel fraction of LDPE should be measurable, particularly since the author mentions that LDPE requires 150-300 kGy for adequate crosslinking; see Chapiro [1962], Kammel and Wiedenmann [1976], Yu et al. [1992, 1994],

Singh et al. [1995] and Siquin and Chen [1995, 1996], for crosslinking of LDPE at <150 kGy). Moreover, the addition of EPDM to LDPE followed by irradiation resulted in the materials showing: (i) better corona resistance; (ii) reduced voids; (iii) higher hot modulus; (iv) more flexibility; and (v) better sliceability. Three principal factors that influence the radiation response of EPDM are polymer molecular weight, ethylene content, and polymer unsaturation. It is the unsaturation (carbon-carbon double bonds) that promotes crosslinking by EPDM and the crosslinking agents, e.g., through analogues of the reaction (R-30) followed by (R-12). The addition of EPDM to EPR leads to crosslinking at a faster rate. The blends irradiated with 0-300 kGy showed increased tensile strength with increasing dose. The tri-functional trimethacrylate (TMPTMA) is a more effective crosslinking promoter than EDMA, and the addition of TMPTMA increases the tensile strength. The addition of a non-aromatic plasticizer may also increase crosslinking by facilitating free radical diffusion.

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Use of High-Energy Radiation in Polymer Blends Technology

Table 11.26. Gel fraction of irradiated EPDM/LDPE blends [Spenadel, 1979] Dose (kGy) 20 20 20 20 50 50 50 50 100 100 100 100 300 300 300 300 a

Vistalon 3708 (%)a 100 30 20 0 100 30 20 0 100 30 20 0 100 30 20 0

LDPE (%)a 0 70 80 100 0 70 80 100 0 70 80 100 0 70 80 100

Gel fraction (%)a 76 0 0 0 78 72 70 0 92 80 72 0 96 86 82 84

wt%?

Spenadel [1979] investigated a blend containing equal parts of EPDM, LDPE and carbon black (TEDX-15), to meet automobile fascia requirements. Although the peroxide-cured product was deficient, the radiation-cured product was better (see Table 11.27). The EPDM/PP blends containing carbon black showed reduced elongation as the dose increased — the carbon black evidently promoted crosslinking. In the blend of EPDM/PP with TMPTMA, the addition of carbon black followed by irradiation resulted in extremely high tensile strength (Table 11.28). 11.4.2.1.2 EPR/PP Blends Lee and Chen [1983] investigated gamma-radiation-induced changes in the dynamic properties of PP, EPR and PP/EPR blends (Table 11.9). EPR pellets were ground in a Thomas-Wiley laboratory mill, and the EPR powder was dried at RT. PP powder was dried at 80°C in a vacuum oven for 2 h. The dried resins were mixed in a kneader, extruded using a single-screw machine and pelletized. The pellets were dried in the vacuum oven

before use. The PP/EPR blend was hot-pressed (190°C, 2 MPa) in a compression molding machine to produce films. For irradiation, the specimens were sealed in a copper cylinder, with the air replaced by nitrogen. The samples were irradiated to doses of 50, 90 and 150 kGy. Dynamic testing of PP showed two damping peaks, corresponding to Tg and a secondary transition caused by the rotation of -CH3 around the polymer chain. Dynamic analysis of PP/EPR blends showed that the magnitude of both the storage modulus, E’, and loss modulus, E”, decreased with increasing EPR content. It also showed that the loss tangent, tan δ, of the irradiated blends had a new transition peak at T = -18°C for doses of 50-90 kGy. This new transition suggested the presence of a third phase, an interface, formed as a result of radiation-crosslinking between PP and EPR, which improved the impact strength and mechanical performance. SEM analyses showed that as the content of the EPR increased, the fracture surfaces became finer and the degrees of tearing were increased. The impact strength

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Table 11.27. Properties of irradiated and peroxide-cured TEDX-15a [Spenadel, 1979] Property

Additive (phr)

Dose (kGy) 0

Tensile strength (MPa) - do - do - do - do -

EDMA(1) EDMA(5) TMPTA(1) TMPTA(5)

Elongation (%) - do - do - do - do -

EDMA(1) EDMA(%) TMPTA(2,5) TMPTA(5)

Gel fraction (%) - do - do -

EDMA(1-5) TMPTA(1-10)

a b

11 11.5 11.5 11.7 18.1

Peroxidecuredb

10

30

12.8 13.9 14.1 14.2 14.5

16.3 17.3 17.6 17.9 18.1

18.1 18.1 18.1 18.1 18.1

325 360 345 305 330

325 330 280 250 240

250 240 215 180 170

280 280 280 280 280

2 3 3

70 76 76

83 82 82

88 88 88

Contains equal parts (by weight?) of EPR, LDPE and carbon black. At T = 205°C.

Table 11.28. Irradiation of EPDM/PP blend (60% vistalon 2504/40% PP) a [Spenadel, 1979] Property

Tensile strength (MPa) - do - do - do - do Elongation (%) - do - do - do - do Gel fraction (%) - do - do a

wt%?

Additive (phr)

EDMA(1) EDMA(5) TMPTA(1) TMPTA(5) EDMA(1) EDMA(5) TMPTA(2.5) TMPTA(5) EDMA(1-5) TMPTA(1-10)

Dose (kGy) 0

10

30

11 11.5 11.5 11.7 18.1 325 360 345 305 330 2 3 3

12.8 13.9 14.1 14.2 14.5 325 330 280 250 240 70 75 76

16.3 17.3 17.6 17.9 18.1 250 240 215 180 170 83 80 82

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Use of High-Energy Radiation in Polymer Blends Technology

changed with the composition of the blends and on irradiation, as shown by the data in Table 11.29. The impact strength for two of the blends (PP/EPR, 70/30 and 80/20, wt%) peaked at a dose of 50 kGy. 11.4.2.1.3 EPDM/PE Blends Blends of EPDM with PE were made (90 rpm, full water cooling) by using upside down loading, and the batch was discharged at 140°C in 4-5 min. [Grossman, 1977] (Table 11.9). The material was sheeted on a mill, then extruded (3/4 inch extruder, L/D = 10/1, with a 1.5/1 single rubber-type screw) as a 20 x 5 mm tape. For the heat cure, the materi-

als were mixed in two passes, with the curatives added in the second pass. After cooling the press, curing was conducted in an electric press and an ASTM D15-66 slab mold [Grossman, 1977]. It was suggested that for the best performance EPDM should contain >70% of ethylene and >3.5% of diene. The effects of irradiation as well as heat-cure, are shown in Table 11.30. The blends #1 and 2 are typical for flameresistant wire insulation. Blend #2 required a dose of 150 kGy for a cure comparable to that of heat-cured blend #1 with DCP. The heat-aging adversely affected properties of all the blends, more in the case of the irradiated than the heat-cured

Table 11.29. Effect of gamma irradiation under nitrogen on the impact strength of PP/EPR blends [Lee and Chen, 1983] Blend compositiona

a b

PP

EPR

100 100 100 100 95 95 95 95 90 90 90 90 80 80 80 80 70 70 70 70

0 0 0 0 5 5 5 5 10 10 10 10 20 20 20 20 30 30 30 30

wt%? originial data in kgf-cm/cm; converted, 1 kgf-cm/cm = 9.81 J/m

Dose (kGy)

Impact strengthb (kJ/m)

0 30 50 90 0 30 50 90 0 30 50 90 0 30 50 90 0 30 50 90

28 25 25 21 47 42 41 38 67 68 68 61 137 482 510 427 595 683 693 673

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A. Singh and K. Bahari

Table 11.30. Properties of heat-aged EPDM/PE blends [Grossman, 1977] Blenda # Dose (kGy) Cured sample tensile strength (MPa)c Elongation (%) (ii) Heat-aged sample (7 d at 121°C) Residuald tensile strength (%) Residuald elongation (%) (iii) Heat-aged sample (7 d at 150°C) Residuald tensile strength (%) Residuald elongation (%) (i)

a

b c d

1 0b

2 150

3 100

4 100

5 100

6 100

12.0 310

11.3 380

8.8 560

12.3 340

9.1 520

9.3 500

96 92

87 82

98 90

110 48

96 78

95 84

92 90

75 36

82 62

25 18

70 45

78 60

Composition of the blend was (in phr): Epsyn 5508 = 100; LDPE = 25; silane-treated clay = 90; Litharge = 7; Sb2O3 = 22, and paraffin oil = 10. The additives were (1) Age Rite MA = 1.5; DCP = 3.5; ZnO = 7. (2) Age Rite MA = 1.5; TMPTM = 4; ZnO = 7. (3) same as (2). (4) TMPTM= 4; ZnO = 7. (5) Age Rite MA = 1; ZnO = 7; TMPTM= 4. (6) Age Rite MA = 1; ZnO = 15; TMPTM= 4. Heat-cured for 20 min at 160°C. Assuming that the values given by Grossman [1977] are in psi. Percentage of the value for the cured sample.

blends. This was attributed to radiation-induced degradation of the antioxidant (Table 11.9). The data for the 100 kGy irradiated blends containing different levels of the antioxidant show that as the level of the antioxidant increased, the retention of properties on heat aging improved. The presence of antioxidants inhibits oxidative degradation during heat aging, in a manner similar to the inhibition of post-irradiation oxidation [Kashiwabara and Seguchi, 1992; Williams, 1992]. One may expect that antioxidants would reduce the level of crosslinking, as well as of oxidative degradation. Though the properties of the blends changed with increasing antioxidant concentration, without gel fraction measurements and infrared spectroscopy it is difficult to say whether the changes are due to crosslinking or oxidative degradation. The results also show that the presence of zinc oxide reduces the adverse effect of heat aging. The addition of EPDM to HDPE (Table 11.9) increases low-temperature flexibility; it also helps maintain high strength with high elongation values at room temperature [Brooks, 1983]. The pres-

ence of polyfunctional acrylates can be used to enhance the radiation crosslinking of polyolefins. Table 11.31 shows the values for the modulus and ultimate elongation obtained for the LDPE/EPDM and HDPE/EPDM blends containing two different multifunctional acrylates [Brooks, 1983]. As the data show, the modulus of the blend increases with increasing dose. However, it increases faster in the presence of the multifunctional acrylates TMPTA and TMPTMA, with the former being more efficient. Irradiation of the HDPE/EPDM blend in the presence of EVAc also leads to a good product, as shown by the reported tensile strength and elongation data [Brooks, 1983]. In the case of a HDPE/CDB/EPDM/TMPTA blend, the tensile strength at break values increase, and the elongation values decrease, as the dose increases from 0 to 125 kGy [Brooks, 1983]. 11.4.2.1.4 LDPE/EPDM Blends Kammel and Wiedenmann [1976] investigated effects of irradiation for various compositions on the LDPE blends with EPDM (Table 11.9).

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Use of High-Energy Radiation in Polymer Blends Technology

Table 11.31.

a

Effect of irradiation on HDPE/EPDM and LDPE/EPDM blends (based on the data of brooks [1983]

HDPE

LDPE

EPDM

Dose

Temperature

Additive

100% Modulus

(%)a

(%)a

(%)a

(kGy)

(oC)

(3%)

(MPa)

50 50 50 50 50 60 60 60 60 60 50 50 50 50 50 50 50 50 50 50 50 50 50 50 50 -

50 50 50 50 50 60 60 60 60 60

50 50 50 50 50 40 40 40 40 40 50 50 50 50 50 50 50 50 50 50 50 50 50 50 50 50 50 50 50 50 40 40 40 40 40

0 25 50 75 100 0 25 50 75 100 0 0 0 25 25 25 50 50 50 75 75 75 100 100 100 0 25 50 75 100 0 25 50 75 100

-40 -40 -40 -40 -40 -40 -40 -40 -40 -40 150 150 150 150 150 150 150 150 150 150 150 150 150 150 22 22 22 22 22 22 22 22 22 22

TMPTA TMPTMA TMPTA TMPTMA TMPTA TMPTMA TMPTA TMPTMA TMPTA TMPTMA -

wt%?

8,275 8,206 8,448 8,682 8,964 11,033 11,172 12,137 12,827 13,102 41 110 97 97 207 187 179 310 248 228 372 276 5,447 5 688 5 723 5 930 5 343 6 205 6 343 7 102 7 447 8 205

Ultimate elongation (%) 650 640 630 600 580 500 480 475 470 390 730 760 1 120 1 180 1 300 845 910 1 030 1 160 1 275

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The polymers were pressed (180°C, 3 min) into 1-mm plates, whereas the blends were formed into 1.6-mm-thick sheets. Antioxidants or stabilizers were not used in the blends to be irradiated. However, the blends prepared as 100-µm-thick film for thermo-oxidative stability studies contained 0.5% of the polymeric 2,2,4-trimethyl-1,2-dihydroquinoline as a stabilizer. The samples were irradiated with a 2-MeV electron accelerator. On irradiation of the two neat blend components, EPDM crosslinked faster than the LDPE; for example, at 100 kGy, the gel fraction for EPDM was 84% and that for LDPE 58%. In the blends of the LDPE containing 20, 40, 60 and 80 mol% of EPDM, the gel fraction increased with EPDM concentration, following the additivity rule. Reduction of the tensile strength and the elongation at break were more complex. The tensile strength of the blends decreased with increasing EPDM concentration, but it increased with increasing dose up to 200 kGy. This suggests that crosslinking of EPDM contributes to the improvement of the tensile strength. In the non-irradiated samples, the tensile strength dropped markedly as the concentration changed from 20 to 40% EPDM. In irradiated LDPE, the tensile strength increased slowly, reaching a maximum at 100 kGy, and then decreased gradually at higher doses. In the blends, however, this maximum was shifted to 50 kGy. A similar behavior was observed for ethylene copolymers [Lyons and Vaughn, 1967]. This was attributed to increased molecular weight preceding the crosslinking [Charlesby, 1960]. The 30-min values of the thermo-mechanical elongation of the LDPE/EPDM samples under a strain were found to decrease with the concentration of EPDM but increased with dose. However, the concentration effect was not linear — an initial addition of the EPDM markedly reduced the elongation (e.g., for 0 to 20% EPDM, at 100 kGy dose, from 500 to 100%), but further increases had smaller effect (e.g., for 20 to 50% EPDM, from 100 to 50%). The elongation at break at RT also decreased with increasing EPDM concentration, though less markedly (e.g., for 0 to 20% EPDM, at 100 kGy dose, from 520 to 480%) [Kammel and Wiedenmann, 1976]. In conclusion,

the authors noted that the blends containing 30% EPDM (which contained 8% ENB (5-ethylidene-2-norbornene)) and irradiated to a dose of 80-100 kGy should meet the requirements of the Verband Deutscher Elektrotechniker tests #0208 and 0209, as well as have the optimum thermooxidative stability. Replacement of the ENB in the EPDM with dicyclopentadiene reduced the thermo-oxidative stability of the blend by a factor of 3. The relative behavior of these two dienes was consistent with the later observations by Brooks [1983]. Abdel-Aziz et al. [1992] studied effects of gamma irradiation of EPDM blended with LDPE (Table 11.9). The mixtures were prepared in a rubber mill at 70°C. The sheeted materials were compression-molded (2-mm-thick) in an electrically heated hydraulic press at 150°C for 5 min. Increasing the LDPE content resulted in an increase of the tensile strength and modulus but a decrease in elongation at break. The tensile strength of the blends also increased with dose up to 100 kGy, beyond which there was a gradual decrease (see Table 11.32). Even in the case of LDPE alone, an increase (caused by crosslinking) with dose up to 100 kGy was observed. The presence of EPDM improved the tensile properties of the irradiated blend. The elongation at break decreased with the radiation dose as well as with LDPE concentration. However, at doses exceeding 100 kGy, no significant change of elongation at break with increasing LDPE concentration was observed. For doses lower than 100 kGy the radiation-induced crosslinking took place at a higher rate than the degradation. Though the bulk of the sample seems to show predominantly crosslinking, since the gamma irradiation was in air, there would be a thin layer at the surface where oxidative degradation is more important than crosslinking. The permanent set of LDPE/EPDM blends increased with LDPE concentration at all radiation doses, and at lower doses the rate of increase was higher. This could be attributed to the fact that permanent set values are measured after rupture of the sample. However the magnitude of the increase in permanent set was a function of

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Use of High-Energy Radiation in Polymer Blends Technology

the elongation at break — at higher radiation doses crosslinking of both EPDM and LDPE resulted in a decrease of the elongation at break. Hamza and Elwy [1995] investigated the effect of gamma irradiation on LDPE/EPDM blends containing 30 phr HAF carbon black (Table 11.9). The blends were prepared in a two-roll mill with a gear ratio of 1:4 (temperature, 70°C; mixing time, 5 min). The relationship between stress and strain at the four doses used (0, 200, 300 and 400 kGy) appeared to be parabolic (Table 11.33). The fracture stress, yield stress and elastic modulus also increased, whereas the ductility decreased, with increasing dose. These changes in the mechanical properties of the LDPE/ EPDM blends were attributed to radiationinduced crosslinking.

11.4.2.1.5 LDPE/EPDM/CB Blends Jia and Chen [1994] investigated the effect of irradiation on the positive temperature coefficient (PTC) function of LDPE/EPDM and LDPE/EEA (ethylene-ethyl acrylate) blends containing carbon black. They used LDPE, EPDM, EEA and CSF carbon black (Table 11.9) in the desired ratios and mixed them in a two-roll mill in an internal mixer at 130°C for 5 min. The samples were molded under pressure at 160°C for 10 min. On irradiation the PTC intensities of the LDPE/ EPDM/CB and LDPE/EEA/CB blends increase by several orders of magnitude (Table 11.34). In the case of the LDPE/EPDM/CB blend, the negative temperature coefficient that is seen in the case of the non-irradiated samples, disappears on irradiation. Jia and Chen [1994] suggest that

Table 11.32. Tensile strength (MPa) of EPDM/PE blends [Abdel-Aziz et al., 1992] PE Fraction a

0.0 0.32 0.483 0.737 1.0

Dose (kGy) 50

100

200

300

400

1.36 4.53 6.30 8.76 11.40

1.95 4.85 7.60 10.50 13.40

1.1 3.8 5.4 8.0 10.0

0.82 3.55 4.90 7.70 8.50

0.75 3.45 4.30 7.30 6.00

Table 11.33. Mechanical properties of gamma irradiated EPDM/LDPE blendsa [ Hamza and Elwy, 1995] Dose kGy

Fracture stress (MPa)

Yield stress (MPa)

Elastic modulus (MPa)

Ductility (%)

0 200 300 400

0.29 0.44 0.50 0.58

0.16 0.22 0.29 0.35

390 1180 1570 1960

216 157 122 89

a LDPE:EPDM = 1:1 by weight

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A. Singh and K. Bahari

Table 11.34. Effect of irradiation on PTC values of LDPE/EPDM/CB and LDPE/EEA/CB blends [Jia and Chen, 1974] Blend composition ratiosa

a

Temperature (°C)

LDPE

EPDM

EEA

CB

3 3 3 3 3 3 3 3 3 3 3 3 3 3

1 1 1 1 1 -

1 1 1 1

0.18 0.18 0.18 0.18 0.18 0.20 0.20 0.20 0.20 0.20 0.19 0.19 0.19 0.19

60 80 100 110 120 60 80 100 110 120 60 80 100 120

Log (volume resistivity) (ohm.cm) 0 kGy

150 kGy

5.5 6.9 8.7 6.0 8.1 9.5 7.5 6.5 7.0 7.9 8.1

5.2 6.4 9.2 9.1 7.1 8.5 11.5 13.1 6.2 7.5 9.2 9.7

by weight?

due to radiation-induced crosslinking, the CB agglomerates get bound more strongly in the polymer matrix, leading to the increased PTC. In a later publication, Jia et al. [1996] report further on the improvement of the PTC of the LDPE/EPDM/CB blends, which were irradiated in air at room temperature with either a Co-60 gamma source, or a 10-MeV electron accelerator (most probably, it is 1-MeV, rather than 10-MeV). The PTC intensity increased with gamma irradiation and the negative temperature coefficient (NTC) disappeared at a dose of 40 kGy. This effect was again attributed to inhibition of the reagglomeration of the CB particles because of the crosslinking of the polymer blend. The NTC reappeared at the very high doses of 2000 and 4000 kGy (Table 11.35); this was attributed to oxidative degradation of the blend. In comparison, the samples irradiated to doses of 500, 2000 and 3000 kGy with electrons showed no NTC and greatly improved PTC. Based on the WAXD and

DSC data, the authors concluded that the crystallinity of the samples decreased considerably at the high doses on gamma irradiation, but less so on electron irradiation. 11.4.2.1.6 PP/EPDM Blends To improve impact properties of PP, the resin is usually modified by incorporation of an elastomer. Since the performance depends on morphology, radiation crosslinking was used in its stabilization [van Gisbergen et al., 1989a] (Table 11.9). The blends were made either in a two-roll mill, at 185°C, then compression molded into 1-mm-thick sheets, or in a co-rotating twin-screw extruder, then pelletized. The irradiations (100 kGy) were done using a 3-MeV electron accelerator. Irradiated and non-irradiated pellets were injection-molded. For the DSC measurements, the samples were melted at 200°C, quenched to 110°C, and then heated at a rate of 10°C/min.

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Use of High-Energy Radiation in Polymer Blends Technology

Table 11.35. Effect of irradiation on PTC values of EDPE/EPDM/CB (3:1:1) blend a [Jia et al., 1996] Dose (kGy)

Temperature (°C)

Log(volume resistivity) (ohm.cm)

0 0 0 0 30 30 30 30 40 40 40 40 1000 1000 1000 1000 4000 4000 4000 4000

60 80 100 120 60 80 100 120 60 80 100 120 60 80 100 120 60 80 100 120

4.79 6.49 4.40 5.19 6.72 10.84 6.72 6.27 8.52 11.95 12.11 6.04 7.98 11.89 11.75 6.33 12.04 8.82 7.96

500 500 500 500 3000 3000 3000 3000

60 80 100 120 60 80 100 120

5.78 7.10 12.43 12.25 6.40 7.60 12.88 12.88

Gamma irradiation

Electron irradiation

a

by weight?

Irradiation of PP reduced the MW and viscosity, as expected (Section 11.3.2). The dynamic mechanical analysis of irradiated blends showed strong interactions between the PP matrix and the dispersed EPDM particles. Because of the reduced viscosity of PP, caused by irradiation,

the reorganization from imperfect to perfect crystals was faster, resulting in two melting peaks. In blends, this behavior was observed only at higher doses. As shown in Table 11.36, irradiation of the blends caused crosslinking of EPDM, as well as a pronounced change of the

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A. Singh and K. Bahari

Table 11.36. Gel fractions of irradiated EPDM and PP/EPDM blends [van Gisbergen et al., 1989a] PP/EPDM a

0/100 50/50 70/30 a

PP/EPDM (K514)

PP/EPDM (K578)

20 kGy

50 kGy

20 kGy

50 kGy

38 48 44

69 82 74

34 40 41

70 90 70

by weight

rheological behavior. In addition to the crosslinking, there was evidence of PP-grafting onto the EPDM particles. The morphology fixation could be enhanced by this phenomenon, which may improve flow behavior, low-temperature properties, and impact strength. In another study, electron irradiation and crosslinking of the dispersed EPDM phase was used to improve the processability of the high molecular weight polypropylene, HMWPP, blends with EPDM [van Gisbergen et al., 1989b]. The radiation-induced crosslinks stabilized this optimum morphology during injection molding, ascertaining good mechanical properties. The SEM micrographs showed that the dispersed EPDM particles in the HMWPP matrix were small and uniformly distributed. After injection molding of the nonirradiated blends, EPDM particles remained small. The extruded 70/30 blends were irradiated with a dose of 44 kGy, which resulted in a decrease of the degree of orientation. The EPDM particles dispersed in low molecular weight matrix, LMWPP, were irregular and large. The particles in the injection-molded blends were more regular and smaller than in the extruded specimens. The rheological behavior of the blends was found to be more complex than that of PP. For blends, there was a strong dependence of viscosity on frequency. It appears that the reactions occurring at the PP/EPDM interface play an important role, both in the determination of the overall viscosity and in the morphology fixation process.

The impact strength and mechanical properties of the irradiated (44 kGy) HMWPP/EPDM blends were superior to those of the LMWPP/ EPDM blends. However, comparison of the non-irradiated and irradiated HMWPP/EPDM blends indicated that irradiation caused a small decrease of the impact strength, modulus, and yield strength, attributable to the crosslinking of the EPDM, and chain scission of the PP matrix. In a follow-up to the earlier work [van Gisbergen et al., 1989a,b], van Gisbergen et al. [1991a] investigated melt rheology of electron-irradiated blends of PP and EPDM. EPDM is widely used to improve the impact properties of PP; for best results, the EPDM particles should be small (diameter ~0.3-0.5 µm). Because the viscosity of EPDM is high, this morphology induction requires the use of high molecular weight (high viscosity) PP. On the other hand, easy processing (injection molding) requires low viscosity. The aim of this work was to see if irradiation of these blends would result in a lower viscosity. Their experimental approach to sample preparation was as described earlier [van Gisbergen et al., 1989a,b]. The materials used included a highdiene-containing EPDM (EP-H), a low-dienecontaining EPDM (EP-L), and one containing no diene (EP). The EPDM or EP content of the blends made from high molecular weight PP (13E10) varied from 5 to 30%. The measured viscosities depended on the frequency used during the DMA (dynamic mechanical analysis) measurements. At high frequencies

Use of High-Energy Radiation in Polymer Blends Technology

(102-103 rad/s), the viscosities of the irradiated blends were lower than the viscosities of the corresponding non-irradiated blends. Thus in principle, injection molding of the irradiated blends should be easier than that of the non-irradiated blends. At low frequencies (~10-1), the viscosities of the irradiated blends depended on the EPDM content of the blend. The viscosities of the blends with low EPDM content were lower than those of the non-irradiated blends. However, the viscosities of the blends with high EPDM content were higher than those of the non-irradiated blends. Some of the data reported by the authors are shown in Table 11.37. For the PP/EP-L 70/30 blend irradiated to 50 kGy dose, there was hardly any effect of temperature in the range 190 to 290°C, which suggests that the irradiated blend demonstrates typical network behavior. van Gisbergen et al. [1989ab; 1991a] determined the inter-particle distance in the blends by SEM measurements, and concluded that for inter-particle distances above a critical value (~2.5 µm), the viscosity is mainly determined by the PP matrix. With decreasing inter-particle distances and increasing viscosity of the rubber, the network behavior becomes dominant. Harnischfeger et al. [1990] also investigated the effect of electron irradiation on the mechanical properties of blends of iPP and PP-PE, with EPDM (Table 11.9). Immiscible blends of PP and EPR exhibit high impact strength because of

811

the presence of dispersed rubber particles in the thermoplastic matrix. The aim of the investigation was to see whether radiation-induced crosslinking of the components of the blends improves the mechanical properties of the blends. According to the authors, earlier work by Nakamura et al. [1986] had successfully demonstrated interphase crosslinking of immiscible blends. The semicrystalline polymers used were iPP, iPP-PE copolymer, and EPDM. The blends contained either 35% or 65% EPDM. X-ray analyses showed the blend samples to be highly isotropic. Radiation-induced crosslinking of iPP and PP-PE is inefficient, with the gel doses (gel dose is defined as the lowest dose at which the formation of a gel, that is due to crosslinking, can be observed) being 150 and 200 kGy, respectively [Harnischfeger et al., 1990]. In the case of EPDM, on the other hand, the crosslinking is very efficient with the gel dose being 3.4 kGy. Even the presence of the small amount of the diene, ENB (~3 double bonds per 1000 carbon atoms) in the ethylene-propylene copolymer dramatically improves the efficiency of radiation-induced crosslinking, as the data just quoted show. The gel dose in the blends containing 35% or 65% EPDM was a bit higher (iPP/EPDM, ~3.7 kGy; PP-PE/EPDM, ~5.8 kGy); however, it was still very much lower than the gel doses of the two polyolefin polymers. As discussed in Sections 11.2 and 11.3, the presence of carboncarbon double bonds in the system enhances the

Table 11.37. Viscosities of irradiated and unirradiated PP/EP-H blends [van Gisbergen et al., 1991a] PP (wt%)

EP-H (wt%)

Dose (kGy)

Frequency (rad/s)

Viscosity (kPa.s)

95 95 70 70 95 95 70 70

5 5 30 30 5 5 30 30

0 50 0 50 0 50 0 50

102 102 102 102 10-1 10-1 10-1 10-1

0.80 0.10 1.00 0.14 30.0 0.20 8.00 70.0

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probability of both the free radical initiated, and the cation initiated, crosslinking reactions. In the case of both the mechanical properties studied (tensile strength and elongation), on comparison of the properties of the component polymers and the blends, the non-irradiated blends showed negative deviation from the linear additivity of the properties. However, on irradiation, this negative deviation was changed into positive deviation. Some of the mechanical

properties of the polymers and their blends reported by Harnischfeger et al., [1990] are given in Table 11.38. The authors interpreted the negative deviation from the linear additivity as an indication of the introduction of mechanical weakening sites in the blends. The positive deviation on irradiation was interpreted as an indication of mutual crosslinking of the component polymers at the interface.

Table 11.38. Effect of electron irradiation of polyolefin/EPDM blends, under vacuum [Harnischfeger et al., 1990] Blend composition (wt%) iPP

PP-PE

EPDM

100 100 100 100 65 65 65 65 35 35 35 35 0 0 0 0 0 0 0 0 0 0 0 0 0 0 0 0

0 0 0 0 0 0 0 0 0 0 0 0 0 0 0 0 100 100 100 100 65 65 65 65 35 35 35 35

0 0 0 0 35 35 35 35 65 65 65 65 100 100 100 100 0 0 0 0 35 35 35 35 65 65 65 65

Dose (kGy)

Tensile strength (MPa)

Elongation (%)

0 36 105 390 0 36 105 390 0 36 105 390 0 36 105 390 0 36 105 390 0 36 105 390 0 36 105 390

36 34 10 1 27 23 20 13 19 21 20 10 15 11 5 2 37 32 24 3 32 22 16 20 20 22 13 21

450 0 0 0 600 400 300 0 620 510 430 110 130 700 320 100 540 0 0 0 600 300 510 10 760 600 490 140

Use of High-Energy Radiation in Polymer Blends Technology

11.4.2.1.7 EPR/PBT Blends The effects of gamma irradiation and of DCP crosslinking of polybutyleneterephthalate, PBT, blends with EPR were investigated [Lee et al., 1982] (Table 11.9). PBT/EPR blends were prepared either by reactive or mechanical blending. In the first case, BDO and DMT, at a molar ratio = 2, with 300 ppm catalyst, were charged into a N2-flushed, stainless steel, stirred reactor. After 1 h at 190°C, methanol started to be collected for an hour. Next, the temperature was increased to 210°C and 30 wt% of EPR particles were added into the reaction mixture. The polymerization of PBT in the presence of EPR resulted in a semi-interpenetrating network. TGA showed that the weight loss of EPR was negligible; thus EPR was found to be thermally stable during polymerization. For the mechanical blending in an extruder, particles of PBT and EPR were dried in a vacuum oven at, respectively, 120°C and at RT for 3 h. The components were first mixed in a high-speed blender, then extruded in a single-screw machine and pelletized. During the extrusion, a small amount of DCP was added to crosslink EPR. DCP did not affect the properties of the blend, although it eliminated the delamination. The surface of the sample was coarse. Gamma irradiation of the blend prepared by reactive blending did not show any significant effects on PBT and the blends’ properties. Since the EPR/PBT blends were irradiated after injection molding, the delamination phenomenon was still observed on the surface of the sample. The extruded blends were irradiated before injection molding. The impact strengths of the samples in the BI series (without crosslinking) were higher than those of the BR series (crosslinking by gamma irradiation). The reason was that the soft domains of EPR became hard regions through crosslinking and lost their impactmodifying ability. On the other hand, the tensile strength and modulus of the BR series had higher values than the BI series. The extruded BI-blends of PBT/EPR showed higher toughness than those made by reactive blending. In the latter case, the average domain size of EPR was much smaller than in the former. The tensile strength, modulus

813

and flexural strength of the blends decreased as their EPR content increased. Observations with a polarizing microscope showed that the blends without any crosslinking displayed delamination — addition of DCP eliminated it, but owing to local crosslinking by DCP it produced a coarse surface. The gammairradiated blends showed lack of delamination and a smooth surface. In the later study, the effect of gamma irradiation, under nitrogen, on the glass-fiber-reinforced EPR/PBT blend was investigated [Chen and Lee, 1986] (Table 11.9). In two-phase EPR/ PBT blends, the physical properties were affected by poor interactions between the two components. Addition of 15 wt% EPR increased the impact strength of PBT by a factor of 3, but the tensile strength and modulus were reduced to less than half. During injection molding, the high melt viscosity of EPR caused the rubber particles to agglomerate and elongate in the flow direction, which resulted in lamellar peeling of the moldings. Radiation-induced crosslinking made EPR more rigid, thus it decreased the impact strength. The microscopic observations of the injection molded blends showed that the EPR particles were crosslinked, and thus were insoluble in xylene. It was also found that addition of glass fibers to EPR/PBT blends improved the modulus and tensile strength, and minimized the peeling. 11.4.2.1.8 EPDM/BR Blend Hamed [1982] investigated the effect of morphology and radiation-induced crosslinking on the energy to break of a blend of EPDM and polybutadiene (BR) (Table 11.9). The mixing of EPDM and BR was done in an internal mixer, and the blended rubber was pressed into a flat sheet and irradiated with an electron beam. The effect of irradiation on the energy to break for the 75/25 (by weight) EPDM/BR blend varied with the mixing time. At 20 kGy dose, the energy at break increased only slightly as the processing time for the blend varied from 2 to 60 min. However, the increase was much larger (by about a factor of 3) at the 160 kGy dose, as shown by the data in Table 11.39. As the blend processing time increases, the size and shape of the

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A. Singh and K. Bahari

dispersed phase changed from large and fibrous, to small and spherical, as shown by the transmission electron spectroscopy data. The largest increase in the energy to break occured between the processing times of 2 and 5 min, when the dispersed phase particle size decreases from about 100 µm to about 1 µm. There was somewhat smaller increase in the energy to break as these particles obtained at 5 min are changed into elongated particles at 15 min. At doses higher than 60 kGy, the tensile samples were found to undergo a slight stress-whitening upon stretching. Hamed [1982] also investigated the interfacial adhesion between EPDM and BR, as a function of the irradiation dose (Table 11.40). The modes of failure were attributed as follows: (i) cohesive failure in BR at low doses (0-20 kGy); (ii) slip-stick behavior in which the locus of failure alternated between cohesive and adhesive (~30 kGy); and

(iii) interfacial failure ((50 - 140 kGy). The author suggested that since EPDM and BR are quite immiscible, interfacial crosslinks are unlikely to be formed on irradiation. As irradiation proceed, at low doses, crosslinking leads to an increase in their tear strength, but, at high doses, a decrease in the tear strength results from extensive crosslinking, hence a tight network unable to dissipate much energy. Hamed [1982] suggested that at low doses the interfacial bond is stronger than the cohesive strength of the BR, hence the BR tears. With increasing irradiation the BR becomes stronger, comparable to the interfacial strength, leading to the observed slip-stick behavior. At the high doses the loss of dissipative capacity in the two rubbers results in weak interfacial failure.

Table 11.39. Effect of irradiation and mixing time on the energy to break of EPDM/BR (75/25) blendsa [Hamed,1982} Polymer

BR EPDM EPDM/BR BR EPDM EPDM/BR BR EPDM EPDM/BR EPDM/BR EPDM/BR EPDM/BR EPDM/BR EPDM/BR EPDM/BR EPDM/BR EPDM/BR EPDM/BR a

wt%(?);

b

Dose (kGy)

Mixing time (min)

log Eb (Energy to break, J/m3)

20 20 20 80 80 80 160 160 160 20 80 160 20 80 160 20 80 160

2 2 2 5 5 5 15 15 15 60 60 60

6.48 7.18 7.08 6.25 5.91 6.14 6.06 5.77 5.86 7.13 6.35 6.05 7.15 6.45 6.26 7.15 6.47 6.33

Eb = energy to break (J/m3)

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Use of High-Energy Radiation in Polymer Blends Technology

Table 11.40. Effect of irradiation on the peel strength of EPDM-BR adhesion [Hamed, 1982] Dosea

Peel energy (kJ/m2)

(kGy) 0 8 16 20 28 50 85 110 140 a

0.44 0.75 1.22 1.11 1.06 0.67 0.41 0.38 0.32

Mode of peeling

Cohesive Cohesive Cohesive Stick-slip Stick-slip Interfacial Interfacial Interfacial Interfacial

Rounded data from Figure 7 of Hamed [1982]

11.4.2.2

Butyl Elastomer Blends

11.4.2.2.1 PE/Butyl Elastomer Blends The usefulness of crosslinked PE is well known [Charlesby, 1960; Chapiro, 1962; Dole, 1972,1973; Silverman, 1992]. Unfortunately, crosslinked PE

is difficult to fuse, without special treatment. Irradiation of various blends of PE with butyl rubbers leads to crosslinking of PE and degradation of the butyl rubbers — bonding of these irradiated blends is no better than the bonding of crosslinked PE. However, bonding of irradiated PE blends with conjugated diene butyl polymers, CDB, was found to be excellent — they could be easily heat-sealed (Table 11.9). The optimum content of the C5 conjugated dienes in the butyl elastomer was 1-2%, and the useful range of the CBD was 30-70 wt%. Other ingredients, such as pigments, fillers, antioxidants and plasticizers, could also be added [Berejka, 1981]. CDB containing 1-2% conjugated diene acted as a crosslinking agent for PE (by participation of its C=C bonds in cationic and free radical reactions, Sections 11.2 and 11.3). Although neat PE requires 150 kGy dose, the blends require only 20-50 kGy for the same level of crosslinking. The irradiations were done using a 3-MeV electron accelerator. The compositions were prepared in a standard internal mixer (see Table 11.41). Conventional press and molds were used to produce slabs (1.9 mm) and thin film (0.13 mm) for irradiation and testing. Slabs of sample E

Table 11.41. Composition of PE/CDB blends [Berejka, 1981] Ingredients (wt%)

CDB Butyl 268 HDPE Irganox 1010 Chemlink 30 Microcel E CDB: Butyl 268: HDPE: Irganox 1010: Chemlink 30: Microcel E:

Composition A

B

C

D

E

F

70 30 -

50 50 -

30 70 -

49.5 49.5 1.0 -

50 50 -

49.4 49.4 0.9 0.3

Conjugated diene butyl polymer (C5 conjugated diene content 1.2-1.5 mol %; Exxon). Butyl rubber (isoprene/isobutylene copolymer). High-density polyethylene (density ~ 960 kg/m3) Hindered phenolic antioxidant (Ciba-Geigy). Trimethyolpropane trimethacrylate (TMPTA). Diatomaceous earth silica filler.

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A. Singh and K. Bahari

(Table 11.41) irradiated to a dose of ≥100 kGy, showed a phase separation, attributed to the chain breaks in the butyl rubber, BR, and immiscibility of BR with crosslinked HDPE [Berejka, 1981]. In the slabs of sample F at doses greater than 100 kGy, some immiscibility could be attributed to the presence of the TMPTA additive. These effects were not observed in the films because of predominant oxidative degradation in the thin films through reaction (R-22) and related reactions (Section 11.3.1)). The six samples (Table 11.41) were tested for ultrasonic bonding. Up to a dose of 50 kGy, samples A and E showed good bonding, whereas bonding of samples B and D was good up to a dose of 70 kGy. The physical properties of these blends are listed in Table 11.42. Evidently, the response of samples A, B, C and D to irradiation was excellent. Some physical properties of the irradiated samples were also determined at T = 140-175°C, viz. the thermal expansion and recovery, heat-aging properties, and effects of the carbon black addition. The blend comprising: 48 wt% HDPE, 48 wt% CDB, 1 wt% Irganox 1010 and 3 wt% carbon black, irradiated to 25 kGy, was found to be heat-sealable [Berejka, 1981]. 11.4.2.2.2 PP/CDB Blends Berejka and Bradley [1981; Bradley, 1983] also studied electron irradiation of PP/CDB blends, prepared as described in the preceding section (Table 11.9). The blends’ performance was in between those of PP and CDB. One would expect that irradiation will cause crosslinking of CDB and degradation of PP. The results are shown in Table 11.43. The modulus and tensile strength values of PP/CDB = 0/100, 20/80 and 50/50 are consistent with crosslinking. However, the increase of the MFI suggests that there is also a considerable degree of PP degradation (because of radiation-induced chain scission, see Section 11.3.2). 11.4.2.2.3 PP/Polybutylene Blends Rolando and Krueger [1992] developed blends of mesomorphous PP and polybutylene (PB), to overcome the deficiencies of crystalline PP (radia-

tion degradation), crystalline PB (post-irradiation degradation) and mesomorphous PP (poor tear and puncture resistance), for use in medical products (Table 11.9) [also see Rolando, 1993]. They also studied the radiation resistance of some mesomorphous PP/Exxelor PA-23 blends, for possible use in medical products. The blends did not require any stabilizers or antioxidants, for the required levels of radiation resistance, and for the maintenance of their post-irradiation mechanical properties. Their work also showed that when crystalline PP was used in the blends, in place of mesomorphous PP, the irradiated blends degraded during the post-irradiation period, whereas those made with mesomorphous PP did not. PP and PB were dry-blended, and the mixture was fed into a 5.1-cm (2 inch) Rheotec extruder. For some of the samples a Berlyn extruder was used. The extruder was operated at a flow rate of 9 kg/h with a screw speed of 20 rpm and staged temperature zones to ensure a melt temperature of 244°C. The blends were extruded into 4 mil (0.102 mm) films on a chrome-plated casting roll in which the quench temperature was controlled at 12.2°C, so that the structure of PP is mesomorphous. The materials used for the various samples, and the processing conditions used, are given in Table 11.44. It is important to cool the extruded form of the blends rapidly so that mesomorphous PP is obtained. Based on both the elongation at break and tensile strength measurements (Table 11.45), the blends containing 25 and 50 wt% PB appeared to be the best for use in medical products. The heatseal tear-strength measurements suggested that the blends containing 40-60% PB were satisfactory (Table II of Rolando and Kruger [1992]). The blends containing PB also showed better puncture resistance than mesomorphous PP or LDPE (Table IV, Rolando and Krueger [1992]). Blends of mesomorphous PP and PA-23, the elongation at break data (Table 11.45; see also Table V in Rolando and Krueger [1992]) shows that even at the high dose of 100 kGy, the blends containing 20-80% of PA-23 show good maintenance of properties, up to 12 months after irradia-

817

Use of High-Energy Radiation in Polymer Blends Technology

Table 11.42. Physical properties of PE/CDB blends at 23°C, AS a function of dose [Berejka, 1981] Sample

Dose (kGy)

Shore A hardness

Modulus 100% (MPa)

Modulus 300% (MPa)

Ultimate tensile strength (MPa)

Elongation (%)

A A A A

0 25 75 150

74 75 75 75

2.52 3.34 3.45 4.31

2.76 4.69 5.75 —

— 6.76 6.58 5.76

450 590 380 200

B B B B

0 25 75 150

94 94 94 94

6.65 7.62 7.10 7.83

— 8.14 8.58 9.86

6.83 12.48 11.65 11.41

300 630 500 375

C C C C

0 25 75 150

97 96 97 97

9.20 10.24 11.24 11.48

— 10.34 11.55 12.20

10.93 15.58 18.82 15.10

340 640 605 415

D D D D

0 25 75 150

95 94 94 94

6.31 6.96 6.69 7.34

6.48 7.24 7.45 8.76

6.79 12.82 12.14 10.79

500 720 600 420

E E E E

0 25 75 150

95 95 95 95

6.24 6.24 6.21 6.14

6.41 6.45 — —

7.79 6.31 6.38 6.21

700 250 230 165

F F F F

0 25 75 150

92 93 93 93

6.83 8.48 9.17 9.17

— — — —

8.10 9.38 10.24 10.20

210 175 170 165

tion (typically, the dose required for sterilization of medical devices and disposables is 25 kGy [Morrissey and Phillips, 1993]). Rolando and Krueger [1992] also studied the effects of grafting of various compounds, e.g., acrylic acid (AA), dimethyl acrylamide (DMA), N-vinyl-2-pyrrolidone (NVP) and a copolymer of NVP and TMPTA. The grafting was done by electron irradiation (0.175 MeV) under nitrogen.

They found that the grafted blends do not suffer degradation because of electron beam irradiation. The grafted blends maintained their structural integrity, with the advantage of improved surface properties. 11.4.2.2.4 Polybutylene/CDB Blends Berejka and Bradley [1981] reported that blends of polybutylene, PB, and a CDB (Table 11.9) can

818

A. Singh and K. Bahari

Table 11.43. Electron irradiation of polypropylene (PP) and conjugated diene butyl polymer (CDB) [Berejka and Bradley, 1981] PP/CDB

Dose

Modulusa at 100%

(kGy)

elongation (MPa)

Ultimate tensile strengtha (MPa)

Elongation (%)

(%)

Melt index (g/10 min) at 238°C, 5 kg

Set (%)

0/100 0/100 0/100 0/100 0/100

0 10 20 30 40

1.8 1.8 1.9 2.2 2.8

4.1 4.5 5.2

600 650 450

0 0 0 0 0

0-3 0-3 0-3 0-3 0-3

20/80 20/80 20/80 20/80 20/80

0 10 20 30 40

2.0 2.2 2.7 2.3 2.2

2.1 3.6 5.2 4.7 5.5

300 500 725 600 620

12.0 11.6 10.8 13.0 18.4

42 35 30 20 10

50/50 50/50 50/50 50/50 50/50

0 10 20 30 40

8.7 8.2 10.2 9.8 8.8

8.7 8.4 8.8 10.2 9.9

115 230 130 300 175

5.2 3.0 12.0 12.5 15.4

100 100 100 100 100

80/20 80/20 80/20 80/20 80/20

0 10 20 30 40

0 0 0 0 0

21.0 21.6 23.3 21.0 21.6

70 70 60 30 55

32.0 54.0 72.0 96.0 92.0

100 100 100 100 100

100/0 100/0 100/0 100/0

10 20 30 40

0 0 0 0

34.1 32.9 33.7 35.1

42 25 35 28

58 100 124 208

100 100 100 100

a

Values of the modulus and the tensile strength have been rounded off to the first decimal place.

be electron-cured to give products with excellent resistance to set, good strength properties, and adequate melt flow properties to be re-workable after irradiation as a thermoplastic. The blends were made by mixing PB and CDB with Irganox 1010, in a conventional rubber mill, internal mixer, extruder, kneader, or other blender, at 149 to 169°C. The blend could then be electron-cured either before or after forming it into the desired

shape. The irradiations were done using a 3-MeV electron accelerator in the dose range of 20 to 40 kGy. Typical examples of the products made from these blends include athletic shoe soles, electrical connectors, sealing tapes and molded automobile bumper components. The blends comprising: 20-70 wt% PB and 80-30 wt% CDB were compression-molded (177°C, under 14 MPa) into 1.5-mm thick slabs. Their values of modulus

819

Use of High-Energy Radiation in Polymer Blends Technology

Table 11.44. Processing conditions of polypropylene-based blends [Rolando and Krueger, 1992] Sample No.

1 2-4 5 6-10 11-15 16-19 20-22 23-25

Material

PBa PPb/PBa PPc PPb/PBa PPb/LDPd PPe/PA-23f PPb/PBa PPb/PA-23f

Flow rate (kg/h)

Screw speed (rpm)

Melting temp. (°C)

Film thickness (mm)

9 9 5 9 * 9 ** ***

20 20 40 37 * 20 ** ***

244 244 203 221 * 232 ** ***

0.10 0.10 0.04 0.09 * 0.10 ** ***

Quenching temp. (°C) 12.2 12.2 6.7 0 * ≤ 22 55 57.2

a

Shell 0400, melt index = 20 Coseden 8771, melt index = 9 c Exxon PP-3014, melt index = 12 d Dowlex 752 e Fina 8771, melt index = 9 f Exxelor PA-23 from Exxon * Prepared in the same manner as the samples 1-5 ** Prepared in the same manner as the samples 1-4 *** Prepared in the same manner as the samples 16-19 b

at 100% elongation, ultimate tensile strength, elongation, MFI and set are given in Table 11.46. 11.4.2.3

Elastomers

11.4.2.3.1 PE/CR Blends Coran and Patel [1982] patented a process for electron irradiation, of PE (10-75 wt%) blends with polychloroprene, CR (90-25 wt%) (Table 11.9). Irradiation improved such properties as tensile strength, stress at break, tension set, toughness, solvent resistance and high-temperature behavior. The improvements were attributed to crosslinking of both PE and CR. The blends were useful for making such products as tires, hoses, belts, gaskets and moldings. The blends were prepared by charging the required proportions of PE and CR to an internal mixer, and masticating the mixture at 190°C for 3 min after the PE had melted, then remasticating the mixture for 2 min. Sheets 3 to

3.5-mm-thick were irradiated with a Dynamitron electron accelerator, at doses up to 150 kGy. The irradiated sheets were masticated again in an internal mixer, removed, cooled and compressionmolded at 225°C. Some of the reported properties of the blends are listed in Table 11.47. Blends having low PE content exhibit better tension set, whereas those with higher PE content show higher stress-strain properties, including stress at break. It is important that the elastomer be dispersed as small particles (d ~ 50 mm), otherwise the composition will either be weak or not processable as a thermoplastic. The properties may be modified by the addition of carbon black or silica. Radiation crosslinking promoters are beneficial, particularly at low radiation doses, 10-50 kGy. The compositions containing crosslinking promoters were irradiated with a dose of 10 kGy. The data show that the additives significantly improved the properties. The amounts of the crosslinking promoters were 1-10 phr by weight

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Table 11.45. Mechanical properties of PB/PP and PA-23/PP blends [Rolando and Krueger, 1992] Property

Time after irradiation (months)

Elongation at break (%)

Elongation at break (%)

Yield tensile strength (MPa)

Blend composition (wt%)

PP Type

PB

PP

0 9 0

100 100 75

0 0 25

Ma

9 0 9 0 9 0 6b 0 9 0 9

75 50 50 25 25 0 0 50 50 25 25

25 50 50 75 75 100 100 50 50 75 75

PA-23

PP

80 80 60 60 40 40 20 20 80 80 40 40 20 20

20 20 40 40 60 60 80 80 20 20 60 60 80 80

PB

PP

100 100 75 75 50

0 0 25 25 50

0 12 0 12 0 12 0 12 0 12 0 12 0 12

0 9 0 9 0

Dose (kGy)

0

30

60

200 200

200 125

200 0

M M M M M M M Cc C C C

270 270 370 390 500 500 420 >400 400 380 500 500

270 270 370 340 500 500 >400 >400 400 270 500 420

270 270 370 410 500 450 — — 400 100 500 15

M M M M M M M M C C C C C C

669 700 649 693 622 630 580 590 681 662 628 632 642 635

637d 683 599 667 620 613 560 605 440e — 575 40 610 65

612d 758 594 745 590 694 563 585 410e 10 540 20 580 21

M M M M M

— — — — —

— — — — —

30.3 — 15.2 14.4 14.2

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Use of High-Energy Radiation in Polymer Blends Technology

Table 11.45. Continued Property

Time after irradiation (months)

Blend composition (wt%) PB

PP

PP Type

Dose (kGy)

0

30

60

Yield tensile strength (MPa)

9 0 9 0 9

50 25 25 0 0

50 75 75 100 100

M M M M M

— — — — —

— — — — —

15.0 19,8 20.9 17.2 17.9

Break tensile strength (MPa)

0 9 0 9 0 9 0 9 0 9

100 100 75 75 50 50 25 25 0 0

0 0 25 25 50 50 75 75 100 100

M M M M M M M M M M

— — — — — — — — — —

— — — — — — — — — —

31.7 — 31.0 30.2 27.9 26.9 24.1 23.6 18.3 19.6

a b c d e

M-Mesomorphous Values for 9 months not reported C-Crystalline The values of the irradiated samples are for 1 month after irradiation, instead of 0 month The values of the irradiated samples are for 2 months after irradiation, instead of 0 month.

Table 11.46. Electron irradiation of polybutyiene, PB, and conjugated diene butyl polymer, CDB [Berejka and Bradley, 1981] Blend

Dose

Modulusa

PB/CDB (wt%)

(kGy)

at 100% elongation (MPa)

Ultimate tensile strengtha (MPa)

Elongation (%)

Melt index (g/10 min) at 238°C, 5 kg

Set (%)

0/100 0/100 0/100 0/100 0/100

0 10 20 30 40

1.8 1.8 1.9 2.2 2.8

4.1 — — 4.5 5.2

600 — — 650 450

0 0 0 0 0

0-3 0-3 0-3 0-3 0-3

822

A. Singh and K. Bahari

Table 11.46. Continued Blend

Dose

Modulusa

PB/CDB (wt%)

(kGy)

at 100% elongation (MPa)

Ultimate tensile strengtha (MPa)

Elongation (%)

20/80 20/80 20/80 20/80 20/80

0 10 20 30 40

0.7 1.1 1.2 1.1 1.1

1.1 3.1 4.5 4.8 4.2

770 735 565 600 600

0.5 0.013 0.0046 0 0

45 34 28 25 25

30/70 30/70 30/70 30/70 30/70

0 10 20 30 40

1.8 2.0 2.0 2.1 1.8

2.9 4.1 4.9 5.6 4.7

770 770 770 745 700

0.58 0.02 0 0 0

35 40 40 40 35

40/60 40/60 40/60 40/60 40/60

0 10 20 30 40

2.5 3.3 3.2 3.3 3.7

7.3 7.8 8.0 8.2 10.3

770 770 770 770 680

0.6 0.3 0.4 0.5 3.13

52 48 45 45 42

50/50 50/50 50/50 50/50 50/50

0 10 20 30 40

4.4 5.0 4.8 5.4 5.4

8.5 10.1 10.3 12.5 12.5

770 600 770 705 640

0.5 0.3 0.4 1.28 1.20

58 54 50 50 50

60/40 60/40 60/40 60/40 60/40

0 10 20 30 40

2.8 5.8 5.8 5.8 5.9

12.8 13.3 13.1 14.8 13.1

770 770 770 760 720

2.8 1.0 1.4 2.0 3.6

60 60 60 60 60

70/30 70/30 70/30 70/30 70/30

0 30 20 30 40

7.3 6.2 7.4 7.3 7.3

18.9 20.0 18.6 17.4 18.5

770 770 770 770 750

4.5 3.6 4.9 10.2 13.0

60 62 68 70 75

100/0 100/0 100/0 100/0 100/0

0 10 20 30 40

2.2 2.1 1.7 1.9 1.4

24.8 — — 22.8 24.8

330 — — 420 430

40.0 — — 49.0 100.0

99-100 99-100 99-100 99-100 99-100

a

Values of the modulus and tensile strength have been rounded off to the first decimal place.

Melt index (g/10 min) at 238°C, 5 kg

Set (%)

823

Use of High-Energy Radiation in Polymer Blends Technology

Table 11.47. Properties of electron irradiated and chemically crosslinked PE/CR = 1:1 blenda [Coran and Patel, 1982].

a b

Sample

Dose (kGy)

Tensile strength (MPa)

Ultimate elongation (%)

Tension set (%)

1 2 3 4 5

0 10 50 100

11.8 14.0 17.5 17.5

560 530 450 380

71 64 62 53

b

15.9

320

50

by weight Chemically crosslinked with m-phenylene bismaleimide (2 phr), benzothiazyne disulfide (0.5 phr) and ZnO (2.5 phr).

Table 11.48. Properties of electron irradiated (10 kGy) PE/CR = 1:1 blenda containing trimethylolpropane triacrylate (TMPTA) [Coran and Patel, 1982]. Sample

TMPTA (phr of PE+CR)

Tensile strength (MPa)

Ultimate elongation (%)

Tension Set (%)

1 2 3 4

0 2.0 3.8 7.4

11.8 13.4 16.0 14.8

360b 350 440 390

64 65 58 58

a b

by weight Patent does not state why this value is not the same as that given in Table 12.38.

of PE + CR, and included TMPTMA, TAC, polyethylene glycol dimethacrylate (PEGDMA) and triallyl phosphate (TAP). The data obtained in the presence of TMPTMA are shown in Table 11.48. Coran and Patel [1982] suggested that the effects would be similar for electron, gamma, and X-ray irradiation. However, as discussed elsewhere in this chapter, the probability of oxidative degradation would increase as the dose rate is decreased [Kashiwabara and Seguchi, 1992]. Therefore, the improvements in the properties of the blends reported in this patent may not be the same with gamma or X-ray irradiation.

The authors also suggested that because of the limited penetration of electrons, it would be difficult to irradiate samples thicker than 30 mm with electrons. However, nowadays 100-mmthick samples of most polymers can be easily irradiated using a 10-MeV electron accelerator [Woods and Pikaev, 1994; Singh et al., 1996]. 11.4.2.3.2 HDPE/NR Blends Akhtar et al. [1986] investigated gamma irradiation of HDPE/NR blends (Table 11.9). NR and HDPE were blended in an internal mixer with a cam-type mixer rotor, at 60 rpm, and T = 150°C. The HDPE was melted for 2 min and then the

824

A. Singh and K. Bahari

NR was added; mixing continued for 4 min when DCP was added, and the system was mixed for additional 2 min. The blends were sheeted using a laboratory mill at 1.25-mm nip setting. The sheets were cut into small pieces and re-mixed at 150°C for 1 min and sheeted out again in the laboratory mill. The blended sheets were compression-molded into the dumbbell specimens (2 ± 0.2 mm) in an electrically heated hydraulic press at 150°C for 3 min. The specimens under compression were quenched in cold water (which should keep the crystallinity at a low level). The irradiations were done at T = 25°C, in the γ-chamber-900 at BARC, Mumbai. The irradiated HDPE/NR blends (30-70 wt% HDPE) showed a broad minima in tensile strength vs. dose, at 100-250 kGy. According to the authors [Akhtar et al., 1986] irradiation had two effects on the mechanical properties of the blends: (i) it produced crosslinks between carbon atoms, which increased the stiffness; and (ii) it reduced crystallinity; they claimed that high radiation doses ultimately eliminate the crystallinity leaving a flexible polymer with good elastic recovery. However, no data on the crystallinity are presented. It is doubtful that at the doses used by the authors, the crystallinity of HDPE would decline drastically. As the radiation dose is further increased, irradiated HDPE becomes stiffer and more brittle, when the degree of crosslinking increases with no compensating reduction in the degree of crystallinity [Akhtar et al., 1986]. Thus, after 150 kGy, the tensile strength of HDPE is negligible. The degree of orientation of the molecular chains of the two polymers is hindered by crosslinks, leading to a decrease in tensile strength. The authors claim that when the blends were exposed to gamma radiation at the higher dose rate, they underwent mainly crosslinking and, as the dose increased so did the extent of crosslinking, resulting in a brittle behavior. However, the difference in the two dose rates used was too small to warrant any dose rate-dependent effects. The two dose rates (1.47 and 2.1 kGy/h) had almost the same effect on the tensile strength of the blends. Up to 10 kGy there was little change, and then the tensile strength started to decrease.

Blends rich in HDPE showed greater reduction of the tensile strength. The blend containing 70 wt% NR showed an increase in the ultimate elongation at the doses below 10 kGy — at higher values there was a continuous decline in the ultimate elongation. Akhtar et al. [1986] claim that the small number of crosslinks produced by the addition of DCP had a negligible effect on the radiation resistance of the blend. Up to 10 kGy, at a dose rate of 2.1 kGy/h, the energy rupture values remained constant, and a broad minima in the tensile strength was observed in the range of 40-150 kGy. In the presence of DCP, the initial tensile strength of the samples is higher, presumably because of crosslinking during the blending and molding processes (thermally initiated by the peroxide, analogues of reactions (R-23) and (R-37)). On irradiation the tensile strength follows the same pattern as for the samples that do not contain the DCP, except that the tensile strength remains higher than the samples without the DCP, up to a dose of about 300 kGy. In comparison, the addition of the carbon black does not affect the tensile strength of the non-irradiated samples. However, it does increase the radiation resistance of the samples (Table 11.49). All the high rubber blends show a similar pattern of change in the energy at rupture values; a minima in the values is observed in the dose range 40-150 kGy, as in the case of the tensile strength data (Table 11.49). The energy at rupture increases at higher doses (150-250 kGy), and then decreases again as the dose increases further. The authors relate the minimum in the energy at rupture values with the phenomenon of brittle distortion. However, at least partially it could be related to the dose required to crosslink NR, which may well be quite a bit higher than that for HDPE. On irradiation in nitrogen NR undergoes crosslinking (also see Section 11.4.2.3.8); however, the crosslinking is inhibited by oxygen [Bradley, 1984]. Under nitrogen atmosphere, at a dose rate of 2.1 kGy/h, the samples undergo crosslinking in the bulk, and the chain scission is minimal. The tensile strength slightly increased because

825

Use of High-Energy Radiation in Polymer Blends Technology

Table 11.49. Effect of gamma irradiation on HDPE/NR blendsa [Akhtar et al., 1986] Blend composition (wt%) HDPEb

Additives

Dose (kGy) DCPe

Ultimate tensile

Elongation at Break

strength (MPa)

(%)

NRc

CBd

70 70 70 70 70

30 30 30 30 30

-

-

0 10 100 150 250

13.5 12.7 3.7 3.2 3.7

47 39 0 0 0

70 70 70 70 70

30 30 30 30 30

-

1 1 1 1 1

0 10 150 250 250

nr nr nr nr nr

20 0 0 0 0

50 50 50 50 50

50 50 50 50 50

-

-

0 10 100 150 250

9.0 8.1 5.3 4.7 5.7

400 231 50 36 33

30 30 30 30 30

70 70 70 70 70

-

-

0 10 100 150 250

8.4 7.9 3.3 3.4 5.0

638 639 458 320 309

30 30 30 30 30

70 70 70 70 70

-

1 1 1 1 1

0 10 100 150 250

13.0 11.0 4.3 3.9 5.2

501 477 322 253 174

30 30 30 30 30

70 70 70 70 70

40 40 40 40 40

-

0 10 100 150 250

13.7 7.9 5.7 5.7 5.8

466 489 247 312 154

30 30 30 30 30

70 70 70 70 70

40 40 40 40 40

1 1 1 1 1

0 10 100 150 250

12.8 9.3 6.3 6.3 6.8

300 306 219 205 170

826

A. Singh and K. Bahari

Table 11.49. Continued Blend composition (wt%) HDPEb

a

b c d e f

nr

Additives

Dose (kGy) DCPe

Ultimate tensile

Elongation at break

strength (MPa)

(%)

NRc

CBd

100 100 100 100 100

-

-

-

0 10 100 150 250

32.3 13.2 6.0 4.2 2.2

nr nr nr nr nr

30f 30f 30f 30f 30f

70f 70f 70f 70f 70f

-

-

0 10 100 150 250

8.4 8.1 10.6 12.8 14.6

638 576 549 450 440

Irradiations in air, except for one sample (see below). Minimum values of tensile strength at ~150 kGy, except for HDPE, which declines sharply from 10 to 250 kGy. Hostalen GA7260 from Polyolefin Industries Ltd., Bombay. Crumb rubber, ISNR 5 grade from Rubber Research Institute of India, Kottayam, Kerala. Carbon black: high abrasion furnace black (N330) from Phillips Carbon Black Limited. Dicumyl peroxide (40%) from Bengal Waterproof Ltd., Panihati, West Bengal. DCP loading given relative to the rubber phase. Sample irradiated under nitrogen; minimum value of tensile strength, 5.8 MPa at 40 kGy, after which the value rises sharply to 250 kGy. Not reported

of the onset of crosslinking at a small dose but decreased as the crosslinking in bulk increased at higher doses. Under nitrogen these phenomena were observed at a lower dose, 40 kGy, than for irradiation in air. The ultimate elongation at the dose rate of 2.1 kGy/h was found to decrease with dose. The blends with high rubber content had higher ultimate elongation, compared to blends with high HDPE content. Up to the dose of 10 kGy, the ultimate elongation of the blends remained constant. The initial increase in elongation at break was followed by a gradual decrease, caused by increased crosslinking. The dependence was similar when the samples were irradiated in nitrogen. The SEM data of the tensile fracture surface of the high rubber blend show a rough surface

with few tear lines. The SEM data of the same blend after irradiation to a dose of 250 kGy show a number of tear lines curving inward from the surface and joining the straight horizontal fracture path extending from one end of the surface to the other. Further degradation is seen at a dose of 1000 kGy. These inward curving lines suggest that their origin lies on the surface of the sample, implying thereby extensive degradation in the form of chain scissions on the surface [Akhtar et al., 1986]. The SEM data in the presence of DCP (1000 kGy) are similar. In contrast, the SEM data on the sample irradiated under nitrogen do not give any indication of degradation at 250 kGy, which confirms that the tear lines seen in the samples irradiated in air are due to oxidative degradation. In the work on irradiation there would

Use of High-Energy Radiation in Polymer Blends Technology

be DCP-sensitized crosslinking [Singh, 1995], as discussed in sections 11.4.1.2 and 11.4.1.3. 11.4.2.3.3 PE/Elastomer Blends Ivchenko et al. [1981] investigated the possibility of producing elastic films and coatings for synthetic leathers by introducing elastomers into PE (Table 11.9), followed by irradiation to increase their strength and heat resistance, and prevent ply separation (delamination). Six different elastomers (butyl rubber, IIR, chlorinated-PE, CPE, chlorsulphonated-PE, CSR, polyisobutylene, PIB, and two polybutadiene rubbers, BR, were added to PE. Films were made on a roll mill at 105°C, and were then pressed at 130°C. The samples were irradiated at 20°C with electrons (dose rate = 0.4 kGy/s). The blends containing PIB, IIR and the two BR were poor, either because of adhesion to the processing equipment, or poor performance. The blends with either CPE or CSR were found to be satisfactory. The results for the CSR/PE blends are summarized in Table 11.50. The optimum composition of CSR/PE was 60:40 (by weight); on irradiation the blend crosslinked, its tensile strength and modulus increased, while the elongation at break and rigidity decreased. The heat tolerance of the blend films also increased by about 30°C, on irradiation. Addition of a plasticizer, unsaturated oligoesteracrylate TGM-3, facilitated the processing. The plasticizer was crosslinked on exposure to heat and electrons. The optimum properties were obtained for the composition CSR/PE/TGM-3 = 60:40:10. The composition was also satisfactory as far as the light- and heataging behavior was concerned. 11.4.2.3.4 PVC/TPE Blends Thomas et al. [1986] studied the effect of gamma irradiation (1, 10, 100 and 500 kGy) on PVC/TPE = 100:0, 75:25, 50:50, 25:75 and 0:100 blends (Table 11.9). As TPE a thermoplastic, elastomeric polyether-ester block copolymer of 1,4-butanediol-polybutylene and glycol-terephthalic acid copolymer (Hytrel™), was used. The blends were prepared in an internal cam-type mixer at 80 rpm, 180°C, mixing time 7 min. TPE was added first,

827

followed by PVC. The blends were sheeted in a laboratory mill with a 2.00-mm nip. To ascertain uniformity, the sheets were cut into small pieces and re-mixed for 1 min, then sheeted again. The sheets were compression-molded at 200°C for 3 min, then quenched under pressure. At a dose greater than 1 kGy the PVC blends developed color, increasing in intensity with increasing PVC content and dose, from pale yellow to dark brown. At a dose of 500 kGy, the blend containing 75 wt% TPE showed surface cracks. At 1 kGy, the tensile strength of the three blends (TPE content 25-75%) first slightly decreased, at 10 kGy returned to the initial value, and then decreased sharply at 500 kGy. The sharp decrease was caused by the embrittlement of the TPE phase because of crosslinking. It may be noted that the oxidative degradation of the blend as well as dehydrochlorination of the PVC, both of which would be expected to occur in this case (see Section 11.3), could also cause a decrease in the tensile strength. The modulus at 100% elongation first decreased and then increased. In the case of 75% TPE, the value at 500 kGy was about 20% lower than that of the non-irradiated sample. The two other blends at 500 kGy showed slightly higher moduli than those for the corresponding non-irradiated samples. The results for elongation at break, tear strength, tensile set after failure, and energy at rupture in tensile failure, showed a similar pattern — the values remained the same up to 10 kGy after which they sharply declined with dose increasing to 500 kGy. SEM showed that irradiation to 500 kGy of the TPE 75% blend resulted in surface cracks, though these were smaller than those in neat TPE specimen. Decreasing the TPE concentration progressively eliminated the surface cracks; instead surface foldings and corrugations, associated with the PVC degradation on irradiation, were observed. The cracks originated in crosslinking and surface oxidation of TPE. The tensile and tear fracture surfaces were intermediate between those of TPE and PVC.

828

A. Singh and K. Bahari

Table 11.50. Effect of electron irradiation on PE/CSR blends [Ivchenko et al., 1981]. Blend composition (wt%

Dose (kGy)

Tensile strength (kPa)

Elongation at Break (%)

Gel Fraction (%)

PE

CSR

100 100 100 70 70 70 40 40 40 40 40 30 30 30 40a

100 100 100 30 30 30 60 60 60 60 60 70 70 70

10 10 -

0 50 100 0 50 100 0 50 100 0 50 100 0 100 0 50 100

140 160 175 40 88 68 112 135 160 85 120 125 55 110 145

650 660 660 2000 980 700 720 660 540 700 600 480 780 715 600

1 45 1 65 1 56 1 65 -

60 60 60 60 60 60 60 60 60 60 60 60

5 5 5 10 10 10 15 15 15

0 50 100 0 50 100 0 50 100 0 50 100

62 100 120 65 105 120 53 130 145 40 150 185

300 460 370 120 340 340 120 340 390 100 410 390

-

40a 40a 40a 40a 40a 40a 40a 40a 40a 40 40

TGM-3

a Appears to be from a different batch since the mechanical properties are different from those given earlier in the Table for 40 wt% PE blend.

11.4.2.3.5 TPU/POM Blends Thermoplastic polyurethane, TPU, and polyoxymethylene, POM were mixed in the desired ratios using a 30-mm single-screw extruder (L/D = 20) with a mixing head, at 170-200°C (Table 11.9).

The test specimens for the HDT and tensile strength measurements were prepared by using an injection molding machine at 180-200°C. To relieve any residual stresses, the specimens for HDT were

Use of High-Energy Radiation in Polymer Blends Technology

annealed at 70°C. The tensile strength specimens were gamma-irradiated to a dose of 25 kGy. Irradiation reduced the yield stress and elongation at break but increased the elastic modulus. The yield stress and elongation decreased because of oxidative degradation, but the slight rise in the value of the modulus could not be explained. On irradiation, the blends developed yellow coloration, but less intensive than that of the neat POM [Kumar et al., 1993]. 11.4.2.3.6 PP/Elastomer Blends Magerramov et al. [1981] studied the effect of gamma irradiation on the electrical strength (Est) of the blends of iPP with Polymethylvinylsiloxane (SKTV-1) and ethylene-propylene elastomers (SKEP-60) (Table 11.9). The blends were made by mixing the components in a blade mill at 200-220°C for 20 min, followed by granulation. The elastomer content of the blends varied from 0.5 to 15 wt%. Films of iPP and the blends (0.1 mm) were prepared by compression molding at 220°C, under 10 MPa pressure, between sheets of aluminum foil. The films were treated as follows: (A) quenching of the film in the aluminum foil by ice-water mixture; (B) annealing of the film produced in A for 5 h at a temperature of 5-10°C below its melting point; and (C) cooling the film to 20°C slowly (2°C/min) in the mold. Optical examination of the films showed that the films treated as A and B did not contain any iPP spherulites; however, spherulites were present in the slowly cooled films (treatment C). The addition of the elastomers to iPP resulted in the reduction of the size of the spherulites by about a factor of 2. The slowly cooled films (treatment C) showed higher degrees of crystallinity and densities, than the quenched films (treatments A and B). The Est of the films decreased with increasing crystallinity. However, the presence of the elastomers in the films contributed to the increase of the Est values in some of them. The authors attributed the increase to two structure-forming roles of the elastomers: (i) reduction in the size of the spherulites, and (ii) prevention of the formation of microcracks by being located preferentially in the interspherulitic regions.

829

The films were gamma-irradiated to doses of 50-1000 kGy, with a PX-γ-30 device either in air or under water at room temperature (details not given; the environment under water would be close to an inert atmosphere, but some of the oxygen dissolved in water would still be available for reactions with the free radicals formed in the film, e.g., reaction (R-22). Futhermore, the free radicals formed in water (such as H• and •OH) would be able to react with the surface of the films which could result in a coating of oxygenated groups (formed via reactions of the type (R-8), (R-22), and (R-26) to (R-29)). Except for few films that showed small increases in their Est values at 50 kGy, in general, the Est values of all the films decrease with increasing dose (up to 1000 kGy). The Est values for all the films are generally lower on irradiation in air, as compared to irradiation under water. Apparently, radiolytic oxidation of the films on irradiation in air (e.g., reaction (R-22)) lowers their Est values. The films produced by treatment C are more sensitive to the dose-dependent reduction of their Est values, than the films produced by treatments A or B. The slowest dose-dependent decrease of the Est values is seen in the case of the iPP/SKEP-60 blends. Some of these trends are shown by the data on the D30 values (the dose required to reduce the Est values by 30%) for the various films, given in Table 12.51. 11.4.2.3.7 EVAc/PP Blends Minkova and Nikolova [1989 a,b] investigated the effect of electron irradiation, on the thermodynamic and mechanical properties of thin films (100 µm) of EVAc/iPP (Table 11.9) (see details given earlier in Section 11.4.1.5). They observed segregational melting of the blend components. The peak melting temperature of the EVAc and iPP, in the blends, did not depend on the blend composition. The temperature of melting of the two polymers, whether pure or in the blends, decreased with increasing dose (Table 11.16).

830

A. Singh and K. Bahari

Table 11.51. Dose required to reduce the EST values of the blend films by 30% [Magerramov et al., 1981]

1 2 3 4 5

Film

Irradiation

treatment

environment

D30 (kGy) PP

1

2

3

4

5

A B C

Water Water Water

85 50 10

70 35 30

60 30 25

20 15 15

90 75 45

85 70 40

A B C

Air Air Air

30 20 10

35 20 15

25 20 15

20 10 10

55 60 35

50 50 30

PP + SKTV-1 (0.5 wt%) PP + SKTV-1 (1.5 wt%) PP + SKTV-1 (15 wt%) PP + SKEP-60 (3.5 wt%) PP + SKEP-60 (10 wt%)

The authors reported rather complex behavior for the enthalpies of melting of the blends, as follows: (i) The enthalpies of melting for pure EVAc and for EVAc in the blend containing 40 wt% EVAc decreased with increasing dose, linearly up to 100 kGy and then more rapidly to 160 kGy; on the other hand, in the blends containing 50-90% EVAc, the EVAc enthalpies of melting first increased almost linearly to 100 kGy and then decreased sharply to 160 kGy (Table 11.16). (ii) The enthalpy of melting of pure iPP decreases with increasing dose. In the case of the blend with 60% iPP (EVAc/iPP) the decrease is sharper, and the values are lower than for pure iPP. However, as the proportion of EVAc increases, the values of the enthalpy of melting for iPP increase and the shape of the melting curve approaches that of pure EVAc. Minkova and Nikolova [1989a,b] attributed the observed effects to crosslinking in the amorphous regions and crystallite surfaces of EVAc and scission in iPP. They also suggested that low levels of iPP in the blends (10%) sensitize scis-

sion in EVAc. It should be realized that even though electron irradiation was used, based on the half-value layer data for oxidation, given by Kashiwabara and Seguchi [1992], 10% or more of the film would undergo oxidative degradation. The effect of the presence of the unspecified stabilizer is also difficult to predict. It is possible that the change in the trends of the effects at 100 kGy dose is related to possible depletion of the 0.15% stabilizer. Even though the compositions of the blends varied from 30-70% of EVAc, and the rest of iPP (plus 0.15% stabilizer), the gel content remained between 50-63% at doses from 50 to 120 kGy. Unfortunately, detailed data are not given. It is obvious that as crosslinking of EVAc would increase with dose, scission of iPP will increase. There would also be a complex role played by the stabilizer; the 50 kGy dose lower limit may well be the dose where the stabilizer gets used up. Further, these films could develop about 10-20% oxidation layer on their surface, with the electron irradiation used in this work (based on the data given by Kashiwabara and Seguchi, [1992]). Minkova and Nikolova [1989a,b] also report that the gel content

Use of High-Energy Radiation in Polymer Blends Technology

is insensitive to the blend compositions (this is surprising; with increasing iPP, one would expect increasing levels of chain scission). In the EVAc/iPP blends E (elastic modulus, which is proportional to crosslink density) increases with increasing dose. The blend composition and irradiation have rather small influence on the tensile strength; however, the elongation at break decreases with increasing dose (Table 11.17). Thomas et al., [1987] studied the effect of gamma irradiation on the stress-strain behavior, elongation at break, energy at rupture, hardness, and morphology of EVAc/iPP blends (Table 11.9). The blending was done in an internal mixer, using a cam-type mixer with a rotor speed of 80 rpm, and the mixer temperature of 180°C, with a mixing time of 7 min. The iPP was first melted in the mixer followed by the addition of the EVAc, and the mixing continued for 2 min; the required curatives were then added and the mixing continued for another 4 min. The blend was then sheeted through a laboratory mill at a 2.00-mm nip setting. The sheeted blend was cut into small pieces and again mixed for 1 min, and sheeted through the mill. The sheeted blend was then compression-molded at 200°C for 3 min, in a specially designed mold such that the mold containing the sample can be cooled immediately after molding, while the sample is still under pressure. This would keep the crystallinity at a low level. The samples for tensile strength tests were gamma-irradiated at room temperature (dose rate 3.21 kGy/h) in air to total doses of 10, 100, 250, 500 or 1000 kGy. EVAc samples and the blend containing 70 wt% EVAc developed brown color at the higher doses (500 and 1000 kGy). The iPP samples developed microcracks at these higher doses. All the other samples (irradiated and non-irradiated) remained colorless. The stress-strain curve of non-irradiated iPP is typical of high modulus brittle materials [Thomas et al., 1987]. At the doses of 100 and 1000 kGy, iPP shows brittle behavior. The non-irradiated blend containing 30% EVAc shows necking rupture, whereas the dynamically crosslinked sample (with DCP) shows no neck formation and yielding. On irradiation to 100 kGy, the necking

831

and yielding behavior of these two samples is replaced by brittle behavior. In contrast, the blend containing 70% EVAc shows an initial increase in stress, followed by a decrease until failure occurs. The stress-strain curve of the irradiated sample is quite different; the sample shows very high elongation (625%). The dynamically crosslinked (with DCP) sample containing 70% EVAc shows elastic type failure; on irradiation to 100 kGy, the behavior is similar with slightly more elongation at break. EVAc itself shows very low values of stress up to 600% elongation, followed by small increase at higher strains (elongations greater than 600%). This is attributed by the authors to the strain-induced crystallization of the PE segments of EVAc. The behavior is similar on irradiation to 100 kGy, but the elongation is considerably reduced. On irradiation to 1000 kGy, all the samples show brittle behavior. The tensile strength of the blends decreases with decreasing iPP content. The tensile strength of iPP decreases sharply with increasing dose, beyond 10 kGy. In the blends containing 30 and 50% EVAc, the tensile strength gradually decreases as the dose increases to 500 kGy. The values of the dynamically crosslinked blends are lower up to 500 kGy. In the case of EVAc, the values are similar up to 100 kGy, after which they decrease and then level out. The tensile strength data are briefly given in Table 11.52. The elongation at break also decreases with increasing iPP content of the blends. The elongation values decrease with increasing dose. In the case of the samples containing 70% EVAc, a maximum is reached at 100 kGy, followed by decrease of the values (Table 11.52). The values for energy at rupture in tensile failure generally decrease with increasing dose. In the case of the EVAc-rich blends, the high elongation values led to initial increase in the energy at rupture, whereas in the case of EVAc the energy at rupture could not be measured up to a dose of 250 kGy because of the high elongation values. The hardness of most of the samples was unaffected by irradiation. In the case of EVAc (around 10 kGy) and the blend containing 70% EVAc (up to 500 kGy), the hardness values showed some increase.

832

A. Singh and K. Bahari

Table 11.52. Effect of irradiation on the tensile strength and elongation at break of PP/EVAc Blends [Thomas et al., 1987] Blend PP

Composition EVAc (vol%)

Dose (kGy)

Tensile strength (MPa)

Elongation at break (%)

100 100 100 100 100 70 70 70 70 70 50 50 50 50 50 30 30 30 30 30 0 0 0 0 0

30 30 30 30 30 50 50 50 50 50 70 70 70 70 70 100 100 100 100 100

0 100 250 500 1000 0 100 250 500 1000 0 100 250 500 1000 0 100 250 500 1000 0 100 250 500 1000

42.2 16.8 11.3 3.6 4.7 21.3 15.0 13.0 8.1 6.9 13.7 6.7 4.9 4.2 3.9 2.5 3.2 3.4 3.4 2.8 3.9 4.3 1.4 1.4 1.3

22 9 1 3 3 100 26 25 15 12 100 26 26 28 15 100 625 175 130 55 1500 650 150 130 50

SEM examination shows continuous brittle cracks on the surface of iPP irradiated to 1000 kGy; these cracks are attributable to chain scission and oxidative degradation of the iPP (see Section 11.3.2). In comparison, the surface morphology of the irradiated blend (1000 kGy) containing 50% iPP is quite different; it shows mainly holes and cavities associated with the detachment of the dispersed rubber particles from the continuous iPP matrix, along with minute discontinuous cracks. The irradiated surface of EVAc (1000 kGy) shows brittle cracks, which are more intense and wider than those in iPP. The tensile failure surface of non-irradiated and irradiated iPP exhibits

brittle failure. The tensile failure surface of the non-irradiated blend containing 50% iPP shows ductile-type failure; however, on irradiation to 1000 kGy, it changes into a brittle failure. In the case of EVAc, the non-irradiated sample shows channel cracks, but the fracture of the irradiated samples (1000 kGy) shows brittle failure. 11.4.2.3.8 NR/EVAc Blends Koshy et al. [1992] investigated the effect of gamma irradiation of NR/EVAc blends as a function of the blend composition and the cure system (sulfur, peroxide or mixed) (Table 11.9). Blends of NR and EVAc were prepared in a laboratory

Use of High-Energy Radiation in Polymer Blends Technology

mixer, set at 80°C and a rotor speed of 60 rpm. The NR was masticated for 2.0 min and then blended with the EVAc for 2.5 min. The resulting temperature of the blend was in the range 110-128°C, depending on the blend ratio. The blends were compounded in a two-roll laboratory mill, with three cure systems, the mixtures and their ratios being blend (100.0), zinc oxide (5.0), stearic acid (1.5), styrenated phenol (1.0), dibenzothiazyl disulfide (0.8, only in the sulfur and the mixed cure systems); DCP (4.0, 40% active ingredient, only in the DCP and the mixed cure systems); and sulfur (2.5, only in the sulfur and the mixed cure systems). The compounded blends were molded to optimum cure in a steamheated hydraulic press at 160°C to obtain the test samples. The blends containing ≥50 wt% EVAc could not be fully cured with sulfur. The samples were irradiated with a Co-60 source at a dose rate of 3.21 kGy/h to doses of 100-500 kGy. Based on benzene extraction of the NR from the samples, SEM of the remainder allowed the authors to conclude that the EVAc forms the dispersed phase when its proportion is <40% and that NR forms the dispersed phase when its proportion is ≤40%. In the NR/EVAc, 60/40 and 50/50 blends, both the polymers remain as continuous phases leading to a sponge-like structure for the blend. On irradiation, NR undergoes degradation, whereas EVAc crosslinking. This is consistent with the modulus and tensile strength data (see Table 11.53). The radiation resistance of the EVAc-rich blends is better. Of the blends cured by the three systems, those cured with the mixed system showed lower retention of tensile strength but higher retention of the modulus. 11.4.2.3.9 cPI/tPI Blends Bhowmick et al. [1986] investigated properties of cis- and trans-polyisoprene blends (cPI/tPI), with a rather cursory look at the effect of gamma irradiation (Table 11.9). The blends consisted of guayule rubber (cPI) and synthetic tPI (TransPIP from Polysar). The blends were prepared by mixing the two polymers, along with the typical

833

additives for rubber, e.g., zinc oxide, stearic acid and antioxidant, in an internal mixer at 70°C. The mechanical blends were taken out hot and 1 phr DCP mixed with the samples in an open mill. At this stage sulfur and an accelerator were also added, where appropriate. The resulting samples were compression-molded at 150°C for 1 h and slowly cooled to room temperature. For irradiation, the blends were prepared by mixing the solutions of the two polymers in toluene. The blends were crosslinked by gamma irradiation for 24 h (no details given of the source, dose rate, etc.). Characterization of the crosslinked mechanical and solution blends showed that the level of crosslinking in the gamma-irradiated blends was lower. 11.4.2.3.10 Heterogeneous Elastomeric Blends Bohm and Nelson [1981] claimed that the morphology of heterogeneous blends, containing two or more radiation crosslinkable elastomers, can be stabilized by irradiation to the gel dose level (in this case, about 5% gel) (Table 11.9). The irradiated blends with stabilized morphology can be conventionally processed, e.g., by milling, extruding and injection molding; the final morphology of the blends after processing is substantially the same as the initial morphology. Examples of such blends are NR/EPDM, SB/EPDM and BR/EPDM. Processing can significantly alter the morphology of heterogeneous elastomer blends. Because of this, it is difficult to obtain the desired morphology for the products made from heterogeneous elastomer blends, such as tires. The morphology of these blends can be stabilized by moderate levels of vulcanization or curing. However, such stabilized blends can not be further processed. The irradiation method [Bohm and Nelson, 1981] gives stabilized blends that can be conventionally processed. During the processing of the radiationstabilized blends, conventional additives, such as antioxidants, accelerators, carbon black, etc., can be used, as required. Bohm and Nelson [1981] point out that these radiation-crosslinked blends have an additional advantage; they are more resistant to ozone than

834

A. Singh and K. Bahari

Table 11.53. Effect of γ-irradiation on the modulus and tensile strength of NR/EVAc blends [Koshy et al., 1992] Blend composition (wt%)

a

NR

EVAc

100 100 100 100 80 80 80 80 60 60 60 60 40 40 40 40 20 20 20 20 0 0 0 0

20 20 20 20 40 40 40 40 60 60 60 60 80 80 80 80 100 100 100 100

Dose (kGy)

100 200 300 500 100 200 300 500 100 200 300 500 100 200 300 500 100 200 300 500 100 200 300 500

% Retention, 300% modulus

% Retention, tensile strength

Sulfura

DCPa

Mixeda

Sulfura

DCPa

Mixeda

93 90 67 67 107 108 111 86 107 118 120 101 -

66 72 80 52 76 86 103 63 91 88 110 81 97 105 111 115 95 106 117 121 97 103 111 118

103 95 69 43 110 103 85 84 109 115 119 103 107 112 117 121 99 105 114 120 101 105 115 120

69 48 25 24 70 51 35 25 70 57 48 41 -

56 47 43 38 63 58 38 36 82 77 49 40 91 83 63 50 94 84 66 60 96 81 61 60

50 45 42 38 45 43 40 37 42 40 35 32 44 41 36 33 68 61 54 50 89 75 70 68

Refer to the cure system used.

the conventionally manufactured blends. This is due to the small domain or dispersed particle size achieved in the radiation-crosslinked blends. In the conventionally manufactured blends, the domain size tends to grow, particularly during curing, leading to greater sensitivity to ozoneinduced degradation. The required irradiation dose tends to increase with decreasing molecular weight of the elastomers in the blend. The required dose would also depend on the additives present, such as the antioxidants. Bohm and Nelson [1981] suggest a dose range of 1 to 200 kGy being sufficient for

this application. The irradiation can be conducted at ambient, or even lower temperatures (down to liquid nitrogen temperature), presumably in air (liquid oxygen tends to collect in liquid nitrogen open to air; therefore, irradiation of samples frozen in liquid nitrogen requires careful handling to avoid liquid oxygen mediated explosions). The lower temperatures would be particularly useful for blends that might change morphology on standing at ambient temperature. Irradiations using electron accelerators are preferred. In one of the examples given by Bohm and Nelson [1981], equal parts of a BR and EPDM

Use of High-Energy Radiation in Polymer Blends Technology

were mixed, along with 1% of DBPC (dibutylpara-cresol), for 15 min. The mix was then extruded at 150°C using a 0.8-inch extruder followed by rapid quenching in liquid nitrogen and storage in dry ice. The samples were later irradiated, presumably at room temperature with electrons, to doses of 0, 2.4, 4.7 and 140 kGy, and then annealed for 1 h at 100°C. The measured domain sizes in these blends were 16, 6.0, 4.8 and 4.5 µm for the doses of 0, 2.4, 4.7 and 140 kGy, respectively. The gel dose for the two individual polymers, and for the blend, is about 14 kGy. These data show that the domain size at the rather low dose of 4.7 kGy is only slightly larger than that obtained at the full cure at 140 kGy. Similar data have also been given for another example, in which case ozone resistance was also demonstrated by Bohm and Nelson [1981]. 11.4.3

Blends with Polystyrene, PS

Saunders and Pelletier [1976] claim in a patent to have produced polymer blends having high impact resistance, by blending together a nonelastomeric monovinylidene aromatic polymer, such as polystyrene, with a radiation-crosslinked elastomeric copolymer, such as a block copolymer of styrene and butadiene. However, since in this work the blends were not irradiated, they will not be discussed here. 11.4.3.1

PS/PMMA Blends

Blends of PS (Shell HR high purity) with PMMA (prepared by bulk polymerization initiated by azobisisobutyronitrile at 60°C) were dissolved in chloroform and co-precipitated in methanol (Table 11.9). The dried polymers were pressed into 3-mm sheets between stainless steel glazing plates at 190°C. Samples were evacuated for 50 h and sealed in glass tubes under nitrogen, then gamma-irradiated. PS and PMMA form immiscible, opaque blends. For the ESR measurements, the blends were prepared by freeze-drying from either benzene or naphthalene, and by coprecipitation and blending [Garrett et al., 1979, 1985].

835

A copolymer of styrene and methylmethacrylate, SMMA, showed higher flexural strength than the homopolymers. However, blends of PS/PMMA showed markedly lower strengths, the decreases being greatest at about 1:1 ratio. There was a decrease in modulus, but an increase in elongation at break produced by the two-phase morphology. Irradiation of blends resulted in decrease of the flexural strength and modulus — the effect was intermediate between that of PS and PMMA. The PS/PMMA blends containing 100, 90 and 70 wt% PS showed increasing radiation sensitivity. These systems could be considered dispersions of PMMA in a PS matrix, where the flexural strength is determined by the matrix. Addition of PMMA did not greatly reduce the radiation resistance of PS and addition of 10 wt% PS did not increase the radiation resistance of PMMA. In the blends, a protective effect of higher concentrations of PS was observed — this effect depended on the MW of the component polymers. The ESR spectra of the SMMA at -196°C showed a distinct gradation with composition. A study of the variation of the G(R•) value for the copolymer, with the copolymer composition, showed a marked decrease of G(R•) at low styrene compositions. This protective effect is analogous to that observed in gaseous product yields from liquid mixtures of benzene/cyclohexane (Figure 11.1), and in the SMMA copolymers. The protective effect is the result of transfer of charge or energy or both along the polymer chain to the aromatic group of the styrene unit. On warming the copolymer from -196°C to 0°C, chemical reactions occur in which MMA radicals are converted to styryl radicals. Schulz and Mankin [1976] investigated the degree of intermixing of PS and PMMA, by freeze-drying of their co-solutions in naphthalene (Table 11.9). They used two anionically polymerized polystyrenes, with narrow molecular weight distribution and two PMMA samples from Rohm and Haas Co. supplied in 1959 (L1489 and L1490). The L1489 PMMA was irradiated with 1.5-MeV electrons to give three different samples of lower molecular weight PMMA. For preparing

836

A. Singh and K. Bahari

the freeze-dried samples, 5 g of the polymer or the polymer mixture was dissolved in 95 g of naphthalene at 100°C, under nitrogen. The hot solution was poured slowly into ice-water undergoing high-speed stirring in a blender, to rapidly freeze the naphthalene solution. After filtration the naphthalene solution (obtained as a powder) was dried in an air stream, at room temperature, for 24 h. From this powder the naphthalene was removed at room temperature, by keeping it under vacuum for 1 week. The blends were irradiated with electrons, then the radiation-induced chain scission (PMMA) and crosslinking (PS) were measured. The results show that in the freeze-dried blends, PS protects PMMA against radiation-induced chain scission, and PMMA inhibits the radiation-induced crosslinking of PS. The authors concluded from these results that there was a considerable amount of intermixing of these two immiscible polymers, in their freeze-dried blends. Torikai et al. [1994] compared the effects of gamma irradiation of films of PS/PMMA blends and PS-PMMA copolymer (co-PS-PMMA) (Table 11.9). Polymer films were cast from methylene chloride solutions and were dried under vacuum. Based on the UV and FTIR spectroscopy, and viscosity measurements, Torikai et al. [1994] concluded that whereas the presence of PS in the copolymer provided protection against radiationinduced degradation to the PMMA units, similar

protection was not provided in the PS/PMMA blends. Katare et al. [1994] studied the microhardness of PS/PMMA blends, made from solutions of the two in benzene, and concluded that the increase in the microhardness of the blends on irradiation (Table 11.9) was due to crosslinking of PMMA. Since PMMA is known to degrade on irradiation (see Section 11.3, and the authors Introduction) the results obtained by these authors’ seem to require alternate explanation. 11.4.3.2

PS/EPDM Blends

PS blends with EPDM were compatibilized by addition of either SB (polystyrene/polybutadiene di-block copolymer) or SEP (polystyrene/ethylene propylene di-block copolymer) and were blended in a corotating twin-screw extruder T = 200°C (Table 11.9). The strands were quenched and pelletized. Some of the pellets were injection-molded (200°C) into 4 mm-thick specimens and irradiated to 50 kGy (3-MeV electron accelerator). Another part of the pellets was similarly irradiated and then injection-molded. The samples for notched Izod impact testing (ASTM D256) were annealed for 24 h at 80°C, or were left at RT for 3 weeks, before testing [van Gisbergen et al., 1990, 1991b]. The impact properties of PS/EPDM blends can be improved on electron irradiation with a suitable compatibilizer (see Table 11.54). It was

Table 11.54. Izod impact strengtha for irradiated and non-irradiated blends [van Gisbergen et al., 1991b]

a b c

PS (wt%)

EPDM (wt%)

SB (wt%)

SEP (wt%)

80 78 78 78 78

20 21 19 17 19

1 3 5 -

3

kJ/m2 at 23°C Irradiated before injection molding. Irradiated after injection molding.

Nonirradiated 2.5 4.0 7.9 11.0 11.0

Irradiated (50 kGy) 1.7b 3.2 b 13.0 b 21.0 b 9.6 b

9.1c 8.0 c

Use of High-Energy Radiation in Polymer Blends Technology

shown that a change in interfacial adhesion, the interaction between the butadiene part of the copolymer and the EPDM, increases because of irradiation to such an extent that the ultimate impact properties are determined by the interaction between the PS of the block copolymer (PSb) and the PS matrix. The Izod impact values of nonirradiated PS/EPDM blends gradually increased with increasing concentration of SB. Addition of only 1 wt% of SB improved the performance. Using 3 wt% of SEP in the non-irradiated blend gave a somewhat higher Izod impact value than that obtained for the blend with 3 wt% SB. This was due to the combined effect of increased adhesion and reduced particle size. The SEM micrographs showed that the average EPDM particle size decreased with SB or SEP concentration — the particle size with SEP was larger than with a comparable concentration of SB. Evidently, the adhesion properties of the SEP copolymer are good enough to compensate for the larger particle size. Irradiation before injection molding only slightly affects the EPDM particle shape and size. For specimens irradiated before injection molding, a strong increase of the impact strength was obtained for the blends with 3 and 5 wt% of SB, whereas addition of SEP had a negative effect. The impact strength of the blend irradiated after injection molding, and comprising 3 wt% of SB, slightly increased, but it was reduced when SEP was used as the compatibilizer. Using higher doses hardly changed the impact value. van Gisbergen et al. [1990, 1991b] concluded that the effect of crosslinking of the dispersed phase did not substantially contribute to the increase in Izod impact strength. Upon irradiation, these values remained constant even though the degree of crosslinking gradually increased. van Gisbergen et al. [1990, 1991b] considered a model for crack propagation in the blends, based on their results. According to the model, the block copolymer ought to be located at the interface between the PS matrix and the dispersed EPDM particles. The PS part of the block copolymer (PSb) is situated at the PS side of the interface and the polybutadiene (PB) or ethylene propylene

837

(EP) part at the EPDM side. If crack propagation of the blend during impact testing occurs via the interface, the ultimate impact value will be determined by the weakest part of the PS/EPDM interface. For non-irradiated blends, the fracture occurs at the EPDM/PB or EPDM/EP interface. Upon irradiation, the EPDM phase and the butadiene part of the SB block copolymer become covalently bonded. This implies that the interaction between the EPDM and PB may strongly increase through co-crosslinking of the two phases (interpenetrating network formation) or grafting of PB onto EPDM. The interaction between the PS and the PSb will remain constant since PS is insensitive to electron irradiation, as discussed in Section 11.3.4. By contrast, in the blend with the SEP, the interaction between the EP and the EPDM hardly changes. The authors quote Böhm and Tveekerm [1982] and suggest that in the EP part crosslinking or grafting occurs at a slow rate and is partly compensated for by chain scission [van Gisbergen et al., 1991b]. The fracture could thus occur at the EPDM/EP interface in the irradiated samples. The observed decrease of the impact strength for the SEP blends upon irradiation was due to chain scission of the SEP block copolymer. Some evidence for the increased adhesion was obtained by FTIR, NMR, and Tg measurements. In the IR spectra for irradiated blends, peaks at 700 and 758 cm-1 were clearly visible, characteristic of the styrene part of the SB compatibilizer. For the non-irradiated PS/EPDM samples, the IR spectra did not show these peaks. Evidently, removal of the block copolymer from the EPDM film was no longer possible after irradiation, since the interaction between the two had strongly increased. However, in the case of the irradiated SEP blends, the authors found no evidence of its crosslinking with EPDM. The 13C NMR analyses supported the FTIR results; the irradiated blends showed peaks at 127 and 144 ppm for the blend containing 5 wt% of SB. For the blend with SEP, almost no aromatic groups are detected. Again, this was an indication that the SB is more reactive than the SEP copolymer. Irradiation induced the largest

838

A. Singh and K. Bahari

changes at the interface in the blends with 5 wt% SB. This was also reflected in their Izod impact values. Compared to the blend without a compatibilizer, a clear shift of Tg to a lower value was visible for the blends with 3 and 5 wt% of SB; the blend with the SEP copolymer had Tg comparable to that of uncompatibilized blend. The shift in the Tg reflects the degree of interaction of the compatibilizer with the EPDM phase [Morbitzer et al., 1976; Booij, 1977; van Gisbergen et al., 1991b]. 11.4.3.3

PS/HDPE Blends

Meijer et al. [1988] and Elemans et al. [1988] investigated the potential of electron irradiation to stabilize the morphology of immiscible polymer blends. Their concept was that selective crosslinking of a dispersed phase in a matrix that remains unaffected, or degrades, should help fix the morphology of the blend. Elemans et al. [1988] investigated the protective effect of PS on HDPE during electron (3 MeV) or gamma irradiation (up to 25 kGy) (Table 11.9). PS and PE blends were made in a co-rotating twin-screw extruder. Results for these two types of irradiation were similar. The PS/HDPE blends showed that the co-continuous part is larger and that the imaginary lines indicating the cocontinuous area are not parallel. When the blends were subjected to a small deformation in compression, the elongation forces occurring in the surface layer during injection molding were insufficiently fixed, because of the low crosslinking efficiency of the PE phase. The shielding effect of the PS matrix made the PE phase less responsive to irradiation. The authors did not indicate the sample size. Since the irradiations were done in air, both the radical scavenging by oxygen (e.g., analogues of the reaction (R-27)) and the protective effect of polystyrene (e.g., analogues of the reactions (R-14) and (R-15)) would contribute to the apparent low levels of crosslinking of the PE phase.

11.4.3.4

PMMA/SAN Blends

Nguyen and Kausch [1984] found that the presence of phenyl groups in poly(styrene-co-acrylonitrile), SAN, protects PMMA in the blend, during radiolysis (Table 11.9). SAN and PMMA were dissolved in dimethyl formamide, 50-µm-thick films were cast, and then dried under vacuum. Transparent bar specimens were compressionmolded from the film. Mixing the polymers by co-precipitation from methanol resulted in opaque samples. Gamma irradiations were done in evacuated and sealed glass tubes, at a dose rate of 3 kGy/h. Comparison of freshly irradiated samples with irradiated and annealed ones showed the absence of any post-irradiation effects. Mechanical tests, molecular weight, and solubility measurements on irradiated PMMA/SAN blends showed that phenyl substitution in one of the polymers can offer partial radiation protection to the other component. Protective effects of the phenyl group are short range and occur only in miscible blends where mixing at the molecular level exists. However, the protection is not complete since some 15% of the yield for main chain scission in PMMA could not be suppressed, even at high styrene (in SAN) concentrations. The protection of PMMA by SAN against degradation is attributable to the charge and excitation energy transfer from PMMA to SAN (analogues of reactions (R-14) and (R-15)). Nguyen and Kausch [1984] explained the difference on the basis of polymer miscibility; PMMA and SAN are miscible, which facilitates many of the energy and charge transfer, and free radical/grafting reactions (e.g., analogues of reactions (R-14), (R-15), (R-30) and (R-31)), whereas the lack of miscibility of the PMMA/PS blends is not conducive to such reactions. Non-irradiated blends of PMMA/SAN showed composition dependence of the flexural strength. It was found that adding as little as 20 wt% of SAN to PMMA reduced the degradation effect of irradiation by a factor of 4. Irradiated PMMA showed random chain scission without crosslinking — G(scission) in µmol·J-1 for PMMA decreases from 0.12 at 100% PMMA, to 0.24

Use of High-Energy Radiation in Polymer Blends Technology

at 50% PMMA/50% SAN. The radiolytic yield of scission should decrease linearly with the PMMA level, if there was no influence of SAN. The protection efficiency of the phenyl group in SAN was slightly larger than in the copolymer isobutylene-co-styrene, since the cyano group in acrylonitrile acted as an electron scavenger. It was found that on irradiation SAN crosslinks, and a small fraction of PMMA is grafted onto the gel (analogues of reactions (R-30) and (R-31) involving radicals from PMMA). IR spectra of the insoluble fraction in irradiated blends showed the presence of carbonyl groups from PMMA incorporated in the network. When the two polymer chains were in close contact, a grafting reaction occurred by recombination between the macro-radicals of PMMA and SAN. Since under irradiation PMMA exclusively degrades, its presence in the gel can only be explained if a grafting reaction had occurred. The IR absorption coefficient of the carbonyl group in PMMA was 18 times stronger than the absorption coefficient of the cyano group in SAN. Using this ratio, it was determined that the gel fraction of 70 and 50 wt% SAN blends contained, respectively, 8.6 and 13.8 wt% of grafted PMMA. 11.4.3.5

PS/LDPE Blends

The microrheology of PS/LDPE blends, as influenced by electron (3-MeV electrons) irradiation-induced crosslinks was investigated by van Gisbergen and Meijer [1991] (Table 11.9), who attempted to relate the irradiation-induced changes in viscosity and elasticity, to the observed changes in the breakup and coalescence rates of the LDPE thread-like structures within the PS matrix. The blends (PS/LDPE, 80/20 wt%) were prepared in a co-rotating twin-screw extruder at an average barrel temperature of 200°C. The strands produced with non-equilibrium morphologies were quenched in water and irradiated to the desired doses (500, 100 or 200 kGy). The irradiated and non-irradiated strands were then heated to 200°C for selected periods of time in a Fontijne press at very low pressure. The changes in the morphologies of these blends were examined by SEM, using samples microtomed at -196°C.

839

On heating the extruded blend samples to 200°C for about 120 s, LDPE droplets (~1.5 µm diameter) were formed in the matrix. Irradiation of the samples (40, 100 or 180 kGy) before heating resulted in slower growth of the droplets. Special samples of the blends were prepared for more detailed investigation of the process of droplet formation: threads of LDPE were spun (radius, ~100 µm) and placed between PS plaques about 1 mm thick (for details see Elemans et al. [1991]). On heating these samples at 200°C for 40 min, Rayleigh distortions developed on the LDPE threads. The samples were cooled down, freezing in the distortions, before complete breakup of the threads occurred. The samples were then irradiated to the desired doses and reheated to see the effect on the distortions. van Gisbergen and Meijer [1991] expected that on heating, irradiation-induced crosslinking of the LDPE in the blends would prevent the LDPE from thread breakup and coalescence. They found that irradiation of the blend at low doses (40-180 kGy) did delay the breakup and coalescence. As the dose was increased to 472 kGy, the heat-induced thread breakup and coalescence of LDPE were prevented on heating for 10 min. However, the complete preservation of the original thread-in-matrix morphology of the PS/LDPE blend was not achieved, even at the high dose of 472 kGy. The authors concluded that in the case of these blends, the irradiation-induced crosslinking of the dispersed phase was less effective in delaying thread breakup and coalescence, than it would be expected for Newtonian materials with the same viscosity. The authors considered two possible reasons for the incomplete preservation of the morphology, by irradiation of the blends: (i) non-uniform crosslinking of the LDPE, and (ii) lower efficiency of the crosslinking of the LDPE in the PS matrix, than that expected for the LDPE alone. The latter effect is well known in many systems, as discussed earlier, because of reactions of the type (R-14) and (R-15). Depending on the sample thickness, radiation-induced oxidative degradation of the LDPE may also have contributed to the observed effects.

840

11.4.3.6

A. Singh and K. Bahari

PS/Polybutadiene Blends

Witt [1959] studied under vacuum gamma-radiation-induced crosslinking in butadiene-styrene copolymers, homopolymers and mixtures of these homopolymers, (Table 11.9). The behavior of the styrene units in the copolymers and in the physical mixtures, was different. Gel fraction measurements showed that in the copolymer, the styrene units did inhibit the crosslinking of the polybutadiene. However, there was no evidence of such inhibition in the mill- and latex-prepared physical mixtures of the two homopolymers. The polymers were prepared by emulsion polymerization. The monomers were removed from the latex obtained, and the polymer was coagulated with salt-alcohol, washed with water, and dried in a vacuum oven at 50°C. The physical mixtures of the two polymers, BR and PS, were made by combining the latex of the two and then coagulating the mixture. A set of mixtures was also prepared by combining the two polymers in a 4-inch mill at a roll temperature of 50°C. The polymeric samples were sealed in glass ampoules under vacuum and irradiated. The lack of inhibition of BR crosslinking by PS, in the mixtures, is due to the fact that these mixtures were not really the blends of the two polymers, because of the processing at these low temperatures. As a result, the intermolecular distances between the BR and the PS were most likely too large for efficient energy or charge transfer or intermolecular free radical reactions. Schrijver et al. [1988] patented a process that permits the preparation of impact-resistant blends with radiation-resistant polymer as the matrix, a radiation-crosslinkable rubber phase and a suitable compatibilizer, so that the final product does not exhibit delamination (Table 11.9). The process yields a mixture in which the optimum size and shape of the particles, as present in the matrix, are essentially fixed, and thus are retained throughout further processing, e.g., injection molding. The dispersed rubber phase becomes wholly or partly cross-linked as a result of the irradiation and remains anchored in the matrix, while the matrix itself is hardly

affected by irradiation, and therefore remains thermoplastic. The preferred compositions comprize 10 to 25 wt% rubber. Though many different thermoplastics can be used [Schrijver et al., 1988], the authors preferred polymers or copolymers of styrene; all the data given are for Styron 686 E (RTM). A wide variety of rubbers can be used, such as polybutadiene (BR), and butadieneacrylonitrile copolymer [Schrijver et al., 1988]. The rubbers used by the authors are listed in Table 11.55, along with their data on the irradiated and non-irradiated blends. Schrijver et al. [1988] found that the impact resistance of the material increases as the particle size of the rubber decreases, with the preferred size being 0.2 to 0.5 µm. Suitable additives can help reduce the size of the rubber particles, promoting adhesion between the rubber and the thermoplastic. The authors suggested that a part of the additive should be soluble in the thermoplastic and another part in the rubber, such as ethylene-styrene copolymer, or the commercially available Kraton G-1652 (RTM). The recommended concentrations of the additives are in the range of 0.1 to 10 wt% of the thermoplastic-rubber mixture. The blend prepared from these compositions can be processed into the desirable form, e.g., granules, and electron or gamma irradiated. The blends were made in a batch kneader with a starting temperature of about 110°C, which increased to about 200°C during kneading. The kneaded mixture was granulated using an extruder. The irradiated and non-irradiated granules were injection molded into test bars according to ASTM standard 256/56. The results obtained by Schrijver et al. [1988] (Table 11.55) show that (i) the non-irradiated blends delaminated on deflection, whereas the irradiated samples did not delaminate; (ii) the Izod impact values of the samples decreased with increasing temperature of injection molding; (iii) the effect of irradiation on the blend Izod impact values was very small, if any; (iv) Keltan 378 and 578 as the matrix thermoplastics showed the highest Izod impact values; and (v) with Keltan 578 as the matrix thermoplastic, Kraton G1652 (RTM) as the rubber

841

Use of High-Energy Radiation in Polymer Blends Technology

Table 11.55. Izod impact values of thermoplastic/rubber blends [Schrijver et al., 1988] Styron

Rubber

Compatibilizer

686E (wt%)

(wt%)

(wt%)

Keltan 378,20 Keltan 378,20 Keltan 378,20 Keltan 378,20 Keltan 378,20 Keltan 378,20 Keltan 378,20 Keltan 578,16.5 Keltan 578,16.5 Keltan 312,16.5 Keltan 512,16.5 Keltan 712,16.5 Keltan 812,16.5 Keltan 312,20 Keltan 312,20 Keltan 578,16.5 Keltan 578,16.5 Keltan 578,16.5

Kraton G1652,5 Kraton G1652,5 Kraton G1652,5 Kraton G1652,5 Kraton G1652,5 Kraton G1652,5 Kraton G1652,5 Kraton G1652,3.5 Kraton G1652,3.5 Kraton G1652,3.5 Kraton G1652,3.5 Kraton G1652,3.5 Kraton G1652,3.5 Kraton G1652,5 Kraton G1652,5 Cariflex TR1107,3.5 Cariflex TR1107c,3.5 Cariflex TR1107d,3.5

75 75 75 75 75 75 75 80 80 80 80 80 80 75 75 80 80 80 a b c d

Injection molding

Dosea

temperature (°C)

(kGy)

Izod valueb (kJ/m2)

Delamination

220 240 240 240 300 300 300 200 200 200 200 200 200 220 220 200 200 200

60 0 50 75 0 50 75 30 60 50 50 50 50 3 7 40 40 40

15.0 8.3 8.2 8.3 3.7 4.5 5.2 9.8 10.0 6.2 4.4 4.5 4.6 12.2 12.8 4.4 6.4 7.1

No Yes No No Yes No No No No No No No No No No No No No

Irradiation temperature, 50-53°C (most likely); electron irradiation (atmosphere not given) Experimental error, ± 0.1 to ± 0.3 Hydrogenated to 80% Hydrogenated to 90%

phase gave higher Izod impact values than the Cariflex TR-1107 (as supplied by Shell, and the two hydrogenated samples). 11.4.3.7

Polystyrene/Poly(vinyl methyl ether) Blends, PS/PVME

PS and PVME are a rather unique pair of polymers that are dissimilar, both chemically and in their response to irradiation, yet they form miscible blends over the entire range of composition. McHerron and Wilkes [1993] report on the effects of dose, blend composition and the state of miscibility on the gel content obtained on electron irradiation of PS/PVME blends (Table

11.9). About 75 µm thick films were made by casting toluene solutions of the polymers onto Teflon sheets. The films were dried for 24 h at room temperature and then for 48 h under partial vacuum and elevated temperatures (80 or 120°C), to remove all solvent. Phase separations of the film samples were carried out isothermally at 160°C for either 0.5 or 2 min. The samples were irradiated at room temperature at the rate of 100 or 200 kGy per pass and were exposed to air in between the passes. The small amount of heating occurring during irradiation, due to the very high dose rates, was not measured. The gel fraction of the blends increases with increasing dose, and with increasing concentration

842

A. Singh and K. Bahari

of the PVME in the blends. At 500 kGy dose there is no gel observed until the PVME concentration has reached about 15%; however, at 1000 kGy dose, the gel fraction of the sample with 15% PVME is 30%. The reasons for this difference are not clear, though the authors seem to attribute it to mutual grafting of the PS and PVME. The gel fraction values of the phase separated blends were generally lower, as shown by the data in Table 11.56. The composition of the gels was compared to the composition of the corresponding blends, based on FTIR analyses. The PVME content of the gel increased with increasing PVME content of the blends, but not quite linearly. The Tg values of the blends were found not to be significantly affected by irradiation.

11.4.3.8

Deuterated Polystyrene/Poly(vinyl methyl ether) Blends, PSD/PVME

Briber and Bauer [1988] investigated the effects of radiation-induced crosslinking on the phase diagram, spinodal temperature, and the small angle neutron scattering (SANS) function for a miscible polymer blend of PSD and PVME (Table 11.9). They found that the gamma radiation-induced crosslinking increased the single phase region of the phase diagram and that the main features of the results were consistent with the predictions given by de Gennes [1979]. The samples for SANS were prepared by dissolving 30 wt% PSD and 70 wt% PVME in toluene. A thin film (0.1 mm) was cast and dried under vacuum. Disks were cut from the film, stacked and hot pressed at 120°C into spacer rings. The disks were sealed under high vacuum and gamma irradiated to 1250 kGy at ~45°C.

Table 11.56. Gel fraction of PS/PVME blends on electron irradiation [McHerron and Wilkes, 1993] Blend composition (wt%) PS

PVME

100 100 95 95 85 85 70 70 50 50 30 30 15 15 0 0

0 0 5 5 15 15 30 30 50 50 70 70 85 85 100 100

a Phase separated at 160°C for 2 min, before irradiation

Dose (kGy)

Gel fraction (%)

Gel fractiona (%)

500 1000 500 1000 500 1000 500 1000 500 1000 500 1000 500 1000 500 1000

0 1 0 33 0 56 43 70 68 85 78 88 84 88 90

16 50 35 63 50 73 68 87 76 92 -

843

Use of High-Energy Radiation in Polymer Blends Technology

The samples for crosslinking kinetic studies were cast into thicker films (0.5 mm), sealed under high vacuum and gamma irradiated to 40 kGy. GPC (gel permeation chromatography) analyses of the irradiated samples suggested that chain scission for both PSD and PVME was minor. FTIR analyses indicated mutual grafting of PSD and PVME. The inverse of the spinodal temperatures obtained (Table 11.57) varies linearly with dose, and inversely with Nc (number of monomers between crosslinks), as predicted by de Gennes [1979]. 11.4.4

of the LCP are required to lower the coefficient of expansion, resulting in low elongation, and thus the resulting blends are unsuitable for use in dimensionally recoverable articles. The work was aimed at getting around this problem by the use of irradiation. Toy et al. [1994] recommend the following compositions of the blends: (i) 85-96 wt% of the matrix polymer (either PVDF, or another fluoro-polymer or copolymer containing at least 80 mol% of vinylidene fluoride homopolymer, or HDPE); (ii) 2-4 parts by weight of a LCP (such as, copolymers of: p-hydroxybenzoic acid (HBA) and PET; HBA and 6-hydroxy-2-naphthoic acid (HNA); HBA, terephthalic acid (TA) and aminophenol; and TA, 2,6-naphthalenedicarboxylic acid and 4-(4’-hydroxyphthalimido)phenol; (iii) 1-10 parts by weight of a radiation crosslinking promoter (such as TAC, TAIC and TMPTMA); and (iv) 0.05-1 part by weight of a thermal degradation inhibitor (such as, Goodrite 3125TM, Irganox 1010TM and Topanol CATM), or acid acceptors such as calcium carbonate and zinc oxide, or both.

Miscellaneous Blends

11.4.4.1

Liquid Crystalline Polymer/ Polyvinylidene Fluoride or High Density Polyethylene Blends, LCP/PVDF or HDPE

Toy et al. [1994] investigated the effect of irradiation on blends of liquid crystalline polymers (LCP) with either polyvinylidene fluoride (PVDF) or HDPE (Table 11.9); such blends are useful for making dimensionally recoverable articles. This patent contains many claims but few results. Blending of PVDF with LCP improves its mechanical properties and reduces the coefficient of linear expansion. However, high concentrations

The blends were prepared by melt blending the components in a single- or twin-screw extruder

Table 11.57. Variation of spinoidal temperature (Ts) and number of monomers between crosslinks (Nc) with dose [Briber and Bauer, 1988] Dose (kGy)

Nc (PVME)

Nc (PSD)

Method (Nc)

Ts (°C)

0 250

780 1200 390 610 160 250

170 270 90 140 40 60

GPC FTIR GPC FTIR GPC FTIR

149 240

500 1250

GPC: Gel Permeation Chromatography FTIR: Fourier Transform Infra Red Spectroscopy

255 430

844

A. Singh and K. Bahari

(or other similar device). The recommended processing temperature is slightly above the melting point of the LCP (if the melting point of the LCP gets depressed in the blend, then the depressed melting point should be used as the reference temperature). They recommended irradiation doses up to 500 kGy, preferably lower than 250 kGy. SEM of the samples showed that the LCP was dispersed in the matrix polymer as microfibrils, creating a reinforcing effect. As shown by the data in Table 11.58, the irradiated samples have reduced shrinkage, but the radial expansion is unaffected by irradiation. 11.4.4.2

Polysulfone/Polyvinylpyrrolidone Hollow Fiber, PSF/PVP Blends

In a novel application in which a blend is first formed in a liquid solution, Kobayashi and Tanaka [1992] developed a selectively transmissive hollow polysulfonic (PSF) fiber, containing a hydrophilic polymer (e.g., polyvinylpyrrolidone, PVP) that was crosslinked by gamma irradiation into a water-insoluble form, for the medical or food industry use (Table 11.9). The hollow fiber was spun by the use of a spinning solution containing a polysulfonic resin (10 to 30 wt%) and a hydrophilic polymer (1 to 20 wt% of molecular weight 10,000 to 500,000). The hollow fibers can then be crosslinked in the wet state, either before or after

being made into a module, by gamma irradiation (10-50 kGy). The solution in the fibers, or the module, was discharged and the fibers are washed with water to remove uncrosslinked hydrophilic polymer. Detailed steps for a typical sample are as follows: (i) mix 18 wt% PSF and 9 wt% PVP (molecular weight, 40,000); (ii) add 43 wt% of dimethyl acetamide (DMAc), 29 wt% of dimethyl sulfoxide (DMSO) and 1 wt% of water; (iii) stir the mixture for 15 h to dissolve the components, maintaining it at 85°C; (iv) spin the hollow fiber [Kobayashi and Tanaka, 1992]; (v) replace the solution in the fiber with a solution of glycerin (70 wt%); (vi) prepare a fiber module with a membrane area of 0.7 m2; (vii) wash the module with warm water (35°C) and charge it with water; (viii) gamma irradiate the module (presumably in air; PVP crosslinking would be more efficient in the absence of oxygen [Rosiak, 1991]) to a dose of 25 kGy; (ix) wash the module and seal it in a polyethylene bag; (x) gamma irradiate the sealed module to a further dose of 25 kGy, for sterilization. The other compositions of the polymers used by Kobayashi and Tanaka [1992] were (i) PSF (P-3500)/PVP, 18/18; (ii) poly(ether sulfone) (PES, VICTREX 4800P, ICI)/PVP, 18/9; (iii) PSF(P-3500)/PSF (P-1700)/PVP, 4/12/6; and (iv) PSF (P-3500)/PVP, 15/9.

Table 11.58. Effect of electron irradiation (60 kGy) on Kynar 450/Vectra A950 blends [Toy et al., 1994] Blend composition (wt%) Kynar 450a 100.0 98.0 96.0 a

b

Longitudinal shrinkage (%)

Vector A950b 0 2.0 4.0

Non-irradiated 40 24 0

Irradiated 7 5 0

Elongation (%)

Radial expansion (%)

Non-irradiated

Irradiated

490 440 400

400 400 400

Contains by weight, Irganox 1010 (0.10), diallyl phthalate (2.30) and triallyl cyanurate (2.30) for 91.3 by weight of Kynar 450. This mixture was melt-prepared in a mixer (such as Banbury mixer) at 199°C and pelletized. The blends of the mixed and pelletized Kynar 450 with Vectra A950 were prepared in a co-rotating twin-screw extruder (screw speed 150 rpm; die pressure, 3.3 MPa (500 psi); melt temperature, 304°C).

Use of High-Energy Radiation in Polymer Blends Technology

11.4.4.3

Mesophase PP-Based Blends

Wilfong and Rolando [1993] described the preparation of a large number of mesophase PP-based blends (Table 11.9) that are softer or tougher or both than comparable blends based on crystalline PP, and are capable of maintaining substantial structural integrity after exposure to ionizing radiation (for sterilization or grafting or both). The suggested uses of these blends include the production of films, fibers, tubes, and pouches for the medical or food industry. They demonstrated radiation-induced grafting of another polymer on the films of some of these blends and suggested that the properties of the blends, such as surface adhesion, friction, oxygen permeability, and moisture permeability, could be enhanced by the grafted layers. Medical devices and their packagings often require heat sealing. The packaged medical devices are also being sterilized by exposure to high-energy radiation [Williams, 1992; Morrissey and Phillips, 1993]. Polypropylene is a polymer of choice for many medical uses because of, for example, its lack of toxicity, inertness to drugs, low cost, and ease of processing. However, there are two inherent disadvantages to the use of the crystalline PP: its inability to be heat-sealed against other materials, and gamma-irradiation-induced degradation (which continues in the post-irradiation period, e.g., through analogues of reactions (R-22) and (R-23)) of its properties, such as embrittlement and discoloration. Many approaches have been taken to overcome these disadvantages [Williams, 1992; Morrissey and Phillips, 1993; Wilfong and Rolando, 1993]. One such approach has been to use mesomorphous PP [Insley et al., 1990a] because of its resistance to irradiation [Insley et al., 1990b]. However, the heat sealability and the radiation resistance of even the mesomorphous PP remain unsatisfactory. Another approach has been to develop polymer blends of mesomorphous PP with a miscible polymer, which resulted in a product with radiation resistance similar to that of mesomorphous PP [Krueger et al., 1991]. The aims of Wilfong and Rolando’s work [1993]

845

included overcoming these weaknesses. Many different though blends are mentioned in the patent, the ones for which mechanical property measurements are given, are listed in Tables 11.9, 11.59 and 11.60, and include CO, EAA, EVAc and PB. They claim that the mesopolymer blends described in this patent substantially maintain their structural integrity for a useful period after radiation sterilization and provide other improvements in properties such as lower modulus of elasticity and increased fracture strain. They also claim that (i) conventional additives, such as pigments, dyes and plasticizers, can be added to the blends (”10%); (ii) the blends do not require any stabilizers or antioxidants to protect their mechanical properties from damage during irradiation and in the post-irradiation period; (iii) the mesopolymer blends are substantially softer, and tougher, than comparable crystalline blends; and (iv) the mesopolymeric blends exhibit increased radiation resistance (dose range 1-200 kGy). Their work also includes property modification of the surface of the films of the developed blends, by radiation-induced grafting of another polymer. Wilfong and Rolando [1993] extruded dryblended polymer mixtures. To obtain mesopolymeric blends, it is necessary to rapidly cool the extruded product, and the shapes and thicknesses of the extruded materials are limited by the ability to cool them rapidly below 24°C. The methods of quenching included plunging in, or, spraying with a cold liquid, e.g., ice water, or running the extruded material over a cooled drum or roll or both. The Young’s modulus values of the non-irradiated mesomorphous and crystalline blends are compared in Table 11.59. Wilfong and Rolando [1993] claim that on electron irradiation these values are lower by 1-99%, but the data are not given in the patent. The data supporting the authors’ claims on the improvement in adhesion, obtained on irradiation of two of the blends, grafted with either acrylic acid (AA) or dimethylacrylamide (DMA) are shown in Table 11.60. The data on fracture strains obtained by the authors do show that in most cases the fracture strain values of the unirradiated mesopolymer blends are higher than the values expected on

846

A. Singh and K. Bahari

Table 11.59. Young’s modulus valuesa of the unirradiatedb mesomorphous (quenched) and crystalline (not quenched) PP blends [Wilfong and Rolando, 1993] Blend

CO1/CO2 CO1/PP2 CO1/EVAc PP1/EAA PP1/EAA PP1/EAA PP1/PB/EAA PP1/PB/EAA PP1/EVAc PP1/EVAc PP1/EVAc PP1/EVAc PP1/EVAc PP1/PB/EVAc PP1/PB/EVAc PP1/PB/EVAc PP1/PB/EVAc a b

Composition (weight ratios)

Young’s modulus (MPa) Mesomorphous

Crystalline

1/1 1/1 1/1

248 136 306

380 187 408

3/1 1/1 1/3 3/3/2 1/1/2

259 203 124 125 113

346 238 137 181 130

1/0 3/1 1/1 1/3 0/1

352 275 189 109 23

431 315 223 132 42

1/1/0 3/3/2 1/1/2 1/1/6

147 115 93 72

186 161 118 73

Using 50 x 25 mm film (ASTM D-882-88) Values for irradiated samples were not given.

CO1: PP/PE copolymer (#7371, Fina) CO2: PP/PE copolymer (#6571, Fina) EAA: Ethylene/acrylic acid copolymer (PRIMACORTM, #3340, Dow) EVAc: Ethylene/Vinyl acetate copolymer (#UE 656, Quantum) PB: Polybutylene (#PB400, Shell) PP1: Polypropylene (#PP3576, Fina) PP2: Polypropylene (#5A95, Shell)

the basis of simple additivity of the component polymers. However, no data are given on the effect of irradiation on the fracture strains of these blends. Wilfong and Rolando [1993] studied the formation and decay of free radicals in some of the

blends and found that the free radical decay in the mesopolymer blends was much faster than that in the corresponding crystalline blends (by a factor of about 10). This is consistent with the authors’ claim that the mesopolymer blends are more radiation-resistant. Lower post-irradiation

847

Use of High-Energy Radiation in Polymer Blends Technology

Table 11.60.

Peak peel adhesion of blend films with a layer of monomer grafted by electron irradiation [Wilfong and Rolando, 1993] Monomer

Dosea (kGy)

Peak peel valueb (g/25 mm)

PP1/EVAc, 1/1 PP1/EVAc, 1/1 PP1/EVAc, 1/1 PP1/EVAc, 1/1

none none AA DMA

0 50 50 50

4.4 7.3 13.7 13.7

PP1/EAA, 1/1 PP1/EAA, 1/1 PP1/EAA, 1/1 PP1/EAA, 1/1

none none AA DMA

0 50 50 50

5.4 8.4 13.2 13.1

Blend

AA: Acrylic acid (Aldrich) DMA: Dimethylacrylamide (Chem. Services, Inc., Winchester, PA) a Electron irradiation under nitrogen, presumably at room temperature b 180° peel adhesion measurements using 2.5-cm-wide, 20.3-cm-long strip of pressure-sensitive adhesive tape (ScotchTM, #8411, 3M)

concentration of free radicals would result in lower post-irradiation oxidative degradation of the blends, based on the mechanisms of such degradation discussed earlier (e.g., through reaction (R-22)). 11.4.4.4

Transparent Polymer Alloys

Hauser et al. [1993] claim that articles for medical and other uses made from a polymer alloy of a clear styrenic acrylate, a ductile styrenediene polymer, and a rubbery block styrene-diene copolymer may be sterilized by gamma irradiation, without significant increase in its yellowness index or significant deterioration of the alloy’s mechanical properties (Table 11.9). There are many applications, e.g., urine sample bottles, items in space probes, and cosmetic containers, that are made from transparent plastics and need to be radiation-sterilized. However, many clear plastics do develop color on irradiation, and their mechanical properties tend to deteriorate, particularly on irradiation in air. Hauser et al.

[1993] targeted the development of the styrenic acrylate alloys that on gamma sterilization would not change their yellowness index by more than 5 and would retain at least 80% of their tensile elongation values. The polymer alloys (ZYLARs containing styrene, methyl methacrylate and butadiene-based polymers) are listed in Tables 11.9 and 11.61, along with the data on the radiation effects on these alloys. The alloys can be made by conventional extrusion methods. As the data in Table 11.61 show, the yellowness index of the four alloys increase by 1.6 or less, for the maximum dose of 50 kGy (the conventional dose for gamma or electron sterilization is 25 kGy). In comparison, the increase in the yellowness of four other plastics (a gamma grade of PVC, PVC(G), KODAR PETG 6763 (presumably, a gamma grade of PET), a gamma grade of polycarbonate (PC(G)), and a PC was much larger, up to about 50. Barring a few exceptions, the overall trend of the tensile strength and flexural modulus was to increase with increasing dose, with simultane-

848

A. Singh and K. Bahari

ous decrease in the Izod impact strength. These changes can be attributed to the crosslinking reactions in the alloys, mainly because of the presence of the dienes in them. One would expect a decrease in the elongation values with increasing dose, both because of the crosslinking, as well as because of any oxidative degradation (if the irradiations were done in air). However, although there are hints of such decrease for ZYLARs ST-94-560 and 568 (Table 11.61), there does not seem to be a clear trend.

11.4.4.5

Heat-Shrinkable Film

Yoshimura et al. [1989] give a lengthy description of their invention, which was primarily aimed at producing film (Table 11.9) particularly suited for shrink wrapping, with many examples and extensive sets of results. The film is made from the following three groups of polymers: (i) LDPE and a copolymer of ethylene, either with vinyl ester monomer, or with an unsaturated aliphatic monocarboxylic acid or its aliphatic ester (EVAc being the most desirable copolymer); (ii) one of the

Table 11.61. Effect of gamma irradiationa on the yellowness index and mechanical properties of optically clear plastic alloys [Hauser et al., 1993] ZYLARb ST-94-560

Property

Dose (kGy)

Haze (%)

0

3.1

3.0

2.9

2.8

10 30 50 0 10 30 50 10 30 50 0 10 30 50 0 10 30 50

0.2 0.2 1.6 29.5 29.8 30.2 30.5 89 95 84 2.15 2.12 2.25 2.30 214.1 198.7 198.7 184.8

0.4 0.8 1.3 24.9 24.9 25.1 25.2 99 100 100 1.71 1.68 1.59 1.90 336.3 351.8 315.4 336.3

0.4 0.8 1.3 21.0 21.0 20.8 21.0 100 100 100 1.53 1.57 1.59 1.72 500.7 494.2 485.2 476.7

0.1 0.6 1.5 41.0 41.3 41.7 42.2 100 100 92 2.63 2.59 2.86 2.73 36.3 34.7 35.8 34.7

∆ Yellowness index

Tensile strength (MPa)

Retained tensile elongation (%)

Flexural modulus (MPa)

Izod impact strength (J/m)

a b

ZYLAR ST-94-561

ZYLAR ST-94-562

ZYLAR ST-94-568

Presumably in air, presumably at room temperature; details were not given. ZYLAR series of alloys of brittle polymer, ductile polymer and rubbery polymer are available from Novacor Chemicals, Inc. The components are brittle polymer, styrene (70 wt%) and methyl methacrylate (30 wt%) (Novacor’s NAS30); ductile polymer, block copolymer derived from styrene (75 wt%) and butadiene (25 wt%); and rubbery polymer, a tapered polymer derived from styrene (43 wt%) and butadiene (57 wt%). No specific gamma stabilizers were added to these polymer alloys.

Use of High-Energy Radiation in Polymer Blends Technology

crystalline polymers, PP, HDPE or polubutene-1; and (iii) a mixture of elastomers: an ethyleneα-olefin copolymer (e.g., ethylene-propylene or ethylene-butene-1 copolymer), and a block copolymer containing a monovinyl substituted aromatic hydrocarbon (e.g., styrene-butadiene or styreneisoprene block copolymer), and a polyolefin. The film consisting of these components was made by melting and blending the selected composition, extruding the blend through an annular die, cooling the tubular sheet with a refrigerant, irradiating the sheet and heating it up to 110°C, and then cold stretching the sheet into a film to an area stretching ratio of 5 to 30. They recommended either irradiation or use of peroxide additives to crosslink the films to obtain gel fractions up to 60% (preferably, 20%). The doses of 500-keV electrons used were between 30 and 100 kGy (many experimental details not given). Yoshimura et al. [1989] claim that the films produced by using this invention possessed outstanding low-temperature shrinkability, strength and optical properties. 11.4.4.6

Microstructure of Miscible Blends

Vesely and Finch [1988] investigated an electron microscopy-based method to study mixing (miscibility, solubility) of polymers at the molecular level, for the use as resists. They selected five polymer blends, which were reported to be miscible, at least over a part of the composition range: PVC/SAN, PVC/PMMA, PVC/PCL, POM/PVC and POM/PMMA (Table 11.9). PVC stabilized with a 1% organotin stabilizer was blended with SAN, PMMA and PCL by mixing on a two-roll mill. Polymer blends of POM with PVC and PMMA were prepared from a solution in dimethyl formamide by casting onto a glass slide at 50 and 165°C, respectively. The solvent was removed by placing the samples on the glass slides in a vacuum oven at 50°C for 1 d. Vesely and Finch used two radiation effects for contrast enhancement of the samples, which were examined by scanning transmission electron microscopy: (i) differential formation of double bonds (through reactions such as (R-13)) in a component polymer, and (ii) differential mass

849

loss of a component polymer through formation of volatile products. They concluded that in the PVC/SAN blend there was no preferential mass loss (suggesting that there was no energy or charge transfer between the two polymers leading to the degradation or protection from degradation of the PVC by SAN). In the PVC/POM and PMMA/POM blends, preferential degradation of the PVC and the PMMA was observed, suggesting that there was energy or charge transfer from POM to the other component in the blends, leading to degradation reactions. In the case of the PVC/PCL blend, the degradation of the PVC seemed to be prevented by the presence of the PCL, particularly at levels >20%. They interpreted preferential degradation, or prevention of degradation, as a sign of intermolecular interaction between the components of the blends. 11.4.4.7

Heat-Shrinkable Fluoro-Polymer Blends

Cheng et al. [1990] refer to several possible polymeric blends, consisting of fluorinated and non-fluorinated, elastomeric and non-elastomeric, polymers and copolymers, which may also contain crosslinking agents and other additives, for use in cable jackets and heat-recoverable (e.g., heat-shrinkable) articles. They point out that PVDF, and copolymers of ethylene and tetrafluoroethylene (ETFE) have many desirable properties but their high stiffness can be undesirable for certain applications. However, their blends with some thermoplastic elastomers show markedly reduced stiffness, while substantially retaining their other desirable properties. Cheng et al. [1990] preferred irradiation of many of the blends in the presence of crosslinking agents such as TAC and TAIC (Table 11.9). They used two different ETFE samples, with melting points of 220 and 270°C. The lower melting ETFE and the TPE, along with TAIC and unspecified additives were blended in a mixer for 3-4 min at about 240°C. The samples were molded into 25-mil-thick slabs and irradiated to 50 kGy. Some of the irradiated samples were annealed at 150°C for 2 h (the conditions of cool-

850

A. Singh and K. Bahari

ing of the annealed samples were not specified). Similar blends with the higher melting ETFE were made in a single-screw extruder at 300°C. Some of their results obtained using these blends are given in Table 11.62. As the data in the Table show, (i) the tensile strength decreases with increasing ETFE content; (ii) at ETFE content greater than 10%, there is marginal increase in the tensile strength on irradiation to 50 kGy; in fact, the increase on irradiation seems to increase with the increasing content of the TPE; (iii) the elongation of the blends increases with decreasing ETFE content, and on irradiation to 50 kGy. 11.4.4.8

EPDM/Polyacetylene Blends

Lee and Jopson [1983] prepared blends of EPDM and PAC (polyacetylene), to impart flexibility to the otherwise stiff PAC (Table 11.9). EPDM samples were purified by dissolution in heptane followed by precipitation in methanol. The blends were made by two methods, starting with a solution of the EPDM in freshly distilled toluene to which Ziegler-Natta catalyst had been added: (i) deposit an uniform film of the EPDM, containing the catalyst, from the solution under vacuum

in a rotating flask and then introduce acetylene gas into the flask; and (ii) introduce acetylene gas into the toluene solution of EPDM and the catalyst. The first procedure produced a film that was removed and dried overnight under vacuum; the second procedure gave a gel. The toluene was removed under vacuum, and the gel was pressed and dried under vacuum. IR and SEM analyses showed that the blends contained both EPDM and PAC regions, with no evidence of grafting of one to the other. The PAC was formed predominantly in the trans-configuration. The films could be crosslinked either by treatment with sulfur monochloride in toluene solution, or by irradiation under vacuum. The chemically crosslinked film could not be doped with iodine. However, the radiation-crosslinked film could be doped with iodine, resulting in a product with conductivity as high as 10/Ω mm. IR analysis suggested that only the EPDM in the blends was crosslinked by irradiation. 11.4.4.9

Polycarbonate/Polyester Blends

Polycarbonate (PC) is well known for its superior impact strength, clarity, heat resistance and

Table 11.62. Effect of irradiation on polymeric blends [Cheng et al., 1990] Blend composition (wt%) ETFEa 97 77 50 10 97e 77e 47e 10e a b c d e

TPEb 20 47 87 20 50 87

TAIC 2.1 2.1 2.1 2.1 2.1 2.1 2.1 2.1

Tensile strength (MPa) at kGy Addc 0.9 0.9 0.9 0.9 0.9 0.9 0.9 0.9

MP = 220°C Dai-el T-530 from Daikin Industries additives, unspecified irradiated samples annealed at 150°C for 2 h MP = 270°C

0 32.3 31.5 21.2 4.1 39.5 32.2 15.8 5.1

50 34.9 31.4 26.3 22.5 37.2 31.9 26.5 19.2

50d 32.9 30.9 27.1 22.2 37.1 32.6 26.9 21.2

Elongation (%) at kGy 0 320 366 420 1040 290 315 290 955

50 280 305 335 400 220 240 305 380

50d 243 280 330 405 180 220 290 390

Use of High-Energy Radiation in Polymer Blends Technology

dimensional stability. This combination of properties makes PC the resin of choice for many applications requiring tough, transparent, and complex parts, including medical devices such as blood pumps, connectors and syringes [Mason, 1995]. However, it develops yellow coloration on sterilization by irradiation. In this paper Mason [1995] describes work with a developmental transparent blend (DB, Table 11.9) which seems to have acceptable yellowness index. The blends were prepared in a twin-screw extruder. The spiral flow was determined at a melt temperature of 288°C (550°F) for the PC and 266°C (510°F) for the transparent DB. The mold release energy was determined by measuring the force required to eject a disc measuring 76 mm in diameter and 3.2 mm thick with 6.4 mm high ribs. The chemical resistance was determined by measuring the maximum allowable outer fiber strain on 2.0 mm tensile bars, which did not reduce tensile elongation. The blends of PC with PET, PBT and PETG were opaque whereas the blends with PCT and PCTG were transparent. The blends of an unidentified copolycarbonate, (coPC) with PET, PBT, PETG and PCTG were transparent. The opacity of the blends was attributed to their crystallinity. Mason [1995] found the irradiated (30 kGy, gamma irradiation) DB to show more acceptable yellowness index than irradiated PC (3 for DB as compared to 15 for PC), good injection molding characteristics (spiral flow similar to that of PC), similar part ejection force for PC and DB, and reasonable chemical resistance. 11.4.4.10 PMMA/CO Blends The use of PMMA as a resist is favored by researchers because of its high resolution and contrast [Jubinsky et al., 1986]. However, its sensitivity is too low for commercial use. Jubinsky et al. [1986] describe an approach to improve its sensitivity, by blending it with a radiationsensitive polymer, CO (poly(epichlorhydrin), HydrinR) (Table 11.9). Three blends were made by dissolving CO and PMMA in tetrahydrofuran (CO/PMMA: 20/80, 33/67 and 50/50), in the

851

presence of cyclohexanone (to optimize viscosity and volatility for casting films). The films were spun at 1250 rpm for 1 min on 7.5-cm silicon wafers, which were then baked in a forced-convection air oven for 1 h. Jubinsky et al. [1986] report that the sensitivity of the PMMA to X-rays was enhanced by the presence of CO in the blends. Films with 20-50% CO in the blend require only 50-25% dose of synchrotron radiation to produce patterns equivalent to those produced in PMMA films. 11.4.4.11 PMMA/PVDF and PMMA/PCTFE Blends Mishra et al. [1994] and Bajpai et al. [1994] determined the microhardness of PMMA/PVDF and PMMA/PCTFE blends (Table 11.9) made by solution casting from dimethyl formamide solutions. The solutions containing the two polymers were heated at 110°C for 3 h and were poured into an optically plain glass mold to prepare pellets of the blends. The poured material was annealed at 75°C for 3 h. The samples were cut from the slowly cooled (24 h) pellets for this work. At 10 kGy dose, the Hv values (Vickers hardness number) were higher for the irradiated blends (Table 11.63). The Hv values decreased with increasing fraction of PVDF, and also decreased with increasing dose beyond 10 kGy. The authors attributed the increase in the Hv values to crosslinking of PMMA. Though PVDF is known to crosslink [Zhong et al., 1993] on irradiation, PMMA is known to degrade on irradiation (Section 11.3). In general, similar results were obtained for the system PMMA/PCTFE. PCTFE is also expected to degrade on irradiation. Further work is needed to determine alternate mechanisms to explain the results reported in these papers. 11.4.4.12 Miscible Blends Lindberg et al. [1989] investigated the effect of electron irradiation (in an electron microscope operated at 200 keV, presumably under vacuum) on the blends of PVC with PMMA, PCL and

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Table 11.63. Microhardness of PMMA/PVDF blends [Mishra et al., 1994; Bajpai et al., 1994] PMMA (wt%)

PVDF (wt%)

Dose (kGy)

Hv (kg/mm2)

Load (g)

100 100 100 95 95 95 79 79 79 95 95 79 79

0 0 0 5 5 5 21 21 21 5 5 21 21

10 10 10 10 10 10 10 10 10 50 200 50 200

3.88 4.94 5.06 5.54 10.75 12.37 4.71 8.77 10.50 11.84 10.15 10.16 8.23

20 60 100 20 60 100 20 60 100 100 100 100 100

PMMA

PCTFE

99 99 99 94 94 94 90 90 90

1 1 1 6 6 6 10 10 10

10 30 40 10 30 40 10 30 40

SAN; these blends are miscible. Commercial grade PVC, PMMA, and SAN were used asreceived (Table 11.9). Polymer blends were made by using an internal mixer. Very thin samples (100 to 200 nm thick) were cut at room temperature using an ultramicrotome. The samples were mounted on copper or nylon grids and coated on one side with 5 to 20 nm of carbon to reduce charge and shrinkage during irradiation in the electron microscope. In the PVC/PMMA blend, the chlorine and the mass loss of the PVC increased in the presence of PMMA. Similarly, in the PVC/PCL blend also

100 100 100 100 100 100 100 100 100

the chlorine loss from the PVC was enhanced by the presence of PCL. In contrast, in the PVC/SAN blend, the dechlorination of the PVC seemed to be unaffected by the presence of SAN. 11.4.4.13 Blends of Plastic Wastes Numata and Fujii [1995] investigated the use of gamma irradiation in the presence of a crosslinking agent to improve the mechanical properties of blends made from plastics in municipal wastes. The plastics content of industrial and municipal wastes in Japan amounts to 6 MT annually.

853

Use of High-Energy Radiation in Polymer Blends Technology

Recycling of plastic wastes as polymer blends is attractive when sorting of different plastics is avoided. However, the mechanical properties of the resulting blends are generally poor. They also studied the free radical levels formed in the irradiated blends by ESR. The authors used virgin materials for their work (Table 11.9). Powdered TAC (2, 10 or 20%) was mixed with powdery polymer mixture (HDPE, 50; PP, 30; PSGP, 10; PSHI, 10 wt%) and test specimen were molded at 220°C with a manually operated injection molding machine. Test specimen of PE and PP were also similarly prepared, for comparison and were gamma irradiated. For ESR work, the molded specimen were crushed into small particles (~2 mm, diameter) before irradiation. The flexural strength and modulus of elasticity of the irradiated blends containing 10 and 20 wt% TAC were higher than that of the non-irradiated blend, and also higher than those of the constituent polymers (Table 11.64). The results show that irradiation (150-350 kGy) of such blends, in the presence of 10 wt% TAC, would improve the mechanical performance of the blends. Though the authors did not point this out, the increases in the modulus values and the free radical levels for blends were somewhat lower than those in the cases of PE and PP.

This is attributable to the protective effect of the PS in the blends. Further work on this interesting topic, using other crosslinking agents, mixtures of other plastics that need recycling, and electron irradiation, is warranted.

11.5

Concluding Remarks

Since Charlesby’s discovery of crosslinking of PE and the memory effect [1960], industrial irradiation of polymeric systems [Saunders, 1988; Singh and Silverman, 1992] has evolved into a multibillion dollar industry [Cook, 1990]. About 1000 electron accelerators and over 100 60Co gamma irradiators are being used to irradiate polymeric systems. Irradiation offers industry an array of options in modifying the properties of the products made from polymer blends, analogous to irradiation of individual polymers, as discussed in this chapter. In particular, the use of polymer blends is being made, or recommended, in medical products [Fennimore, 1984; Pal et al., 1992; Ramachandran and Dinunzi, 1996; Chen et al., 1996; Ding and Qin, 1996], wire and cable insulation [Sintic and Pasler, 1989; Tran, 1996], elastomers [Bohm and Nelson, 1981; Böhm and

Table 11.64. Improvement of flexural properties of polymer blends by addition of a crosslinking agent and gamma irradiation [Numata and Fujii, 1995] Dose (kGy)

0 100 150 200 350 400 600

Flexural strength (MPa) PE

PEa

PP

PPa

Blend

18.2 20.8 22.7 23.8

11.8 32.6 36.7 37.5

28.5 29.2 28.5 28.6 25.0

19.1 30.0 37.4 35.9 26.5 18.5

28.6 31.0 30.8 30.9 32.0

a containing 20% TAC b HDPE:PP:PSGP:PSHI = 50:30:10:10

Blenda,b 28.4 36.4 35.1 25.5

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A. Singh and K. Bahari

Tveekerm, 1982; Silverman, 1981; Singh and Silverman, 1992], magnetic media [Santosusso, 1985], and in crosslinked film and many PE products [Berejka, 1981; Saunders, 1988, Silverman, 1981; Singh and Siverman, 1992]. In all these products there is potential for a large increase in the use of irradiation to modify the properties of the products. Relatively low cost of radiation processing [Cleland, 1992; Feng and Ma, 1992], and the fact that the processing can be done at ambient temperature after the product has been formed, add to the appeal of this technology. The information on the chemical details of the effects of radiolysis decreases as one goes from pure chemical compounds to individual polymers and then to polymer blends. However, those working on blends can effectively use the information available for chemical compounds and individual polymers, at least for the preliminary design of their radiation processing work.

11.6

Acknowledgments

One of us (KB) would like to thank AECL for the hospitality shown during his short stay as a visiting scientist at Whiteshell. The authors thank Ms. Cheryl Rohrig for enthusiastic word processing of this manuscript. Helpful discussions with Drs. A. Berejka, G. R. Freeman and J. Silverman are gratefully acknowledged.

11.7

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Academic Press, New York (1973). Partridge, R. H., in The Radiation Chemistry of Macromolecules, Volume I, Dole, M., Ed., Academic Press, New York, (1972). Ramachandran, R., and Dinunzi, S. A., U. S. Pat., 5589545 A, to Union Carbide (1996). Randall, J. C., Zoepfl, F. J., and Silverman, J., Radiat. Phys. Chem., 22, 183 (1983); Org. Coating Appl. Polym. Sci. Proc., 48, 211 (1983). Raval, H., Devi, S., Singh, Y. P. and Mehta, M. H., Polymer, 32, 493 (1991). Rizzo, G., Spadaro, G., Acierno, D., and Calderaro, E., Radiat. Phys. Chem., 21, 349 (1983). Rolando, R. J. and Krueger, D. L., U.S. Pat. 5,140,073, to Minnesota Mining and Manufacturing Company, St. Paul, Minn. (1992). Rolando, R. J., Tappi Journal, 76, 167 (1993). Rosiak, J. M., in Radiation Effects on Polymers, Clough, R. L. and Shalaby, S. W. (Eds.), ACS Symposium Series 475, American Chemical Society, Washington, D.C. (1991). Salovey, R., in The Radiation Chemistry of Macromolecules, Vol. II, Dole, M., Ed., Academic Press, New York (1973). Santosusso, T. M., Radiat. Phys. Chem., 25, 557 (1985). Sasaki, T., Hosoi, F., Hagiwara, M., Araki, K., Saito, E., Ishitani, H., and Vesugi, K. Radiat. Phys. Chem., 14, 821 (1979). Sauer, Jr., M. C., Werst, D. W., Jonah, C. D., and Trifunac, A. D., Radiat. Phys. Chem., 37, 461 (1991). Saunders, C. B., Radiation Processing in the Plastics Industry: Current Commercial Applications, Atomic Energy of Canada Limited Report AECL-9569 (1988). Saunders, F. L. and Pelletier, R. R., U.S. Pat. 3,936,365, (1976); see also Canadian Pat. 1,078,990, (1980) that seems to describe the same work, to Dow Chemical Company, U.S.A. (1980). Scholes, G., and Simic, M., Nature (London), 202, 895 (1964). Schrijver, J., Hoen, N.G.M., Lemstra, P.J. and Van Gunst, C. A., EP Pat., 146,172, to STAMICARBON B.V., NL. (1988). Schroeder, K. F., Beta-Gamma, 1, 10 (1990). Schulz, A. R. and Mankin, G. I., J. Polym. Sci., Polym. Symp., 54, 341 (1976). Seguchi, T., and Yamamoto, Y., Diffusion and Solubility of Oxygen in γ-ray Irradiated Polymer Insulation Materials, Japan Atomic Energy Research Institute report, JAERI 1299, Takasaki, Japan, 1986. Shalaby, S. W., Macromol. Revs., 14, 419 (1979).

Silverman, J., J. Chem. Educ., 58, 168 (1981). Silverman, J. in Radiation Processing of Polymers, Singh, A., and Silverman J., Eds., Hanser Publishers, Munich (1992). Singh, A., Radiat. Res. Rev., 4, 1 (1972). Singh, A., in CRC Handbook of Free Radicals and Antioxidants in Biomedicine, Vol. 1, Miquel, J., Quintanilha, A. T. and Weber, H., Eds., CRC Press, Inc., Boca Raton, Florida (1989). Singh, A. and Silverman, J., Eds., Radiation Processing of Polymers, Hanser Publishers, Munich (1992). Singh, A. and Singh, H., in Isotopes and Radiation Technology in Industry, Rao, S. M., and Kulkarni, K. M., Eds., National Association for Applications of Radioisotopes and Radiation in Industry, BARC, Bombay (1994). Singh, A., in Proc. RadTech Asia’95, December 1995, Chulalongkorn University, Bangkok, (1995). Singh, A., Lopata, V. J., Kremers, W., and Sze, Y.-K., Crosslinking of Commercial Polyethylenes by 10 MeV Electrons, Atomic Energy of Canada Limited report AECL-11259 (1995). Singh, H., Dose Rate Effect in Food Irradiation: A Review, Atomic Energy of Canada Limited report AECL-10343 (1991). Singh, A., Saunders, C. B., Barnard, J. W., Lopata, V. J., Kremers, W., McDougall, T. E., Chung, M. and Tateishi, M., Radiat. Phys. Chem., 48, 153 (1996). Siquin, D. and Chen, W., Proc., 6th Japan-China bilateral Symp. Radiat. Chem., JAERI-Conf-95-003, p. 414 (1995). Siquin, D. and Chen, W., J. Appl. Polym. Sci., 62, 75 (1996). Sintic, K., and Pasler, N., Proc. PVC - Production, Processing, Application and Testing Conf., Yugoslavia, paper 46, p. 1 (1989). Soebiantio, Y. S., Katsumura, Y., Ishigure, K., Kubo, J., Hamakawa, S., Kudoh, H. and Seguchi, T., Radiat. Phys. Chem., 48, 449 (1996). Spadaro, G., Rizzo, G., Acierno, D., and Calderaro, E., Radiat. Phys. Chem., 23, 445 (1984). Spadaro, G., Acierno, D., Calderaro, E., and Valenza, A., Radiat. Phys. Chem., 40, 249 (1992). Spadaro, G., Valenza, A., Calderaro, E., and Brucato, V., Thermochim. Acta, 227, 75 (1993). Spadaro, G., Acierno, D., Dispenza, C., Calderaro, E. and Valenza, A., Radiat. Phys. Chem., 48, 207 (1996). Spenadel, L., Radiat. Phys. Chem., 14, 683 (1979). Spinks, J. W. T., and Woods, R. J., Introduction to Radiation Chemistry, 3rd edition, John Wiley & Sons, Inc., New York (1990).

Use of High-Energy Radiation in Polymer Blends Technology

Stannett V. T., Silverman, J., and Garnett, J. L., in Comprehensive Polymer Science, Vol. 4, Eastmond, G. C., Ledwith, A., Russo, S., and Sigwalt, P., Eds., Pergamon Press, Oxford (1989). Stone, J. A. and Dyne, P. J., Can. J. Chem., 42, 669 (1964). Stone, J. A., Quirt, A. R., and Miller, O. A., Can. J. Chem., 44, 1175 (1966). Tabata, Y., Beta-Gamma, 1, 6 (1990). Tabata, M., Sohma, J., Yokota, K., Yamaoka, H., and Matsuyama, T., Radiat. Phys. Chem., 36, 551 (1990). Tabata, Y., Ito, Y., and Tagawa, S., CRC Handbook of Radiation Chemistry, CRC Press, Boca Raton, Florida (1991). Takehisa, M., Levey, G., and Willard, J. E., J. Amer. Chem. Soc., 88, 5694 (1966). Tenorth, U., Beta-Gamma, 1, 32 (1990). Thomas, S., Gupta, B. R., De, S. K., and Thomas, K. T., Radiat. Phys. Chem., 28, 283 (1986). Thomas, S., Gupta, B. R. and De, S. K., Poly. Degrad. Stab., 18, 189 (1987). Tokuda, S., and Kemmotsu, T., Radiat. Phys. Chem., 46, 905 (1995). Torikai, A., Geetha, R., Nagaya, S., and Fueki, K., Polym. Degrad. Stab., 16, 199 (1986). Torikai, A., Harayama, K.-I., Hayashi, N., Mitsuoka, T. and Fueki, K., Radiat. Phys. Chem., 43, 493 (1994). Tran, C., Makromol. Chem., Macromol. Symp., 112, 137 (1996). Toy, L. T., Lau, A. N. K. and Leong, C.-K., U.S. Pat. WO 9,414,890 A1, to Raychem Corporation, Menlo Park, CA (1994). Valenza, A., Spadaro, G. and Acierno, D., SPE Techn. Pap., 38, 1895, (1992). Valenza, A., Gallo, L., Spadaro, G., Calderaro, E. and Acierno, D., Polym. Eng. Sci., 33, 1336 (1993a). Valenza, A., Spadaro, G., Calderaro, E. and Acierno, D., Polym. Eng. Sci., 33, 845 (1993b). Valenza, A., Calderaro, E., and Spadaro, G., Radiat. Phys. Chem., 43, 315 (1994). van Gisbergen, J. G. M., Meijerink, J. I., and Overbergh, N., in Proc. Integration of Fundamental Polymer Science and Technology, Lemstra, P. J. and Kleintjens, L. A., Eds., Elsevier Science Publishers, Barking (1989 a). van Gisbergen, J. G. M., Meijer, H. E. H., and Lemstra, P. J., Polymer, 30, 2153 (1989 b). van Gisbergen, J. G. M., Borgmans, C. P. J. H., van der Sanden, M. C. M. and Lemstra, P. J., Polym. Commun., 31, 162 (1990). van Gisbergen, J. G. M. and Meijer, H. E. H., J. Rheol.,

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35, 63 (1991). van Gisbergen, J. G. M., Hoeben, W. F. L. M. and Meijer, H. E. H., Polym. Eng. Sci., 31, 1539 (1991 a). van Gisbergen, J. G. M., Van der Sanden, M. C. M., de Haan, J. W., van de Ven, L. J. M., and Lemstra, P. J., Makromol. Chem., Macromol. Symp., 41, 153 (1991 b). van Gisbergen, J. and Overbergh, N., in Radiation Processing of Polymers, Singh, A. and Silverman, J., Eds., Hanser Publishers, Munich (1992). Vesely, D. and Finch, D. S., Makromol. Chem., Macromol. Symp., 16, 329 (1988). Walter, L., Hirayama, F., and Lipsky, S., Int. J. Radiat. Phys. Chem., 8, 237 (1976). Waterman, D. C., and Dole, M., J. Phys. Chem., 74, 1913 (1970). Wilfong, D. L. and Rolando, R. J., Canadian Pat., 2,083,750, to Minnesota Mining and Manufacturing Company, St. Paul, Minn. (1993). Williams, T. F., in Fundamental Processes in Radiation Chemistry, Ausloos, P., Ed., Interscience Publishers, New York (1968). Williams, J. L., Dunn, T. S., and Stannett, V. T., Radiat. Phys. Chem., 19, 291 (1982). Williams, J. L., in Radiation Processing of Polymers, Singh, A., and Silverman, J., Eds., Hanser Publishers, Munich (1992). Wilski, H., Radiat. Phys. Chem., 29, 1 (1987). Wilson, J. E., Radiation Chemistry of Monomers, Polymers, and Plastics, Marcel Dekker, Inc., New York (1974). Winslow, F. H., Hellman, M. U., Matreyek, W., and Stills, S. M., Polym. Eng. Sci., 6, 237 (1966). Witt, E., J. Polym. Sci., 41, 507 (1959). Woods, R. J., and Pikaev, A. K., Applied Radiation Chemistry: Radiation Processing, John Wiley & Sons, Inc., New York (1994). Xanthos, M., Polym. Eng. Sci., 28, 1392 (1988). Yoshimura, I., Hata, H. and Kaneko, T., U.S. Pat., 4,835,218, to ASAHI Dow, Japan (1989). Yu, D. W., Xanthos, M., and Gogos, C. G., Proc. ACS, Div. Polym. Mat.: Sci. Eng., 67, 313 (1992). Yu, D. W., Xanthos, M. and Gogos, C. G., J. Appl. Polym. Sci., 52, 99 (1994). Zhang, L., Zhang, W., Zhang, Z., Yu, L., Zhang, H., Qi, Y., and Chen, D., Radiat. Phys. Chem., 40, 501 (1992). Zhang, H. and Xu, J., Radiat. Phys. Chem., 42, 117 (1993). Zhong, X., Yu, L., Zhao, W., Zhang, Y. and Sun, J., Polym. Degrad. Stab., 39, 399 (1993).

CHAPTER 12

PROPERTIES AND PERFORMANCE OF POLYMER BLENDS

S. F. Xavier

Indian Petrochemicals Corporation Ltd., Vadodara, Gujarat, India

12.1

Abstract

This chapter presents an overview of properties and performance of polymer blends. It is structured into six sections dealing with aspects required for assessing the performance of a polymer blend. These are: 1. low speed mechanical properties, viz. tensile, compressive, and flexural strength, rigidity and hardness, fatigue failure, 2. high speed mechanical properties and fracture mechanics, viz. Izod and Charpy impact strengths, fracture toughness, toughening mechanisms (such as crazing, shear yielding and cavitation), 3. chemical and solvent effects, 4. thermal and thermodynamic properties, viz. thermal conductivity, heat capacity, heat deflection temperature, the thermodynamic interaction parameter, 5. flammability, 6. others, such as electrical and optical properties, viz. the methods of determining electrical conductivity, resistivity, break-down voltage, optical transparency, haze and refractive index of polymer blends, etc. Each section opens up with standard test methods such as ASTM, BS, DIN and ISO for each property evaluation and is summarized. Since presentation of all test methods for each property is beyond the scope of this chapter, one popular test method is described in detail while others are discussed with reference to it. The factors controlling each property are also examined. Each section concludes with an outline of the state of the art pertinent to the aspect in focus. Definitions of all terms from each section are grouped together in an Appendix. Toughening plays an important role in designing polymer blends. Due emphasis has been given to this aspect by presenting different methods of determining blends toughness, specially using ductile fracture mechanics, the mechanisms of toughening and also factors influencing toughness. Flammability aspect assumed a great deal of interest ever since the USA Federal Trade Commission’s (FTC) action in 1972. Commercial exploitation of a polymer blend is regulated, since then, by its flammability characteristics. A brief review on factors affecting flammability is presented and a list L.A. Utracki (Ed.), Polymer Blends Handbook, 861-950. © 2003 Kluwer Academic Publishers. Printed in the Netherlands.

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of fire retardant chemicals is provided in an Appendix. The chapter concludes with perspectives for the future developments. 12.2

Introduction

Modern technology thrusts challenging demands on the performance capabilities of materials, including polymers and their blends. A new approach to the science and technology of polymer blends has emerged recently, i.e., polymer blends by design, rather than by availability. These polymeric materials must perform under strenuous mechanical, chemical, thermal and electrical conditions imposed by the requirements of a specific application. Service in these applications usually involves several criteria to be fulfilled without a loss of economic advantage. Indeed, performance requirements of polymer blends are often at the limit of the properties that can be achieved. Moreover, these materials are expected to endure complex environmental conditions for extended time. All these factors stress the need for in-depth studies of the properties and performance of polymer blends. There are three aspects to material properties and performance: 1. The origin — identification of the mechanisms responsible for given performance characteristics. 2. Methods of determination — the most reliable way the properties should be measured. 3. The numerical values of the characteristic parameters of the material. In principle, the entire Handbook is dedicated to discussion of these three aspects. For example, in Chapter 2 on thermodynamics, the molecular aspects of polymer-polymer interactions, the methods of characterization, as well as the numerical values of the thermodynamic parameters are given. Similarly, in Chapter 3 on crystallization the three aspects vis-à-vis nucleation, melting etc. are presented. In Chapter 7 on flow, generation of flow-imposed lamellar morphology was discussed. This morphology has been used to control permeability through polymeric membranes.

Another important aspect of the material performance characteristics that is growing in significance, is the balance of properties — for many applications it is not so important what is the value of a single parameter, but how well the material combines a number of characteristic properties. Take, for example, the use of polymer alloys for automobile fenders. Here the general requirements (viz. weight reduction, part consolidation, cost, design flexibility, increased impact, crashworthiness), translate into requirements for: stiffness, strength, impact resistance, low coefficient of linear thermal expansion, weight reduction, chemical, corrosion and heat resistance, finish, oven paintability and cost. These aspects are well presented in Chapter 13 on the Application of Polymer Alloys and Blends. The third aspect of the properties, the numerical values of measured parameters represent enormous challenge to authors of a Chapter entitled “Properties and Performance”. The polymer blends industry produces well over 500 generic name blends, each in dozens of grades [Utracki, 1994, 1998]. Physically it is impossible and practically useless to attempt reproducing the datasheets of these thousands of blends. Furthermore, the industry is dynamic, continuously adding and/or removing grades from the market. The modern source of the numerical information must also be active, changing along with the variability of materials, viz. the Internet. Most of the major resin producers offer updated data sheets on the Internet, for example: www.allied.com, www.amoco.com, www.basf.com, www.dow.com, www.dupont.com, www.eastman.com, www.ge.com, www.hoechst.com, www.solutia.com, etc. This chapter presents an overview of properties and performance of polymer blends, focusing on these aspects that are outside the main domain of the other Chapters in this Handbook. Such properties as: mechanical, chemical and solvent effects, thermal, flame retardancy, electrical and optical properties are discussed. The response of a polymer blend to tensile, compressive and flexural stresses is examined in the initial section. Its rigidity, fatigue and failure characteristics are also studied. The toughened

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polymers have enhanced the status of polymer blends, thus the toughening mechanisms have received considerable attention from researchers. The fracture mechanics approach of testing and the fascinating toughening mechanisms prevailing in these wonderful materials are also examined. Insufficient chemical resistance of a blend at times leads to its rejection for use in an aggressive chemical environment, although it possesses an excellent combination of mechanical properties. Thus chemical and solvent effects on polymer blends are important factors that frequently determine blends’ applicability. Attention has been given to chemical resistance of blends starting from the fundamental concept of the solubility parameters. Apart from the chemical and environmental restrictions, thermal resistance of a polymer blend is often a major criterion for its applicability. Thus, the thermal conductivity, heat capacity and heat deflection temperature of polymeric materials are discussed in separate sections. Flame resistance has become a legal requirement for commercial utilization of polymers and their blends in many applications. Innumerable test methods for flammability have been developed in different countries and several books and handbooks are exclusively dealing with this subject. Discussion of the test methods that are en vogue in various countries is beyond the scope of this chapter, thus only the most popular test methods are discussed. The fire retardant chemicals and their suppliers are tabulated in Appendix 1. The use of polymer blends in electrical as well as electronic applications has been increasing rapidly. The electrical insulation properties of these materials cannot be ignored. Moreover, conducting polymers are likely to make an industrial breakthrough. These types of blends are also briefly discussed. Optical clarity has received considerable attention from the research community, as well as industry, especially, since transparent ABS was introduced to the market. Although success in this area has been limited, nevertheless this property frequently is pertinent when considering blend suitability to a particular application. The chapter concludes with a discussion of this aspect.

12.3

Low speed Mechanical Properties

12.3.1

Tensile Strength, Tensile Modulus, Elongation

The tensile stress-strain test is most widely used. Owing to the viscoelastic nature of polymers, the test is only a rough guide to how a polymer will behave in a finished product. Often results of a single test conducted at one temperature and speed of testing are published. To get a clear understanding of a polymer it is required to have the tests at several temperatures, rates of testing and other conditions [Dukes, 1966]. There is no universally accepted set of definitions with regard to the tensile tests. The terms listed in Appendix 2 have been taken from the widely accepted norms (ISO/DIS 527, BS 2782, Methods 320A to F, ASTM D638-72). The most frequently applied stress-strain measurements are made in tension by stretching the specimen as shown in Figure 12.1. A tensile stress can thus be defined as: σ1 = F1/Ao

(12.1)

where σ1 is the tensile stress, F1 is the tensile force, and Ao the cross-sectional area of the specimen. If the tensile stress stretches the specimen to length l1, the tensile strain, ε1, is defined as: ε1 = (l1-lo)/lo = kt/lo

(12.2)

where lo is the initial length of the specimen, k is the rate of extension and, t, is the time. Continuing the stressing operation to the ultimate, i.e., measuring the force until the material breaks, tensile strength, known as the ultimate tensile stress:

Figure 12.1. Material in tension.

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σ = F/A

(12.3)

where F is the force at failure and A is the crosssection at failure. During the process of stretching, the specimen’s dimensions orthogonal to the axis of applied force decrease and thus the area of cross-section decreases. For experimental convenience, however, tensile strengths are usually based on the original cross-section (Ao) which is easily measured at the beginning of the experiment. Elongation at break (%) = 100 (l - lo)/ lo

(12.4)

From the point of view of mechanical performance, four types of materials have been identified. They are best discussed in terms of the stress-strain dependence: 1. Brittle, showing proportionality between stress and strain up to the point of rupture. Here the modulus, E = σ/ε is constant, independent of strain, ε. 2. Semi-ductile, showing decreasing proportionality between stress and strain up to the point of rupture. Here the modulus, E = σ/ε decreases with strain, ε. 3. Ductile, initially showing similar relationship between stress and strain as the semi-ductile materials. However, these materials deform further, causing the stress pass through a maximum (yield). The rupture takes place at lower values than the yield stress. 4. Ductile with flow. These materials show still greater deformability than the typical ductile materials. Initially, the stress-strain dependence resembles that described for ductile resin, but before the rupture there is a zone of deformation where the stress remains about constant. Within this zone there is “flow” of material that usually leads to molecular alignment and/or to changes to the crystalline structure (viz. deformation of polyolefins). 12.3.1.1

Methods of Measurement

Stress-strain measurements for polymer blends can be conducted in one of two modes: using

a constant rate of loading or a constant rate of stretching. The first method is very often used in adhesives testing. The latter method is the most extensively employed in polymers and blends testing. In tensile testers, a sample is clamped between grips or jaws that are pulled apart at constant strain rates varying from 0.5 to 500 mm/min. The stress on the sample is monitored with the load cells ranging between 2 g to 5000 kg or more. The elongation must avoid errors arising out of sample slippage from the grips. There are a variety of jaws that can hold different samples. A review of grip systems is presented in ‘Handbook of Plastics Test Methods’ [Brown, 1981]. Jaw design and specimen shape and preparation are selected so as to minimize the introduction of extraneous stress or strain. Tensile tests on different polymer blends employ specimens of different sizes. To conduct a tensile test, a specimen capable of being gripped at both ends is required. The basic types of dumbbell configurations and dimensions recommended by ISO are illustrated in Figure 12.2. The American specifications differ only in number of dimensional details and are based essentially on Imperial units. In case of rigid polymer blends (such as engineering blends), the specimen can be molded, machined on a lathe, or simply cut out from thin, flat sheets. The dumbbell specimen with narrow-waist (Figure 12.2.a) is generally preferred in testing rubbers and their blends with thermoplastics. The dumbbell specimen with broad-waist (Figure 12.2.b) is used for polymers and blends exhibiting low-to-moderate elongation at break. Dog-bone test specimens (Figure 12.2.c) are used for polymer blends that do not allow any elongation measurement to be made. The parallel strip specimens (Figure 12.2.d) are used for reinforced thermoplastics or blends. In order to avoid the problem of sample fracturing near the grips, the end pieces of the grips are normally bonded.

Properties and Performance of Polymer Blends

12.3.1.2

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Standard Test Methods for Tensile Properties

The standard test methods for tensile properties are listed in Table 12.1 and the recommended test speeds in Table 12.2. 12.3.1.2.1 Measurements of Strain The elongation of specimen is followed by using gauge marks and measuring the distance between them preferably continuously, or by making use of clip-on type of extensometer. In the case of blends exhibiting strains in excess of 50%, optical extensometers are to be used. The merits and demerits involved in different methods of strain measurements are discussed in detail elsewhere [Brown, 1981]. 12.3.1.2.2 Measurement of Modulus The standard test methods, calculate the tensile modulus by drawing a tangent to the initial linear part of the stress-strain curve and calculating the slope of the line. In cases where no clearly defined linear portion exists, the ‘secant modulus’ should be determined. 12.3.2

Figure 12.2. (a) Tensile test specimen: Narrow-waisted dumbbell (ISO/DIS 527 Type A). (B) Tensile test specimen: Broad-waisted dumbbell (ISO/DIS 527 Type B). (c) Tensile test specimen: Dog-bone dumbbell (ISO/DIS 527 Type C). (d) Tensile test specimen: Parallel-sided strip (ISO/DIS 527 Type D).

Compressive Strength

Stress-strain curves developed during tensile, flexural and compression tests may be quite different from each other. The moduli determined in compression are generally higher than those determined in tension. Flaws and sub-microscopic cracks significantly influence the tensile properties of brittle polymeric materials. However, they do not play such an important role in compression tests as the stresses tend to close the cracks rather than open them. Thus, while tension tests are more characteristic of the defects in the material, compression tests are characteristic of the polymeric material as it is. The ratio of compressive strength to tensile strength in the case of polymers is in the range 1.5 to 4.0 [Dukes, 1966].

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Table 12.1.

Standard test methods for the determination of the tensile properties of polymer’s

Test method

Materials

Specimen

Test speeds

ISO R 1184

Plastics

See Table 12.2

BS 2782 Method 320A

Flexible plastic sheets Blends, filled or reinforced

1 mm thick < 1 mm thick Stamped from sheets Injection molded Machined from sheets

Method 320B Method 320C

Table 12.2.

Composites Rigid thermoplastic, thermosets

500 mm/min 1, 5, 25, 50 and 100 mm/min As in Method 320A

Recommended test speeds according to DIS 527 [ISO/DIS 527]

Speed A Speed Al Speed B Speed C Speed D* Speed E Speed F Speed G* Speed H

1 mm/min 2 mm/min 5 mm/min 10 mm/min 20 or 25 mm/min 50 mm/min 100 mm/min 200 or 250 mm/min 500 mm/min

± 50% ± 20% ± 20% ± 20% ± 10% ± 10% ± 10% ± 10% ± 10%

* Both the mentioned speeds are allowed because they are popularly used throughout the world.

12.3.2.1 Standard Methods for Compressive Tests The standard method for compressive tests are listed in Table 12.3. For example, the ISO Standard 604 allows four types of test specimens: 1) the right square prism, 2) the right rectangular prism, 3) the right cylinder and 4) the right circular crown tube. The test specifies for each of these test specimens, the load-bearing surfaces be parallel to each other within 0.1% of the height of the test piece.

12.3.2.2

Plain Strain Compression Test

Willams and Ford developed ‘plain strain compression test’ which was initially applied to metals [Williams and Ford, 1964]. It was based on the fact that strain is easier to measure in compression test. The same test method may be used for polymer blends to obtain total deformation curves up to high levels of strain that may be encountered in engineering applications. Williams had further explained the application of this technique to polymers [Williams, 1964].

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Table 12.3.

Standard test methods for compressive tests

No. Test method

Specimen

Materials

1.

ISO 604 (1973)

1) The right square prism, 2) the right rectangular prism, 3) the right cylinder and 4) the right circular crown tube.

The load bearing faces are to be parallel to each other within 0.1 % of the test piece height

2.

British standard BS 2782, Method 345A Type 1

For sheets of thickness not less than 12 mm.

Thermoplastics, including polymer blends and thermosets. Thermoplastics, including polymer blends, and thermosets.

Type 2

Type 3 Type 4 Type 5 3.

4.

USA standards ASTM D695-96* The method is the same as in BS, but specimen dimensions differ. German standards DIN 53454(1971)

For sheets of thickness less than 12 mm. (For sheets of thickness >12 mm, the test piece is machined only on one face so as to reduce the thickness to 12 mm). For casting and lamination resin systems without fibrous reinforcement. Test pieces are identical to ISO 604. For flat injection molded or compression molded thin sheets materials. For a right prism the preferred dimensions are 12.7 mm square by 25.4 mm height.

Neat resins. Glass fiber reinforced laminates. For rod slenderness ratio should be in the range 11 to 15:1.

Similar to British Standards in the form as well as dimensions of test specimens.

Note: * ASTM standard test methods are available on web: http://enterprise.astm.org/

12.3.3

Flexural Strength and Flexural Modulus

Flexural tests may be carried out in tensile or compression test machines. In standard tests, three-point bending test is preferred, although it develops maximum stress localized opposite the center point (support). If the material in this region is not representative of the whole, this may lead to some errors. Four-point test, offers equal stress distribution over the whole of the span between the inner two supports (points) and gives more realistic results for polymer blends (Figure 12.3). Expressions for the calculation of flexural strength and modulus for differently shaped specimens are given in Table 12.4.

Standard test methods for flexural properties are listed in Table 12.5. They may be carried out in tensile or compression test machines. Three-

Figure 12.3. Four point bending.

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Table 12.4.

S. F. Xavier

Expressions for calculation of flexural strength and modulus

No. Geometry 1.

Three-point bending ASTM D790-99 (support span-to-depth ratio < 16:1)

Table 12.5.

2.

3.

4.

S = PL/bd2

E = 0.21 L3m/bd3

S = 3PL/4bd2

E = 0.17 L3m/bd3

S = 3PY L/2bd2

E = L3m/4bd3

Comments P = Load (N) at the moment of break. PY = Maximum load (N) at yield or break. L = Support span (mm). b = Width of beam (mm). d = depth of beam (mm). m = slope of the secant (N/mm). • For Unreinforced and reinforced plastics. • For fiber reinforced pultruded rods use ASTM D4476-97.

Standard test methods for flexural properties

No. Test method 1.

Modulus (MPa)

Four-point bending ASTM D6109-97 (load span = 1/3 support span) ASTM D6272-98 (load span = 1/2 support span)

2.

Strength (MPa)

Specimen

Material

Standard test specimen dimensions: 80 x 10 ± 0.5 x 4 ± 0.2 mm. Length is20 times the thickness and width is between l0 and 25 mm. For single-phase materials. Length is 20 times the thickness and width is between For materials containing 20 and 50 mm. fillers. British standards Identical to ISO 178. It employs cantilever-bending mode. BS 2782 Method Standard test specimen (molded) dimensions are 335A Method 336B 70 x 25.4 x l.5 mm.A hole of diameter 2.0 to 2.02 mm is to be drilled centrally. USA standards Three-point loading system. Procedure-A is for materials ASTM* D790 that fracture at small Method I Test specimen dimensions and deflections. Procedure-B Procedure-A rate of cross-head motion are to be selected based on support is for materials that Procedure-B span-to-depth ratios (l/d = 16 to 1, 32 to 1, 40 to 1 or 60 to 1). undergo large deflections. Method II Procedure A Four-point loading system.Recommended test specimen Procedure B dimensions and rate of crosshead motion are given based on support span-to depth ratios. German standards DIN 53452 DIN 53435 Similar to three-point loading for ISO 178 For four-point loading (using Dynstat apparatus) ISO standards ISO 178

Note: *ASTM standard test methods are available on web: http://enterprise.astm.org/

point bending is often used. Four-point test offers equal stress distribution over the whole of the

span between the inner two supports (points) and it is preferred for polymer blends. The curvature

Properties and Performance of Polymer Blends

of the bearing rods is also important as too sharply curved rods lead to fracture of the specimen. 12.3.4

Rigidity and Rockwell Hardness

‘Hardness’ is not a fundamental property. Its measurement is subjected to the effects of temperature, time and other test variables. Hardness measurement of plastics is similar to the traditional methods applied to metals. It usually employs a standard indenter (often a hardened steel ball), forcing it under known load into a flat surface of the plastic, then measuring the resultant degree of indentation. Unlike other Rockwell scales, Rockwell parameter, Rα, correlates with the hardness as determined by ball indentation. Fett [1972] has shown that: H = [441.4/(150 - Rα)]1.23

(12.5)

where H = hardness by ball indentation, and Rα is Rockwell hardness parameter.

Table 12.6.

2.

3.

4.

The standard test methods for determining Rockwell hardness are listed in Table 12.6 and the Rockwell scales are given in Table 12.7. For example, ISO 2039 employs a hardened steel ball, 5 mm diameter. The ball is pressed into the specimen under a specified load selected to give an indentation between 0.07 and 0.10 mm (Method A) or between 0.15 and 0.35 mm (Method B). The recommended thickness of the specimen is 4 mm and the suggested time of application of the load is 30 s before the depth reading is taken. 12.3.5

Fatigue Characteristics

Fatigue failure may occur when a specimen fractured into two parts, was softened or its stiffness significantly is reduced by thermal heating and/or cracking. Sometimes, for different reasons, a large number of cycles elapses from the first formation of microscopic cracks to complete fracture. In this case, the fatigue failure is arbitrarily defined as having occurred when the specimen

Standard test methods for determining rockwell hardness

No. Test Method 1.

869

ISO standards: ISO 2039 Method A Method B Revised procedure follows that of ASTM D785-98. British standards BS 2782 Method 365A Method 365D Method 1001. Part 3 (Method 365C). USA standards ASTM D785-98* Procedure A Procedure B German standards DIN 53456

* See: http://enterprise.astm.org/

Specimen

4 mm thick specimen.

Indentation

Between 0.07 and 0.10 mm. Between 0. 15 and 0.35 mm. Time of load application is 30 s.

8 to 1 0 mm thickness For ‘Softness Number’. As ISO 2039. As ASTM D2583-95. Measures Rockwell Hardness. Uses Rockwell hardness tester, and scales (see Table 12.7). 6 mm thick specimen As in ISO 2039.

Minor load is applied for 10s, then major load for 15s; the hardness reading is taken off the scale 15s after the major load is removed. Indentation is recorded 15s after application of major load, but with minor load still on. Same as ISO 2039 except the major load must be selected from 49N, 132N, 358N and 96IN, with a minor load of 9.81N in all cases.

870

Table 12.7. Rockwell hardness

S. F. Xavier

Rockwell scales [ASTM D785-98 (http://enterprise.astm.org/)] Minor load (kg)

Major load (kg)

Indenter diameter (inch)

R L M E K

10 10 10 10 10

60 60 100 100 150

can no longer support the applied load within the deflection limits of the apparatus. Plastics, including polymer blends, are relatively high damping and low thermal conductivity materials. Thus, repeated straining of an article leads to a temperature rise within and throughout its body. Rapid stress-strain cycling can significantly heat up the article and thereby induce thermal failure — the phenomenon is frequencydependent. Where the thermal effect is to be minimized, much lower frequencies, of the order of a few Hz, should be employed. Fatigue data are usually presented in the form of S-N curves, in which stress amplitude S is plotted vs. log Nf, where Nf is the number of cycles to fracture an unnotched specimen, either in bending or in tension. A typical example of such a curve is shown in Figure 12.4. Here, the S-N curve for rubber-toughened PMMA provides the ‘endurance limit’, which is defined as the lowest stress amplitude at which fracture occurs. In general, most S-N curves flatten out at Nf = 107. Rubber toughening, in the case of styrene polymers, is found to reduce fatigue resistance and causes a decrease in the endurance limit [Sauer and Chen, 1983, 1984]. This is due to promotion of crazing and reduction of stresses by the rubber particles. The cyclic loading then degrades the crazes into cracks. Fatigue-induced deformation mechanisms can be studied by measuring the volume changes. Another sensitive method is to monitor the hysteresis loops under tension-compression loading.

0.5000 ± 0.0001 0.2500 ± 0.0001 0.2500 ± 0.0001 0.1250 ± 0.0001 0.1250 ± 0.0001

(mm) 12.7000 6.3500 6.3500 3.1750 3.1750

± 0.0025 ± 0.0025 ± 0.0025 ± 0.0025 ± 0.0025

This is illustrated in Figure 12.5. All specimens, initially show a small elliptical loop, indicating the viscoelastic response of the polymer [Bucknall and Stevens, 1980; Sauer and Chen, 1984; Bucknall and Marchetti, 1983]. Plots of tensile vs. compressive loop area reflect the proportions of shear yielding and crazing. This method has been used to detect the onset of crazing in fatigue tests on ABS. Following the trend in metallurgy, fracture mechanics techniques are now widely used to study fatigue in polymers. Standard test methods for fatigue testing are listed in Table 12.8. As an example, ASTM D671 makes use of the constant amplitude of force approach [Satake, 1970]. Two dumbbell-type test pieces, both of triangular form, are described.

Figure 12.4. S-N curve for rubber-toughened PMMA [Bucknall, 1988].

871

Properties and Performance of Polymer Blends

Figure 12.5. Hysteresis loops developed during fatigue tests of ABS and HIPS [Bucknall, 1988].

Table 12.8.

Standard methods for fatigue testing

No. Test method

Approach

Specimen

Frequency of testing

1.

USA Standards ASTM D671-93*

Constant amplitude of force

2.

German Standards DIN 53442

Constant amplitude of deformation

Two specimens of triangular Fixed frequency of 30 Hz form with rectangular cross-section Dumbbell shaped Variable frequency (tensile dumbbell) test specimens

* See: http://enterprise.astm.org/

12.3.6

They are to be selected according to the thickness and stress range over which the measurements are to be made. The apparatus described, operates only at a fixed frequency of 30 ± 5% Hz.

Low-speed Mechanical Properties of Blends

The low-speed stress-strain dependence for PS and HIPS is shown in Figure 12.6. These data well illustrate the change induced by incorporation of elastomeric particles into PS matrix. As shown,

872

S. F. Xavier

Figure 12.6. Tensile stress-strain curves for PS and HIPS [Bucknall, 1988].

upon toughening, PS brittle behavior changes into ductile with flow. The low-speed mechanical properties of polymer blends have been frequently used to discriminate between different formulations or methods of preparation. These tests have been often described in the literature. Examples of the results can be found in the references listed in Table 12.9. Measurements of tensile stress-strain behavior of polymer blends is essential [Borders et al., 1946; Satake, 1970; Holden et al., 1969; Charrier and Ranchouse, 1971]. The rubber-modified polymer absorbs considerably more energy, thus higher extension to break can be achieved. By contrast, an addition of rigid resin to ductile polymer enhances the modulus and the heat deflection temperature. These effects are best determined measuring the stress-strain dependence.

12.4

High-speed Mechanical Properties

The subject of impact strength of plastics has received considerable attention in official standards, material data sheets and literature at large.

The result of an impact test is basically no more than one point on the general curve of studying strength properties as a function of speed of testing. One advantage an impact test can offer is a ready measure of the actual energy required to break a test piece, which information can also be calculated from stress-strain diagrams in tensile or flexural tests with some effort. All materials tend to fracture if stressed severely enough. Some materials fracture more easily than others, and are thereby said to be “brittle”. Brittleness is the property of a material manifested by fracture without appreciable prior plastic deformation. In ductile fracture significant plastic flow occurs before fracture. Strain at fracture is more than a few per cent, unlike brittle fracture, and may be several hundred per cent. However, a sharp distinction cannot be made between brittle and ductile fracture since even in glassy materials some deformations take place. Further, a given material will fail in a brittle manner under some conditions and a ductile manner under other conditions. Thus, brittle fracture is favored by the low temperature, fast loading and when the state of stress approaches a uniform, i.e., triaxial or dilatational, state. Materials with low Tg are more likely to suffer ductile fracture, but the mode of fracture also depends on the fracture conditions. Furthermore, any structural feature that raises Tg such as bulky side groups or crosslinking, promotes brittle fracture. In a given material fracture may undergo brittle/ductile transition depending upon the testing temperature. At this transition temperature, the mode of fracture changes from brittle to ductile fracture [Kinloch and Young, 1983]. 12.4.1

Impact Strength

In many applications a satisfactory resistance to impact loading is an important performance requirement and, indeed, impact toughness is often the deciding factor in materials selection [Vincent, 1971; Bucknall et al., 1972; Turner, 1973; Bucknall, 1977; Reed, 1979; Savadori, 1985; Kinloch and Young, 1983; Brostow and Corneliussen, 1986; Havriliak et al., 1996]. For

Keitz et al., 1984; Kurauchi & Ohta, 1984; Koo et al., 1985;

PC/SAN

Tensile properties

Tensile properties

Mechanical properties

Fatigue behavior

Fatigue behavior

Fatigue behavior

Mechanical properties

Tensile & impact properties

PC/HDPE,

PC/LDPE & PC/PS PC/ABS

PA/PE/ionomer

Polymers/blends

Rubber-toughened

PMMA Rubber-toughened PS

HDPE/LCP

PP/LCP, PVC/LCP

Weber & Schoeps, 1985; Gregory et al., 1987; Chiang and Hwung,

Yamaoka et al., 1989

Seppala et al., 1992; Lee, 1988

modulus and HDT. Decreased elongation at break, and impact strength Mechanical properties varied with interphase adhesion and orientation of LCP domains

Sauer and Chen, 1983, 1984

the fatigue resistance & endurance limit Increased tensile strength,

Bucknall, 1988

Saur & Richardson, 1980; Andrews, 1969

Takemori, 1984; Radon 1980; Hertzberg & Manson, 1980;

Rubber toughening reduced

Figure 12.4

Review of fatigue behavior

and Han, 1984, 1985; Macknight et al., 1985; Han and Chuang, 1985

Fisa, 1991; Fairley, 1990; Fairley and Prud’homme, 1987; Chuang

elongation at break Reduction in particle size of dispersed phase

Suarez et al., 1984; Chiang and Hwung, 1987

Kunori and Geil, 1980

1987; Skochdopole et al., 1987; Quintens & Groeninckx, 1990 immiscible/miscible blends Tensile strength, modulus and

Mechanical properties of

on mechanical properties

Effect of blend composition

modulus etc. Tensile properties of different

PP/Crosslinked rubber

compositions

Dao, 1982

modulus, yield stress, etc.

D’Orazio et al., 1983

Stress-strain curves, strength,

compositions Tensile yield strength, tensile

Nolley et al., 1980; Utracki, 1995

Greco et al., 1980; Dumoulin, 1988

Robertson and Paul, 1973; Deanin and Sansone, 1978;

Bucknall, 1988

Tensile properties for different

See Figure 12.6

HDPE/PP + two EPR’s

Tensile properties

HDPE/PP

References

EPR acts as a compatibilizer

Tensile stress-strain

HIPS

Results

HDPE/PP/EPR

Test

Sources for low-speed mechanical properties data of polymer blends, examples

Blend

Table 12.9.

Properties and Performance of Polymer Blends

873

Mechanical properties

Mechanical properties

Mechanical properties

Tensile and other

PA-6/LCP

PA-12/LCP

Amorphous PA/LCP

PC/LCP

Flexural properties

Tensile properties

Flexural properties

Tensile & impact properties

Mechanical properties

Mechanical properties

Tensile & impact properties

PET/LCP

PET/LCP, PBT/LCP

PES/LCP

PPS/LCP

PSU/LCP

PEI/LCP

PEEK/LCP

mechanical properties

Test

continud

Blend

Table 12.9.

Siegmann et al., 1985

Isayev & Modic, 1987; Kiss, 1987; Weiss et al., 1989;

increased with increasing LCP content Tensile properties improved

Ramanathan et al., 1988; Seppala, et al., 1992; Nobile et al., 1990

modulus, but decreased strength Mechanical properties depended on

due to LCP content

Properties varied with anisotropy

Kiss, 1987; Cogswell, 1981; Mehta and Isayev, 1991

Blizard et al., 1990; Nobile et al., 1990; Kiss, 1987

Nobile et al., 1990

Kiss, 1987; Cogswell et al., 1981; Yazaki et al., 1994

as compatibilizer was used Increasing LCP content increased

miscibility, LCP orientation, etc.

Poli et al., 1996

Joseph et al., 1984; Sukhadia et al., 1990; Seppala et al., 1992

Zhuang et al., 1988; Amano and Nakagawa, 1987; Brostow et al., 1988;

Zhuang et al., 1988; Shin & Chung, 1989

PET-LCP copolymer

with LCP content

Flexural modulus increased

Kiss, 1987; Ramanathan et al., 1987; Blizard and Baird, 1986

with morphology Tensile strength and flexural modulus

Blizard & Baird, 1986; Malik et al., 1989; Blizard et al., 1990;

La Mantia et al., 1989

in terms of morphology Mechanical properties correlated

References

Changes in mechanical explained

Results

874 S. F. Xavier

Properties and Performance of Polymer Blends

the past several decades, ductile grades of usually brittle polymers, such as PS, PMMA, PVC or PA have been used. The most widely used test for impact assessment is the notched-Izod impact test. A single operator can run up to 150,000 tests in a year [Havriliak et al., 1996]. The test is also an important material or product specification for toughness — it is often one of few material constants specified as a product development objective. It appears that the earliest reference to the subject of impact testing was in 1734 by Swedenborg who wrote that iron bars were tested by throwing them against a sharp edge [Lethersich, 1948]. If the blow marks the bar without breaking any part of it, this was a sign of tenacity. Further experimental and theoretical works pointed out the dependence of the resistance of metals on the test speed and notches. Two devices were introduced by Charpy (in 1901) and by Izod (in 1903) for analyzing the impact performance of materials (see Figure 12.7). The impact resistance is evaluated in energy terms, i.e., by evaluating the difference between the potential energy before and after impact, the energy absorbed by the specimen during the impact process is obtained. From the physical point of view, the Izod equipment is equivalent to that of Charpy. However, the main differences between the two are: the clamping system, the notch, the hammer speed and its weight. Charpy adopted the keyhole form of notch. At such an early stage in the history of impact testing, Charpy found that correlations between static and dynamic tests were obtainable provided a notched bar was used. In 1925, the Izod and Charpy tests were extended to plastic materials and many results on plastics were published a year later [Werring, 1926]. The growth of fracture mechanics has placed greater emphasis on tests that use sharply notched specimens. These results were found to provide more fundamental information. Instrumented impact testing is a recent development that provides information on force-deflection curves. In addition to these notched bar tests, extensive use of falling dart tests is also being made. However, interpretation of the data from the latter is far from straight

875

Figure 12.7. Cantilever Beam (Izod Type) impact machine [ASTM D256].

forward. The impact strength is not a fundamental material property. The results depend on the specimen geometry, test method, and the employed parameters. Thus, it is difficult to correlate the results obtained from different test techniques and extremely difficult to correlate the results from impact tests on specimens of the material to the impact performance of the manufactured article. Correlation of results from one test to another for a given material becomes difficult because of different stress states of the specimen and the associated strain rates in different tests. In the tensile-impact test, the stress state is uniaxial and it measures the tensile property at a high strain rate. In Izod and Charpy tests, the presence of notch gives a triaxial state of stress. The falling-

876

S. F. Xavier

weight test is always in the forefront of high speed testing for evaluating the strain rate sensitivity of materials. In this case, the stress is biaxial. Several attempts were made to relate fracture mechanics theories to impact test results [Brown, 1973; Marshall et al., 1973; Plati and Williams, 1975]. The topic was also reviewed [Kinloch and Young, 1983; Brostow and Corneliussen, 1986]. 12.4.1.1

Izod Impact

One of the most often used tests for impact assessment is the notched-Izod impact test. The basic principle of the test is to allow a pendulum of known mass to fall from a known height and strike a standard specimen at the lowest point of its swing, and to record the height to which the pendulum continues its swing (see Figure 12.7). If the striking edge of the pendulum coincides with the center of the percussion of the pendulum, the bearings of the pendulum are frictionless and there is no loss of energy to windage, then the product of the mass of the pendulum and the difference between the fall distance and the height it reaches after impacting the test specimen is the impact strength of the latter. The test may be carried out on plain rectangular bars, but most often a carefully defined notch is molded or machined into the face to be struck (Figure 12.8). The impact tests are often regarded as a means of assessing the resistance of a material to shock where notches or ‘stress raisers’ generally are present. The ratio of impact strength of unnotched to that of notched specimen is sometimes regarded as a measure of the notch sensitivity of a material. Despite the popularity of the test, it is still poorly understood in terms of generating an actual ‘material’ property. The test reveals little about molecular dynamics and is not related to molecular structure. It is often criticized by fracture mechanics experts because of uncertainties about gauge length, complex state of stress, its dependence on thickness as well as a wide range of shear rates during the experiment, and the relationship of these factors to real situations [Havriliak, 1996].

Figure 12.8. Relationship of vise, specimen and striking edge to each other for Izod test methods A and C [ASTM D256].

The test, nevertheless, does have several important features. First, it is accepted by a large technical audience and is in common use. Second, it is a reproducible test, mostly because of the work of the ASTM. Finally, the impact results for various materials are spread over two orders of magnitude. When this spread is compared with the signal to noise ratio, the material range is impressive. Attempts are made to set up this ubiquitous test method on a firm platform based on the principles of fracture mechanics. Standard test methods for the determination of Izod impact strength are listed in Table 12.10. For example, ISO R180 normalizes the notch length. The velocity of the striker on impact has been standardized at 3.5 ± 10 per cent m/s with impact energies of 1.0, 2.75, 5.5, 11.0 and 22.0 J. Four types of test pieces are permitted. 12.4.1.2

Charpy Impact

The Charpy test is similar to the Izod impact test. In both the tests, flexural impact takes place by a pendulum (Figure 12.9) striking a barshaped test piece (Figure 12.10). However, as described before, there are quite significant dif-

877

Properties and Performance of Polymer Blends

Figure 12.9. Simple beam (Charpy-Type) impact machine [ASTM D6110-97].

Figure 12.10. Relationship of anvil, specimen and striking edge to each other for Charpy test method [ASTM D6110-97].

Table 12.10. Standard test methods for the determination of Izod impact strength No. Test method

Impact energies

Test specimen

Notch

1.

ISO R180

2.

BS 2782 Method 306 A

1.0, 2.75, 5.5, 11.0 and 22.0 J. 1.36, 4.07 and 13.6 J.

Four types permitted.Type 4 is preferred (80 ± 2 x 10 ± 0.2 x 4 ± 0.2 mm). (63.5± 2 x 12.7 ± 0.2 x 12.7 ± 0.3 mm) or (63.5 ± 2 x l2.7 ± 0.2 x 6.4 ± 0.3 mm)

3.

ASTM D256-00

A range of pendulum energies from 2.710 to 21.680 J.

Length: 63.50 mm (max.), 60.30 mm (min) Width: 12.7 mm (max), 3.00 mm (min). Breadth: 12.70 ± 0.15 mm.

Two types of cut notches allowed. Molded notch allowed in Type A. Cut notches allowed in Type B and C. Cut notches allowed.

them and no general correlation relating the data obtained from each have been developed.

Standard test methods for the determination of Charpy impact strength are listed in Table 12.11. For example, BS 2782 (method 351A) uses a

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S. F. Xavier

Table 12.11. Standard test methods for the determination of Charpy impact strength No. Test method

Impact energies

Test specimen

1.

BS 2782, Method 351 A

0.5, 1, 4, 15 and 50 J

Preferred test dimensions are: 120 x 15 x 10 mm

2.

ISO R179

Two striking energy levels

3.

DIN 53453 (Similar to BS Method) ASTM D6110-97

As in BS Method

4.

2.710 ± 0.135 J.

Type A (standard), square section, 2 mm wide and one-third specimen thickness in depth. Type B and C are V-shaped with base radii 0.25 mm and 1.00 mm respectively. The depth of these notches is set to one-fifth thickness. Four types of test pieces are allowed. Same notch types as above. Molded First three are as in British notches are permitted.Machined notches Standards. Fourth type is 125 mm are preferred. long by 13 mm square As in BS Method Type A as given in BS Method (Type B and C are not specified). (127.00 to 124.50) x (12.70 ± 0.15) The included angle of the notch is 45 ± 1°, x (12.70 to 3.00 mm) with a radius of curvature at the apex of 0.25 ± 0.05 mm.

rectangular, notched or unnotched, bar supported at both ends in such a position that the pendulum strikes it in the center, directly behind the notch. The energy absorbed in the impact is read directly by means of a pointer from a scale calibrated to allow for frictional and windage errors. Three standard test pieces are defined — the preferred being 120 x 15 x 10 mm with a span of 70 mm. The standard impact energies are 0.5, 1, 4, 15 and 50 J. 12.4.2

Notch

Fracture Mechanics

Griffith [1920] showed that brittle solid materials fail at lower strengths because of the presence of flaws acting as stress concentrators. The hypothesis has become the basis of ‘fracture mechanics’, used to interpret the fracture of many solids, including polymers and their blends. The theoretical background is presented in standard texts [Kinloch and Young, 1983; Williams, 1984; Broek, 1986; Brostow and Corneliussen, 1986]. Linear elastic fracture mechanics (LEFM), which has grown out of the work of Griffith,

provides the most satisfactory basis for characterizing the fracture process of polymer composites and multiphase polymers. It enables each of the factors contributing to fracture to be considered separately. Results obtained from the fracture mechanics analyses have thrown considerable light upon the behavior of polymers under tensile, impact and fatigue loading. Unfortunately, the conditions of LEFM are very severe and require that the fracture process is nearly totally elastic. Only under such conditions the data can be used for large-scale predictions. However, elastic behavior cannot always be obtained in a laboratory-scale test and some experimental or theoretical tools need to be available to help in predicting the brittleness of large-scale article. Furthermore, in case of rubber-toughened plastics, extensive yielding usually precedes fracture even in the presence of a sharp crack, so that LEFM techniques are unsuitable. A crack in a solid may be stressed in three different modes [Kinloch and Young, 1983; Brostow and Corneliussen, 1986]. The cleavage or tensileopening mode (Mode I) is technically the most

Properties and Performance of Polymer Blends

879

important one since it is commonly encountered and usually results in failure. Two closely related approaches have been used [Williams, 1984]: 1. The first is an energy criterion that supposes that fracture takes place when sufficient energy is released (from the stress field) during crack growth to supply the energy requirements of the new fracture surfaces created [Orowan, 1948]. The fracture of a material is thereby characterized by the material property Gc known as the ‘strain energy release rate’ or ‘fracture energy’ [Kinloch and Young, 1983; Brostow and Corneliussen, 1986]. 2. Rivlin and Thomas [1953] developed the second approach. They showed that the stress field around a sharp crack in an elastic material could be uniquely defined by a parameter known as the ‘stress intensity factor’, K. When K reaches a critical value Kc (which is a material property often called the ‘fracture toughness’), fracture takes place. The criterion for crack propagation is that KI > KIc. For plane strain in Mode I, values of GIc and KIc are related: K2Ic = E G Ic / (1 - ν2)

(12.6)

where E is Young’s modulus and ν is Poisson’s ratio. In SI units, KIc is usually given in MPa.m0.5, and GIc in kJ m-2. To make valid fracture mechanics measurements, it is necessary to ensure that specimen dimensions are large in comparison with the plastic zone surrounding the crack tip. For metals, according to ASTM E399: (w-a), a, B > 2.5 (KIc /σy)2

Figure 12.11. Schematic diagrams of various specimens used for fracture mechanics testing of flexible polymeric materials. P = Applied load [Young, 1988].

(12.7)

where w is the width of the specimen, a is the crack length, B is the thickness, and σy is the uniaxial yield stress of the specimen. In the case of polymer blends, it is preferable to experimentally determine the effects of specimen dimensions upon Gc and Kc rather than rely upon the applicability of the above conditions. The main experimental problem is to prepare specimens in which the plane strain/plane stress

conditions are satisfied [Williams, 1984; Kinloch and Young, 1983]. Apparent toughness values of Kc and Gc are higher in thin specimens than in thick ones. Measurements on several rubbertoughened plastics have shown a decrease in KIc with increasing thickness, B. The minimum value of B required, in the case of HIPS, for a valid determination of KIc at 23°C (296K) appears to be about four times higher than that given by Eq. 12.7 [Yap et al., 1983]. Unfortunately, there is no reliable criterion for crazing under plane strain, and therefore it is difficult to suggest an alternative to the standard approach that has been used successfully for metals. The toughness observed for rubber-toughened plastics is determined by the mechanism of deformation created by the plane-strain at the tip of a sharp crack. There are two factors to be considered:

880

S. F. Xavier

Table 12.12. Expressions for GC for fracture mechanics of crack growth in flexible polymers [Kinloch and Young, 1983] Geometry

Expressions for GC

Comments

References

Single-edge crack

GC = 2k1aWC k1 = πλc-1/2

Rivlin and Thomas, 1953; Greensmith, 1963; Lake, 1970.

Pure-shear Trouser-tear

Gc = lWC GC = (2PCλC/b) – 2wWC when λC = 1 GC = 2PC/b

λc = Extension ratio at onset of crack growth. WC = Critical stored elastic strain energy density. 1 = Initial length PC = Load at onset of crack growth w = Width of specimen arms λC = Critical extension ratio in arms WC = Strain-energy density in arms b = Specimen thickness and are critical extension ratios in regions A and B respectively; P1 and P2 are loads respectively transverse and in the split direction.

(See Fig.12.11)

Split-tear

Angled-tear

GC =

λAC + λBC 2b

(

P12C + P22C − P2 C

)

GC = (2PC/b)sin (α/2)

Rivlin and Thomas, 1953. Rivlin and Thomas, 1953.

Lake et al., 1979.

Thomas, 1960.

Table 12.13. Typical values of the fracture energy Gc and the fracture toughness Kc. for various materials [Kinloch and Young, 1983] Material Rubber Polyethylene Polystyrene High impact polystyrene PMMA Epoxy Rubber-toughened epoxy Glass-reinforced thermoset Glass Wood Aluminum alloy Steel-mild Steel-alloy

Young’s modulus, E (GPa)

Gc (kJm-2)

0.001 0.15 3.0 2.1 2.5 2.8 2.4 7.0 70 2.1 68 210 210

13 20 0.4 15.8 0.5 0.1 2.0 7.0 0.007 0.12 20 12 107

1. The fall in yield stress that occurs in ductile polymers as the temperature increases or the

Kc (MNm-3/2) --1.1 -1.1 0.5 2.2 7.0 0.7 0.5 37 50 150

strain rate reduces. This fall results in a planestrain to plane-stress transition, as indicated in

Properties and Performance of Polymer Blends

Eq. 12.7 and therefore an increase in fracture resistance. 2. The release of constraint that results from cavitation in the matrix or void formation in the rubber particles [Bucknall, 1988; Young, 1988]. Linear elastic fracture mechanics studies on toughened brittle plastics at room temperature concentrated on thermosetting resins, which have sufficiently high yield stresses to meet the requirements of Eq. 12.7. There has been increasing emphasis on ductile fracture mechanics in testing the toughened thermoplastics. An alternative approach is to determine the parameter, JIc, which is the quantity corresponding to GIc in linear elastic fracture mechanics, as discussed below. 12.4.3 12.4.3.1

Fracture Mechanics Testing Determination of Gc

The ductile fracture mechanics is gaining importance in testing polymer blends, especially the toughened thermoplastics. Some of the types of specimen which have been used to study the failure of ductile polymers whose deformation is elastic (but non-linear) are shown in Figure 12.11 [Kinloch and Young, 1983]. Formulae have been developed to determine Gc for these specimens and examples are given in Table 12.12. Again, the parameter defining the fracture process Gc, is a function of applied load, crack length and geometrical factor [Young, 1988]. Typical values of Gc and Kc are listed in Table 12.13. The problem of defining conditions for crack propagation becomes more difficult when the material is sufficiently ductile to form a large plastic zone at the crack tip. The problem is encountered particularly while testing rubbermodified plastics. Two ductile fracture criteria have been developed for metals, one based on crack tip opening displacement (CTOD), and the other upon the energy line integral (J-integral) around the crack tip. The crack opening displacement (COD) criterion has been applied to a number of polymers and blends, but the J-integral method is receiving more attention. Physically,

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COD is measured between the outer edges of the crack whereas CTOD is defined as the distance between two crack walls at the end of the fatigue crack. Thus, while COD is easy to measure, the determination of CTOD is not simple. The CTOD δ is the relative displacement of the two fracture surfaces at the crack tip. Critical values of CTOD δc may be measured by means of a mechanical clip gauge or recorded photographically. Furthermore, a cine camera may be used to follow the crack initiation and propagation as in HIPS (at 293 K) [Ferguson et al., 1973]. The authors reported that both LEFM and CTOD criteria were applicable at different stages of the fracture. On application of load to the specimen, the crack began to extend at a fixed value of KI, to give a value of KIc, for crack initiation. Then, at the crack tip, began to form a stress-whitened yield zone. The load on the specimen continued to increase as both crack and yield zone extended, and a load maximum was observed at a fixed value of CTOD. Precise determination of CTOD is often difficult. Furthermore, these measurements are unsuitable for use in design. For these reasons, models that enable CTOD to be calculated in terms of stresses have been developed [Dugdale, 1960]. In a wide plate with a central crack of length 2a with a narrow planar plastic zone of length L, extending from each of the crack tips, the applied stress σ is given by [Bucknall, 1978]: σ = a/(a + L) = cos (πσ/2 σy)

(12.8)

The CTOD δ can be expressed as: δ = [8 σy a/ (πE)] ln sec (πσ/2 σy)

(12.9)

For small values of applied stress (σ < 0.3σy), the plastic zone is small compared with the crack length and Eq. 12.9 can be simplified to read: δ = πσ2 a /(Eσy)

(12.10)

Under these conditions, LEFM analysis is applicable to the specimen.

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In the case of a brittle fracture in a wide and thick plate containing an edge crack of length a, the critical applied stress at fracture σc is related to the Young’s modulus E, the Poisson’s ratio, ν, and the fracture surface energy GIc, the critical stress can be expressed from the Griffith equation as: σ2c = E GIc /πa(1 - ν2)

(12.11)

From Eqs 12.10 and 12.11, for plane strain deformation: GIc = σy δc (1 - ν2) (12.12) Tests are conducted, in normal practice, on compact tension specimens rather than wide center-notched plates. Some allowance must be made for geometrical effects, including the finite width of the specimen, the difference between edge and center notches, and any rotations occurring at the grips. In general, it is not possible to measure CTOD, but rather the crack-opening displacement (COD). The latter quantity can be determined at the outer end of the notch with a suitable clip gauge. Thus, for a notched three-point-bend specimen, it was shown that a “plastic hinge” can form around the tip of the crack [Brostow and Corneliussen, 1986]. If the center of rotation is known, the CTOD can be calculated from the measured COD. A Standard has been published [BS 5762]. 12.4.3.2

demonstrated that the value of J was independent of the integration path. In practice, J can be determined from changes in load displacement diagram with changes in crack length [ASTM E813]. This method has been reviewed [Williams, 1984; Landes and Begley, 1979; Pascoe, 1986]. The J-integral is given by [Kinloch and Young, 1983]:

J-Integral Techniques

12.4.3.2.1 Single Specimen Method If there is extensive plasticity in a material under tension, it gives rise to a relatively large size of the crack-tip plastic zone and solutions for elastic-plastic behavior are not readily available. J-integral provides a means of determining the energy release rate for such cases. This integral was applied to crack problems [Cherepynov, 1967; Rice, 1968]. Rice showed that J-integral describes the flow of energy into the crack tip region and that the dominant term in the description of stress and strain singularities at the crack tip could also be written in terms of J. He

  ∂u   J = ∫  Z∂y − T   ds   ∂x    Γ

(12.13)

where Γ is a closed contour in a stressed solid (Figure 12.12), T is the tension vector perpendicular to the contour in an outward direction, u is the component of displacement of the contour in x- direction, ds is an element of the contour Γ, and Z is the strain energy (plastic and elastic) per unit volume. In Figure 12.13 a closed contour is taken as two curves surrounding the tip of the crack, one DEF inside the other ABC which are joined by two portions of the crack surface AF and CD. The integral around the contour is zero. Along the parts AF and CD which lie parallel to the X-axis and which have no normal stress on them, T = 0 and dy = 0. Therefore, the integral along ABC is equal and opposite in sign to that along DEF. For outward directed vectors, T, therefore, the integral is path independent. The J-integral method [ASTM E813] of fracture toughness measurement has been applied to a variety of polymers [Theuer et al., 1988; Rimnac et al., 1988] and rubber-toughened polymers [Huang and Williams, 1987; Huang, 1988, 1989; Hashemi and Williams, 1985; Takemori and Narisawa, 1989]. The multi-specimen J-integral technique [ASTM E813] also provides a method for determination of JIc, a measure of fracture toughness. A critical evaluation of ASTM E813-81 and E813-87 has been published [Narisawa and Takemori, 1989; Huang, 1989, 1990, 1991]. 12.4.3.2.2 Multiple-Specimen Method A major problem often encountered in the above described ‘single-specimen method’ is that the crack growth measured from a side-view may not

Properties and Performance of Polymer Blends

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1983; Hashemi and Williams, 1986; So and Broutman, 1986; Huang and Williams, 1987; Narisawa, 1987; Rimnac et al., 1988; Narisawa and Takemori, 1989; Huang and Williams, 1990; Huang, 1990; Moskala and Tant, 1990].

Figure 12.12. Contour for definition of J-integral [Brostow and Corneliussen, 1986].

Figure 12.13. Contour surrounding crack-tip Brostow and Corneliussen, 1986].

be accurate, as the crack front may vary from the central region to the sides. A ‘multiple-specimen method’ was developed to bypass this problem. The method has been applied to numerous ductile polymeric materials [Begley and Landes, 1972; Landes and Begley, 1974]. The critical J values obtained by using single specimen method were greater than those obtained from the standard multiple-specimen method [Westerlind et al., 1991]. Many workers have used the ASTM standards of E813-87 to characterize the fracture toughness of polymers [Chan and Williams, 1981,

12.4.3.2.3 Hysteresis Energy Method When a pre-cracked specimen of a toughened polymer is under load, viscoelastic and inelastic micro-mechanisms such as crazing, cavitation, debonding and shear yielding are expected to take place mainly around the crack tip. These micro-mechanisms occur during the process of crack tip blunting (pre-crack) and during crack propagation. A portion of the storage energy is therefore consumed and a relatively large crack tip plastic zone is formed, which can be quantified by the corresponding hysteresis energy. For rubber-toughened polymeric materials, the crack tends to propagate within the plastic zone. A new J-integral method based on hysteresis properties of polymeric materials was proposed [Lee and Chang, 1992; Lee et al., 1992]. In case of a cracked specimen, the material surrounding the crack tip can be divided into three parts: (1) the first plastic zone, (2) the second plastic zone, and (3) the elastic fracture surface [Lu et al., 1996]. The specific energy balance equation for a cracked specimen can be expressed as: (1/B) (dU/da - dUe/da - dUk/da) = (1/B) (dUpppz/da + dUpspz/da) + 2 γs

(12.14)

where U is the input energy at different displacements, Ue is the elastic energy, Uk is the kinetic energy, Upppz is the plastic energy for the primary plastic zone, Upspz is the plastic energy for the secondary plastic zone; and γs is the fracture surface energy. The energy dissipated of the system is given by: (1/B) d(HE)/da = (1/B) (dUpppz/da + dUpspz/da) + 2 γs where HE is the hysteresis energy.

(12.15)

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This approach assumes that there is a region surrounding the crack tip with local energy dissipation. This arises from viscoelasticity, plasticity and bond rupture and can be considered the characteristic of the fracture process. For polymers the characteristic of this localized energy dissipation is considered to be independent of geometries. The JIc values obtained based on hysteresis energy method were close to these obtained using E813-81 method, but significantly lower than those from E813-87 method. Experimentally the hysteresis energy method is relatively simple, because the tedious measurement of crack growth length is not necessary. Figure 12.14 shows variations of Jc values for PC/ABS blends obtained using different J-integral methods — as shown, the spread is ± 40% [Lu et al., 1996].

Figure 12.14. Plots of critical JIC value vs. the strain rate [Lu, et. al., 1996].

The J-integral by hysteresis energy method was applied to elastomer-modified PC [Lee and Chang, 1992; Lee et al., 1992], HIPS [Lee et al., 1992, 1993], ABS [Lu et al., 1995], PC/ABS blend [Lu and Chang, 1995; Lu et al., 1996], and PC/PBT blend [Lu and Chang, 1995]. 12.4.3.2.4 Essential Work of Fracture Method The theoretical analysis of J-integral are well

established [Rice, 1968; Begley and Landes, 1972; Landes and Begley, 1974] and the experimental procedure is standardized [ASTM E813-89]. However, some aspects of the method still remain controversial [Hashemi and Williams, 1986; Huang and Williams, 1987, 1990; Narisawa and Takemori, 1989, 1990; Swei et al., 1991]. For example, the procedure for J-R curve construction restricts the application of the J-integral method to only static loading tests. J-integral method is usually difficult and expensive [Bramuzzo, 1989]. The specimen size required by the J-integral method makes it impossible to characterize the toughness of polymeric thin films. It is also well recognized that using a blunting line to define the critical value of J-integral may not be proper for some ductile polymers [Hashemi and Williams, 1986; Narisawa and Takemori, 1989]. The J-integral analysis based mainly on metals is not fully appropriate to polymers. This is particularly true when the heterogeneous and toughened polymer blends are involved. To overcome the above drawbacks, a new method based on “essential work of fracture” concept was introduced [Broberg, 1971, 1975]. In this method, it is proposed that when a cracked ductile solid, such as a toughened polymer blend is loaded, the fracture process and the plastic deformation take place in two different regions, viz. the inner process zone and the outer plastic zone. Much of the fracture work during crack propagation, dissipated in the plastic zone, is not directly associated with the fracture process. Only that work that goes into the fracture process zone is a material constant. Hence, the total fracture work, Wf, should be separated into two parts, i.e., the essential work of fracture (i.e., the work required to create two new fracture surfaces, We), and a non-essential work of fracture (Wp): Wf = We + Wp

(12.16)

We is essentially a surface energy, and for a given thickness it is proportional to ligament length (l = W - a), while Wp is a volume energy and proportional to l2. Thus, the total fracture work is rewritten as:

Properties and Performance of Polymer Blends

Wf = wetl + βwptl2

(12.17)

where we and wp are the specific essential work of fracture and nonessential work of fracture (or specific plastic work), respectively, β is the plastic zone shape factor, while t, W and a are thickness, width and initial crack length. Then the specific total fracture work, wf is: wf = Wf /tl = we+ βwpl

(12.18)

There are two kinds of specific essential work of fracture available, according to the stress state of the ligament area, viz. plane-stress specific essential work of fracture (we) and plane-strain specific essential work of fracture (wIc) [Wu and Mai, 1996]. The we can be obtained if l/t ratio is large enough to ensure plane-stress condition in the ligament area and it is proved to be a material constant for a given sheet thickness [Mai and Cotterell, 1986; Mai et al., 1987; Mai and Powell, 1991]. With a reduction of l/t ratio, plastic constraint increases and the plane-stress/planestrain fracture transition may occur at a certain l/t ratio. Theoretical analysis shows that the specific essential work of fracture method is equivalent to the J-integral method for all three fracture modes [Mai and Powell, 1991; Mai, 1993]. The essential work of fracture approach has been applied to characterize the fracture properties of toughened polymer blends, such as PBT/PC/IM (where IM is the impact modifier) and ABS/PC. It is successfully used to determine the fracture toughness of a ductile LLDPE film [Wu and Mai, 1996], and single edge and double edge notched different polymeric films [Hashemi, 1993; Hashemi and Yuan, 1994; Chan and Williams, 1994; Karger-Kocsis and Czigany, 1996; KargerKocsis and Varga, 1996]. It is also applied to study the effect of specimen size, geometry and rate of tests in case of PBT/PC blends [Hashemi, 1997]. 12.4.4

Mechanisms of Toughening

Early investigations of the fracture of solids assumed that fracture involved only the creation

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of new surfaces [Griffith, 1920; Kinloch and Young, 1983]. However, since measured values of Gc were well in excess of the surface energy of the material, it was soon realized that significant amounts of energy were also dissipated through other processes such as localized plastic deformation in the vicinity of the crack [Kinloch and Young, 1983]. In general, two mechanisms are responsible for this plastic deformation in rigid polymers; namely ‘crazing’ [Kausch, 1983; Kambour, 1973] and ‘shear yielding’ [Ward, 1983]. The two mechanisms are not mutually exclusive. Under certain conditions both operate simultaneously. 12.4.4.1

Crazing

Crazing is an important source of toughness in rubber-modified thermoplastics. A craze can be described as a layer of polymer a nanometer to a few micrometers thick, which has undergone plastic deformation approximately in the direction normal to the craze plane as a response to tension applied in this direction [Kambour, 1986]. Crazing occurs without lateral contraction. As a result, the polymer volume fraction in the craze is proportional to 1/λ, where λ is the draw ratio in the craze. The reduction in density occurs on such a small scale that the refractive index is markedly reduced, which accounts for the reflectivity of the craze [Kramer, 1983]. Several methods of studies have been developed. Osmium-staining technique, pioneered by Kato [1967], is one of the most successful methods for observing crazing in rubber toughened plastics. It depends upon a reaction between osmium tetroxide, OsO4, and double bonds in PBD and other unsaturated polymers. However, it is not suitable for saturated rubbers. Ruthenium tetroxide, RuO4, is more reactive staining agent that can be used to differentiate between rubber and matrix when the former is essentially saturated. For example, clean glass slides (50 x 10 mm) were dipped in 2 wt% solutions of polymer and the solvent was subsequently evaporated under vacuum at 50°C (323 K) for 24 h. The films were removed from the glass

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substrate by immersing the slides in distilled water and then lifting the floating film from the water surface onto copper microscope grids. A 0.5 wt% solution of RuO4 in distilled deionized water was used for staining. The aqueous solution, (golden yellow when fresh), was found to be effective for a considerable time (up to 6 months if kept in a firmly sealed glass container in a freezer). Film-covered grids were vapor stained in a glass-covered dish [Trent et al., 1981, 1983]. Transmission electron micrographs, TEM, can be taken to illustrate detailed morphological features (at an accelerating voltage of 80 KV). Recently, methods have been developed to characterize the relative amount of crazing and non-crazing that occurs in rubber toughened glassy polymers, using the invariant obtained from small angle X-ray scattering (SAXS) analysis. SAXS not only overcomes the disadvantages of transmission electron microscopy (e.g., the use of ultra-thin samples), but the use of high intensity synchrotron radiation permits in-situ deformation studies [Paredes and Fischer, 1979; Brown and Kramer, 1981]. The new method of SAXS analysis leads to quantification of the contribution from crazing and non-crazing to the total deformation [He et al., 1998]. 12.4.4.2

Shear Yielding

Yielding is a mechanism, in which a thin layer of polymer deforms in shear at constant volume. It is characterized by regions of sheared polymer oriented approximately at 45°C to the tensile or compression stress. Unlike crazing, shear flow is essentially a process continuous in space, i.e., one that may spread through a much greater volume fraction of the stressed body and thus consume much more energy in total. Shear yielding is much less sensitive to environmental effects. In short, shear deformation is better than crazing, but crazing is better than no deformation at all [Kambour, 1986]. Shear bands are highly birefringent and are most clearly observed in transmitted polarized light [Bucknall, 1977]. They are also visible as reflecting planes in ordinary transmitted light at glancing

incidence, owing to refractive index differences between the band and the adjacent undeformed polymer [Kramer, 1974, 1975]. Both crazing and shear yielding involve the absorption of energy and most methods of toughening polymers involve modifying the polymer such that more crazing and shear yielding take place. The rubber-modified polymer absorbs considerably more energy in a tensile test because of its higher extension to break, which can be achieved only as a result of yielding in the matrix. The rubber particles play only a secondary role but, nevertheless, a vital one. Firstly, they accelerate yielding by acting as stress concentrators initiating deformation in the matrix; secondly, they respond to the hydrostatic component of stress by cavitating and increasing in volume, thus allowing the strain in the matrix to increase; and thirdly, in their cavitated and extended state, they stabilize the yielded polymer by carrying a share of the applied stress [Bucknall, 1988]. All three functions appear to be necessary for effective toughening, although their relative importance varies, depending upon the mechanisms contributing to toughening, and the kinetics of deformation, which in turn depend upon the material and the type of loading. Various types of response of the rubber particles have been observed experimentally as the polymer yields. They include (a) debonding between rubber and matrix [Haward and Bucknall, 1976], (b) cavitation within the particle [Breuer et al., 1977; Kinloch, 1985; Yee and Pearson, 1986], (c) craze like fibrillation of the rubber phase [Beahan et al., 1976; Donalad and Kramer, 1982] and (d) crazing within sub-inclusions [Seward, 1970]. Many polymers are toughened by blending or copolymerizing with a rubber [Kinloch and Young, 1983; Bucknall, 1977]. This method of toughening is now well established for many thermoplastics, thermosets and even adhesives. 12.4.4.3

Other Mechanisms

A new route for achieving a substantial lowering of stresses for craze growth without relying on potent craze initiators, involves controlled local plasticization of a polymer by a low molecular

Properties and Performance of Polymer Blends

weight diluent, distributed in a heterogeneous fashion throughout the material [Gebizlioglu et al., 1990; Argon and Cohen, 1990]. This mechanism is schematically shown in Figure 12.15.

[Bucknall, 1977]. The impact modified materials are usually the blends of a rigid matrix polymer with an elastomer. The composition of the constituents, their miscibility and the morphology influence the deformation and failure mechanism in the blend. Particle size of the elastomer, its dispersion and its adhesion with matrix are also the important factors determining the toughness. 12.4.5.1

Figure 12.15. Schematic rendering of craze moving in a field of encapsulated PB pools draining their content onto the craze surfaces when tapped by the advancing craze [Argon et al., 1990].

The advancing craze, nucleated from free surfaces or other occasional imperfections cuts into the dispersed population of PB-2.76 K pools (which at this low molecular weight acts like a relatively low viscosity liquid) and drains the contents of these pools onto the surfaces of the craze. Although the solubility of the PB 2.76 K into PS under standard conditions of room temperature and atmospheric pressure is negligibly small (of the order of 4 × 10-3) this should increase greatly in the presence of a negative pressure [Argon and Cohen, 1990]. The plasticization due to the increased sorption of the low molecular weight PB diluent into the craze surfaces is a highly interactive and complex process. The new mechanism offers considerable promise for practical industrial applications as only very small quantities of the plasticizing substance are needed and thus subsidiary properties such as optical transparency and tensile modulus are less affected compared to the other methods of toughening. 12.4.5

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Composition

The composition of individual constituents of a blend plays an important role in modifying the impact strength of the blend. The impact strength of polypropylene block copolymer (PPBC) blends with different concentrations of EPDM is shown in Figure 12.16 [Xavier et al., 1994]. Upon incorporation of the elastomer, the impact strength increases. EPDM was found to reduce the crystallinity of PPBC and significantly influence its failure mechanism. Both crazing and shear yielding were found to be responsible for the observed increase in impact strength. As shown in Figure 12.16, above 10 wt% of EPDM the increase in impact strength is more prominent. However, it was observed that such significant rise in impact strength adversely affected other mechanical properties, such as flexural and tensile moduli of the blends.

Factors Affecting Blend Toughness

Rubber toughening is the most often used method of improving the impact resistance of polymers

Figure 12.16. Influence of EPDM concentration on Izod impact strength of PPBC-EPDM blends.

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In the case of PVC/ABS blend, the addition of ABS improved the impact strength of the blend [Sharma et al., 1988]. At low concentrations of ABS a small number of rubber particles (i.e., the butadiene particles in ABS), is insufficient to significantly improve the impact strength (Figure 12.17). Increasing ABS concentration up to 50 wt% increased impact strength. The maximum impact strength obtained at the optimum blend composition is considerably higher than that of neat ABS itself. Since the particle size [Kulshreshtha et al., 1989] of the dispersed PBD (of ABS) phase is unlikely to change with blend composition it is evident that there exists a critical volume fraction of rubber phase necessary for the maximum improvement in impact strength. When this critical concentration of rubber (or ABS) is exceeded, impact strength drops.

12.4.5.2

Morphology

The performance characteristics of a blend depend on its morphology, which in turn depends on the thermodynamic and rheological properties of the components [Plochocki, 1983; Karger-Kocsis et al., 1984; Howe and Wolkowicz, 1987; Wu, 1987; Utracki, 1989]. However, due to nonequilibrium nature of the highly viscous polymer mixtures, often the processing conditions strongly influence the product morphology. The topic is discussed in the last part of Section 12.3.6. Further details of the morphology-processing conditions can be found in Chapters: 7. Rheology, 8. Morphology, 9. Compounding, and 10. Forming in this Handbook. The properties of PPBC/EPDM blends strongly depend on the crystalline micro-morphology of PPBC, as well as on the particle size and degree of dispersion of EPDM [Xavier et al., 1994]. The DSC cooling thermograms indicated that the degree of crystallinity in PPBC decreased with increasing concentration of EPDM (Table 12.14 and Figure 12.18).

Table 12.14. Crystallinity indices (A/m values) from DSC [Xavier et al., 1994] No. Blend 1. 2. 3. 4. 5.

Figure 12.17. Influence of PVC/ABS blend composition on Izod impact strength.

PPBC (neat polymer) PPBC + EPDM (2.5 wt%) PPBC + EPDM (5.0 wt%) PPBC + EPDM (10.0 wt%) PPBC + EPDM (15.0 wt%)

A/m (arbitrary units) 232 216 208 206 197

As observed under optical microscope with crossed polarizers, incorporation of up to 10 wt% EPDM into PPBC does not affect the nucleation density or ultimate size of PPBC spherulites (Figure 12.19). The EPDM particles act as inert inclusions, constituting geometrical obstacles to the PPBC spherulites’ growth, thus changing their morphology. Nevertheless, some interfacial inter-

Properties and Performance of Polymer Blends

Figure 12.18. DSC thermograms recorded during cooling cycle for PPBC and its blends with EPDM (10 wt%).

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actions are observed in the case of a blend with 10 wt% EPDM. The spherulites of PPBC are found to nucleate from the interface with EPDM (Figure 12.19c). This resembles the transcrystalline structure observed in several glass or carbon fibers reinforced, semicrystalline polymers, such as: PP, PE, PA-6, etc. [Xavier, 1991]. Such a structure was considered an indication of good interfacial interaction between the two constituents. The EPDM particles were found to either initiate crazes or to terminate them, depending on the interfacial bonding, the particle size, concentration and the interparticle distances. The variation of notched Izod impact strength of PPBC blends with different EPDM concentrations is shown in Figure 12.16. The fracture surface of a blend with 10 wt% EPDM, as examined under SEM, is shown in Figure 12.20. The hemispherical embeddings and hollows (representing the removed

Figure 12.19. Optical micrographs of (a) PPBS spherulites, (b) PPBC spherulites in the presence of EPDM and (c) PPBC spherulites nucleating from interface with EPDM.

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Figure 12.20. SEM micrograph of PPBC+EPDM (10 wt%) blend fracture surface.

EPDM particles) are clearly visible on the fracture surface. The ribbon-like structures visible on the fractured surfaces are probably the micro-shear bands in the blends.

Approximate ranges of the experimental techniques to study different blend morphologies are summarized in Table 12.15. See also Chapter 8. Morphology of Polymer Blends in this Handbook.

Table 12.15. Approximate ranges of experimental techniques to study blend morphology of: (1) Inter-atomic; (2) Molecular, spherulites; (3) Filler aggregates, compatibilized blends; (4) Reinforcements, immiscible blends; (5) Voids [Utracki, 1989] Domain Size Scale (µm)

l

2

10-4

10-3

3 10-2

10-1

100

Microscopy

4

5

101

102

Optical SEM TEM

Spectroscopy Thermal Mechanical Dielectric Diffraction

IR DSC TMA DS WAXS SAXS SANS Light

Properties and Performance of Polymer Blends

12.4.5.3

Particle Size

Elastomeric particle size plays a prominent role in controlling the toughening mechanisms of a polymer. It has been shown that particle size of an elastomer significantly influences the deformation and failure processes: small particles favor shear yielding, while coarser dispersions promote crazing [Jang et al., 1985]. There is an optimal particle size resulting in maximum impact resistance [Speri and Patrick, 1975; Stehling et al., 1981; Karger-Kocsis et al., 1981]. The tensile stress-strain curves of PPBC and its blends with EPDM of different particle sizes (for concentration equal 10 wt%), are shown in Figure 12.21 [Xavier et al., 1994]. The particle size of EPDM has considerably affected the postyield behavior of the blends. Although the yield stress initially increases with reduction of EPDM particle size (from 4.35 to 3.45 µm), with further reduction of size ( to 0.60 µm) it decreases. Earlier studies of the tensile properties of heterogeneous polymer blends have shown that Young’s modulus poorly reflects morphological changes [Pukanszky et al., 1989; Pukanszky and Tudos, 1990]. Accordingly, the particle size dependence of Young’s modulus is weak.

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By contrast (see Figure 12.22), in the case of PPBC blend with EPDM, the flexural modulus was found to significantly decrease when large EPDM particles (d • 4 mm) were used. The notched Izod impact strength was strongly affected by the particle size. As usually is the case, reducing it caused the impact strength to increase. With the reduction in EPDM particle size, the number of particles has increased and the interparticle distance was reduced. Thus, multiple crazing in the blend was enhanced further with reduction of the particle size.

Figure 12.22. Influence of EPDM particle size on flexural modulus and Izod impact strenght of PPBC/EPDM (10 wt%).

Figure 12.21. Influence of EPDM particle size on tensile stress-strain curves of PPBC blends. EPDM particle sizes (1) 0.60 mm, (2) 3.45 mm, (3) 4.35 mm, (4) 5.50 mm.

The shear yielding also is influenced with change in particle size resulting in increased impact strength. The notched Izod impact fracture surfaces examined under SEM are shown in Figure 12.23. Change in EPDM particle size had significantly changed the fracture morphology. The fibrous sheet-like structures are probably the microshear bands in the blends. The number of these bands increases with reduction of the rubber particle size as the smallest particle (at constant loading) corresponds to their maximum number. The formation and break down of the microshear bands absorb enormous energy and hence it increases the Izod impact strength.

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Figure 12.23. SEM micrographs of the impact fracture surfaces of (a) PPBC and its blends with EPDM (10 wt%) with different (b) 5.50 mm, (c) 4.35mm, (d) 3.45 mm and (e) 0.60 mm particle sizes.

Properties and Performance of Polymer Blends

Both principal fracture mechanisms, shear yielding and crazing, are influenced by the particle size. In PPBC matrix, where spherical elastomeric particles are chemically bonded, the energy absorption takes place mainly by deformation of the matrix. In such systems, a large amount of shear yielding is to be expected. The shear yielding becomes more prominent upon increasing the concentration of EPDM as well as reduction of their particle size. The microshear bands in the fracture surface (Figure 12.23e) clearly support these expectations. 12.4.5.4

Miscibility

The notion of polymer miscibility implies intimate mixing on the segmental level. A miscible blend is expected to have a homogeneous composition throughout. The understanding of chemical principles in polymer miscibility is getting refined as a result of the appearance of several reviews and books on the topic [Krause, 1972; Olabisi et al., 1979; Paul and Barlow, 1980; Paul, 1982; Ottenbrite et al., 1987; Utracki, 1989]. The level of molecular mixing existing in polymer blends that exhibit macroscopic properties indicative of single-phase behavior is commanding considerable attention. More detailed information on this topic can be found in Chapter 2 Thermodynamics of Polymer Blends in this Handbook. The best commercial advantages of a polymer blend can be summarized by the word ‘versatility’ [Olabisi et al., 1979]. Unfortunately, miscible polymer-polymer blends usually show additivity of the component polymers properties, thus their versatility is limited. Furthermore, like any other single-phase resin, for most applications miscible blends need to be toughened and/or reinforced. Thus, with the exception of PMMA/PVDF blends (primarily used for coatings) there are no miscible blends on the market. The interest in miscible polymer blends is for the purpose of compatibilization and judicious selection of the processing conditions that may lead to the spinodal decomposition-type morphology (see Chapter 8 Morphology in this Handbook). Immiscibility dominates polymer blends. It reveals itself as opacity, delamination, double

893

glass transition, or combination of these properties. Most immiscible polymer blends require compatibilization and toughening. Owing to low values of the combinatorial entropy mixing, miscibility in polymer-polymer systems requires the existence of strong specific interactions between the components, such as hydrogen bonding [Olabisi et al., 1979; Solc, 1982; Walsh and Rostami, 1985; Utracki, 1989]. The thermodynamic characterization of the interactions in miscible polymer blends has been the subject of extensive studies [Deshpande et al., 1974; Olabisi, 1975; Mandal et al., 1989; Lezcano et al., 1992, 1995, 1996; Farooque and Deshpande, 1992; Juana et al., 1994]. Based on the Huggins-Flory theory, the polymer-polymer interaction parameter, χ23, has been used to describe interactions between the two components. As a consequence, this “parameter” takes into account the enthalpic and noncombinatorial entropy of mixing contributions. Calorimetry, differential scanning calorimetry (DSC), Fourier transform infra-red spectroscopy (FTIR), inverse gas chromatography, microscopy, etc., are used to investigate the miscibility and morphology of the blends [Zhong and Guo, 1998; Lezcano et al., 1998]. Comprehensive surveys of miscible polymer systems along with various methods of miscibility determination have been published [Olabisi et al., 1979; Utracki, 1989; Coleman et al., 1991]. 12.4.5.5

Other Factors

Temperature strongly influences the impact behavior of toughened plastics. Charpy impact energy measurements at different temperatures in the case of HIPS containing various concentrations of PBD showed two transitions, at 233 and 273 K [Bucknall, 1988]. At these temperatures, the material exhibited transitions from brittle to semi-ductile and then to ductile. Newman and Williams [1978] carried out sharp-notch Charpy tests for ABS at: 193 ” T (K) ” 333 and showed that linear elastic fracture mechanics was applicable only up to 233 K. Above 273 K, the energy absorbed in impact

894

S. F. Xavier

was proportional to the fracture area, and correlated well with the volume of the whitened zone. Mixed behavior occurred at the intermediate temperatures. More detailed study of the notched Izod impact behavior of ABS was carried out using instrumented tests [Rink et al., 1978]. The authors found that the force at peak load, Fm, decreased slowly with increasing temperature from 133 to 353 K, increased by a factor of two between 193 and 273 K, and then decreased again. Because of the oxidative degradation of the main-chain double bonds, the plastics toughened by diene-type elastomers (e.g., PBD or other rubbers) are susceptible to aging. The UV radiation breaks the chemical bonds, initiating chain reactions in which the polymer is attacked by the atmospheric oxygen, becoming crosslinked or chemically degraded. Embrittlement of the surface has a similar effect to the introduction of a sharp crack. The effects are clearly seen in Charpy and Izod impact tests. Although the sub-surface polymer is unaffected by the aging, a crack initiated at the surface can accelerate throughout the degraded layer and cause lowenergy fracture of the specimen (see Chapter 14. Degradation and Aging). Geometry of a toughened plastic specimen also influences the impact strength and its mode of failure. The specimen’s length, width and thickness may affect the fracture behavior. Whether the specimen is notched or not, as well as the dimensions of the notch may also influence the impact behavior [Kinloch and Young, 1983]. As discussed in Part 12.4.3 Fracture Mechanics Testing, it is important to determine the material parameters (the initiation and the propagation energies) using the specimen geometry that reduces the effects of geometry to an acceptable level.

12.5

Miscibility and Solubility (by V. N. S. Pendyala)

The influence of chemicals and solvents on a polymer blend depends on the nature of solvent and the blend components, as well as on morphol-

ogy of the blends. The chemical/solvent resistance of an amorphous polymer is improved by the presence of semicrystalline polymer(s). For the best solvent resistance, the latter polymer should be the matrices. The semi-crystalline polymers such as PA, POM, PBT or PET contribute the solvent and chemical resistance, high processability, and rigidity, while amorphous polymers (ABS, PC and PSF) provide impact strength and elongation, and often cost reduction. Amorphous polymers are typically used to impact-modify crystalline polymers [Toensmeier, 1994]. For example, PEEK is a semicrystalline polymer with excellent chemical resistance, while PEI is an amorphous polymer and soluble in many solvents. The chemical resistance of the miscible PEEK/PEI blend is improved. It is observed that the solvent uptake of PEEK/PEI depends on such factors as: the solubility parameter, hydrogen bonding and structural factors of the solvent. PEEK will uptake solvent with a solubility parameter within a small range centered about 19.38 (J/mL)1/2 and a hydrogen bonding index of about 3.1 [Stober and Seferis, 1988]. Liquids with solubility parameters of about 20.40 (J/mL)1/2 were found to produce the maximum amount of load drop and cause crazing. The mechanical and other properties of PEEK/PEI (50/50 wt%) blend, after immersing in solvent, are given in Table 12.16 [Browne et al., 1997]. The crystallization in PEO/PMMA blends is influenced by the nature of the casting solvent. Phase segregation in the blend and PEO crystal growth inhibition by PMMA in certain directions have been observed in these miscible blends [Radhakrishnan and Venkatachalapati, 1996]. In PET/PC blend (Makroblend™ 4T 1018), the chemical and solvent influences are improved, compared to constituent polymers. The chemical resistance of this blend is shown in Table 12.17 [Mark et al., 1988]. Miscibility plays an important role in influencing chemical and solvent effects on a polymer blend. In a blend, where strong intermolecular interactions exist, the free energy of mixing, ∆Gm, is given by the relation [Painter et al., 1988, 1989]:

895

Properties and Performance of Polymer Blends

Table 12.16. Effect of immersion in acetone on the mechanical performance of amorphous PEEK, PEI and their 50/50 blend [Browne et al., 1997] Conditions

Specimen Tensile strength (MPa)

Yield stress (MPa)

Modulus (xl00 MPa)

Untreated

PEEK PEI Blend PEEK PEI Blend PEEK PEI Blend

41 50 45 27 0.6 15 32 5.0 24

13.4 ± 14.0 ± 13.4 ± 9.6 ± 0.83± 5.0 ± 1.0 ± 0.3 ± 1.0 ±

Immersed at 25°C to equilibrium Immersed at 45°C to equilibrium

56 91 70 33 1.5 16 38 5 28

± ± ± ± ± ± ± ± ±

2 4 4 2 0.9 0.9 4 3 3

± ± ± ± ± ± ± ± ±

1 2 2 2 0.6 0.9 2 3.0 5

0.5 0.9 0.07 0.5 0.9 0.3 0.2 0.2 0.2

Strain at break (%) 245 ± 15 14 ± l 50 ± 50 286 ± 35 4 ± 2 17 ± 6 6.8 ± 0.5 304 ± 12 5 ± 1

Weight increase (%)

13.1 20.2 17.6 12.1 16.9 19.4

± ± ± ± ± ±

0.2 0.2 0.4 0.2 0.3 0.4

Table 12.17. Chemical resistance of an impact-modified PET/PC Makroblend® 4T 1018, after one day1 [Mark et al., 1988] No. Medium

1. 2. 3. 4. 5. 6. 7. 8. 9. 10. 11. 12. 13. 14. 15. 16. 17. 18. 19. 20. 21. 1

Isopar G cable oil Isopar G cable oil Norpar 12 cable oil AW 32 cutting oil AW 32 cutting oil Harmony 53 hydraulic fluid Harmony 53 hydraulic fluid Unleaded gasoline, premium Optikleen windshield-washer fluid 10 % Sulfuric acid Power-steering fluid Automatic transmission fluid Motor oil Zerex antifreeze Tri-chloro ethylene Xylene Toluene 1 Methyl ethyl ketone Acetone Plastic adhesive3 Freon TF solvent

except otherwise stated; Chemical Co.

2

Temperature (°C)

Rating2

Strain limit (%)

23 49 23 23 82 23 82 23 23 23 23 23 23 23 23 23 23 23 23 23 23

R R R R R R R R R LR R R R R R R R§ R§ R§ R R

1.6 1.2 1.6 1.6 1.6 1.6 1.6 1.6 1.6 1.0 1.2 1.2 1.4 1.6 1.6 1.6 1.6 1.4 1.4 1.4 1.4

R = resistant; LR = limited resistance;

3

Manufacturer of medium Exxon Exxon Exxon Arco Arco Gulf Gulf Amoco GM Wynn’s Wynn’s Valvoline Conoco

3M M-S Chem. §§

after 6 days; §: stress whitened; §§: Miller-Stephenson

896

S. F. Xavier

∆Gm / RT = ∆GH / RT +

(V / Vr ){(φ1 / N1 )ln φ1 + (φ2 / N2 )ln φ2 + χ12′ φ1φ2 }

(12.19) where R is the gas constant; T is temperature (Kelvin), φi (where i = 1 or 2) is the volume fraction of polymers 1 or 2, Ni is the degree of polymerization (i = 1, 2), V is the total molar volume and Vr = 100 (mL/mol) is the reference volume (see Eq 2.39 in Chapter 2 and the following discussion). Excepting the new term, ∆GH (introduced to express the effects of specific interactions) the relation is the same as published independently in 1941 by Huggins and Flory. The new term has been used to describe all types of specific interactions (hydrogen bonding, ion-ion, ion-dipole, charge transfer, π−π-electron interactions, etc.) that provide negative contribution to the free energy of mixing. The interactions of the van der Waals type are accounted for by the χ′12φ1φ2 • 0 term with χ′12 given by Hildebrand’s solubility parameter. For more details see Chapter 2 Thermodynamics of Polymer Blends. 12.5.1

Solubility Parameters

In systems where the interactions are dominated by the van der Waals forces, solubility is favored by chemical similarity of solvent and polymer. Molecular weight, chain branching and crosslinking of individual polymers slightly influence the solubility parameter, δ [Hildebrand, 1916]. In a strict sense, the molecular interactions should be non-specific, without forming associations or orientation, hence not of the hydrogen or polar type. The intermolecular interactions 1-2, are assumed to follow the geometric mean of the intramolecular interactions, 1-1 and 2-2:  ∆E v ∆E2v  2 ∆E v ∆Hm / Vφ1φ2 =  1 − 1/ 2 121/ 2 + = V V V V2  1 2  1 2

 ∆E v  1/ 2  ∆E v  1/ 2   1  −  2   = (δ1 − δ 2 )2 ≥ 0  V2    V1 

(12.20)

where ∆Hm is the enthalpy of mixing, ∆Εiv is the molar energy of vaporization of substance i, and δi is the solubility parameter. It is important to note that Eq. 12.20 is limited to the enthalpic part of binary interaction parameter, i.e.: χ12′ = χ S + χ H = χ S + (Vr / RT )(δ1 − δ 2 )2

(12.21) where χS, originates from local configurational as well as non-configurational entropic effects. The biggest drawback of the solubility parameter approach has been omission of the entropic and specific interactions’ effects. Furthermore, the fundamental dependencies do not take into account either the structural (isomeric), orientation, or the neighboring group effects. However, since the contributions that are included in the solubility parameters are indeed detrimental to miscibility, minimizing their value must but help the miscibility [Van Krevelen and Hoftyzer, 1976]. To obtain a reasonable prediction of miscibility, the values of the solubility parameters are to be measured to an accuracy better than δ1 − δ2 < 0.2 (J/mL)1/2. The measured values of δi (where i = 1 or 2) far exceeds the magnitude of the critical difference of these parameters, δ1 − δ2 < 2 (J/mL)1/2. The calculated values of δi are claimed to be precise within ± 0.8 (J/mL)1/2 [Coleman et al., 1990]. The solubility parameter approach is applicable to amorphous polymer systems. In order to adopt to highly crystalline polymers, the heat and entropy of fusion (∆hmf and ∆smf, respectively) is to be dealt with in free energy of mixing equation [Van Krevelen and Hoftyzer, 1976]: ∆Gm = {∆Hm + ∆hmf } - T{ ∆Sm + ∆smf} (12.22) Highly crystalline polymers, viz. PE or PTFE are insoluble at room temperature, but they obey the solubility principles at T • 0.9 Tm, i.e., at temperatures not more than 10% (in Kelvin) above their melting temperature [Leung et al., 1987].

Properties and Performance of Polymer Blends

12.5.2

Binary Interaction Parameters

The nature of polymer-solvent interactions plays an important role in deciding the influence of chemical and solvent effects on blends. The Hildebrand’s solubility parameter, has been extended to systems that have dispersive (subscript d), polar (subscript p) and hydrogen bonding (subscript h) interactions, i.e., δd, δp, and δh, respectively [Hildebrand and Scott, 1949; Burrel and Hansen, 1975]. The miscibility behavior of blends is affected by the presence of solvent, due to difference in the interaction parameters. For, e.g., PS/poly(α-methyl styrene) blends have shown immiscibility in tetrahydrofuran and miscibility in cyclohexane for blends comprising either 0-25 wt% or 73-100 wt% of poly(α-methyl styrene). Similarly PS/PMMA blends are immiscible in chloroform and miscible in tetrahydrofuran at all compositions, excepting 50/50 wt% blends. It was observed that the degree of coil expansion, that depends on the solventpolymer interactions, is responsible for these effects [Van et al., 1987]. The phase separation of two polymers in organic solvent is due to repulsive interactions between the chain segments giving an endothermic heat of mixing. However, the mixing of polymers in an aqueous solution can lead to phase separation (immiscibility) or formation of insoluble complexes of homogeneous solution (miscibility) [Hefford, 1984]. As discussed in Chapter 2 of this Handbook, the osmotic pressure can provide one of the most direct methods of determination of interaction parameter. The estimation of these interaction parameters is applied to PS/PVME blends in toluene and ethyl benzene respectively [Shiomi et al., 1985]. It is observed that in toluene χ12 varied with composition from -0.044 to +0.0093, while in ethyl benzene it increased with PVME content from -0.06 to 0.027. Various radiation scattering techniques have been employed in order to estimate values of the interaction parameters, χ12. The techniques used, are small angle neutron scattering (SANS) and small angle X-ray scattering (SAXS) [Wendorff, 1980, 1982]. Coleman et al., extensively discussed the application of solubility parameters

897

to calculate interaction parameters of blends based on PVC and other polymers; the limitations of solubility parameters evaluation etc.[Coleman et al., 1990]. 12.5.3

Effect of Crystallinity

Polymer crystallization from the melt takes place when system is cooled below the equilibrium melting point, Tm,e, i.e., to the crystallization temperature, Tc < Tm,e. The difference: Dc = Tm,e Tc, depends on the cooling rate and the nucleation. There are three mechanisms of the crystallization nucleation [Utracki, 1989]: 1. Spontaneous, homogeneous nucleation — it rarely occurs in the supercooled homogeneous melt. 2. Orientation-induced nucleation, caused by alignment of macromolecules, e.g. in extensional flow field. 3. Heterogeneous nucleation on the surface of a foreign phase. In thermoplastic blends (2) and (3) are most important mechanisms. It has been known for more than a century that impurities reduce the melting point. This observation has been used to determine the molecular weight of the contaminant by Raoult in 1882-5. Nearly a hundred years later, this concept was used to calculate the thermodynamic binary interaction parameter, χ12, from the melting point depression of a crystalline polymer in miscible blend with low concentration of another polymer. The relation is usually used in the simplified form for very high molecular weight components [Nishi and Wang, 1975, 1977]: Tm/Tm,n = 1 + (χ12 RT/V1) (V2/ ∆Ηm) φ2

(12.23)

where Tm is the melting point of crystalline polymer in miscible blend, Tm,n is melting point of crystals of neat polymer, Vi is molar volume (i = 1, 2), ∆Ηm is the heat of fusion [Utracki, 1989]. More detailed description of problems associated with nucleation, crystallization, and the use of the melting point depression for the

898

S. F. Xavier

characterization of χ12 ” 0 in polymer blends can be found in Chapter 3 Crystallization, Morphological Structure and Melting of Polymer Blends in this Handbook. 12.5.4

Effect of Molecular Weight

As discussed in Chapter 2 of this Handbook, by definition, the miscible polymer blend is characterized by negative value of the free energy of mixing, ∆Gm § ∆Hm ” 0, and positive value of its second derivative ∂2∆Gm/∂φ2 > 0. Substituting into these relations the expression for ∆Gm makes it possible to define the critical conditions for miscibility. The simplest expression for ∆Gm (for binary systems that contain ingredient i = 1 or 2) is that derived by Huggins and by Flory (H-F): ∆Gm / RTV = (φ1 / V1 )ln φ1 + (φ2 / V2 )ln φ2 + ( χ12 / V1 )φ1φ2 ∆Gm = RTV [(φ1 / V1 )ln φ1 + (φ2 / V2 )ln φ2 ] + Bφ1φ2

χ12 / V1 ≡ χ12′ ;

B ≡ χ12 RT (V / V1 )

(12.24)

conditions at constant temperature and pressure for systems with species of different molecular weights. Three cases must be distinguished: 1. Solvent solutions: V1 = V2 or N1 = N2 ≅ 1, are miscible if χ12 < χ12,cr = 2 2. Polymer solutions: V2/V1 → ∞ or N1 ≅ 1 and N2 → ∞, are miscible if χ12 < χ12,cr = 1/2 3. Polymer blends: V2 and V1 → ∞ or N1 ≅ N2 → ∞, are miscible if χ12 < χ12,cr = 0 More details of these aspects are presented elsewhere [Olabisi et al., 1979; Paul, 1985; Utracki, 1989] as well as in Chapter 2 of this Handbook. It may be interesting to note that assuming equal degrees of polymerization, Ni, of both blend components, Eq. 12.25 yield a simple relationship between the binary interaction parameter and the molecular weight (as expressed by Ni): BcrN/RT = 2. Thus, within the framework of the Huggins-Flory theory, system will be miscible when BcrN/RT < 2, and immiscible when BcrN/RT > 2. This is schematically shown in Figure 12.24.

where φi is the volume fraction and Vi is the molar volume of the specimen “i” (where i = 1 or 2). The first two logarithmic terms give the combinatorial entropy of mixing, while the third term the enthalpy. For polymer blends Vi is large, thus the combinatorial entropy is vanishingly small, hence miscibility depends primarily on the value of. Substituting H-F relation into the isobaric and isothermal thermodynamic conditions for phase separation: (∂2∆Gm/∂φ22)P.T = (∂3∆Gm/∂φ23)P.T = 0, and treating the socalled binary interaction parameter, χ12, as a constant, the critical conditions for the phase separation have been expressed as: 2

2

χ12′ ,cr ≡ χ12,cr / V1 = (V1−1/ 2 + V2−1/ 2 ) / 2 ≅ ( N1−1/ 2 + N2−1/ 2 ) / 2 2

Bcr = ( RTV / 2)(V1−1/ 2 + V2−1/ 2 ) ≅ ( RT / 2)( N1−1/ 2 + N2−1/ 2 )

2

(12.25) where Ni is degree of polymerization of polymer i = 1 or 2. This dependence provides the miscibility

Figure 12.24. Free energy of mixing for polymers 1 and 2 with the same molecular weight M calculated from from Eq 12.25 [Paul, 1985].

Properties and Performance of Polymer Blends

12.5.5

Miscibility of Copolymers

In blends of random copolymers, or in blends of a polymer with random copolymer, the presence of repulsive forces among segments (other than specific interactions discussed before) may lead to miscibility [Hammer, 1971; Kambour et al., 1983; ten Brinke et al., 1983; Ellis, 1990; Zhu et al., 1990]. The excimer fluorescence has been used to characterize the miscibility of the guest and host polymers [Xie et al., 1993]. Since the excimer forming site concentration depends on the extent of the guest polymer aggregation, the ratio of excimer to monomer fluorescence intensity, IE/IM, is related to blend miscibility. If the blend is miscible, the guest polymer dilutes the matrix of the host polymer. This lowers the excimer-forming site concentration, thereby leading to a decreased intensity of excimer emission. It has been found in the study of PVME and SBS triblock copolymer that solubility of PVME in PS block copolymer domains is larger than in PS homopolymer. This may indicate that the mixing enthalpy has an effect on the blend miscibility [Xie et al., 1993]. The behavior has been attributed to the effect of PB segments in SBS. The phase equilibria and miscibility in polymer blends containing random or block copolymer was reviewed [Roe and Rigby, 1987]. More recent data are presented in Chapter 4 Interphase and Compatibilization by Addition of a Compatibilizer in this Handbook. 12.5.6

Standard Methods of Evaluation

Table 12.18 provides a list of standard test methods for the evaluation of solvent and chemical susceptibilities of polymers and their blends. These tests are to help in: (1) designing of fabricated parts, (2) comparison of different materials and (3) characterization of materials for long term performance, etc. The materials performance, summarized after the immersion in reagents for 7 days at elevated temperatures according to ASTM D 543, based on changes in weight, dimensions, appearance and mechanical properties, is shown in Table 12.19.

12.5.7

899

Factors Influencing Chemical and Solvent Effects

The influencing factors for these effects, viz., structure, composition, nature of copolymer in blends are presented in this section. 12.5.7.1

Structure

Chemical structure of polymers constituting the blend determines its solvent resistance. The polymers, having the backbone linkages involving oxygen, sulfur, and silicone exhibited enhanced chemical and solvent resistance. Thus, enhanced chemical and solvent resistance has been reported for polymer blends that comprise polymers with: • ether, thioether, and oxymethylene linkages [Mizutani et al., 1991; Hergenrother and Havens, 1990; Nakajima and Misoo, 1991; Brown et al., 1992], • siloxane and/or imide groups [Kinoshita and Ishikawa, 1992; Rock et al., 1991], • carbonate, with other polymers, viz. semicrystalline polyesters [Kosmo and Dawn, 1990; Skochdopole and Wright, 1990; Boutni, 1991; Tera et al., 1992], • fluorine [Siol et al., 1990; Abe, 1991; Endo et al., 1990], • certain block polymers [Mullins et al., 1990; Ota, 1990; Rein and Heinz, 1990], • PPE/PS blends [Okabe et al., 1990; Hirose, 1990], • blend of polybenzimidazole and PI [Leung et al., 1987], • polymer blends of hydroxyl terminated polyester and PPE-copolymer [Falk and Khait, 1990], • ternary blends of maleic anhydride modified PPE, melamine-formaldehyde-methanol copolymer and a POE [Maruyama and Mizumo, 1990], • blends of PPS having -SH end-groups and PPE with unsaturation and/or carboxy/anhydride groups [Kawaki et al., 1991], • 50:50 blends of PPE grafted with glycidyl methacrylate with PP-MA (MA = 0.6%) [Arashiro et al., 1992],

ASTM D1239-98

ASTM D1712-89

ASTM D2990-95

ASTM D2292-82

DIN 53476

BS 4618-Sec. 4.1 ASTM D2299-92

2.

3.

4.

5.

6.

7. 8

Note:

Test specimens should be immersed for 7 days in several selected solvents.

(1/8 in); molded slabs diameter: 50.8 mm (2 in) (surface area 45.6 cm2). 2. For mechanical properties: standard tensile bars of D 638, type I or II. 3. For molded & extruded specimens (as described in 1). Cut specimens’ edges shall be made smooth by fine sand paper or emery cloth. In the form of squares with 50 ± 0.25 mm (2 in) on each side.

*ASTM standard test methods are available on web: http://enterprise.astm.org/

application of the material. Applicable to specimens having thickness > 0.1 mm.

Resistance determined in ca. 50 liquids. Test conditions depend on desired

For instrumental evaluation, as per required size for that instrument. Follows ISO 175-81.

Exposure period & solvents as suggested by ISO R175. Determination of stain resistance of plastics.

Susceptibility of stain and the extent of staining are determined.

sections are 12.7x 12.7 mm (0.50 x 0.50 in) or 12.7 mm diameter; plane surface. Flexural creep: 63.5xl2.7x3.18 (2.5x0.5x0.l25 in) or 127xl2.7x6.4 mm (5.0x0.5x0.25 in) with close tolerances. Flat, smooth surface having 25 mm2 for visual examination.

and after exposure to chemical environment.

Tensile, creep rupture, flexural & compressive creep tests are conducted before

Unconfined compressive creep: right cylinder or prism cross

Determination of sulfide staining.

Length 100 ± 25 mm (4 ± l in) by width 13 ± 6 mm (0.5 ± 0.25 in) by the thickness of composition being tested. Tensile creep: Type II, or I as specified in D638-99.

50% ethanol, etc., for plastic films by weight loss before & after immersion.

For disk type: 50 ± 0.25 mm (2 in) diameter (total area 41.5 cm2) or a tension specimen 100 by 25 mm (4 by 1 in). Size and shape of test specimens are relatively unimportant.

Chemical resistance using soap solution, cottonseed oil, mineral oil, kerosene,

Similarity between testing & end use conditions should be maintained.

1. For weight & dimension changes: thickness: 3.175 mm

ASTM D543-95

1.

Comment

Specimen

No. Test Method

Table 12.18. Standard test methods* for chemical/solvent effects

900 S. F. Xavier

901

Properties and Performance of Polymer Blends

Table 12.19. Classification of test results of materials’ chemical and solvent effects [Agranoff, 1983] Category

Significant Significant, but not conclusive Not significant

Change in %

Change in appearance

Weight

Dimension

Tensile and flexural strengths

> 1.0

> 0.5

> 20

0.5 - 1.0 < 0.5

0.2 - 0.5 < 0.2

10 - 20 < 10

• PEI/PC blends comprising amine terminated siloxane [Rock et al., 1991]. The structure of repeat units of individual polymers constituting a blend, as well as the nature of interactions between polymers in a blend are the factors that influence solubility characteristics of a blend. Thus, solubility is affected by: • cross linking [Yoshimura and Fujimoyo, 1968; Zapp, 1973], • hydrogen bonding [Cangelosi and Shaw, 1983; Djadoun et al., 1977; Pearce et al., 1984; Isasi et al., 1993; Lee et al., 1988], • formation of donor-acceptor complexes [Pugh and Percec, 1986; Rodriguez-Parada and Percec, 1986], • dipole-dipole interactions [Prud’homme, 1982; Woo et al., 1983], • ion-dipole interactions [Hara and Eisenberg, 1984], • ion-ion interactions [Domard and Rinaudo, 1980; Natansohn et al., 1988], • acid-base interactions [Zhou and Eisenberg, 1983; Weiss et al., 1990], • segmental interactions as observed in SAN/ NBR or PS/PVME blends [Cowie et al., 1992, 1993]. The methods for estimation of solubility parameter, δi, of copolymers can be extended to polymer blends [Pendyala, 1988; Ho et al., 1991]. As it was mentioned before, three contributions to the solubility parameter, δi, have been

Softened, crazed, warped, distorted Optically discolored Slightly discolored

identified: dispersive δd, polar δp and hydrogen bonding δh [Hansen, 1967]. The values of δi for PC/ABS blends, shown in Table 12.20, have been calculated using Hansen’s approach [Pendyala and Xavier, 1992]. These values are found to be comparable with those obtained by other techniques (see Table 12.21) [Small, 1953; Van Krevelen and Hoftyzer, 1967; Hoy, 1970]. The three values of δi are important when considering the solvent susceptibility of a particular polymer blend. 12.5.7.2

Composition

The chemical and solvent effects are also related to the blend composition. Three different grades of ABS (differing in the A, B and S content) were blended in the full composition range with two grades of PC (PC1 and PC2 having different molecular weights). Their solvent susceptibility is illustrated in Tables 12.22 and 12.23. It is shown in Table 12.22, that chloroform dissolved all these blends within the entire composition range. The PC/ABS = 90/10 blends have shown different behavior for different grades of ABS. Thus, solubility of PC-1/ABS-3 in tetrahydrofuran, cyclohexanone, 1,4-dioxane was found to be different. This has been attributed to the difference in the degrees of interaction between the two polymers. The basic difference in the weight fractions of A, B and S in the ABS resin was the influencing factor. However, with the increase in ABS content in the blends, the blends showed

902

S. F. Xavier

Table 12.20. Partial (δd, δp, δh) and total (δ) solubility parameters and radius of interaction sphere (CR) of polymers/blends [Pendyala and Xavier, 1992] Blend PC1 and PC2 ABS3 PCl/ABS3-1 PCI/ABS3-3 PCl/ABS3-5

δd

δp

δh

δ

CR

18.3 17.6 18.6 18.2 18.7

7.0 7.7 4.1 7.8 9.1

5.7 4.4 5.3 9.3 8.8

20.4 19.7 19.7 21.8 22.5

5.2 4.6 2.3 6.0 7.2

Where PCl/ABS3-1 = 90/10; PCl/ABS3-3 = 50/50; PCl/ABS3-5 = 30/70 All values of the solubility parameters are in (MJ/m3)1/2

Table 12.21. Solubility parameters of polymers and their blends obtained using different methods [Pendyala and Xavier, 1992] Polymer PC1 and PC2 ABS3 PCl/ABS3 and PC2/ABS3

δ (Small)

δ (Van Krevelen)

δ (Hoy)

20.1 24.7 21.2

21.1 26.7 21.6

21.6 25.0 21.6

δ (Hansen) 20.4 19.7 19.3-22.5 19.7-20.4

All the values of solubility parameters are in (MJ/m3)1/2

increased and similar affinity to all used solvents (see Table 12.23). Reduction of molecular weight of PC also enhanced the solvent susceptibility of these blends. 12.5.8

Review of the Domain

The solvent effects on polymer blends can be estimated using the solubility parameter, δ. The concept was originally used to characterize the strength of interactions in simple liquids, but later on it was extended to polymer/solvent as well as to polymer/polymer systems. For example, in the coating and rubber industries the solubility parameter concept has been used respectively, to select solvents or to study swelling of cured rubber by solvents. Bohn has extended the solubility parameter approach to polymer-solvent systems [Bohn, 1968]. Solubil-

ity parameter approach enables one to identify the potential soluble components, without experimental study. The approach was also useful in calculating the effects of pressure and temperature on free energy of mixing. A generalized approach is suggested to calculate δ from the internal pressure of a system [Olabisi and Simha, 1977]. Other methods of calculation include heat of solution [Guillet, 1973, Newman and Prausnitz, 1972], solution behavior [Manning and Rodriguez, 1973; Song and Brownawell, 1970], etc. Hildebrand applied the solubility parameter approach to mixing of two polymers, to estimate the enthalpy of mixing per unit volume [Hildebrand and Scott, 1964]. The concept was extended to incorporate the polar and hydrogen bonding interactions [Hansen, 1967, 1969; Nelson et al., 1970]. Thus, the three-dimensional solubility parameter was expressed as:

903

Properties and Performance of Polymer Blends

Table 12.22. Solubility characteristics of PC/ABS blends with different A, B, and S content of ABS [Pendyala and Xavier, 1992] Solvent No.

1.

2.

3.

4.

5.

6.

PC1/ABS Blend composition

ABS used

90/10

70/30

50/50

30/70

10/90

+ + + + + + + * + + -

+ + + + * + + + * + + + *

+ + + + * + + + + + + + + * -

+ + + + * + * + + + + + * * + * -

+ + + + + * + + + + + * + + * -

ABS1 ABS2 ABS3 ABS1 ABS2 ABS3 ABS1 ABS2 ABS3 ABS1 ABS2 ABS3 ABS1 ABS2 ABS3 ABS1 ABS2 ABS3

Where 1 = Chloroform; 2 = Tetrahydrofuran; 3 = Cyclohexanone ; 4 = 1,2-Dichlorobenzene; 5 = 1,4-Dioxane; 6 = Methylene chloride. + Indicates soluble, * partially soluble and - insoluble ABS1 = A(24-26%), B(20%), S(56-54%) ABS2 = A(24-26%), B(12-13%), S(62-64%) ABS3 = A(26%), B(24%), S(50%) PC1 = Viscosity av. mol wt. Mv = 33,243 [in CHCl3 at 25°C]

δ2 = δd2 + δp2 + δh2

(12.26)

For several polymer-polymer systems the interaction energies were estimated from structure of polymer repeat units using the solubility concept [Shaw, 1974]. The methods of calculation of δ for various polymers and blends was also published [Van Krevelen and Hoftyzer, 1976; Olabisi et al., 1979]. Using Hansen’s method the values of δi for polymers and polymer blends can be estimated similarly as these of pigments. The “spherical

volume of interactions” for those systems can be represented by the three-dimensional plots. The points in three-dimensional space are shown in three separate plots as projections in XY, YZ and XZ planes. The “sphere of interaction” is projected as circles. The center of the sphere corresponds to the three-dimensional solubility parameter of that polymer or blend. This method of estimation of δ was applied to polymethacrylonitrile (PMAN), polymethacrylic acid (PMA) and poly(methacrylonitrileco-methacrylic acid) (MAN/MAA copolymer)

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S. F. Xavier

Table 12.23. Solubility characteristics of PC/ABS blends different A, B, and S contents of ABS [Pendyala and Xavier, 1992] Solvent No.

1.

2.

3.

4.

5.

6.

PC2/ABS Blend composition

ABS used

90/10

70/30

50/50

30/70

10/90

+ + + + + + + +

+ + + * * + + * + + + * +

+ + + + + + + + + * + + + * * +

+ + + + + + * * + + + + + + + + +

+ + + + + + + * + + + + + * + + +

ABS1 ABS2 ABS3 ABS1 ABS2 ABS3 ABS1 ABS2 ABS3 ABS1 ABS2 ABS3 ABS1 ABS2 ABS3 ABS1 ABS2 ABS3

Where: 1 = Chloroform; 2 = Tetrahydrofuran; 3 = Cyclohexanone; 4 = 1,2-Dichlorobenzene; 5 = 1,4-Dioxane; 6 = Methylene chloride. + Indicates soluble, * partially soluble and - insoluble ABS1 = A(24-26%), B(20%), S(56-54%) ABS2 = A(24-26%), B(12-13%), S(62-64%) ABS3 = A(26%), B(24%), S(50%) PC2 = Viscosity av. molecular wt. Mv = 14,079 [in CHC13 at 25°C]

[Ho et al., 1991]. The estimated values are shown in Table 12.24, and the three-dimensional representation of the spherical volume of interaction is shown in Figure 12.25. These values were compared with others, calculated using group contribution method [Van Krevelen and Hoftyzer, 1976; Fedors, 1974; Hoy, 1970]. Ho et al. concluded that the calculated values of δ were in good agreement with the experimental results for PMAN, but appeared to be low for PMAA and MAN/MAA. The methods of predicting solubility parameters of the latter copolymer from known values of homopolymer constituents

were ‘justified’ by comparing with their experimental results. The three-dimensional solubility parameter, δ, and its individual contributions were applied to PC/ABS blends [Pendyala and Xavier, 1992]. The values are listed in Table 12.20, and the sphere of interaction represented in three different circles is shown in Figure 12.26. In this case, the δ values calculated using Hansen’s method were in agreement with those obtained from other techniques. Hansen’s method was found suitable to provide optimum values of AN-content in PC/ABS blends [Huang and Wang, 1989].

905

Properties and Performance of Polymer Blends

Table 12.24. Comparison of experimental and calculated solubility parameters of PMAN, PMAA and MAN/MAA copolymers [Ho et al., 1991] Method PMAN Experiment Group contribution

PMAA Experiment Group contribution

MAN/MAA copolymer Experiment Group contribution

δd

δp

δh

δ

18.0 --

15.9 --

8.0 --

--

--

--

25.3 25.7 a 26.1 b 21.2 c

17.3 --

12.4 --

15.9 --

-17.3

-6.3

-12.2

17.7

14.2 --

11.8 --

-16.8

-11.9

-9.1

26.7 25.5 a 21.2 b 24.9 c 18.6 d 25.9 25.6 a 23.7 b 23.0 c 21.6 d

Note: a Calculated by Fedors method; b Calculated by Van Krevelen method; c Calculated by Hoy method; d Calculated from solubility parameter group contribution method. The values of solubility parameter are in (MJ/m3)1/2

Numerous blends exhibiting good solvent and/ or chemical resistance (in addition to other properties) have been described, viz., based on polyolefins [Yamamoto, 1991; Aleckner et al., 1986; Sezaki et al., 1990; Takada et al., 1991], fluoropolymer-PA type [Kawashima and Yoshida, 1991], PPE [Brown et al., 1992; Arashiro et al., 1992; Nakajima and Misoo, 1991], PC [Tera et al., 1992; Endo et al., 1987, 1989], PEST [Ohmae et al., 1992]. 12.5.9

Data on Blends

Table 12.25 provides information on the publications pertinent to solvent and chemical resistance of polymer blends. The chemical and solvent susceptibilities of the commercial polymer blends (as reported by the manufacturers) are listed in Table 12.26.

12.6

Thermal Properties

1.6.1

Thermal Resistance (R)

For a flat slab, it is calculated as [ASTM C177]: R = A (T1-T2)/Q = 1/Γ = D/λ

(12.27)

where R = thermal resistance (Km2W-1), A = area measured on a selected isothermal surface (m2), T1 = temperature of warm surface of specimens (K), T2 = temperature of cold surface of specimens (K), Q = heat flow rate (W), Γ = thermal conductance (Wm-2K-1), D = thickness of specimen measured along a path normal to isothermal surfaces (m), λ = thermal conductivity, (W/mK). The reciprocal of thermal resistance is known as thermal conductance, Γ.

906

S. F. Xavier

Figure 12.25. The solubility regions of PMAN, PMAN and their copolymer in δd- δp, δd- δh and δp- δh planes respectively [Ho, et al., 1991].

1.6.2

Thermal Conductivity (λ)

Thermal conductivity of materials can only be defined for homogeneous materials, where the thickness is greater than that for which the apparent thermal resistivity of the material does not change by more than 2% with further increase in thickness. The thermal resistance must be sufficiently independent of the area of the specimen and for a flat slab specimen the thermal resistance must be proportional to the thickness. When all these conditions are met: λ = QD/A (T1-T2) = D/R

(12.28)

where the symbols are the same as in Eq. 12.27.

The reciprocal of the thermal conductivity is called thermal resistivity (r). The most common units for thermal conductivity are: cal/cm.s.°C and Btu.in/ft2.h.°F. The SI unit for conductivity is W/mK. Since a variety of units has been in practice for thermal properties, the conversion factors are given in Table 12.27. ASTM C177 and BS 874 recommend guarded plate method for materials of low conductivity. Two different types of guarded hot plate apparatus are described in ASTM C177. The low temperature guarded hot plate is the most suitable method for determining the thermal conductivity of polymeric solid materials including foams. It is generally used for measurements where the tem-

Properties and Performance of Polymer Blends

Figure 12.26. (90/10) blend.

907

Typical plots of three dimensional solubility parameters in δd- δp, δd- δh and δp- δh planes for PC/ABS

perature of the heating unit is not above 500 K. The second method is the high-temperature guarded hot plate which is ordinarily used for measurements where the heating unit temperature is greater than 550 K but less than 1350 K. The schematic arrangement of the apparatus is shown in Figure 12.27. Three possible configurations to restrict edge heat flux are illustrated in Figure 12.28. The apparatus consists of a heating unit, a cooling unit and an edge insulation. The heating and cooling units may be either round or square. The heating unit consists of a central metering section and a guard section. The metering section consists of a metering area heater and metering area surface

plates. The guard section consists of one or more guard heaters and the guard surface plates. The working surfaces of the heating unit and cooling plates should be smoothly finished to conform to a true plane. The heating unit has a separation (or gap) not greater than 4 mm between the surface plates of the metering area and the guard. Two specimens should be selected from each sample with their surfaces made plane. The temperature difference between the hot and cold surfaces of the specimens should be not less than 5 K [De Ponte and Di Filippo, 1974]. The central heat source and the guard should have independent power supplies. The cold sur-

908

S. F. Xavier

Table 12.25. Review of the domain Blend

Test

Result

Miscible and immiscible polymer pairs PC/PMMA

Phase separation behavior

Thought to be hydrogen bonding factor Phase separation, solvent induced crystallization in PC, modulated bi-phase structure, single Tg

Solvent casting in THF, cyclohexane, methylene dichloride

at 47°C PS, PMMA/ Solution casting — excimer Phase separation before poly(N-vinyl-2-pyrrolidone) fluorescence solidification PBT/polyarylate Blends’ crystallization behavior Solvent induced crystallization of in acetone over a range of PBT temperature (in the range ~30, 50-60 wt%) PS/PVME Light scattering in various solvents Miscibility concluded Styrene copolymer/ polymethacrylate PVC-poly(caprolactone)

PET/polyarylate

Cellulose and PAN or PA

Poly(2,6-dimethyl-1,4phenylene ether)/epoxy Polyester polyol blocked polyisocyanates Fluoro-olefin copolymer/ polyisocyanate PC/ABS

ABS/SAN PVC/Polyamide

Miscibility detected by the nonradiative energy transfer (NET) Interaction parameter by inverse gas chromatography at 120°C Dissolved in PhOH, 1,1,2,2-tetrachloroethane; effect of solvent composition on miscibility Phase separated morphology studies in solutions (DMAc) containing LiCl Solution mixing of two resins

Miscibility strongly depends on the casting solvent/procedure

Reference Hefford, 1984 Saldanha and Kyu, 1987

Semerak and Frank, 1987 Waywood and Durning, 1987 Ould-Kaddour and Strazielle, 1988 Chen and Morawetz, 1989

Vapors of poor solvents interacted; Hattam et al., 1989 better solvents penetrated them slowly Long term effect of non-equilibrium Kim et al., 1990 state of phase separation; difficult transesterification Solvent induced phase separation

Properties dominated by PPE. Solvent is converted to dispersed phase. Blend coating applied on steel Resistant to acids, methyl plate and baked at 220’C. ethyl ketone Fluoro polymer (F > 5%), solvents Good resistance to solvent, with > 30 % paraffins or staining and chemicals; naphthenes selected re-coatability Complete range of blends Solvent effects identified. δ of behavior in different solvents; blends correlated to mechanical behavior mechanical behavior 50/50 blend melt kneaded to plate Xylene/ethanol resistant (1.5-2.0)/1 blend with Superior chemical resistance functionalized ethylene compatibilizer

Roeder et al., 1993

Venderbosch et al., 1994

Maruyama and Herarudo, 1996 Hanamura et al., 1996

Pendyala and Xavier, 1992; 1997 Ito et al., 1997 Hofmann, 1997

909

Properties and Performance of Polymer Blends

Table 12.25. continued Blend

Test

Result

PC/PET 75/25 and 60/40 blends with core-shell copolymer PMMA/ poly(Na methacrylate) SMA/epoxidized NR

Melt kneaded and specimens prepared

Good heat, impact, solvent resistance and moldability

Fujita and One, 1997

Blend melt kneaded and specimens prepared Turbidity measurement for δ

Tensile strength, elongation, and excellent ethanol resistant δp prediction

Sumitomo Chemical Co, 1997

Solvent uptake by 50/50 blend at 25, 35 and 45°C 60/40 blend injection molded for property testing

Solvent plasticized specimens and induced crystallization of PEEK Good solvent resistance and bond strength

Investigation of miscibility

Strong interactions between amino and sulfonic end groups Tensile strength at yield = 66-77; elongation at break = 11-52%; good resistance to alkali, diesel fuel and chloroform

PEEK/PEI SAN/acrylonitrile-Buacrylate-ethylene glycoldimethylacrylate styrene graft copolymer Aminated PS/sulfonated PEG PC/PP

Tensile properties of complete range of blends before and after immersing in different solvents

face heaters are to be adjusted so that the temperature drops through the two specimens do not differ by more than 1%. To attain a correct value for properties, the time required should be adjusted — its magnitude depends on the specific apparatus, control system and its operation, the test temperatures, the thermal diffusivity and thickness of the specimens [Shirtliffe, 1974]. The conductivity is calculated using the Eq. 12.32. The attainment of equilibrium is important, especially for polymer blends that have low conductivity. The equilibrium times, for example, for cellular materials, are in the order of hours or tens of hours. For this reason, stable over long time-periods power supplies are necessary. 1.6.3

Heat Capacity

Specific heat of polymer blends is usually measured by differential thermal analysis (DTA) [Slade and Jenkins, 1966] or differential scanning calorimetry (DSC) [Strella and Erhardt, 1969; Richardson and Burrington, 1974]. DTA measures

Reference

Ng and Chee, 1997 Browne et al., 1997 Ito et al., 1997

Haraguchi et al., 1996 Manna et al., 1997

the difference in temperature between the sample and a standard for the same rate of heat input, while DSC compares the rate of heat inputs for the same rate of temperature rise. The results of DSC are easier to analyze as they give a direct measure of the rate of heat input. Measurement of specific heat is made by heating a test specimen at a known and fixed rate [Blaine, 1973]. Once dynamic heating equilibrium of the specimen is reached, the heat flow is recorded as a function of temperature. This heat flow, normalized to specimen mass and heating rate, is directly proportional to the specimen’s specific heat capacity. In practice, two thermal experiments are required for each measurement. In the first, a baseline run is performed on only the empty pan and lid. In the second run, the test specimen is enclosed in the pan and lid. The specific heat capacity information is derived from the difference between the two resulting thermograms. Heat flow calibration of the apparatus is also required. This is obtained by running baseline

910

S. F. Xavier

Table 12.26. Chemical and solvent effects of most of the commercialized polymer blends [Utracki, 1994; Toensmeier, 1994] No. Blend’s Name (Manufacturer)

Composition

1.

Alcryn (Du Pont, USA)

2.

Arloy 100 (Arco, USA)

3. 4.

Chemigum TPE (Akron, USA) Delrin (Du Pont, USA)

5.

Durethan (Miles Inc., USA)

6.

Fomion (A. Schulman, Inc., USA)

7.

Geolast (Advanced Elastomer System, USA) Grilamid (Ems-Chemie AG, Switzerland) Hostalloy (Hoechst Celanese, USA)

Halogenated ethylene inter-polymer alloy Resistant to acids, alkalis, oils, gasoline and alcohols Mechanical blends of PC with Good resistance towards acids, alkalis, oils copolymer grades and alcohols; poorly resistant in gasoline NBR/CPE or NBR/TPU or NBR/PA Resistant to oil and gasoline POWPU (some grades contain Resistant to alcohols, gasoline and oils; polytetrafluoroethylene) varying resistance in acids, alkalis PA-6/copolymer of ethylene-butyl Resistant to alcohols, gasoline, oils and acrylate/acrylic acid < 15% (BC 30, 40) alkalis; poor resistance in acids PA-6/ polymer of methyl acrylate, 1,3-butadiene and n-butyl acrylate < 50% (BC303, 402) Ethylene-methacrylic acid or ethyleneResistant to oil and gasoline acrylic acid copolymer/polyethylene EAA/PE Nitrile rubber/EPDM Resistant to acids, alkalis, oils, alcohols and gasoline Polyamid-12/aromatic, aliphatic Resistant to acids only polyamide PE having proprietary molecular Resistant to acids, alkalis, oils, gasoline and weight distribution (731 grades) alcohols Elastomer modified and filled blend. Resistant to acids, alkalis, alcohols and oils; PPBC/EPDM, fillers: 2-10 phr; poor resistant to gasoline initiator, sulfur and additives 0.01 to 0.1 phr each PP/EPDM Resistant to acids, alkalis, oils, gasoline and alcohols PA-6/PTFE and PPS/PTFE Resistant to solvents, alkali and acids Resistant to acids, alkali and solvents PBT/ABS 80/20 to 60/40 Resistant to acids, alkalis, gasoline and oils; poor resistance to acetone, tetrahydrofuran Polyethylene/Polyisobutylene Resistant to acids, alkalis, alcohols; poor (LDPE:PIB = 1:1) resistance in oil and gasoline Styrene-acrylonitrile copolymer/grafted Resistant to acids, alkalis, oils, gasoline and acrylic ester elastomer, SAN/acrylic alcohols elastomer PPE/HIPS Resistant to acids, alkalis and alcohols; poor resistance in oils and gasoline Modified PPE / HIPS or PS Resistant to acid, alkali, alcohols; poor in oil and gasoline PPE/PA (mainly with PA-66; Resistant to oil, gasoline, alkali, alcohol compatibilizer and impact modifier and acids often added)

8. 9.

10. IPCL JSA-86 (IPCL, India)

11. Kelburon (DSM Polymers Intl., Netherlands) 12. Lubricomp PA-6 (LNP Eng. Plastics, USA.) 13. Lumax (Lucky Co. Ltd., Korea) 14. Lupolen 0 250 H, (Rheinische Olefinwerke GmbH, Germany) 15. Luran S (BASF AG, Germany)

16. Luranyl (BASF AG, Germany) 17. Noryl (GE Plastics, USA) 18. Noryl GTX (GE Plastics, USA)

Comments

Properties and Performance of Polymer Blends

911

Table 12.26. continued No. Blend’s Name (Manufacturer)

Composition

Comments

19. Novalloy-S (Daicel Chemicals, Japan) 20. Polypur (A. Schulman, Inc. USA) 21. Polystyrol (BASF AG, Germany)

PC/ADS and a compatibilizer Thermoplastic polyurethane alloys PS/Polybutadiene

22. Prevex (GE Plastics, USA)

Modified PPE/HIPS

23. Rovel (Uniroyal, USA)

SAN copolymer/saturated olefinic elastomer, SAN/EPR PET/Elastomer

Resistant to acid, oil, alcohol and gasoline; fair-to-poor resistant to alkalis Resistant to oils, gasoline and alkalis Resistant to acid, alkali, alcohol and oil; poor resistance in gasoline Resistant to acid, alkali, alcohol and oil; poor resistance in gasoline Resistant to acid, alkali, alcohol and oil; poor resistance in gasoline Resistant to alcohol, gasoline and oils; good-to-poor in acid, alkalis Resistant to acid, alkali, alcohol and oil; poor resistance in gasoline Resistant to acid, alkali, oil and alcohol; poor in gasoline Resistant to acid, alkali, alcohol, oil and gasoline Resistant to oil and alcohol; poor in acid, alkali and gasoline Resistant to alkali, oil, gasoline and alcohol; good-to-poor in acids Resistant to acid, alkali, oil, gasoline and alcohols Resistant to acid, alkali, oil, gasoline and alcohols Resistant to acid, alkali, oil, gasoline and alcohols Resistant to acid, alkali, oil, gasoline and alcohols Resistant to acid, alkali, oil, gasoline and alcohols Resistant to acid, alkali, oil and alcohols; poor in gasoline Resistant to alcohol

24. Rynite (Du Pont, USA) 25. Stapron N (DSM Polymers Intl., Netherlands) 26. Styroblend (BASF AG, Germany) 27. Sumiploy (Sumitomo, Japan) 28. Terblend S (BASF AG, Germany)

ADS/PA-6 PS/poly(2,6-dimethyl-1,4-phenylene ether) and a compatibilizer PES/PEEK

30 Ultem LTX (GE Plastics, USA)

Acrylonitrile-styrene-acrylate copolymer/PC, ASA/PC PA-6/ABS; PA-66/ABS; PA-6 and PA-66/ABS PEI/PC

31. Ultrablend S (BASF AG, Germany)

PDT/ASA

32. Ultraform (BASF AG, Germany)

Polyacetal/polyurethane

33. Vinidur (BASF AG, Germany) 34. Xenoy (GE Plastics, USA)

PVC impact modified by acrylate graft copolymer (with compatibilizer) PC/PBT; PC/PET

35. XT (Cyro Industries, USA)

Acrylic based multipolymer compounds

36. Zyler (Novacor, USA)

Styrene methyl methacrylate copolymer/ elastomer block copolymer, SMMA/SBS PA-6,6/ionomer, PA 6,12/ionomer, Resistant to alcohol, gasoline, oils; PA-6/PA-66, PA-6 or PA-6,6 excellent-to-poor resistant in alkali, or PA-6,12/elastomer limited in acids

29. Triax 1000 (Monsanto, USA)

37. Zytel (Du Pont, USA)

912

S. F. Xavier

Table 12.27. Thermal conductivity units [Ives et al., 1971] Units

Cal/ (cm s °C)

W/ (cm °C)

W/ (m °C)

Kcal/ (m h °C)

Cal cm/(cm2 s °C) (or) cal/(cm s °C)

1 0.230 0.00239 0.00278 0.000345

4.19 1 0.01 0.0116 0.00144

419 100 1 1.16 0.144

360 86 0.86 1 0.124

J cm/(cm2 s °C) (or) W/cm °C J cm/(m2 s °C) (or) W/m °C Kcal m/(m2 h °C) (or) Kcal/m h °C Btu in/ft2 h °F

Btu in/ (sq ft h °F) 2900 693 6.93 8.06 1

Figure 12.28. Possible configurations to restrict edge heat loss and gain [ASTM C177].

and experimental traces for a material whose specific heat capacity is well known. Sapphire is the calibration material of choice since it is easily available and its specific heat capacity is accurately known. The relationship for calculation of specific heat capacity is given by: Figure 12.27. General features of the metal-surfaced hot plate apparatus [ASTM C177].

Cp = KEq/mβ

(12.29)

Properties and Performance of Polymer Blends

where Cp = specific heat capacity (J/kg K), E = calibration coefficient (dimensionless), q = heat flow (mW), β = heating rate (K/min), m = specimen mass (mg), and K = 60,000 is the conversion constant. The calibration coefficient E is expressed as: E = Cp(lit)/Cp(obs)

(12.30)

where Cp (lit) is the reported literature value of specific heat capacity of the standard sample, while Cp (obs) is its experimentally observed value under the same conditions. The coefficient, E, is to be used for the determination of the unknown. It is wise to calculate E for several temperatures over the region of interest — E should be constant. If it is not, particularly at low temperatures, it indicates that dynamic temperature equilibrium is not attained at the temperature and that the experimental temperature program needs to be started at a lower temperature. The overall accuracy of the method is found to be ±5.5%. Precision can be improved with large samples and higher heating rates, provided dynamic temperature equilibrium is achieved. 12.6.4

Review of the Domain

Owing to the absence of electronic effects in most polymers, heat conduction occurs as a result of lattice vibrations, similar to dielectrics. It is known that the thermal conductivity of an amorphous polymer increases to Tg with increasing temperature while it decreases above Tg [Godovsky, 1992]. Thermal conductivity is a fundamental and important factor in processing polymer blends [Agari, 1992]. Influence of miscibility on thermal conductivity and diffusivity was studied [Agari, 1993, 1994]. In the blend of low molecular weight PS with coumarone-indene resin, which showed miscibility over all blend compositions, the thermal conductivity was approximately linearly dependent on composition [Agari, 1993]. Thermal conductivity, thermal diffusivity and heat capacity of PMMA/PC blends were studied with respect to temperature and blend composition. The specific

913

heat capacity of the two phase 50/50 blend was larger than that of the one-phase blend. The thermal diffusivity and the conductivity of the 50/50 blend slightly decreased with the increase of temperature up to 450-460K (LCST) and then decreased abruptly with increase of temperature [Agari et al., 1997]. Several investigators [Krause et al., 1982, Schultz and Young, 1983, Rodriguez-Parada and Percec, 1986] had used the specific heat increment (∆Cp) to investigate the polymer - polymer miscibility by DSC. The effect of molecular weight of PMMA on the miscibility of PMMA/PS blends was examined by studying the specific heat increment ∆Cp at Tg [Burns and Kim, 1988]. Using Couchman’s equation, Cp for PMMA was calculated and was found to decrease with the composition of PS [Couchman, 1978]. The Cp for PS similarly decreased with PMMA composition. From these results the authors inferred that some of the PMMA dissolved in the PS phase and vice versa. Thus, the blends were found to be partially miscible. This result was found to be consistent with the polymer-polymer interaction parameter values. The authors also studied the PC/SAN blends miscibility by the thermal analysis [Kim and Burns, 1988]. The values of the specific heat increment ∆Cp at Tg for PC and SAN in PC/SAN blends were measured. For PC, Cp decreased linearly with addition of SAN. For SAN, Cp also linearly decreased with addition of PC. This suggested that some of the PC dissolved in the SAN-rich phase and vice-versa. 12.6.5

Heat Distortion Temperature (HDT)

ASTM, D648 describes the determination of temperature of deflection under load for plastics and ebonite. ISO 75, BS 2782 Method 121 A and 121 B are equivalent. DIN 53461 is similarly related to the ISO method. Since these standards are similar, only the ASTM method will be described. The heat distortion temperature (HDT), the deflection temperature under load (DTUL), or the softening temperature are practical and important parameters of a polymeric material. They denote

914

S. F. Xavier

the upper temperature limit up to which the material can support a load for any appreciable time. ASTM D648 provides a method for determining DTUL of plastics under flexural load. The method is applicable to molded and sheet materials available in thickness ≥ 3 mm, which are rigid at room temperature. The specimen is taken in the form of a rectangular bar with the load applied at its center to give maximum fiber stresses of 455 or 1820 Pa (see below). The metal supports (rounded to a radius of mm) for the specimen are provided 100 mm apart (see Figure 12.29) allowing the load to be applied on top of the specimen vertically and midway between the supports. The specimen is immersed under load in a heat-transfer medium. The temperature is raised at 2 ± 0.2 K/min. The load applied on the specimen to obtain a maximum fiber stress of 1820 ± 2.5% Pa is calculated as follows: P = 2 S b d2 / 3 L

(12.31)

where P = load (N), S = maximum fiber stress in the specimen (1820 or 455 Pa), b = width of specimen (m), d = depth of specimen (m) and, L = 0.1 m is width of span between supports. The load of 1820 Pa is usually used for rigid polymers (e.g., PS) while 455 Pa is used for softer crystalline materials that have Tg < 298 K (e.g., PE). ASTM D1637 provides tensile HDT test for plastic sheets. In this test a load of 345 Pa is applied to a strip, and the temperature is increased at a rate of 2 K/min. The HDT in this case is defined as the temperature at which the elongation becomes 2%. ASTM D1525 provides a third type of the softening temperature test. A flat-ended needle of 1 mm circular cross-section is pressed into a thick sheet of the polymer with a load of 1.0 kg. The polymer is heated at a rate of either 50 or 120 K/h. The Vicat softening temperature is the temperature at which the needle has penetrated the polymer to a depth of 1 mm. For such a depth of penetration, the material must be very soft, hence the Vicat softening temperature is higher than other HDT data.

Figure 12.29. [ASTM D648].

Apparatus for deflection temperature test

Heat deflection temperature is influenced by: (i) blend composition, (ii) fillers/reinforcing agents, (iii) annealing, and (iv) applied stress: • HDT of a blend is influenced by its composition. Figure 12.30 shows variation of HDT (measured according to ASTM D648) with PC/ ABS blend composition [Xavier and Pendyala, unpublished]. The observed variation is caused by changes in flexural modulus, which is also shown.

Figure 12.30. Variation of heat deflection temperature with PC/ABS blend composition (curves 1 and 2 are obtained with loads 4.6 and 18.2 kg.f/cm2, respectively).

915

Properties and Performance of Polymer Blends

• Inorganic fillers (e.g., talc, mica or CaCO3) or reinforcements (e.g., glass or carbon fibers) increase HDT of neat polymers or blends [Nielsen, 1974; Xavier and Sharma, 1986]. Table 12.28 illustrates the influence of talc concentration on HDT for a blend.

Table 12.28. Variation of HDT with talc concentration in PPCP-EPDM blend [Xavier et al., 1994] Blend composition Neat PPCP PPCP + EPDM (10 wt%) PPCP + EPDM (10 wt%) PPCP + EPDM (10 wt%) PPCP + EPDM (10 wt%)

Talc (wt%)

HDT (K)

0 8 10 12 15

353 355 356 359 360

• Annealing of a crystalline polymer, either neat or in a blend, increases the degree of crystallinity, changes the crystallite morphology and relieves built-in stresses in the amorphous phase. Table 12.29 shows the gradual increase of HDT with annealing time. Flexural modulus also increases. Similar effects were reported for amorphous polymers such as PS [Nielsen, 1974].

Table 12.29. Effect of annealing time on HDT and flexural modulus of PPCP - EPDM (10 wt%) blend (HDT was tested at 455 Pa) [Sarkar, 1989] Test

Heat deflection temperature (K) Flexural modulus (MPa)

Value obtained after annealing (hrs). 0

3

6

330.0 805.2

343.0 882.2

349.4 889.8

• Polymer HDT decreases with applied stress. The major cause of this effect is the decrease of modulus with temperature with the consequent greater deformation at the higher temperature for a given load. HDT occurs, by definition, at a constant deformation. The deformation is proportional to the load and inversely proportional to the modulus. HDT values of some commercial polymer alloys are given in Table 12.30.

Table 12.30. HDT of some alloys with and without reinforcement for car body panels and bumpers [Moro et al., 1988] No. Blend

1. 2. 3. 4. 5. 6.

HDT (K); ASTM* tests at 455 Pa

PPE/PA (NORYL GTX 900) PBT/Elastomer PBT/Elastomer, Glass Reinforced PP/Elastomer PP/Elastomer, Glass Reinforced PC/PBT

456 348 421 333 403 393

Note: *ASTM D648-01; see: http://enterprise.astm.org/

12.7

Flammability

12.7.1

Standard Methods of Measurement

There are two types of tests, viz. burning and combustion toxicology tests. The burning tests aim at determining either the burning characteristics or the burning rate. The combustion toxicology tests aim at measuring the types and quantities of toxic gases that evolve during burning and smoldering of plastics, and their effects on men and animals. A list of ASTM tests, the specimens, the purpose of each test, etc. is provided in Table 12.31. Underwriters’ Laboratories (UL) provide Standards UL 94 Tests for Flammability of Plastic

Specimen/sample

Purpose of the test

11. ASTM D3801-96 for measuring the comparative extinguishing characteristics of solid plastics in a vertical position.

8. ASTM D2863-97 for measuring the minimum oxygen concentration to support candle-like combustion of plastic. The ratio O2/(O2 + N2) when multiplied by 100 is designated as the oxygen index [Imhof and Steuben, 1974]. 9. ASTM D3014-99 for flame height, time of burning and loss of weight of rigid cellular plastics in a vertical position. 10. ASTM D3713 for measuring response of solid plastic to ignition by a small flame. A set of specimens of identical composition and geometry. Solid plastic material. A set of specimens with identical composition and geometry.

Rigid cellular plastics.

Various forms such as films, etc.

4. ASTM D757 for incandescence resistance of rigid Rigid plastic. plastics in a horizontal position. 5. ASTM D1433 for rate of burning of flexible thin plastic Flexible plastic in the form sheeting supported in a 45° incline. of film or thin sheeting. 6. ASTM D1929-96 for ignition properties of plastics (Setchkin technique). 7. ASTM D2843-99 for density of smoke from the burning or decomposition of plastics.

Discontinued in 1991, not replaced. This method combined with the best features of UL 94 resulted in writing of ASTM D3801. Discontinued in 1966, no replacement. Discontinued in 1987, replaced by D4549.

Comments

Determination of comparative extinguishing characteristics.

To characterize the response of a plastic to a small flame of controlled intensity for quality control.

Combination of the best features of UL 94 and ASTM D635.

Discontinued in 2000, no replacement.

Relative resistance to incandescent surface at 1223 ± 10 K (1742 ± l8°F). Relative rate of burning and/or extent and time of burning. Determination of self-ignition, flash-ignition temperatures and self-ignition by temporary glow. To measure smoke- density across a 12 inch light path. For materials that excessively drip, auxiliary burner is used. To determine relative flammability of plastic by Useful for determining measuring the minimum concentration of oxygen the ‘Limiting Oxygen in a flowing mixture of oxygen and nitrogen that Index’ of plastics.It has will just support flaming combustion. gone through several The apparatus is shown in Fig.12.33. modifications. Determining relative extent and time of burning. Revised to accommodate thermosets.

Flat sheet or plate form. Relative comparison of the ignition resistance of materials and the extent of burning. 2 ASTM D568 for rate of burning and/or extent of burning Flexible thin sheets or films. Relative comparison of rate of burning and/or of flexible plastics in a vertical position. extent and time of burning (of plastics). 3. ASTM D635-98 for rate of burning and/or extent of Bars either molded or cut Relative comparison of average burning rate, burning of self-supporting plastics in a horizontal from sheets, plates average time of burning and average extent position. or panels. of burning.

1. ASTM D229-96 Testing rigid sheet and plate materials used for electrical insulation.

No Test method

Table 12.31. Summary of ASTM test methods*

916 S. F. Xavier

Specimen/sample

12. ASTM D3894 for evaluation of fire response of Rigid cellular plastic. rigid cellular plastics using a small corner configuration. Full-scale Corner Wall Test. 13. ASTM D4100 for gravimetric determination of smoke Plastic material in particulates from combustion of plastic materials. a slab configuration. 14. ASTM E84-00 for surface burning characteristics Any building material of of building materials. dimensions 24 ft × 20.25 in 15. ASTM E 119-00 Methods of fire tests of building Full-size wall section. construction and materials. 16. ASTM E136-99 for behavior of materials in Building material test specimens of size a vertical tube furnace at 705 °C. 1.5 × 1.5 × 2 in 17. ASTM E162-98 for surface flammability of materials Specimen of dimensions using radiant heat energy source (Radiant Panel Test). 6 × 18 in.

No Test method

Table 12.31. continued

Determination of flame spread index of a material.

Determination of fire resistance of walls, floors, ceilings, roofs, etc. Determination of combustion characteristics of building materials.

Gravimetric determination of smoke particulate matter produced from the pyrolysis of plastics. Determination of surface burning characteristics, e.g., of foam insulation.

Prediction of performance of a Factory Material

Purpose of the test

Intended for research and development only.

Discontinued in 1994 and not replaced. Discontinued in 1997 and not replaced. Suffers from several limitations. Also known as Steiner Tunnel Test. It is similar to UL 263 and NFPA 251.

Comments

Properties and Performance of Polymer Blends

917

918

S. F. Xavier

Materials for Parts in Devices and Appliances. The standard is important for classifying polymeric materials (including polymer blends and alloys) for the use in electrical applications. It is widely used and the results are reported in the literature and in company catalogues. However, the requirements are not applicable to polymeric materials used in building construction or finishing. The tests conducted under this standard are summarized below. 12.7.1.1

Horizontal Burning Test for Classifying Materials, 94 HB

The test uses small bar specimens: 127×12.7 mm. It is similar to ASTM D635, mentioned in Table 12.31. Materials classified under this test shall not have a burning rate exceeding either 38 or 75 mm/min over a 75 mm span, for specimen thickness of respectively 12.7 or 3 mm. The materials must cease burning before reaching the 100 mm mark. 12.7.1.2

Figure 12.31. Test layout for classification in 94 V-0, 94 V-1 and 94 V-2 according to UL 94 [Troitzsch, 1983].

Vertical Burning Test for Classifying Materials, 94 V-0, V-1, V-2.

The test specimen (127x12.7 mm, with maximum thickness 12.7 mm) is supported vertically by its upper end and is ignited at its lower end for 10 s by Bunsen/Tirril burner, in a draft-free area (see Figure 12.31). If flaming or glowing combustion stops within 30 s after removal of the flame, the specimen is re-ignited for 10 s. The duration of the flame is again noted. In case the specimen drips flaming particles, they are allowed to fall into a layer of surgical cotton 0.3 m below the sample. The particles are considered significant if the cotton ignites. Flammability ratings are based on the specimen behavior during the test, materials rated 94 V-0 being the most while those rated 94 V-2 being the least resistant to burning. Table 12.32 summarizes the test requirements.

Table 12.32. UL 94 Vertical burning test [Landrock, 1983] No. Requirement

1. Total flaming combustion time for 10 ignitions, maximum (sec) 2. Individual flaming time, maximum (sec) 3. Glowing combustion time (sec) 4. Flame drippings

Classification 94 V-0

94 V-1

94 V-2

50

250

250

10

30

30

30

60

60

None

None

Allowed if burns briefly

Properties and Performance of Polymer Blends

12.7.1.3

Vertical Burning Test for Classifying Materials 94 5V

This test is more stringent than UL 94 V-0, V-1 and V-2. Here, 127 mm ignition flame is applied on specimen bars of dimensions 127×12.7 mm, with maximum thickness 12.7 mm. In Method A, a Tirril burner is positioned 20 degrees from the vertical and the overall height of the flame is adjusted to 127 mm. The flame is applied for 5 s and removed for 5 s. The procedure is repeated 5 times. After the 5th removal of the flame, the duration of flaming and glowing, the distance the specimen burned, dripping of particles from the specimen (during the test) and deformation of physical strength of the specimen immediately after burning (and when cooled) are observed. When the specimen shrinks, elongates, melts, etc., additional tests are carried out using test plaques (152×152 mm) under Method B. These tests are conducted in various positions both vertical and horizontal positions of the plaques with flame applied to different places in the plaques, using the procedure as in Method A. The observations focus on the same items as in Method A. Materials are classified 94-5V when: • No specimen burns with flaming and/or glowing combustion (after the fifth flame) for more than 60 s. • None specimen drips particles.

12.7.1.4

Flame Spread Index Test Using Radiant Panels

This test is conducted in accordance with ASTM E162 mentioned in Table 12.31. 12.7.1.5

Vertical Burning Test for Classifying Materials 94 VTM-0, -1 or -2

Some materials due to their thickness, distort, shrink or get consumed up to the holding clamp, when tested according to the methods described above. VTM means “very thin materials” — test specimens are cut to 200×50 mm dimension.

919

Each specimen is supported from the upper 6 mm of its length, with the longitudinal axis vertical using a heavy spring clamp. The lower end of the specimen is placed 9.5 mm above the top of the Bunsen burner tube and 0.3 m above a horizontal layer of dry surgical cotton. The test flame is placed under the lower end of the test specimen for 3 s. Then the flame is taken away from the specimen and the duration of specimen flaming is noted. When flaming of the specimen ceases, the test flame is applied once again for 3 s and then withdrawn. The deviation of flaming and glowing of the specimen is noted. In case the specimen drips molten or flaming material, the burner may be lighted to angle up to 45°. The following are observed after the removal of the flame: • Duration of flaming after first flame application. • Duration of flaming after second flame application. • Duration of flaming plus glowing after second flame application. • Whether or not specimens burn up to 127 mm. • Whether or not specimens drip flaming particles that ignite the cotton. 94 VTM-0 classifies materials under most severe conditions. Table 12.33 summarizes the test requirements.

12.7.1.6

UL 746A-78 Polymeric Materials Short Term Property Evaluations

Test procedures for seven major areas are given in this standard. They are all applied for the determination of resistance of polymeric materials to ignition from electrical sources. The individual tests are: Resistance to Hot Wire Ignition, Resistance to High-Current Arc Ignition, Resistance to High-Voltage Arc Ignition, and Resistance to Hot-Bar Ignition. 12.7.1.7

UL 746B-79 Polymeric Materials Long Term Property Evaluations

This standard deals with long-term tests for the evaluation of materials and parts of end products.

920

S. F. Xavier

Table 12.33. UL 94 Vertical burning test for classifying very thin materials [Landrock, 1983] No. Requirement

1. Total time of flaming combustion for 10 ignition, max. (sec) 2. Individual flaming time, max. (sec) 3. Glowing combustion time (seconds) 4. Flame drippings

Classification 94 VTM-0 (most severe)

94 VTM-1 (intermediate)

94 VTM-2 (least severe)

50

250

250

10 30 None

30 60 None

30 60 Allowed if burns briefly

Along with UL 94, UL 746A and UL 746C, these tests provide data regarding the physical, electrical, flammability, thermal and other properties of the materials under consideration. 12.7.1.8

UL 746C-78 Polymeric Materials Use in Electrical Equipment Evaluations

This is the test procedure, including flammability, for parts of polymeric materials used in electrical equipment. It provides a table of shortterm and long-term properties to be considered during evaluation of polymeric materials used in electrical equipment. 12.7.1.9

UL 746D-80 Polymeric Materials Fabricated Parts

This standard is for polymers, blends of polymers, copolymers, terpolymers and alloys. It considers plastic parts that have been produced under a material identity control system. Molders/fabricators are required not to employ such additives/flame retardants that would adversely affect critical material properties. A detailed discussion on national and international fire protection regulations and test methods for plastics is presented by Troitzsch [1983].

12.7.1.10 ASTM D2863 Measuring the Minimum Oxygen Concentration The method provides means for the determination of relative flammability of plastics by varying the oxygen to nitrogen concentration (see Table 12.31). The oxygen indexer is shown in Figure 12.32. 12.7.2

Factors Affecting Flammability

Application of a heat source, such as flame, raises the temperature of polymer and ultimately causes it to burn. Burning of a solid polymer has been divided into four stages: i) heating, ii) decomposition, iii) ignition, and iv) combustion [Landrock, 1983]. In the first stage, a thermoplastic material softens or melts and begins to flow. The temperature at which it melts can have a significant effect. In the second stage, gases or the volatile fragments of degraded polymer are removed. The temperature and the rate at which this occurs depend on the thermal stability of polymer and the chemical reactions occurring under those conditions. Ignition takes place as the flammable gases combine at appropriate ratios with oxygen from the air. Sustaining the burning depends on the transfer of sufficient heat from the flame to polymer, capable to maintain supply of flammable decomposition byproducts. Supply of oxygen is also essential. If decomposition of the polymer requires

Properties and Performance of Polymer Blends

921

to a higher tendency for condensation into aromatic chars and therefore, to the less flammable products [Fenimore and Martin, 1966]. Van Krevelen had confirmed the empirical relationship between polymer structure, charformation and polymer flammability. A mathematical formula was proposed that (based on structural units) allows calculation of the oxygen index and char residue values for a wide variety of hydrocarbon polymers. The very existence of such a relationship indicates that pyrolytic condensedphase processes are of primary importance in determining polymer flammability at least in the studied cases [Van Krevelen, 1975]. A relationship between the polymer structure and its flammability was related to unsaturation for co-polyterephthalates and co-polycarbonates [Quinn, 1977]. This work is an excellent illustration of the importance of condensed-phase pyrolytic mechanisms upon polymer flammability. Figure 12.32. Oxygen index apparatus [Troitzsch, 1983].

12.7.3

Prevention Methods

more heat than it is supplied by the flame, or if solid non-flammable residues coat the surface and insulate the remainder of the flammable part, a continuous propagating flame will not be obtained. Thus, the last stage of the burning sequence, very much depends on the polymer characteristics. It may be correlated with such energy factors as cohesive energy, hydrogen bonding, heat of combustion and dissociation energy [Einhorn, 1972]. An interesting relationship between polymer structure and polymer flammability has been observed. Commercial polymers that possess aromatic groups in the main chain (e.g., PPE, PC, PSF, phenolic resins) undergo char-forming condensed-phase reactions — as a result they have low flammability. Higher oxygen index of PC and PPE was apparently related to their higher charring tendency in comparison to the aliphatic hydrocarbon-type polymers. The greater thermal stability of aromatic-type polymer backbone leads

Since a thorough review of fire-retardant methods is beyond the scope of this chapter, only a brief summary is given below. The readers interested in a more detailed discussion are referred to pertinent reviews [Einhorn, 1972; Hilado, 1972, 1981; Vendersall, 1971]. Four general methods for reduction of polymer flammability have been identified [Kuryla and Papa, 1978]: • A nonflammable coating that prevents the normal pyrolytic or combustion mechanism, is either applied to the polymer surface or produced in the presence of a flame. • Appropriate chemicals are incorporated during polymer processing. Their role is to either alter the rate of pyrolytic fuel generation, or inhibit the exothermic gas-phase reactions. • Gas-phase flame reaction can be prevented by the generation of non-flammable gases, which dilute the fuel gases below the flammability limits. • Incorporated solid components consume sufficient heat during pyrolytic decomposition that they sufficiently cool the substrate to a temperature below the ignition point.

922

S. F. Xavier

In any given fire retardant one or more above methods may be used. The effect of a fireretardant strongly depends on the basic chemical structure of the polymeric material. Owing to complexity of the processes and the experimental limitations, it is difficult to predict which mechanism is most important or operative for any system. A list of commercially available fireretardants is given in Appendix-1. These materials are classified as organic, inorganic and reactive types. A fact to be kept in mind is that for blends or alloys, the fire retardancy behavior is usually between those of the base resins; for example, consider Arylon and Kydene (acrylic/PVC) [Landrock, 1983]. 12.7.4

Review of the Domain

The concept of fire-retardancy is remarkably old. The Greek historian, Herodotus, in 484-431 BC recorded that the Egyptians imparted fire-resistance to wood by soaking it in a solution of alum (potassium aluminum sulfate) [Browne, 1958]. The Romans added vinegar to the alum for the same purpose. Vitruvius in the first century BC described the natural fire-retardant properties of the larch tree and some military applications of fire retardant materials such as plaster of clay reinforced with hair [Vitruvius, 1960]. In 1638, Circa recorded that Italian theaters were painted with a mixture of clay and gypsum (potassium aluminum silicate and hydrated calcium sulfate) to protect them from fire. Wild was issued a British patent in 1735 for his process of treating wood with a mixture of alum, ferrous sulfate and borax (sodium tetraborate decahydrate). And Gay-Lussac in 1821 showed that a solution of ammonium phosphate, ammonium chloride and borax acts as a fire-retardant for wood. In all these processes the key ingredients are the elements from group III (B and Al) of the periodic table. Now, at the end of the 20th century, with so much of research activity for better fire retardants, the most effective elements are still found in groups: III (B and Al), V (N, P and Sb) and VII (Cl and Br). Research efforts to find new and improved fire-retardant agents for synthetic

polymers and their blends have been concentrated on the same three groups of the periodic table, with the same seven elements. The search is for new ways of incorporating them into polymers [Chamberlain, 1978]. Certain compounds based on Ba (group II), Zn (group II B), Sn (group IV) are claimed to be effective in some polymers, especially when used in conjunction with one or more of the seven key elements mentioned above. The burning or non-burning characteristics of plastics have been given a great deal of attention by the scientific community. After the Federal Trade Commission (FTC) announced inquiry into flammability of plastics in October 1972, the suppliers started more carefully to describe flammability of their products. ASTM and other standards developing groups have given considerable effort to develop more meaningful tests and have dropped or modified certain tests. Thus, ASTM D1692 was discontinued [Hendersinn, 1977]. A theory, that certain flame retardants vaporize and produce an effect by acting as free-radical chain stoppers to extinguish the flame or to inhibit the flame speed of the burning gases, was proposed. It is based on extensive studies for 30-40 years [Kuryla and Papa, 1978]. The research efforts devoted to understand the mechanisms of combustion and inhibition for solid materials burning with a diffusion flame in an air environment has multiplied rapidly in the last two decades. 12.7.5

Data on Blends

The flame retardancy properties of some commercial polymer blends are given in Table 12.34.

12.8

Other Properties

12.8.1

Electrical Properties

12.8.1.1

Standard Methods of Measurement

12.8.1.1.1 Resistivity of Insulating Plastics Measurements of insulating polymers or polymer blends are usually carried out using a sheet specimen in form of a disc or a square [ASTM D257].

923

Properties and Performance of Polymer Blends

Table 12.34. Data on blends Blend (Trade Name, Grade and Manufacturer)

UL 94 Flame class HB rating (inch)

PC/ABS Cycoloy, GEC ASA/PC Geloy, XP 1001, XP 2003, XP 4001, GEC PC/PBT Xenoy, 6120, 6240, 6123, 6370, 6620, 6380, GEC. PC/ABS Bayblend, Miles; FR 1439 FR 1440, FR 1441

0.060 0.063 0.061 — —

V-0 rating (inch) 0.060 0.130 V-1 — 0.062 0.062

5 VB rating (inch) 0.098 0.130 VA — — —

Oxygen index ASTM D2863-97

< 21.0 — — 28 30

Commercially available resistance meters can measure resistance in the range from 106 to 1015 Ω. In the case of plastics, the method can be applied if the resistance values are of the same order or lower than the volume resistance, and if the volume resistivity is ρv >108 Ω. Flat plate metal electrodes, preferably guarded (Figure 12.33), are used for testing flexible and compressible materials (at room or elevated temperatures). 12.8.1.1.2 Voltmeter-Ammeter Method The D-C voltmeter and the D-C amplifier (or electrometer to increase the sensitivity) are connected to the voltage source and the specimen. The applied voltage, Vx, is measured by a D-C voltmeter. The current, Ix, is measured in terms of voltage drop across a standard resistance Rs. The voltage drop is amplified by the D-C amplifier and read on an indicating meter as Vs. The resistance Rx, or the conductance Gx is calculated as: Rx = 1/Gx = Vx/Ix = (Vx/Vs)Rs

(12.32)

The time of electrification, unless otherwise specified, should be 60 s and the applied direct voltage Vx = 500 ± 5 V. 12.8.1.1.3 Volume Resistivity or Conductivity Measure the dimensions of the electrodes and width of the guard gap, g, accurately. Unless otherwise specified, the time of electrification should be 60 s, and the applied direct voltage

Figure 12.33. Top view and side view of flat plate guarded metal electrodes [ASTM D257].

500 ± 5 V. Volume resistivity is expressed as (Ω -cm): ρv = A/t Rv

(12.33)

924

S. F. Xavier

where A = the effective area of the measuring electrode (see Figure 12.33), t = average thickness of the specimen, Rv = measured volume resistance in Ω. Volume conductivity is calculated as (s/cm): νv = t/AGv

(12.34)

where Gv = measured volume conductance in Siemens. 12.8.1.1.4 Surface Resistivity or Conductivity The electrode dimensions and the distance between the electrodes, g, are to be measured accurately. The surface resistance or conductance between electrodes No. 1 and 2 is measured with a suitable device [Brown, 1981]. The time of electrification should be 60 s and the applied direct voltage shall be 500 ± 5 V. Surface resistivity, per square cm is given by: ρs = P/g Rs

(12.35)

where P = the effective perimeter of the guarded electrode (see Figure 12.33), Rs = measured surface resistance in Ω, g is as indicated in Figure 12.33. For specimens of square, rectangular and tube forms, appropriate electrodes and mathematical relations are given in ASTM D257. BS 4618 recommends preconditioning of the test specimens at not more than 1% relative humidity (RH) for the study of effect of temperature. ASTM D257 covers resistivity measurements for insulating materials. Electrode sizes are not stipulated (round, square and rectangular types are permitted). The gap between guard ring and center electrode is made approximately equal to twice the specimen thickness. The test voltage is usually 500 V applied for 60 s, as in the British test. DIN 53482 uses methods similar to some of those in IEC 93, using silver or graphite painted electrodes for volume resistivity. A different electrode system was suggested for the measurements of surface and volume resistivity. A narrow guard gap of 1 mm makes it difficult to avoid short-circuiting the electrodes.

12.8.1.1.5 Power Factor and Permittivity Measurements of dielectric constant and loss in polymeric solids and melts over a wide frequency range was described in detail elsewhere [Porter and Boyd, 1972]. The measurement of power factor and permittivity and the related parameters such as dissipation factor, phase angle, etc., may need to be carried out over a wide range of frequencies from a few Hz to several tens of GHz. However, most measurements are made between 50 Hz and 100 MHz [Brown, 1981]. Insulating materials, such as polymers and polymer blends, are used as dielectrics at commercial frequencies between approximately 50 Hz and 100 GHz. Two different techniques are adopted to study the dielectrics in two ranges, i.e., below and above 100 MHz. Dielectric measurements at ultra-low frequencies are of some interest, as they reveal the basic structure of the material [McCrum et al., 1967]. Bridge circuits are invariably employed for the measurements of power factor and permittivity. The test specimen, whose dielectric constant and loss factor are to be measured, is cut or molded to a suitable shape and thickness determined by the material specifications or the test method. The thickness of the specimen must be accurately measured. The electrodes are selected, based on convenience and whether or not the specimen must be conditioned at high temperature and high relative humidity. The test specimen with its attached electrodes is placed in a suitable measuring cell and its capacitance and a-c loss are measured using a suitable bridge. For routine work when either the highest accuracy is not required, or when neither terminal (of the specimen) is grounded, it is not necessary to place the solid specimen in a test cell. In the Schering Bridge, one sets the resistance ratio R1/R2 (range) and varies Cs and C1 to obtain a balance [ASTM D150]: Cx = (R1/R2) Cs

(12.36)

Dx = ωC1R1

(12.37)

Properties and Performance of Polymer Blends

where Cx = capacitance of the specimen, Cs = standard capacitance, Dx = dissipation factor (tan δ) of the specimen, ω = (2π × frequency) and C1 = variable capacitance. The test method covers dielectric measurements from 1 Hz to several hundred MHz. It has few recommendations about the procedure or apparatus. Size of the electrodes is not suggested and it is recommended that the guard gap should be as small as possible and the guard width should be at least twice the thickness of the specimen. An Appendix is provided, which describes number of bridges and their circuits. The German standard for dielectric measurements is DIN 53483. 12.8.1.1.6 Dielectric Strength Low-level conduction in insulating materials can originate in a variety of ways. Often it is attributed to impurities that provide small concentrations of charge carriers in the form of ions and/or electrons. At high fields, the electrodes may also inject new carriers into the polymer, causing the current to increase more rapidly with voltage. At very high fields, these and other processes inevitably lead to complete failure of the polymer as a dielectric. This localized, sudden and catastrophic phenomenon is known as the dielectric breakdown [Ku and Liepins, 1987]. In many cases the dielectric breakdown or dielectric strength of a material can be the determining factor in the design of an apparatus in which it is to be used. A method for determination of dielectric strength of solid electrical insulating materials at commercial power frequencies was developed [ASTM D149]. The voltage can be applied at a fast uniform rate or step-by-step or at a slow rate-of-rise. At the dielectric breakdown voltage an abrupt rupture through the specimen results in a visible puncture and decomposition of the material — the occurrence can be seen and heard. This form of breakdown is irreversible. Dielectric strength is calculated (in terms of KV/mm or V/mil) by noting the gradient at the highest voltage step at which breakdown did not occur. BS 2782 Method 201 deals with plastics breakdown voltage. Other British standards such as BS

925

3784, BS 3816, BS 3924 and BS 5102 directly invoke BS 2782. 12.8.1.2

Factors Affecting Electrical Properties

Factors such as improper mixing of polymer blends leads to variations in test specimens. Such compositional and/or structural changes from specimen to specimen often lead to widely divergent data. 12.8.1.2.1 Effect of Temperature Resistivity depends on temperature — for nonmetallic materials it invariably decreases with T. Volume resistivity is markedly more sensitive to temperature, than surface resistivity. In any measurement it is important to ensure that temperature is maintained constant during the test. Temperature fluctuations produce changes in measured current and lead to significant errors [Brown, 1981; ASTM D257]. In the case of power factor or permittivity determinations, the effects of temperature and frequency are interrelated. With nonpolar polymeric materials the changes in properties with temperature and frequency are small while with polar materials very large changes may take place. 12.8.1.2.2 Effect of Humidity The insulation resistance of solid dielectric materials decreases with increasing humidity. Surface resistance is particularly sensitive to humidity changes. Insulation resistance or conductance is a function of both the volume and surface resistance or conductance of the specimen. Surface resistance changes almost instantaneously with a change of RH. Therefore, it is essential to maintain both T and RH within close limits during the conditioning and measuring of the specimens. In a humid environment, absorption of water into the volume of the insulating material as well as the formation of an ionized water film on the specimen surface takes place. These factors lead to significant rise in the insulating materials’ permittivity and loss index. The process of dielectric breakdown in the case of neat polymers, is not completely understood

926

S. F. Xavier

and many unknowns are still remaining [Ku and Liepins, 1987]. Still more uncertainties exist for polymer blends and alloys. However, it is clear that the chemical structure, the solid-state structure, degree of plasticization, the nature and concentration of filler, molecular weight and morphology of a polymer, etc., influence the electrical properties. For example, the dielectric breakdown of EVAc and its blends was studied in the low temperature region (193 K) [Nagao et al., 1976, 1977]. The dielectric breakdown voltage of EVAc and blends was found to rise with an increase in VAc content. This may be caused by increased electron scattering associated with the decrease of crystallinity and the increase of polar groups. The process of polymer blending was used for improving the mechanical properties of electric wire insulation, as early as 1968 [Ku and Liepins, 1987]. In 1978, it was found that the use of a HDPE/LDPE blend enhanced the electrical treeing inception voltage [Nitta and Funayama, 1978]. This approach to electrical treeing inhibition and some of the experimental results on the use of SB/LDPE and HDPE/LDPE blends are given in Table 12.35 [Wu and Chen, 1983]. 12.8.1.3

Review of the Domain

Most polymers have high electrical resistivity, are inexpensive in comparison to other known insulating materials are heat resistant and sufficiently durable. Owing to their sensitivity to oxidation and solvents, they are frequently blended to generate better electrical insulating alloys. In the past two decades, there has been serious effort to modify the electrical properties of polymers and their blends. The electrically conductive polymers can be broadly categorized as (i) electrostatic dissipating polymeric compositions, and (ii) electrically conductive polymer blends. Utracki has reviewed evolution of these materials [Utracki, 1998]. The electrostatic dissipating polymeric compositions (ESD) are developed to overcome problems related to the accumulation of surface charge and its rapid discharge leading to shocks, fire, explosions and damage to electronic components, etc.

These compositions must provide surface resistivity 105 < R < 1012 (Ωcm). The early efforts, to achieve the optimum surface resistivity, such as coating the polymeric parts with electrostatic dissipative materials, addition of either graphite, metal particles or fibers, incorporation of low molecular weight antistatic agents etc., did not yield fruitful results [Kozlowski, 1995]. Antistatic properties are observed for materials having either -SH or -OH groups (e.g., phenolic, alcoholic, or acidic). Since the mid-1980’s the most frequently used ESD has been a copolymer of ethylene oxide and epichlorohydrin, EO-CHR. Many chemical companies (such as Borg-Warner, B.F. Goodrich, Asahi, and General Electric) came up with several ESD formulations containing EO-CHR for improving the electrostatic properties of PVC, CPC, PC, PEST, epoxy, phenolics, etc. [Federl and Kipouras, 1986; Kipouras and Federl, 1988; Barnhouse and Yu, 1988; Yu, 1988; Lee, B. L., 1993; Shimamura and Suzuki, 1991; Giles and Vilasagar, 1994]. The second variety of materials, viz., electrically conducting polymer blends, ECPB, have been known since the early 1980’s. These are prepared, in principle, by synthesis of the conducting polymer within the host or by simple blending [Billingham and Calvert, 1989]. Polyacetylene, PACE, blends were prepared by the polymerization of acetylene in LDPE films doped with Ziegler catalyst [Galvin and Wnek, 1982, 1983; Galvin et al., 1984]. PACE formed a particulate second phase with a size ranging from 60 to 200 nm. Addition of 7% PACE to LDPE increased the yield point of the latter resin from 7 to 10 MPa, but the extension to break was reduced. The effect was greater in blends produced by polymerization in solid PE. Acetylene was polymerized in the presence of polybutadiene rubber and the blends were investigated for their electrical conductivity [Rubner et al., 1983 and Sichel and Rubner, 1985]. The electrical properties of these blends were explained in terms of the morphological features [Tripathy and Rubner, 1984]. In these investigations, a conductivity of 10 S/cm was achieved at PACE loadings above 30%. Polymerization

927

Properties and Performance of Polymer Blends

Table 12.35. Effect of polymer blending on tree growth and length [Wu and Chen, 1983] Parameter Modifier (wt%) Tl (min) T2 (min) T3 (min) Tl + T2 + T3 (min) Length §(µm)

LDPE 100 -0 300 2300 1300

SBR 10 198 100 2180 2478 300

20 127 200 2200 2527 600

HDPE 30 54 230 2284 1400

10 25 156 75 250 450

15 70 170 30 270 500

20 90 160 2100 2350 400

Notes: T1 - induction period for tree initiation in minutes; T2 - growth period of the tree in minutes; T3 - the saturation period of tree development in minutes; § - The length of the tree after saturation period.

of acetylene in EPDM resulted in tough, conductive films [Lee and Jopson, 1983, 1984]. The conductivity of these films was found to be significantly enhanced by stretching. Mechanical blending of various conducting polymers with thermoplastics was studied [Wessling and Volk, 1987]. These materials although processable, showed low electrical conductivity (<10-5 S/cm). They were prepared by blending an electrically conductive polymer (such as polyacetylene or polypyrrole) with a polymer having strong anionic group (such as sulfonated-PE, sulfonated-SEBS, sulfonated-PS or sulfonatedpolyacryl-amidomethyl-propane) [Cross and Lines, 1995]. Electrically conducting polymer blends are also produced by blending another conducting polymer (e.g., poly-3-octyl thiophene) with a matrix polymer (e.g., PP, PVC, PS, PE, EVAc, PVC/ABS etc.) introducing a dopant (e.g., iodine) [Kokkonen et al., 1994]. Several strategies were adopted in preparing ECPBs. In one example, polyaniline was blended with dodecylbenzene sulfonic acid, mixed with PS, PE or PP and then melt processed. In another case, polyaniline was mixed with protoning acid metallic salt. The conductive material was melt mixed with PE, PS, PP or ABS [Karna et al., 1994 a & b]. ECPBs were prepared by melt blending a matrix polymer (selected from between PE, PP, PB, PIB, PMP, EPR, CPE, CSR, PS, poly-alkanes containing styrene, acryl, vinyl, or fluoro-ethyl

groups and their blends) and an electrically conducting thermotropic liquid-crystal polymer, containing either ferric chloride or iodine as a dopant. The blend conductivity, that could be manufactured into fibers and films, was reported as 10-12 - 102 S/cm [Ho and Levon, 1995]. Amineterminated polyaniline was first grafted onto a thermoplastic polymer comprising a functional group capable of reacting with -NH2, e.g., maleic anhydride. They were further compounded with polymers, fillers and/or electrically conductive solids [Jongeling, 1993]. Blends of PVC with “doped” polyaniline and at least one other additive (e.g., impact modifier, plasticizer, acidic surfactant) were developed to give electrically conductive blends. ECPBs were also prepared compounding polyaniline and a thermoplastic polymer [selected from PA-6, PA-66, PA-11, PA-12, PET, PC TPU, CPE, etc.] and 0-10 wt% carbon black [Kulkarni and Wessling, 1992, 1994]. Electrically conductive materials were prepared by dispersing pre-blends of aniline tosylate particles and poly (ethylene terephthalate glycol), PETG, and then diluting the pre-blend with PETG. The blends were useful for the manufacture of films, printing inks, and coatings, in shielding, antistatic and adhesive applications [Shacklette et al., 1993]. Curable blends, with good flame retardancy, comprising of fluorinecontaining polycyanurates and a thermoplastic polymer (e.g., PSF, PPE, PEEK) were patented by IBM in 1992. The filled materials were useful in

928

S. F. Xavier

several applications (fabrication of printed circuit boards, semiconductor chip carriers, metal-core boards, chip modules, and multi-layer thin film circuits) [Ardakani et al., 1994]. 12.8.1.4

Data on Blends

Electrical properties of selected commercial polymer blends are listed in Table 12.36. 12.8.2 12.8.2.1

Optical Properties Methods of Measurement

12.8.2.1.1 Haze and Luminous Transmittance ASTM D1003 provides two methods for measuring light transmittance and haze in planar sections of transparent plastics. The method making use of a hazemeter is briefly described. A spherical hazemeter, which is pivotable about a vertical axis through the entrance port

(where the specimen is placed), is shown in Figure 12.34. In the normal position the collimated incident light passes straight through the sphere, leaving through the exit port, which is closed by an absorbent light trap. When light is scattered either by the instrument alone or by the specimen and the instrument (when the specimen is loaded) it is reflected from the region around the edge of the exit port and finally collected by the photocell after multiple reflections from the highly reflecting walls of the sphere. When the integrating sphere is rotated slightly so that the incident light hits the opposite highly reflecting wall of the sphere adjacent to the exit port, the measurements with and without a specimen gives a measure of the total transmittance. The total transmittance, Tt, is given by: Tt = T2/T1

(12.38)

where T2 = the total light transmitted by the

Table 12.36. Data on some commercial polymer blends [Martino, 1994] No. BLEND (Trade name, composition and manufacturer)

1. 2. 3. 4.

5.

Volume resistivity (ohm-m) ASTM D257-99

CYCOLOY C2950 (PC/ABS) 1.0E15 GE Thermoplastics GELOY GY1220(ASA/PVC) 1.4E12 GE Thermoplastics XENOY-2230, 2735, 5220, 2.1E14-9.5E14 5230, 5720.(PC/PBT) GEC MAKROBLEND (PC/PET) UT 1018 > 2.6E13 UT 400 & UT 403 MILES > 7.4E13 BAYBLEND(PC/ABS) MILES T45MN, T65MN, T44, T64 > 1.0E14 T85MN, T84 > 1.0E14

Surface resistivity (ohm) ASTM D257-99

1.0E15-1.8E15

Dielectric strength oil, 0.125”; V/mil (3.2 mm; kV/mm) ASTM D149-97

Dielectric constant (100 Hz), ASTM D150-98

Dissipation factor (100 Hz), ASTM D150-98

495 (19.3)

3.0

0.0049

476 (19)-490 (17) (in air) 705 (28)-1096 (43) (in air)

3.4-3.5

0.0007-0.0021

2.93-3.60

0.001-0.002

3.06 (l MHz) 3.10 (l MHz)

0.014 (1 MHz) 0.009 (1 MHz)

396 406 > 1.0E14 > 1.0E14

600 600

Note: *ASTM standard test methods are available on web: http://enterprise.astm.org/

0.007 (l MHz) 0.008 (l MHz)

Properties and Performance of Polymer Blends

929

Figure 12.34. Pivotable-sphere hazemeter [ASTM D1003].

specimen, T1 = the incident light without the specimen. Diffuse transmittance, Td, and %-of haze are calculated as: Td = [T4- T3 (T2/T1I)]/T1

(12.39)

Haze (%) = 100 (Td/Tt)

(12.40)

where T4 = light scattered by instrument and specimen, T3 = light scattered by instrument (without specimen). BS 2782 Method 515 A is the British equivalent, which deals with only haze of films. 12.8.2.1.2 Refractive Index When a ray of light passes from one isotropic medium into another, the sine of the angle of incidence makes a constant ratio to the sine of the angle of refraction (both measured with respect to the normal) for all angles of incidence. This dimensionless ratio, while dependent on the wavelength of the light and temperature, it characterizes the two media concerned. ASTM D542 describes two methods, viz., i) refractometric and ii) microscopic, for the measurement of refractive index of transparent organic plastics. Both the methods require optically homogeneous specimens of uniform index. The ASTM recommends that refractometric method is to be

preferred wherever possible, since it is capable of providing results with greater precision (up to four significant figures). The microscopic method depends on operator’s skill and it yields results accurate up to only three significant figures. For the refractometric method, the apparatus consists of an Abbe refractometer, a suitable source of white light and a small quantity of suitable contacting liquid. The test specimen for refractometer method should be 12.7x6.3 mm, with one flat face and one perpendicular surface. The two surfaces (preferably polished) shall intersect along a sharp line (without a rounded edge). The test specimen is attached to the prism of the refractometer with a drop of liquid of refractive index higher than the test specimen by at least 0.01 and it should not soften or dissolve the specimen. ASTM D542 suggests a list of liquids for a variety of plastics. Measurements are to be carried out at specified conditions, 296 ± 2 K, and 50 ± 5 per cent RH. Temperature is to be accurately controlled. For maximum accuracy, Sodium D lines are recommended. ISO R489 suggests two methods based on Abbe refractometer and the Becke line methods. It does not recommend any specimen conditioning procedure prior to the test. DIN 53491 provides practical details relevant to refractometer measurements. It recommends test temperature

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S. F. Xavier

as 293 ± 5 K, like ISO R489. For maximum accuracy it suggests the use of Sodium light. All these methods provide lists of liquids suitable for different plastics. 12.8.2.2

Review of the Domain

Lack of transparency is a significant drawback in the commercially important toughened polymers such as HIPS or ABS [Manson and Sperling, 1981]. Transparency in these materials is lost due to the light scattering at the interface between the phases. The degree of light scattering (turbidity) was found to be a function of the amount of dispersed phase present, its particle size, the ratio of refractive indices of the phases and the wavelength of light. In typical polymer pairs, at a given dispersed phase level, the maximum turbidity was observed in the range of particle sizes considered to be necessary for good impact strength [Conaghan and Rosen, 1972]. If the refractive indices are matched at a particular temperature, small particle sizes greatly increase the temperature range over which scattering is minimized. In other words, a clear blend can be obtained if both phases have identical refractive indices, regardless of the details of the phase morphology [Rosen, 1967]. Thus, for example, clarity in ABS and toughened acrylics was achieved by matching refractive indices of the continuous and dispersed phases [Gesner, 1967]. Impact modifiers for PVC that can impart toughness as well as clarity were described by several workers [Petrich, 1972; Souder and Larson, 1966; Ryan and Crochowski, 1969; and Ryan, 1972]. Recently, formation of transparent blend films cast from solutions of PVC and (PC-PDMS)n multiblock copolymer has created concern [Gorelova et al., 1992]. The finding is interesting because the refractive indices of PVC, PC and PDMS are different and consequently, transparency of the blend films may suggest miscibility of their constituents, despite the fact that PVC is immiscible with PC and PDMS [Krause, 1978]. In other words, these blends are pseudo-miscible and their transparency is caused by the very small size of the dispersed block copolymer phase.

PVC was also found to form transparent blends with other multiblock copolymers [Papkov et al., 1995]. Several examples of the formation of transparent blends were reported. The constituents of these blends are homopolymers of various chemical compositions and flexibility; viz., PS, PMMA and PVC; and the multiblock copolymers are PC-PDMS, PSF-PDMS, PSF-PB, and polytetra- methylene oxide-PB [Papkov et al., 1998]. For copolymers and homopolymers of various chemical structures, the composition range for each type of block copolymer, within which the formation of transparent blend film takes place, is relatively narrow. The phenomenon appears to be similar to the so-called ‘miscibility window’ in some blends with random copolymers [ten Brinke et al., 1983]. Microphase separation and the evolution of the multiblock copolymer phase in the form of smallsized particles (up to 100 nm in diameter) is the physical basis of their transparency. Thus, these transparent blends are considered microheterogeneous systems [Papkov et al., 1998]. Detailed theoretical and experimental investigations are required to understand details of this phenomenon [Sikora and Karasz, 1993].

12.9

Conclusions

This chapter focused on the characteristic properties of polymer blends. Their measurements have reached a mature stage, both in terms of necessary theories and the methods of testing. However, efforts made in enhancing one property are often mitigated by the loss of another property. Intensive study of interrelations between processing, microstructure and properties would certainly enable one to have a good control on performance of an ultimate product. Looking at the likely advances in understanding of these materials and their integration with developing technology, polymer blends will continue to receive increasing acceptance for a variety of applications.

Properties and Performance of Polymer Blends

12.10

Acknowledgments

We gratefully acknowledge the help and encouragement given by Dr. M. Ravindranathan and IPCL (R&D) Management. Dr. S. F. Xavier is thankful to Dr. K. V. Rao for the morphology work and Dr. V. N. S. Pendyala for his efficient handling of the script on PC.

12.11

References

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Properties and Performance of Polymer Blends

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Properties and Performance of Polymer Blends

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941

Appendices

Deca-bromodiphenyl- oxide

Tri-cresyl phosphate

Tri-butyl phosphate

Tri-butoxy ethyl phosphate

Halogenated hydrocarbons

2.

3.

4.

5.

6.

A. Organic 1. Phosphate esters

No. Name

A, E, EC, N, PVA, PVC, UF.

A, CA, CAB.

CA, CAB, CN, EC, PVC.

A, CN, E, EC, PET, PS, PVA, PVC.

ABS, A, E, P, PET, PE, PP, PS, PVA, UF, UR, PBT, N, (EPDM).

A, CA, CAB, CN, E, EC, N, P, PE, PP, PS, PVA, PVC, UF, UR.

Recommended for

Monsanto Ameribrom Elf Atochem Ethyl Great Lakes Chem. Akzo, Harwick Albright & Wilson FMC Miles Akzo, Harwick Albright & Wilson FMC Focus Chem. Akzo, Harwick. Albright & Wilson FMC Focus Chem. Argus Elf Atochem Dover Ferro Harwick Occidental

Albright & Wilson FMC

Akzo, Harwick

Supplier

Lindol, Phosflex TPP 362, 370, 387, 390, 710. Pliabrac TCP, TXP, 521, 519, 524. Kronitex 50, 100, 200, 3600, 100B, 20OB; TCP, TXP25. Santicizer 141, 143, 148, 154, TPPa. FR-1210 Thermoguard 505 Saytex 102E Great Lakes DE - 83R. Lindol, Lindol XP Plus Pilabrac TCP Kronitex TCP Disflamoll TKP. Phosflex 4 Albrite TBPO TBP TEP Phosflex TBEP Amgard TBEP KP-140 TBEP Flexchlor, CPP, FLX. Electrofine S-70 Chlorez, Paroil Kloro series Plastichlor, CPW100 Chlorowax.

Trade name

Appendix 12.1. Principal flame retardants, their trade names and suppliers [Modern Plastics, 1993]

12.12 942 S. F. Xavier

PVC

Trioctyl phosphate

Tri-phenyl phosphate

Halogenated organic phosphate

7.

8.

9.

ABS, A, CA, E, EC, N, PET, PA, PE, PP, PS, PVA, PVC, UR, UF, (PBT, EVA, TPR)

PE, PP, UF, UR, PBT, EVA, TPU, PET

10. Halogenated organics

11. Non- halogenated organics

A, CA, CAB, CN.

CA, CAB, CN, PVA, PVC.

Recommended for

No. Name

Appendix 12.1. continued

Dover Ferro Harwick Occidental Quantum, USI. Stanchem. Great Lakes Chem. Hoechst Celanese Monsanto 3-V Chem.

Elf Atochem

Great Lakes Chem. Argus

Albright & Wilson

Akzo FMC Miles Akzo

Albright & Wilson FMC Miles Rhone-Poulenc

Supplier

Phosflex TPP Kronitex TPP Disflamoli TP Fyrol CEF, DMMP, EFF 38, 25, 6, FR2, PCF 99, 51. Antiblaze 78, 80, 100, 125, 150, 175, 190. Firemaster 836, 642, HP 36 Flexchlor CPF, FLX, Fyarestor 100, 102, 205, 104 Electrofine S-70 & Thermoguard XS 70 T Chlorez & Paroil Kloro-check series Plastichlor, CPW 100 Chlorowax & Dechlorane Plus. Spectratech Cereclor 42, 545, 562, 70. NH-1511, CN-1197. Exolit IFR Spin-Flam MF-82 Plastisan B.

Amgard TOF TOF Disflamoli TOF TOF

Trade name

Properties and Performance of Polymer Blends

943

Recommended for ABS, A, CA, E, EC, N, PET, PS, PVA, PVC, UF, UR.

ABS, A, CA, E, EC, N, PET, PS, PVA, PVC, UF, UR (Polybutylene)

ABS, E, PA, PC, PO, PS (PBT, PET)

No. Name

12. Chlorinated paraffin

13. Chlorinated hydrocarbon

14. Brominated organic

Appendix 12.1. continued

Quantum, USI Santech

FMC Ferro Great Lakes Chem.

Ethyl

Dover

Argus Dover Ferro Harwick Occidental Stanchem Ameribrom Argus Elf Atochem

Argus Dover Ferro Harwick Occidental Quantum USI Elf Atochem

Elf Atochem

Supplier Electroline S-70, Thermoguard XS-70-T Flexchlor CPF, FLX Chlorez solids & Paroli liquids. Keil CW Series. Plastichlor Series, CPW 100 Chlorowax Spectratech. Electrofine S-70, Thermoguard XS-70-T Flexchlor, CPF, FLX Chlorez solids & Paroli liquids Kloro 6001 Plastichlor series CPW 100 Chlorowax Cereclor 42, 545, 552, 70. FR-1034, FR-1025, FR-913 Fycerestor series Thermoguard 200 series, BBH 44, Pyronil series DD 8426, DD 8133, DD 8207, DG 8410. Saytex: BN-451, BCL-462, BT-93, BT-93W, 120, 8010 Kronitex PB-460, PB-370. Bromoklor 50/70 Pyro-chek BC-52, BC-58, FF-680, PE-68, PO-64P, DP-45, FR-756, FB-72, PDBS-10 and 80. Spectratech Santechem 17-184.

Trade name

944 S. F. Xavier

2.

Antimony oxide

ABS, A, CA, CAB, CN, E, EC, N, P, PC, PET, PE, PP, PS, PVA, PVC, UF, UR.

ABS, A, E, N, P, PC, PET, PE, PP, PS, PVA, PVC, UF, UR.

ABS, A, CA, CAB, CN, E, EC, N, P, PET, PE, PP, PS, PVA, PVC, UF, UR.

15. Chlorinated organics

B. Inorganic additives 1. Alumina tri- hydrate

Recommended for

No. Name

Appendix 12.1. continued

Harwick, Laurel Ind.

Asarco Elf Atochem.

Anzon

Amspec

Georgia Marble R. J. Marshall Solem, Harwick

Alcan Alcoa AluChem. Climax Custom Grinders Franklin Industrial

Akzo Albright & Wilson Argus Elf Atochem. Ferro Dover Harwick Occidental Quantum, USI

Supplier

H, FRF, SF & UF series C-series, Hydral series AC-series Hydrax ATH series Polyfil series Custom grinds H-series Custom grinds of hydrate & carbonate. KC-series A-100 series, A-200 series SB-series surface modified aluminas. Micral series. Amstar HP, KR-High and LTS-Low Tint. TMS, Oncor 75 RA & 55, TMS-HP Trutint 50. Microfine, Trutint 80 Very high, high, and low tint, ultra pure. Thermoguard S.L., CPA, Low Dust Series, Antimony Halogen Series. Fire shield H-L & Ultrafine grades; LSFR, Pentoxide TP2, TPL.

Fyrol PCF, CEF, Fr-2, 38. Antiblaze 80, 100, 125, 150, 175, 195. Flexchlor, CPF, FLX Electrofine S-70, Thermoguard XS-70-T Kloro 3000, 3100. Chlorez solid & Paroil liquids. Plastichlor series, CPW 100. Chlorwax, Dechlorane Plus. Spectratech.

Trade name

Properties and Performance of Polymer Blends

945

ABS, A, CA, CAB, CN, E, EC, N, P, PC, PET, PE, PP, PS, PVA, PVC, UF, UR

Antimony oxide dispersions

Magnesium hydroxide

Organic- inorganic additive

Sodium antimonate

Tin compounds

Zinc Borate

3.

4.

5.

6.

7.

8.

ABS, El N, P, PC, PA, PET, PE, PP, PS, PVA, PVC, UFF, UFR ABS, A, E, N, P, PC, PET, PE, PP, PS, PVA, PVC, UFF, UFR

ABS, A, CA, CAB, CN, E, R, EC, N, P, PC, PA, PE, PP, PS, PVA, PVC, UFF, UFR. (EVA, PBT, SAN, TPE) ABS, A, CN, E, N, PET, PE, PP, PS, PVC, UFF, UFR.

ABS, A, E, P, PE, PP, PS, PVC, UF, UR (TPE, EPR, EVA.)

Recommended for

No. Name

Appendix 12.1. continued

Climax Harwick U.S - Borax

Amspec Anzon Elf Atochem Holtrachem Alcan

D. J. Enterprises Solem Anzon D. J. Enterprises

Aluchem Climax

ZB-112, 325, 237, 233, 467 Firebrake ZB, Zb-Fine, Fire brake 500, Fire brake 415.

S. A. 100 Pyrobloc SAP Thermoguard FR Montana Flamtard H.S.

Regular, Red & White, Treated. Spectratech Montana high tint, low tint, Micropure treated blends. Amsperse Environstrand, Fyrebloc Thermoguard S, L, NF. Montana DIDP treated, high, low, tints, Micro pure. ACM-MW, ACM-MH 93. Hydramax HM-B8, HM-B8S, HM-C9, HM-C9S. ACM-MW, ACM-MH93 Zerogen. F.R.C. Enviro- strands, Fyrebloc Micro P, Sillum 200, PL-200, Q/P

Miljac Quantum, USI Holtrachem Amspec Anzon Elf Atochem Holtrachem.

Trade name

Supplier

946 S. F. Xavier

Ethylene oxide adduct of TBBPA Phosphorus containing polyol. Tetra-chloro- phthalic anhydride Hexa-chloro cyclopentadiene

Ameribrom Ameribrom Ethyl Ameribrom Ethyl Great Lakes Chem. Ameribrom

ABS, PE, PP, PS. A, CA, E, PET, UFF, UFR.

CA, CAB, CN, EC.

E, P, PET, UFR.

E, P, UFF, UFR.

ABS, E, PET, UFF, UFR.

Velsicol

Monsanto

Albright & Wilson

Akzo

CN, P, PET, UFF, UFR.

ABS, E, P, PC, PET, UFF, UFR.

Ameribrom

Supplier

A, CA, CN, PET, UFF, UFR.

Recommended for

PCL

Tetrathal

Vircol 82

FR-522 FR-1138 1524 Saytex RB-100 Great Lakes BA-59P FR-1525

FR-1033

Victastab HMP

FR-513.

Trade name

Note: A = acrylic; CA = cellulose acetate; CAB = cellulose acetate butyrate; CN = cellulose nitrate; E = epoxy; EC = ethyl cellulose; N = nitrile; P = phenolic; PC = polycarbonate; PA = polyamide; PE = polyethylene; PET = polyester, thermoset; PP = polypropylene; PS = polystyrene; PVA = polyvinyl acetate; PVC = polyvinyl chloride; UFF = urethane flexible foam; UFR = urethane rigid foam.

9.

8.

7.

6.

Reactive types 1. Tri-bromo- neopentyl alcohol 2. Di-(polyoxy- ethylene) hydromethyl phosphonate 3. Tri-bromo- phenyl maleimide. 4. Di-bromo- neopentyl glycol 5. Tetra-bromo bisphenol A

No. Name

Appendix 12.1. continued

Properties and Performance of Polymer Blends

947

948

S. F. Xavier

Appendix 12.2. Definition of terms used in this chapter Autoignition Autoignition temperature

Burning behavior Char Chemical resistance

Combustion

Decomposition temperature Dielectric breakdown voltage

Dielectric constant Dielectric Strength Dissipation factor Elongation Elongation at break Elongation at yield Fatigue

Fatigue crack growth rate

Fire Fire resistance

Ignition caused solely by heat without application of a flame. (Also called self-ignition.) The minimum temperature to which a substance must be heated, without application of a flame, in order to cause that substance to ignite. The physical and chemical changes that take place when materials, products, or structures burn and/or exposed to fire. Carbonaceous material formed by pyrolysis or incomplete combustion [ASTM E 176-81a]. The ability of a material to resist chemical attack (the attack is dependent on the method of test and its severity and is measured by determining the changes in physical properties [ASTM, 1982]. Reaction of a substance with oxygen with release of heat, generally accompanied by flaming and/or emission of smoke. Any chemical process that produces light and heat, either as glow or flame. This is the temperature range associated with the decomposition of the polymer in the presence of oxygen. The voltage at which electrical failure or breakdown of a dielectric occurs when energized between two electrodes under prescribed test conditions. For a given configuration of electrodes, the ratio of the capacitance with the material as the dielectric to the capacitance with vacuum. The average voltage gradient at which electric breakdown occurs under specific conditions of test. The ratio of the loss index to its relative permittivity, the tangent of its loss angle, δ or the cotangent of its phase angle, Θ. The strain produced in the test specimen by a tensile stress, expressed as a percentage with respect to the gauge length. The percentage elongation produced in the gauge length of the test specimen at the break point. The percentage elongation produced in the gauge length of the test specimen at the yield stress. The process of progressive localized permanent structural change occurring in a material subjected to conditions that produce fluctuating stresses and strains at some point or points that may culminate in cracks, complete fracture or thermal softening after sufficient number of fluctuations [Borders et al., 1946] Crack extension caused by constant amplitude fatigue loading and expressed in terms of crack extension per cycle of fatigue, da/dN [Borders et al., 1946]. A process of combustion characterized by the emission of heat, accompanied by smoke and/or flame [Mark et al., 1975]. The property of a material or an assembly to withstand fire or give protection from it.

Properties and Performance of Polymer Blends

949

Appendix 12.2. continued Fire retardant (flame retardant) The quality of a substance of suppressing, reducing or delaying markedly the combustion of certain materials. A fire retardant causes a material to resist burning when exposed to a high-energy source [Sanders, 1978]. Flexural deflection The distance over which the top or the bottom surface of the test piece at mid-span has deviated during flexure from its original position. Flexural stress The flexural stress at a deflection equal to 1.5 times the thickness (at conventional deflection) of the test piece. Flexural stress at maximum load The flexural stress developed when the load reaches the first maximum. Flexural Stress at rupture The flexural stress developed at the moment of rupture. Fracture A break in the mechanical continuity of a material caused by stress exceeding the strength of the material, including joints and faults. Fracture toughness A conventional fracture mechanics strength parameter indicating the resistance of a material to crack extension. Gauge Length The original length between two marks on the test piece over which the change in length is determined. Haze Percentage of transmitted light that passes through the specimen and deviates from the incident beam by forward scattering [ASTM D 1003]. Haze, reflection The scattering of reflected light in directions near that of specular reflection by a specimen having a glossy surface. Haze, transmission The scattering of light within a specimen or at its surface responsible for the cloudy appearance of objects observed through the specimen. Homogeneous specimen A specimen in which every geometrically identical portion has the same apparent thermal conductivity. Impact strength The property to resist physical breakdown when subjected to a rapidly increasing applied force [ASTM, Standard Definitions, 1982]. Layered specimen A specimen that, if sliced parallel to the faces, has one or more slices with a significantly different apparent thermal conductivity than the other slices. Limiting Oxygen Index (LOI) This is a measure of the minimum concentration of oxygen in an oxygen-nitrogen atmosphere that is necessary to support a flame for at least 3 minutes under specified test conditions [Mark et al., 1975]. Refraction Change in the direction of propagation of radiation determined by change in the velocity of propagation in passing from one medium to another. Refractive Index The ratio of the velocity of light (of specified wavelength) in air, to its velocity in the substance under examination. Rockwell Hardness A number derived from the net increase in the depth of impression as the load on a penetrator is increased from a fixed minor load to a major load and then returned to the minor load [ASTM D 785]. Secant modulus The ratio of stress to strain, in general, at any given point on the stress-strain curve. Smoldering Combustion of a solid without flame. The combustion of a material without light being visible and generally evidenced by smoke and an increase in temperature [ASTM E 176-81a].

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S. F. Xavier

Appendix 12.2. continued Specific heat Strain Surface resistance (Rs)

Surface resistivity Tensile modulus (also elastic modulus in tension or Young’s modulus)

Tensile strength (nominal) Tensile stress (nominal)

Tensile stress at break Tensile stress at yield

Thermal conductance, (Γ) Thermal Conductivity, (λ) Thermal resistance, R Thermal resistivity (r) Toughness Toxicity

Transmittance (of light) Volume resistance, (Rv )

Volume resistivity

The heat capacity, C, per unit mass or per unit volume; usually the term refers to mass specific heat [Brown, 1981]. The change in length per unit original length of the measured gauge length of the test specimen. It is expressed as a dimensionless ratio. Resistance between two electrodes that are on the surface of a specimen is the ratio of the direct voltage applied to the electrodes to that portion of the current between them which is primarily in a thin layer of moisture or other semiconducting material that may be deposited on the surface. The ratio of the potential gradient parallel to the current along its surface to the current per unit width of the surface. The ratio of tensile stress to corresponding strain below the proportional limit. Many polymers/blends do not obey Hooke’s law through out the elastic range but deviate therefrom even at stresses well below the yield stress. However, stress-strain curves almost always show a linear region at low stresses, and a straight line drawn tangent to this portion of the curve permits calculation of tensile modulus. The maximum tensile stress (nominal) sustained by a test piece during a tension test. The tensile force per unit area of the original cross-section within the gauge length carried by the test piece at any given moment. The standard unit is mega-Pascal (MPa = MN/m) The tensile stress at which break of the test specimen occurs. The tensile stress at which the first marked inflection of the stress-strain curve occurs. Where any increase in strain occurs without any increase in stress, this point is taken as the tensile stress at yield or yield stress The reciprocal of thermal resistance. The heat flux per unit temperature gradient in the direction perpendicular to an isothermal surface, under steady-state conditions. The temperature difference required to produce a unit of heat flux through the specimens under steady state conditions [ASTM C177]. The reciprocal of the thermal conductivity. That property of a material by virtue of which it can absorb work. The amount of a substance that produces detrimental effects in an animal. It is expressed as a dose divided by the body weight of the animal, i.e., in mg/kg [Chamberlain, 1978]. That fraction of the incident light of a given wavelength which is not reflected or absorbed, but passes through a substance. The volume resistance between two electrodes that are in contact with, or embedded in a specimen, is the ratio of the direct voltage applied to the electrodes to that portion of the current between them that is distributed through the volume of the specimen [ASTM, Standard Definitions, 1982]. The volume resistance (in ohm-cm) between opposite faces of one centimeter cube of the material.

CHAPTER 13

APPLICATIONS OF POLYMER ALLOYS AND BLENDS

J.J. Scobbo, Jr.1 and Lloyd A. Goettler2

1

General Electric Plastics, Mt. Vernon, IN, USA

2

University of Akron, Department of Polymer Engineering, Akron, OH, USA

13.1

General Principles for the Use of Polymer Blends

Since their widespread commercialization, the applications of polymer blends have been directed at replacement of traditional materials, most commonly, metals. Although plastic raw materials can be more costly than metals on a weight basis, they are often more economical in terms of final manufactured cost [Paul and Newman, 1978; Legge et al., 1987; Utracki, 1989, 1998; Datta and Lohse, 1996; Louise, 1997]. This is because plastic parts can consolidate many functions into fewer parts, usually require less complex assembly (e.g., they are amenable to snap fitting and ultrasonic welding) and can be easily formed (by injection molding) into complex finished shapes, even incorporating textured or high gloss surfaces. In use they are more corrosion resistant as well as lighter in weight than metals, which is especially important for fuel economy in automotive applications. There is a trend toward specialization in the polymer products industry. Since the industry is expanding globally, a sufficient market is available for these products. Blending is a convenient route to time-efficient and cost-effective upgrading of commodity resins and to tailoring these resins to specific performance profiles for the desired application. The time to commercialization can now be reduced to less than one year for PAB’s vs. 8-10 or more years for the synthesis of new polymers. The development of the latter can exceed $10 MM. After development, modifications can also be more easily implemented. For example, flame retardant (FR) grades of PAB’s have been developed for the business machine and electronics markets. These materials combine high modulus, heat resistance, and impact strength in addition to FR. Costs can be reduced by blending engineering thermoplastics (ETPs) with less expensive commodity resins. Also, the distinction between plastics and elastomers can be breached by PAB’s and is in fact narrowing. Advantages of PAB’s blends for meeting application requirements can be summarized as follows: • lower development costs • faster development • recycling issues (post-consumer) • regulatory considerations L.A. Utracki (Ed.), Polymer Blends Handbook, 951-976. © 2003 Kluwer Academic Publishers. Printed in the Netherlands.

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• reduced manufacturing costs, e.g., lower capital required (vs. metal conversion) • JIT (just-in-time) inventory • part size reduction, thinner walls • parts consolidation • tailored properties to meet performance specifications. Furthermore, the beneficial PAB properties for general applications include: • better processing: e.g., PA/ABS vs. PA for blow molding • enhanced performance over single resin systems • mechanical (creep, impact, stiffness, strength) • heat resistance • lower cost • optical • electrical properties • flammability • better flow for thin pathways in intricate molds (especially by blending with LCP resins)

• • • • • •

chemical resistance barrier (permeation) weathering surface characteristics sound deadening (acoustics) dimensional stability.

This chapter on applications of PAB’s focuses on polymer systems giving synergistic and generally high performance properties. Low performance PAB’s of commodity plastics, rubber toughened plastics, copolymers, and interpenetrating networks are excluded. Some of the more common PAB’s are described elsewhere in this book. Markets for engineering alloys and blends have been growing steadily and currently the total market is in excess of 400 kton. Growth rates are between 3 and 4% overall. Table 13.1 and Figure 13.1 summarize the approximate market penetration for several classes of blends and alloys considered in this chapter. Volume

Polycarbonate 36%

PPE 39%

Polyesters 11%

Polyamides 14%

Figure 13.1. Blend and alloy market volumes.

Table 13.1. Blend and alloy market volume Basic Resin PPE Polyamides Polyesters Polycarbonate

Volume, kton/y 100 36 27 91

Applications of Polymer Alloys and Blends

Blend families, i.e., blends of polyphenylene ether (PPE), blends of polyesters (PEST), blends of polycarbonate of bisphenol-A (PC), tend to be appropriate for certain types of applications which require the special attributes of the major blend component. For this reason, it is possible to group blend families together under the major resin component and to list the types of applications that have been commercialized. PPE: with HIPS, PS, PA, PTFE PPE/HIPS (high impact polystyrene) or PPE/PS (polystyrene): good dimensional stability, high resistance to moisture, impact resistance, low temperature impact strength, low creep, improved processing. Applications are in automotive instrument panels, interior trim, business equipment chassis, appliances, and electrical applications including those requiring flame retardancy, appliances, building and construction, medical. PPE/PA (polyamide): good chemical resistance, high deflection temperature under load (DTUL), high continuous use temperature, ease of processing, high impact, high modulus. Applications are in injection molded automotive body panels, thermoformed truck body panels, mirror housings, trim components, wheel covers, under the hood electrical connectors, fluid handling, resonators, conductive grades for electrostatic painting. Polyester (PBT and PET) with elastomers, polysulfones PET/PBT (polyethylene terephthalate/polybutylene terephthalate): good colorability, excellent surface aesthetics, gloss, chemical resistance, impact, and electrical properties. Applications include appliances, electrical applications, building and construction. PET/PSF (polysulfone; usually fiber reinforced): warp resistance, dimensional stability, stiffness, high temperature performance, chemical resistance. Applications include industrial process equipment, electrical connectors, and food processing equipment. PBT/elastomers: chemical resistance, notch Izod impact strength, heat resistance (for paint ovens). PET/elastomers: stiffness, impact strength, and

953

processability. Applications are in automotive (body parts, steering wheels, under-the-hood components). PC with ABS, PBT, PET, SMA, ASA, TPU PC/ABS: good heat resistance, toughness, tensile strength, creep resistance, color stability, flame resistance, moldability, surface appearance. Used in business machine housings, exterior automotive trim, wheel covers, IP’s, interior parts. PC/ABS replaces over-engineered straight PC in applications where a higher heat or impact resistance than straight ABS is required. Eliminating painting, e.g., with a low gloss PC/ABS, provides cost reduction, elimination of volatile organics and better recyclability. PC/ASA (acrylic rubber styrene acrylonitrile) is essentially a more weatherable offset to PC/ABS. PC/PBT: low temperature impact, dimensional stability, and chemical resistance. Applications in automotive bumpers. PC/PET: transparent, low temperature performance, chemical resistance, abrasion resistance, impact resistance. Applications are in automotive (bumpers and exterior trim), furniture (chair arm rests, support members), appliances (housings), hand tools (housings), lawn and garden (motor housings, guards for electric hedge trimmers, protective housings), industrial (fluid handling equipment like pump housings, valve bodies and handles), electrical/electronic (coil bobbins and connectors). PC/SMA: heat resistance, high stiffness, and impact resistance. Used in automotive, electrical/ electronic, industrial and consumer (lawn and garden). PC/TPU: exterior motor vehicle for chemical resistance and strength. PC/Acrylic: improved toughness/modulus balance over PC with good toughness retention at low temperature in thick parts. These alloys also process better than PC and toughened PC/PET blends. Applications are in the following markets: Sporting Equipment (helmets, protective gear), Telecommunications (housings for portable equipment), Computer and Business machine (housings) and Industrial (casings and housings for power tools).

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PA alloy with ABS, PE, PTFE, PPE, PP, ethylene copolymers (EPR, EPDM, EVAc, ethylene acrylic acetate, ethylene butyl acrylate) PA/ABS: high temperature warp resistance, chemical resistance, good processability, and superior appearance. automotive (body panels, under-the-hood connectors and components, fuel tanks, bearings, bushings, cams). PA/PE: chemical resistance to gasoline; used in fuel tanks. PA/PTFE: low coefficient of friction and high temperature resistance (used with glass reinforcement). Used in mechanical parts like bearings, bushings, thrust washers, also construction hardware like door latches. PPS alloys with PTFE, PEI, PPE PPS/PTFE (polyphenylene sulfide/polytetrafluoroethylene): high durability and wear resistance, tensile strength, hydraulic stability, heat resistance. Applications include seals valves and bushings. PPS/PEI (polyether imide) and PPS/PPE: high heat, high flow, low flash during injection molding, Applications are in electrical and electronic connectors. Miscellaneous: • PSF/ABS: toughness, heat resistance, chemical resistance. • TPU/ABS: low temperature toughness, chemical resistance, abrasion resistance, and superior toughness POM (acetal): used in O-rings • POM/elastomer: toughness, notched Izod resistance, and chemical resistance. Used in interior automotive, gears, cams, etc. • POM/PTFE: lubricity; used in bearings • Fluoroplastics: are used for o-rings • Polyolefin: transportation, industrial, electrical and electronic, construction; generally for low cost applications; can be easily tailored for specific applications. • PE/PIB: used in sheet and film for barrier properties. • PE/ionomer: impact resistance, stiffness, heat resistance, easy processing. • PP/EPDM, PE/EPDM: low temperature perfor-

• •

•

• •

•

mance and toughness; rubber-like; used in wire and cable insulation, hose and tubing, interior automotive trim and mechanical components. PVC: with CPE or EVAc for construction (siding, pipe, conduit). PVC/CPE or PVC/EVAc: impact resistance and UV stability are used in the construction industry as siding, pipes, conduit, foamed material and rigid sheets. PVC/NBR (nitrile rubber): oil and other organic fluid resistance; used in hoses, belts, wire and cable insulation, appliance parts. PVC/TPU: thermal stability, chemical resistance, low temperature flexibility for footwear. PVC/PMMA: sometimes include small amount of MMA copolymer to improve melt strength and aid processing. Used in business machine housings, furniture, mass transit seating and food service trays. MBS or MABS (methacrylate acrylonitrile butadiene styrene) added to improve low temperature toughness. Improved grades used in siding, pipe and fittings. SAN/EPDM: weatherability, impact strength, chemical resistance, and good processing. Used in construction products and sporting goods.

13.2

Automotive Applications

Blending and alloying technology has enhanced the applicability of plastics in automotive applications that had previously been the domain of homopolymers and metals. The resulting new materials frequently boast synergistically enhanced strength, low temperature impact resistance, high temperature capability and good paintability. These are all important, if not essential, performance attributes for automotive applications. The last two also ensure compliance with current manufacturing practices that often require exterior plastic parts to be painted in assemblies with steel parts and dried in high temperature paint ovens. Current trends in the automotive industry include: • Globalization of automotive platforms, and therefore of materials sourcing • Increased safety, • Increased economy/fuel efficiency,

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Applications of Polymer Alloys and Blends

• Increased comfort and ergonomics, and • Increased quality, driven by global competition. This is the largest and fastest growing market for plastic alloys (about 2/3 of total market). General application areas are exterior parts (body panels, seals, weather-stripping, bumpers, air dams), interior parts (IP’s, steering wheels, seat belt parts, decorative pieces) and under-the-hood components. Plastics account for 11-13% of a car’s weight, or 500 kg in 1990. Usage is growing 8% per year and could reach 20% by year 2000. The benefits of polymer alloys include: good strength, weight savings, low temperature resistance, paintability, heat resistance and low cost. Factors having an impact on choice of materials for automotive applications include: • safety • environmental • fuel efficiency • quality functions

• • • • •

cost short design cycle design for disassembly/recycling production efficiency parts consolidation.

Major resin components in blends for automotive applications are shown in the Table 13.2: PPE is top performer due to a large number of resins it can be alloyed with: PS, PA, PTFE. They all exhibit chemical resistance, corrosion resistance and paintability, which are useful for mirror housings, body panels and wheel covers. The motor vehicle market is one of the fastest growing for PAB’s due in part to their recent introduction into exterior applications and growing use in underhood applications. Other typical applications are given in Table 13.3. Details of the automotive applications can be found in numerous references [Claar, 1992; Miller, 1995; Eller, 1995; and Whalen, 1988].

Table 13.2. Major resin components in PAB’s for automotive applications Blending resin

Performance advantages

Typical applications

PPE

chemical corrosion paintable chemical resistance impact strength

mirror housings body panels wheel covers exterior trim “bumpers “fuel system” components, bearings, bushings, cams

Polyester Polycarbonate Polyamides

Table13.3. Automotive applications for polymer alloys Under the hood

Exterior

Interior

Connectors Resonators Ignition components Fluid handling Power distribution Sensors

bumpers wheel covers vertical body panels mirror housings cowl vent grilles fascias, spoilers

trim instrument panels door assemblies

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13.2.1

J.J. Scobbo and L.A. Goettler

Under the Hood Applications

The engine compartment of an automobile is one of the most demanding environments for plastics. The requirements include an ability to withstand extremes in heat, corrosive fluids, vibration, and mechanical loads. These must be balanced, however, with an ever present desire for low weight and low cost. The following summarizes some of the critical requirements for under-hood materials: • high temperature resistance • chemical resistance • good processability • low specific gravity • good dimensional stability • high dielectric strength • good heat aging • high modulus at elevated temperatures. One of the growing areas for PAB blend applications under the hood is in electrical and electronic connectors and systems for power distribution and control. Such applications often require low melt viscosity at injection molding temperatures to fill the intricate tooling that may be designed for up to hundreds of individual electrical connections. Practical toughness is required for snap-fit of the connections. Furthermore, glassreinforcement of these PAB’s may be required to improve modulus and dimensional stability. Typical blends suitable for these applications include PBT/PC, PA/PPE and PPS/PEI. Various sensors for temperatures, air flow, etc., all require exceptional dimensional stability, property retention at elevated temperatures, fluid resistance and creep resistance. Ignition components require many similar properties but have the additional need for good electrical properties, i.e., high dielectric strength, and good adhesion to epoxy potting compounds. Materials that see these applications include PBT/PC and PPE/HIPS. The fluids in an under-hood environment are typically quite corrosive, but that does not preclude the use of PAB’s in the these applications. Pump impellers and components in close proximity to or in direct contact with fuel can be fabricated out of PPE/PA and alloys containing PET/PC

or PBT/PC. For the most severe environments, PPS/PPE blends recently introduced in Japan have utility. The following fluids are common in underhood environments and material compatibility with each must be understood: • antifreeze • brake fluid • gasoline • methanol-based fuel • motor oil • power steering fluid • transmission fluid • windshield washer fluid. Resonators and air ducts can be manufactured by blow molding, and therefore require excellent melt strength, good surface appearance, low weight and excellent noise reduction (damping). For these applications, unfilled alloys such as PPE/PA are often used. Table 13.4 summarizes the property profile of typical PBT/PC and PPE/ PA blends used in under-hood applications. For comparison, these data are normalized to the performance of PBT/PC in Figure 13.2. The vibration damping capability of such systems as PA modified with an elastomer, modified polyolefin, or modified PS, are useful for the engine covers. Tackifiers such as cumarone resin, terpene resin or petroleum resin could be included in the blend formulation along with glass fibers. TPO elastomeric blends (blends of polypropylene, PP; with ethylene propylene rubber, EPR; or ethylene propylene diene monomer based rubber, EPDM) are finding growing uses in automotive applications, such as: bumper guards, air dams, wheel well liners, step pads, bumper fascia, rocker-panel and side cladding, airbag covers, skins for IP’s, door panels and consoles, softtouch parts, weather-stripping and under-the-hood blow-molded air intake components, boots and bellows. Advantages are in low cost, low density, wide modulus range (70 to 2,000 MPa), low temperature impact strength, flow, durability and recyclability. They are replacing some PC/PBT blends in these markets. As one example, a clean-air duct resonator made off a PP/EPDM elastomeric alloy.

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Applications of Polymer Alloys and Blends

Table 13.4. Property summary for typical 30% glass filled resins used in under-hood applications Property

Units

PBT/PC

PPE/PA

Flexural modulus Heat deflection temp. Tensile strength Notched Izod impact Density

GPa (ksi) °C (°F) MPa (psi) J/m (ft-lb/in) kg/m3

7.6 (1100) 207 (405) 119 (17300) 85 (1.6)

8.6 (1250) 241 (465) 150 (22000) 107 (2.0)

CTE, x105, at T = -40 - 135°C

m/m/°C (in/in/°F)

1530 8.8 (4.9)

1330 9.5 (5.3)

Flexural modulus 1,5 1

5

CTE, x10 , -40 to 135°C

Heat deflection temp.

0,5 PBT/PC PPE/PA

0 Density

Tensile strength Notched Izod impact

Figure 13.2. Physical properties for 30% glass-filled PPE/PA resins normalized to PBT/PC.

Automotive manufacturers have recognized the inherent benefits of PAB’s for exterior automotive applications. These include: • weight reduction • part consolidation/integration/reduction • reduced tooling costs • increased design flexibility • increased robustness to minor impact

• • • • • • • • • •

To deliver on the advantages indicated above, there are several important factors that must be considered for proper material selection. These can be summarized as: • finish • cost

The largest potential applications are in the areas of vertical body panels, which require “Class A” surface aesthetics. Due to the extreme requirements

13.2.2

Exterior Applications

strength stiffness impact resistance weight coefficient of linear thermal expansion (CLTE) chemical resistance crashworthiness heat resistance corrosion resistance in-line paintability (required > 200°C heat resistance, which only PA/PPE could meet, but oven temperatures now being reduced to, say 120°C).

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for creep resistance of horizontal body panels, polymer blends are currently not used for these applications. However, for vertical panels, trim pieces, wheel covers, etc., creep behavior is not as critical due to the lower unsupported weight of the parts. PC/ABS is one of the most widely used PAB’s for plastic body components because it combines low density (40% less weight than metal in the same application) with high toughness (less than 20% damage in 5 km/h impacts). Relative toughness in ball drop tests and tensile elongation are both relatively high in comparison with thermosetting sheet molding compound and RIM urethane, as seen in Table 13.5. One of the largest markets for plastic body panels is the General Motors Saturn car. Exterior applications in the Saturn include: • PC/ABS for doors • TPO for fascia • PA/PPE for front fender and rear quarter panels Typical PAB applications for the Saturn are listed in Table 13.6. The PC/ABS blend contains approximately 8% of a mineral filler to aid in dimensional stability. It is an excellent door material owing to its impressive dent resistance. For example, an impact energy of 7.2 J is required to produce a 1 mm dent as compared to only 0.9 J for steel. PPE/PA is the material of choice for fenders and quarter panels because of its excellent chemical resistance to gasoline spills and other automotive fluids found under the hood in close proximity to the front fenders. It should be pointed out that such extensive use of PAB’s for exterior panels, in particular the quarter panels, would not be possible in conventional car designs where the rear panel serves as an integral, structural member. Typical properties of PC/ABS and PPE/ PA for exterior body panel applications are given in Table 13.7. A graphical representation of these data is provided in Figure 13.3. Second generation PC/ABS has a proprietary filler system that is said to provide outstanding dimensional stability for applications in body panels. The product was designed specifically for

paint systems that are cured at unusually low temperatures (below 150°C), without an E-Coat. For the applications requiring solvent resistance, where PPE/PA is commonly used, PC/PEST alloys are expected to show better dimensional stability due to less water absorption. The limiting factor for these applications is its lower heat resistance. One of the most important factors determining the applicability of a PAB for body panel use is the material’s ability to withstand high temperature. Plastic body panels usually see one of three possible paint oven environments: • Up to approximately 200°C in “on-line painting,” where the plastic panel is assembled to the vehicle prior to the high temperature of the postcorrosion resistance coating oven (E-coat). • Approximately 140 to 180°C in “in-line painting,” where the above step is circumvented because the part is assembled to the vehicle afterwards, but the plastic part must go through subsequent ovens. • Approximately 120 to 140°C in “off-line painting,” where the plastic panels are painted independently of the vehicle and are assembled as painted parts. For the first two applications PPE/PA alloys have essentially been the exclusive materials of use. For the third application, PC/ABS and PPE/PA have been used. Blends of PPE and PA-66 are suited for exterior automotive applications, where they produce class “A” surface on body panels, because they combine the processing and solvent resistance of polyamides with the moisture resistance and dimensional stability of PPE. Therefore, they have good heat resistance, impact resistance, moldability, paint adhesion and chemical resistance. Comparing PPE/PA, PC/ABS and PEST/ABS the following observations can be made: • none is suitable for horizontal panels • CLTE is still too high - need to have one end free to move • PA/ABS and PEST/ABS have too low DTUL • PC/ABS has best low T impact • PC/ABS has poorer chemical resistance • all have sufficient impact strength

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Applications of Polymer Alloys and Blends

Table 13.5. Exterior body applications Property Ball Drop at -18°C, J Ultimate elongation, % DTUL, 0.45 MPa, °C

PC/ABS

SMC

PU

51 80 95-125

1.7 1 -

2.30 40 90

Table 13.6. PAB applications on the GM Saturn Part

Material

Front fender Rear quarter panel Side door (coupe) Rocker panel Front fascia Bumper beam Mirror shell Rear bumper cover

PA/PPE PA/PPE PC/ABS PC/ABS PP/EPDM PBT/PC PS/PPE PP/EPDM

Table 13.7. PC/ABS and PA/PPE properties for body panel applications Property

Units

PC/ABS

PA/PPE

Density Elongation at break Falling dart impact-30°C Falling dart impact 23°C DTUL, 1.8 MPa DTUL, 0.45 MPa

kg/m3 % J J °C °C

1180 90 51 51 110 126

1100 60 40 51 143 193

Density 2 1,5 DTUL, 0.45 MPa

1

Elongation at break

0,5 PC/ABS PA/PPE

0

DTUL, 1.8 MPa

Falling dart impact-30°C

Falling dart impact 23°C

Figure 13.3. Body panel material performance normalized to PC/ABS.

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• PEST/ABS has moderate chemical resistance • PPE/PA has best DTUL and chemical resistance • PA/ABS is too moisture sensitive Body panels with somewhat less demanding aesthetic requirements include truck fairings. Extremely large parts such as these are often made by thermoforming as opposed to injection molding. For this reason, good melt strength is essential. Currently PPE/PA materials are used in these applications. Bumper fascias (TPO) are increasing in surface area and becoming more integrated into the body design. Thinner bumper fascia will require plastic to be stiffer, have better melt flow characteristics and still meet -30°C impact strength requirements. PC/PBT bumpers are re-used for station wagon tail lamp housings, replacing high-heat ABS. Advantages are cost savings and reduction of landfill requirements. A rear bumper beam injection molded from PC/PBT resin for the 1997 Saturn incorporates 13 previously separate parts. It attaches under a flexible fascia and contains molded-in towers that absorb energy by crushing upon impact. The bumper beam passes the Federal Motor Vehicle Safety Standards and Insurance Institute for Highway Safety rear pole tests. The bumper system saves 1.4 kg of weight over even the lighter weight aluminum/polypropylene foam system adopted by Saturn in 1991. In addition, the bumper requires less labor for attachment to the vehicle, and it can be recycled. PU/PC blends used for soft bumpers have high impact strength over the range from -29 to 66°C, plus flexibility and paintability. A PC/ASA blend is used on exterior parts such as the cowl vent grilles, trim, front grilles, head lamp bezels, pillars, license plate brackets and window surrounds. The prime advantage of this weatherable alloy is the elimination of painting, which reduces process solvent emissions. Fuel tanks with better barrier properties than HDPE alone are fabricated using barrier resin blow molding technology. HDPE is blended with PA/PE barrier resin A urethane/ABS blend can be used in place of straight ABS in the lower portion of a two-piece

side chassis fairing to increase fairing fatigue resistance. The top part of the fairing is made of fiber reinforced polyester. Thermoplastic alloys based on styrene maleimide with PA have good thermal stability, chemical resistance, abrasion resistance, paintability and weatherability. Those blended with ABS combine thermal resistance with very good processability. The alloys find use in various automotive components such as defroster grills, door handles, and headlamp surrounds. Alloys made by combining polyarylamide with PA and polypropylene (which could be lightly glass reinforced) offer low water absorption, even at high PP levels, good impact resistance, and lower density, which reduces price per volume. Targeted applications include automotive body panels, and electrical components. Blends of PMMA (polymethyl methacrylate) and ABS that are strong, resistant to heat and weathering and have good flow properties during processing are used in automobile rear light cluster frames. PC/ABS blends find use in automotive exterior and interior parts, including bumper beams and fascias, wheel covers, cowl vent grilles, body side moldings, vertical body panels, spoilers, instrument panel components, air vents, decorative trim, glove compartment doors, support post covers and mirror housings. PC/ABS blends having good flow properties, high heat resistance, and low temperature impact resistance are being used for automotive front indicator housings. Wheel covers on front-wheel-drive cars must withstand temperatures up to 150°C. In contrast, rear-wheel-drive wheel covers only have a temperature-resistance requirement of 105°C and can utilize PC/ABS wheel covers which are less costly. Plating grade PC/ABS alloy can replace metal automotive parts due to its higher impact and heat resistance than traditional platable ABS and lower cost and weight than stainless steel and die cast zinc. Other exterior applications include mirror housings, door handles, spoilers and wheel covers. Each of these applications is painted off-line, so that heat resistance requirements allow for a relatively broad material selection, as seen in Table 13.8.

Applications of Polymer Alloys and Blends

961

Table 13.8. Common exterior automotive applications for PAB’s Part

Materials

Rear bumper cover Rear spoiler Bumpers Bumper beams Trim Mirror housing Side molding Grille Wheel covers Headlamp bezels

PC/PBT PPE/PA PBT/PC, PP/EPDM PBT/PC PC/ABS, PC/ASA, PA/PPE, PS/PPE PPE/PA, PPE/HIPS, PC/ABS, PS/PPE PC/ASA, PPE/HIPS PC/ABS, PPE/PA PPE/PA, PPE/HIPS

There has been a trend in recent years to move towards the electrostatic deposition of paints for exterior automotive applications. Reasons for this include improved paint transfer efficiency. Normally, plastic parts need to be painted with a conductive primer prior to the electrostatic painting of base and clear coats. One of the most recent developments in the area of external automotive applications is that of intrinsically conductive resins for electrostatic painting. Currently there are commercial applications for body panels that utilize PPE/PA that contains carbon black as the conductive component. In the USA, a mirror shell application uses PPE/PA where the conductivity is achieved through the incorporation of graphite nano-tubes. The benefits of conductive PAB’s for electrostatic painting can be summarized as follows. A significant reduction in base coat and clear coat usage is seen when one changes from traditional painting to electrostatic painting of ordinary plastics. However, a conductive priming layer must first be applied. When electrostatic painting of a conductive plastic is employed, excellent paint transfer efficiency is maintained, but the extra process step for conductive priming is eliminated. This means that the paint booths and labor associated with conductive priming can be turned over to base and clear coat operations. Another significant benefit of electrostatic painting of conductive

plastics is that the reduction in coatings usage dramatically reduces the emissions of volatile organic compounds (VOCs). If a painting operation is near the legal VOC limits, there is little that can be done to expand capacity. A change in technology to conductive plastics can reduce VOCs to the extent that capacity may be effectively increased, without significant expense for plant and equipment. PPE/PA blends have been developed that provide the benefits described above while also providing the flexibility of injection molding and the physical performance of an engineering resin. 13.2.3

Interior Applications

Factors affecting material selection for interior applications include: • cost • esthetics • fit • finish • acoustics • mechanical properties • dimensional stability • heat resistance • adhesive bonding • colorability • processability

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Two of the more common applications for PAB’s are instrument panels and consoles, trim, pillars, etc. Each of these categories presents special issues for the material designers: Instrument panels: • cost • safety (knee bolster, energy management) • styling • temperature resistance • parts consolidation. Consoles, panels, pillars, door trim: • cost • quality functions • heat resistance • dimensional stability. ABS/PP alloys have high acceptability in automotive applications because their parent components are already extensively used for interior applications, so they do not introduce any new components into the recycle stream. Thus, the ABS/PP alloy supply would derive partially from recycled interior scrap that has been blended and compatibilized in a reactive extrusion. Virgin materials would probably be added to upgrade performance. Pricing is critical, which may limit the options for compatibilization, which is not straightforward. One of the most highly used alloys in the interior of automobiles is PC/ABS. The move from the typical material, ABS, to PC/ABS is due to the need for both higher heat capability and higher impact resistance. There is also a change from PC traditionally used in areas that require high impact, to PC/ABS, for reasons of improved melt flow. In this case the PC/ABS alloy also provides improved chemical resistance, especially to hydrocarbons, which is of great importance in automotive environments. Furthermore, improved impact resistance is achieved in large crosssection part geometries where PC is known to be deficient. A special use of PAB’s in interior components is the instrument panel (IP). The IP is unique in the car, first because of its location, where it can be heated by the sun up to temperatures in excess

of 120°C. Secondly, it is unique because of the high load of attached components. Lastly, and perhaps most importantly, it is unique because of the extreme impact requirements. Design parameters can exceed the performance requirements of the typical materials being used, as glassreinforced SMA or glass-reinforced PPE/PS. Glass reinforcement is needed to achieve the high stiffness (elastic modulus >5.2 GPa) needed to maintain the stability of the part across the entire span of the vehicle width, typically up to 2 m. Parts of that size require strict attention to injection molding tool design, especially gating and the position of weld lines. There is a multitude of exposed parts attached to the IP retainer, which in the event of an accident, are likely to strike an occupant. Of great concern are the knee bolsters and glove box, which are likely points of occupant knee impact. For impact resistance, the performance requirement is the retention of sufficient toughness at low temperature, i.e., -30 to -40°C. At such low temperatures, the presence of PC in PC/ABS blends helps to boost impact above that for ABS alone. Another reason PC-based alloys are used in car interiors is the improved chemical resistance over the styrenics, such as ABS. In the IP system, impact resistance and crack resistance is necessary for all components. The entire assembly is tested in various simulated impact configurations at temperatures as low as -40°C. PC/ABS has significantly better performance than ABS and has been found to be more durable than PC in the environment of the car. Also critical is the ability of the IP to withstand the forces of airbag deployment without shattering or unacceptable permanent deformation or damage to other components. The trend in the industry is toward high modulus IP’s. High stiffness correlates with greater resonant frequency and superior noise, vibration and harshness (NVH) characteristics. As auto makers add more components to the IP such as music systems, phone, and navigation systems, the load bearing capability of the entire IP structure must be increased, while still maintaining toughness to withstand crash and airbag deployment impacts. In comparison to the conventional monolithic

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Applications of Polymer Alloys and Blends

materials, such as PP and ABS, in many of the IP applications the move to PC/ABS and PPE/PS increases the heat capability and impact resistance, as seen in Table 13.9 and Figure 13.4. The incorporation of polyesters into interior applications is driven largely by its compatibility with other materials in the car such as upholstery and carpeting. As recycling is becoming increasingly important, trends toward the use of compatible materials such as these will undoubtedly advance. Polyester-based alloys can be formulated to the desired property balance needed for molding

and the performance in interior parts. The lower viscosity of the crystalline component of the alloy provides for improved flow characteristics and the ability to reproduce the mold surface texture well, especially where low gloss is desirable. The replacement of ABS with PC/ABS for interior applications in recent years is partly driven by gloss reduction. Easy or high flow materials are important for such applications where the parts could be large, requiring multi-gate injection, and good appearance at weld lines is desirable so that

Table 13.9. Comparison of different IP materials Property

Units

ABS

PC/ABS

PPE/PS

Tensile strength Tensile modulus Tensile strain at break Flexural modulus Izod impact, 73°F Izod impact,-20°F HDT, 0.45 MPa HDT, 1.8 MPa Density

MPa MPa % MPa J/m J/m °C °C kg/m3

48 2210 — 2550 290 90 98 87 1020

59 2210 75 2380 638 270 127 110 1140

47 — 40 1960 400 — 121 116 1050

1,6 1,4 1,2 1

ABS PC/ABS PPE/PS

0,8 0,6 0,4 0,2

Figure 13.4. Instrument panel material comparison.

D en sit y

at br ea Fl k ex ur al m od ul us Iz od im pa ct ,7 3° Iz od F im pa ct ,-2 0° F HD T, 0. 45 M Pa Hd t, 1. 8 M Pa

m od ul us

el on ga tio n

Te ns ile

Te ns ile

Te ns ile

st re ng th

0

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J.J. Scobbo and L.A. Goettler

painting can be eliminated. When paint is to be eliminated, UV stability is necessary. For blends containing ABS, the butadiene portion of the ABS tends to degrade over time, hence the surface appearance and toughness suffer. For this reason PC/ASA blends are seeing increased usage both for interior and exterior applications. The only drawback of the alloys in these applications, where high heat or impact is not necessary, is the cost, which is typically higher than ABS or PP. Other material systems are finding increased usage. A blend of SMA and HIPS is used in automotive dashboard parts. PVC/ABS blends are used for powder slush molding of IP skins. For interior applications in public transport, another blend has gained widespread use — acrylic/PVC parts made by thermoforming. Its extruded sheets see use as benches, seat backs, armrests and window reveals for subways, railroad cars, buses, monorails and trolleys. Key properties that this alloy possesses include scratch and scuff resistance, and the ability to be cleaned easily using harsh detergents.

13.3

Lawn and garden

There are two performance classes of plastic products for the lawn and garden market. True structural materials, the first class, must boast a very high stiffness with tensile modulus in the range of at least 35 GPa. In some cases this can be as high as 140 GPa. These modulus values must be maintained over the temperature range of use (-30 to 50°C), with chemical resistance to oil, fertilizers and soil. Creep compliance under expected loading must be very low, but impact resistance would not be critical. These materials, which are used in structural members, must be easily bondable by welding or adhesion to the fascia, which is an example of the second performance class. While the modulus for this latter application need only be in the 700 to 2000 MPa range with moderate creep resistance, good impact resistance would be of prime importance down to -34°C. Also, the fascia must have good appearance (molded-in color or paintable), excellent resistance

to ultraviolet radiation and excellent chemical resistance to garden and household chemicals. Materials in the intermediate modulus range (7 to 10 GPa) are usually applied to hoods in ride-on mower or tractor applications, while the low modulus range qualifies for applications providing cosmetic cover without structural requirements, such as shrouds, fenders, and grilles. PAB’s could further be differentiated on the basis of molded-in color vs. paintability. Painted low modulus (i.e., non-reinforced) performance can be provided by PPE/PA alloys for high quality applications while painted ABS would suffice for lower end applications. Unpainted applications would require UV resistance as well as capability for molded-in color. PC/PBT blends are suitable for many of these applications, while PC/PCTG miscible blends may be required for higher end applications. High modulus PAB’s are achieved through reinforcement, usually by compounding glass fiber into any suitable blend. “Excellent” UV resistance can be interpreted to mean “as good as painted metal”, which is less than 3 delta E overall color change following 12 months of outdoor Florida exposure. Allowable loss in gloss would be on the order of a small change from approximately 80% to 50% reflectance after the same exposure. Acceptable mechanical property loss following weathering should not exceed 25%. Scratch resistance would also be important for appearance parts. It must be better than PE or PP. Heat resistance requirements differ depending on the particular application under consideration. In all cases, the modulus should not drop by more than about 25% between room temperature and the maximum rated temperature for the application. These are given in Table 13.10. Larger hoods require higher stiffness to minimize sag over larger spans. This can be accomplished either by using a higher modulus reinforced blend (modulus of approximately 7 GPa) or through part design. The requirements for shrouds are less than for hoods because they are smaller and typically not moveable. Through proper design, the material modulus requirement should be able to be kept below 3.4 GPa at room temperature, as shown above. In terms of toughness, the target for all fascia

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Applications of Polymer Alloys and Blends

Table 13.10. Modulus temperature relationship for lawn and garden applications Part

Shroud Hood

Flexural modulus at T = 23°C (MPa)

Maximum use temperature (°C)

3450 2410 to 6900

49 177

applications should be the ability to withstand an impact at least 1.5 times that of sheet molding compound (SMC) at -30°C. In one example, over 150 parts of a steel riding mower have been incorporated into two injection molded thermoplastic panels made from a PC/polyester blend. These are adhesively bonded to a lower fiber reinforced PU lower member. Although the plastic material cost is ten times that of the sheet metal it replaces, economies in assembly, painting and part inventory give the final product a 10% lower cost. The aluminum die-cast deck of a mower tractor has been replaced with PC/PBT alloy incorporating molded-in color, saving 3.6 kg of weight and 25% of cost. The same material finds application in tractor hoods. PC/ABS has also been used on the inlet of a vacuum/chipper. Important properties for these applications can be summarized as: • resists breakage • scratch resistance • UV resistance • injection moldability • good all-weather durability and impact resistance • very good chemical resistance • lighter that aluminum in many applications • easily colorable, molded-in color • dimensionally stable • resistance to chipping, denting corrosion, cracking. PC/PET and PC/PBT blends combine PC toughness with the chemical resistance of polyester to offer a desirable combination of high strength, toughness (including at low temperature), chemical resistance and excellent UV stability along with improved flame retardancy and colorability.

They find use in consumer products for items such as snow thrower engine covers and control panels, tractor shrouds and consoles, lawn mower housings and grilles, and lawn and garden tools. PC/PBT blends are replacing metal stampings in lawn mower housings. Advantages are lighter weight and elimination of painting while providing needed chemical resistance, heat resistance (near engine and exhaust), strength, stiffness and toughness as well as innovative aesthetics. A motorized tree pruner’s housing sections, the trigger, and safety button are all injection molded of PC/PET alloy for weight savings. Weight reduction is also achieved in a snow thrower by substituting high-strength plastics for metal in a number of components, including the impeller injection molded of PC/PET blend. Additionally, the use of plastics allowed for parts and reduced secondary operations, such as welding, assembly and painting. A PC/PEST blend has replaced an aluminum alloy in rotary gas powered mowers. The blend provides impact and fuel resistance, yet has a good, durable surface finish.

13.4

Electrical and Electronics Applications

Electrical applications have long been the domain of thermosets. However, thermoplastic blends and alloys possess clear advantages: • due to lower density and the ability to reduce wall thickness, weight savings of 30 to 50% • elimination of disposal logistics • reduced scrap rates • reduced secondary operations • parts consolidation

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• increase design freedom • cycle time reduction. Typically PPE/PS, PC/ABS, and PBT/PC blends are used in electrical applications such as lighting, meters, ballasts, electrical boxes, safety yokes on power lines, smoke and carbon monoxide detectors, etc. Important material properties for these applications include: • good tensile strength • creep resistance • high HDT • good volume resistivity • high continuous use temperature • high flash ignition temperature • non corrosive • halogen-free • paintable • color/UV stable • thin-wall moldable • good property retention after aging • good weatherability • chemical resistance • good impact performance at low and high temperatures. Polymer blends for packaging and materials handling of electrical components would require a degree of electrical conductivity to provide for static charge dissipation. This can be achieved either by a conductive additive, such as carbon black, carbon nano-tubes or metal fibers, or via blending with a conductive polymer like polyaniline. Blending of the latter is necessary to provide flow

characteristics necessary for typical thermoplastic fabrication processes. Flammability ratings are required for certain electrical applications, such as the internal components of electronic devices. Flame retardancy can be imparted by blending an inherently flame retardant polymer such as PVC with one providing improved processability, such as a styrenic. Primary application areas are in business machine and appliance housings. Good processability may be required in fabrication operations such as injection or blow molding, extrusion and thermoforming. High mechanical performance may not always be needed in non-load bearing situations. Examples of PAB’s for various electrical applications are shown in Table 13.11. One of the fastest growing market segments is that of electronics. The critical performance criteria for PAB’s in electronic connectors can be summarized as follows: • dimensional stability • high heat resistance • high flow • low flash • solvent resistance • IR solderable • flame retardancy • toughness, ductility, impact resistance. In recent years the industry has been driving toward surface mount technology over throughhole mount technology. This is more demanding thermally and requires materials to be formulated

Table 13.11. PAB’s for electrical applications Blend/Alloy

Performance Features

Applications

PVC alloys PMMA/ABS PVC/Styrenic

Flame retardancy; heat resistance Strength, heat resistance weather resistance, processability Flame retardancy, processability

PVC/Polyaniline

Conductivity, EMI shielding processability, coatability

Low load (replacement for PC/ABS, PPE) Auto light frames, business machine housings, appliance housings Business machine housings appliances, construction, electronics, automotive Packaging for electronic components

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Applications of Polymer Alloys and Blends

with high heat distortion temperature. Furthermore, product miniaturization is creating greater demands for reduced melt viscosity. Currently the market is penetrated mostly by glass reinforced engineering resins such as LCP, PPS and PCT. LCP’s are favored for their high flow and high temperature performance, but high cost and knit line weakness limit their use. PPS is another high flow resin but ductility and mold flashing can be problematic. Finally, PCT is a low cost, low flash material but is limited by a very narrow processing window and relatively low use temperature. One of the critical performance criteria for electronic applications is infra-red (IR) solderability. The “high heat” nature of PPS (Tm = 285°C) makes it suitable for these applications. However, PPS is a very low viscosity resin, and injection molded parts tend to flash very easily. This is a particular problem for the tight geometries and high fill pressures inherent in connector molding. Another need for products in this market is flame retardancy (UL 94 V-0). Furthermore, for mechanical/physical performance, glass reinforcement is required. Recently, blends have been commercialized comprising PPS and PEI to overcome the inherent drawbacks of PPS mentioned above. The high glass transition temperature of the PEI phase (about 220°C) provides a “meltable” reinforcement effect up to IR soldering temperatures. Also, the relatively high viscosity of the PEI phase reduces the mold flashing tendencies of the PPS. Blends of PPS with PPE have been developed in Japan for similar applications (PPE Tg = 210°C). Typical properties of 30 wt% glass fiber reinforced blends are given in Table 13.12. Automotive connectors have many of the same demands of high heat surface mount technology connectors. Examples of blends and alloys in these applications are largely based on PBT and PA-66. The high heat and high solvent resistance requirements make a semicrystalline resin a suitable choice. These systems include: • PBT with polyolefins or ASA • PA with PPE or polyolefins or styrenics including SAN.

A typical glass reinforced PA blend for these applications has the property profile as given in Table 13.13. Table 13.12. Typical properties of 30% glass reinforced PPS/PEI Property

Units

Value

Tensile strength Strain at break Flexural modulus Izod impact HDT at 1.8 MPa Density

MPa % MPa J/m °C kg/m3

130 1.3 11580 64 248 1550

Table 13.13. Typical automotive connector property profile, 15% glass reinforcement Property

Units

Value

Tensile strength Strain at break Flexural modulus Izod impact HDT at 1.8 MPa Density

MPa % MPa J/m °C kg/m3

105 3.5 4400 90 200 1170

13.5

Medical Applications

The medical market consists of several different types of applications. Table 13.14 summarizes some of the performance criteria and applications. Blends and alloys that are most appropriate for the above applications include PPE/PS and PC/ABS on the lower to mid performance end, with PEI blends at the high end. For obvious reasons, FDA compliance is critical. Most applications of blends and alloys in medical devices are best suited to the impact and toughness performance profiles of amorphous systems. Since many of the applications require clarity and transparency, PC is one of the most widely used resins.

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One of the most important performance criteria for medical applications is that of autoclaving. This is a common method to sterilize instruments and devices that uses a combination of high pressure and steam. In an aggressive cyclic environment such as this, retention of impact properties and optical properties is important. Another essential sterilization technology that is less severe than autoclaving is the use of ethylene oxide, against which blends and alloys must display chemical resistance. Gamma radiation is growing as an accepted alternative to ethylene oxide sterilization. Although it is cleaner than ethylene oxide in that it does not leave any residue, there is a tendency for resins to discolor or turn yellow. PAB’s are also used in applications employing chemical vapor sterilization. This sterilization process, common in dental applications, uses an

unsaturated vapor that will not corrode metal instruments. It is common in dental applications. Besides sterilization, chemical resistance is generally an important attribute that blends and alloys in medical applications must possess. Key chemical agents that must be evaluated include saline solution, isopropyl alcohol, ethyl alcohol, betadyne, chlorine bleach, and lipids. Some of the more common blends and alloys used in medical applications are summarized in Table 13.15. Blends of plasticized PVC with NBR have lower O2, CO2 and N2 permeability with reduced leaching of the plasticizer. Blending of expensive thermoplastic urethanes with less costly PVC resins produces a material for use in the medical field that offers good clarity and heat stability, and high strength and abrasion resistance. A 70 Shore A blend has 25%

Table 13.14. Applications vs. performance criteria for medical applications Application

Devices

Properties

Housings:

oxygenators, intravenous spikes, pregnancy kits intravenous devices, trocars dialysis tubes stopcocks, IV sets

opaque, non-load bearing, medium impact resistance transparent, gamma radiation stable, low melt viscosity transparent, lipid resistance, high loads/stress, gamma radiation stable transparent or opaque, good chemical resistance, high cycle autoclave- up to 2000 cycles

Gamma-radiation resistant devices: Lipid resistant: Autoclavable:

trays, pumps, reusable devices, fluid filtration

Table 13.15.

Common alloys used in medical devices

Alloy

Attributes

PPE/PS

good hydrolytic stability high, heat resistance, broad range of heat/viscosity available, steam, gamma, ethylene oxide sterilizable, UL94 V-0 flame rating excellent heat distortion, high temperature stability, high impact resistance, non-brominated flame retardancy available, high melt flow for thin wall applications gamma radiation stable, autoclavable, very high heat resistance >200°C, excellent lipid resistance, excellent chemical resistance

PC/ABS modified PEI

Applications of Polymer Alloys and Blends

more tensile strength, and a 14°C lower brittle temperature than straight PVC. It can be used for example as a blow molded cushion for an anesthesia face mask. Thermoplastic polyurethane blends are also used to replace PVC in the large and growing medical tubing market, where they circumvent the plasticizer contamination issues associated with PVC.

13.6

Building and Construction

For exterior building and construction applications, weatherability is the most important set of performance criteria that differentiates candidate PAB’s. The typical applications that PAB’s are specified for include exterior siding, door and window casings, and other trim components. The material systems comprise blends of styrenics such as SAN with saturated rubber such as butyl acrylate. Also used are alloys of PS and acrylate rubber. The material characteristics and design characteristics can be summarized as in Table 13.16. Roof panels are another relatively new application for polymer blends. For this purpose, panels are injection molded to resemble cedar shakes, with a single plastic panel replacing several individual shakes. One obvious benefit is the reduced labor involved in installation. Furthermore, the panels are interlocking and light in weight. Because they are plastic, they resist mildew and are suitable for all climates. PPE/PS has been

969

used in this application for several years. Specific advantages of this type of roof system include: • weatherability — resistance to mildew, fungus and moisture; meets wind-driven rain, wind uplift criteria • fire resistant — the inherent char-forming characteristics of the material minimize the spread of flames • durability — high impact, good property retention, ability to be walked on without damage • installation — can be installed over existing roofing materials, light weight interlocking panels. For interior building applications, there has recently been the market introduction of mineral filled PBT/PET blends for sinks and counter tops. This material has the look and feel of ceramic material due to its high loading of mineral filler, which is typically 38 to 63%. The loading of dense filler produces PAB compositions that have specific gravity ranging from 1.7 to 2.4. Extrusion grades exist for the production of sheets that can be laminated to substrates for counter tops. Other grades are available for injection molding. Different aesthetic treatments have been incorporated into the technology such that granite and marble looks can be achieved. FDA compliant grades are available for applications that come into food contact. Typical properties for these materials are given in Table 13.17. Figure 13.5 presents data normalized to the 38% filler content.

Table 13.16. Materials and design characteristics for exterior building and construction applications Material characteristics

Design Characteristics

Resistance to chalking Resistance to fading Color retention Molded-in color Good property retention Good practical impact at high and low temperatures Good processability and high line rates

extrusion profiles good thermal insulating properties easy to clean dent resistance warp and shrink resistance high condensation resistance factor (CRF)

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J.J. Scobbo and L.A. Goettler

Business Machines and Communications

Plastics in electronic equipment enclosures have grown to 160 kton of resins and additives. Gains of 6.4 percent/year are expected to follow through 1997 when the category reaches 215 kton. In the USA alone, plastics in electronic enclosures will gain by 5.5 percent/year to 97 kton in 1997. ABS, the former market leader, is being replaced by new ABS/PC blends. Specifically, ABS/PC today accounts for greater than 30 kton/year of the market. Major opportunities for blends include tough plastics for cases, enclosures and keyboards and conductive plastics for component shielding and packaging. Applications are in consumer items such as computers (housings), home entertainment centers (TVs, Radios, VCRs), camcorders, telephones, home security systems, microwave ovens and home office

equipment, as well as in conventional electronics including switches, sockets and connectors, circuit and outlet boxes, wire and cable jackets, battery cases and light fixtures. Two special property requirements in certain business machine and electrical/electronic applications are conductivity and flame resistance. Blends of intrinsically conductive polymers with conventional thermoplastics provide a means to process the former without sacrificing their conductivity. Uses are in electromagnetic interference (EMI) shielding applications. Blending a polyether with PS or HIPS imparts sufficient conductivity for ESD (electrostatic discharge) protection in micro-electronics packaging applications. Highly elongated domains, such as are produced in the skin areas of injection molded parts, enhance conductivity by increasing the likelihood of domain overlap and interconnectivity. Higher viscosity matrices provide for better mixing and elongation of the domains, while also promoting

Table 13.17. Typical physical properties of PBT/PET alloys for kitchen applications Property

Units

38% filler

45% filler

63% filler

Tensile strength Tensile elongation at break Flexural modulus Izod impact Density

MPa % MPa J/m kg/m3

55 2 3270 27 1700

103 3 6900 53 1880

45 1 5450 21 2400

Figure 13.5. Physical Properties of PBT/PET alloys for kitchen applications normalized to 38% filler.

Applications of Polymer Alloys and Blends

the migration of the lower viscosity polyether to the molded surfaces. Styrenic/PVC-based polymers and alloys are said to have better flame retardancy and processability characteristics than traditional PVC compounds and flame retardant ABS. Targeted end use markets are in business machine housings, appliances, general construction, electronics, and in the automotive industry. A flame resistant grade of PP/EPDM elastomeric is used for extrusion-coating electrical cables because it contains only 5.5 percent acid-gas materials, compared to PVC that contains 25 percent. Large concentrations of acid-gas can corrode the surface of sensitive electronic equipment if exposed to fire. ABS/PC blends can be made flame retardant (UL flammability rating of V-0 at 1.6 mm) for use in the business electronics market without the use of chlorine or bromine additives or loss of processing window. Deflection temperatures can be maintained at 91-102°C under 1.82 MPa load. PC/ABS blends used for the housing for palm-pad computer are weatherable and ignition-resistant. In one example application, a notebook computer requires that the selected resin offers impact toughness to endure outdoor and physical abuse that occurs in use by truck drivers, railroad conductors, and building and civil contractors. The computer’s base and inner and outer frame are molded with a black PC/ABS blend. A loading magazine for the automated computer cartridge filling of 18 track cartridges is made possible with plastic blends. The magazine itself is injection molded from a lubricated, glass, fiber reinforced ABS/PC compound. Fatigue life is greatly extended over previous materials. Housings for computer network analyzers are designed in PC/ABS rather than sheet metal mounted on a cast aluminum frame because of its ability to withstand the same abusive handling. While the metal housing brought the units weight to 22 kg, the redesigned PC/ABS housed unit is tougher than the metal-over metal frame unit, and weighs only 11 kg. A small and light ultra-compact cellular phone uses ABS/PC blend for light weight and impact resistance. Its good moldability permits fabrication

971

of a housing with a wall thickness of only 1.0 mm. The thin wall permitted reduction of molding cycle times, allowing cost reduction and increased productivity. Window lenses for testing equipment are molded from clear polysulfone and then insert-molded in an ABS/PC meter case. Both the case and lens resist fuel, oil and common solvents, and meet military shock and vibration standards. An alloy made from liquid crystal polyester and PPS can be used in electrical and electronics connectors, thin walled parts, and multi-gated parts. The two materials complement each other, as the LCP offers good reinforcing strength and allows for good flow into thin sections without flash. The PPS reduces the alloy’s cost, cuts warpage, and increases weld line strength. Also, the material can be processed in LCP molds. A PC/acrylic alloy is used in battery-pack housings and cellular phones. PC/PET and PC/PBT blends offer high strength, toughness (including at low temperature), chemical resistance, excellent UV stability along with improved flame retardancy and colorability. They find use in telecommunications equipment for marine radio housings, antenna housings and bases, and cellular telephone handsets. Blending an electrostatic discharge (ESD) polyether with PS or HIPS imparts a level of conductivity sufficient for ESD protection in micro-electronics packaging applications [Mass and Fahey, 1991]. For a given matrix thermoplastic, the efficiency of the ESD polyether increases with increasing viscosity of the matrix material. This effect is a direct result of better ESD polyether dispersion in the higher viscosity matrices. Distortion of the ESD polyether domains by the flow fields encountered during injection molding has an effect on blend conductivity. The highly elongated domains in the skin areas of molded parts enhance conductivity because they increase the potential for domain overlap and interconnectivity. In injection-molded parts, the ESD polyether is more concentrated in the interior areas of parts based on lower viscosity matrixes, while the opposite holds for parts made from higher viscosity matrixes.

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Blends of intrinsically conductive polymers (ICP’s) with conventional thermoplastics offered a practical means of processing ICP’s without sacrificing conductivity [Kulkarni, et al., 1991]. One of the major use for these blends was electromagnetic interference shielding. In the specific area of business machines, there are two categories: • document handling: copiers, fax machines, scanners, printers • computers: notebook and desktop computers, and internals such as hard drives. For document handling applications, there are several components for which PAB’s are suited. These include bases and chassis, which require high modulus; pen bodies and optical devices, which require excellent dimensional stability; moving parts that require lubricity and wear resistance; and paper paths that require electrostatic dissipation (ESD). Similar requirements occur for related components in computers. Electromagnetic Interference (EMI) shielding is required for computer enclosures and printed circuit boards. Among the more common PAB’s specified for these applications are PC/ABS and PPE/PS. For ESD/EMI applications, one of the key performance criteria is surface resistivity, in units of ohms/cm2. The suitability of materials for different applications can be based on surface resistivity as shown in Table 13.18.

There are several technologies for achieving ESD/EMI shielding performance. These are usually based on fillers and fibers that can include: • graphite and carbon fibers • aluminum flake • carbon black • graphite nano-tubes • inorganic salts. The important physical characteristics of these fillers and fibers within the PAB matrix resin are the aspect ratio, degree of dispersion, distribution, loading, and surface area per unit volume. These are all related to the ability of the additive to establish a conductive network. Other specific applications already mentioned that often require filler and fiber technology are the business machine internal components such as chassis. In these applications glass fiber, carbon fibers, mica and glass beads are all common. There are certain important features related to the reinforcement/filler and it’s ability to provide modulus enhancement within the PAB. These include aspect ratio, packing fraction, particle size, particle size distribution, and degree of dispersion. Fibrous reinforcements are expected to increase strength, stiffness, heat distortion temperature (HDT), wear resistance, impact resistance (for brittle materials, typically semi-crystalline resins) and specific gravity. Melt flow, elongation at break, and impact resistance (for ductile materials, typically amorphous) usually decrease with fiber

Table 13.18. Surface resistivities for various applications log (surface resistivity)

Application

Material characteristics

<5

EMI shielding

5-9

ESD

>7

Antistatic

highly conductive, grounds charge quickly can shield against electric fields, does not generate a charge, does not allow a charge to remain localized safely bleeds charge to ground, does not generate a charge, does not allow a charge to remain localized does not generate a charge, does not allow a charge to remain localized

Applications of Polymer Alloys and Blends

addition. However, fiber reinforced PAB’s can be tougher than their pure component resins. Nonfibrous reinforcements such as mica or glass beads increase strength, HDT, and specific gravity while decreasing impact resistance and melt flow. Combinations of fibrous and non-fibrous reinforcements are often used to overcome the inherent orientation effects when injection molding fiber-containing PAB’s. This is done for applications where dimensional consistency is critical, such as ink jet printer pen bodies. One of the most important performance criteria for these applications is mold shrinkage and it’s degree of isotropy. Furthermore, low coefficient of thermal expansion (CTE) is desirable. The principles for formulating plastics for these applications are simple. First, fillers have a much lower CTE than do plastics, so an elementary model for CTE based on volume fraction filler/ reinforcement is adequate to predict the improved dimensional stability for high loadings. Second, by judicious choice of non-fibrous to fibrous loading ratio, the degree of flow and cross-flow anisotropy can be controlled. Taking filler/reinforcement technology one step further, it is possible to capitalize on the inherent wear resistance provided by these reinforced formulations. For applications where component parts are moving, this characteristic must be designed into the formulation. Wear can be characterized by abrasion, adhesive wear and material transfer, chemical wear, and material loss by thermally induced degradation and melting. One of the most important measures for these applications is the wear rate, characterized by loss in volume of material per unit time. Another important measure for these applications is coefficient of friction that includes static coefficient of friction and dynamic coefficient of friction. Several approaches can be used to reduce the coefficient of friction. These include adding fibers such as glass and carbon, as previously mentioned, as well as fillers such as PTFE, graphite, and molybdenum disulfide, and fluids such as silicones. Silicones present special challenges because they are generally incompatible with the components of the PAB. In business machines and communications another key technological element apart from filler

973

and reinforcement technology is flame retardancy. Underwriters Laboratory (UL) has several rating criteria for flame resistance performance depending on the severity of the applications (i.e., V-0, V-1, V-2, 5VA, 5VB). In many of the business machine housings and other similar applications for PAB’s such as PC/ABS and PPE/PS, low flame out time and resistance to flaming drips is essential. One of the more recent developments in this area is related to the increasing tendency toward miniaturization. The drive from desktop personal computers to notebook computers to personal digital assistants means that housings must be lighter, yet still provide stiffness, impact resistance and resistance to elevated temperatures. Designers are therefore moving toward thin wall enclosures for computing devices and cellular phones. The PAB’s for these applications must therefore have improved melt flow. Recently PC/ABS blends have been introduced to the marketplace for these applications. They are characterized by an increase in melt flow rate (g/10 min) of from 50% to more than 100%.

13.8

Packaging

Some of the more common applications are in films, containers, closures, and shipping containers. For example, PET/EMAA (ethylene-methacrylic acid copolymer) blends and some of their Zn and Na salts have high toughness in films for permeability barrier applications against hydrocarbons and other organic solvents. The Zn-ionomers are toughest. Crystallinity also affected toughness [Subramanian, 1985; Subramanian and Mehra, 1987]. An alloy based on PVC and SMA (styrene maleic anhydride) mixed with other transparent PVC compounds produces a material that allows clear PVC to be microwaved without becoming sticky by withstanding exposure to temperatures up to 110°C for 30 min. It is thus ideal for microwaveable hot-fill food packaging. PA MXD6/PA-6 blends have comparable gasbarrier properties to PVDC-coated, oriented PA-6 films. The blends, which can be processed on the same machines used for ordinary film, can be

974

J.J. Scobbo and L.A. Goettler

used for retort pouches for food, or packaging for meat, including casings for sausages. PA-6/amorphous-PA blends provide good barrier protection with only a five layer structure. Typical EVAl structures contain seven layers. Thus, the PA blends can be run on machinery that is less costly and more readily available than EVAl packaging machinery. A container of PP/tie layer/PA blend/tie layer/heat seal provides similar protection to a structure of PP/tie layer/PA/EVAl/ PA/tie layer/heat seal. An extrusion grade of PP and PA-6 designed for the production of thin-walled packaging has high moisture and oxygen barrier properties and is suitable for contact with foodstuffs. The material, which is heat resistant to 130°C, provides high flow and can be used in injection blow molding applications and in thin wall moldings. Furthermore, it can be used to thermoform containers with much thinner walls than possible with PP. Other applications include PA/PE blend is used for permeation resistance and HDPE/EVAc copolymer - polybutylene blends, which are used in heat sealable film. Expandable beads made from blends of PS and modified PPE have thermal and mechanical properties significantly better than those of EPS. They are intended for high-performance applications in the packaging industry. Maximum use temperatures vary from 105 to 120°C. The smallest bead size, ranging from 300 to 500 mm can be used to make parts with wall thickness between 1 and 2.5 mm. Typical applications include microwave-resistant drinking cups and steamcleanable, reusable packaging.

13.9

Miscellaneous

13.9.1

Power Tools

A cordless drill housing molded from PC/ABS combines impact strength, UV resistance, processability, chemical resistance and cost effectiveness. It also provides a molded-in and chemically bonded soft polyurethane grip that transfers less vibration to the operator. The material retains 70% of its original properties following 1,000-hr accelerated weathering. 13.9.2

Applications including automotive air conditioning hoses, tubing, seals, gaskets, fasteners, ski boots, and other products that require impact resistance can utilize a flexible PA. They also exhibit paintability, colorability, and resistance to oils, chemicals, and high temperatures and can be injection molded, blow molded, or extruded on conventional thermoplastic equipment. Like any thermoplastic alloy, they can also be remolded as regrind. Mechanical performance is comparable to thermoplastic urethanes or plasticized PA-11 or -12 with a hardness from 50 to 65 Shore D and a flexural modulus in the range from 275-410 MPa. Institutional cleaning devices use molded PA/ABS housings for its superior resistance to stress cracking. 13.9.3

Recreation

PC/ABS blends can withstand temperatures as low as minus 40°C, and the abuses associated with recreational skiing. Snowmobile throttle and brake handles are fabricated from PA/ABS alloy to take advantage of the material’s gasoline resistance combined with low temperature impact strength about equal to PC. ASA is coextruded with ABS/PC blend for pleasure boat hulls. 13.9.4

Desirable blend properties are impact strength (withstand drop test at -29°C), solvent resistance and weatherability. Blends predominantly used include impact-modified PA (6 and 66), PC/PEST and PC/ABS.

Institutional/Consumer

Appliances

For these applications cycle times have been reduced by between 18 and 36% with the use PET/PC blends. The blends having a tensile strength of 56 MPa, a tensile modulus of 2.3 GPa, and a heat distortion temperature of 113°C at 1.8 MPa offer good colorability, gloss and

Applications of Polymer Alloys and Blends

surface appearance. Possible applications include appliance housings. Blends of PMMA/ABS that are strong, resistant to heat and weathering and have good flow properties during processing are used in business machine and appliance housings and furniture components. FR-PC/ABS blend is used to foam the flameretardant front panel on air conditioners. A PC/ABS blend is used to make light weight, small room-use air conditioners for the following reasons: • lighter weight. A 5,100 Btu unit measuring 350x480x280 mm deep weighs only 21 Kg and can thus be easily moved from room to room. • easier to manufacture than competing metal based models • flame retardancy. The unit’s top cover is injection molded from ABS/PC, while a retractable handle is injection molded from straight PC. 13.9.5

Furniture

Polyurethane/ABS resins combine elastomeric toughness with the scratch- and scuff-resistance, and processability of ABS. Flexural moduli ranging from 344 to 1000 MPa and the resin’s low temperature impact strength qualify this blend for applications in hard-sided luggage and furniture armrests. Furniture components represent another possible application for PET/PC blends, for the reasons cited earlier. 13.9.6

Recycling

One of the more recent applications of polymer blends and alloys is in plastics recycling. The current solid waste crisis has resulted in public demands for industrial solutions that would result in a reduction of all landfilled solid waste, while minimizing alternate negative environmental impact. It has been perceived by the public that plastic materials account for the overwhelming majority of the landfilled solid mass, when plastics account for only approximately 7 wt% of 18 vol%.

975

Of the plastics that are currently found in the solid waste stream, commodities are by far the greatest component. Seventy eight % of this waste is comprised of commodity thermoplastics while 16% is thermosets and only 6% is engineering thermoplastics. Therefore the contribution of high performance blends and alloys to the waste stream is rather negligible. There are several levels of technology for the recycling of plastics. These can be categorized as primary, secondary, ternary and quaternary with respect to the complexity of the recycling technology. In primary recycling, the plastics are reprocessed with standard processing machinery into materials for similar applications. Secondary recycling is similar to primary except that the materials are used in less demanding applications. Tertiary recycling involves recovery of chemical products, such as monomers, from the waste products. Quaternary recycling involves the recovery of energy. The major problem in the reprocessing of commingled plastics, and therefore of utmost importance for PAB’s, is the immiscibility of most polymers so that poor mechanical properties develop when they are blended together. Ways to enhance the mechanical behavior of these mixtures include adding a compatibilizer, adding a rubbery phase to impact modify the blend, or adding a reinforcement to restore mechanical properties. In the early 1970s the potential for recycling plastics as polymer blends was recognized. More recently, however, the trends have turned toward “life-cycle management”, to mitigate the effects of mixed recycle streams. In this scenario the parts made from the PAB’s are disassembled and segregated. This approach has been very successful in the recycling of automotive PBT/PC bumper beams. Currently, Europe is leading the way with this approach, which logistically and economically, seems to make the most sense today.

13.10

References

Claar, R. E., ChemTech, Sept. 1992, pp 546-9. Datta, S., and Lohse, D., “Polymeric Compatibilizers”, Hanser

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J.J. Scobbo and L.A. Goettler

Pub., Munich (1996). Eller, R., Plast. World, Oct. 1995, pp. 49-52. Kulkarni, V. G., Mathew, W. R, Campbell, J. C., Dinkins, C. J., and Durbin, P. J., SPE ANTEC Techn. Pap., 49, 663 (1991). Legge, N. R., Holden, G., and Schroeder, H. E., “Thermoplastic Elastomers”, Hanser Pub., Munich (1987). Luise, S., “Applications of High Temperature Polymers”, CRC Press, Boca Raton (1997). Mass, T. R., and Fahey, T. E., SPE ANTEC Techn. Pap., 49, 551 (1991). Miller, B., Plast. World, Oct. 1995, pp. 42-48. Paul, D. R., and Newman, S., “Polymer Blends”, Academic Press, San Diego (1978). Subramanian, P. M. and Mehra, V., Polym. Eng. Sci., 27, 663 (1987). Subramanian, P. M., Polym. Eng. Sci., 25, 483 (1985). Utracki, L. A., “Commercial Polymer Blends”, Chapman and Hall, Ltd, London (1998). Utracki, L. A., “Polymer Alloys and Blends”, Hanser Verlag, Munich (1989). Whalen, D., in AIChE Symposium Series, 84 (260), “New Polymer Technology for Auto Body Exteriors” (1988). www.allied.com www.amoco.com www.basf.com www.dow.com www.dupont.com www.eastman.com www.ge.com www.hoechst.com www.solutia.com

13.11

List of Abbreviations

CLTE

coefficient of linear thermal expansion; or CTE, coefficient of thermal expansion deflection temperature under load engineering thermoplastic electromagnetic interference electrostatic discharge flame retardant heat distortion temperature intrinsically conductive polymers instrument panel infra-red just-in-time noise, vibration and harshness polymer alloy and/or blend Underwriters Laboratories ultraviolet volatile organic compounds

DTUL ETP EMI ESD FR HDT ICP IP IR JIT NVH PAB UL UV VOC

CHAPTER 14

AGING AND DEGRADATION OF POLYMER BLENDS

J. M. G. Cowie1, I. J. McEwen1 and R. McIntyre2

1

Heriot-Watt University, Edinburgh EH14 4AS, Scotland

2

Millenium Inorganic Chemicals, Linconshire England

14.1

Abstract

The selection of polymers and polymer blends for use as specific materials requires the consideration of how these will withstand the environmental conditions to which these will be subjected. The long term stability of a polymer will depend on its aging characteristics both physical and chemical. Physical aging is important mainly when the materials have a substantial amorphous content as it is manifest in volume contraction of the solid caused by small scale relaxation processes in the glassy state. This can be followed by enthalpy and volume relaxation methods or mechanical measurements such as creep, stress-relaxation and dynamic mechanical methods. The systems have been modeled in a number of ways and those of Moynihan, Hodge, KAHR, Petrie-Marshall and Cowie-Ferguson are described. In blends, the general observation is that aging rates differ from those of the individual components. Both homogeneous and heterogeneous blend (and composites) behavior is described. The external effects of the environment on polymer blends are chemical in nature, and normally lead to degradation of the polymers. Chain scission, depolymerization and reactions on the side-chain substituents all contribute to overall deterioration of blend properties. These are described for blends containing polyvinyl chloride, polystyrene, acrylics and polyolefins mixed with a variety of other polymers. The general features of radiation damage and the detrimental effects of processing are reviewed. These degradation reactions contribute to the ‘weathering’ of blends. This is best studied by long term exposure to appropriate environmental conditions and the subsequent methods of assessing the results using ESR, FTIR, XPS and chemical methods are covered. Because many of the degradation reactions involve radicals, ESR is a particularly powerful approach and as surfaces are the first areas of attack techniques such as FTIR and XPS are being used extensively in this work. Examples of the application of these techniques to coatings are detailed.

L.A. Utracki (Ed.), Polymer Blends Handbook, 977-1021. © 2003 Kluwer Academic Publishers. Printed in the Netherlands.

978

14.2

J. M. G. Cowie, I. J. McEwen and R. McIntyre

Introduction

The long-term stability of polymeric materials is a matter of considerable importance, both to materials scientists and to engineers. Two types of aging occur that lead to changes in the properties of polymers; chemical and physical. Chemical aging normally leads to modification of the polymer chain and may involve chain scission, oxidation, dehalogenation, loss of pendant groups, hydrolysis, and crosslinking, all of which are chemical reactions. Physical aging is a manifestation of small scale relaxation processes that take place in the amorphous regions of a glassy polymer, causing volume contraction and densification of the sample. Thus while both are time and temperature dependent, the chemical aging results in an alteration of the chemistry of the polymer and usually leads to degradation of the polymer with a concurrent deterioration of the properties (e.g. coloration, embrittlement), while in the case of physical aging the polymer structure remains unchanged but the local packing of the chains alters. This leads to dimensional changes and alteration of physical properties such as brittleness, tensile strength, and the glass transition temperature, Tg. Thus the main difference between the two effects is that physical aging is a reversible process, whereas chemical aging is not.

14.3

Physical Aging

When a polymer is cooled from the melt to a temperature below its glass transition, the glass that is formed is not in equilibrium with its surroundings. Physical aging is then the slow structural reorganization of the glass, on annealing at a given temperature, Ta, as it attempts to achieve the equilibrium state. The continuous, slow, relaxation of the glass from the initial non-equilibrium state towards a final thermodynamic equilibrium state, produces time-dependent changes in the physical properties of the polymer. As the extent of physical aging increases there are corresponding decreases in the enthalpy, the specific volume and the fracture toughness, while

increases in glass transition temperature, Tg, the yield stress and tensile modulus of the material may also be observed. The volume and enthalpy relaxations that are features of this phenomenon, as are the time-dependent small strain mechanical properties, may then be used to follow the progress of the physical aging process. 14.3.1

Fundamental Principles

14.3.1.1 Enthalpy Relaxation Physical aging involves segmental relaxation processes that occur in the temperature range Tβ < T < Tg [Struik, 1978], where Tβ is the temperature of any secondary relaxation process that can occur in the glassy state. This restricted temperature range is disputed [Johari, 1982; McCrum, 1992] and it is believed that it is a phenomenon affecting all viscoelastic relaxation processes. As the annealing temperatures Ta drop further away from Tg [i.e., (Tg - Ta) increases], the aging process slows down and the time scales involved become quite long. Consequently many studies are carried out under thermally accelerated conditions and the relaxation of the enthalpy of the glass is a convenient parameter to follow when monitoring the physical aging process. Measurements are readily carried out at temperatures below Tg using differential scanning calorimetry, (DSC), if suitable precautions are taken. It is important to define a reproducible thermal history for the sample before meaningful comparisons can be made. The first step is to ensure that the influence of any previous thermal history is erased by annealing the sample at a temperature in excess of Tg + 50°C, i.e., at the point A in Figure 14.1. The polymer is then cooled from the melt at temperature T2, (A), at a rate q1, into the glass at temperature Ta, (B). The distance of the sample from its “equilibrium state” at Ta will depend on the rate of cooling of the sample and so this thermal treatment should remain unchanged for all measurements that are to be compared. Annealing the sample at temperature Ta, for a specified time ta, results in an enthalpy loss HB - HC (along the lines B-C), the extent of

Aging and Degradation of Polymer Blends

979

which depends on the magnitude of ta. When measurements are carried out using DSC, the sample is then quenched from Ta after ta, and on reheating the polymer at a rate q2, the enthalpy overshoots the equilibrium curve, as shown in Figure 14.1 by an amount proportional to that lost during the aging process. This is represented by the line (C-A). The enthalpy difference: (HB - HC) is then given by: Tβ

HB - HC = ∫ [Cp (aged) - Cp (un - aged)]dT

(14.1)

Tα

and represents the area (X-Y) in Figure 14.2. However, the methods of calculating enthalpy changes can depend on the model used to describe physical aging, as described next. 14.3.1.2 The Multiparameter Phenomenological Models A number of theoretical treatments have been developed that attempt to model the aging process in organic and inorganic glasses. Notable examples are the phenomenological models developed by Moynihan et al. [1976] and Hodge et al. [1981, 1982, 1983] and more recently Gomez et al. [1995]. As the relaxation processes in the glassy state and glass-transition region are nonexponential and non-linear, the theories must take account of the thermal history of glass formation and the asymmetry of the relaxations, which depend on how the system departs from equilibrium. Use is then made of the convenient concept of a ‘fictive’ temperature, Tf, first proposed by [Tool, 1946] defined as the temperature at which the non-equilibrium value of a macroscopic property of the system would be an equilibrium one. This may be better understood by referring to Figure 14.3 where the enthalpy difference H(T2)-H(T) can be expressed in terms of the heat capacity Cp and leads finally to: Tf

Figure 14.1. Schematic diagram of cooling, aging, and heating cycles using a fixed thermal history.

T2

∫ [Cp (T’) - Cp glass (T’)]dT - ∫ [Cp liquid (T’) - Cp (T’)]dt’ = 0 T

Tf

(14.2) where T’ is a dummy variable. From this a normalized heat capacity can be formulated in terms of experimental quantities as:

Figure 14.2. Schematic diagram of DSC curves for an aged (full line) and un-aged (broken line) polymer sample.

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J. M. G. Cowie, I. J. McEwen and R. McIntyre

Figure 14.3. is defined.

dTf dT

=

Representation of how the fictive temperature

CP(T) - CP, glass(T)

Figure 14.4. Schematic diagrams of the areas that are measured from DSC scans when estimating the enthalpy

(14.3)

CP, liquid(Tf) - CP, glass(Tf)

When the heat capacities of the glass and the liquid are defined by linear functions of temperature, the value of Tf can be calculated, from numerical integration of the DSC curves to obtain H(T2) - H(T). Gomez et al. [1995] did not use the fictive temperature, but considered instead the temporal evolution of the configurational entropy, Sc. This made it possible to examine the raw Cp data without having to resort to normalization procedures. This also removes the restriction of having to calculate the limiting Cp of the aged glassy polymer from a linear extrapolation of the liquid Cp curve. In other words, the authors assumed that the polymer may not be able to attain such a value because of the physical restraints imposed by chain entanglements and inefficient packing in glass. This concept has also been used by Cowie and Ferguson [1986, 1989]. The annealing of the glass at the aging temperature Ta results in a relaxation of the enthalpy towards the equilibrium value H(∞). This can be described by a relaxation function Φ(t), which is defined by: Φ(t) = [H(t) - H(∞)]/[H(0) - H(∞)]

(14.4)

The experimental measurement of these quantities is illustrated schematically in Figure 14.4. The areas A1, A2 and A3 can then be related to the quantities in Eq 14.4 as follows: H(t) - H(∞) = (A3 - A2)

(14.5)

H(0) - H(∞) = (A3 - A1)

(14.6)

Φ(t) can also be expressed in terms of a semiempirical function introduced originally by Kohlrausch [1847] and revived and redeveloped by Williams and Watts, [1970], abbreviated as the KWW equation: Φ(t) = exp [-(t/τ)β]

(14.7)

where β is related to the breadth of the distribution of relaxation times and has values 0 < β < 1. A value of β = 1 would imply an infinitely sharp distribution with only one relaxation time. The latter parameter is represented by τ and the effects of physical aging are then analyzed in terms of β and τ. The relaxation time τ can be expressed in terms of the fictive temperature Tf following Narayanaswamy, [1971] and Moynihan et al., [1976] who used form:

981

Aging and Degradation of Polymer Blends

x∆h (1 - x)∆h ln τ = ln A +  + RT RTf 

(14.8)

which assumes the system to be thermo-rheologically simple with only a single relaxation time. Here ∆h is the activation energy, x is a structural parameter with values 0 < x ” 1 and A is a fitting constant. In this model ∆h and x have no clear physical meaning, but Hodge, [1987] has used the Adam-Gibbs theory [Adam and Gibbs, 1965] to show that x = (Tf ’/T2) where Tf ’ is the value of the fictive temperature in the glass and T2 is the “equilibrium” Tg value postulated by Gibbs and Di Marzio, [1958]. The parameters in Eq 14.8 can be estimated by assuming ∆h has a fixed value and allowing ln A, x and β to vary until the best fit to the experimental, normalized Cp data is obtained. A more informative assessment can be made by comparing ∆H(ta, Ta) values from the model with experimental data. Alternative analytical methods are the KAHR isothermal single relaxation time model [Kovacs et al., 1979] and the peak shifting technique, [Hutchinson, 1992]. None of these attempt to use the data to predict long-term aging effects for which a different approach is necessary. 14.3.1.3 Predictive Models for Long-term Aging The descriptions of physical aging given in the previous section do not have the facility to predict long-term aging effects and for this one must use either the Petrie-Marshall (P-M) [1975] or the Cowie-Ferguson (C-F) [1986; 1989] models. In both, the enthalpy lost on aging is given by: ∆H(ta, Ta) = ∆H(∞, Ta) {1 - Φ(ta)}

(14.9)

where the P-M approach uses Eq 14.7 to define Φ(ta) with β = 1, but C-F express Φ(ta) as: Φ(ta) = exp[- (t/tc)β]

(14.10)

where tc is a characteristic time, such that ta = tc when the polymer glass has aged to 63.2% of the fully aged glass. The method of determining ∆H(∞, Ta) also differs. In the P-M model this

is estimated, as illustrated in Figure 14.4, by a linear extrapolation of the heat capacity (liquid) into the glassy state. The relaxation time is then related to the departure from equilibrium of the enthalpy (δH) by: 1 τ

=

 ∆H(ta, Ta)  ln 1 ta  ∆H(∞, Ta) 

1

(14.11)

where δH = ∆H(∞, Ta) - ∆H(ta, Ta) (14.12) The C-F approach treats ∆H(∞, Ta) as an adjustable parameter as it is considered that the linear Cp extrapolation is inaccurate. Thus a direct measure of the area (A2 - A1) is used to calculate ∆H(ta, Ta). The data are then analyzed by curve fitting plots of ∆H(ta, Ta) against log10ta, to assess the thermodynamic aspects from the ∆H(∞, Ta) parameter, and the kinetic aspects, embodied in Φ(ta), both of which are obtainable from this approach. Also considered in the C-F approach is the prediction of te , which is the time to reach 99.9% of the thermodynamic equilibrium state of the infinitely aged glass. This can be predicted from short-term experiments and is illustrated in Section 14.3.2.1. 14.3.2

Physical Aging in Polymer Blends

Remarkably few studies of physical aging in polymer blends have appeared in the open literature, considering the importance of the phenomenon. There have been many investigations of single polymer systems and this is useful as it provides a means of comparing blend behavior with that of the original components. While the data described in detail here will refer to blend systems only, the relevant homopolymer data will also be mentioned. Some of the work reported is purely qualitative and is summarized in Table 14.1 for homopolymers and Table 14.2 for blends and composite systems. When an attempt at quantitative analysis of the data has been made these references are gathered in Table 14.3 (Homopolymers) and Table 14.4 (Blends and Composites). Wholly theoretical work is listed in Table 14.5. More specific data for homopolymers are given in Tables 14.6 to 14.8, derived from the Hodge models. In Table 14.9,

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J. M. G. Cowie, I. J. McEwen and R. McIntyre

the parameters derived from the CF model are shown for a range of homopolymers, while Table

Table 14.1.

14.10 contains the parameters calculated using the Petrie model.

Physical aging in homopolymer systems: Qualitative treatment

Polymer Techniques Arom. Polyethers BPA-Epoxy resin Epoxy resins Epoxy resins Epoxy resins P(alkyl-MA)s PA-66 PC PC PC PC PC PC PC PC PC PC PC PEEK PEEK PEEK PEI PEI PEI & PEEK PET PET blends PHB PiBA PMA PMMA PMMA PMMA PMMA

References

DSC J(t) DSC E(t) Rho DSC Creep, Vol. Dil Dil DSC E’ and E” E(t) Fluorescence Impact Strength Mech, Positron Mech. NMR Rho Creep DSC G’ & G” Creep, DSC DSC DSC DSC DSC, Mech DMTA Fluorescence DETA Dil, ESR DSC DSC Fluorescence

Hatakeyama, 1990 Plazek, 1989 Jo, 1991 Jo, 1991 Jo, 1991 Yoshida,1982 Spinu, 1994 Muller, 1988 Pixa, 1986 Muller, 1988 Ricco, 1990 Yee, 1988 Torkelson, 1993 Paul, 1992 Hill, 1990 Lundberg, 1996 Smith, 1988 Pixa, 1986 Carfagna, 1988 Carfagna, 1988 Ogale, 1989 Echeverria, 1995 Hay, 1991 Hay, 1990 Montserrat, 1989 Woo, 1990 Scandola, 1989 Torkelson, 1993 Gomez-Ribelles, 1985 Simha, 1986 Diaz-Calleja, 1989 Stewart, 1987 Torkelson, 1993

Polymer

Techniques

References

PMMA PMMA PMMA PP PP PP PP PPSS PS PS PS PS PS PS PS PS PS PS PS PTMSP PTMSP PTMSP PVAc PVAc PVAc PVAc PVC PVC PVC PVC PVC PVC PVC

IR PAL, DMTA, DETA TBA DSC Mech Rho X-ray Mechanical Creep Dil Dil Dil DSC DSC DSC Fluorescence Fluorescence G(T) Modulated DSC DMTA gas-permeability Ultrasonic Fluorescence Fluorescence Pos. Anhil. TSD Creep Dil DSC E(t) ε(t), Creep, Strain G(T) Rho

Joss, 1985 Davis, 1998 Gillham, 1992 Vittoria, 1989 Min, 1989 Vittoria, 1989 Vittoria, 1989 Kobayashi, 1993 Arnold, 1993 Goldbach, 1966 Goldbach, 1967 Lee, 1991 Privalko, 1986 Richardson, 1977 Tsitsilianis, 1993 Torkelson, 1988 Torkelson, 1993 Schoenhals, 1985 Hourston, 1996 Tanaka, 1992 Tanaka, 1992 Tanaka, 1992 Torkelson, 1992 Torkelson, 1992 Kobayashi, 1989 Schrader, 1989 Tomlins, 1994 Lee, 1990 Rabinovitch, 1992 Rabinovitch, 1992 Lee, 1991 Schoenhals, 1985 Rabinovitch, 1992

983

Aging and Degradation of Polymer Blends

Table 14.2. Physical aging in blends, copolymers and composites: Qualitative treatment System

Polymer

Techniques

References

Blend Blend Blend Blend Blend Blend Blend Blend Blend Blend Blend Blend Blend Blend Blend Copoly Copoly Comp Comp Comp Comp Comp

ABS/PC Aromatic Polyamides PA-66/PPE PMMA PMMA/SAN blends Poly(lactide) blends PS/P(2-VP) PS/P(2-VP) PS/PPE PSF/CPSF PVC/PiPMA PVC/PiPMA PVC/PMMA PPE/PVME PMMA/PEG P(HBA-stat-ETP) PS-b-PMMA PE PEEK PEEK Polyamideimide PPS

G(t) DSC E(t) TSD DSC DSC DSC DSC Dynamic Shear DSC DSC DSC DSC PAL, Mech. PAL, Mech. DSC DSC Creep Impact toughness Creep Creep Impact toughness

Maurer, 1985 Ellis, 1990 Laverty, 1988 Wendorff, 1988 Mijovic, 1989 Wilkes, 1988 ten Brinke, 1989 ten Brinke, 1989 Perez, 1986 Lau, 1993 ten Brinke, 1989 ten Brinke, 1988 ten Brinke, 1988 Chang, 1997 Chang, 1997 Gedde, 1990 Tsitsilianis, 1992 Kubat, 1988 Chen-Chi, 1992 D’Amore, 1990 Geil, 1990 Chen-Chi, 1992

Table 14.3. Physical aging in homopolymers: Quantitative treatment Polymer

Techniques

Model

References

ABS Epoxy Epoxy Adhesives Epoxy network Epoxy Resin P(γ-Bz-L-glutamide) P(γ-Bz-L-glutamide) PBT PC PC PC PC PC

J(t) & G(t) DSC Mech. Torsional dilatometer DSC DETA DSC D(t) DSC DSC DSC Mech. Pos. Anhil.

Booij KAHR & N-M W. Watt KAHR & N-M W. Watt W. Watt Hodge AGV Read Bauwens G-M Hodge W. Watt Free Volume

Booij, 1979 Hutchinson, 1995 Vleeshouwers, 1989 McKenna, 1995 Montserrat, 1994 Colomer, 1989 Colomer, 1989 Read, 1992 Bauwens-Crowet, 1986 Gomez, 1995a Hodge, 1987 LeGrand, 1990 Vleeshouwers, 1992

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J.M.G. Cowie, I.J. McEwen and R. McIntyre

Table 14.3.

Continued.

Polymer

Techniques

Model

References

PC PC PEI PEI & PEEK PEI & PEEK PEK-C PES PHB PMA PMMA PMMA PMMA PMMA PMMA PPE PS PS PS PS PS PS PS PS PS PS PS PS (substituted) PSF PVA PVAc PVAc PVAc PVAc PVAc PVAc PVC PVC PVC PVC PVC PVME Theory

Positron Annih. SHG DMTA DETA DMTA DSC DSC DSC DSC DMTA DSC DSC DSC SHG Mech. Dil DSC DSC DSC DSC DSC DSC DSC DSC DSC DSC DSC HGR DSC Dil Dil DSC DSC DSC Volume D(t) DSC DSC DSC DSC DSC Volume, Mech.

W. Watt W. Watt WLF WLF WLF CF C-F C-F G-M Perez C-F Hodge Hodge, CF W. Watt W. Watt Free volume 2nd ord. kinetics CF G-M Hodge Hodge KAHR KAHR KAHR Petrie C-F C-F VTF C-F Chow Ngai CF, Tau eff. Hodge Moynihan Stochastic Read CF, Tau eff. Hodge KAHR KAHR CF Booij

Hill, 1990 Torkelson, 1989 Hay, 1991 Hay, 1990 Hay, 1990 Zhang, 1993 Jong, 1997 Briddlestone, 1996 Gomez, 1995b Diaz-Calleja, 1989 Cowie, 1995 Hodge, 1987 Cowie, 1993 Torkelson, 1989 LeGrand, 1990 Schwarzl, 1989 Chang, 1989 Agrawal, 1989 Brunacci, 1996 Hodge, 1987 Privalko, 1986 Hutchinson, 1988 Hutchinson, 1989 Hutchinson, 1992 Petrie, 1972 Brunacci, 1997b Brunacci, 1997a Wang, 1994 Cowie, 1997 Chow, 1984 Rendell, 1987 Cowie, 1987 Hodge, 1987 Moynihan, 1978 Robertson, 1984 Read, 1992 Cowie, 1987 Hodge, 1987 Hutchinson, 1992 Hutchinson, 1992 Cowie, 1989 Espinoza, 1990

985

Aging and Degradation of Polymer Blends

Table 14.4.

Physical aging in blends: Quantitative treatment

Polymer blend

Techniques

Model

References

PEEK/PEI PES/Epoxy PES/Epoxy PMMA/PVDF PMMA/SAN PMMA/SAN PMMA/SAN PMMA/SAN PS/PPE PS/PPE PS/PVME

DSC DSC DSC E(t) DSC DSC E(t) E(t) DSC E(t) DSC

W. Watt C-F C-F W. Watt Hodge Hodge Stress Rel. W. Watt Hodge W. Watt C-F

Hay, 1992 Breach, 1992 Hay, 1992 Mijovic, 1991 Mijovic, 1990 Mijovic, 1993 Mijovic, 1990 Mijovic, 1991 ten Brinke, 1992 Mijovic, 1991 Cowie, 1989

Table 14.5. Theoretical treatment of physical aging Topic

Techniques

References

Inorg. Glasses Inorg. Glasses PC PEI PS PVAc Review Review Theory Theory Theory Theory Theory Theory Theory Theory Theory Theory Theory Theory Theory Theory Theory Theory

DSC DSC DSC, Mech Mech DSC Fluorescence DSC E(t) Blends Dil Dil Dil DSC DSC DSC DSC DSC DSC DSC Fractal Dynamics J(t) Ordering params Viscoelastic Volume

Moynihan, 1976 Moynihan, 1984 Liu, 1992 LeGrand, 1990 Hutchinson, 1988 Jamieson, 1990 Oleinik, 1987 Oleinik, 1987 Koenigsveld, 1993 Ibar, 1982 Kovacs, 1982 Matsuoka, 1986 Moynihan, 1976 Moynihan, 1976 Moynihan, 1974 Hutchinson, 1990 Matsuoka, 1986 Hodge, 1985 ten Brinke, 1990 Chow, 1992 Cavaille, 1992 Aklonis, 1987 Aklonis, 1993 Oleinik, 1984

986

J.M.G. Cowie, I.J. McEwen and R. McIntyre

Table 14.6.

Enthalpy relaxation in homopolymers studied by DSC Hodge model (Narayanaswami parameters)

Polymer

ln (A/sec)

∆h*/R, kK

x

β

PVAc PVC PS PMMA PC PMMA

-277.5 -662.0 -211.2 -357.8 -355.8 -206.0

88 225 80 138 150 83

0.27 0.10 0.49 0.19 0.19 0.37

0.51 0.23 0.74 0.35 0.46 0.40

Tr/K986 †

References

317.12 361.74 378.79 385.69 421.59 403.02

Hodge 1987 “ “ “ “ Cowie 1993

Tr/K †

References

318.54 363.69 380.49 386.86 430.69 305.78

Hodge 1987 “ “ “ “ Colomer 1989

Hodge model (Adam Gibbs Vogel parameters) Polymer

ln (A/sec)

PVAc PVC PS PMMA PC PγBLG

-66.60 -59.74 -100.30 -55.45 -70.30 -40.93

†

β

T2/K

D/R, kK 6.23 2.61 17.10 3.43 7.43 3.02

225 320 210 325 325 232

0.55 0.28 0.74 0.34 0.34 0.25-0.36

Here, Tr is defined as the temperature at which τreq. = 1 second. For the Hodge N model, Tr = -1000(Dh*/R)/lnA, and for the Hodge AGV model, Tr = T2 is more useful than the lnA parameter for the interpretation of the results.

Table 14.7.

Mn x 10-3 9 17 110 233 196‡ † ‡

Effect of molecular weight on enthalpic aging of PS. Cp data analyzed by Hodge model, Narayanaswami parameters [Privalko (1986)] ln A -198.8 -234.4 -263.0 -287.8 -211.6

∆h*/R, kK 76.1 88.8 101.0 110.0 80.5

from Tr = -1000(∆h*/R)/Ln A, A in seconds. broad molecular weight sample.

x

β

Tr/K†

0.40 0.43 0.42 0.41 0.48

0.63 0.53 0.50 0.47 0.58

382.80 378.84 384.03 382.21 380.43

987

Aging and Degradation of Polymer Blends

Effect of inert filler (glass) on enthalpic aging of PS (196 kg mol-1); Cp data analyzed by Hodge model, Narayanaswami parameters [Privalko (1986)]

Table 14.8.

Wt% filler 0 1 20 50 70 †

ln A

∆h*/R, kK

x

β

Tr/K†

80.5 80.9 84.0 91.0 98.0

0.48 0.49 0.48 0.36 0.28

0.58 0.50 0.45 0.43 0.38

380.43 378.04 376.01 373.41 375.77

-211.6 -214.0 -223.4 -243.7 -260.8

from Tr = -1000(∆h*/R)/Ln A, A in seconds.

Table 14.9.

Data derived from the C-F model

Polymer

Ta/K

Tg-Ta/K

PMMA

375.0 377.5 380.0 382.5 385.0 387.5

19.9 17.4 14.9 12.4 9.9 7.4

1.597 1.460 1.338 1.244 1.219 1.062

2.791 2.490 2.428 2.208 2.103 2.149

0.45 0.42 0.43 0.37 0.35 0.30

4.669 4.498 4.376 4.489 4.488 4.959

Cowie 1993 “ “ “ “ “

PVAc

303.0

10.0

2.813

1.751

0.34

4.191

Cowie 1987

PVC

346.0

8.0

1.469

2.365

0.30

5.163

Cowie 1987

PVME

5.6 8.1 10.6 15.6 20.6

1.08 2.43 2.83 3.24 3.77

1.23 1.77 1.82 2.44 2.84

0.91 0.51 0.66 0.50 0.62

2.15 3.41 3.09 4.12 4.20

Cowie 1989 “ “ “

PEK-C

250.0 247.5 245.0 240.0 235.0 “ 483.2

2.55

1.724

0.55

3.261

Zhang 1993

PS 17.5 k

362.2

8.0

1.527

3.140

0.57

4.613

Agrawal 1989

PS 206 k

365.2 367.2

8.0 6.0

1.517 1.037

3.185 2.778

0.54 0.63

4.739 4.110

“ “

PS 1260 k

370.2 372.2 373.2

7.4 5.4 4.4

1.651 1.133 0.912

3.320 2.740 2.519

0.51 0.64 0.69

4.966 4.051 3.735

“ “ “

Epoxy**

323.2 333.2 343.2 353.2 363.2

10.321 7.963 6.160 4.349 2.747

0.18 0.22 0.25 0.28 0.21

15.010 11.744 9.464 7.304 6.706

∆H∞(Ta)/Jg-1

19.8 15.1 10.8 6.7 3.0

log10(tc/min)

β

log10(te/min)*

References

Montserrat 1994 “ “ “ “

* te = time here for ∆H(ta,Ta) to approach 99.9% of the equilibrium value. ** here, ∆H∞(Ta) was estimated from extrapolation of the liquid enthalpy curve of the unaged sample, and NOT treated as an adjustable parameter.

988

J.M.G. Cowie, I.J. McEwen and R. McIntyre

Table 14.10. Parameters derived using the Petrie model Polymer PS PVME Epoxy PET PMMA PC PVC PEI

ln (A/min-1) 366.3 192.3 337.1

EH/kJ mole-1 1192 407 1020 860 507 1731 1409 1470

C/gJ-1 2.92 1.14 2.10

References Petrie 1972 Cowie 1989 Montserrat 1994 “ “ “ “ Hay 1991

Figure 14.5. DSC thermograms for aged polymer blends (a) polyvinylchloride/poly isopropyl methacrylate, immiscible blend, aged at a temperature of 60°C, and (b) polyvinyl chloride/polymethylmethacrylate, miscible blend, aged at 80°C. Time of aging, ta in hours, is shown alongside each curve. Broken lines represent the un-aged samples for comparison.

14.3.2.1 Homogeneous Blends — Thermal Methods One of the most commonly used criteria for establishing the phase behavior in amorphous binary polymer blends is the presence of one or more Tgs. If the blend is one-phase, a single Tg

lying between the values for each component, is detected and characterizes the mixture. If the blend is two-phase then two Tgs are observed close to or matching those of the two components. On occasions the component glass transition temperatures are so close together that it is difficult to discriminate between the presence of one

Aging and Degradation of Polymer Blends

or two Tgs in a DSC measurement at normal scan rates. Both Bosma et al. [1988] and Jorda and Wilkes [1989] demonstrated that the use of isothermal aging experiments could overcome this problem. It was argued that in a homogeneous, one-phase, blend the kinetics of the aging process would be an average representing the blend and as such would exhibit only one enthalpy recovery peak. If, on the other hand, the blend was twophase the components would age at their own individual rates and this would be manifest in the appearance of two enthalpy recovery peaks on the DSC trace. The technique was demonstrated for several systems. Blends of poly(vinyl chloride), PVC, (Tg = 80°C) and poly(isopropyl methacrylate), PiPMA, (Tg = 82.5°C), are believed to be immiscible but because of the closeness of the Tg values this is difficult to confirm. A 50/50 blend was annealed first at 195°C, to erase previous thermal history, then quenched to 60°C, i.e., (Tg - Ta) ~20°C and aged for various times ta [Bosma et al., 1988]. The thermograms shown in Figure 14.5(a) that were obtained demonstrate clearly the development of two distinct enthalpy recovery peaks that increase with increasing aging time ta. Good separation of the peak maxima are obtained as Tmax tends to increase with aging time for both components but at different rates. This indicates that phase separation has occurred and the enthalpy recovery peaks are characteristic of each distinct phase in the mixture. A miscible blend of PVC and atactic polymethylmethacrylate, PMMA, was also treated in a similar fashion but only one enthalpy recovery peak could be detected, Figure 14.5(b), indicating a single phase system. This work was extended by Grooten and ten Brinke [1989] to include the immiscible blend of polystyrene, PS, with poly(2-vinyl pyridine), P2VP. The authors concluded that the most appropriate range of aging temperature for this type of experiment was Tg to ~(Tg - 20°C). The method was also applied to distinguish the two-phase nature of PS-P2VP diblock copolymers [ten Brinke and Grooten, 1989]. The sensitivity of this method was demonstrated by Jordan and Wilkes [1989] who showed, after aging, that blends of racemic polylactide and

989

its optically active L-form were two-phase, a convincing demonstration that two stereo-regular forms of a polymer may be immiscible. Similar levels of sensitivity were observed by Ellis [1990] in his studies on the phase behavior of blends of aromatic polyamides. Having predicted the phase behavior of some of these blends using the mean field binary interaction model, [ten Brinke et al., 1983], [Paul and Barlow, 1984] proof was required for blends with components having similar values of their Tg. The aging experiments confirmed the predicted phase behavior. The first comprehensive study of physical aging in a miscible blend system using enthalpy relaxation was reported by Cowie and Ferguson [1989] who followed the enthalpic relaxation in a series of blends of PS and poly(vinyl methyl ether), PVME. Comparison of the blend behavior with that of the two components by analyzing the data on the basis of both the P-M and C-F models led to the conclusions that, the blends aged more slowly than PVME when aging was carried out at a comparable temperature below Tg, hence the component with the lower Tg, i.e., the more mobile component in the blend, PVME, was responsible for most of the aging effects seen in this blend (see Table 14.11). The PS component did not appear to contribute significantly to the total aging because the relaxation processes were much slower than PVME at the aging temperature. This suggests that even though the blend can be regarded as a miscible, one-phase system, the components can largely relax independently although the relative rates of each will be influenced by the second component. In the PS/PVME blends the Tgs of the two components are quite different and the blends can be regarded as comprising a flexible polymer (PVME) mixed with a relatively stiff polymer (PS), a situation which is not favored thermodynamically. It is interesting to note that if an even more rigid analogue, poly(α-methyl styrene) replaces PS in the blend, a two-phase system is obtained. Thus PS/PVME may not be a typical miscible blend system, as its glass transition region is broad, indicating a mixture close to phase separation.

990

J. M. G. Cowie, I. J. McEwen and R. McIntyre

Table 14.11. PS/PVME blends [Cowie (1989)] Petrie model wt% PVME

ln (A/min-1)

EH/kJ mole-1

C/g J-1

0 50 100

366.6 388.2 192.3

1192 929 407

2.92 4.08 1.14

C-F - model Blend

Ta/K

Tg-Ta/K

∆H∞(Ta)/Jg-1

log10(tc/min)

β

log10(te/min)

PS/PVMEà

270.0 265.0 260.0 255.0 250.0

11.8 16.8 21.8 26.8 31.8

1.08 1.39 1.65 1.88 2.09

2.15 2.46 2.47 2.74 3.25

0.449 0.316 0.373 0.412 0.364

3.83 5.12 4.72 4.77 5.56

Oudhuis and ten Brinke [1992] have addressed these points by examining the aging of blends of PS with poly(2,6-dimethyl-1,4-phenylene ether), PPE, where now both components are relatively rigid, although there is still about 100°C difference between the Tg values. These authors found that while the amount of enthalpy relaxation observed in the blends was lower than that for either component there was no evidence for faster relaxation by the component with the lower Tg, viz. PS (Table 14.12). This observation has been confirmed by Cowie and Elliot [1990] for this system using three different blend compositions. Oudhuis and ten Brinke suggested that since the enthalpic definition of Tg was used as proposed by Cowie and Ferguson [1989] and the blend showed a broad glass transition region covering about 23°C, the use of the onset Tg instead of the mid-range temperature might be a more accurate reference for selection of the aging temperatures. When a broad transition region is present, some aging temperatures close to the enthalpic Tg might impinge on the onset region of the glass transition process and show an accelerated aging for the more flexible component. This illustrates some of

the problems associated with accelerated aging studies where the precise location of Tg for a weakly miscible blend may be difficult to define. However, an alternative explanation may be found in the “Sequential Aging Theory” proposed by Chai and McCrum [1980]. The authors postulated that at a given Ta and ta, the viscoelastic elements with relaxation times equivalent to ta will be aging, but that elements with τ < ta will already have reached equilibrium and those with τ > ta will not yet have begun to move towards equilibrium. Thus in the PVME/PS case the more flexible PVME, which will also be closer to its own Tg, will possess more elements with shorter τ than the PS. Consequently the PVME relaxation spectrum will tend to move more rapidly towards equilibrium than the PS and so would age more rapidly. This idea will require further testing. Blends of PMMA with a commercial sample of poly(styrene-stat-acrylonitrile), SAN, containing 25 wt% AN, have been studied using enthalpy [Mijovic et al., 1989] and stress relaxation [Mijovic et al., 1990] measurements. Data are shown in Tables 14.12 and 14.13. The authors observed, in common with all other workers, that aging is

991

Aging and Degradation of Polymer Blends

Table 14.12. Enthalpy relaxation in blends of PS/PPE and SAN/PMMA analyzed using the Hodge model (N-parameters) Blend

ln (A/sec)

w2

SAN/PMMA*

PS/PPE**

100 80 60 40 20 0 0 50 100

-328.9 -358.5 -370.8 -392.6 -377.6 -359.8 -288.1 -316.6 -328.8

∆h*/R, kK 125.0 134.7 138.6 146.2 140.3 132.2 140 135 126

x

β

Tr/K

References

0.221 0.137 0.147 0.253 0.227 0.338 0.37 0.33 0.28

0.48 0.46 0.42 0.35 0.35 0.26 0.58 0.56 0.57

380.05 375.73 373.79 372.39 371.56 367.43 485.94 426.41 383.21

Mijovic 1989 “ “ “ “ “ Oudhuis 1992 “ “

w2 = weight % first homopolymer component in blend. * values for unaged ** values for ta = 120 minutes

Table 14.13. Physical aging in polymer blends, studied by Stress relaxation, Mijovic equation* parameters System SAN/PMMA

PS/PPE

w2**

Cw

0 20 40 60 80 100 0 20 40 60 80 100

-7.72 -7.65 -7.49 -7.25 -7.02 -6.68 -8.08 -10.34 -14.55 -14.35 -10.42 -8.07

CT

Ct

3.325 3.325 3.325 3.325 3.325 3.325 3.61 4.24 5.15 5.15 4.31 3.68

1.013 1.013 1.013 1.013 1.013 1.013 0.357 0.683 0.885 0.936 0.838 0.849

References Mijovic 1990 “ “ “ “ “ Ho 1991 “ “ “ “ “

* Mijovic Equation: E(t)/Eo = exp[-(t/τ)β] where ln τ = Cw + CT ln(Tg - Ta) + Ct ln ta ** w2 = weight % first homopolymer component in blend, Note: Ct = ƒ(ln τ)/ƒ(ln ta) = µ

Table 14.14. Aging parameters calculated from the C-F and P-M analysis of data obtained from PMMA, SAN (26.6 wt% AN) and their (50/50) blend PMMA/SAN at Ta = (Tg - 10 K) Polymer

PMMA SAN (26.6 wt% AN) PMMA/SAN (50/50)

C-F model

P-M model

∆H(∞)

log tc

β

log te

∆H(∞)

0.820 2.365 1.766

1.780 1.864 1.678

0.290 0.422 0.498

4.678 3.853 3.362

2.381 3.386 4.142

log te 11.479 4.697 6.008

992

J. M. G. Cowie, I. J. McEwen and R. McIntyre

faster at higher temperatures. No comparison was made with the component polymers but it was observed that blends rich in SAN relaxed faster than PMMA- rich blends. The aging times used in that work were no more than 150 min. A more comprehensive study of this system has been carried out by Cowie and Ferguson [1991]. These authors followed the enthalpy and stress relaxation of a series of PMMA/SAN blends with SAN compositions spanning the miscibility window, i.e., from 13.3 to 30 wt% AN. It was found that the blends relaxed faster than either of the components, when aging temperatures were (Tg - Ta) = 10°C, but that this was no longer the case at (Tg - Ta) = 20°C where blend aging was intermediate to both components. The data were analyzed using both the P-M and C-F approaches and examples of ∆H(ta, Ta) vs. log ta plots are shown in Figure 14.6. The C-F model

provided a better estimate of long-term aging effects where over-estimates of the time taken to reach equilibrium were made by the P-M model. The parameters derived from each analysis are shown in Table 14.14 for PMMA/SAN (26.6 wt% AN) at Ta = (Tg - 10°C). Enthalpy relaxation studies have also been used to assess the aging of polyether ether ketone blends with polyetherimide, PEEK/PEI = 50/50 [Hay, 1992]. The preparation of the blend produced an amorphous system with Tg ~215°C, but crystallization of the PEEK occurred after raising the temperature above Tg. Enthalpic relaxation could only be observed in the temperature range Tg to (Tg - 50) and no aging could be detected at temperatures below 150°C (Table 14.15). The system was analyzed using the KWW Eq 14.7 that yielded values of β = 0.4, intermediate between those of the component polymers.

Figure 14.6. Comparison of the experimental enthalpy change on aging for ta minutes, PMMA, SAN containing 26.6 wt% AN, and a 50/50 blend of these two polymers, with the theoretical curves derived from the Petrie-Marshall (PM) and Cowie-Ferguson (CF) models.

Aging and Degradation of Polymer Blends

When the blends were examined using dielectric relaxation measurements which probe the dipole relaxation spectrum, values of β were found to be much lower (0.1-0.22). This was interpreted as indicating the development of heterogeneity at the molecular level caused probably by the crystallization of the PEEK component. Polyether sulfone, PES, can be blended with epoxy resins in certain combinations that do not lead to phase separated systems, and Breach et al. [1992] have investigated the aging characteristics of Epikote 828 (Shell) and Victrex 5003P. Comparison of aged and un-aged samples allowed the enthalpy relaxation to be calculated from the peak areas. Data are collected in Table 14.16. It was found that the blends aged at a faster rate than the components, and that this increased with increasing PES content. It was observed that β also increased and these effects were explained by considering that the incorporation of PES loosened the epoxy resin network thereby increasing the free volume. As this would change the size distribution of the free volume “holes”, the presence of larger holes would accelerate the relaxation process and narrow the relaxation time distribution. 14.3.2.2 Homogeneous Blends Mechanical Methods Since there is a shift in the viscoelastic relaxation spectrum to longer times with aging time, aging can be followed using stress relaxation, creep or volume relaxation measurements. It has been shown that for aging experiments [Struik, 1978], momentary creep curves have a universal shape and a master curve can be constructed using either time-aging (t - ta) or time-temperature (t - T) superposition. The stress relaxation modulus can then be expressed in terms of the KWW Eq 14.7 and the equation proposed originally by Struik for analyzing creep compliance: G(t) = Go exp [- (t/τ)β]

(14.13)

where Go and G(t) are the stress relaxation moduli at zero time and time t respectively.

993

Hence Φ(t) = G(t)/Go. Several homopolymer systems have been studied and the results are collected in Tables 14.17, 14.18, and 14.19. As long as the measuring time is short compared with the relaxation or retardation times the aging process can be studied effectively, and short-term stress-relaxation measurements have been carried out on blends of SAN/PMMA and PS/PPE [Ho et al., 1991]. Eq 14.13 was used to fit the data and the authors found that τ could be expressed by: ln τ = A + CT ln (Tg - Ta) + Ct ln ta

(14.14)

where A is a function of the weight fraction of one of the components in the blend, while CT and Ct are adjustable temperature and time coefficients. Data are listed in Table 14.13. Ho et al. also observed that the values for β fell within a narrow range centered on what was called a “universal value” β = 0.41. While many results suggest β does not vary significantly, it is more likely that β is an increasing function of (Tg - Ta), but that it may eventually reach a plateau for large values of (Tg - Ta). 14.3.2.3 Positron Annihilation Lifetime Spectroscopy (PALS) Three blends have been studied by Chang et al. [1997] using stress relaxation measurements and positron annihilation lifetime spectroscopy (PALS). It was observed that for blends of PS with PPE and PS with PVME the stress relaxation rates were faster for the blends in comparison with PS alone, whereas the opposite was true for a PMMA/PEG blend when compared with neat PMMA. As aging tends to lead to changes in the packing density of the system, PALS was used to study the blends as it will provide qualitative estimate of the free volume in the system. The stress relaxation results were confirmed qualitatively using PALS. It was found that PS/PPE and PS/PVME blends were less dense than PS, while PMMA/PEG was denser than PMMA, thereby making chain relaxation easier in the former and more difficult in the latter.

994

J. M. G. Cowie, I. J. McEwen and R. McIntyre

Table 14.15. PEEK/PEI blends analyzed using the C-F model [Hay (1992)] Ta/K

∆H∞(Ta)/Jg-1

log10(tc/min)

β

457-482 429-445 389-410

not given not given not given

not given not given not given

0.35 0.40 0.55-0.60

wt% PEEK 0 50 100

Table 14.16. Epoxy resin blends with polyethersulphone data derived using the C-F model from enthalpy relaxation measurements [Breach (1992)] Blend E828DDS/PES† 0 E828DDS/PES 20 E828DDS/PES 30

Ta/K

Tg-Ta/K

453.2 453.2 453.2

... ... ...

∆H∞(Ta)/Jg-1

β

log10(tc/min)

3.35 3.13 2.52

3.672 3.515 2.881

log10(te/min)

0.26 0.36 0.52

6.900 5.846 4.495

† Epikote 828 resin (Shell) cured with 4,4’-diaminodiphenylsulphone (DDS) and blended with Polyethersulphone (Victrex 5003P).

Table 14.17. Physical aging data for several homopolymers [LeGrand (1990)] Stress relaxation Material PC PEI PPE PPE

History

Strain

T/°C

Er

∆E1

∆E2

Quenched Quenched Quenched Annealed

0.5 0.5 0.5 0.5

50 145 145 145

1.38 0.8 0.5 2.3

0.72 0.55 1.45 0.3

0.49 1.4 0.2 0.15

α 0.15 0.15 0.33 400

β

τα

0.33 0.33 0.33

400 400 3000

τβ 90000 30000 100000

Strain recovery Material

History

T/°C

Ep

δE1

δE2

α

β

τα

τβ

PPE PPE

Quenched Annealed

30 30

0.1 0

0.3 0

0.2 0.118

0.33 0.15

0.15 400

16000

400

τα

%V2

β

τβ

%V0

0.15

0.15

400

-0.35

Volume relaxation Material

History

T/°C

%V

α

PC

Quenched-SR

50

0.8

0.35

350000

Thermal activated strain recovery Material

History

PC

Quenched-SR

T/°C

ε1

ε2

εp

50 70 90

3 11 11

4 2 2.2

1 0.1 0.1

995

Aging and Degradation of Polymer Blends

Table 14.18. Physical aging data for PVAc derived using stress relaxation [Jamieson (1990)] Ta/K

ta/h

log ta

log to

log Go

Slope

Intercept

293.5

0.25 0.5 1 2 4 7 23.5

2.954 3.255 3.556 3.857 4.158 4.401 4.927

2.45 2.67 2.93 3.20 3.48 3.70 4.28

9.198 9.217 9.214 9.219 9.217 9.217 9.236

0.9236

-0.3328

295.7

0.25 0.5 1 2 3.7 6 22.5

2.954 3.255 3.556 3.857 4.124 4.334 4.908

2.34 2.54 2.77 3.06 3.30 3.44 4.02

9.202 9.206 9.209 9.215 9.214 9.218 9.223

0.8625

-0.2586

298.2

0.25 0.5 1 2 5 70

2.954 3.255 3.556 3.857 4.255 5.401

2.03 2.34 2.60 2.82 3.17 4.12

9.210 9.200 9.201 9.208 9.208 9.216

0.8434

-0.4256

Table 14.19. Physical aging Williams Watt parameters derived for PC using positron annihilation [Hill (1990)] Polycarbonate Calibre (as molded)

Calibre (annealed above Tg)

Lexan (aged at 120C for 120 hr)

Ta/K

to/min

β

∆Ea/kJ mole-1

303.16 273.16 263.16 78.43

36.67 64.27

0.8505 0.8172

12.552

0.6891

303.16 293.16 283.16

44.16 55.2 70.32

0.5732 0.8358 0.8921

16.736

303.16 293.16 285.16

4037.4 8979.6 9332.4

0.6223 0.5626 0.4667

34.309

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J. M. G. Cowie, I. J. McEwen and R. McIntyre

14.3.2.4 Heterogeneous Systems Maurer et al. [1985] have examined the two-phase blend of acrylonitrile-butadiene-styrene copolymer, ABS, (Tg = 110°C) and polycarbonate of bisphenol-A, PC, (Tg = 151°C) using stress relaxation measurements. Four regimes of behavior were found. Below 70°C both (t - T) and (t - ta) superpositions were possible, because the aging rates µ defined as: µ = d(log τ)/d(log ta)

(14.15)

of both components were equal. Between 70°C and 100°C only (t - ta) superposition worked. However, close to the Tg(ABS), and between the Tgs of both components neither were valid. While Eq 14.13 has been used by other workers Booij and Palmen [1978] found that a more precise form was: G(t, τ) = Go(t/τ)α exp [-(t/τ)β]

(14.16)

where β = 0.4 was regarded as a constant (although this might be structure dependent) and α was normally a very small number. This form was used to analyze the ABS/PC blend data. The behavior of the two-phase systems is complex and responses on aging can be affected by the thermal history and aging temperature. This was illustrated by aging ABS/PC blends in regime 3 and quenching from regime 4. Quenching to 95°C followed by aging gave curve ‘a’ in Figure 14.7 whereas slow cooling to 95°C then heating to 115°C and quenching back to 95°C produced curve ‘b’. For the first (rapid quench) both components age from the glass at a similar rate, but in the second case the thermal history ensures ABS is largely relaxed whereas PC still has to age fully from the glassy state. This change in the relative aging rates of the two components allows a modulus plateau to develop in the modulus-time curve, and illustrates the need for care with thermal histories [Booij and Palmen, 1978]. Polycarbonate has been found to lose its toughness when aged at temperatures below its Tg and impact modifiers are often added to counteract this effect of physical aging. Blends of PC with

several core-shell methacrylate-butadiene-styrene impact modifiers have been studied after aging at 125, 130 and 135°C [Cheng et al., 1992]. Tests for changes in impact strength, tensile strength, dynamic mechanical response and fracture morphology were complicated by the simultaneous chemical degradation of the samples when aged in air, but the modifiers did appear to slow down the sample embrittlement caused by physical aging. The increase in the toughness lifetime of the modified PC depended on the amount and nature of the impact modifier used, and the aging environment conditions. The effectiveness of the modifiers was reduced by their chemical degradation at elevated temperatures, particularly in the region of the sample surface. A number of composites of carbon fiber with engineering polymers have become increasingly important in the aerospace industry. Physical aging is a crucial feature in their long-term stability and this has stimulated studies of this effect in PEEK,

Figure 14.7. The shear modulus plotted as a function of aging time (in seconds) at 95°C, for a (34/66) ABS/PC blend, prepared using two different thermal histories.

997

Aging and Degradation of Polymer Blends

polyphenylene sulfide PPS, and polyamideimide, PAI, based composites. D’Amore et al. [1990] reported that while the crystalline regions in PEEK and the carbon fibers themselves do not participate in the aging process, they are likely to constrain the mobility of chains in the amorphous regions near the inclusion boundaries. This led to an increase in the composite Tg and lowered the material sensitivity to physical aging. Creep compliance (J) measurements were made at various aging times but the curves were modeled using the Findley relation [Findley et al., 1976]: J = Jo + A (t/to)n

(14.17)

rather than Eq 14.13. The time-temperature shift factor aT was calculated from the curves at various aging times ta, and was found to increase with increasing ta. The rate of aging was obtained from: µ = d(log aT)/d(log ta)

(14.18)

It was observed that crystallinity in PEEK slowed the aging rate. This can also be seen in Figure 14.8, where µ is plotted as a function of the undercooling (Tg - Ta). Close to Tg the rates of aging for the amorphous, semi-crystalline PEEK and PEEK composites were similar but,

Figure 14.8. Comparison of the shift rate for crystalline, amorphous, and composite PEEK as a function of the extent of under-cooling (Tg-T°C).

as the aging temperature dropped further into the glassy state, retardation of the aging process in the composite and the semi-crystalline PEEK became more obvious. Thus: µamorph > µcryst > µcomp indicating that the aging process becomes less self-retarding as the amount of filler increases. The influence of physical aging on the comparative and intrinsic toughness of PEEK and PPS carbon-fiber composites was determined by Ma et al. [1992]. Both types exhibited a loss of toughness with increase in aging time and temperature. The PEEK composites tended to retain more impact toughness but after aging were less ductile than the PPS composites. Nichols et al. [1990] examined carbon-fiber composites where a PAI was used as the thermoplastic matrix (Amoco A1696/T650-2). At high loadings the non-recoverable creep strain in the composite depended on the length of time the material remained under load and this effect was attributed primarily to physical aging of the specimen. These viscoelastic creep measurements were made after aging for 14 h at 225°C (Tg = 250°C). The data for the aged and un-aged samples are compared in Figure 14.9. It can be seen that the plateau strain is lower in the aged sample. Differences were also exemplified in the steady state creep rates that could be up to 50 times greater in the aged, compared with unaged, samples. This was explained by assuming that the increase in stiffness due to aging leads to development of a smaller strain under load for the aged material compared with the un-aged sample where the chains slip past one another more easily. These authors also highlighted the fact that the reduction in free volume on aging must lead to shrinkage of the sample during a creep measurement. Consequently, this effect opposes the deformation and negative creep may be seen in low-stress tests. When these were carried out the master curve constructed from the data at various Ta deviated from that expected, see Figure 14.10. Because aging effects led to an over-prediction of the strain when short-term data

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J. M. G. Cowie, I. J. McEwen and R. McIntyre

Figure 14.9. Plot of creep strain as a function of time for aged and un-aged samples of polyamideimide-carbon fiber composites, at a temperature of 225°C.

Figure 14.10. Master curve for short-term creep tests on polyamideimide-carbon fiber composites, made at various temperatures. This shows the over-prediction for creep at higher temperatures, caused by physical aging having stiffened the polymer.

were extrapolated to long-term responses. Nichols et al. concluded that while physical aging reduced the undesirable creep effects in thermoplastics at elevated temperatures, the reduction in fracture toughness was too great and tended to negate

any advantages a thermoplastic might have over a thermoset in these composite materials. Work on Epon 828/Epoxy adhesives has been reported [Jamieson 1989] and the data are given in Table 14.20.

999

Aging and Degradation of Polymer Blends

Table 14.20. Physical aging data for Epon 828/Epoxy adhesives [Jamieson (1989)] 1) Creep data System

Ta/K

Epoxy A

325.2 10 20 40 80 110 315.2 20 40 80 17000 315.2 10 20 40 80 160 320

Epoxy B

Epoxy C

β 0.3 590000 440000 330000 1500000 1900000 0.45 4700 9900

ta/min

te/sec

log ta

log te

5 2.778 3.079 3.380 3.681 3.820 10 3.079 3.380

100000 5.771 5.643 5.518 6.176 6.279 2200 3.672 3.996

2.477

5.000

2.778

3.342

4.230 5 2.778 3.079 3.380 3.681 3.982 4.283

5500 4.041 4.322 4.415 4.740 4.857 5.279

2.477

3.740

ta/min

te/sec

log ta

log te

5 2.778 3.079 3.380 3.681 3.820 5 2.778 3.079 3.380 3.681

5200 5.771 5.643 5.518 6.176 6.279 310 2.613 2.820 2.944 3.342

2.477

5.000

2.477

2.491

3.708 5

1000

αstress

βcreep

µcreep

µstress

0.23 0.45 0.4

0.3 0.45 0.45

1 1 0.8

1 0.9 0.8

3.681 0.45 11000 21000 26000 55000 72000 190000

2) Stress relaxation data System

Ta/K

Epoxy A

325.2 10 20 40 80 160 315.2 10 20 40 80 160 5100 315.2 10 20 40 80 160

Epoxy B

Epoxy C

α 0.23 46000 70000 140000 260000 270000 0.45 410 660 880 2200 3.982 0.40 2000 3100 4900 11000 15000

3) Summary data Epoxy A B C

to creep/to stress 11 8 6

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J. M. G. Cowie, I. J. McEwen and R. McIntyre

14.4

Chemical Aging

14.4.1

Thermal Degradation of Polymer Blends

Polymer degradation reactions can be divided into two classes, (a) main chain reactions leading to scission and/or depolymerization and (b) substituent reactions at side chains. During the course of the reaction in the former unaltered parent material will be present whereas this is not so in the latter. Degradation may be induced thermally, by high energy radiation, or during melt processing where high shear rates can cause bond breaking by mechanical means. Shear-induced backbone scission introduces reactive radical centers which may then initiate further reactions. An extensive literature exists detailing the degradative behavior of many homopolymers and copolymers [Grassie, 1964, 1966; Conley, 1970; Jellinek, 1978; Schnabel, 1981; Grassie and Scott, 1985; Rabek, 1987]. Polymer blends are generally assigned as miscible (single phase) or immiscible (multiphase) on the basis of some convenient physical test such as the presence or absence of a single glass transition (Tg) [Olabisi et al., 1979]. A blend with a single Tg is deemed miscible on a scale at or below the total number of segments cooperatively relaxing at the glass transition. Miscibility on the single-segment level is probably unattainable (due to the constraints of chain flexibility) so a characteristic domain size will be present in all blends; this can range from several repeat units (nominally miscible) to several microns (definitely immiscible). If the degradation of one component is to be modified by the presence and possibly degradation of another, then diffusion of active species across phase boundaries, or direct reaction at the boundary, must occur and so to some extent any interaction between degrading blend components will depend on the degree of dispersion of the phases. More important is whether or not these interactions are synergistic in the sense that they lead to accelerated degradation of one, or all, of the blend components with consequent loss

of properties, or whether overall retardation of degradation results. Sacrificial degradation of one component may be desirable in that it offers a protection mechanism for the other(s) and thus fulfills the role of a polymeric stabilizer. An earlier review of polymer blend degradation is available [McNeill, 1977]. Since then, increasing interest in polymer blends as an economical route to better service properties has continued to generate interest in the thermal behavior of miscible and immiscible blends [Lizymol and Thomas, 1993] and a guide to the best choices of components for enhanced blend stability has been compiled [Wypych, 1992]. 14.4.1.1 Polyvinyl Chloride The least stable of the high tonnage polymers is PVC in which dehydrochlorination and subsequent polyene formation is the major feature. The rate of dehydrochlorination is accelerated by the presence of oxygen and radicals. PVC is probably the most widely studied polymer and continuing interest has been expressed in its blends. The interaction of diffusing reactive species was studied some time ago in heterogeneous PVC blends with PS and with polyacrylonitrile, PAN [Dodson and McNeill, 1976; McNeill et al., 1978]. Since then many blends have been investigated; with poly(dimethyl siloxane), PDMS, and PC [McNeill and Basan, 1993], with PS and graft copolymers with acrylonitrile and butadiene, NBR [Braun and Kommerling, 1992], with cis-poly(butadiene) [Naqvi and Sen, 1991], with polyester [McNeill et al., 1991], with poly(methacrylates) and poly(acrylates) [Braun et al., 1990], with poly(α-methyl styrene-coacrylonitrile) [Go, 1990; Goh, 1993; Kovacic et al., 1993], with rubbers [Varughese, 1988; Singh and Singh 1984] and with various copolymers [Kaminska and Kaczmarek, 1984]. The fate of organotin stabilizers in PVC blends has also been followed [Muller, 1989] and the problems of PVC stabilization in its blends has been discussed [Minsker et al., 1990]. The effects of miscibility in PVC blends have been specifically examined [Moskala and Lee, 1989]. The auto-catalyzed dehydrochlorination

Aging and Degradation of Polymer Blends

of PVC is retarded in miscible mixtures (with poly(ε-caprolactone), PCL, and poly(ethylene adipate)) but not when the blends are phaseseparated [Goulet and Prud’homme, 1986]. The requirement of miscibility is apparently not general though since the two-phase blend system with poly(isobutylene), PIB, was also reported to suppress dehydrochlorination [Kolawole and Olugbemi, 1985]. PVC blends with other chlorinated materials, although generally compatible, offer little advantage over pure PVC due to catalysis by HCl [Dodson et al., 1974; Jamieson and McNeill, 1974; Singh and Singh 1984]. Blends of plasticized PVC with PMMA are reported to lose plasticizer at relatively low temperatures (60°C) [Razinskaya et al., 1988]. Unplasticized PVC blends with PMMA can form miscible foamed structures above 200°C with good mechanical properties, but since the foaming agent is HCl this has little practical application [Jayabalan, 1982]. In the main, however, no clear picture emerges and certainly no worthwhile increases in stability have been obtained from PVC blends compared with homopolymer. A single study [Minsker et al., 1990] with a range of blends summarizes the general conclusion; degradation depends on the chemical nature of the second polymer, its content in the blend and the conditions of degradation. 14.4.1.2 Polystyrene The thermal degradation of PS involves an initial molecular weight loss and a significant yield of monomer. Studies suggest that the α-hydrogen is involved in transfer processes. PS forms well-characterized miscible blends with PVME in which weak specific interactions have been proposed [Garcia, 1984]. Thermal degradation of PS blends with PVME is characterized by an induction period after which oxidation and molecular weight decrease of the vinyl ether component is rapid. A hindered phenol is an effective antioxidant [Kim et al., 1992]. The same effect can be obtained in PVME blends with poly(styrene-co-vinyl phenol), where additional stabilization from cross-termination was postulated

1001

[Park et al., 1990]. Phase separation occurred at the end of the induction period, but low molecular weight PS remain miscible and increased stability was obtained [Pearce et al., 1990] which seems to confirm the involvement of a radical transfer process between the components. Polystyrene also forms miscible blends with PPE in which the hydrogen-transfer reaction of PS is rendered less effective due to the PPE component [Jachowicz et al., 1984]. In contrast, poly(α-methyl styrene) blends with PPE, although also reported miscible, show no evidence of any interactions during degradation. Head-to-head and head-to-tail PS structures were found to behave differently [Kryszewski, et al., 1982]. PS blends with PMMA as well as PS with polybutadiene, PB, and PVC have been studied [Kaminska et al., 1991; 1989]. These blends showed decomposition temperatures above those of the individual polymers. The relative increase in decomposition temperature was greatest when the fraction of the second component in the blend was low. For example 1-5% of PVC was effective and this possibly reflects the extent of miscibility of the blends and the consequent effectiveness of the second component as a radical “sink”. In a separate study [Kaminska and Kaczmarek, 1990], PS blended with PB was found to be effectively stabilized by as little as 1% of the diene and increased amounts had little further effect. Here the origin of the stabilization is both crosslinking and oxidation of the diene component. In blends with polyisoprene an element of enhanced stability was also achieved. Although chain scission of the PS increased, volatiles were produced less readily by the blend than in each component alone [McNeill and Gupta, 1980]. Polyvinylacetate, PVAc, was also established as providing a stabilizing effect on PS [Jamieson and McNeill, 1976]. Here, and in the diene blends, diffusion of radical-scavenging volatiles into the PS phase was suggested to be important [McNeill et al., 1978]. 14.4.1.3 Acrylics Thermal degradation of acrylic homopolymers can proceed by different routes depending on

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J. M. G. Cowie, I. J. McEwen and R. McIntyre

the nature of the α-substituent. PMMA unzips to yield virtually 100% monomer, whereas polymethylacrylate, PMA, yields only a trace of monomer. Monomer evolution from PMMA in PAN blends was retarded [Grassie et al., 1979] whereas PVAc had a destabilizing effect (due to evolved acetic acid) [Jamieson and McNeill, 1976]. More recently PMMA blends with PC have been studied [Goh, 1989; Rincon and McNeill, 1987]. These blends were found to be more or to be less stable than the component polymers depending on the method of preparation or their composition. Generally immiscible blends were found more stable than miscible ones. Blends with cellulose have attracted some attention [Flaque and Montserrat, 1991]. The degradation characteristics of blends with PAN were found to be superimposable with those of the pure components, however grafted products showed reduced thermal stability [Nishioka et al., 1992]. Somewhat similar results were obtained with cellulose graft/blend systems with PMA, PMMA and poly(2-hydroxyethyl methacrylate) [Nishioka and Yoshida, 1992] and could be correlated with the degree of compatibility. In one case the thermal stability of blends containing grafted cellulose was 100°C less than those without grafted product [Nishioka et al., 1993]. Bromine is an effective flame retardant and bromine-containing blends poly(di-bromo-propyl acrylate) with PMMA and PMA) have been studied with this in mind [Grassie et al., 1987; Diab, 1986]. Though the degradation products are those expected from the individual components it has recently been found that high temperature pyrolysis of blends (600°C) containing bromine flame retardants can generate detectable (ppm) amounts of para-dioxins [Luijk and Govers, 1992]. PMMA is compatible with poly(ethylene glycol), PEG, through strong polar interactions, and is found to protect the latter against oxidation by a mechanism which appears to promote increased PMMA chain scission in proportion to the amount of polyether in the blend [Makhija et al., 1992]. Even stronger specific H-bonded interactions exist between poly(acrylic acid) or poly(methacrylic acid) and poly(vinyl alcohol), and both well-

defined complexes as well as one-phase blends may be prepared. The polyacids degrade via mechanisms which involve initial anhydride formation and both blends and complexes produce similar end products, although the degradation temperatures of the complexes are marginally lower due to the absence of crystallinity [Zhang et al., 1992]. Polymer blends from the structurally-related monoesters of itaconic acid have also been investigated [Radice and Gargallo, 1991]. 14.4.1.4 Polyolefins The accelerated degradation of polyolefinic wrapping materials in the environment has been achieved by using blends with natural products such as starch [Pettijohn, 1992]. However relatively little has appeared on the thermal responses of polyolefin-based blends, although their radiation chemistry (below) is well enough represented and combinations of low and linear low density PE have been examined in these terms as to their resistance to photooxidation [Daro et al., 1990; Trojan et al., 1990]. Data on the kinetics of thermal degradation of this same PE blend combination have also been obtained [Bhardwaj et al., 1988] and the stability of PE blends during heat processing and short term photooxidation has been noted [Ye, 1993]. From a study of the behavior of injection molding castings it is suggested that a simple summation of individual homo-polyolefin thermograms can be used to predict the degradative weight loss behavior of their blends [Wright et al., 1989]. The degradation of ethylene-propylene copolymers, EPDM, blended with butyl rubber has been studied [Deuri et al., 1992; 1993] in order to evaluate tire inner tube performance. EPDM was found to inhibit, advantageously, rubber degradation via main chain scission. The mechanical properties (tensile strength, modulus) and electrical properties (resistivity, power factor) of EPDM blends with polypropylene, PP, have been shown to deteriorate during thermal aging [Roy et al., 1990]. Interestingly, cross-linked samples were more susceptible to property loss.

Aging and Degradation of Polymer Blends

Ethylene-propylene copolymers are useful models for homopolymer blends in that they also are biphasic, and their thermo-oxidative response [Singh et al., 1993] indicates that degradation occurs mainly in the crystalline PP phase and not in the amorphous. Ethylene/vinyl acetate copolymer-rubber blends have also been studied [Koshy et al., 1992]. Rheological anomalies in maleic anhydridePE graft blends with polyamide-6, PA-6, have been noted. These were explained on the basis of chemical reactions occurring at the interface of the biphasic melt [Kim et al., 1991]. Chlorinated PE has been found to be an effective modifier for polyurethane systems where somewhat surprisingly up to 3% of the added chlorinated material improved the thermal resistance [Koscielecka, 1993]. It was speculated that the stabilization derives from the specific interaction of chloroethylene groups with urethane carbonyls at the interface, which stabilizes the relatively labile urethane link. 14.4.1.5 Miscellaneous Blends Structures involving the ether linkage are often employed as blend components. The thermal stability of polyoxymethylene, POM, blends with polyurethane have been studied and these were found to have higher activation energies for decomposition than either of the components [Kumar et al., 1993]. The opposite effect occurred in miscible PEG blends with Phenoxy (the polyhydroxy ether of bisphenol-A); although the degradation process was affected by composition, the addition of the Phenoxy had a negative effect on blend stability [Iriarte et al., 1989]. PEEK showed extreme thermal stability, but when blended with a liquid crystalline poly(hydroxybenzoic/naphthoic acid) copolymer, LCP, there is a measurable reduction in this property [Mehta and Isayev, 1991]. The reversible nature of ester link formation implies that blends of different polyesters should undergo transesterification, or “scrambling” reactions resulting in interchange of repeat units between blend components. Blends of Phenoxy with polyethyleneterephthalate, PET,

1003

and polybutyleneterephthalate, PBT, have been studied in this regard [Eguiazabal and Iruin, 1987; Eguiazabal et al., 1986]. Such interchange reactions were reported to lead to the compatibilization of polyester/Phenoxy blends at temperatures of ca. 300°C [Kodama, 1991]. PC blends also undergo interchange during thermal treatment [Montuado and Puglisi, 1992] at as low a temperature as 160°C, according to one infrared study [Yang et al., 1992]. PC - PBT blends have been studied by direct pyrolysis in a mass spectrometer [Montuado et al., 1993]. Here exchange reactions were found to occur below 300°C and it was concluded that the overall thermal reaction evolved towards the formation of the most thermally stable polymer. This contrasts with an earlier report of a PC-PET blend in which little, if any, interchange was detected under melt processing conditions [Murff et al., 1984]. The extensive reprocessing (five cycles at 320°C) of a commercial blend of PC with polyphenylsulfone, PSF, indicated that degradation was mainly thermal and molecular weight reduction was faster in the blend than for the separate components [Sanchez et al., 1992]. In another sulfur-containing blend, poly(phenylene sulfide), PPS, with a polyarylate (bisphenol A/phthalic acid copolymer), PAV, the evolution of SO2, H2S and H2O from PPS caused accelerated hydrolytic degradation and molecular weight reduction of the second component [Golovoy, et al., 1989]. Imide blends are usually cross-linked and inherently stable structures. The additivity rule seems to be appropriate for blends with Novolak resins [Kundu et al., 1986] and the addition of a phosphine oxide to bismaleimide lowered the curing temperature without sacrifice of thermal properties [Varma and Mittal, 1989]. 14.4.2

Radiation Degradation of Polymer Blends

In accordance with a general rule vinyl polymers substituted with hydrogen in the α-position should crosslink on irradiation, and any other substituent should lead to chain degradation [Miller et al., 1954]. This is the case for most systems, exceptions occur when side chains containing long methylene

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J. M. G. Cowie, I. J. McEwen and R. McIntyre

sequences, or halogen atoms, are present. PMMA exemplifies the second type of structure and its sensitivity to radiation is significantly enhanced in poly(epichlorohydrin) blends but is reduced with styrene-acrylonitrile copolymers, SAN, when subjected to low dose radiolysis [Jubinsky et al., 1987]. A detailed study of the latter system showed the responses of both components to be altered by blending [Daro et al., 1989]. The photodegradative behavior of PMMA blends with PS [Torikai et al., 1988] and with styrene-methyl methacrylate copolymer [Torikai et al., 1992] has also been studied. Although the rate of formation of oxygenated products increased with the PMMA content in these blends, no interaction at the phase boundaries could be observed. PE-based blends have received a degree of attention, for example with EPDM copolymers [Abdelaziz et al., 1992] and with ethylenevinylacetate copolymers [Zhang et al., 1990]. The degradative radiolysis of PE blended with butadiene-styrene copolymer shows that oxidation occurs mainly in the diene component, but that effective stabilization is still possible using tin compounds [David and Zabeau, 1985]. A more detailed study, on PE blends with styreneisoprene-styrene triblock copolymer, revealed that PE degradation is initiated by the rapidly oxidized unsaturated block. This could also be obtained by blending with low molecular weight photoinitiators instead of unsaturated copolymer [David et al., 1993]. Ethylene-vinyl acetate copolymers have been used to enhance the thermal, radiation and ozone resistance of natural rubber, although to some extent this is dependent on the type of cross-linking system used [Koshy et al., 1992]. The radiative-oxidative stability of low density PE, LDPE, with PP blends were found to be more stable than pure PP [Gorelik et al., 1992]. This was explained as due to a decrease in crystallite size and possible interfacial crosslinking. The thermal stability after irradiation was considerably impaired in comparison with untreated material [Minkova et al., 1992]. This was due to the presence of free radicals within the PP component (which itself is significantly more susceptible to thermal degradation after irradiation than LDPE).

Blends of different grades of PE, are used as insulators, and the effects of electrical discharge degradation have been of interest. In blends of low molecular weight PE with both linear and branched high molecular weight PE, the time to breakdown decreased with increasing content of low molecular weight material [Gustafsson et al., 1991]. The level of crystallinity is thought to be of importance. Irradiation offers more than just overall crosslinking or degradation. By utilizing the specific sensitivities of blend components a variety of effects such as matrix stabilization, interfacial reactions, changes in rheology and mechanical properties, can be achieved. These areas have been discussed in review articles [van Gisbergen, 1991; van Gisbergen and Overbergh, 1992]. An interesting application of the understanding of radiation effects in blends is radiation absorbing melanin hydrogels for ocular devices [Chirila et al., 1992]. More detailed discussion on the radiation degradation of polymer blends is presented in Chapter 11. Use of High-Energy Radiation in Polymer Blends Technology. 14.4.3

Mechanical Degradation

The adverse environments experienced during high temperature processing are exemplified by studies on PS blends with PIB [La Mantia and Nocilla, 1987] and with PE [Komova et al., 1991; 1992], of PE/PP blends [La Mantia et al., 1985], and of poly(hydroxy acids) blends with poly(hydroxybutyrates) [Gogolewski et al., 1993]. The mechano-reactions in the later PS-PE system were found to be extremely complex; possible formation of block copolymer was reported, along with increases in PS molecular weight. Such reactions are, however, put to constructive ends in the developing area of reactive processing, either by making use of mechano-radicals or by employing suitably functionalized blend constituents [Chang and Hwu, 1991]. Graft and exchange reactions can be used to modify compatibility, morphology and interfacial parameters [Braun et al., 1991; Geuskins, 1990; Okamoto

Aging and Degradation of Polymer Blends

and Inoue, 1993]. In many cases compatibilization of the blend components is desirable [Ramesh and De, 1993] and recent developments in the chemistry and techniques involved have been reviewed. Basically compatibilization can be obtained through direct reaction of functionalized blend components, through reactive (co)polymers added separately, or through low molecular weight compounds reactively promoting copolymer formation or crosslinking [Xanthos and Dagli, 1991]. The kinetics of melt mechanical degradation have been discussed along with the kinetics of initiation under reactive processing conditions [Goldberg and Zaikov, 1987]. The maximum initiation rate is heating-rate dependent and for optimum results this must accord with the material residence time. [Turcsanyi, 1993]. A theoretical model on the effects of thermal processing of blends was proposed by Loev et al. [1989].

14.5

Weathering of Polymer Blends

14.5.1

Predicting Durability in Polymeric Blends

Probably the two most common uses of polymeric blends are thermoplastic molded objects such as automotive components and polymeric coatings such as paint. A prerequisite for all these systems intended for outside use is that they should withstand the weather. Manufacturers of such products have facilities available to them for natural exposure testing of their products, however, this is a very expensive, time consuming activity which extends product development cycles considerably. The drive towards more durable systems which are produced with shorter development cycles has over the years resulted in a vast range of artificial testing procedures being offered to the manufacturers all of which claim to give “good” correlation with natural exposure. In fact the performance of simple polymeric blends is very difficult to predict by any artificial, accelerated test procedure. There are so many factors that affect the life span of commercial polymeric materials that it is hardly surprising that the durability performance is difficult to assess. For commercial systems the intrinsic photo

1005

and chemical stability of the polymers is only one of many important parameters that have to be considered by the formulator. Even the application conditions such as temperature, humidity, and substrate preparation can adversely affect the durability performance of a coating system considerably. Other well known factors which are dependent on the composite macro structure can be equally important such as permeability, modulus, extensibility and toughness. Of crucial importance for commercial systems can be the additives necessary to achieve the rheological properties for application and final appearance, such as pigments, surfactants, surface flow promoters, anti-foaming and degassing agents. The more durable the intended product the more effort goes into the design of the additives to achieve the desired durability. Given the number of parameters which can affect the final durability of a polymeric material it should not be surprising that simply accelerating the process of UV exposure, by increasing the amount of incident energy, does not give a good correlation with natural exposure. Although no longer acceptable to discerning customers, this practice is nevertheless still widely used internally, in the attempt to reduce development cycles. 14.5.2

Physical Measurement Techniques

The standard method of assessing the performance of polymers exposed to UV radiation are gloss loss, chalking, color retention, cracking and crazing, and dirt retention. Dirt retention is not strictly a measure of degradation, but can obviously be very important e.g. for architectural and decorative coatings. The remainder, along with several others, measure physical changes due to weathering. Probably the most important of the above is gloss loss, as a pigmented film degrades, the surface gradually roughens due to the body of the binder being lost, this is apparent as a loss of gloss. The disappearance of the binder eventually leads to the pigment being exposed and this remains on the surface as non bound powder particles, a process known as chalking. Chalking and gloss measurements go hand in hand with the gloss loss process

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J. M. G. Cowie, I. J. McEwen and R. McIntyre

preceding the chalking event. Gloss is assessed by monitoring the reflection from the surface of a visible beam of light in relation to a standard black glass surface. Chalking is measured by performing a tape pull off test. A very common specification is to indicate the minimum time a coating will maintain a gloss level above 50% of the original value, for a particular exposure method. 14.5.3

Artificial Exposure Methods

14.5.3.1 Carbon ARC Sources The enclosed carbon arc lamp was the original artificial weathering source, being first employed in 1918 to evaluate textile performance. Its spectral output has little comparison with natural sunlight, the main radiation originating from three high intensity peaks between 350-450 nm. Furthermore, it produces a significant amount below 290 nm (the approximate solar cut off) although this can be removed by the use of suitable filters. Test procedures specifying the use of enclosed carbon arc are becoming fewer in number but it is still used for some American (AATCC) and Japanese (JIS) standards. The sunshine carbon arc is also being progressively phased out of international test specifications but is still widely used in industry for company specifications. It’s output is closer to that of sunlight than the enclosed version, but it still has significant differences, in particular being more intense in the 350-450 nm region and less intense below 350 nm. 14.5.3.2 Xenon ARC Sources This source provides much better simulation of sunlight than carbon arc sources, which explains its use in testing a wide range of materials. Filters are used to remove the radiation that is produced at shorter wavelengths than the solar cut off. Apart from close simulation of the solar spectrum, other advantages of commercially available xenon arc testers include automatic control of the light intensity, temperature, light/dark periods and humidity, giving rise to consistent

and reproducible test methods. Furthermore, the facility to incorporate aggressive atmospheric pollutants into the system make these instruments very useful to mimic certain outdoor conditions. 14.5.3.3 Fluorescent Tube Lamps These sources are much cheaper to operate than a xenon arc and do not produce any unwanted heat. The tubes are available with different spectral outputs, the most common being UVb and UVa bulbs. The UVb spectrum has a peak intensity around 313 nm, with a high intensity at wavelengths below the solar cut off (i.e., 270-290 nm). The consequence of this is that they provide strong acceleration, but they are also likely to induce unnatural chemical reactions. UVa bulbs provide a much closer reproduction of the solar spectrum up to around 350 nm and can therefore be expected to produce a much better correlation with natural exposure. As only negligible heating is produced with these lamps, a separate heating source is necessary which can be used to achieve reasonable temperature control, humidity and light and dark cycles are also normally controlled with these pieces of equipment. The correlation of these artificial test results with natural weathering is often poor particularly for colored materials due to differences in the visible region, although it is generally differences in the UV region which are most likely to cause deviation. 14.5.3.4 Natural Exposure There are a number of exposure sites around the world each with a different climates, which can be used for the purpose of predicting the exterior durability of materials. The site chosen should be the one, which is most representative of the general environment of interest, however, it is more common to choose sites of extreme climates. In this way natural exposure is achieved, even though the UV, temperature and humidity levels may be vastly different from those met in service. Southern Florida with its high levels of UV radiation, humidity and temperature provides

Aging and Degradation of Polymer Blends

a sub-tropical environment which is itself an acceleration of the central European climate (an acceleration of approx. four times) and has been the choice location specified by the automotive industry for many years. 14.5.3.5 EMMAQUA The weathering of materials can be further accelerated by a more efficient use of outdoor exposure. An example of this is the widespread use of EMMA (Equatorial Mounts with Mirrors for Acceleration) and EMMAQUA (EMMA with a water spray cycle) These machines have ten highly polished moveable aluminum mirrors which enhance the rays of the sun focused upon a test sample. Although the total amount of radiation reaching the surface of the sample is about 9 times that of natural exposure the deterioration rate is found to be, on average, only 5-6 times faster. This is a consequence of the fact that the aluminum mirrors reflect the more damaging UV components of sunlight less efficiently than the visible and IR components. As a result of the high radiation intensities focused up on the sample, temperatures of up to 150°C may develop, hence the application of cool air is necessary. A drawback of the technique is that continuous clear sky is necessary for efficient use, hence the Arizona test site with over 4000 hrs of sunshine per annum is ideal. The major assumption of EMMAQA is that gloss loss is directly proportional to the total cumulative exposure, thus EMMAQA results are usually represented as gloss loss versus total energy of exposure, the total energy normally being converted to equivalent Florida years (1 Florida year = 308 MJm-2 radiation). It has been shown that this is not strictly true and that gloss loss is in fact proportional to both the total energy and the flux. In other words a seasonal variation occurs in which samples will exhibit quite different gloss loss characteristics depending on the time of year of initial exposure.

14.5.4

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Detection Methods

14.5.4.1 Electron Spin Resonance (ESR) For many years ESR has been perceived as being a specialist technique used by physicists and physical chemists. However, in more recent times the technique has been employed increasingly by the biochemical, pharmaceutical and polymer industries. A particular example in the polymer and coatings industries has been the use of ESR to monitor the UV durability of polymeric coatings for external exposure applications. The method was pioneered by Bauer et al. [1986] who claimed that, by considering the appropriate steady state equations, quantitative estimations of the UV durability of the coating is possible [Bauer et al., 1988]. Since then, the method has gained more widespread use in the polymer industry. In particular, basing on work by Kampf et al. [1991], Bayer A.-G. in Leverkusen have claimed “an ultra fast method for evaluating the weather resistance of polymers”. Existing test methods for assessing the UV durability of polymeric systems rely on the use of test centers in areas of high natural UV irradiation such as Florida and Arizona. The usual experiment involves sending panels to a test site where they are exposed for 12-60 months, during which the coatings are continuously evaluated for gloss and color retention. The cost of this process can be considerable and therefore it is usual to carry out an initial screening program using weatherometers, employing QUA and or QUB radiation sources. This still involves exposure for periods up to 3000 hrs that extends the product development cycle considerably. The main aim of the ESR method is not to replace the natural exposure tests but to provide a more rapid and accurate screening procedure. The method relies on the fact that during the irradiation of polymeric coatings radicals are produced at very short time scales. For example, exposure of commercial epoxy coatings to a Xe-arc lamp for 15 min (N.B., the number of Langleys delivered is equivalent to 2-3 that of natural sunlight) will produce phenoxy radicals which are very short lived, decaying in minutes

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[Dolby, 1988]. Similarly, exposure of commercial polyester resins can be shown to produce acyl radicals, albeit over somewhat longer time scales and at liquid nitrogen temperatures [Butcher, 1992]. These particular radicals have a short lifetime at room temperature and it is therefore necessary to either irradiate at low temperature or to extend the lifetime of the radical by the inclusion of a spin trapping material to ensure that the radical is sufficiently stable to collect accurate data. Irradiation at low temperature has been successfully used by [Sommer et al., 1991] to investigate 20 different clear 2-pack polyurethane coatings, the coatings were exposed for a period of several hours using a high-pressure mercury lamp at -133°C. Under these conditions the radicals produced were from the primary photolytic process and in this case a good correlation with natural weathering was claimed, enabling both the effectiveness of different stabilizers and different curing chemistries to be assessed. The more elaborate procedures of either incorporating spin trapping agents into the polymer prior to producing the coating or infusing the spin trap into the coating just prior to the irradiation have also been attempted with some success [Howell, 1992; Holmes and McIntyre, 1990]. The advantage of this procedure is that the experiments can be carried out at more realistic temperatures. The rationale for this procedure comes from the fact that Hindered Amine Light Stabilizers, HALS, function by a radical trapping mechanism. The generally accepted mechanism involves initial oxidation to the corresponding nitroxide. Nitroxides are powerful radical scavengers and act as chain breaking antioxidants via the cycle: RNO• + R• ⇒ RNOR RNOR + RO2• ⇒ ROOR + RNO •

Figure 14.11.

Because of the cyclic regeneration of the nitroxides, each has the potential to trap a large number of radicals. The decay in the nitroxide signal, on irradiation, should therefore be some measure of the number of radicals involved in the degradation process. It was shown that measurement of the nitroxide decay rate, as a function of the initial starting concentration of nitroxide in the coating, can be used to determine the Photo Initiation Rate (PIR) of free radicals in the coating [Bauer et al., 1988]. This work demonstrated that acrylic-melamines with low PIRs invariably outperform those with higher PIRs. Nippon Oil and Fats have demonstrated that a similar technique using hindered phenols can be used to predict the durability of a series of silicon modified polyesters with irradiation time of only 20 min [Okamoto et al., 1990]. A number of independent investigators have successfully used ESR as a fast, predictive method for assessing the UV durability of polymeric materials, however, these investigations have been restricted to comparisons with a particular polymer chemistry. The real test for the method would be to compare the UV durabilities across a range of polymer chemistries including different polymer blends. The first investigation involves a comparison of different chemistries. Standard commercial coatings with well known UV durability performance were evaluated using the nitroxide decay procedure [Holmes and McIntyre, 1990], in which the coating samples were doped with a known amount of the volatile nitroxide, 4-hydroxy-2,2,6,6-tetramethylpiperidinyl-N-oxyl. Doping was accomplished by placing the samples, coated onto a 5 mm quartz disc, in sealed containers along with 10 µL of methylenechloride and 0.5 µL of methanol (used as a transfer solvent) containing a known amount of the nitroxide. The containers were kept at 55°C overnight to allow solvent infusion and hence polymer swelling. The samples were then removed, and warmed at 60°C for 5 hr to remove the solvent leaving behind the nitroxide. For each sample the initial ESR spectrum was recorded showing the characteristic features for an immobilized nitroxide free radical. Figure 14.12 shows the spectra obtained for the individual

Aging and Degradation of Polymer Blends

components. It is evident that the degree to which the nitroxide is immobilized looks similar for both components of the blend, c.f. the nitroxide undergoing free rotation. Integration of these signals allows quantitative estimates of the amount of nitroxide in the coating at anytime. The concentration of nitroxide could therefore be determined accurately as a function of exposure time. Each sample was then exposed to UV irradiation in a Hanau Suntest sun center (Xe arc lamp) for a controlled period of time, the process was repeated until the nitroxide signal had decayed by >20% of its original value. The whole procedure was repeated for several concentrations of nitroxide for each coating. For each coating the decay rate

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of nitroxide was calculated. These rates, which are the photo-initiation rates for a given concentration of nitroxide, when plotted against the actual nitroxide concentration, and extrapolated to zero, give the photo-initiation rate of the polymer (see Figure 14.13). A comparison of these data with a widely used industry standard QUV (Quick Ultra Violet) shows good correlation, Figure 14.14. It should of course be mentioned that these three coatings are vastly different in their UV durability and therefore achieving a good correlation could be considered to be an easy test of the ESR method. For this approach to be a valuable fast method, it ought to be capable of discriminating between coatings

Figure 14.12. ESR spectra for nitroxide probes in different environments.

Figure 14.13. Comparison of photo-initiation rates for varying nitroxide contents in epoxy, polyester and lumiflon coatings.

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Figure 14.14. Percentage loss of gloss for epoxy, polyester, and Lumiflon coats as a function of the exposure in Quick Ultra Violet (QUV) standard hours.

Figure 14.15. Percentage gloss retention as a function of exposure for blends of polyester and acrylic copolymers.

with only 10-20% difference in their actual external exposure performance, for different polymer chemistries, and also for materialsconsisting of blends of different chemistries. To make an assessment of this method a series of coatings were produced based on blends of polyester and acrylic copolymers. The actual QUV data (see Figure 14.15) show a gradual transition from the less durable polyester to the more durable acrylic. The photo-initiation rates were calculated from the nitroxide decay curves,

in a the same way as described earlier, and are shown in Figure 14.16. The relative PIRs for the polyester and the acrylic are shown to be in the correct order compared with the gloss loss data, however the PIR values for the intermediate blends do not show such a good correlation. Although there are only a limited number of data points it does appear that the correlation is poor for the intermediate blends. There are a number of potential reasons for this, the most likely being that the coatings

Aging and Degradation of Polymer Blends

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Figure 14.16. Variation of photo-initiation rate with nitroxide content of various polyester-acrylic copolymer blends.

are at least two phase systems which leads to different rates of infusion of the radical scavenger and hence different concentrations in the two phases. A further complication is the fact that the lower surface energy component is known to undergo some degree of interfacial segregation and thus the compositions at the air/coating interface will be enriched in the acrylic component and furthermore the enrichment may well be a nonlinear function of the initial starting composition [McIntyre and Dolby, 1990]. From the above work our data suggests that the PIR values calculated for different chemistries do correspond to the natural weatherability of the polymer systems. This would have to be confirmed with more extensive studies with chemistries that have UV durability performance closer together. However, in the case of polymeric blends the PIR values show poor correlation with the known outdoor performance and this is always likely to be the case where the blends produce complex heterogeneous morphologies. 14.5.4.2 Fourier Transform Infrared Spectroscopy (FTIR) Increasing use is being made of Attenuated Total Reflectance Fourier Transform Infrared Spectroscopy, FTIR-ATR, as a surface sensitive method which is capable of giving quantitative

information in the near surface region (typically 0.2-3 µm). This is the most suitable FTIR method for attempting to assess the changes that take place after relatively short periods of UV exposure. Over the past five years Photoacoustic FTIR has been developed for depth profiling coatings, over the typical mirror velocities available on conventional FTIR machines the penetration depth ranges from about 2 to 50 µm for 90% of the polymeric material used today [Dolby and McIntyre, 1991]. The changes observable after prolonged QUV are restricted to the highest mirror velocities which are the most surface sensitive, and the problems of poor signal to noise at high mirror velocities makes ATR a much better choice for these studies. Typical results for the natural exposure of an epoxy-polyamine coating, a particularly poor coating with respect to UV durability, are shown in Figure 14.17. Even after several days of natural exposure, in a northern European climate, oxidation products are clearly visible in the top 2-3 µm layer. It should be pointed out that quantitative assessment of these changes is possible well before any measurable gloss loss can be recorded [Davies, 1993]. The spectra shown in Figure 14.17 were produced using a ZnSe ATR crystal. At least two discrete bands centered around 1730 cm-1 and 1670 cm-1 and a broad band lifting the spectral intensity over the range 1000-1300 cm-1 appear with exposure.

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Figure 14.17. Attenuated reflectance spectra of epoxypolyamine coatings under northern European climate exposure conditions for 0, 8, and 23 days.

The band at 1730 cm-1 is most likely due to ester formation, the band at 1670 cm-1 is less easily assigned but given that the phenoxy radicals are present in the early stages of photo-degradation (based on ESR work mentioned above) then quinone type species are likely candidates or some other, unsaturated carbonyl moiety. The broad band in the region 1000-1300 cm-1 is likely to be due to various C-O stretching vibrations and is typically observed for acid peroxides [Davies, 1993]. For more UV durable blends accelerated exposure was used to bring about surface oxidation. Typical ATR spectra are shown in Figure 14.18

Figure 14.18. ATR spectra for polyester-acrylic blends under quick ultra violet exposure.

for a polyester/acrylic blend. After 3-4 days QUV oxidation products can be detected in the near surface region of a 50/50 blend, after 6 days the changes are sufficiently large for quantitative estimates of the amount produced in the surface region. It is known from ESR work that the acyl radical is produced upon irradiation of this blend at low temperatures. In this case the formation of acid and anhydride, peaks c and d respectively, is suggested. Although absorptions at 1780-1800 cm-1 could be attributed to aryl and, unsaturated acid peroxides, our hydroperoxide titrations suggest that peroxides are not present at that level and therefore anhydride seems more likely. To accompany the above changes there is also a strong band centered around 1175 cm-1, which is likely to be additional C-O stretching vibrations. We believe this behavior to be largely due to the degradation of the polyester [Johnson, 1993]. Polyurethane-acrylic blends are often used as UV durable coatings, and the ATR spectra of such a coating after accelerated weathering using QUV are shown in Figure 14.19. Over an extended period there are some very reproducible changes, which can be fairly unambiguously assigned. Peaks a and b, centered around 1640 and 1520 cm-1, are characteristic of the amide I and II bands of urea, in this case a polyurea which tends to be formed

Figure 14.19. ATR spectra for polyurethane-acrylic blends under quick ultra violet exposure.

Aging and Degradation of Polymer Blends

to a greater or lesser extent depending on the temperature and humidity on application. There are clearly other exposure dependent changes taking place. At around 1245 and 1040 cm-1 two bands appear which decrease with exposure, and may well be the initial phase of acrylic degradation [Johnson, 1993(a)]. 14.5.4.3 X-Ray Photoelectron Spectroscopy, XPS The use of XPS for the assessment of the UV durability of polymeric coatings has met with little success, the reason for this is, ironically, associated with the sensitivity of the XPS technique. XPS is surface sensitive, and the rate of detection of photoelectrons emitted from a depth of greater than 5 nm into the polymer material is too low to give a useful signal. This causes a problem with the reproducibility of the test conditions, as the slightest loss of degradation products from the surface can be detected and this makes the job of obtaining reproducible quantita-

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tive information difficult. Furthermore, it cannot be assumed a priori that the rate of degradation of the top 5 nm is comparable with the rate of degradation of the material in the near surface region (i.e., 5 µm) because of the natural surface dipole which often leads to order. There is also the possibility of surface contaminants that can act as chromophores, which dramatically alter the rate of degradation of a coating system. For the purpose of investigating the photodegradation of polymer blends, XPS can play a valuable role in determining the composition of the blend at the surface. Normally, as the result of interfacial segregation, the lower surface energy component can be found in excess quantities at the surface and this can have a profound effect on the UV stability. A case in point is that of the acrylic-urethane blend discussed above, and Figure 14.20 shows the XPS spectra of the above acrylic/urethane coating before and after QUV exposure [Johnson, 1993(b)]. The coatings were prepared in an ultra clean environment and

Figure 14.20. XPS spectra for unstabilized polyurethane-acrylic blends for zero (upper) and 43 days (lower) quick ultra violet exposure.

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examined by XPS, prior to and after exposure. Prior to exposure there is no evidence of nitrogen in the top 5 nm despite the fact that the average composition should contain about 5 atomic % which is well within the detection limits of XPS. Nitrogen does, however, gradually begin to appear over an extended period of exposure of up to 60 days, suggesting that the surface layer of the acrylic is degrading preferentially. This view is supported by the FTIR data presented above and further evidence comes from the fact that the addition of a small amount of a commercially available HALS causes a delay in the appearance of the nitrogen signal, on QUV exposure, which can best be explained by stabilization of the acrylic in the near surface region. Acrylic and polyester blends are too similar to determine any specific changes in the chemistry, the only real difference before and after exposure is an increase in the oxygen to carbon ratio, and this is not a large effect. Also, there is no linear increase with irradiation time, in fact the O/C ratio appears to oscillate over a period of 60 days suggesting that surface layers are gradually becoming friable and falling off, exposing more new surface with a lower C/O ratio [Johnson, 1993(b)]. 14.5.4.4 Iodometric Redox Measurements Hydroperoxide determination by iodometric titrations are quite common in the literature, the essential differences in the methods lies in the end point determination [Mielewski et al., 1989]. The propagation mechanism in the standard UV degradation process involves the production of hydroperoxides by the abstraction of a hydrogen atom from the polymer by the peroxy radical. Assuming the reaction goes to completion, which it does in the presence of a small amount of ammonium molybdate, an equilibrium will be established which should behave according to the simple Nernst equation [Johnson, 1993(c)]. Therefore it ought to be possible by monitoring the redox potential of the I-/I3-couple to get a relative measure of the hydroperoxide in the coating:

ROO• + RH ⇒ R• + ROOH H+ ROOH + 3I- ⇒ ROH + H2O + I3To test this hypothesis an epoxy (based on Epikote 1004) was doped with known amounts of t-butyl hydroperoxide, duplicate panels of the conventional epoxy containing 10-2, 10-3 and 10-4 mol/mL of the peroxide were produced and the panels were allowed to cure then stored under N2 for 4-days before analysis to remove oxygen. The coatings were removed from the panel, cryoground (using 6700 Glen Creston Freezer Mill), then placed in CH2Cl2 for 15 min to swell the polymer, followed by refluxing with the KI solution containing isopropyl alcohol and acetic acid. The potential of the redox couple measured at a platinum electrode against a standard calomel reference electrode was recorded. The expected Nernstian response was observed with approximately 30 mV difference between the three coatings, but the calculations showed that only 2% of the hydroperoxide was detected. Nonetheless, the method was shown to be capable of distinguishing between different chemistries with respect to their UV durability performance. Unfortunately, the method was not capable of correctly ranking the UV durability performance of a series polyester and acrylic blends.

14.6

Summary, Conclusions and Outlook

The study of blends by physical aging is still largely unexplored and requires close attention both from the experimental and theoretical standpoint. One major advance would be the ability to crosscorrelate mechanical and thermodynamic physical aging data, but while there is a theory that links the two this is imperfect and requires refinement. One problem is that the aging processes, when followed by these different methods, proceed at different rates. If conditions could be established that led to these proceeding at equal rates then this cross-correlation would be possible.

Aging and Degradation of Polymer Blends

The theoretical approaches to describing the phenomenon in general are also in need of refinement. There is some controversy as to the best theoretical approach to use to explain the phenomenon, but effort would be better spent on elaborating ways of predicting long-term aging from short-term thermally accelerated experiments than in arguing over the finer points of the phenomenological models. Degradation of polymers is a science of long standing and little more may be said about it at this stage but the area of blend weathering is of fundamental importance. There is a considerable amount of work ongoing in the area of theoretical prediction of the heterophase nature of blends, both from the point on statistical thermodynamics building on the original work of Langer [1973] and also statistical mechanics, building on the work of Flory and De Gennes. Statistical mechanical predictions, which have been experimentally verified, are now possible for interfacial segregation occurring in heterogeneous systems [Russell 1993]. These are essential developments if we are to have a chance of accurate prediction of the durability of polymer blends. A better understanding of the degradation pathways for specific chemistries will also have to emerge before industry will have the confidence to use fast answer methodology for polymer blends. For the foreseeable future it appears that depth profiling using methods, which give chemically specific information are the techniques most likely to succeed. This will include the techniques discussed above such as FTIR-ATR and FTIR-PAS but perhaps also more elaborate techniques can be justified since coating materials are becoming more and more UV durable and therefore conventional testing becomes more expensive. Therefore the drive for fast answer procedures becomes more financially viable hence greater expenditure on test procedures becomes possible. There are, for example, 2 and 3-dimensional ESR techniques that could be calibrated to give quantitative information on the concentration of degradation products, this is one way of dealing with the heterogeneity usually found in commercial polymeric compositions.

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Much remains to be done in all of these areas and as each aspect deals with the long-term performance of polymers and blends it would be helpful to have collaboration amongst the various research groups working on their individual areas to obtain a more coherent approach to the study of long range, material performance.

14.7

References

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J., J. Appl. Polym. Sci., 34, 2493 (1987). Struik, L. C. E., “Physical Aging in Amorphous Polymers and Other Materials”, Elsevier, Amsterdam (1978). Tamada, H., Okita, S., and Kobayashi, K., Polymer J., 25, 339 (1993). Tanaka, A., Nitta, K., Maekawa, R., Masuda, T., and Higashimura, T., Polymer J., 24, 1173 (1992). ten Brinke, G., and Grooten, R., Coll. Polym. Sci., 267, 992 (1989). ten Brinke, G., Karasz, F. E., and MacKnight, W. J., Macromolecules, 16, 1827 (1983). ten Brinke, G., Macromolecules, 23, 1225 (1990). Tomlins, P. E., Read, B. E., and Dean, G. D., Polymer, 35, 4376 (1994). Tool, A. Q., J. Amer. Ceram. Soc., 29, 240 (1946). Torikai, A., Hiraga, S., and Fueki, K., Polym. Degrad. Stabil., 37, 73 (1992). Torikai, A., Sekigawa, Y., and Fueki, K., Polym. Degrad. Stabil., 21, 43 (1988). Trojan, M., Daro, A., Jacobs, R., and David, C., Polym. Degrad. Stabil., 28, 275 (1990). Tsitsilianis, C., and Bokaris, E. P., Polymer Bull., 30, 609 (1993). Tsitsilianis, C., and Staikos, G., Macromolecules, 25, 910 (1992). Turcsanyi, B., Polym. Bull. 30, 297 (1993). van Gisbergen, J., and Overbergh, N., Prog. Polym. Process., 3, 51 (1992). van Gisbergen, J., Kunstst. Rubber, 44, 20 (1991). Varma, I. K., and Mittal, H. P., Thermochim. Acta, 144, 339 (1989). Varughese, K. T., Kautschuk Gummi Kunststoffe, 41, 1114 (1988). Venditti, R. A., and Gillham, J. K., J. Appl. Polym. Sci., 45, 501 (1992). Victor, J. G., and Torkelson, J. M., Macromolecules, 21, 3490 (1988). Vittoria, V., J. Macromol. Sci.-Phys., B28, 97 (1989). Vleeshouwers, S., Jamieson, A. M., and Simha, R., Polym. Eng. Sci., 29, 662 (1989). Vleeshouwers, S., Kluin, J.-E., McGervey, J. D., Jamieson, A. M., and Simha, R., J. Polym. Sci: Part B: Polym. Phys., 30, 1429 (1992). Wang, C. H., and Zhang, X. Q., J. Polym. Sci: Part B: Polym. Phys., 32, 569 (1994). Williams, G., and Watts, D. C., Trans. Farad. Soc., 66, 80 (1970).

Aging and Degradation of Polymer Blends

Won, H. J., and Kyoung, J. K., Polym. Eng. Sci., 31, 239 (1991). Woo, L., and Cheung, Y. W., Thermochimica Acta, 166, 77 (1990). Wright, J. K., Evans, J. R. G., and Ederisinghe, M. J., J. Amer. Ceram. Soc., 72, 1822 (1989). Wypych, J., Polym. Networks Blends, 2, 53 (1992). Xanthos, M., and Dagli, S. S., Polym. Eng. Sci., 31, 929 (1991). Yang, X. M., Painter, P. C., and Coleman, M. M., Macromolecules, 25, 4996 (1992). Ye, Y. C., Polym. Degrad. Stabil., 39, 193 (1993). Yee, A. F., Bankert, R. J., Ngai, K. L., and Rendell, R. W., J. Polym. Sci: Part B: Polym. Phys., 26, 2463 (1988). Yoshida, H., and Kobayashi, Y., J. Macromol. Sci.-Phys., B21, 565 (1982). Zhang, M., Mai, K., and Zeng, H., Polymer J., 25, 541 (1993). Zhang, W. X., Liu, Y. T., and Sun, J. Z., Radiation Phys. Chem., Part C, 35, 163 (1990). Zhang, X. Q., Takegoshi, K., and Hikichi, K., Polymer, 33, 718 (1992).

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CHAPTER 15

COMMERCIAL POLYMER BLENDS

M. K. AKKAPEDDI

Honeywell Inc., EAS R & T, Morristown, USA

15.1

Abstract

In this chapter, an overview of the commercially important blends is presented with a particular emphasis on the rationale for their commercial development, the compatibilization principles, their key mechanical properties and their current applications and markets. To facilitate the discussion, the commercial polymer blends have been classified into twelve major groups depending on the type of the resin family they are based on, viz. (i) polyolefin, (ii) styrenic, (iii) vinyl, (iv) acrylic, (v) elastomeric, (vi) polyamide, (vii) polycarbonate, (viii) poly(oxymethylene), (ix) polyphenyleneether, (x) thermoplastic polyester, (xi) specialty polymers, and (xii) thermoset blends. Within each major category, the individual polymer blends of industrial significance have been described with relevant data. Since the discussion is limited only to those blends that are actually produced and used on a commercial scale, the relevant cost and performance factors that contribute to the commercial viability and success of various types of blends have been outlined. In comparing the different blends, the specific advantages of each type, as well as any potential overlap in performance with other type of blends have also been discussed. The fundamental advantage of polymer blends viz. their ability to combine cost-effectively the unique features of individual resins, is particularly illustrated in the discussion of crystalline/amorphous polymer blends, such as the polyamide and the polyester blends. Key to the success of many commercial blends, however, is in the selection of intrinsically complementing systems or in the development of effective compatibilization method. The use of reactive compatibilization techniques in commercial polymer blends has also been illustrated under the appropriate sections such as the polyamide blends. In many commercial blends, rubber toughening plays an important and integral part of the blend design. Combining high impact strength with other useful properties such as heat and solvent resistance can significantly enhance the commercial value of a blend. Hence, the nature of the impact modifiers used and the role of morphology on properties have been discussed under the appropriate cases of commercial blends. The chapter concludes with an outline of the potential trends in the commercial polymer development. L.A. Utracki (Ed.), Polymer Blends Handbook, 1023-1115. © 2003 Kluwer Academic Publishers. Printed in the Netherlands.

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15.2

M. K. Akkapeddi

Introduction

Polymer blends have gained significant commercial growth in the last two decades outpacing the growth rate of existing polymers by at least 2 to 5%. The current worldwide market volume for polymer blends and alloys is estimated to be more than 700,000 metric ton/y, with an average growth rate of 6 to 7%. Although this pace of growth slightly slackened during the ‘90’s, the demand for polymer blends is expected to be maintained due to the possibility of adjusting the cost-performance balance and tailoring the technology to make products for specific end-use applications [Utracki, 1998]. A combination of the following important factors contribute to the continued commercial interest in the polymer blends: 1. The blending of commercially available polymers is a more cost-effective method of developing a new product that meets the customer or market requirements, as opposed to developing a totally new polymer what generally involves prohibitively high research, development and capital costs. 2. Polymer blends can fill the cost-performance gaps in the existing commercial polymers. Several properties can be uniquely combined in a blend that a single resin often cannot provide. In some cases, synergistic improvements in properties such as toughness and heat resistance are achievable. 3. Polymer blending can be done at a relatively low cost using an extruder. Production of new polymers, on the other hand, requires capitalintensive plants and reactors that must operate on a reasonably large scale for reasons of economics. 4. The flexibility of extruder blending enables custom production of different blends in a wide range of production volumes. Polymerization plants are generally not as flexible and not economical for small volume production. 5. Polymer blends provide an avenue for diversifying and expanding the product line for resin producers and suppliers, without significant investment risks.

6. Blends can be formulated, optimized and commercialized generally at a much faster rate (from concept to commercialization) than new polymers, provided there are no major technical hurdles for the compatibility development between the components. 7. The development of an effective compatibilization technology, whenever needed, allows the resin supplier to establish a proprietary and competitively advantageous position. 8. Blends offer useful and economic means of upgrading recycled and off-specification polymers. The terms “polymer blends” and “polymer alloys” are defined in Chapters 1 and 2 of this Handbook. In the trade literature, they have been used interchangeably. In the context of current discussion, we will treat all of them simply as polymer blends, except specifying, where possible, the origin of the technological compatibility between the components in each type of blend. Table 15.1 lists the commercially available polymer blends according to their primary structural categories (Figure 15.1). 15.2.1

Compatibilization Mechanisms in Commercial Polymer Blends

To be useful, most commercial polymer blends are either designed or selected to have some degree of the technological compatibility between the components to resist delamination and loss in ductility. Compatibility is defined here as the ability for the polymer components to co-exist either as molecularly miscible or as morphologically distinct phases, but interfacially stabilized, without a tendency for delamination. The technological compatibility in polymer blends can result from any of the following mechanisms: 1. Thermodynamic miscibility between the components such as in the case of polystyrene (PS) and poly(2,6-dimethyl 1,4-phenylene ether) (PPE) blends. 2. Segmental miscibility between the components, even when they are phase separated, imparting a low interfacial tension and an

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Commercial Polymer Blends

Table 15.1.

Commercial polymer blends

Blend

Producer

Trade Name

Compatibilization

Blend Type

Mechanism(a) Polyolefin Blends Elastomeric Polyolefin blends PP/EPDM PP/NBR PP/EPDM/NBR PP/PBD PP/Butyl Ethylene terpolymer/PVC

Monsanto Novacor Advanced Elastomer Systems Japan Synthetic Rubber

Santoprene® Sarlink 3000® Geolast®

None (b); Dynamic Vulcanization (c) Grafting/Dynamic Vulcanization Grafting/Dynamic Vulcanization None; Dynamic Vulcanization None; Dynamic Vulcanization Partial miscibility(d)

Crystalline/ Amorphous

None

Crystalline/ Amorphous

Pax-Plus 3200®

None

Crystalline/ Amorphous

Triax 2122® Cycoloy EHA® Bayblend® Cycovin® Triax 4000®

Partial miscibility

Amorphous/ Amorphous

-

-

Triax 1000® Cycovin® Malecca K®

Grafting(R) Partial miscibility Partial miscibility

Amorphous/ Crystalline Amorphous/ Amorphous Amorphous/ Amorphous

Dynaflex P®

Monsanto Vyram® Novacor Sarlink 1000® Novacor Sarlink 2000® Advanced Elastomer Systems TPE-3000® Trefsin® DuPont Alcryn®

Thermoplastic Polyolefin blends (TPO) PP/EP or EPDM

BP Performance polymers; Bayer; D&S International; Ferro; Himont Hoechst; ISR; Mitsubishi Petrochem; Republic Polymers; A. Schulman; Teknor Apex; Tonen and others HDPE/Polyisobutylene Paxon Polymer Co. Styrenic blends ABS blends ABS/PC

ABS/PBT

ABS/PA ABS/PVC ABS/SMI

Monsanto General Electric Mobay General Electric Monsanto Daicel (Japan) Monsanto General Electric Denka (Japan)

Crystalline/ Amorphous Crystalline/ Amorphous Crystalline/ Amorphous Crystalline/ Amorphous Amorphous/ Amorphous

1026

Table 15.1.

M. K. Akkapeddi

continued

Blend

Producer Mechanism(a)

Trade Name

Compatibilization

Blend Type

Dow Vista

Rovel® Suprel®

Miscibility Partial Miscibility

Amorphous/ Amorphous Amorphous/ Amorphous

Royalite® Hycar® Denka LCS® Krynac® Vythene®

Partial Miscibility Miscibility occurs with NBR containing >25% AN None

Amorphous/ Amorphous Amorphous/ Amorphous

PVC/polyurethane

Polycast B.F. Goodrich Showa Denka Polysar, etc. D & S International

Acrylic blends PMMA/PV PMMA/PVC PMMA/acrylic core shell elastomer

DFRexham Kleerdex Rohm & Haas CYRO

Fluorex® Kydex® Plexiglas® Cyrex®

Miscibility Partial miscibility Miscibility Miscibility

Santoprene® Sarlink® Geolast® Sarlink® Trefsin® Vyram®

Dynamic Vulcanization Amorphous/ Amorphous

Butyl rubber/PP PBD/PP

Advanced Elastomer Systems Novacor Advanced Elastomer Systems Novacor Novacor Novacor

Polyamide blends PA/ABS PA/Acrylic rubber

Monsanto DuPont

Triax 1000® Zytel FN®

PA/Elastomer

DuPont

PA/ Polypropylene PA/Polyarylate

AlliedSignal EMS-American Grilon D&S Int. Unitika

SAN blends AES (SAN/AES blend) SAN/PVC Vinyl blends PVC/PMMA PVC/Nitrile rubber

Elastomeric blends EPDM/PP NBR/PP

Dynamic Vulcanization Dynamic Vulcanization Dynamic Vulcanization Dynamic Vulcanization

Amorphous/ Amorphous

Crystalline/ Amorphous Amorphous/ Amorphous Amorphous/ Amorphous

Amorphous/ Amorphous Amorphous/ Amorphous Amorphous/ Amorphous Amorphous/ Amorphous

Graft-coupling(e) Grafting/controlled crosslinking Zytel ST® Graft-coupling Zytel Z408® Polar interactions(f) ® Capron 8350, 8351 A28®, BT40X®

Crystalline/ Amorphous Crystalline/ Amorphous

Dexlon®, Dexpro® X-9®

Crystalline/ Crystalline Crystalline/ Amorphous

Grafting

Crystalline/ Amorphous

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Commercial Polymer Blends

Table 15.1.

continued

Blend

Producer

Trade Name

Compatibilization

Blend Type

Mechanism(a) PPE blends PPE/HIPS PPE/PS PPE/PA-66 (and 6)

PPE/PBT

Polycarbonate blends ABS/PC AES/PC ASA/PC PC/Styrene-Acrylic PC/SMI PBT/PC

PET/PC

PA/PC Polyester blends PBT/PC PBT/PET PBT/ASA PBT/Acrylic PBT/Elastomer PBT/SMI Copolyester/PBT PCTG/PC

General Electric AlliedSignal BASF General Electric

Noryl® Gepax® Noryl GTX® Dimension® Ultranyl® Gemax®

Dow Dow BASF Novacor Chemicals CYRO Denka (Japan) General Electric Dow Mobay Mobay General Electric Polysar Eastman Dow Chemical Dexter Plastics

Pulse® Rovel 747® Terblend S® SD 9000® Cyrex® Malecca® Xenoy® Sabre 1628® Makroblend® Makroblend UT® Xenoy 2230® Petsar PD 8100® Ektar EA-001® Sabre 1647, 1664® Dexcarb 500®

Partial miscibility Partial miscibility Partial miscibility Partial miscibility Partial miscibility Partial miscibility Partial miscibility

Amorphous/ Amorphous Amorphous/ Amorphous Amorphous/ Amorphous Amorphous/ Amorphous

Partial miscibility

Crystalline/ Amorphous

General Electric General Electric Hoechst Celanese BASF (Germany) Hoechst Celanese General Electric Denka (Japan) Hoechst Celanese DuPont Eastman

Valox 508®, 553® Valox 815®, 830® Celanex 5200®, 5300® Ultrablend S® Vandar® Valox 357®, 430®, 780® Malecca® Riteflex® Bexloy V® Ektar DA®

Partial Miscibility Miscibility

Crystalline/ Amorphous Crystalline/ Crystalline

PC additive (?) None Miscibility

Crystalline/ Amorphous Crystalline/ Amorphous Crystalline/ Amorphous Crystalline/ Amorphous Crystalline/ Amorphous

Miscibility

Amorphous/ Amorphous

General Electric

Miscibility

Amorphous/ Amorphous

Grafting Grafting Grafting Polycarbonate as compatibilizing additive

Crystalline/ Amorphous

Crystalline/ Amorphous

Amorphous/ Amorphous Crystalline/ Amorphous

Compatibilizing additive Crystalline/ Amorphous

1028

Table 15.1.

M. K. Akkapeddi

continued

Blend

Producer

Trade Name

Compatibilization Mechanism(a)

Blend Type

Specialty Resin blends Polysulfone/ABS

Amoco

Mindel A®

Amorphous/ Amorphous

Polysulfone/PET

Amoco

Mindel B®

Polyarylate/PET Polyetherimide/PC PBT/Acrylic

Unitika General Electric Hoechst Celanese

V-8000® Ultem UT350® Vandar®

Phenoxy resin or compatibilizer Phenoxy resin or compatibilizer Partial Miscibility Partial Miscibility None

Crystalline/ Amorphous Crystalline/ Amorphous Amorphous/ Amorphous Crystalline/ Amorphous

Notes: (a) Compatibilization mechanisms are defined here as the underlying principles responsible for the blend’s desirable properties, delamination resistance and morphology stability. (b) No compatibilizer used. Low interfacial tension between the components is responsible for the inherent compatibility in these systems. (c) Crosslinking of the dispersed phase via dynamic vulcanization (see 16.6) stabilize the blend morphology. (d) Partial miscibility between the components leads to self-compatibilization even though these systems are phase separated. Experimentally, these blends exhibit two Tg’s, but different from the pure components due to a small amount of mutual miscibility. (e) Grafting involves direct chemical reaction between the components during melt mixing, generating a graft copolymer as the compatibilizer. (f) Ionic or chelate complex forming interactions between the components at the interface, may lead to compatibilization.

Figure 15.1. Price/performance profile of commercial polymers.

Commercial Polymer Blends

adequate level of interfacial adhesion, e.g., in ABS/polycarbonate, PBT/polycarbonate, PVC/ nitrile rubber blends. 3. Compatibilizing effects of interfacial agents such as block or graft copolymers that reduce the interfacial tension, stabilize the morphology, and strengthen adhesion at the interface. Although complete miscibility between polymers is generally rare, when this does happen, it provides a unique opportunity to tailor the blend properties by simply adjusting the blend ratio. The immense commercial success of PPE blends with high impact polystyrene (HIPS) is primarily attributed to the miscibility between PPE and PS, that enabled one to combine the low cost, easy processability features of HIPS with the high performance (high heat resistance and strength) features of PPE. There was also some synergistic improvement of the toughness due to the improved ductility of the blend matrix resulting in improved rubber toughening efficiency. Simple melt mixing in extruder type compounding equipment is adequate to make an alloy from such miscible polymers. Even in the phase separated blends, where some degree of partial miscibility or compatibility exists between the components, simple melt blending in an intensive shear mixer is adequate for making a well dispersed, reasonably stable blend product with useful combination of properties, such as polypropylene/ethylene-propylene rubber blend, ABS/polycarbonate blend, etc. The self-compatibilizing nature of these blends stems from partial miscibility and the mutual interpenetration of polymer chains at the interface. Slight modifications of the polymer backbone are often employed, particularly in the case of styrenic and ABS resins to induce partial miscibility with other resins. Compatibilization of the highly immiscible commercial polymer pairs has thus far been a technically more challenging task for the polymer blend technologists in the industry. Significant progress has, however, been made in recent years in utilizing compatibilizers based on graft or block copolymer or other interfacial agents that

1029

effectively reduce the interfacial tension between the components to achieve useful levels of ductility and delamination resistance, while at the same time stabilizing the morphology against processing effects. Interfacial compatibilization in commercial polymer blends is generally achieved through reactive extrusion in which the block or graft copolymer compatibilizer is generated insitu at the interface. The end groups or pendant groups in some commercial polymers such as polyamides (amine or carboxyl), polyesters (carboxyl or hydroxyl) and styrene-maleic anhydride copolymers (anhydride group) are useful for this purpose. In most other cases, the polymer backbones must appropriately be modified to include reactive functionalities by graft reactions with small molecules (e.g., maleation of polyolefins with maleic anhydride) or through copolymerization techniques. The recent advances in reactive extrusion technology, involving reactive modification of polymers and/or reactive blending to form graft or block copolymer compatibilized blends in an extruder, have resulted in several successful commercial blends of otherwise highly immiscible polymer pairs, such as polyamide/olefinic elastomer, polyamide/polyolefin and polyamide/PPE blends. In many cases, the reactive modification of the base polymers and the subsequent reactive blending can be combined into a one-step, sequential operation in a twin-screw extruder, making this an economically attractive process. Reactive compatibilization is still an emerging technology for the development of useful blends from otherwise highly immiscible polymers, and particularly those of crystalline polymers. Specific examples of reactive modification and reactive blending will be discussed later in individual cases. Further details on compatibilization strategies can be found in chapters 4 Interphase and Compatibilization by Addition of a Comptabilizer and 5 Reactive Comptabilization of Polymer Blends of this book.

1030

15.2.2

M. K. Akkapeddi

Rationale for Polymer Blends

Commercial polymer blends belong to one of the following three categories: 1. Blends of amorphous/amorphous polymers 2. Blends of crystalline/amorphous polymers 3. Blends of crystalline/crystalline polymers While the specific advantages of each type of the blends will be discussed later in detail for the individual cases, the general motivations for making commercial polymer blends may be any of the following factors: • lower cost • improved heat resistance • improved toughness • improved solvent resistance • improved moisture resistance • improved dimensional stability • improved processability • improved weatherability • improved flame resistance • improved esthetics and appearance, etc.

Many commercial polymer blends often include an elastomer, to improve the impact strength of the blend under conditions of stress concentration (notched Izod impact strength) and to lower the ductile-brittle transition temperature of the blend. The elastomeric dispersions are judiciously employed either in the matrix phase, in the dispersed polymer phase, or in both phases, depending upon the requirement and the fracture behavior of the blend. As a general rule, the more brittle component in a given polymer blend has a greater need for rubber toughening. However, an overwhelming factor in determining the impact strength of an immiscible or partially miscible blend is the degree and efficiency of interfacial compatibilization that either is inherent in or has been designed into the blend system. If the interfacial adhesion or compatibilization is poor, the elastomer dispersion alone will not improve the toughness. Further details on the role of compatibilization and rubber toughening effects in polymer blends will be discussed later with specific commercial examples. Combining

Figure 15.2. Notched Izod impact strength vs. DTUL (in °C at 1.82 MPa) of various commercial blends.

Commercial Polymer Blends

1031

Figure 15.3. Notched Izod impact strength vs. DTUL (in °C at 0.45 MPa) of various commercial polymer blends.

a high level of impact strength with a high level of heat resistance and/or chemical resistance has been the primary thrust of most commercial polymer blends (Figures 15.2., and 15.3). Polymer blend processing Most commercial polymer blends are produced by melt mixing in continuous compounding equipment such as single-screw or twin-screw extruders and kneader-extruders. Twin-screw extruders are becoming more popular because of their greater versatility and production efficiency. Owing to their segmented barrel and screw designs, twinscrew extruders offer the advantage of multiple processing zones. The degree of shear mixing, residence time, and the temperature in each of these zones can be varied at will by simply changing the order and/or the type of the screw elements and kneading blocks, and using separately controlled heaters. Twin-screw extruders also offer the choice of multiple ports for liquid injection, venting and

downstream feeding purposes. For example, the twin-screw extruders can be readily adapted to combine both the polymer modification and polymer-polymer grafting steps of a reactive blending process into a single-pass, multi-stage reactive extrusion process. The polymer that needs to be functionalized is fed at the throat of a co-rotating twin-screw extruder and melt-mixed with a suitable grafting agent and catalyst in the initial zones. After vacuum venting to remove the unreacted and volatile materials, the ‘functionalized polymer’ melt is then mixed with the second polymer and other components (impact modifier, fillers, additives, etc.) added through the down stream feed ports. The temperature and the shearing conditions in these latter zones of the extruder are controlled to promote intimate mixings. Polymer-polymer grafting, compatibilization and the morphology development in the blend occurs at this stage of extrusion. Thus all the sequences of the complex blending process can be accomplished in an economically viable one-step, extruder process.

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M. K. Akkapeddi

Many commercial blends are often not as complex as above and can be made by simple melt blending without a compatibilizer because of partial miscibility between the components. However, even in these cases, good intimate mixing (dispersive and distributive) between the components is necessary to ascertain a morphologically stable, good quality product. Although twin-screw extruders are preferred for polymer blending, many single screw extruders, with mixing devices, are also employed, for the reason of their low cost. Further discussion on the compounding procedure can be found in Chapter 9 Compounding Polymer Blends, in this Handbook. Although commercial twin-screw extruders can be as large as 300 mm size, capable of compounding up to 40 ton/hr, the actual type and size of the equipment used depends on the type of the polymer blend and the production volume. Normally, for engineering polymer blends, twin-screw extruders of about D = 90 mm size (L/D ≅ 30 to 40) and capable of compounding at 700 to 1000 kg/hr, are used. For blending PVC or elastomer blends other types of compounding equipment are used, e.g. Farrell continuous mixer (FCM), Buss co-kneader, or a batch mixer, such as Banbury. The technological details on commercial polymer blends are kept proprietary by the manufacturers, particularly with respect to the exact compositions of the commercial grades and the processes used to make them. Even when the patent literature is available for a given type of polymer blend, often one cannot infer the actual compositions and processes used for the commercial blends. Hence, in reviewing the technology of commercial polymer blends, certain assumptions and generalizations have to be made regarding the compositional effects on the properties of some commercial blends, where the literature information is lacking. However, in the following sections, all relevant technological principles behind the various commercial blends have been outlined along with a discussion of their key properties, differentiating values and applications based on a reasonable evaluation of the literature available to date.

15.3

Polyolefin Blends

Polyolefins (PO) constitute the largest single group among all the commercial thermoplastics, considering the worldwide volume of usage. The polyolefin family comprises of (a) polyethylenes of various types, viz. high density polyethylene (HDPE), low density polyethylene (LDPE), linear low density polyethylene (LLDPE); (b) polypropylene (PP); (c) ethylene-propylene copolymers, block and random, ranging from tough thermoplastic to rubbery types; (d) ethylene copolymers with various comonomers such as ethyl acrylate, acrylic acid, methacrylic acid and the ionomers; (e) specialty polyolefins such as poly(4-methyl,1-pentene), poly(1-butene), etc. Among these, HDPE, LDPE and LLDPE and polypropylene are the four most widely used, commodity type polymers with a current total consumption in USA of nearly 13 Mton/y [Greek, 1991; Anonym., 1993]. Because of the wide range of properties available within PO’s, generally there has not been any major need to blend polyolefins with other types of polymers. Furthermore, the immiscibility of the commodity polyolefins with other types of polymers has also been a major reason for the lack of commercial interest in such blends, although with the advent of reactive modification and grafting chemistry, significant progress has been made in recent years in compatibilizing such dissimilar polymer systems as PA/PO blends [Ide and Hasegawa, 1974; Epstein, 1979; Hobbs et al., 1983]. Blending within the family of PO has, however, been more common [Plochocki, 1978]. Although they are usually immiscible with each other, there exists some degree of mutual compatibility between them. The similarity of their hydrocarbon backbones and the closeness of their solubility parameters, although not adequate for miscibility, accounts for a relatively low degree of interfacial tension. For example, the solubility parameters of polyethylene, polyisobutylene, ethylene-propylene rubber and polypropylene are estimated to be 16.0, 15.4, 15.5, and 17.0 J1/2cm3/2 respectively, all very close to each other [Van Krevelen 1990]. Similarly, the inter-

1033

Commercial Polymer Blends

facial tension coefficients between PE or PP and EPelastomers are quite small (typically ca. 0.1 MN/m) [Shih, 1990; Wu, 1989]. Hence polyolefin blends have been made by simple melt mixing without a compatibilizer since the early days of polyolefin commercialization. Generally, crystalline polyolefins such as HDPE and polypropylene have been blended with low modulus/elastomeric polyolefins such as LDPE, EP-rubber or polyisobutylene, in order to improve the toughness. Hence toughened PO have traditionally constituted the major volume of polyolefin blends used commercially. Most of the toughened PO blends are simple mechanical mixtures of PO’s and olefinic elastomers melt blended in an extruder without a compatibilizer. However, recent advances in polymerization technology has allowed the production of toughened polypropylenes, through sequential polymerization of ethylene-propylene copolymer in PP matrix leading to blends with some block or graft copolymer exhibiting somewhat improved modulus/toughness balance [Galli and Haylock, 1991]. Another recent development in polyolefin blends is the technology of dynamic vulcanization by which an elastomer is dispersed and cured in the matrix of the thermoplastic polypropylene [Coran, 1987]. Since the latter type blends have the characteristics of thermoplastic elastomers, they will be discussed under the elastomer blends section in Part 15.7. 15.3.1 15.3.1.1

Blends Based on Polyethylenes Polyethylene Blends (LLDPE/LDPE or HDPE Blends)

HDPE, the toughness, puncture resistance and heat sealability of the HDPE films could be improved. It is estimated that in the USA over 30% of PE consumed in films, is usually blended with LLDPE [Kosoff, 1987]. 15.3.1.2

HDPE is a highly crystalline polymer with a good combination of stiffness, strength and toughness, suitable for most packaging applications. However, for more demanding applications such as industrial bagging, an increased toughness, puncture resistance and environmental stress crack resistance was needed. This could be achieved by blending an elastomeric polyolefin such as polyisobutylene (PIB, e.g., Vistanex®, Exxon). The dart impact and tear strength of the film are improved significantly with the elastomer dispersion [Anonym., 1974] (Figure 15.4). Some commercial grades of HDPE (Paxon 3200 series) utilize this technology [Haartman, et al., 1970]. Although HDPE and PIB are not miscible, they have some degree of interfacial compatibility. The fact that their solubility parameters are quite close δ = 16.0 vs. 15.4) leads to low interfacial tension coefficient. The dispersability of the elastomer can be optimized by controlling the molecular weight, rheological parameters and the mixing conditions in the extruder blending process. With the advent of new tougher versions of LLDPE designed for blown film applications, further growth of HDPE/butyl rubber blends may however be limited. 15.3.1.3

Over 60% of the LDPE and LLDPE produced are used in film applications. Because of the continuous nature of the film extrusion, traditionally this process has been more amenable for blending different types of polyolefins to achieve specific improvements in film properties. For example, by blending 30% LLDPE with LDPE, thinner gage films with improved tensile and tear strengths could be produced [Forger, 1982]. Similarly, by blending about 30% LLDPE with

Toughened HDPE (HDPE/polyisobutylene blends)

Polyethylene/polystyrene Blends

A blend of PS in HDPE (interpenetrating network) is commercially sold as expandable beads (Arcel®, ARCO; Neopolen® S, BASF) for making cellular foams [Kosoff, 1987]. Although PS and polyethylene are immiscible, closed-cell foams made from this immiscible pair appear to combine the rigidity of PS with the solvent resistance and abrasion resistance of HDPE. Hydrogenated styrene-butadiene-styrene block

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M. K. Akkapeddi

Figure 15.4. Dart impact and tear strength of films (0.1 mm) made from HDPE/polyisobutylene blends [Anon., 1974].

copolymers are known to compatibilize the PE/PS blends. However there are no commercially significant applications for such compatibilized blends. 15.3.1.4

Polyethylene/polyamide Blends

A graft copolymer compatibilized blend of PA66/PA-6 (75/25) copolymer with HDPE was commercially offered as a barrier resin for making permeation resistant solvent containers (Selar®RB, DuPont) [Subramanian, 1984]. Before melt blending with the PA, the PE backbone must be modified by grafting with such reagents as maleic anhydride [Steinkamp, 1976]. A graft-coupling reaction between the PA and the maleated polyethylene, involving an amine/anhydride addition reaction leads to the graft copolymer formation at the interface, which reduces the interfacial tension and stabilizes the PA dispersion in the HDPE matrix.

By proper choice of the molecular weight, melt rheology of the PA and processing conditions, a platelet like dispersion of the PA in HDPE matrix could be achieved [Subramanian, 1985]. Since PA is a good barrier to hydrocarbons and many organic solvents, the platelets in HDPE provide the desired permeation resistance to solvents that HDPE lacks. On the other hand, polyethylene matrix provides the toughness, moisture resistance and low cost advantages compared to PA. The blend was designated as a ‘laminar barrier’ blend. The PA/PE ‘grafted blend’ was offered commercially as a concentrate (Selar® RB) to be melt blended with HDPE to a final PA/HDPE ratio ≅ 15/85 for subsequent blow molding into containers such as gasoline tanks, solvent containers, etc. This laminar barrier blend of HDPE and PA was reported to provide up to 100 fold improvement in the barrier to permeation of such organic solvents as toluene, relative to pure HDPE, or a similar blend composition containing PA as a uniform spherical dispersion.

Commercial Polymer Blends

The commercial potential of the PA/HDPE laminar barrier blend technology has not been fully established yet. A potential problem could be the high sensitivity of the morphology to the process conditions, which could lead to a lack of reproducibility in achieving the desired platelet morphology. The advent and rapid commercial succes of coextrusion blow molding may also be another factor limiting the blend’s market penetration. Coextrusion assures a more uniform barrier layer of PA that can be relied upon for permeation resistance. Surface fluorination of polyethylene, to improve the resistance to permeation of oxygen and solvents, was also a commercially competing technology. 15.3.2 15.3.2.1

Blends Based on Polypropylene Impact Modified Polypropylene (PP/EPR blends)

Polypropylene is a very large volume thermoplastic with an annual consumption of nearly 4 Mton/y in the USA alone. It is used in a variety of applications such as fibers, films and molded parts. Commercial polypropylene homopolymer is produced by the stereospecific polymerization of propylene in the presence of Ziegler-Natta type catalysts, which gives an isotactic polymer of high crystallinity. Due to this crystallinity and a reasonably high melting point (165°C), isotactic polypropylene exhibits a useful combination of properties such as good stiffness, strength, heat and solvent resistance. The hydrocarbon nature of the backbone imparts it excellent hydrolytic and dimensional stability. In addition to these properties, polypropylene’s low density, 900 kg/m3, and relatively low cost makes it an attractive candidate for many applications and often competing with even the higher priced engineering thermoplastics. The crystallinity of isotactic polypropylene homopolymer, PP, however, leads to its well known brittle behavior at low temperatures or when impacted under conditions of stress concentration, i.e., in the presence of sharp notches. For example, the low temperature drop weight (Gardner) impact strength of unmodified PP is

1035

typically < 2 J (at -30°C) and its notched Izod impact strength is < 40 J/m at room temperature, both significantly lower than that of polyethylene. This brittleness of polypropylene is related to the spherulitic morphology and the intrinsic tendency of PP for crazing followed by unstable craze growth and crack propagation under conditions of stress concentration and/or low temperatures [Kinloch and Young, 1983; Friedrich, 1983]. Commercial impact modified PP were developed during the early 1970’s by melt blending about 5 to 25 wt% ethylene-propylene rubbers (EP or EPDM) with the polypropylene homopolymer via extruder compounding [Holzar, 1966]. Some LDPE or HDPE is often used to assist the dispersability of the EP rubber and enhance the impact/modulus balance of the product [D’Orazio, 1982]. For the toughening of PP, EP rubbers meet the normal criteria for impact modifiers. Toughening of rigid polymers generally requires the formation of fine rubber dispersions with good adhesion to the matrix and thus providing multiple sites for crazing and localized shear yielding as mechanisms for the impact energy dissipation. Although not miscible, there exists a reasonable level of compatibility between EP rubbers and the polypropylene matrix due to their similarity in the hydrocarbon type structures and the closeness of solubility parameters which leads to low interfacial tension and an adequate level of interfacial adhesion [Krause, 1972]. Hence a small particle size dispersion of EP rubber is readily achievable by adjusting the molecular weights and melt viscosity ratio of EPR to PP and through the proper choice of mixing conditions [Yeh and Bisley, 1985; Rifi, 1987]. The control of the rubber-phase morphology in polypropylene also depends on the rubber composition (ethylene/propylene ratio), crystallinity, compounding and processing methods, as well as rheological properties. A slight crosslinking of EPDM rubber leads to better morphology stability during high shear flow conditions such as injection molding [Dao, 1982; 1984]. Commercial impact modified PP based on such blends exhibit excellent notched Izod impact strengths ranging from 80 J/m to 800 J/m and moduli ranging from 1000 to 1700 MPa. Because

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M. K. Akkapeddi

of the large number of commercial suppliers and grades, no attempt will be made to describe the properties of any specific type of commercial impact modified PP blends, except to generalize the trends in the impact/modulus behavior with respect to the content and the nature of the EP rubber. In general, for best impact modification, the ethylene-propylene copolymer must have > 30% propylene and show essentially no crystallinity [Rifi, 1987]. The size and distribution of the EP rubber particles in the polypropylene matrix depend on the molecular weights and the melt viscosity ratio between the two polymers, as well as the mixing conditions [Speri and Patrick, 1975; Jang et al., 1985]. Particle size must be < 3 µm and preferably be < 1 µm for optimum impact strengths. Melt index ratio of PP/EPR, MI ≅ 10 to 60 was found to give acceptable dispersion size. The volume fraction and the particle size distribution of the rubber dispersion affect the modulus, impact strength and melt flow of the blend significantly. Although historically, commercial impact modified PP has been produced by extruder compounding, more recent advances made in the catalyst and polymerization technology allowed the production of these blends via an “in-situ” process in the reactors [Rifi, 1987; Galli and Haylock, 1991]. By using a “super-active” catalyst and a gas phase, fluidized-bed reactor technology, the ethylene/propylene copolymer is allowed to polymerize and grow within the polypropylene homopolymer particle, polymerized earlier in the same reactor without isolation. The reactor technology could be tailored to produce a wide choice of products ranging from impact modified polypropylene to thermoplastic elastomer type PP/EPR blends (TPOs) by a judicious choice of the reaction conditions and the component feed ratios. The particle size of EPR in reactor TPOs is reported to be much smaller and hence more efficient in impact modification. Impact vs. modulus balance of the extrudercompounded blends of PP/EPR is illustrated in Figure 15.5. While the flexural modulus values are comparable, the low temperature impact strength

of the reactor-made blend products is usually better than that of extruder-made blends. Conventional extruder compounding process for making impact modified PP and thermoplastic polyolefin blends (TPOs) is still widely used to date, particularly by the independent compounders because of the versatility of this process for making a variety of specialized products having a wide-range of performance characteristics. Compounding additive packages for improved heat stability, weatherability and paintability, is readily feasible in the extruder blending process. Colors, pigments and fillers can also be blended in during the compounding of the PP/EPR blends. Impact modified polypropylene is used for injection molding automotive, consumer and appliance parts. For example, the medium impact PP is widely used for automotive interior trim. High impact PP is used for battery cases, fender and truck liners. Impact PP is also used extensively in the houseware and appliance markets. 15.3.2.2

Thermoplastic Polyolefin Blends (TPOs)

Thermoplastic polyolefin blends (TPOs), also sometimes referred to as thermoplastic polyolefin elastomers (TPE’s) or olefinic TPE’s, are blends of olefinic thermoplastics (primarily polypropylene) with elastomers with properties ranging from flexible elastomers to tough rigid materials, each specifically formulated to meet the needs of a particular application [Srinivasan, 1991; Spencer, 1990]. Essentially they are not much different from the impact modified polypropylenes, except that higher levels of EP rubber (up to 60%) may be used for additional low temperature impact strength. In some cases, the dispersed rubber may also be partially crosslinked during the mixing without losing the thermoplastic character of the matrix. However, the dynamically vulcanized elastomeric alloys with high rubber content are considered as a separate class of materials and hence will be discussed separately under elastomer blends. Most TPO alloys are compounded in the extruders whereby additional components such

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Commercial Polymer Blends

Figure 15.5. Effect of EPR content on the modulus and the low temperature impact strength of Polypropylene/EPR blends [Rifi et al., 1987].

as fillers, stabilizers and adhesion promoters may also be employed. For improved paint adhesion without primers, the blends may contain maleic anhydride or acrylic acid grafted polyolefins. Commercially important TPOs may have flexural moduli ranging from 20 to 2000 MPa and service temperatures ranging from -40 to 130°C. They are noted for their excellent toughness, high notched Izod impact strengths, and reasonable level of chemical resistance. The largest market for TPO is in the automotive area because of the material’s low cost, low specific gravity, low temperature toughness and weatherability. TPOs with molded-in color are used for bumper fascias, grilles, air dams, side moldings, etc. Under-thehood applications include ducting, protective shielding and sound abatement systems. The current annual consumption of TPO blends in transportation alone is 50 Mton/y in the USA and is projected to grow at >10%/yr [Spencer, 1990]. The use of TPOs in automotive bumper covers is expected to grow even more at an annual rate of about 30% [Evans and Mosier, 1991].

Lower cost and recyclability are the primary reasons for the rapid growth of TPOs in bumper applications, relative to SMC, polyurethane RIM and steel. Typical properties of TPOs used in bumper applications [Evans, 1991]:

Density (kg/m3) Flexural Modulus (MPa) at -30°C at 22°C at 70°C Tensile yield strength (MPa) Elongation at break (%) Hardness (Shore D) Instrumented Impact (-30°C) (Energy at Max. load, J) Notched Izod

930 2380 700 175 13.5 400 60 19 (with ductile failure) No break

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M. K. Akkapeddi

Reactor TPOs, reportedly offering better combination of toughness and stiffness, are gaining increasing penetration into the automotive market, in both the soft bumper fascia applications as well as other rigid applications replacing the more expensive engineering resins. 15.3.2.3

Elastomeric Polyolefin Blends

Thermoplastic elastomer blends comprising fully cured elastomer dispersions in a matrix of thermoplastic polyolefin such as PP, have been commercial for some time. These blends have been made by the technology of dynamic vulcanization [Speri and Patrick, 1975]. The process consists of melt mixing and dispersing a high volume fraction of an elastomer such as EPDM rubber or nitrile rubber (NBR) into a thermoplastic matrix such as PP, using a compatibilizer if necessary, and then selectively crosslinking the dispersed elastomer during the extrusion, with specific curing agents. The resulting elastomeric blends display the typical properties of cured rubbers such as high elastic recovery, low compression or tension set, yet process like thermoplastics, due to the presence of PP thermoplastic matrix. Commercial, elastomeric polyolefin blends produced the dynamic vulcanization include EPDM/polypropylene blend (Santoprene®); Nitrile/polypropylene blend (Geolast®); Butyl rubber/polypropylene (Trefsin®). The technology and properties of these blends will be discussed in Part 15.7 Elastomeric blends. 15.3.3

Ionomer/polyolefin Blends

Commercial ionomers are ethylene-methacrylic acid copolymers and terpolymers in which the carboxylic acid moiety is partially neutralized with sodium or zinc, to promote interchain ionic bonding. Ionomers exhibit excellent low temperature toughness, chemical resistance and adhesion. However they lack in stiffness and heat resistance. Hence ionomer blends with polyolefins such as polyethylene have been developed which, upon reinforcing with suitable fillers, seem to give a unique combination of high strength, excellent low temperature toughness, and moderate stiff-

ness and heat resistance (Surlyn® HP, DuPont). Key to this technology appears to be the selection of a suitable surfactant that must be added in small amounts (0.3%) to aid the dispersability of glass fibers, yet retain high toughness [Murphy, 1986]. The blend is reported to have an interpenetrating network (IPN) type morphology, with the ionomer as the continuous phase. An unusual feature of this blend is the high notched Izod impact strength of >1000 J/m at 23°C and >760 J/m at -29°C, even in the presence of 15% of chopped glass fiber. In the absence of the surfactant additive, the blend with the same level of glass fiber showed somewhat poorer impact properties, e.g., notched Izod of < 500 J/m at -29°C and elongation at break of < 3%. In the presence of the surfactant additive, an ionomer/polyolefin blend reinforced with 15% glass (Surlyn® HP) exhibits a flexural modulus of 1600 MPa, a high notched Izod of > 1000 J/m and a moderate DTUL of 82°C (at 0.45 MPa). These properties were claimed to be good enough to enable this blend to compete with the impact modified polyamides in some applications. 15.3.4

Reactor-made Polypropylene/ Non-olefinic Polymer alloys

A new class of ‘reactor-made’ alloys or ‘in-situ’ graft copolymer compatibilized blends of polypropylene with other amorphous, non-olefinic polymers have been commercially introduced by Montell recently under the trade name of Hivalloy® [DeNicola, 1994]. The commercial Hivalloy G series consisting of polypropylene/polystyrene alloys were first launched on a pilot scale in the mid-1994 and later fully commercialized in 1996. Their Hivalloy W series are developmental grades of polypropylene/ acrylic alloys, while the Hivalloy T series are experimental grades of polypropylene/ styrene-maleic anhydride (SMA) copolymer alloys. All these reactor-made PP alloys are produced by Montell’s proprietary ‘Catalloy’ or ‘reactor granule’ technology [Galli and Haylock, 1991; Galli et al., 1994], which is a multi-stage, multimonomer polymerization process. The fundamental basis for the reactor granule

Commercial Polymer Blends

technology starts with the formation of a highly porous polypropylene particle first by the polymerization of propylene monomer using a ‘superactive’, third-generation, Ziegler-Natta initiator system which consists of a MgCl2-supported, electron donor-modified TiCl4/AlR3 catalyst in a very high surface area, spherical particle shape (‘Spheripol’ process). A unique feature of this catalyst system is the high uniform porosity that is maintained during the polymerization in the growing PP particles. As the polymerization takes place, a growing skin of PP polymer is formed from the active sites on the surface of the expanding initiator particle. Proper control of the morphological structure of the initiator particles and the polymerization process results in the formation of highly porous, lattice-structured PP granules. These highly porous PP granules or beads are then used as the ‘reactor bed’ for subsequent polymerization of one or more non-olefinic monomers such as styrene and methyl methacrylate. Since the latter polymerization is generally free radical, it results in the simultaneous formation of the non-olefinic polymer as well as its graft copolymer with PP uniformly distributed in micronscale domains within the individual PP granules. This ‘in-situ’ generated graft copolymer then effectively compatibilizes and stabilizes the PP blend morphology during subsequent melt processing such as injection molding. The blends exhibit typical multi-phase morphology behavior with the non-olefinic polymer generally forming the fine dispersed phase in the continuous PP matrix phase [DeNicola, 1992]. The PP matrix may contain additional rubber particle (EPR type) dispersions for impact toughening of the blend. The mechanical properties of these reactormade alloys offer a balance of stiffness and toughness not generally attainable through the simple melt blending of the same polymer systems, primarily due to the effective graft copolymer compatibilization and good interfacial adhesion between the component phases. The Hivalloy reactor blends have been positioned by Montell to compete against ABS alloys and other low-end engineering resins based on on some comparable

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properties, as well as better chemical resistance and a lower specific gravity. The higher Tg of the amorphous, non-olefinic polymer dispersion (PS, Acrylic, SMA) is expected to reinforce the PP matrix yielding a somewhat higher stiffness, strength and heat distortion temperature, while the high crystalline melting point, ductility, chemical resistance and high melt flow characteristics of PP are maintained. The following is a summary of the properties and applications of the two commercially significant reactor alloys viz., PP/PS and PP/acrylic polymer alloys. 15.3.4.1

Commercial, Polypropylene/ Polystyrene (PP/PS) reactor alloys

A series of reactor alloys of PP and PS are available from Montell under the Hivalloy G trade name, in the unreinforced, rubber toughened and glassreinforced forms. Typical properties are illustrated in Table 15.2. A key feature of these PP/PS alloys is the impact strength/stiffness envelope that reportedly exceeds the performance of conventional PP and approach in some aspects with those of other engineering resins such as acetals, PC/ABS, PC/PBT. Depending on the PS content in the blend, the flexural modulus and DTUL increase predictably above that of PP while maintaining a high level of ductility and ultimate elongation. The notched Izod impact toughness could be raised to a ‘no-break’ level particularly with the incorporation of some ethylene-propylene (EP) or styrene-ethylene/butylene-styrene block copolymer (SEBS) elastomeric modifier. Compounding varying levels of glass fibers and reinforcing mineral fillers provides the desired balance of stiffness and toughness in these PP/PS alloys. The reinforced grades reportedly exhibit improved stiffness and creep resistance compared to PP alone and may compete against reinforced polyamides and polyesters, particularly in applications that do not require high temperature performance. Typical applications in development with these PP/PS alloys include automotive bumper beams, pillars, sporting and recreational equipment, sledge hammer handles and other consumer tools & appliance components. The lighter weight (lower

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M. K. Akkapeddi

Table 15.2.

Properties of some commercial, ‘reactor-made’ polypropylene/polystyrene and polypropylene/acrylic polymer alloys

Blend Type

PP/PS

PP/PS/EP

PP/PS, 35 wt% GF

PP/PMMA

PP/PMMA

UNITS

Hivalloy G120 Montell

Hivalloy G170 Montell

Hivalloy GXPA072 Montell

Hivalloy WXPA018 Montell

Hivalloy WXPA012 Montell

0.94 1.3 <0.2

0.92 1.2 <0.2

1.2 0.3 <0.05

0.98 1.4

0.95 1.4 <0.05

1520 (220) 43 (6.2) 33 (4.8) 45

1200 (170) 33 (4.8) 27 (3.9) 200

7580 (1100) 165 (24) 113 (16.4) 2.5 114

1930 (280) 54 (7.9) 36 (5.2) 20

1380 (200) 39 (5.7) 30 (4.3) 180 90

135 (2.5)

No break

107 (2.0)

59 (1.1)

160 (3.0)

36(27)

40 47

93 60

88 55

PROPERTY

ASTM Method

PHYSICAL Density

D792

kg/m3

Mold Shrinkage Water Absorption, 24 h

D955 D570

% %

MECHANICAL Flexural Modulus Flexural Strength Tensile Strength at Yield Elongation at Break Rockwell Hardness

D790A D790A D638 D638 D785

IMPACT Izod Impact, Notched at 23°C Instrumented Impact Total Energy, at 23°C Total Energy, at -20°C THERMAL Heat Deflection Temp. at 0.45 MPa at 1.82 MPa

MPa (kpsi) MPa (kpsi) MPa (kpsi) % R

D256

J/m (ft-lb/in)

D3763

J (ft-lb)

D648

°C

density) to stiffness & toughness balance of these alloys compared to the PC alloys is claimed to be a key advantage in these applications. 15.3.4.2

44 (32)

Commercial, Polypropylene/ Acrylic (PP/PMMA) Reactor Alloys

A series of reactor alloys based on polypropylene and methyl methacrylate copolymer have been commercialized by Montell under the Hivalloy W series. Typical properties of such an alloy are shown in Table 15.2. These alloys have been claimed to exhibit excellent weatherability,

160 150

96 63

90 57

surface appearance and colorability in addition to a good stiffness/ toughness combination. In comparative SAE J1960 exterior weathering tests, PP/acrylic alloy showed superior color & gloss retention compared to ASA and PC/PBT [Sherman, 1997; Hivalloy Data sheets, 1997]. The UV resistance of the blend must originate from the acrylic polymer, which may be enriched on the surface of the part. The attractive refractive index of the blend also makes it easy to color with molded-in colors. Typical applications in use or development with the PP/ acrylic polymer reactor alloys include

Commercial Polymer Blends

automotive exterior mirror housings, interior trim and handles, truck wheel fenders, marine/ outdoor recreational equipment, building and construction applications such as the outer layer for siding etc. The apparent outstanding weatherability, good colorability and mechanical properties of these alloys are positioned to compete against the well established engineering resins such as ABS, ASA, PC/ABS and PC/PBT. 15.3.5

Metallocene Polyolefin Blends

Recent progress made in the development of novel metallocene catalyst systems for polyolefin polymerizations represents a true technological breakthrough with wide ramifications in the plastics industry [Wigotsky, 1995; Schut, 1996; Zamora, 1997]. As the latest generation catalysts, the metallocenes differ from the traditional Ziegler-Natta catalysts in that they have well defined single catalytic sites and well understood molecular structures. Typically they consist of a transition metal such as zirconium sandwiched between suitably substituted cyclopentadienyl ring structures to form a sterically hindered or ‘constrained geometry’ metal catalyst site. With these new generation of catalysts, a wide variety of monomers can be polymerized or copolymerized in high efficiency but more importantly various parameters such as comonomer distribution, polymer molecular weight, molecular weight distribution, molecular architecture, stereospecificity, degree of linearity or branching can be independently and precisely controlled. While the range of the new metallocene based polymers includes such specialty polymers as cyclo-olefin copolymers (COC), syndiotactic polystyrene, ethylene/styrene copolymers, which are still in the developmental stage, commercially the most prominent candidates are the ethylene/ α-olefin copolymers such as ethylene/butylene or hexene copolymers (Exxon’s Exact) or ethylene/ 1-octene copolymers (Dow’s Engage and Affinity). Depending on the comonomer content these copolymers have been classified as plastomers or elastomers. At comonomer levels of > 25%, the copolymers exhibit the characteristics of thermoplastic elastomers such as high softness, toughness,

1041

flexibility and resilience and hence been referred to as polyolefin elastomers (POE). As a result of the controlled long-chain branching in otherwise linear polymers, the processability of these resins is claimed to be significantly enhanced compared to the standard LLDPE and EPR materials. Environmental stress crack resistance of metallocene polyolefin is also claimed to be significantly better. A wide range of densities (860 to 930 kg/m3), flexural modulus (10 to 100 MPa), melting points (60-120°C) and melt flow index (0.5 to 125) is available in the commercial metallocene-based ethylene copolymers. Thermoplastic polyolefin (TPO) blends of metallocene-based polyolefin elastomers (POE) with polypropylenes are beginning to gain commercial significance because of the improved melt flow and toughness compared to the conventional TPOs based on PP/ EPR or EPDM blends [Toensmeier, 1994]. In comparative tests with 70/30 PP/elastomer blends, the blends with POE maintained ductile behavior at -29°C even with high melt flow index PP (MFI = 35), while the corresponding EPR based blends were brittle with PP of MFI = 20. In addition, they showed improved knit-line strengths. Because of their high shear-thinning and melt elasticity, plastomers and POEs disperse well in PP matrix yielding fine particle dispersions leading to the better properties. TPOs are the fast growing segment of thermoplastics market in the ‘90s with more than 10% growth rate/yr. Typical applications of these TPOs include automotive bumper fascias, interior instrument panel skins, air-bag doors, etc. Blends of metallocene polyolefin elastomer/ plastomer materials compounded with EPDM (typically 70/30) are finding several extrusion applications such as flexible hoses, cords and wire jacketing because of much higher extrusion rates and better heat aging characteristics than EPDM [Sherman, 1997]. Blends of POEs with ethylenevinyl acetate copolymers are used as cushioning in sports shoes because improved resiliency and durability. Blends of POE with PVC are employed in some extruded profiles for refrigerator gaskets, window and garage door seals. Other applications of POE blends include soft-touch handles for

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M. K. Akkapeddi

hand-held tools and for automotive noise, vibration and harshness (NVH) dampening. As the metallocene olefin copolymers become increasingly more cost-effective, the application of their corresponding blends are expected to proliferate.

15.4

Styrenic Blends

Styrenic resins, a family of commercially significant polymers and copolymers derived from styrene, rank among the major volume thermoplastic materials used, with an annual consumption of nearly 4 Mton/y in the USA alone [Greek, 1991]. Their low cost, ease of processability and good balance of properties account for widespread use. Commercial styrenic resins may be classified into the following types: 1. Polystyrene (PS) and high impact polystyrene (HIPS) 2. Styrene-acrylonitrile copolymer (SAN) and its impact modified versions, viz., ABS (polybutadiene rubber grafted SAN), ASA (acrylate rubber grafted SAN), AES (EPDM rubber grafted SAN) 3. Styrene-maleic anhydride copolymer (SMA) and terpolymers with methyl methacrylate (SMA-MMA) and acrylonitrile (SMA-AN) 4. Styrene-methyl methacrylate copolymer (S-MMA) 5. Styrene-butadiene block copolymers [di, tri and radial block (S-B)n] Among these, polystyrene is the lowest cost, commodity type resin followed by SAN, SMA and the other copolymers. All the styrenic resins are essentially amorphous polymers with glass transition temperatures ranging from about 100 to 130°C, and heat distortion temperatures ranging from about 80 to 120°C, depending upon the comonomer and impact modifier content. Although the unmodified styrenics, viz. polystyrene, SAN and SMA copolymers, exhibit good clarity, strength and rigidity, they are invariably brittle for many applications. Hence the rubber modified styrenics such as HIPS and ABS, which combine a good level of impact strength with

moderate heat resistance, have become more widely accepted in many molding and extrusion applications. HIPS and ABS may themselves be considered as blends, since they normally contain • 5% polybutadiene rubber as a discrete phase, dispersed as 0.1-5 µm size particles in the matrix of polystyrene or SAN copolymer [Echte, 1989]. However, the rubber phase in these resins is incorporated during the free radical polymerization of styrene or S-AN monomer mixture via a mass, suspension or emulsion polymerization process that results in the graft-coupling of the rubber phase to the matrix phase. In the context of the current discussion on blends, HIPS and ABS are considered not blends, but more impact modified resin systems made in a reactor, although several grades of ABS are produced routinely by captively melt blending with SAN to adjust the rubber levels to the desired property specifications. HIPS and ABS are themselves used, however, as components for blending with other thermoplastic resins to make new blends with desired combination of properties. The general motivation for blending styrenic resins with other polymers, particularly with the higher priced engineering resins, such as polycarbonate or polyphenylene ether, is primarily to lower the cost and improve the processability of the latter resins. As far as styrenic resins are concerned, some of the reasons for blending stem from the need to improve their property deficiencies, viz. solvent resistance, impact strength, heat resistance and flame resistance. 15.4.1

Polystyrene and High Impact Polystyrene (HIPS) Based Blends

Because of its inherent brittleness, polystyrene homopolymer itself has limited application in blends. However, its impact modified version, viz., HIPS, is more widely used. HIPS itself is a reactor-made multiphase system with 5 to 13% polybutadiene (‘cis’-rich) dispersed as discrete particles in the polystyrene phase, with an optimum particle size of mean diameter of 2.5 µm. The rubber in HIPS is chemically grafted to some extent to the polystyrene. The effective volume of

Commercial Polymer Blends

the rubber dispersion is actually increased through the occlusion of some polystyrene. To optimize the impact strength, the rubber particle size (• 2.5 µm) and the distribution is normally controlled by the agitation and the proper choice of other process conditions during the polymerization. The property improvements in HIPS, viz., increased impact strength and ductility, are accompanied by the loss in clarity and a decrease in the tensile strength and modulus compared to the unmodified polystyrene. 15.4.1.1

Blends of Polystyrene with S-B Block Copolymers

In order to improve the toughness, yet maintain a sufficient level of clarity in the polystyrene homopolymer, it is often melt blended with some styrene-butadiene block copolymers containing • 40% styrene blocks. Such polystyrene blends are commercially employed in the blow molding, sheet extrusion and thermoforming applications [Traugott, 1985]. Both the multiblock type (Firestone Synthetic Rubber and Latex Company’s Stereon® resins) and the radial block type (Phillips Petroleum Company’s K-resins) S-B block copolymers are used for blending. These blends are normally made by the fabricator, e.g., during the sheet extrusion or blow molding [Salay, 1991]. Some commercial precompounded grades of blends are also available. Typical applications include thermoformed food packaging containers, blister packs, shrink-wrap films, etc. 15.4.1.2

Polystyrene or HIPS Blends with Poly(phenylene ether) (PPE/HIPS blends)

Blends of polyphenylene ether (PPE, also known as PPO®) with HIPS are, by far, the most successful of all the commercial blends. Currently, more than 100 kton/y of PPE/HIPS blends are produced in the USA [Levy, 1991] almost exclusively by General Electric company, which originally introduced this blend commercially in 1964 under the trade name, Noryl®. There are now other producers of this blend in Europe and Japan.

1043

PPE is the acronym used for poly (2,6-dimethyl 1,4-phenylene ether), a high Tg (205 to 210°C) polymer produced by the oxidative coupling polymerization of 2,6-dimethyl phenol [Hay, 1959; 1976]. Sometimes a minor amount of 2,3,6 trimethyl phenol is used as comonomer. Although PPE exhibits a good level of ductility and toughness along with a high heat distortion temperature, its high softening temperature and high melt viscosity precluded it from being used as a commercial molding resin, by itself. However, the discovery that blending PPE with high impact styrene could lead to improved processability and impact properties resulted in the successful commercialization of these blends [Cizek, 1968]. By simply adjusting the blend ratio, a wide spectrum of blend products with the desired combinations of DTUL, impact strength, processability and cost balance could be produced. This versatility of tailor-making the blends for various levels of performance led to their rapid commercial success. Key to this success, however, was the observed miscibility between the PPE and the PS phases resulting in a single phase matrix with a single glass transition temperature that can be varied at will with the blend ratio. This thermodynamic miscibility between PPE and PS was established by various characterization techniques viz. glass transition temperature [Schulz and Gendron, 1972; Fried et al., 1978], electron microscopy [Kambour et al., 1980], small angle X-ray scattering and calorimetric methods [Weeks et al., 1977], the latter showing a negative heat of mixing over the entire composition range. The presence of polybutadiene rubber particle dispersions in such a homogeneous matrix was found to lead to significant synergistic enhancement of the toughness of the blend, due to both the crazing and shear yield mechanisms of toughening [Bucknall, 1977; Yee, 1977]. The particle size and the amount of rubber, however, affect the impact, modulus and tensile properties of the blend. Commercial PPE/HIPS blends generally span a range of blend ratios from, 25/75 to 60/40 of PPE/HIPS. Typical properties of commercial PPE/ HIPS blends are shown in Table 15.3. Depending

1044

M. K. Akkapeddi

upon the PPE content, the heat distortion temperature can vary from 90°C to about 150°C. The melt rheology and flow characteristics again depend predominantly on the ratio of PPE to HIPS, with the molecular weights of PPE and HIPS also playing an important role [Schmidt, 1979; Priest and Porter, 1972]. All the blends exhibit good ductility and impact strength. The notched Izod values range from 250 to 500 J/m. The rubber particles in HIPS contribute to the enhanced toughness of the blend. Since PPE is inherently more ductile than polystyrene, the

Table 15.3.

efficiency of rubber toughening increases as the PPE content in the blend increases. Shear yielding process also contributes to the overall toughening effect, in addition to the usual craze-toughening mechanism. In blends containing • 50% PPE, shear yielding is the dominant mode of energy dissipation [Yee, 1977]. Unlike the case of polystyrene, the blend needs smaller size (” 2 µm) rubber particles for optimum impact strength properties [Bucknall, 1972]. For higher impact strength, additional blending of styrene-butadiene-styrene (S-B-S) block copolymer type elasto-

Key properties of commercial PPE/HIPS blends ASTM TEST METHOD

UNITS

NORYL SPN410 GEC

NORYL SPN420 GEC

NORYL 731 GEC

NORYL N300 GEC

PHYSICAL Density Mold Shrinkage Water Absorption, 24 hrs.

D792 D955 D570

kg/m3 % %

1,060 0.6 0.07

1,070 0.6 0.07

1,060 0.6 0.07

1,060 0.6 0.06

MECHANICAL Flexural Modulus Flexural Strength Tensile Strength at Yield Elongation at Break

D790 D790 D638 D638

MPa (kpsi) MPa (kpsi) MPa (kpsi) %

1830(265) 51(7.8) 36(5.2) 82

2140(310) 59(8.6) 44(6.4) 60

2530(360) 95(13.5) 67(9.6) 60

2400(350) 104(15) 76(11) 20

D256

J/m(ft-lb/in) 375(7) 190(3.5)

295(5.5) 133(2.5)

270(5) 130(2.5)

272(5) 130(2.5)

96

121 110

137 129

157 149

2.7

2.7

1.0E+15

1.0E+15

21.6 2.6 0.001 1.0E+15

21.6 2.6 0.001 1.0E+15

PROPERTY

IMPACT Izod Impact, Notched at 23°C at -40°C THERMAL Heat Deflection Temp at 0.45 MPa at 1.82 MPa ELECTRICAL Dielectric Strength Dielectric Constant at 1 KHZ Dissipation Factor at 1 KHZ Volume Resistivity

D648

D149 D150 D150 D257

°C

K/V/mm

ohm-m

1045

Commercial Polymer Blends

mers and their hydrogenated derivatives, S-EB-S, are often employed. The ratio of PPE to PS in the blends can be determined from the ratio of the IR peaks at 854 cm-1 and 700 cm-1 respectively [White, 1983]. The rubber particle size is determined by transmission electron microscopy using osmium tetroxide staining [Bucknall, 1977]. Most of the PPE/HIPS blends are utilized by the injection molding process, although in recent years, blow molding and extrusion/thermoforming applications are also increasingly practiced. Low moisture absorption and the good melt stability of PPE/HIPS blends along with their broad range of melt viscosities enable the fabricators a wide choice of processing conditions. PPE/HIPS blends are used in a wide range of applications in the automotive, business equipment, appliance, electrical/electronic and industrial markets. Automotive applications include instrument panel frames, interior trim, glove boxes, fuse boxes, connectors, wheel covers, mirror housings, etc. Flame retarded grades are used for business machine housings which are also often foamed to reduce the specific gravity and blow molded to reduce processing costs. Appliance parts and housings include those for refrigerators, washers, dryers, dishwashers, power tools, etc. The blends exhibit extremely low moisture absorption and good electrical properties suitable for electrical and electronic applications such as TV cabinets, connectors, electrical junction boxes, housings, relays, and bobbins. The low moisture absorption and excellent hydrolysis resistance of PPE/HIPS blends coupled with their high dimensional stability makes them suitable for water meters and pump housings, plumbing parts and a variety of fluid handling equipment parts. Some developing applications include such construction applications as roofing panels, insulation, flooring substrates, etc. The high heat distortion temperature of PPE/HIPS led to its evaluation in microwave packaging. In applications requiring clarity, blends of PPE and crystal PS (Gepax®, GE) have also been evaluated. Trays and packages that can be safely heated in microwave ovens have been made from such blends.

15.4.2

ABS Blends

Among the styrenic resins, ABS enjoys a unique position of being considered an engineering thermoplastic because of its key performance characteristics, viz. high impact strength, solvent resistance and moderate heat resistance. Due to its relatively low cost, ABS can bridge the gap between the commodity plastics and the higher priced engineering thermoplastics. Nearly 700 kton/y of ABS is produced currently in the USA with an estimated 20 kton/y metric tons used in the blends and alloys [Greek, 1991]. ABS itself is a two-phase polymer system with a finely dispersed polybutadiene rubber phase (0.1 to 1 µm) embedded in a continuous matrix of SAN copolymer. The rubber phase is chemically grafted to the SAN during the polymerization of ABS. As mentioned earlier, ABS is commercially produced by the free radical polymerization of styrene/acrylonitrile monomer mixture (usually 3:1 wt. ratio) in the presence of polybutadiene (of high ‘cis’ content) which is added as a solution (ca. 10%) in the mass polymerization process or as a latex seed during the emulsion polymerization (30-60% rubber) [Ku, 1985]. The emulsion grade ABS is usually blended with virgin SAN copolymer to produce an ABS with a desired level of rubber (10 to 25%). The impact strength, modulus, tensile strength, processability and surface gloss of ABS depend on the rubber content and its particle size and distribution that is determined by the polymerization process and its conditions. The high rubber ABS (25 to 50% rubber) grades can be made only by emulsion process and such resins (e.g., ‘Blendex’ grades from General Electric Co.) are often employed for compounding with other more economic thermoplastic resins such as polyvinylchloride (PVC) for impact modification [Dotson and Niznik, 1991]. The bulk polymerization process (continuous or suspension) produces low-gloss, medium impact grades of ABS. While the standard ABS resins, containing 10 to 25% polybutadiene rubber, indeed have a wide range of useful engineering properties, it is possible to extend further and improve these

1046

M. K. Akkapeddi

properties for certain niche applications via blending with other polymers. For example, blending with PVC can improve the flame retardancy of ABS at a low cost. Blending with such higher Tg, engineering resins as polycarbonate and polysulfone can improve the heat distortion temperatures. On the other hand, use of ABS in blends with other resins can bring the advantages of low cost, improved impact strength and processability. These will be discussed with the following examples of blends. 15.4.2.1

ABS/PVC Blends

Incorporation of 10 to 40 wt% ABS into PVC improves its impact strength, processability and hot tear strength. Commercially ABS/PVC blends are available from several sources, but more often these blends are made in-situ by the fabricators of sheet or profile extrusions. High rubber, low modulus grades of ABS made by emulsion polymerization are often used for the impact modification of PVC. Because of the thermal degradation problem of PVC, blending is done typically in Banbury type mixers. The blend exhibits significantly improved notched Izod impact strength (•1000 J/m) over PVC. Although PVC is immiscible with ABS, the interfacial tension between the SAN phase and PVC is low enough to allow enough compatibility. The rubber particles, of course, are responsible for the enhancement of toughness. The primary application of this type of ABS/PVC blend, particularly in Western Europe, is in the manufacture of foil for vacuum thermoforming automotive and mass-transit interiors. The fabricators blend the powders of PVC and ABS and use the blend captively for the foil manufacture. Similar blends of PVC with MBS and ASA as impact modifiers are also used. These will be discussed in Part 15.5.1 PVC/impact modifier blends. Pre-compounded blends of ABS and PVC have also been commercially available for molding and extrusion applications as low cost alternatives to flame retarded grades of ABS and PPE/HIPS blends. Commercial ABS/PVC blends offer high impact strength coupled with flame-

retardant characteristics (self-extinguishing, V-O ratings of UL-94) at a reasonable cost (Table 15.4). Such ABS/PVC blends have been used in appliance and business machine housings, TV cabinets, electrical and electronic component manufacture. Although the low cost, flame retardancy advantages of ABS/PVC blend are attractive, the thermal instability of PVC poses processing problems requiring careful control of processing conditions and temperature. In the injection molding markets, the processing disadvantages of ABS/PVC blend are limiting its growth, while improved versions of flame retardant grades of ABS, ABS/PC and PPE/HIPS blends are steadily gaining competitive advantage due to their superior processability. 15.4.2.2

ABS-Polycarbonate Blends

Blends of ABS and polycarbonate (PC) were commercially introduced by Borg-Warner Co. (now part of GEC) several years ago [McDougle, 1967; Grabowski, 1964]. After an initial sluggish growth, these blends have now gained increased acceptance by resin fabricators and end users. Currently more than 50 Kton/y of ABS/PC blends are consumed in USA, Western Europe and Japan [SRI, 1992]. There are many suppliers of ABS/PC blends, primarily by those involved in producing the polycarbonate and/or ABS. ABS/PC blend is an essentially immiscible blend [Echte, 1989; Suarez et al., 1984; Kim and Burns, 1988] with three distinct phases, viz. the PC phase, the SAN copolymer phase and the grafted polybutadiene rubber phase dispersed within the SAN phase (as in the ABS to begin with). Depending upon the blend ratio, the continuous phase can either be the ABS (or more correctly the SAN phase) or the PC phase. In spite of the immiscible nature, the blends exhibit good toughness, particularly in the region of 30-65 vol% PC [Weber and Page, 1986]. A primary reason for the good properties and delamination resistance in ABS/PC blends is partial miscibility between the PC and SAN phases, which leads to low interfacial tension and high interfacial adhesion, particularly when the SAN contains • 25% acrylonitrile [Keitz et

1047

Commercial Polymer Blends

Table 15.4.

Properties of some commercial styrenic/PVC blends

Blend Type

ABS/PVC ASTM TEST METHOD

UNITS

PHYSICAL Density

D792

kg/m3

Mold Shrinkage Water Absorption, 24 hrs.

D955 D570

% %

MECHANICAL Flexural Modulus Flexural Strength Tensile Strength at Yield Elongation at Break Rockwell Hardness

D790 D790 D638 D638 D785

MPa (kpsi) MPa (kpsi) MPa (kpsi) %

D256

J/m(ft-lb/in)

PROPERTY

IMPACT Izod Impact, Notched at 23°C at -40°C THERMAL Heat Deflection Temp at 0.45 MPa at 1.82 MPa Flammability Rating

D648

ASA/PVC

SMA/PVC

CYCOVIN KAF B.F. Goodrich

KANEKA ENPLEX Kanegafuchi

GELOY GY1220 GEC

TRIAX CBE Monsanto

1,200 0.4 0.09

1,200 0.4 0.09

1,200

1,200 0.5

2310(330) 67(9.6) 40(5.8) 20 R100

2740(390) 61(8.7) 44(6.3) 20 R105

2140(310) 60(8.5) 46(6.6) 25 R96

2800(400)

570(10)

201(3.7)

1080(20) 107(2)

320(6)

79 V-0

80 71 V-0

85 76 V-0

0.1

41(6.1) R106

°C

UL94

al., 1984]. The partial miscibility between PC and SAN has been demonstrated by the small yet finite increases in the Tg of SAN phase and similar decreases in the Tg of PC phase [Kim and Burns, 1988; Morbitzer et al., 1985]. The polymer-polymer interaction parameter (χ) for this blend was calculated to be slightly positive (χ = 0.03). Some interpenetration of the chains at the phase boundary is responsible for the increase in the interfacial strength. Properties of ABS/PC blend depend on the blend ratio. Blends containing major amounts of ABS (e.g., Cycoloy EHA) show improvements in DTUL and tensile properties relative to ABS,

76 V-0

but not in the notched Izod impact strength (Table 15.5). But as the polycarbonate level in the blend increases (particularly at PC • 50%), the notched Izod impact strength improves significantly, even at low temperatures. The ductilebrittle transition temperature in the latter blends is shifted to significantly lower temperature [Weber and Paige, 1985], even compared to polycarbonate alone. The high dimensional stability, the excellent impact toughness at low temperatures, low gloss and easy processability features of ABS/PC blends have led to their application in automotive interior instrument panels and, more recently, in exterior body panels, wheel covers, etc.

1048

M. K. Akkapeddi

Table 15.5. Properties of commercial grades of ABS/polycarbonate vs. polycarbonate Blend Type

ABS/PC ASTM TEST METHOD

UNITS

PHYSICAL Density

D792

kg/m3

Mold Shrinkage Water Absorption, 24 hrs.

D955 D570

% %

MECHANICAL Flexural Modulus Flexural Strength Tensile Strength at Yield Tensile Strength at Break Elongation at Break Rockwell Hardness

D790 D790 D638 D638 D638 D785

MPa (kpsi) MPa (kpsi) MPa (kpsi) MPa (kpsi) %

D256

J/m(ft-lb/in)

PROPERTY

IMPACT Izod Impact, Notched at 23°C at -29°C Instrumented Impact, Energy at 23°C at -29°C THERMAL Heat Deflection Temp at 0.45 MPa at 1.82 MPa Vicat Softening Point ELECTRICAL Dielectric Strength Dielectric Constant at 1 KHZ Dissipation Factor at 1KHZ Volume Resistivity

D3763

D648

PC

CYCOLOY EHA GEC

BAYBLEND T65MN Miles

PULSE 710 Dow

TRIAX 2153 Monsanto

LEXAN 141 GEC

1,090 0.5

1,120 0.6 0.2

1,120 0.6 0.15

1,130 0.7

1,200 0.6 0.15

2550(370)

2100(305) 75(10.9) 50(7.2) 45(6.5) 80 R118

2412(350) 83(12) 48(7) 45(6.5) 60 R110

2390(340) 87(12.4) 52(7.5) 56 R113

2300(340) 97(14) 60(9) 70(10) 130 R118

370(6.9)

534(10) 320(6)

534(10) 481(9)

800(13) 250(4.6)*

694(13) 100(2)

47(35) 47(35)

45(33) 40(30)

57(42) 57(42)

104 95 113

124 104 121

118 102

14 3.3 0.03 5.0E+13

19.8 2.9 0.03 9.5E+14

45.5(6.6) 40 R111

J (ft-lb) 62(46)

°C

D1525

°C

D149 D150 D150 D257

K/V/mm

ohm-m

115 97 128

140 130 154

* at -40°C

ABS/PC is also used in business machine housings, snow throwers, snowmobiles and other such equipment housings. A continued high growth

rate (ca. 10%) is expected for the ABS/PC blends because of their balance of high impact and heat resistance properties and moderate cost.

1049

Commercial Polymer Blends

15.4.2.3

ABS/Polyamide Blends

Blends of ABS with commercial polyamides such as polyamide-6 and -66 are highly incompatible due to their dissimilar polarity. Such simple blends of ABS and PA exhibit poor delamination resistance and have no practical value. Recent efforts to improve the compatibility led to several different approaches. In one approach, modification of ABS through copolymerization of acrylamide was reported to improve the compatibility between ABS and PA-6, presumably through hydrogen bonding interaction [Grant, 1988]. In another approach, the SAN backbone of ABS was modified through copolymerization with maleic anhydride. This modification introduced controlled amounts of an anhydride functionality on ABS, which upon subsequent melt blending with a PA reacts to form a graft copolymer of SAN and PA which effectively compatibilizes the blend. Commercial blends of ABS with PA-6 and PA-66, introduced by Monsanto under the trade name of Triax® 1000, utilize this reactive compatibilization technology [Lavengood et al., 1987]. In another technique of reactive compatibilization, commercial grades of ABS were directly modified by reactive extrusion with maleic anhydride or fumaric acid and then melt blended with PA-6 and optionally adding small amounts of functionalized EP rubber [Akkapeddi et al., 1990]. Key to the successful compatibilization involves the formation of a graft copolymer of SAN and PA at the interface, which reduces the interfacial tension and improves the interfacial adhesion. Blends with stable, dispersed morphology are achievable. They exhibit high delamination resistance and toughness. Since PA-6 and PA-66 are crystalline polyamides with high melting points, it is desirable to keep them as the continuous phases and ABS as the dispersed phase, for better heat and solvent resistance. The phase morphology of the blend, of course, depends upon the blend ratio and the relative viscosity of the individual components.

The advantages of blending ABS with PA are primarily reduced moisture sensitivity, improved toughness and reduced shrinkage and warpage. Typically the blends exhibit high impact strengths with the notched Izod values of > 600 J/m and better solvent resistance and heat resistance than ABS (Table 15.6). Comparative properties of ABS/polyamide blends and other polyamide blends will be discussed in Part 15.8.1 Polyamide/ Elastomer blends. ABS/PA blends are relatively new and have not yet developed a significant market growth. 15.4.2.4

ABS/Thermoplastic Polyurethane Blends (ABS/TPU)

A recent commercial blend of ABS contains thermoplastic polyurethane elastomer as the main blend component. The blend was introduced in 1990 by Dow Chemical Co., under the trade name Prevail®. These blends characteristically exhibit low modulus (340 to 1000 MPa) and high impact strength at low temperatures, e.g. notched Izod values of 370 to 1500 J/m at -29°C. The TPU component of the blend imparts high toughness and also allows paintability without a primer. ABS component imparts heat resistance (for paint ovens) and good tensile strength in the blend. The blend is projected to find applications in the automotive markets, particularly as paintable, soft bumper fascias. Typical properties of commercial ABS/TPU blends are shown in Table 15.6. 15.4.3

Acrylic-Styrene-Acrylonitrile (ASA) Terpolymer Based Blends

Acrylic-Styrene-Acrylonitrile (ASA) resins, commercialized initially by BASF and subsequently by GEC and others, are normally produced by the graft copolymerization of styrene-acrylonitrile copolymer (SAN) onto an acrylic rubber (usually polybutylacrylate) via emulsion polymerization. The properties of ASA polymers are similar to those of ABS, exhibiting high impact strength, good chemical and heat resistance. However, unlike the ABS resins, the ASA resins exhibit outstanding weatherability, lasting at least 10

1050

Table 15.6.

M. K. Akkapeddi

Comparison of various commercial blends of ABS

Blend Type ASTM TEST METHOD

UNITS

PHYSICAL Density

D792

kg/m3

Mold Shrinkage Water Absorption, 24 hrs.

D955 D570

% %

MECHANICAL Flexural Modulus Flexural Strength Tensile Strength at Yield Elongation at Break Rockwell Hardness

D790 D790 D638 D638 D785

MPa (kpsi) MPa (kpsi) MPa (kpsi) %

D256

J/m(ft-lb/in)

PROPERTY

IMPACT Izod Impact, Notched at 23°C at -29°C Drop Weight Impact at 23°C at -29°C THERMAL Heat Deflection Temp at 0.45 MPa at 1.82 MPa CHEMICAL RESISTANCE

ABS/PVC

ABS/PC

ABS/TPU

ABS/PA-6

ABS/PBT

CYCOVIN K25 B.F. Goodrich

PULSE 710 Dow

PREVAIL 3150 Dow

TRIAX CYCLOLIN 1120 GCM 1900 Monsanto GEC

1,200 0.4 0.09

1,120 0.6 0.15

1,090 0.6

1,060 0.9 0.1

1,120

2310(330) 67(9.6) 40(5.8) 20 R100

2412(350) 83(12) 48(7) 60 R110

1034(150)

2070(300)

28(4) 180 D71

46(6.6) 270 R99

1800(260) 72(10.5) 41(6)

640(12)

534(10) 481(9)

NB 374(7)

850(16) 108(2)

57(42) 57(42)

52(38) 64(48)

>45(33) 44(33)

118 102 Fair

78 63 Good

92 78 Excellent

R103

587(11)

J (ft-lb)

D648

°C

times longer than ABS. The butadiene rubber in ABS is subject to oxidative degradation by the combined effect of atmospheric oxygen and solar radiation. Hence ABS gradually loses its impact strength upon exposure to sunlight. On the other hand, ASA polymers retain their impact strength upon outdoor exposure significantly better than ABS, due to the saturated nature of the acrylic elastomer phase. ASA resins are generally produced as powders for easy blending with other resins such as PVC. ASA is also supplied as

83 79 Fair

93 74 Excellent

pellets for extrusion or injection molding applications. Compared to ABS, the ASA resins are relatively more expensive, hence used only for specialty applications requiring outdoor weatherability and aging resistance. 15.4.3.1

ASA/PVC Blends

Most of the ASA produced is used for blending with a lower cost PVC resin by the end users. The fabricator blends powders of ASA and PVC

1051

Commercial Polymer Blends

in the desired ratio using conventional PVC processing equipment such as Banbury mixers. The blend is used in profile extrusions and coextrusions. Applications for ASA/PVC blend include siding, mobile home skirts, window profiles, vending machine trim, automotive exterior trim, etc. The blend has significantly better impact strength, better heat distortion temperature and better color retention than PVC. The blend has cost advantages over the neat ASA. Typical properties of ASA/PVC blends are described in Table 15.7. 15.4.3.2

ASA/PC Blends

In order to take advantage of the high impact

Table 15.7.

strength and outstanding weatherability of ASA resins, specialty blends based on ASA are being developed recently for exterior automotive applications. Blending polycarbonate with ASA enhances the notched impact strength and DTUL, while maintaining the weatherability characteristics of ASA. Black, pigmented parts such as automotive mirror housings, trim, cowl vents, grilles, etc., are some of the target applications for this blend. It must be remembered that the properties and morphology of ASA/PC would be similar to ABS/PC blends, except for the added feature of outdoor weatherability of ASA. Typical properties of ASA/PC are shown in Table 15.7. Owing to the relatively recent nature of this

Properties of some commercial ASA blends

Blend Type

ASA/PVC ASTM TEST METHOD

UNITS

PHYSICAL Density Mold Shrinkage Water Absorption, 24 hrs.

D792 D955 D570

MECHANICAL Flexural Modulus Flexural Strength Tensile Strength at Yield Elongation at Break Rockwell Hardness

PROPERTY

IMPACT Izod Impact, Notched at 23°C at -40°C THERMAL Heat Deflection Temp at 0.45 MPa at 1.82 MPa

ASA/PC

GELOY 1220 GEC

SUPREL 9301 Vista

GELOY XP4001 GEC

TERBLEND S KR2861 BASF

kg/m3 % %

1,200

1,200

1,110 0.29

1,150 0.45 0.3

D790 D790 D638 D638 D785

MPa (kpsi) MPa (kpsi) MPa (kpsi) %

2140(310) 60(8.5) 46(6.6) 25 R96

2760(400)

1910(295) 68(9.7) 48(7) 70 R99

2250(320) 61(11.2) 53(7.5) 100 R109

D256

J/m(ft-lb/in) 1080(20) 107(2)

425(6)

700(13) 190(3.5)

610(11.2)

110 99

127 108

D648

0.1

43(6.3) R106

°C 85 76

76

1052

M. K. Akkapeddi

technology, little has been published regarding the characterization of this blend. Market volume for this blend has also been thus far, relatively small. Competition from lower cost ABS/PC blend will limit the use of ASA/PC blend to niche applications requiring weatherability and UV resistance. 15.4.4

Styrene-Maleic Anhydride (SMA) Copolymer and Terpolymer Based Blends

Commercial SMA resins (Dylark®, ARCO; Stapron®-S, DSM) are amorphous, random copolymers of styrene and maleic anhydride (10 to 30%) that exhibit higher glass transition temperatures (20 to 60°C higher) and higher heat distortion temperatures than polystyrene. The maleic anhydride comonomer imparts rigidity to the styrenic backbone raising its Tg by 2°C with each 1 wt% increase in the maleic anhydride content. In addition, the maleic anhydride comonomer imparts higher polarity to the styrenic backbone and hence, increasing the modulus, strength, solvent resistance and adhesion to reinforcing fillers and glass fibers [Wambach, 1991]. The polarity and the reactivity of the SMA resins have also been utilized for the compatibilization of its blends with other polymers. Unmodified SMA resins are quite brittle, like the PS and hence they are invariably toughened by incorporating suitable rubber dispersions, either by grafting of a polybutadiene rubber during the polymerizations in a manner similar to the HIPS technology, or by blending with styrene-butadiene multiblock copolymer type elastomers. Rubber toughened SMA resins (Table 15.8) exhibit somewhat higher heat distortion temperatures than HIPS or ABS, yet have an equivalent toughness. Hence such elastomer modified SMA copolymers find good application niches in automotive interior instrument panels, headliners, and in some appliance and business machine housing applications. Terpolymers of SMA containing small amounts of methyl methacrylate comonomer were found to exhibit good miscibility with SAN [Hall, 1982], hence this technology was used to develop blends

of SMA-MMA terpolymers with ABS (Cadon®, Monsanto). These blends exhibit higher DTUL/ impact balance compared to ABS (Table 15.8) and reportedly offer processing and cost advantages over PPE/HIPS and ABS/PC blends [Kosoff, 1987]. Applications for these blends are also in similar markets, e.g. automotive interior, appliance, and business machine housings. SMA copolymers and terpolymers have also been used for blending with PVC to improve the heat distortion temperature and processability of PVC. These blends also contain a rubbery component for impact modification that is usually a high rubber ABS or a polymethyl methacrylate grafted styrene-butadiene rubber (MBS). For improved weatherability, acrylic rubber modified PVC has been used for blending with SMA copolymers and terpolymers (Table 15.4). The market for SMA/PVC blends is still relatively low in volume with only a few applications such as in business machine housings as a low cost replacement for flame retarded ABS. A commercial blend of SMA and polycarbonate (Arloy®, ARCO) was offered for some time, but recently it was discontinued. The polarity of SMA copolymer may account for the good degree of compatibility between the two resins. The blend contained the polybutadiene rubber normally used in SMA resins for impact strength. It exhibited good low temperature impact strength (Table 15.8) and in many properties was similar to ABS/PC blend. A commercial blend of SMA and polybutylene terephthalate (Table 15.8) was also offered for applications requiring solvent resistance, which the crystalline PBT provided. This blend lacked adequate compatibility and hence its impact strength is relatively poor. Some imidized derivatives of SMA resins, containing N-phenyl maleimide units in place of maleic anhydride comonomer units have been offered commercially (Denki Kagaku, Japan) as blend components to modify the heat distortion temperature of styrenics (HIPS and ABS). Blends of such resins with ABS have been commercialized in Japan (Malecca®, Denki Kagaku).

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Commercial Polymer Blends

Table 15.8.

Properties of some commercial blends of SMA copolymer resins

Blend Type

SMA/S-B

SMA/PC

SMA/PBT

SMA-MMA/ ABS

DYLARK 250 Arco

ARLOY 1000 Arco

DYLARK DPN 520 Arco

CADON 127 Monsanto

ASTM TEST METHOD

UNITS

PHYSICAL Density

D792

kg/m3

Mold Shrinkage Water Absorption, 24 hrs.

D955 D570

% %

1,060 0.6 0.12

1,130 0.7 0.2

1,160 0.5

1,070 0.6 0.3

MECHANICAL Flexural Modulus Flexural Strength Tensile Strength at Break Elongation at Break

D790 D790 D638 D638

MPa (kpsi) MPa (kpsi) MPa (kpsi) %

2390(340) 62(8.9) 35(5) 25

2200(330) 80(12) 45(6.6) 80

2760(400) 105(15) 50(7.3) 21

2410(350) 67(9.5) 35(5) 40

D256

J/m(ft-lb/in) 180(3.4)

640(12) 320(6)

38(0.7)

218(4)

113 118

108 138

89

120 112

PROPERTY

IMPACT Izod Impact, Notched at 23°C at -40°C THERMAL Heat Deflection Temp at 0.45 MPa at 1.82 MPa Vicat Softening Point

15.5

D648

D1525

°C

°C

Vinyl Resin Blends

Polyvinylchloride (PVC) is commercially the most significant member of the family of vinyl resins. The other important members of this group are chlorinated-PVC (CPVC) and polyvinylidene chloride (PVDC). PVC is one of the most widely used, commodity type thermoplastics with an annual consumption of over 5 Mton/y in the USA. The excellent versatility of PVC is attributed to its blending capability with a variety of plasticizers, additives and fillers to yield products ranging from very flexible to very rigid types. In addition, PVC has a low cost advantage and a reasonably good balance of properties, which

include (a) general inertness to many chemicals and aqueous fluids, (b) good dimensional stability, (c) good electrical properties, (d) inherent flame resistance and (e) good weatherability. These attributes led to its widespread use in building construction, wire/cable and packaging markets. PVC, however, suffers from an inherent susceptibility to thermal degradation and hence must invariably be processed with heat stabilizers and careful control of processing temperature. The other important drawbacks of PVC are: its brittleness in the absence of a plasticizer (low notched Izod impact strength) and its low heat distortion temperature (ca. 60°C) originating from its low glass transition temperature and essentially

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M. K. Akkapeddi

amorphous character. Hence the primary motivations for blending PVC with other polymers in rigid applications have been to improve its impact strength (notched Izod), DTUL and processability. In flexible, plasticizer-free (free from low molecular type materials) applications, PVC is blended with low modulus polymeric modifiers (as polymeric plasticizers). Impact modified and flexibilized PVC blends together constitute a major fraction of PVC used commercially. 15.5.1

PVC/Impact Modifier Blends

In rigid PVC applications (pipes, building and construction) PVC is compounded with impact modifiers, usually up to 10-15%, to improve the notched Izod impact strength without sacrificing tensile strength and modulus. Various commercial impact modifiers effective in toughening the PVC are illustrated in Table 15.9. These are generally core-shell type rubbers with a controlled particle size, rubbery core (polybutadiene, styrenebutadiene rubber or poly-n-butylacrylate) grafted to a styrene-acrylonitrile (SAN) or styrene-methyl methacrylate (S-MMA) copolymer or PMMA as a rigid outer shell. These are generally made by emulsion polymerization of the corresponding monomers using a pre-made rubber latex as seed or core. The shell (rigid polymer) content in MBS and acrylic core rubbers is usually 20-30% and about 30-50% in ‘high rubber’ ABS. The toughening effect of ABS, MBS and acrylic modifiers is undoubtedly due to the controlled

Table 15.9.

particle size of the elastomer phase [Shaw, 1982; Shur and Ranby, 1976; Walsh and McKeown, 1990]. For high impact strength the particle size should be > 0.1 µm but < 1 µm. The good degree of compatibility or miscibility between PVC and PMMA or SAN phase also plays an important role in the toughening effect. Hence a notched Izod impact strength of > 1000 J/m is readily achieved with these impact modifiers. ABS, MBS and acrylic impact modifiers are available as powders, and are usually blended with PVC powder in high shear intensive mixers prior to fabrication of sheet or profile extrusion [Forger, 1977]. Precompounded grades of PVC are also sold commercially by all the major producers of PVC. When the particle size of the rubber is > 0.2 µm, the toughened PVC is opaque. Clear impact modifiers for PVC are controlled particle size grades of MBS, acrylic core shell rubbers, or ABS. They offer the impact strength improvement as well as maintenance of sufficient optical clarity in the PVC blend. These impact modifiers are designed to match the refractive index of PVC. Controlled particle size (100-300 nm) and sufficient compatibility or solubility of the rigid SAN, S-MMA or PMMA phase with the PVC account for the clarity of the blends. Some grades of MBS have been designed to have cluster-like structures [Saito, 1973], in which the individual rubber particles of small diameter (50-70 µm) are held together by a styrene-methylmethacrylate graft copolymer or terpolymer. When blended with PVC, these rubber particles

Effect of various types of commercial impact modifiers for rigid PVC

Modifier

Advantages

‘High rubber’ ABS (PBD • 50%) Methacrylate-butadiene-styrene (‘MBS’) (PMMA-g-SBR) Acrylate core-shell rubbers (PMMA-g-BuA)

Impact, Processability Clarity, Impact strength Processability Weatherability, Clarity Impact, Processability Weatherability, Impact

Chlorinated polyethylene (CPE)

Notched Izod (J/m) (Modifier level 0, 3, 5, and 12%) 53 53

69 64

203 240

972 1335

53

69

192

1041

54

90

235

1225

1055

Commercial Polymer Blends

are small enough to offer transparency to the blend, but the clusters of these particles in PVC matrix are large enough to cause craze toughening. Typical applications for clear impact modified PVC blends are in clear, calendered sheet or film for packaging and for blown bottles. Acrylic impact modifiers based on poly(n-butyl acrylate) or poly(2-ethyl hexyl acrylate) rubbery cores offer improved weatherability to PVC due to their saturated backbone. These blends are used for outdoor applications such as siding, window profiles, etc. Chlorinated polyethylene (CPE) flexibilizes and toughens PVC more by a miscibility mechanism especially when the chlorine content is above 42% [Donbe and Walsh, 1979]. However, partial miscibility occurring at chlorine levels of 36% leads to higher toughening effects. CPE offers also weatherability advantage, which is the major reason for its commercial use. 15.5.2

PVC/Flexible Modifier Blends

In addition to chlorinated polyethylene, three other major types of modifiers are blended commercially with PVC in order to flexibilize the composition. Usually 30-60% modifier levels are employed. These flexibilizing modifiers usually have some degree of miscibility with PVC. A list of the commercial flexibilizing modifiers and their advantages are shown in Table 15.10.

Ethylene-vinylacetate blends with PVC have been widely used in Europe as permanent “polymeric plasticizers” for PVC. Low cost and good weatherability of these blends permitted their use in window profiles, cable jacketing and other outdoor applications. Ethylene vinylacetate copolymer with 65-75% vinyl acetate content is quite miscible with PVC exhibiting a single Tg for the blend [Hammer, 1971; Ranby, 1975; Rellick and Runt, 1985]. Two phases are apparent when the vinylacetate is ” 45%. Vinylchloride grafted ethylene-vinyl acetate copolymers have also been used for blending with PVC. In PVC/nitrile rubber blends, PVC is added more as an ozone resistant additive. In these thermoplastically processable blends, PVC is flexibilized enough to be used for ‘soft’ goods, wire jacketing, hoses, gaskets and seals. When the NBR contains > 25% acrylonitrile, it becomes miscible with PVC and at ” 20% acrylonitrile level, it is fairly compatible due to partial miscibility [Matso et al., 1969]. 15.5.3

PVC/Styrenic Blends

PVC has been blended with some styrenic resins primarily to achieve some degree of flame retardant characteristics and cost benefits. ABS, SMA, and rubber modified SMA, SMA-MMA copolymers have been used commercially for blending with PVC. These have been discussed under the styrenic blends and illustrated in Table 15.4.

Table 15.10. Effect of various types of flexibilizing modifiers for PVC (Modifier levels: 30 to 60%) Modifier

Advantages

Applications

Ethylene-vinyl acetate (EVAc)

Low smoke generation; Weatherability

Thermoplastic polyurethane (TPU)

Oil and chemical resistance; Low temperature toughness

Construction (windows), sheet/film, cable jacketing Shoes, gaskets, seals, tubing, film

Butadiene-acrylonitrile rubber (NBR): High nitrile type Low nitrile type

Oil and chemical resistance Low temperature flexibility

Cable jacketing, hoses, belting, seals, gaskets, shoe soles, etc.

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15.5.4

M. K. Akkapeddi

PVC/PMMA Blends

PVC blends with polymethylmethacrylate (PMMA) have been commercialized (e.g. Kydex®, Kleerdex Co.) as extruded sheets for thermoforming applications such as chairs, seats, trays, etc. Ease of thermoformability, toughness, resistance to cleaning solvents, and the flame retardancy characteristics of the blends have been the primary features leading to its use. The good level of compatibility between PMMA and PVC is mainly responsible for the toughness characteristics of the blend [Walsh and Cheng, 1984; Tremblay and Prud’homme, 1984; Jager et al., 1983]. Notched Izod impact strength of > 600 J/m has been reported, although it is likely that some acrylic rubber modifier may have been used. The inherent flame retardancy and low smoke-generation characteristics of PVC/PMMA blends meet the aircraft fire safety standards. This factor coupled with the low cost, high toughness and easy processing features of the blend led to its use in aircraft components such as toilet shrouds, floor pans, air diffusers, emergency respirator enclosures, etc.

15.6

Acrylic Blends

Commercial acrylic resins comprise a broad array of polymers and copolymers derived from esters of acrylic acid and methacrylic acid. They range from the homopolymer of methylmethacrylate to a variety of copolymers including both the thermoplastic and thermoset type and ranging from hard and stiff type to soft and elastomeric types. The most common of the thermoplastic acrylic resins are the polymethylmethacrylate homopolymer (PMMA) and the copolymers containing predominantly methyl methacrylate but with small amounts of methyl or ethyl acrylate, acrylonitrile or styrene comonomers added for improved toughness. The commercial PMMA-based acrylic resins are rigid, amorphous polymers (Tg’s ranging from 85 to 105°C) particularly noted for their exceptional clarity and UV resistance. They are therefore widely used for glazing, extruded sheet and

thermoforming applications as well as in several molding applications in which these properties are well utilized. Since most of the applications of the PMMA type acrylic resins are based on their high degree of transparency and UV resistance characteristics, there has been little commercial interest or motivation in developing acrylic blends. This is understandable because unless there is complete, molecular level miscibility between the components, it is not possible to maintain a high degree clarity in the blends. Nevertheless, several examples of commercial blends of acrylic resins are known. These will be discussed under separate headings. 15.6.1

Impact Modified Acrylic Resins

Since the homopolymer PMMA as well as the MMA-rich copolymers are quite brittle, exhibiting low elongation to break (” 5%) and low notched Izod impact strength (typically ” 15 J/m), there was a need to blend suitable impact modifiers that would improve the ductility and impact strength of these resins without sacrificing the transparency, rigidity and weatherability characteristics. Two general types of impact modified acrylic resins have been developed commercially, viz. (a) weatherable, impact modified, transparent acrylic resins for outdoor use in signs and automobiles; (b) non-weatherable impact modified, transparent acrylic resins for medical and food packaging applications. The weatherable grades of acrylics are made by blending ‘all acrylic’ core-shell rubbers viz., PMMA grafted, crosslinked poly(n-butyl acrylate) type rubbers (Paraloid®, Rohm Haas). The non-weatherable grades are made by blending polymethyl-methacrylate-g-butadiene/styrene (‘MBS’) type core-shell rubbers. In both cases, due to the small particle size of these core-shell rubbers (” 100 nm) and the miscibility of the shell (PMMA) with the PMMA matrix, the refractive index could be matched and the transparency could be completely maintained. The rubber particles were found to promote localized shear banding in the matrix and hence, the ductility

1057

Commercial Polymer Blends

and toughness of the matrix improves [Hooley et al., 1981; Bucknall et al., 1984; Wrotecki et al., 1991] Commercial impact modified acrylic resins (Table 15.11) exhibit 5 to 10 fold improvement in the notched Izod impact strength and the ultimate tensile elongation compared to the neat PMMA resin. These impact modified acrylics are

usually blended captively by the manufacturers of the acrylic resins. The base resin in a typical weatherable grade (Plexiglas DR, Rohm and Haas) could be a methylmethacrylate copolymer with ethylacrylate and styrene, while the rubber additive (ca. 10%) could be an emulsion polymerized, PMMA grafted, crosslinked poly(n-butylacrylate) rubber of controlled particle size (” 200 nm).

Table 15.11. Commercial acrylic/impact modifier blends vs. acrylic Blend Type

ACRYLIC/IMPACT MODIFIER

ACRYLIC

ASTM TEST METHOD

UNITS

XT 375 Cyro

PLEXIGLAS DR Rohm & Haas

PHYSICAL Density Mold Shrinkage Water Absorption, 24 hrs.

D792 D955 D570

kg/m3 % %

1,120 0.7 0.3

1,150 0.5 0.4

1,190 0.4 0.3

MECHANICAL Flexural Modulus Flexural Strength Tensile Strength at Yield Elongation at Break Rockwell Hardness

D790 D790 D638 D638 D785

MPa (kpsi) MPa (kpsi) MPa (kpsi) %

2450(350) 77(11) 49(7) 28 M45

1900(270) 72(10.3) 39(5.5) 35 M45

3160(450) 120(17) 72(10.2) 5 M96

D256

J/m(ft-lb/in) 108(2)

60(1.1)

12(0.2)

85

86 79

93 92

2.2 0.03 1.0E+15

2.2 0.03 1.0E+15

PROPERTY

IMPACT Izod Impact, Notched at 23°C at -40°C THERMAL Heat Deflection Temp at 0.45 MPa at 1.82 MPa ELECTRICAL Dielectric Strength Dielectric Constant at 1 MHZ Dissipation Factor at 1MHZ Volume Resistivity

D648

D149 D150 D150 D257

PLEXIGLAS VO44 Rohm & Haas

°C

KV/mm 2.8 0.02 ohm-m

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M. K. Akkapeddi

The non-weatherable impact modified acrylic (XT, CYRO) typically consists of a MMA/S/AN copolymer with MBS (ca. 10%) rubber particle dispersions. Generally, the weatherable impact modified acrylic resin has better color and transparency retention than the non-weatherable grade but the latter shows better toughness. The weatherable grades are used for making outdoor signs, automotive headlight lenses, lighting fixtures, glazing, etc. The non-weatherable, high impact acrylics are used for medical devices, medical and food packaging, refrigerator trays, etc. Polycarbonate is a competitive threat to the impact modified acrylic resin markets. Currently over 30 Mton/y of acrylic/impact modified blends are used in the USA. 15.6.2

PVC/Acrylic Blends

Since PVC is known to be quite miscible with PMMA (miscibility with an LCST behavior) [Jager et al., 1983] and is also low in cost, some blends of PVC and PMMA have been used in sheet extrusion and thermoforming applications. However, the acrylic PVC compositions commercially used are invariably contain an acrylic coreshell rubber (PMMA-g-n-BuA or MBS type) to get high toughness), with some PMMA, to reduce the cost/impact performance balance. The role of PVC in these blends is to reduce cost and impart some degree of flame-retardancy. The acrylics definitely help in the processability of PVC. These blends have already been discussed under PVC heading. 15.6.3

PVDF/Acrylic Blends

Polyvinylidenefluoride (PVDF) and polymethylmethacrylate have been known to be a thermodynamically miscible blend exhibiting a single Tg, an indication of a single phase behavior [Bernstein et al., 1977; Mijovic et al., 1982]. A commercial PVDF/PMMA blend film (Fluorex® A, Rexham Corp.) produced captively by solvent casting is used in specialty applications, such as a protective and decorative overlaminated for

automotive parts. A 50 µm film of the blend is laminated to PVC film which is then laminated to metal, followed by the final fabrication (stamping, rolling) of the automotive parts such as rocker panels, hubcaps, pillar posts, door edge guards, etc. The film is pigmented to a solid or colored metallic look matching the body. The role of PVDF in this blend is to offer chemical and weather resistance to the final surface finish on the part.

15.7

Elastomeric Blends

The blending of different types of rubbers and then curing into the final fabricated parts such as automotive tires has long been known in the rubber industry and will not be discussed here. This discussion will deal with the commercial blends containing a high volume fraction of a rubbery polymer and minor amounts of a rigid, amorphous or crystalline thermoplastic. A major motivation for developing such blends was to combine the elastomeric character of the rubber component with the melt processability of the thermoplastic. Hence blending has been an alternative and somewhat lower cost approach to making thermoplastic elastomers compared to the block copolymer approach. Vulcanized rubbers are distinguished by their characteristically low modulus, high extensibility and high elastic recovery, i.e., by their ability to return to the original dimensions after stretching to high strain levels and then releasing the applied stress. The elasticity behavior in vulcanized rubbers is related to the crosslinking between the polymer chains. However in the thermoplastic elastomers, the elasticity originates from the pseudo-crosslinks formed by the rigid phase, which is either the hard segment of the block copolymer, or the rigid inclusion having high gloss transition temperature, or crystalline polymer phase, blended into the elastomer as a fine dispersion.

1059

Commercial Polymer Blends

15.7.1

Nitrile Rubber/PVC Blends

Blends of butadiene-acrylonitrile copolymer rubber (nitrile rubber or NBR) and PVC are among the oldest known examples of commercial elastomer/ thermoplastic blends. The shortage of natural rubber during World War II stimulated research in the USA on the compounding and modification of synthetic polymers to produce rubber-like materials. An outcome of this research was the commercial introduction of NBR/PVC blends by B.F. Goodrich in 1947 under the trade name of Geon® Polyblends [Pittenger and Cohan, 1947]. The blend showed improved ozone resistance and melt processability compared to the nitrile rubber (Table 15.12). Butadiene-acrylonitrile copolymer rubbers containing > 25% acrylonitrile exhibit good miscibility with PVC as evidenced by single Tg behavior of the blend [Zakrzewski, 1973; Matsuo et al., 1969], although high resolution electron microscopy indicated some degree of microheterogeneity with a very fine dispersion size (< 10 nm) [Matsuo, 1968]. This high degree of miscibility, or partial miscibility, between the components accounts for the blend’s high compatibility and improved mechanical properties. Commercial nitrile rubber/PVC blends are used in both not-vulcanized and vulcanized forms. A descriptive list of commercially available nitrile rubber/PVC blends can be found under a variety of trademarks in the Nitrile Elastomers section of “The Blue Book, Materials, Compounding

Ingredients and Machinery for Rubber,” published annually by Bill Communications, Inc. Currently, the PVC/nitrile rubber blends worldwide consumption is estimated to be 30 kton/y. Nitrile rubbers are known for their oil and chemical resistance and addition of PVC improves the ozone resistance. Use of carboxylated NBR is believed to obviate the necessity of vulcanization. Nitrile rubber/PVC blends have reached a mature stage in their commercial usage. They face increasing competition from other thermoplastic elastomers such as the dynamically vulcanized blends of PP/EPDM and PP/NBR (Santoprene® and Geolast®, Monsanto-Advanced Elastomer Systems). 15.7.2

Dynamically Vulcanized, Alloys of PP and Elastomers

Another type of thermoplastic elastomer blend, which is gaining significant commercial interest is the “dynamically vulcanized” blends of fully cured elastomers in a thermoplastic matrix [Coran, 1987]. Dynamically vulcanized alloys are produced by melt blending a high volume fraction of an elastomer with a thermoplastic in a high intensity mixer, with a compatibilizer if necessary, that a fine dispersion of the elastomer is achieved. The elastomer is then fully cured during the melt mixing through the use of selective crosslinking agents. Since the elastomer is fully cured the blend achieves high rubber elasticity character, but since the thermoplastic still remains as an

Table 15.12. Typical properties of nitrile rubber/PVC blends Property Hardness Modulus (MPa) at 100% at 300% elongation Tensile strength (MPa) Elongation at break (%) Compression set (%) (22 hr, 100°C)

NBR/PVC (70/30) 73A 4.6 12 17 530 48

NBR/PVC (55/45) 82A 7.5 13.5 15.8 460 64

NBR 63A 2 10 22 600 37

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M. K. Akkapeddi

uncrosslinked matrix, the blend can be melt processed like a thermoplastic. In order to distinguish from the simple blends of an elastomer and thermoplastic, the dynamically vulcanized blends have been classified as “elastomeric thermoplastic alloys” (ETA) [Wallace, 1992]. Dynamically vulcanized, elastomeric thermoplastic alloys display properties as good as or even better than the block copolymers, viz. a high degree of rubber elasticity yet good melt processability. The main advantages of the fully cured, elastomeric thermoplastic alloys over the uncured thermoplastic/elastomer blends are: (1) improved tensile strength and elongation, (2) reduced permanent set (compression or tension), (3) improved flexural fatigue resistance, (4) improved chemical resistance, (5) better morphology stability during melt processing, (6) recyclability (regrind reuse), (7) improved melt strength. The dynamically vulcanized blends are very elastomeric in their properties, yet readily melt processable in the conventional injection molding, blow molding, and extrusion processing techniques. Commercially important elastomeric thermoplastic alloys are dynamically vulcanized blends of polypropylene with high volume fractions of EPDM, polybutadiene rubber, nitrile rubber, and butyl rubber (Santoprene®, Vyram®, Geolast® and Trefsin®) all currently sold by Advanced Elastomer Systems, a joint venture of Monsanto and Exxon. Another recent member of the commercial dynamically cured elastomeric thermoplastic alloys is the blend of PVC and a crosslinked ethylene copolymer (Alcryn®, DuPont). The current consumption of all the elastomeric thermoplastic alloys in the USA is over 23 kton/y, with the EPDM/PP blend (Santoprene®) assuming about 90% of the market share. 15.7.2.1

Dynamically Vulcanized EPDM/PP Blends

Owing to adequate level of compatibility between polypropylene and ethylene-propylene copolymers, simple blends of these two polymers have been known for a long time. Impact modified

polypropylene blends containing up to 30% EPR have already been discussed under polyolefin blends. Blends containing high content of an uncured or partially cured EPDM in polypropylene have been known [Kresege, 1978]. However, the advantage of fully cured EPDM/PP blends made by selective crosslinking of the rubber phase during the melt mixing has not been commercially realized until recently (Santoprene®, AES; Sarlink® 3000, Polysar). The technology [Coran et al, 1978; 1980], involved an accelerated sulfur cure of the EPDM rubber (• 70%) in the presence of polypropylene (” 30%) while melt compounding under high shear mixing conditions. The curing agents typically consisted of zinc oxide (5 phr), sulfur (2 phr), tetramethylthiuram disulfide (1 phr), and 2-benzothiazolyl disulfide (0.5 phr). The particle size of cured EPDM dispersions was typically < 2 µm. Increasing the crosslink density of the elastomer dispersion results in improvements of the strength and tension set of the blend. The difference between the earlier commercial grades of partially cured EPDM/PP blends (TPR, Uniroyal) and the more recent commercial grades of completely cured EPDM/PP blend is in the improved elastomeric properties, viz. reduced compression and tension set and improved flexural fatigue. More important, the chemical resistance and resistance to oil swelling is improved. Typical properties of commercial dynamically vulcanized EPDM/PP blend (Santoprene®) are shown in Table 15.13. Nearly all of the dynamically vulcanized EPDM/PP blend’s (Santoprene®) growth has been at the expense of thermosetting rubbers due to the blend’s easy melt processability and recyclability. Much of this growth has been in automotive applications such as rack and pinion steering boots, seat belt sleeves, and air ducts. EPDM/PP blend has also been used in window and door glazing seals, weather stripping and extrusion applications in the electrical/ electronic industry.

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Commercial Polymer Blends

Table 15.13.

Key properties of commercial thermoplastic elastomer blends based on polypropylene/elastomer dynamic vulcanizates

Properties

ASTM Method

Specific Gravity (g/mL) Shore Hardness 100% Modulus (MPa) Tensile Strength (MPa) Elongation at Break (%) Tear Strength (KN/m) Tension Set (%) Compression Set (%) Brittlenss Temp. (oC) Flexural Fatigue1 Oil Swelling2

PP/EPDM (Santoprene®)

PP/PBD (Vyram®)

PP/NBR (Geolast®)

D792 D2240 D412 D421 D624 D624 D412 D395

0.98 73A 3.2 8.3 375 28 14 24

1.04 70A 3.7 7.6 380 31 16 32

1.0 70A 3.3 6.2 265 33 10 29

D746

-63 >3.4 32

-58

-40

32

10

PP/Butyl (Trefsin®) 1.2 70A 5.2 7.6 250 26 12 52 60

Notes 1In megacycles to fail; 2As vol% in ASTM oil #3, after 70h at 100°C.

15.7.2.2

Polybutadiene Rubber/PP, Dynamically Vulcanized Blends (Vyram®, AES)

This class of blends was introduced in 1989, as a lower cost alternative to EPDM/PP blend. Its heat-aging resistance and use temperature limits are, however, inferior to the EPDM/PP blends. 15.7.2.3

Nitrile Rubber/PP, Dynamically Vulcanized Blends (Geolast®, AES)

The oil resistance and chemical resistance of nitrile rubber is generally superior to that of EPDM rubbers. However, the highly polar nature of acrylonitrile comonomer is responsible for the high incompatibility between nitrile rubber and polypropylene. The dispersability and the stability of nitrile rubber dispersions in the polypropylene matrix are poor. Hence a reactive compatibilization technology was used [Coran and Patel, 1983]. It consisted of blending a small amount of a low molecular weight amine-terminated butadiene-acrylonitrile copolymer (ATBN, B.F. Goodrich) with the high molecular weight

nitrile rubber and a small amount of maleic anhydride grafted polypropylene (maleated PP) with the high molecular weight polypropylene and then intimately mixing all the components together. A graft coupling reaction presumably takes place between the ATBN and maleated PP. The resulting poly(butadiene/acrylonitrile)-g-polypropylene copolymer compatibilizes and reduces the interfacial tension between the nitrile rubber and polypropylene phases. Typically a 50/50 blend of nitrile rubber and polypropylene is melt mixed with 5% maleated PP and 1% ATBN respectively and then cured with SnCl2 (0.5%). The resistance to hot oil swell (72 hr, 100°C) of NBR/PP blend is significantly better than that of EPDM/PP blend. Typical properties of the commercial dynamically cured NBR/PP blend (Geolast®) are compared with other such blends in Table 15.13. Commercial applications for this blend include seals and gaskets in the automotive oil, fuel and brake systems. Other automotive uses include air ducts, boots, connectors, etc. Geolast® is also used in gasoline powered lawn and garden equipment, power tools and other industrial applications.

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M. K. Akkapeddi

Butyl rubber/PP Dynamically Vulcanized Blends

The dynamically cured, butyl rubber/polypropylene blends were first developed by Gessler et al. using phenolic type crosslinking agents. Commercial dynamic vulcanizates are based on halobutyl rubbers and polypropylene cured by zinc oxide type curatives. [Hazelton and Puydak, 1987; Kay and Ouhadi, 1991; Anon., 1988]. The blend containing a high volume fraction of butyl rubber dispersion in polypropylene exhibits high elastomeric properties but retains good thermoplastic processability. In addition, this blend exhibits the following advantages: (i) Low permeability to air and moisture (ii) High energy absorption and consequently high vibration damping characteristics (iii) Resistance to thermal and UV aging and weathering (iv) Excellent resistance to common organic fluids. Typical applications are in medical stoppers, bladders and packaging seals. Other automotive and non-automotive applications are expected. 15.7.3

Elastomeric Thermoplastic Alloys of PVC and Ethylene Terpolymers

Recently blends of certain types of ethylene terpolymers with PVC, which have been dynamically vulcanized into highly elastomeric yet thermoplastically processable blends, have been offered commercially (Alcryn®, DuPont). The principle behind this technology appears to be the selection of proper types of ethylene copolymers and terpolymers which exhibit high degree of miscibility with PVC and which can selectively be crosslinked in-situ, during the melt mixing with specific curing agents [Loomis and Statz, 1986]. It has been known that PVC forms miscible or partially miscible blends with certain types of ethylene copolymers such as ethylene-vinyl acetate (EVAc), ethylene-methyl acrylate (EMA) and ethylene-butyl acrylate (EBA) copolymers [Krause, 1989]. The degree of miscibility depends on the structure and ratio of ethylene to the

comonomer units. For example, when the vinyl acetate content is high (e.g., > 65 wt%) the copolymer forms completely miscible blends with PVC [Hammer, 1971; Rellick and Runt, 1985]. At lower levels, the blends are partially miscible. Similarly ethylene-methyl acrylate and ethyl-butyl acrylate copolymers form miscible to partially miscible blends with PVC [Kalfoglou, 1983]. Ethylene copolymers, which exhibit partial miscibility, can be expected to be dispersed into fine particles in a matrix of PVC by simple melt mixing, because of the self-compatibilizing nature of the blends. If a suitable third comonomer is also present in these copolymers, which contains a reactive functionality suitable for crosslinking, one can use selective crosslinking techniques through dynamic vulcanization process to achieve high volume fraction dispersions of cured ethylene copolymers in the thermoplastic matrix of PVC. This principle has been commercially employed in developing thermoplastic elastomer blends of PVC with such ethylene terpolymers as ethylene-n-butylacrylate-carbon monoxide (E-BA-CO) and ethylene-vinylacetate-carbon monoxide (E-VAc-CO) terpolymer systems. Due to the presence of the carbonyl moiety in their backbone, these ethylene terpolymers can be crosslinked selectively through the use of peroxide/bis maleimides or diamine type curatives or through electron beam or gamma radiation techniques [Loomis and Statz, 1986]. In a typical formulation, an ethylene-n-butylacrylate-carbon monoxide (60/30/10) terpolymer (60 wt%) is melt compounded with plasticized PVC (40 wt%) in a twin-screw extruder and the ethylene terpolymer dispersion cured in situ during the mixing by catalytic amounts of a suitable peroxide (0.3%) and a bismaleimide crosslink promoter (0.2%). The extruded pellets of the elastomeric blend can be used in conventional melt fabrication processes such as profile extrusion, extrusion coating, milling and calendering of sheets, injection and/or compression molding. Commercial elastomeric blends of such ethylene terpolymers and PVC have been reported to have the following advantages: (i) outstanding weatherability and ozone resistance, (ii) excellent oil

1063

Commercial Polymer Blends

resistance, (iii) good low temperature toughness, (iv) good melt processability and recyclability. In Table 15.14, the typical properties of a commercial elastomeric blend of ethylene terpolymer and PVC (Alcryn®) are compared with similar elastomeric blends. Commercial elastomeric blends of ethylene terpolymer/PVC system (Alcryn®) are used in outdoor weather stripping, seals and gaskets, coated fabrics, pond linings and a variety of other extruded and molded goods for automotive and industrial applications.

15.8

Polyamide Blends

Commercial polyamides, frequently referred to as nylons, are crystalline engineering thermoplastics exhibiting high performance characteristics such as high melting points, high mechanical strength, ductility, and excellent resistance to solvents, fatigue and abrasion. Nylon is a generic term used for all synthetic polyamides in which the recurring amide groups (-CONH-) are part of the main polymer chain. These amide groups impart strong hydrogen bonding capability and crystallinity in the polyamides, PA, which account for

their outstanding solvent resistance and mechanical properties. Commercial polyamides are generally of two type (a) those derived from diamines and dicarboxylic acids and (b) those derived from amino acids or lactams as monomers. The major characteristics of these two types of polyamides are similar since these are determined largely by the hydrogen bonding structure of the amide groups. However, within these two types, a wide variety of PA are known, varying in their melting points and moisture absorption characteristic, depending on their structure. Among these, PA-66, a polyamide made by polycondensation of hexamethylene diamine and adipic acid and PA-6, a polyamide made by the ring opening polymerization of caprolactam, are the two major nylon engineering thermoplastics produced commercially. Because of their widespread use in fiber, plastic and films, both PA-66 and PA-6 are produced on a large scale with an estimated total volume globally to be > 4.5 Mton/y, 18% of which is used for non-fiber applications [SRI, 1992]. In 1993, an estimated 355 kton/y of PA will be used in the USA for molding applications. PA blends consumed in the USA are currently about 30 kton/y.

Table 15.14. Comparison of the typical properties of the dynamically vulcanized ethylene terpolymer/PVC blends vs. similar elastomeric blends based on polypropylene Property

Density (kg/m3) Shore Hardness Tensile Strength (MPa) Elongation at Break (%) Tension Set (%) Compression Set (%) (at 100°C) Brittleness Temperature (°C) Hot Oil Resistance Weatherability (Ozone & UV Resistance)

Ethylene terpolymer / PVC (Alcryn® 2070)

EPDM/PP (Santoprene® 101-73)

NBR/PP (Geolast® 701)

1,140 73A 8.9 680 9 65 -78 Excellent Excellent

980 70A 7.6 470 14 33 -65 Fair Excellent

1,000 68A 6.2 480 16 28 -40 Excellent Poor

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M. K. Akkapeddi

There are a number of other specialty polyamides produced from a combination of other diamines and dicarboxylic acids and/or lactams of varying number of carbon atoms. PA-11 and PA-12 with 11 and 12 methylene units between each repeat amide group are relatively low melting point (170°C), but exhibit excellent ductility and moisture resistance. Among the newer class of polyamides are the high melting (ca. 300°C) PA-4,6 and PA-6T copolymers with 6 or 66 or 6T monomer units. Because of the high level of crystallinity and high melting points, PA’s generally exhibit high heat distortion temperature at low loads even in unfilled form, and when reinforced with glass fibers, exhibit high heat distortion temperatures at high loads. Most of the commercial polyamides exhibit a common set of property advantages attributable to their crystalline nature and hydrogen bonding character. The advantages offered by the crystalline polyamides (PA-6, PA-66) in blends with other polymers, are: excellent solvent resistance (e.g., gasoline, oils, paint solvents, etc.), heat resistance and melt flow characteristics. On the other hand, the primary motivation to blend other thermoplastic polymers with polyamides, is for reducing the moisture sensitivity of PA and improving its dimensional stability and toughness. Because of their highly polar and hydrogen bonded structure of the backbone, as a general rule polyamides are immiscible with most of the commercially known polymer systems. In addition, the high degree of interfacial tension [Wu, 1989] between polyamides and other classes of polymers leads to highly phase separated blends with poor delamination resistance. Hence simple blends of PA with other commercial polymers generally do not have any practical value. Significant progress has been made in recent years in developing techniques for compatibilizing polyamide blends, particularly utilizing the reactivity of polyamide end groups in forming ionic or covalently linked bonds with other polymers at the blend interface. Several commercial blends are based on such reactive compatibilization technology. Some of the more important

polymer candidates commercially used for blending with PA and the major reasons for blending them are listed in Tables 15.15 and 15.16. The graft copolymer compatibilization technique requires the other polymer component to be already “functionalized”, i.e., modified with functionalities such as anhydride or epoxide groups, which are reactive towards the amine or carboxyl end groups of PA respectively. During the melt blending, the reaction between the functionalized polymers and polyamides leads to graft copolymer formation at the interface, which essentially compatibilizes and stabilizes the blend against delamination. 15.8.1

Polyamide/Elastomer Blends (Impact modified polyamides)

Commercial polyamides such as PA-6 and PA-66 are generally regarded as tough and ductile materials since they exhibit high tensile elongation to break and high drop weight impact strengths. They become even tougher after equilibration with ambient humidity, due to the plasticization effect of the absorbed water. However, under conditions of stress concentration such as in the presence of sharp notches or cracks, polyamides exhibit brittle failure. This property, commonly evaluated as notched Izod or Charpy impact tests, indicates that unmodified polyamides exhibit relatively low energies for crack propagation. To overcome this deficiency, polyamides have been blended with several types of impact modifiers that are typically elastomeric or low modulus type olefinic polymers. However, the inherent immiscibility of polyamides with other polymers such as olefinic rubbers necessitated the development of proper compatibilization techniques to reduce the interfacial tension and improve the dispersability of the rubber for effective impact modification. The technology for impact modification of polyamides has evolved significantly over a period of several years through improved methods of compatibilization and particularly through reactive blending techniques [Kray and Bellet, 1968; Murch, 1974; Epstein, 1979; Mason and Tuller, 1983]. Several commercially successful

1065

Commercial Polymer Blends

Table 15.15. Effect of various blend components on the properties of polyamides Blend Component

Reasons for blending

Compatibilization method

Elastomers (functionalized EPR, acrylate rubbers, etc.) Ethylene copolymers, ionomers

* Improve notched Izod impact strength * Shift ductile brittle transition to lower temperatures * Improve toughness and flexibility * Lower modulus * Lower cost * Improve dimensional stability with humidity * Toughness * Reduce moisture sensitivity (better strength and stiffness retention) * Improve DTUL (at 1.83MPa), unfilled * Improve strength and creep resistance * Reduce moisture sensitivity (better strength and stiffness retention with humidity) * Improve DTUL (at 1.83MPa), unfilled * Lower mold shrinkage * Reduce moisture sensitivity * Improve lubricity and wear * Reduce shrinkage and warp * Improve oxygen barrier at high humidity levels

Graft-coupling reaction

Polyolefins (HDPE, LDPE, PP) ABS

Polyphenylene ether, PPE

Polycarbonate (PC)

Silicon IPN Amorphous polyamide (High Tg, barrier)

Polar interaction Graft-coupling reaction Graft-coupling reaction

Graft-coupling reaction

Maleated-PP, polyether amide elastomer None (IPN cured in situ) None (miscibility)

Table 15.16. Comparative rating of performance characteristics of polyamideswith other key blend components

Melting/softening temperature Glass transition temperature Melt processability Moisture absorption Moisture sensitivity (properties, dimensions) Drop weight impact strength Notched Izod impact Tensile strength Creep resistance Solvent resistance Hydrocarbons, oils Paint solvents (MEK, toluene) Alcohols, glycols Hydrolysis resistance (in acidic or basic aqueous fluids)

PA-6, PA-66

Polypropylene

ABS

PPE

High Low Excellent High High High Low High (dry) Moderate (wet) Good

Moderate Low Excellent None None Moderate Low Moderate

Moderate Moderate Excellent Low None High High Moderate

Very High High Poor None None Moderate Low High

Low

Moderate

Excellent

Excellent Excellent Poor Poor

Moderate Moderate Excellent Excellent

Moderate Moderate Excellent Moderate

Poor Poor Excellent Excellent

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M. K. Akkapeddi

impact modified polyamides are based on blends of polyamides with (a) reactive elastomers such as maleic anhydride-grafted (“maleated”) EPDM, EP and styrene-ethylene/butylene-styrene block copolymer rubbers, and (b) functional ethylene copolymers such as ethylene-ethyl acrylate, ethylene-acrylic acid, ethylene-ethylacrylate-maleic anhydride and ethylene-methacrylic acid ionomers. Compatibilization of an olefinic rubber dispersion in a polyamide melt blend is achievable through a direct chemical coupling reaction between the polymers at the interface such as through the addition reaction between the amine end groups of polyamide and the anhydride functionality of a maleated EP rubber. The graft copolymer formed in situ via this reaction during the melt blending process effectively compatibilizes the blend by reducing the interfacial tension and increasing the adhesion at the phase boundary. Due to the graft copolymer’s capability to act as an interfacial agent, the dispersability of the rubber in the polyamide matrix improves considerably, resulting in well stabilized, reduced particle size (” 1µm), rubber dispersions and thereby substantially increasing the toughness. DuPont’s commercially successful ‘super-tough’ nylon, Zytel® ST801 is based on such a reactive blending technology using PA-66 and a maleated EPDM rubber [Epstein, 1979]. Commercial ‘super-tough’ PA-66 exhibits an excellent notched Izod impact strength (• 900 J/m) that is remarkably insensitive to part thickness and notch radius. Impact modification of PA-6 has also been achieved by blending with such ethylene copolymers as ethylene-ethylacrylate, ethylene-acrylic acid and ethylene-methacrylic acid copolymer based zinc ionomers. It is believed that favorable, associative interactions involving some type of complexation between the amine end groups of polyamide and the zinc carboxylate groups of the ionomer are responsible for the high compatibility and toughening efficiency of ionomers. Some commercial impact modified PA-6 blends (Capron® 8253 and 8350, AlliedSignal) are based either on ethylene copolymer or ionomer type tougheners [Mason and Tuller, 1983].

There are several other routes to compatibilizing polyamide/elastomer blends for the purpose of impact modification such as through the use of anhydride modified ABS rubbers [Baer, 1988], anhydride modified S-EB-S block [Gelles et al., 1988] copolymers, carboxylated core-shell rubbers (MBS or acrylic type) [Liu, 1988] and acylcaprolactam grafted EP rubbers [Akkapeddi and Haylock, 1989]. The outstanding impact toughness of the commercial impact modified polyamides is attributed to the small particle size of rubber dispersion and their good degree of adhesion to the polyamide matrix. Typical morphologies of compatibilized polyamide/elastomer blends are shown in Figure 15.6. In all of these blends the rubber particle size is much smaller than would be obtained with a typical unmodified EPR phase. The toughening mechanism is believed to involve the internal cavitation and debonding of the rubber which induces localized shear yielding of the polyamide matrix as the primary energy dissipation processes [Borggreve and Gaymans, 1989; Ramsteiner and Heckmann, 1985; Borggreve et al., 1988] occurring during the impact deformation. Rubber particle size, distribution and interparticle distance [Wu, 1988] are some of the key parameters that have been correlated to the impact toughness. The ductile-brittle transition of the polyamide blend is shifted to low temperature, both by choosing a low Tg rubber, as well as controlling the rubber particle size (” 1 µm) and volume fraction. At an equivalent content of the compatibilized rubber, PA-6 tends to show higher impact strength at low temperatures than PA-66. Commercial impact modified polyamides typically contain 10 to 25% of the reactive or compatible elastomer to maximize the toughening efficiency while retaining a high level of tensile strength and DTUL. Commercial impact modified PA blends (Table 15.17) indeed offer a unique combination of high notched Izod impact and drop weight impact strengths, coupled with a good balance of modulus, tensile strength, heat, solvent and abrasion resistance characteristics. These properties are suitable for many engineering and metal replacement applications. Because

Commercial Polymer Blends

Figure 15.6. Morphology of typical Polyamide/impact modifier blends — TEM., Phosphotungstic acid stain, top — PA-6/ethylene copolymer/ionomer blend (21,000X); bottom — PA-6/maleated EPR (3:1) blend (30,000X).

of their good melt flow characteristics most of the impact modified polyamides are processed by injection molding, although recently blow molding applications are also emerging. Typical automotive applications include gas cap covers, fan blades, shrouds for cooling systems, gears, fasteners, clips, emission control canisters, oil pans, shields, etc. Toughened polyamides are also used in sports and recreational equipment, lawn and power tool housings and components where impact strength, heat and chemical resistance are the primary criteria for their choice. With an estimated consumption of 25 Mton/y in the USA, impact modified PAs

1067

constitute the largest fraction of PA blends used commercially. Some PA/elastomer blends containing a combination of different types of impact modifiers and elastomers have been specifically designed for low volume, specialty extrusion applications (tubing and profiles). Generally, these blends contain higher levels of these elastomers or impact modifying polymers than the impact modified polyamides. Often the rubbers are used in combination with plasticizers to achieve a flexible, low modulus blend product that still retains much of the polyamide advantages, viz. heat resistance and solvent resistance, particularly permeation resistance to such automotive fluids as gasoline, refrigerants (CFC, HCFC), etc. These blends are often formulated to suit a specific customer or application requirement. The demand for such polyamide based flexible blends seems to be growing. Recently a flexible polyamide blend [Saltman and Varnell, 1988; Saltman, 1992] containing < 50% PA-6 or PA-66 and >50% acrylic elastomers has been introduced by DuPont (Zytel® FN) as a plasticizer-free, low modulus composition with good low-temperature toughness, resistance to thermal aging and solvents (particularly to fluorocarbon refrigerants) (Table 15.17). Although the flexible polyamide blends contain < 50% polyamide, the PA phase is continuous and the major elastomeric phase (> 50%) is kept as dispersion. It is believed that the morphology can be controlled by modifying the viscosity of the elastomer phase during the blending through controlled and selective grafting reactions. By using two mutually miscible (or compatible), co-reactive elastomers, a high viscosity ratio between the elastomer and polyamide phases can be maintained preventing phase inversion even when the total volume fraction of elastomer phase is higher than that of polyamide. For example, by using an ethylene-butylacrylate-methacrylic acid ionomer in combination with an ethylenebutylacrylate-glycidyl methacrylate (E-BA-GMA) terpolymer, a grafting or cross-linking reaction between the two elastomers via the carboxy/ epoxide addition reaction is expected to take place, during the melt blending with polyamide.

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M. K. Akkapeddi

Table 15.17. Properties of some commercial polyamide/elastomer blends Blend Type

PA-66 BASED ASTM TEST METHOD

UNITS

PHYSICAL Density

D792

kg/m3

Mold Shrinkage Water Absorption, 24 hrs.

D955 D570

% %

MECHANICAL Flexural Modulus Flexural Strength Tensile Strength at Yield Elongation at Break Rockwell Hardness

D790 D790 D638 D638 D785

MPa (kpsi) MPa (kpsi) MPa (kpsi) %

D256

J/m(ft-lb/in)

PROPERTY

IMPACT Izod Impact, Notched at 23°C at -40°C Drop Weight Impact at 23°C at -40°C THERMAL Heat Deflection Temp at 0.45 MPa at 1.82 MPa Vicat Softening Point

D3029

D648

ZYTEL 408 Du Pont

ZYTEL ST801 Du Pont

ZYTEL FN716 Du Pont

CAPRON 8253 AlliedSignal

CAPRON 8351 AlliedSignal

1,090 1.5 1.2

1,080 1.5 1.2

1,030

1,090 1.2 1.5

1,070 1.3 1.1

1960(285) 80(12) 62(9) 80 R115

1690(245) 68(9.8) 52(7.5) 40 R112

690(100) 30(4.3) 250 D59

2200(320) 85(12.5) 65(9.5) 150 R78

1800(240) 65(9.5) 55(8) 200 R82

240

937(17) 220(4)

NB 270(5)

135(2.5) 70(1.3)

990(18) 265(5)

170(125) 170(125)

135(100) 100(75)

200(150) 200(150)

216 71

176 60 210

170 60 210

0.5

J (ft-lb)

°C 230 73

D1525

PA-6 BASED

°C

216

* Elastromers are functionalized EPR or ethylene copolymers

This would cause an increase in the melt viscosity of the elastomer phase relative to polyamide, thus preventing a phase inversion [Saltman, 1992; Ohme, 1991]. At the same time, the compatibility between the ionomer and polyamide as well as some degree of reactive compatibilization between E-BA-GMA and polyamide (through carboxy/epoxide or amine/epoxide reactions) may lead to the stabilization of the dispersions. The key advantage of this stabilized, high rub-

ber blend dispersion morphology of PA appears to translate into a significant improvement of impact strength and elongation at break after heat-aging at 150°C for 14 days. Because of the saturated and polar nature of the rubber and the continuous matrix of polyamide, the blend retains a high degree of solvent resistance in addition to its aging resistance. Since flexible polyamide blends are relatively new, their uses are still emerging. Typical applications are expected to be in the

1069

Commercial Polymer Blends

extrusion and blow molding areas, e.g., tubing and hoses; ducts and air-intake systems; oil and grease seals, etc. 15.8.2

Polyamide/Polypropylene Blends

Polyamide and polypropylene are both crystalline polymers, but are significantly different in their structure and polarity. Hence they are immiscible. The primary motivations for blending polypropylene with polyamide appear to be based on some cost advantages and some improvements in dimensional stability in the presence of moisture. Although compatibilization of polyamide and polypropylene via grafting with a maleic anhydride modified polypropylene was known for a long time [Ide and Hasegawa, 1974], commercial interest in such polyamide/polypropylene blends has not developed until recently [Girard, 1990; Moody, 1992]. It is known that polypropylene can be modified by a free radical catalyzed grafting reaction with maleic anhydride (maleation) in the melt phase under extruder processing conditions. The reaction involves hydrogen abstraction from polypropylene backbone followed by chain scission. The polypropylene chain fragments with secondary radical or unsaturated chain ends undergo an addition reaction with maleic anhydride forming a modified polypropylene with anhydride functionality bound to the polymer chains. The process involves some chain degradation that must be controlled by amount of peroxide used and the temperature and mixing conditions used. Typically, about 0.5% maleic anhydride can be grafted to PP. Similarly, other unsaturated anhydrides can be used to modify PP. Melt blending of PA-6 (or 66) with such an anhydride functionalized polypropylene causes a fast graft copolymer reaction between the polyamide and PP at the interface, which subsequently compatibilizes the blend. Some commercial polyamide/polypropylene blends may utilize such types of reactive compatibilization techniques. Properties of commercial PA/PP blends, both unfilled and glass filled grades, are shown in Tables 15.18 and 15.19. Typically, these blends

are in 50/50 to 60/40 ratio. Since both polyamide and PP are crystalline polymers, one would expect notch sensitivity and brittleness under impact loading conditions, despite the compatibilization. Accordingly, the polyamide/PP blends show about the same low notched Izod impact strengths as the individual components. However, using PP copolymers, or an atactic PP of low crystallinity or by using a PP impact modified with EPR, higher notched Izod impact strength in the polyamide/PP blends can be achieved. All other properties such as modulus, strength and DTUL are intermediate between the individual components, as expected. Polyamide/PP blends exhibit a significantly slower rate of moisture absorption compared to the polyamides due to the presence of the moisture resistant polypropylene phase. The morphology of PA/PP blend depends on the molecular weights of polyamide and PP (especially after functionalization and blending) as well as the degree of grafting. Typically, commercial polyamide/PP blends show nylon as the continuous phase and polypropylene as the dispersed phase. The morphology, however, seems somewhat sensitive to the flow induced shear effects during the injection molding. Commercialization of polyamide/PP blends is still at an early stage. One of the commercial sources in the USA (D&S International) offered two kinds of polyamide/PP blends, one rich in PP (DexPro®) and the other rich in polyamide (Dexlon®). The blend’s improved dimensional stability over polyamide has led to some applications in automotive, lawn and power tool markets. The relative advantages of polyamide/PP blends vs. polyamide, PP and other blends have not been clearly identified and further market and application development is actively under investigation with several companies. 15.8.3

Polyamide/ABS Blends

As indicated in Table 15.15, the advantages of blending ABS with PA are primarily in impact strength and moisture resistance. Since ABS is an amorphous polymer, its heat resistance is limited by the Tg of the SAN phase, thus the blend would be expected to exhibit lower heat resistance than

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M. K. Akkapeddi

Table 15.18. Properties of some commercial polyamide/polypropylene blends vs. polypropylene Blend Type

PA-6/PP ASTM TEST METHOD

UNITS

PHYSICAL Density

D792

kg/m3

Mold Shrinkage Water Absorption, 24 hrs.

D955 D570

MECHANICAL Flexural Modulus Flexural Strength Tensile Strength at Yield Elongation at Break

PROPERTY

IMPACT Izod Impact, Notched at 23°C THERMAL Heat Deflection Temp at 0.45 MPa at 1.82 MPa

PA-66/PP

PP

ORGALLOY R6000 Atochem

AKULOY RM NY-75 DSM

DEXLON 602 D&S Int.

PROFAX 6523 Himont

1,030 0.79 0.4

1,030 0.9 0.4

1,090

% %

900 1.25 0

D790 D790 D638 D638

MPa (kpsi) MPa (kpsi) MPa (kpsi) %

2050(300) 68(9.9) 43(6.2) 300

2140(310)

1610(230)

48(7) 150

1725(250) 76(11) 48(7) 90

D256

J/m(ft-lb/in) 90(1.6)

77(1.4)

850(16)

52(0.95)

140 60

152 65

216 65

91 53

D648

1.1

35(5) 200

°C

polyamide. However, by keeping the polyamide as a continuous phase and ABS as the dispersion blends with high softening point (due to the high melting point of polyamide) high HDT can be achieved. In addition, a polyamide matrix would be advantageous for maintaining solvent resistance in the blend. Simple blends of ABS and PA are highly immiscible and hence are of little practical value. Compatibilization of ABS with polyamide was accomplished by several methods, most of which involving structural modification of ABS. In one approach, ABS was modified by copolymerization with acrylamide, during the preparation of ABS by the standard emulsion polymerization. The introduction of polar acrylamide units on the

SAN backbone of ABS in sufficient concentration caused compatibilization with PA-6 when melt blended, presumably due to favorable hydrogen bonding interactions [Grant, 1985]. In another approach, a small amount of maleic anhydride was copolymerized with styrene and acrylonitrile during the preparation of ABS by emulsion polymerization. The ‘anhydride modified’ ABS was then melt blended with polyamide to form a compatibilized ABS/PA blend [Lavengood et al., 1986, 1987; Howe and Wolkowicz, 1987]. Obviously, a reaction between the anhydride functionality of ABS and the amine end group of polyamide leads to an in situ graft copolymer responsible for compatibilizing this blend. Monsanto’s ABS/PA blends (Triax® 1000

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Commercial Polymer Blends

Table 15.19. Properties of 30% glass reinforced polyamide/polypropylene blends vs. 30% glass reinforced polypropylene Blend Type

PROPERTY

PHYSICAL Density

PA-6/PP ASTM TEST METHOD

UNITS

D792

kg/m3

AKULOY RM J-75/30 DSM

MCX-Q QA5030 Mitsui

DEXLON 633HI D&S Int.

1,260 0.3 0.35

1,300

1,400

% %

D790 D790 D638 D638

MPa (kpsi) MPa (kpsi) MPa (kpsi) %

7600(1100) 160(23) 120(17.5) 4

7500(1090) 190(27.5) 130(19) 4

7600(1100)

D256

J/m(ft-lb/in) 90(1.6)

165

Mold Shrinkage D955 Water Absorption, 24 hrs. D570 MECHANICAL Flexural Modulus Flexural Strength Tensile Strength at Yield Elongation at Break IMPACT Izod Impact, Notched at 23°C THERMAL Heat Deflection Temp at 0.45 MPa at 1.82 MPa

PA-66/PP

D648

0.4

PP

AKULOY RM COMALLOY M-2040 130 DSM Comalloy

1,280 0.5

1,130 0.4 0

5340(760)

162(23.5) 3

8100(1170) 190(27.5) 135(19.5) 4

100(1.8)

133(2.5)

115(2.1)

82(1.5)

190

199

242

151 143

70(10) 3

°C

series) are based on this reactive compatibilization technology. The polybutadiene rubber content of the ABS used for this blend technology is about 40%, significantly higher than the standard grade ABS (10-20% PBD). Because of the high rubber content and the compatibilization chemistry, the blends exhibit excellent notched Izod impact strengths (Table 15.20). Standard grade ABS (20% PBD) could also be directly modified by grafting an anhydride functionality via reactive extrusion with maleic anhydride in the presence of a trace amount of peroxide initiator [Akkapeddi et al., 1990]. Such an extruder ‘maleated ABS’ upon subsequent melt blending with an amine terminated PA-6, gave a compatibilized ABS/PA-6 blend with significantly

improved impact strength (Figure 15.7) and a fine dispersion morphology (Figure 15.8). Undoubtedly the chemical coupling between the phases via the amine/anhydride reaction, is responsible for the observed compatibilization. A small amount (5 wt%) of maleated EP rubber preblended with the maleated ABS led to substantial improvement in the notched Izod impact strength of the final ABS/PA-6 blend. It was postulated that migration of the maleated EPR to the vicinity of SAN/PA-6 boundary and subsequent graft copolymer reaction with nylon led to substantial toughening of the interphase region [Akkapeddi et al., 1993]. At present, only Monsanto is commercially offering ABS/PA blends. These are based on PA-6 (Triax 1120) or PA-66,6 copolymer (Triax 1125).

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M. K. Akkapeddi

Table 15.20. Properties of some commercial grades of polyamide/ABS blends Blend Type

PROPERTY

ABS/PA-6 and ABS/PA-66,6 ASTM TEST METHOD

UNITS

TRIAX 1120

TRIAX 1125

TRIAX 1180

TRIAX 1315 15% GF

1,060 0.9 0.1

1,060 0.9 0.1

1,060 0.9 0.1

1,160 0.3

PHYSICAL Density

D792

kg/m3

Mold Shrinkage Water Absorption, 24 hrs.

D955 D570

% %

MECHANICAL Flexural Modulus Tensile Strength at Break Elongation at Break Rockwell Hardness

D790 D638 D638 D785

MPa (kpsi) MPa (kpsi) %

2070(300) 46(6.6) 270 R95

2140(310) 47(6.8) 270 R99

1930(280) 51(7.4) 330 R101

4275(620) 79(11.5) 2

D256 D256

J/m(ft-lb/in) J/m(ft-lb/in)

NB

NB

NB

520(9.4)

850(16)

850(16) 100(1.9)

1030(19)

115(2.1)

>45(>33) 44(33)

>45(>33) 6(4.4)

>45(>33) >45(>33)

92

94

93

°C m/m/°C

197 1.7E-04

180 1.7E-04

210

°C

60

60

K/V/mm

16.7 3.8 0.014 9.7E+13 108

17.7 4.5 0.05 6.4E+13 105

IMPACT Izod Impact, No Notch Izod Impact, Notched at 23°C at -40°C Drop Weight Impact at 23°C at -40°C THERMAL Heat Deflection Temp at 0.45 MPa at 1.82 MPa Vicat Softening Point Coefficient of Thermal Expansion U.L. Thermal Index ELECTRICAL Dielectric Strength Dielectric Constant at 1 KHZ Dissipation Factor at 1KHZ Volume Resistivity Arc Resistance

D3029

D648

D1525 D676 UL 746B

D149 D150 D150 D257 D495

J (ft-lb)

°C

ohm-m sec

188 104

Commercial Polymer Blends

1073

Figure 15.7. Instrumented Impact behavior of ABS/PA-6 (50/50) blends; top — uncompatibilized blend, bottom — compatibilized blend [Akkapeddi et al., 1993].

Figure 15.8. Morphology of ABS/PA-6 blends (TEM, phosphotungstic acid); top — Uncompatibilized blend (5000X), bottom — Compatibilized blend (10,000X) [Akkapeddi, 1993].

Commercial ABS/PA-6 blend exhibits excellent notched Izod impact (• 850 J/m) at room temperature and good drop weight impact (> 40 J) at -40°C. These impact properties are equivalent to the impact modified polyamide (Table 15.20). However, the DTUL (0.45 MPa) of ABS/PA blends are relatively low compared to those of impact modified nylons. The key difference being that in ABS/polyamide blends, due to the significant level of ABS, a substantial drop in the modulus occurs

at the Tg of ABS(SAN), and hence the blend exhibits lower DTUL at 0.45 MPa than the impact modified polyamide. The latter are polyamide rich (75%) and hence maintain a higher level of heat resistance due to their crystallinity. Commercial ABS/PA blends compete for the same type of applications as the impact modified polyamide, primarily in areas where impact strength and chemical resistance are required. Presumably due to their lower heat resistance and slightly inferior

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M. K. Akkapeddi

low-temperature notched Izod impact strengths, their market growth has been somewhat slower than that of impact modified polyamide. However, ABS/PA blends exhibit somewhat better dimensional stability and lower warpage characteristics than impact modified polyamide, which led to some applications in electronic devices like calculators, key pads, battery packs, etc. 15.8.4

PPE/Polyamide Blends

Poly (2,6-dimethyl-1,4-phenylene ether) PPE is a high-Tg, amorphous polymer [Hay, 1976]. It was originally developed by General Electric Co., which until recently has been the sole producer of this polymer [Hay, 1967]. PPE is generally of homopolymer type, although a copolymer with minor amounts of 2,3,6 trimethyl phenol is also produced. Although PPE exhibits high mechanical strength, DTUL and ductility, it is generally difficult to process as a molding resin due to its high softening temperature (T • 300°C), high melt viscosity and tendency for thermo-oxidative degradation at the high melt processing temperatures (330°C). The discovery that PPE can form highly compatible (miscible) blends with polystyrene led to the development of PPE/HIPS blends [Cizek, 1969] that have been the commercially most successful to date. The properties of PPE/HIPS blends are suitable for many applications. However, they lack adequate chemical resistance. In order to improve chemical resistance, blends of PPE with crystalline polymers such as polyamides have been the subject of much investigation. Simple blends of PPE and polyamides are highly incompatible, generally leading to brittle and readily delaminating products of low value. Hence considerable attention has been paid in recent years to develop technology for effective compatibilization and impact modification of these blends. The exact compositions, the nature of grafting agents or compatibilizers used and the impact modifiers used in commercial PPE/PA blends are still kept proprietary, although many patents have been issued. Research efforts have been focused largely on the methods of improving the compatibility

between PPE and polyamides. Several compatibilizer additives have been claimed in the early patent literature [Ueno and Maruyama, 1982], which included the use of liquid diene rubbers, epoxides and unsaturated compounds containing acid, anhydride, amino, imino or hydroxyl groups. Since polyamides have reactive end groups (amine and carboxyl) it is quite conceivable that the compatibilizing additives may react first with polyamide if all the components are blended together all at once, inhibiting any actual coupling reaction between PPE and polyamide. The latter is more desirable for the effective compatibilization of the blend. Hence subsequent investigators focused first on functionalizing PPE with a grafting agent such that suitable functional groups are attached to the PPE backbone [Jalbert and Grant, 1987; Akkapeddi et al., 1988]. Although functionalization of PPE is possible through solution-phase end-capping reaction with trimellitic anhydride acid chloride [Aycock and Ting, 1986], one would prefer to use melt-phase grafting reactions via reactive extrusion techniques for reasons of economics. Melt phase reaction of PPE with unsaturated functional reagents such as maleic anhydride, fumaric acid, acrylic acid and their derivatives, glycidyl methacrylate and other unsaturated compounds were investigated. This functionalization reaction could be done in extruders and the functionalized PPE could be isolated and characterized. The functionalization step introduced a reactive functional group on the PPE chain, that upon subsequent melt blending with polyamide would react with the amine or carboxyl end groups of nylon forming a graft copolymer of PPE and polyamide at the interface. Since the block and graft copolymers are the best interfacial agents for a blend, the dispersability of PPE in polyamide improves due to decreased interfacial tension, and consequently, the tensile properties (strength, elongation) and toughness of the blend improve considerably. The structure of maleated PPE and the characterization of the graft copolymer was reported [Glans, 1991; Akkapeddi, 1993; Campbell, 1990]. Functionalization of PPE with maleic anhydride or fumaric acid could be done by melt

Commercial Polymer Blends

1075

Figure 15.9. Ultimate elongation and melt flow behavior of reactively compatibilized PPE/PA-6/impact Modifier (50/40/10) blends — Effect of amine terminated PA-6 [Akkapeddi, et al., 1993].

blending (ca. 300°C) in an extruder. This anhydride functionalized PPE could either be isolated and reextruded with PA-6 or melt blended in a single-pass through down stream addition of the polyamide in a twin-screw extruder. In either case, pre-functionalization of PPE is a necessary and an important step in order to optimize the selective grafting reaction between PPE and polyamide, and prevent any premature reaction between polyamide and the maleic anhydride or fumaric acid. The graft-coupling reaction itself is an addition reaction between the amine end group of polyamide and the anhydride group of the functionalized PPE. It was also found that the compatibilization efficiency increased when a PA-6 rich in amine end groups (‘amine terminated polyamide’) was used [Akkapeddi, 1992 and 1993]. The tensile elongation to break increased considerably, as did also the melt viscosity, the latter indicative of increased polymer-polymer grafting reaction (Figure 15.9). Transmission elec-

Figure 15.10. Morphology of PPE/PA-6 (60/40) blends (TEM, phosphotungstic acid stain); top — Uncompatibilized blend (10,000X), bottom — Compatibilized blend (20,000X).

tron microscopy of compatibilized PPE/PA-6 blend (Figure 15.10) indicates finely dispersed PPE in a PA matrix. The heat distortion temperature of the PPE/PA blend at high loads (1.8 MPa) increases with the amount of PPE in the blend (Figure 15.11) [Akkapeddi and VanBuskirk, unpublished results]. However, because of the crystallinity of the polyamide matrix, the heat distortion temperature at low loads (0.4 MPa) is relatively less sensitive to the PPE content and is largely determined by the polyamide. This is one significant difference

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M. K. Akkapeddi

Figure 15.11. Effect of PPE Content on the DTUL of PPE/PA-6 vs. PPE/HIPS blends.

between blends of crystalline/amorphous polymers (PA-6/PPE) vs. blends of amorphous/amorphous (PS/PPE) polymers. In the PPE/polyamide blends, the crystalline PA phase is invariably the continuous phase, because of the large melt viscosity difference between PPE and PA. Although the binary blends of PPE and polyamides exhibit good ductility (tensile elongation and drop weight impact) after the reactive compatibilization, the notched Izod impact strengths are still relatively low. This is to be expected since the individual resin components, viz. PA-6 (or PA-66) and PPE, exhibit low notched Izod impact strengths (< 70 J/m). Hence commercial PPE/PA blends invariably include an impact modifier component, the exact composition and content of which is kept proprietary and varied from grade to grade. Commercial impact modified PPE/PA blends exhibit notched Izod impact strengths ranging from 175 to 500 J/m at room temperature. They also differ in their ductile brittle transition temperature and low temperature impact behavior. The type of nylon used (PA-6 or PA-66 or copolymer type), its end group concentrations and molecular weight, and more importantly, the nature of the rubber modifier used (compatible

or reactive type), its content and manner in which it is added determine the efficiency of notched Izod impact strength improvement and the low temperature ductility. Hydrogenated styrene-butadiene-styrene block copolymer (S-EB-S) rubbers have been used as impact modifiers in PPE/PA-66 blends [Grant et al., 1988] since they are expected to be compatible and readily dispersible in the PPE phase. Preblending a functionalized rubber (such as a maleic anhydride modified EP rubber) into the functionalized PPE, followed by melt blending with the polyamide, generally gave better impact strength improvement in PPE/PA-6 blends [Akkapeddi et al., 1988 and 1992]. An impact modified PPE/PA blend developed by Hüls in Europe, contained polyoctenylene as impact modifier [Droescher, 1988]. Commercial PPE/polyamide blends typically contain 40-60% polyamide and 0-10% of the impact modifier, each grade being formulated for specific types of applications. Typical properties of some commercial unfilled and glass-filled PPE/PA blends are illustrated in Tables 15.21 and 15.22. In general, PPE/polyamide blends offer a unique combination of high heat resistance (DTUL and continuous use temperatures), high impact strength, hygro-thermal dimensional stability and ease of

1077

Commercial Polymer Blends

Table 15.21. Properties of some commercial PPE/polyamide blends Blend Type

PPE/PA-66

PROPERTY

ASTM TEST METHOD

PHYSICAL Density

PPE/PA-6

NORYL

NORYL

NORYL

DIMENSION DIMENSION

UNITS

GTX910 GEC

GTX901 GEC

GTX625 GEC

D9000 AlliedSignal

D9300 AlliedSignal

D792

kg/m3

Mold Shrinkage Water Absorption, 24 hrs.

D955 D570

% %

1,100 1.4 0.5

1,100 1.4 0.3

1,100 1.5 0.3

1,080 1.6 0.5

1,090 1.6 0.5

MECHANICAL Flexural Modulus Flexural Strength Tensile Strength at Yield Elongation at Break

D790 D790 D638 D638

MPa (kpsi) MPa (kpsi) MPa (kpsi) %

2135(310) 76(11) 60(8.6) 60

2000(290) 76(11) 54(7.8) 50

2000(290) 83(12) 56(8) 50

1900(275) 70(10.5) 56(8) 50

2340(339) 96(14) 70(10.2) 75

D256

J/m(ft-lb/in) 220(4) 137(2.5)

175(3.3) 125(2.4)

477(9) 156(3)

550(10) 275(5)

190(3.5) 80(1.5)

100(75)

136(100) 34(25)

185 121 210

180 112 210

IMPACT Izod Impact, Notched at 23°C at -40°C Drop Weight Impact at 23°C at -40°C Instrumented Impact, Energy at 23°C THERMAL Heat Deflection Temp at 0.45 MPa at 1.82 MPa Vicat Softening Point

D3029

D3763

D648

D1525

J (ft-lb)

J (ft-lb) 50(38)

38(28)

45(33)

193 143 232

160 121 215

177 132

°C

°C

processability. In addition, the blends exhibit better chemical resistance compared to other high heat, amorphous engineering resins such as PPE/HIPS blend and polycarbonate (Table 15.23) [Akkapeddi et al., unpublished results].

PPE/PA blends have found many applications in the automotive area, such as in fenders, hatchbacks, wheel covers and mirror housings. Because of their high heat resistance, exterior automotive body parts made from PPE/PA blends can be painted “on-line,” attached to the metal frame in

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M. K. Akkapeddi

Table 15.22. Properties of 30% glass reinforced PPE/PA blends vs. PPE/PS blend Blend Type

PPE/PA-6

PPE/PA-66

PPE/PS

DIMENSION D9130 AlliedSignal

NORYL GTX830 GEC

NORYL GFN3 GEC

ASTM TEST METHOD

UNITS

PHYSICAL Density

D792

kg/m3

Mold Shrinkage Water Absorption, 24 hrs.

D955 D570

% %

1,330 0.5 0.2

1,330 0.5 0.2

1,270 0.2 0.06

MECHANICAL Flexural Modulus Flexural Strength Tensile Strength at Yield Elongation at Break Compressive Strength

D790 D790 D638 D638 D695

MPa (kpsi) MPa (kpsi) MPa (kpsi) % MPa (kpsi)

7720(1120) 235(32.5) 160(23) 3.7 77(11.2)

7655(1100) 224(32) 160(23) 3.7 76(11)

7730(1120) 140(20) 119(17) 5 125(18)

D256

J/m(ft-lb/in) 115(2.1)

105(2)

125(2.3)

215 205 215

252 235 249

158 148 158

PROPERTY

IMPACT Izod Impact, Notched at 23°C THERMAL Heat Deflection Temp at 0.45 MPa at 1.82 MPa Vicat Softening Point

D648

D1525

°C

°C

the existing automotive paint oven temperatures of ca. 180 to 200°C. The parts retain good toughness required in these applications. PPE/PA blends have also been used in many industrial and nonautomotive markets such as pumps, water meter housings, lawn and garden tractor hoods, monofilaments, etc. The annual growth for PPE/PA blends is predicted to be about 10%. The cost of PPE/PA blends can, in principle, be lowered by the addition of some polystyrene or HIPS into the PPE phase, taking advantage of their mutual miscibility. However, this results in some sacrifice in DTUL. Most of the commercial PPE/PA blends are based either on PA-66, that gives slightly higher heat resistance to the blend or PA-6, which

gives somewhat higher impact strength to the blend at any given blend ratio. Although other specialty polyamides such as PA-12 and PA-4,6 have also been investigated in PPE blends, they have not gained commercial significance due to their higher cost. 15.8.5

Polyamide/Polycarbonate Blends

Polycarbonate has been blended with commercial polyamides (PA-66 and PA-6), in order to improve its poor solvent resistance while maintaining a reasonable level of heat resistance and toughness. However, simple blends of polycarbonate and polyamides are highly incompatible and hence

1079

Commercial Polymer Blends

Table 15.23. Comparison of the chemical resistance of PA-6, PPE/PA-6, PE/HIPS and PC(a) (% change in properties(b)) Solvent

Water Antifreeze (50% aqueous glycol) Gasoline Gasohol (15% Methanol) Brake fluid Transmission fluid Power steering fluid Motor oil Trichloroethylene

PA-6

PPE/PA-6

PPE/HIPS

Polycarbonate

Y.S.

Eb

Y.S.

Eb

Y.S.

Eb

Y.S.

Eb

-40 -20

200 200

0 0

0 0

0 0

-30 -70

0 -75 Stress Cracked

0 -50

200 300

0 -50

100 0

Stress Cracked Dissolved

-80 -100 Stress Cracked

0 0 0 10 0

21 -25 -50 -50 200

10 0 0 0 -70

-25 0 0 0 100

-80 -100 Stress Cracked Stress Cracked Stress Cracked Dissolved

Stress Cracked 0 0 Stress Cracked Stress Cracked Stress Cracked

(a) Tensile bars, 3 mm thickness, tested at 6% strain after one week immersion in a solvent. (b) % change in properties, increase (+) or decrease (-) relative to the dry as molded properties in air; Y.S. = Yield Stress; Eb = elongation at break.

not useful. Several different additives such as phenoxy resins, polyester amide elastomers in combination with maleated polyolefins, polyetheramide block copolymers, and polyamidepolyacrylate block copolymers have been used as compatibilizers and impact modifiers. However, polyamide/polycarbonate blends have not been commercialized to any significant extent yet. The commercial grades of polyamide/polycarbonate blend (Dexcarb®) exhibit high notched Izod impact strength comparable to that of impact modified polyamides and polycarbonate (Table 15.24). It is believed that an elastomeric impact modifier was included in these compositions. Patent claims the use of a polyetheramide and a maleated polypropylene or EPR as the compatibilizing/ impact modifying additives [Perron, 1988]. The commercial grades contain varying amounts of PC, PA-66 (or 6) and the impact modifier. Polyamide/polycarbonate blends have been evaluated for exterior automotive applications such as bumper beams. Their dimensional stabil-

ity, resistance to paint solvents and automotive fluids, and a relatively low mold shrinkage may have been some of the reasons for their consideration. However, polyamide/polycarbonate blends have not yet gained any significant commercial volumes. These blends must face significant competition from PPE/polyamide blends, since the latter offer higher heat resistance viz., higher DTULs at both low and high loads, while also offering an equivalent or better toughness and solvent resistance. Polyamide/polycarbonate blends must also face stiff competition from impact modified PA’s that offer higher impact strengths at low temperatures and higher DTULs at low loads. 15.8.6

Polyamide/Silicone Blends

A commercial blend consisting of a thermoplastic polyamide (PA-6 or PA-66) and 5-25 wt% of a crosslinkable silicone, which forms a semiinterpenetrating network (semi-IPN) upon curing, has been offered under the trade name of Rim-

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M. K. Akkapeddi

Table 15.24. Comparison of different types of commercial polycarbonate blends Blend Type

PC/ABS

PC/S-MMA

PC/PBT

PC/N6

PULSE 710 Dow

SD9101

XENOY 5220 GEC

DEXCARB 507 D&S Int.

ASTM TEST METHOD

UNITS

PHYSICAL Density

D792

kg/m3 % %

1,250 1.6 0.14

1,100

D955 D570

1,120 0.6 0.15

1,150

Mold Shrinkage Water Absorption, 24 hrs. MECHANICAL Flexural Modulus Flexural Strength Tensile Strength at Yield Elongation at Break

D790 D790 D638 D638

MPa (kpsi) MPa (kpsi) MPa (kpsi) %

2275(330) 83(12) 45(6.5) 80

2490(361)

1900(275)

1725(250)

56(8.2) 125

45(6.5) 175

41(6) 100

D256

J/m(ft-lb/in) 534(10) 481(9)

1015(19)

850(16)

908(17)

102

92

106 99

127 102

CHEMICAL RESISTANCE

Fair

Fair

Good

Excellent

plast® (Petrarch, div. of Hüls) [Arkles, 1985]. The blends are produced by extruding the polyamide with a vinyl terminated, polydimethylsiloxane and silicone hydride terminated dimethyl siloxane in the presence of a platinum catalyst. The siloxanes react with one another forming cured silicone thermoset in the thermoplastic matrix phase of the polyamide. At the low concentrations 5-10 % silicone, the blend is reported to retain much of its thermoplasticity and processability. This silicone semi-IPN reportedly improves the lubricity, wear and biocompatibility of polyamide as well as reducing the shrinkage and warpage although

with some sacrifice in strength and elongation. Some applications in medical tubing and industrial uses have been projected. No significant market for this technology has yet been developed, primarily because of the high cost of this blend system.

PROPERTY

IMPACT Izod Impact, Notched at 23°C at -40°C THERMAL Heat Deflection Temp at 0.45 MPa at 1.82 MPa

D648

Novacor

°C

15.8.7

Polyamide/Polyamide Blends

Blends of PA-66 with some aliphatic-aromatic polyamides of higher Tg such as poly (hexamethylene isophthalamide) (PA-6I) have been evaluated as fibers, in order to achieve improved modulus and heat resistance in PA-66 tire yarn and con-

Commercial Polymer Blends

sequently reduce the tire flat-spotting tendency [Zimmerman et al., 1973]. Although the rate of amide interchange in these polyamide/polyamide blends was found to be quite slow (< 3% interchange based on the aromatic polyamide), the formation of some block copolymer was not ruled out. Nevertheless, the observed improvements in the modulus and Tg were primarily attributable to some degree of compatibility or miscibility between the components of the blend. There have been several other research investigations on the compatibility and phase behavior of polyamide/ polyamide blends [e.g., Takanayagi et al.,1980; Ellis, 1990; Hirakawa et al., 1985]. However, no commercially significant polyamide/polyamide blends have been developed until recently. There has been some commercial interest particularly in the blends of PA-6 with aliphatic-aromatic polyamides exhibiting good oxygen barrier properties suitable for food packaging applications. Certain types of aliphatic-aromatic random copolyamides such as those derived from the polymerization of (a) hexamethylene diamine and isophthalic/terephthalic acids (6I/6T) or (b) m-xylylenediamine, isophthalic/terephthalic acids and caprolactam (MXDI/T,6), exhibit high glass transition temperature, amorphous character and a high barrier to oxygen permeation [Akkapeddi and Gervasi, 1989]. The oxygen barrier property of these amorphous polyamides is retained even under high humidities because of their high Tg’s. In contrast, semicrystalline polyamides such as PA-6 and PA-66 characteristically exhibit low Tg (ca. 50°C) that becomes even lower after moisture absorption. As a consequence, the permeability of oxygen in PA-6 and PA-66 increases with humidity. There are several amorphous polyamides commercially available that exhibit good oxygen barrier properties, suitable for packaging applications e.g., Selar(R) PA (duPont), Novamid(R) (Mitsubishi Chemical). To improve the oxygen barrier properties of PA-6 (and PA-66), high Tg, barrier type amorphous polyamides such as PA-6I/6T (Selar® PA) have been blended [Krizan et al., 1989; Blatz, 1989]. These blends exhibit improved barrier to oxygen permeation even at high relative humidity

1081

(Figure 15.12), presumably due to a good degree of miscibility. These commercial PA-6/amorphous polyamide blends have been developed particularly for applications in flexible food packaging markets and hence captively produced as films.

Figure 15.12. Oxygen permeability and water vapor transmission rate (WVT) behavior of PA-6/Amorphous PA-6I/6T blends at 95% RH.

Similarly there has been some commercial activity in the blends of poly(m-xylylene adipamide) (MXD6) with PA-6 and PET. MXD6, produced by Mitsuibishi Gas Chemical Co., in Japan, has been used both as a molding resin (Reny®) as well as an oxygen barrier film resin particularly in multi-layer coextrusions [Harada, 1988]. In blends of MXD6 with aliphatic polyamides, interchange reactions were found to play a major role in the observed phase homogenization during melt-mixing, particularly at long residence time [Takada and Paul, 1992].

1082

15.9

M. K. Akkapeddi

Polycarbonate Blends

Commercial polycarbonate is an amorphous engineering thermoplastic characterized by a high glass transition temperature (ca. 150°C) and an excellent balance of properties such as high toughness, clarity, heat resistance, dimensional stability, good electrical and ignition-resistance characteristics. Because of this outstanding combination of properties, polycarbonate has become one of the most successful of engineering thermoplastics with an estimated global consumption of over 595 kton/y [Kircher, 1990]. Standard polycarbonate, (PC), is made from bisphenol A and phosgene via an interfacial polymerization process. The polymer backbone has an aromatic polycarbonate structure with a recurring carbonate, moiety which, uniquely accounts for the outstanding toughness of the polycarbonate and the rigid aromatic unit contributes to its high glass transition temperature. Although a large number of applications of polycarbonate have been based on its unique combination of high impact strength, heat resistance and clarity there are still a few property deficiencies

in the neat polymer which can be overcome by blending with other polymers or additives. Some deficiencies of polycarbonate are: 1. high notch sensitivity and part thickness sensitivity in impact strength (Figure 15.12); 2. lack of an adequate low temperature notched Izod impact strength; 3. lack of an adequate solvent resistance and stress crack resistance (Table 15.23); 4. limited long-term hydrolytic stability at elevated temperatures; 5. relatively high melt viscosity compared to each crystalline polymers as PA-6, PBT, etc. Hence the development of polycarbonate blends was primarily market driven, with a motivation to extend the applications of polycarbonate into areas where improved chemical resistance and processability are required while still retaining high impact strength. In the development of all the polycarbonate blends the common goal was to maintain a very high level of impact strength while improving the properties and cost balance (Table 15.24). Commercial blends of polycarbonate and the major reasons for their development are listed as follows:

Blend

Reasons for blending

1. Impact modified polycarbonates (PC/LDPE, PC/elastomer, etc.)

* Improve notch sensitivity * Improve low temperature toughness * Improve thermal-aging resistance * Improve low temperature toughness * Improve processability * Lower cost * Improve processability * Weatherability * Improve solvent resistance * Improve processability * Improve solvent resistance * Lower cost, improved processability

2. ABS/Polycarbonate

3. Styrenic/PC blends 4. Thermoplastic polyester/polycarbonate (PBT/PC, PET/PC) 5. Polyamide/PC blend 6. Polyetherimide/PC blend (PEI/PC)

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Commercial Polymer Blends

15.9.1

Impact Modified Polycarbonates

Although polycarbonate is exceptional among engineering resins in exhibiting an outstanding level of toughness, its ductile-brittle transition depends on the temperature, notch sharpness, sample thickness and thermal aging effects. A sharp ductile-brittle transition [Carhart, 1985] for polycarbonate occurs at 0-10°C, hence its

notched Izod impact strength at low temperatures is low, e.g., ” 100 J/m at -30°C. Another important deficiency of polycarbonate is the sensitivity of its notched impact strength to part thickness and notch radius. The notched Izod impact strength of polycarbonate is reduced from 900 J/m to about 100 J/m when the thickness is increased from 3.2 to 6.4 mm [Jones, 1985] These effects are due to the changes in the deformation behavior at

Table 15.25. Properties of commercial impact modified polycarbonate vs. polycarbonate Blend Type

PC/LDPE

PC/ELASTOMER

PC

MAKROLON T7700 Miles

MAKROLON T7855 Miles

MAKROLON 2800 Miles

ASTM TEST METHOD

UNITS

PHYSICAL Density Mold Shrinkage Water Absorption, 24 hrs.

D792 D955 D570

kg/m3 % %

1,190

1,190 0.6 0.15

MECHANICAL Flexural Modulus Flexural Strength Tensile Strength at Yield

D790 D790 D638

MPa (kpsi) MPa (kpsi) MPa (kpsi)

2200(320) 86(12.5) 59(8.5)

2200(320) 86(12.5) 62(9)

2275(330) 86(12.5) 62(9)

D256

J/m(ft-lb/in) 740(14) 510(9.5)

800(15) 640(12)

908(17) 117(2.4)

132 150

134 127 143

138 131 157

16 3 0.01 1.0E+15 100

16 3 0.01 1.0E+15 100

16 3 0.01 1.0E+15 120

PROPERTY

IMPACT Izod Impact, Notched at 23°C at -40°C THERMAL Heat Deflection Temp at 0.45 MPa at 1.82 MPa Vicat Softening Point

D1525

°C

ELECTRICAL Dielectric Strength Dielectric Constant at 1 MHZ Dissipation Factor at 1MHZ Volume Resistivity Arc Resistance

D149 D150 D150 D257 D495

KV/mm

D648

°C

ohm-m sec

1084

M. K. Akkapeddi

the crack tip from shear yielding (plane stress) to crazing and unstable crack propagation. By blending small amounts of elastomeric or low modulus polymers, the impact strength of polycarbonate could be readily improved, largely by modifying the crack tip plastic deformation. The rubber particles facilitate the localized matrix yielding through internal cavitation and debonding mechanisms. Polycarbonate modified with small amount of low density polyethylene (ca. 5%) exhibits substantially improved notched Izod (> 500 J/m) in thick sections (6.4 mm) [Dobkowski, 1988; Dobrescu and Cobzaru, 1978; Gardlund, 1984]. Some commercial impact modified grades of PC (Table 15.25) are based on this technology [Freitag et al., 1985]. GE’s Lexan EM (energy management) grades are believed to contain 5-8% of a polyolefin [Kosoff, 1985]. Addition of polyethylene is also reported to improve the stress crack resistance [Freitag et al., 1985]. In order to improve both the thickness sensitivity and low temperature sensitivity of the impact strength, polycarbonate has been blended with a variety of low Tg, elastomeric impact modifiers. More important among these are the core-shell rubbers like PMMA-g-polybutadiene, PMMA-g-SBR (MBS), PMMA-g-n-butylacrylate (acrylic core-shell), all normally composed of ” 0.1µm crosslinked rubbery core particles. These modifiers improve both the thick-section (6.4 mm) and low temperature notched Izod impact properties of polycarbonate [Witman, 1981; Neuray and Ott, 1981; Bussink et al., 1977] (Figures 15.13 and 15.14). The blends are of course opaque. Impact modified polycarbonate also shows better retention of impact strength with heat-aging. Commercial impact modified (elastomer blended) polycarbonates are used for the production of protective headgear, sporting goods, bobbins for textile industry and automotive components requiring high toughness. Impact modified PC has to compete with the more solvent resistant impact modified PBT/PC, PET/PC blends and polyamides in some of the applications.

Figure 15.13. Effect of sample thickness on the notched impact strength of unmodified (I) vs. impact modified (II) polycarbonates [Freitag et al., 1985].

Figure 15.14. Low temperature impact strength of unmodified (I) vs. impact modified PC (II) [Freitag et al., 1985].

14.9.2

ABS/Polycarbonate Blends

Since their introduction in 1967, ABS/polycarbonate blends [Grabowski, 1964] have enjoyed a dramatic growth in market volume. They are by far the largest volume polycarbonate blends

Commercial Polymer Blends

1085

Figure 15.15. Ductile brittle transition temperature of ABS/PC Blends vs. ABS and PC [Freitag et al., 1985].

used with the global consumption approaching 100 kton/y by 1996. The growth rate of ABS/PC blends is estimated to be > 12%, which is faster than that of polycarbonate or ABS. The major reason for the success of ABS/PC blends is their overall better cost/performance balance relative to PC and impact modified PC. Particularly noteworthy is their unique, synergistic improvement in the low-temperature notched impact strength [Morbitzer et al., 1983], which is better than the individual components (Figure 15.15). The partial miscibility between the SAN and the polycar-bonate phases, the intrinsic ductility of polycarbonate matrix, and the presence of small particle size polybutadiene dispersions are the key factors contributing to the low temperature toughness of the blend. A discussion of the ABS/PC blends comparing with other ABS blend, may be found under the ABS blends section. The properties of the ABS/PC blends, primarily the DTUL and impact strength, are determined by the ratio of ABS to polycarbonate. The morphology is also dictated by the blend ratio. In blends containing • 50% polycarbonate, the continuous phase is formed by the polycarbonate with ABS as the dispersed

phase. The rubber particles are primarily located in the SAN phase. Typical properties of some of the commercial ABS/PC blends are illustrated in Table 15.5. The blends containing higher levels of polycarbonate exhibit better low temperature impact strengths. ABS/PC blends are used in a variety of automotive components such as instrument panels, flaps for glove compartments, dashboards, cladding for steering wheel columns, ventilation ports, spoilers, wheel covers and protective side trims. Flame retardant ABS/PC blends are used in office equipment and business machine housings. Chlorine or bromine-free flame retardant ABS/ PC compositions contain organophosphate/PTFE powder dispersions [Freitag et al., 1991]. ABS/PC blends have been used in automotive exterior body panels for GM’s Saturn models. The assembly line for this car was equipped to handle the new water-borne paint systems, which cure at lower temperatures than the normal ‘E-coat’ paint ovens. Due to their excellent low temperature toughness, dimensional stability and cost advantages, ABS/PC blends are replacing some of the PPE/HIPS applications and are also competing against some polyamide applications. However,

1086

M. K. Akkapeddi

they cannot compete with the impact modified polyamides and the PPE/PA blends in all those applications which require higher heat and chemical resistance such as automotive fuel emission canisters, fasteners, connectors and exterior parts like fenders, etc. 15.9.3

and PC are miscible, exhibiting a lower critical solution temperature behavior (LCST ≅ 180°C). S-MMA/PC blends are relatively new and their advantages, if any, over ABS/PC blends have not been clearly identified. They will, no doubt, be more expensive than the ABS/PC blends but will have better UV resistance.

Styrenic/Polycarbonate Blends

Blends of polycarbonate with other styrenic resins are relatively new and therefore their current market volume is low. They have been developed primarily to upgrade the performance of such styrenic resins as styrenic-maleic anhydride (SMA), styrene-methyl methacrylate (S-MMA), acrylic-styrene-acylonitrile (ASA) resins primarily for impact strength and to some extent for DTUL improvement (Tables 15.7, 15.8 and 15.23). These blends uniquely combine high notched Izod impact strengths (ranging from 500 J/m to > 1000 J/m) with better UV resistance or weatherability than ABS. SMA/PC blend. This blend (Arloy®, Arco) contained SMA grafted with polybutadiene as the impact modifier. The properties of SMA/PC blend were similar to ABS/PC blend with slightly higher heat distortion temperatures (107 to 117°C) but comparable impact strength (> 500 J/m). However, it was discontinued from the market due to unfavorable economics relative to ABS/PC blend. The partial miscibility between the styrene-maleic anhydride and polycarbonate accounts for the adequate compatibility of this blend as evidenced by the high level of tensile and impact strengths. S-MMA/PC blend. This blend (Novacor® SD-9101) was reported to have better flow, surface finish and scratch resistance than PC/polyester blends and an equivalent level of impact toughness (Table 15.24). It is believed that these formulations also include some acrylic rubber (core-shell type) for impact modification. One would expect a sufficient level of partial miscibility for self-compatibilization between the styrenemethyl methacrylate copolymer (S-MMA) and the polycarbonate especially at high MMA content of the copolymer, since the binary blends of PMMA

ASA/PC blends. ASA polymers are similar to ABS except that the polybutadiene rubber phase in the SAN copolymer matrix is substituted by the more weatherable acrylic-rubber, viz., the crosslinked n-butyl acrylate rubber particles grafted with SAN copolymer. The weatherability advantages of ASA polymers can be extended also into the polycarbonate blends, [Sakano, 1980] as long as sufficient ASA is used in the blend (• 50%). ASA/PC blends find niche applications in exterior automotive parts such as cowl vents, grills, mirror housings, trim, etc., where the parts can be pigmented (black or color) instead of being painted, unlike the case of ABS/PC. Since the withdrawal of EPDM-g-SAN copolymer (AES, Rovel®, Dow) from the market, ASA and ASA blends have assumed increased commercial significance in outdoor, weatherable applications. AES/PC blends. AES polymers are similar to ABS except that the polybutadiene rubber is replaced with EPDM rubber, which is grafted to the SAN copolymer. The saturated backbone of the rubber makes the AES or EPDM-g-SAN (Rovel®, Dow) more weatherable. Blends of AES/PC are similar to ABS/PC with the same level of high impact strength and good DTUL but with the added advantage of weatherability. The compatibility between the SAN phase of the AES and the polycarbonate is governed by the same principle of partial miscibility. Since the withdrawal of AES from the market, other suppliers with ASA based blends are actively following this market. 15.9.4

Thermoplastic polyester/PC Blends (PBT/PC, PET/PC, PCTG/PC)

The second most important class of commercial polycarbonate blends is derived by blending with

1087

Commercial Polymer Blends

commercial thermoplastic polyesters such as polybutylene terephthalate (PBT) and polyethylene terephthalate (PET). Both PBT and PET are crystallizable polymers and hence offer the expected chemical resistance advantages of the crystalline polymers in blends with polycarbonate. Among the thermoplastic polyester/polycarbonate blends, the PBT/PC blend has the major commercial volume, followed by the PET/PC blend. A copolymer of 1,4-cyclohexanedimethanol, ethylene glycol and terephthalic acid (PCTG) forms a miscible blend with polycarbonate. PCTG/PC blend is commercially offered by Eastman Kodak (Ektar®) for specialty applications.

PBT/polycarbonate blend, first introduced in 1980 by General Electric (Xenoy®) has enjoyed a fast growth in automotive applications, particularly for bumpers. The blend was developed to meet the low temperature impact strength, dimensional stability and paintability requirements of rigid bumper fascias [Bertolucci and Delany, 1983], which accounted for the bulk of the current market volume for the PBT/PC blend (estimated 25 kton/y in the USA). The development of commercial PET/PC blends followed shortly after the initial success of PBT/PC blends. Currently there are several commercial blends of both PBT/PC and PET/PC blends available,

Table 15.26. Properties of some commercial grades of PBT/polycarbonate blends ASTM TEST METHOD

UNITS

XENOY 1102 GEC

XENOY 5220 GEC

XENOY 6120 GEC

SABRE 1628 Dow

PHYSICAL Density Mold Shrinkage Water Absorption, 24 hrs.

D792 D955 D570

kg/m3 % %

1,200 0.9 0.08

1,210 0.9 0.12

1,250 1.6 0.14

1,200 0.7 0.1

MECHANICAL Flexural Modulus Flexural Strength Tensile Strength at Break Elongation at Break

D790 D790 D638 D638

MPa (kpsi) MPa (kpsi) MPa (kpsi) %

1964(285) 82(11.9) 54(7.9) 150

2040(296) 85(12.3) 53(7.7) 120

1900(275) 72(10.5) 45(6.5) 175

2136(310) 81(11.7) 56(8.1) 160

D256

J/m(ft-lb/in) 800(15) 640(13)

710(13) 300(5.6)

850(16)

800(15) 534(10)

54(40)

54(40)

54(40)

110 91

106 99

121 57

PROPERTY

IMPACT Izod Impact, Notched at 23°C at -40°C Drop Weight Impact at 23°C THERMAL Heat Deflection Temp at 0.45 MPa at 1.82 MPa

D3029

D648

J (ft-lb)

°C 119 83

1088

M. K. Akkapeddi

Table 15.27. Properties of some commercial PET/polycarbonate blends ASTM TEST METHOD

UNITS

PHYSICAL Density Mold Shrinkage Water Absorption, 24 hrs.

D792 D955 D570

MECHANICAL Flexural Modulus Flexural Strength Tensile Strength at Yield Elongation at Break

PROPERTY

IMPACT Izod Impact, Notched at 23°C at -40°C Drop Weight Impact at 23°C Instrumented Impact, Energy at 23°C THERMAL Heat Deflection Temp at 0.45 MPa at 1.82 MPa

MAKROBLEND UT1018 Miles

SABRE 1647 Dow

AlliedSignal

kg/m3 % %

1,220 0.8 0.16

1,220 0.1 0.16

1,180 1.1 0.15

D790 D790 D638 D638

MPa (kpsi) MPa (kpsi) MPa (kpsi) %

2070(300) 75(10.9) 48(7) 160

1930(280) 77(11.1) 51(7.4) 130

1590(230) 60(8.7) 44(6.4) 185

D256

J/m(ft-lb/in) 970(18) 700(13)

750(14) 214(4)

800(15) 140(2.6)

D3029

J (ft-lb)

D3763

J (ft-lb)

135(175)

D648

IMPACT

100(135)

20(27)

26(35)

22(30)

115 88

119 81

110 80

°C

and their properties are compared in Tables 15.26 and 15.27. All the commercial PBT/PC and PET/PC blends also contain typically 10-20 wt% of an additional elastomeric impact modifier. The exact nature and the content of the impact modifier is kept proprietary and often forms the basis for a particular blend patent. Typically core-shell rubbers such as polymethylmethacrylate grafted butadiene-styrene rubber (MBS) or an all acrylic core-shell rubber such as poly (MMA-g-n-BuA) are used [Nakamura, 1975; Chung, 1985]. ABS (with high polybutadiene content • 50%) or ASA

rubber (• 50% acrylate rubber) have also been used. The presence of such a rubber component is definitely needed to obtain high notched Izod impact strengths (• 500 J/m) in these blends. The binary blends of polycarbonate with polybutylene terephthalate (PBT/PC) or polyethylene terephthalate (PET/PC) are now known to be essentially phase separated blend systems exhibiting two glass transition temperatures in each case, one for the polycarbonate-rich phase, and another for the polyester-rich phase. [Murff et al., 1984; Huang and Wang, 1986; Wahrmund et al., 1978]. The evaluation of the amorphous phase miscibility

1089

Commercial Polymer Blends

in these blends was often complicated by the potential of a transesterification reaction between the two polymers during the melt blending, which may in principle lead to a block copolymer and eventually to a random copolymer with a single phase, single glass transition temperature behavior [Deveux et al., 1982; Hobbs et al., 1987; Porter and Want, 1992]. Sometimes different conclusions have appeared in the literature regarding the phase behavior of PET/PC and PBT/PC blends, because different methods were used to prepare the blends. 15.9.4.1

Polybutylene Terephthalate/ Polycarbonate Blends (PBT/PC)

Early investigations of the melt blends of PBT and PC showed two glass transition temperatures indicative of two amorphous phases [Murff et al., 1984]. However the Tg’s did not correspond to those of pure components and in addition, there was a slight depression in the melting point of PBT. A melt-phase reaction was hypothesized to take place. Subsequent studies showed that the melt phase transesterification reaction between PBT and PC can indeed take place, at these high temperatures (T • 260°C), as followed by NMR spectral changes as a function of melt residence time [Deveux et al., 1982]. The presence of titanium catalyst normally present in PBT also catalyzed this reaction. Although NMR and IR techniques were used in characterizing these interchange reactions, these techniques are often insufficient to detect small changes in the structure that occurs when only one or few interchanges per chain take place. For example, NMR can detect only gross structural changes that take place after long reaction times (• 30 min at 250°C). Blends made by solution technique or by normal extruder melt blending process under short residence time exhibited a two phase structure [Hanrahan et al., 1985; Dekkers et al., 1990]. However, the Tg of the PC phase in the melt blended product was usually lower (by 20°C) than the PC phase of solution blended product or neat PC. This was attributed to the effect of a small, although not readily detectable level of interchange reaction in the melt blended product.

The transesterification reactions in PBT/PC melt blends could be suppressed by using organophosphites and phosphonates which probably function by deactivating the titanium or antimony type polymerization catalyst residues present in PBT [Golovoy et al., 1989]. Even in the presence of phosphite stabilizers, PBT/PC blends showed dual phase behavior. However, a partial miscibility was evident since the Tg of PC phase was still reduced from the normal 150°C to about 140°C. This partial miscibility between PBT and PC which occurs even in the absence of an exchange reaction is responsible for the good compatibility and interfacial strength of the blend. 15.9.4.2

Polyethylene Terephthalate/ Polycarbonate Blend (PET/PC)

Early work on the PET/PC blends indicated that in blends containing > 70% PET, PC was miscible and in blends with < 70% PET, the components separated into two phases. Subsequent investigations concluded that the blend was essentially immiscible over the entire composition range. The transesterification reaction between PET and PC in the melt phase was found to be slower than in the case of PBT at 270°C. However, when the reaction does occur, the newly formed ethylene carbonate linkages, -CH2-CH2O-CO-O- in the polymer chain appear to degrade more rapidly than the butylene carbonate units generated with PBT. The divergent conclusions in the literature on the phase behavior of PET/PC blend are undoubtedly caused by the differing degrees of transesterification reaction occurring under the conditions of melt mixing (temperature, time) and the amount of catalyst residues [Goddard et al., 1986]. Under the normal extrusion blending conditions PET/PC blend forms a 2-phase blend morphology. The melt residence times are short enough such that the interchange reaction does not occur, and the PC can be quantitatively extracted out with methylene chloride. From the observed glass transitions of the PET-rich phase and PCrich phase, measured by DSC or DMA, the blend can be classified as partially miscible with an estimated interaction parameter of slightly positive

1090

M. K. Akkapeddi

(χ12 = 0.04). From the measured Tg’s, it appears that more PET dissolves in PC phase than PC does in the PET-rich phase [Kim and Burns, 1990]. Nevertheless, the existence of a partial miscibility, even in the absence of transesterification accounts for the self-compatibilization effect in the blend. Mutual interpenetration of the components at the phase boundary accounts for the high interfacial strength. The binary blends of PET/PC and PBT/PC exhibit good ductility and tensile strengths but the notched Izod impact strengths are still low at all the blend compositions containing ” 80% PC. This is a significant difference from ABS/PC blends, in which the grafted polybutadiene rubber particles of ABS phase contribute to the toughness at all the ABS/PC blend ratios. Hence commercial PBT/PC and PET/PC blends, by necessity, include a proper level of an effective impact modifier, usually a core-shell rubber of small particle size, such as MBS, ABS, etc. It has been shown that during the melt blending, these rubber particles preferentially migrate to the PC phase due to the known compatibility of the PMMA or SAN component of these rubbers (shell structure) with the polycarbonate [Dekkers et al., 1990]. Hence the blend morphology indicates that the rubber particles are predominantly located in the PC phase (Figure 15.16). Commercial PBT/PC and PET/PC blends contain about 15-20 % of such core-shell rubber impact modifiers for maximum toughness, i.e., notched Izod impact strengths of typically • 700 J/m which is maintained even at low temperature (Tables 15.26 and 15.27). The ratio of PBT/PC or PET/PC is usually kept between 50/50 to 40/60 to optimize the ductility in the blend, while still maintaining a continuous or co-continuous phase of the polyester. A continuous phase of PBT or PET with PC as dispersed phase, would be preferred for solvent resistance. The crystallization of PBT and PET in these blends is somewhat suppressed by the partial miscibility of the PC. However, since PBT crystallizes intrinsically faster than PET, blends of PBT and polycarbonate after injection molding show a crystalline PBT phase in their morphology.

Figure 15.16 Morphology of PET/PC/MBS core-shell elastomer (40/40/20) blends - TEM; RuO4 stain; 15000X [Akkapeddi et al., 1993].

Injection molded PET/PC blends, on the other hand, generally lead to an essentially amorphous PET phase. Short term annealing at elevated temperatures (120-150°C) causes this PET phase to crystallize and leads to some loss in ductility especially when the polycarbonate content in the blend is ” 30% [Akkapeddi et al., unpublished results]. Long term thermal aging of PBT/PC blends is known to lead to severe embrittlement [Bertillson et al., 1988]. Phase segregation, secondary crystallization and changes in the amorphous phase free volume with aging are some of the key factors attributable to the embrittlement phenomenon. Embrittlement of PET/PC blends upon heat-aging is even more likely to occur due to the more brittle nature of crystalline PET phase. Recent work has focused on the use of reactive rubber tougheners to improve the embrittlement resistance of PET/PC blends upon heat-aging [Akkapeddi and Mason, 1991; Akkapeddi et al., 1993]. The excellent low temperature toughness and solvent resistance of PBT/PC and PET/PC blends have found good application in automotive exte-

1091

Commercial Polymer Blends

rior parts such as bumper beams and fascias (mainly PBT/PC), air dams, rocker panels, wheel covers, and mirror housings. Non-automotive applications included instrument housings, lawnmower chutes, snow blower components, etc. With the increasing competition from low cost paintable thermoplastic polyolefins (TPO’s) in the automotive bumper and fascia markets, further growth for PBT/PC and PET/PC blends may be limited in this area. Hence these blends must find other applications where their dimensional stability, chemical resistance, high toughness and moderate DTUL (T ” 100°C) are considered suitable. 15.9.4.3

Poly(1,4-cyclohexanedimethyleneterephthalate) Copolymer/ Polycarbonate Blend (PCTG/PC)

The commercial copolyester derived from 1,4-cyclohexane dimethanol, ethylene glycol and terephthalic acid (PCTG, Ektar® DN003, Eastman Kodak) is an amorphous polymer exhibiting excellent toughness (notched Izod impact strength of • 800 J/m), high clarity and good chemical resistance characteristics. However, it has a relatively low glass transition temperature (ca. 85°C) and a low heat distortion temperature (ca. 66°C at 1.8 MPa). Hence a commercial blend of PCTG and polycarbonate was developed, which maintained the high toughness of both components (notched Izod import strength • 800 J/m), yet had a useful combination of DTUL (ca. 95°C) and chemical resistance. More importantly, the blend maintained a high degree of clarity because of the thermodynamic miscibility between the PCTG and polycarbonate, although some transesterification cannot be ruled out. The miscibility was confirmed by a single Tg, single phase behavior of the blend [Mohn et al., 1979; Smith et al., 1981]. Commercial PCTG/PC blends (Ektar® DA series, Eastman Kodak) have been used in lawn and garden equipment, floor care appliance parts, sterilizable medical equipment, etc., where their combination of clarity, toughness, chemical resistance, heat resistance, UV and gamma radiation resistance have been well utilized. Molded parts

generally showed excellent surface finish and hence molded-in-color could be used. 15.9.5

Polyamide/Polycarbonate Blends

Unlike the thermoplastic polyesters (PBT and PET), the commercial polyamides such as PA-6 or PA-66 are more polar and hence highly immiscible with polycarbonate. Hence simple blends of polyamide and polycarbonate are expected to delaminate readily, unless a suitable compatibilizer is used. Because of the lack of an efficient compatibilization technique, blends of PA-6 (or PA-66) with polycarbonate have not yet reached a commercial significance, although there are several patents claiming improvement in properties. At present, there is only one commercial source of the polyamide/polycarbonate blends (Dexcarb®, Dexter Corp.). According to their patent, the blend was compatibilized by using a combination of a polyesteramide elastomer and a maleated olefinic polymer, such as maleated polypropylene or EP rubber [Perron, 1984; 1988]. However the degree or the efficiency of compatibilization achieved is unknown, since the added components are not known to be miscible or compatible with the polycarbonate. Nevertheless, the data sheet indicated good properties including a high notched Izod impact strength of > 700 J/m (Table 15.24). At the present time, the commercial volume of polyamide/PC blend is small due perhaps to its unfavorable cost/performance balance. The product has to compete with the more established impact modified polyamide and PPE/PA blends, which offer higher DTUL/impact strength balance. 15.9.6

Polyetherimide/PC Blends (PEI/PC)

Polyetherimide (Ultem® 1000, GEC) is a high performance engineering thermoplastic with high heat distortion temperature (> 200°C), high mechanical strength and inherent flame-retardancy characteristics. Recently blends of polyetherimide with polycarbonate have been commercially offered as thermoformable sheets and as molding compounds (Table 15.28). The primary reason for

1092

M. K. Akkapeddi

Table 15.28. Properties of commercial polytherimide/polycarbonate vs. polytherimide and polycarbonate Blend Type

PEI/PC ASTM TEST METHOD

UNITS

PHYSICAL Density

D792

kg/m3

Mold Shrinkage

D955

MECHANICAL Flexural Modulus Tensile Strength at Yield Elongation at Break

PROPERTY

IMPACT Izod Impact, Notched at 23°C Drop Weight Impact at 23°C THERMAL Heat Deflection Temp at 1.82 MPa

PEI

PC

ULTEM LTX100A GEC

ULTEM LTX100B GEC

ULTEM 1000 GEC

LEXAN 141 GEC

%

1,310 0.07

1,310 0.07

1,270 0.07

1,200 0.06

D790 D638 D638

MPa (kpsi) MPa (kpsi) %

2900(420) 93(13.5 90

3380(490) 99(14.3) 80

3370(480) 106(15) 60

2390(340) 63(9) 110

D256

J/m(ft-lb/in) 53(1.0)

48(0.9)

54(1.0)

820(15)

D3029

J (ft-lb) 36(27)

36(27)

185

193

210

132

D648

°C

blending polycarbonate appears to be to lower the cost of the PEI while retaining a high level of heat resistance. The blend has about the same impact strength as the polyetherimide, although not as good as polycarbonate [Mellinger, 1985]. However, the DTUL and chemical resistance of the blend is better than that of polycarbonate. Extrusion grade, PEI/PC blend was developed for aircraft applications to meet the federal aviation standards for low flammability, smoke and toxic gas generation [Skeist, 1991]. PEI/PC thermoformable sheet is used for the fabrication of transport aircraft window housings, air ducts, consoles and other components. PEI/PC sheet is also used as high temperature paint mask in the automotive industry. PEI/PC molding compounds have also been evaluated for microwaveable cookware. Due to the high cost of PEI,

blends of PEI and PC will continue to be used only for specialty niche applications in low to moderate volumes.

15.10

Polyoxymethylene Blends

Polyoxymethylene polymers, POM, commonly known as polyacetals or ‘Acetal’ resins are linear thermoplastic polymers containing predominantly the -CH2-O- repeat unit in their backbone. There are two types of acetal resins available commercially: (1) homopolymers made by the polymerization of formaldehyde, followed by endcapping, (2) copolymers derived from the ring opening polymerization of trioxane (a cyclic trimer of formaldehyde), and a small amount of a comonomer such as ethylene oxide. Acetal resins are

1093

Commercial Polymer Blends

highly crystalline polymers with melting points of 160-175°C and heat distortion temperatures of 110-125°C at 1.82 MPa, unfilled. Due to highly crystalline nature, POM resins exhibit excellent rigidity, hardness and resistance to creep, fatigue and chemical attack. They also exhibit low wear and friction, high dimensional stability and good electrical properties. Because of their excellent mechanical properties and moderate cost, acetal resins are among the more widely used engineering resins with an estimated current global consumption of over 370 kton/y. Owing to their high crystallinity, POM resins are not miscible with any of the commercially available polymers. Unmodified POM resins tend to be brittle, particularly when notched. Due to their lack of reactivity, POMs are generally not amenable to any chemical modification by postreactions with grafting agents. Hence there are very few commercial blends based on POM resins, with the exception of impact modified POMs that are simple blends containing elastomeric/ impact modifiers. In order to improve the notched Izod impact strength, several types of impact modifiers have been employed, which included core-shell rubbers of acrylic type [Kusumgar, 1987] polymethylmethacrylate-g-styrene/butadiene [Schuette et al., 1986] or polymethylmethacrylateg-polybutadiene type [Burg, 1985]. Some commercial medium impact grades with notched Izod impact strengths of 100-150 J/m, may contain such impact modifiers. A commercial grade of high impact (notched Izod > 900 J/m) POM resin (Delrin® 100 ST, DuPont) is believed to be a blend of POM with • 30 wt% of a thermoplastic poly(esterurethane) elastomer derived from poly(1,4-butane adipate) diol and methylene-bis-(4,4’-diphenyl diisocyanate) (MDI) [Flexman, 1989]. This blend is reported to have a co-continuous or semi-interpenetrating network of the elastomer in a matrix of the polyacetal [Flexman, 1990]. The toughening effect in such a blend of IPN type morphology was interpreted to occur partly through a rubber band mechanism by which the fracture energy is absorbed. The bands of rubbery domains were believed to span the crack and participate in the

deformation process. Typical properties of commercial impact modified POM resins are shown in Table 15.29. With the increase in the impact strength of these blends, there is a corresponding decrease in the modulus, strength and DTUL relative to the neat POM resin. Impact modified POM blends are still low volume in usage relative to unmodified POM. About 80% of the impact modified POMs are used in the automotive area in typical applications such as electrical switches, fuel system components, gears and hardware. Industrial applications include cams, gears, valves, impellers, pumps and a variety of plumbing and appliance parts.

15.11

Polyphenyleneether (PPE) Blends

PPE is the generic name for the homopolymer, poly(2,6-dimethyl,1,4-phenylene ether) derived from the oxidative coupling polymerization of 2,6-dimethyl phenol [Hay, 1976]. Developed in the early 1960’s, the polymer had many desirable properties such as a high Tg (205°C) and DTUL (174°C at 1.8 MPa), high strength and dimensional stability, moisture and hydrolysis resistance and inherent flame retardancy. However, because its extremely high melt viscosity dictated melt processing temperatures of well above 300°C, at which the polymer tends to degrade or crosslink in the presence of air, its use as a molding resin by itself was severely limited. Hence there was a major motivation for blending PPE with other thermoplastic polymers, to take advantage of its high performance properties and yet combine some useful melt processability features. Several blends of PPE have been investigated stemming from the initial success of PPE/polystyrene blends. Table 15.30 lists some of the currently commercial blends of PPE of different types, comparing their key properties. 15.11.1 PPE/PS or HIPS Blends Since the early discovery of miscibility between the low-cost polystyrene and PPE, several commercial grades of PPE/HIPS have been developed

1094

M. K. Akkapeddi

Table 15.29. Properties of commercial acetal/elastomer blends vs. acetal Blend Type

POM/TPU ASTM TEST METHOD

UNITS

PHYSICAL Density

D792

kg/m3

Water Absorption, 24 hrs.

D570

%

MECHANICAL Flexural Modulus Flexural Strength Tensile Strength at Yield Elongation at Break Rockwell Hardness

D790 D790 D638 D638 D785

MPa (kpsi) MPa (kpsi) MPa (kpsi) %

D256

J/m(ft-lb/in)

PROPERTY

IMPACT Izod Impact, Notched at 23°C THERMAL Heat Deflection Temp at 0.45 MPa at 1.82 MPa Coefficient of Thermal Expansion ELECTRICAL Dielectric Strength Dielectric Constant at 1 MHZ Dissipation Factor at 1MHZ Volume Resistivity Arc Resistance

D648

POM/RUBBER

POM

DELRIN 500T Dupont

DELRIN 100ST Dupont

ULTRAFORM N2640X BASF

DELRIN 500 Dupont

1,390 0.3

1,340 0.4

1,360 0.25

1,420 0.25

2410(350) 69(10) 58(8.4) 260 R119

1240(180) 40(5.8) 45(6.5) 260 R108

1690(240) 48(6.8) 70

2880(410) 100(14.3) 70(10) 40

112(2.1)

908(17)

152(2.8)

76(1.4)

169 100 1.1E-04

145 90 1.2E-04

155 85 1.9E-04

172 136 1.0E-04

15.8 3.6

18.9 4.1 0.007 2.3E+12 120

19.6 4.2 0.009 1.0E+12

3.7 0.005 1.0E+13

°C

D676

m/m/°C

D149 D150 D150 D257 D495

KV/mm

ohm-m sec

which offer a wide choice of heat resistance (DTUL), impact strength and melt processability [Cizek, 1969; Fried et al., 1978]. This versatility of PPE/HIPS blends led to their unparalleled commercial success, accounting for nearly 50% of market volume of all engineering polymers commercial blends.

2.0E+12 120

PPE/HIPS blends filled the price-performance gap between the styrenic resins (HIPS, ABS) and the engineering resins such as polycarbonate, polyarylate and polysulfones. The technology and applications of PPE/HIPS blends have already been discussed under the styrenic resin blends section (Table 15.3).

1095

Commercial Polymer Blends

Table 15.30. Comparison of different types of commercial PPE blends Blend Type

PROPERTY

ASTM TEST METHOD

UNITS

D792

kg/m3

D955 D570

% % %

D790 D790 D638 D638

MPa (kpsi) MPa (kpsi) MPa (kpsi) %

D256

J/m(ft-lb/in)

PHYSICAL Density Mold Shrinkage Water Absorption, 24 hrs. Water Absorption at Saturation MECHANICAL Flexural Modulus Flexural Strength Tensile Strength at Yield Elongation at Break IMPACT Izod Impact, Notched at 23°C at -40°C THERMAL Heat Deflection Temp at 0.45 MPa at 1.82 MPa Vicat Softening Point

D648

D1525

PPE/HIPS

PPE/PA-6,6

PPE/PA-6

PPE/PBT

NORYL 731 GE

NORYL GTX910 GE

DIMENSION D9000 AlliedSignal

GEMAX MX4315 GE

1,060 0.6 0.07

1,100 1.4 0.5 2.8

1,080 1.6 0.5 2.9

1,150

2530(360) 95(13.5) 60(8.6) 60

2135(310) 76(11) 60(8.6) 60

1900(275) 70(10.5) 56(8) 50

1860(270) 69(10) 42(6.1) 40

270(5) 130(2.5)

220(4) 137(2.5)

550(10) 275(5)

640(12) 130(2.5)

137 129 140

193 143 232

185 121 210

152 93

Poor

Excellent

Excellent

Good

°C

°C

CHEMICAL RESISTANCE

15.11.2 PPE/PA Blends Commercial PPE/PA blends were developed by the motivation to combine the high heat resistance characteristics of PPE with the chemical resistance characteristics of the crystalline polyamide polymers (PA-66 and PA-6). Because of the inherent incompatibility between PPE and polyamides, suitable methods of compatibilization and toughening have not been developed until recently. The technology of compatibilized, impact-modi-

fied PPE/PA blends have already been discussed under polyamide blends [see 15.8.4]. Commercial PPE/PA blends are based primarily on the lower cost polyamide (PA-6 and PA-66) and most often include a rubbery impact modifier as the third blend component, added for a desired level of toughness (Table 15.21). The unique difference between PPE/PA blends and PPE/HIPS blends is illustrated by (a) the differences in DTUL at 0.45 MPa and the DTUL at 1.82 MPa in glass reinforced compositions (b) the

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M. K. Akkapeddi

difference in their relative sensitivity or resistance to chemicals, e.g., some common solvents and automotive fluids (Table 15.23). These differences arise from the facts that (a) polyamide is a crystalline polymer unlike HIPS which is amorphous and (b) due to the large melt viscosity difference between polyamide and PPE at the normal blend ratios, the polyamide forms the continuous phase. Hence in the molded parts, the polyamide surface offers resistance to solvent permeation, and high softening temperature. In PPE/HIPS blends due to the single phase amorphous character of the matrix, the solvent resistance is limited and the heat resistance is limited by the Tg of the blend. Because of these differences, PPE/PA blends have found significant application niches.

partial miscibility and therefore contributes to the over-all toughness. It has also been reported that polycarbonate encapsulates the PPE dispersions in the PBT matrix, and consequently improves the compatibility of the blend [Hobbs et al., 1992]. The addition of PC to PPE/PET or PBT blends indeed leads to some sacrifice in heat resistance and some loss in tensile and impact performance during thermal cycling or reprocessing. Although General Electric Co. announced the properties of a developmental grade of PPE/PBT blend (Gemax®), no commercial grades are available as yet. The typical properties of the developmental (semi-commercial) grades of PPE/ PBT are illustrated in Table 15.30. 15.11.4 PPE/Polyolefin Blends

15.11.3 PPE/Polyester Blends Thermoplastic polyesters, PET and PBT also offer the advantages of easy melt processability and good solvent resistance owing to their semicrystalline nature. In addition, PET and PBT are relatively moisture insensitive under ambient conditions, unlike the polyamides (PA-6 and PA-66) which can absorb significant levels of moisture. Moisture absorption in polyamides leads to significant growth in dimensions and loss in modulus and strength. Although PPE/PA blends exhibit lower moisture and dimensional growth than the standard polyamides, they are still not suitable for many exterior automotive applications such as door (PA-6 and 66) panels and many electrical and electronic applications where their moisture absorption and dimensional growth are still unacceptable. Several research investigations have been made to compatibilize PET or PBT with PPE both by reactive and non-reactive routes of compatibilization [Brown et al., 1990 and 1991; Akkapeddi and VanBuskirk, 1992]. Compatibilized binary blends of PPE/polyesters still lacked adequate toughness and invariably required the addition of rubbery impact modifiers (reactive or compatible type) and polycarbonate. The addition of polycarbonate presumably suppresses the crystallization of the thermoplastic PET or PBT phase, due to its

Several approaches to compatibilizing PPE blends with commercial polyolefins (polypropylene, etc.) have been reported in the literature [Lee, 1990; Kirkpatrick, 1989]. However, no commercial blends of PPE/polyolefins have been offered to date. Compatibilization and impact modification of PPE/polypropylene can be achieved by choosing a selected type of styrene-ethylene/butylenestyrene block copolymer and PPE of low molecular weight [Akkapeddi and VanBuskirk, 1992]. 15.11.5 PPE/PPS Blends Some developmental blends of PPE and PPS [Noryl® APS4300 GEC] have been reported to be in the test marketing phase [Gabrielle, 1992]. These alloys have been made from specially functionalized versions of each resin and proprietary compatibilizers. The addition of the amorphous PPE was reported to improve the ductility of PPS while reducing the flash and shrinkage during molding. The blends were primarily developed in glass fiber reinforced forms. 15.11.6 PPE/Epoxy Blends PPE has been blended with commercial epoxy formulations [Chao et al., 1993] presumably to improve dielectric properties (lower dielectric

Commercial Polymer Blends

constant), toughness and moisture resistance of the cured epoxy resins. These formulations have been evaluated with fiberglass reinforcements for the printed circuit boards and other electronic applications. In one formulation of epoxy resin (Epon 825, Shell) cured with aluminum alkoxide, incorporation of ca. 30% PPE increased the tensile elongation at break for the epoxy thermoset from < 2 % to 17%. The heat distortion temperature was increased from 60 to 195°C. The dissolution of PPE in the epoxy raised its viscosity from 0.2 to 400 Pa.s [Anon., 1991].

15.12

Thermoplastic Polyester Blends

Polyethylene terephthalate (PET) and polybutylene terephthalate (PBT) are two of the most important members of a family of commercial thermoplastic polyesters. Other members of this family include specialty polyesters such as PET copolymers (PETG) and poly(1,4-cyclohexanedimethylene terephthalate) (PCT). Commercially important PET and PBT resins are known for their high crystalline melting points (265°C and 225°C), good mechanical properties and solvent resistance characteristics. Between the two, PET is the largest volume commercial polyester with an estimated global consumption of in excess of 8 Mton/y in fiber applications and in excess of 1 Mton/y in thermoplastic applications such as in the injection blow molding of bottles and containers and, extrusion of film and thermoformable sheets [SRI, 1992]. The dramatic growth in the use of PET for soft drink bottles and in other packaging applications has also spurred considerable activity in the recycling of PET. Hence PET is a relatively low cost polymer both because of large scale economics as well as its availability as a recycled material, thus providing a cost incentive for blending with other polymers. However, due to the slow rate of crystallization, PET has not been used in injection molding applications until recently [Burke and Newcombe, 1982; Hecht and Ford, 1985]. Through the use of specific types of nucleators (e.g., sodium stearate, sodium ionomers) and other crystalliza-

1097

tion promoters, PET molding compounds are gradually gaining acceptance as injection molding resins [Deyrup, 1982; Legras, 1986; Kozielski, 1988]. On the other hand, polybutylene terephthalate (PBT) produced in much smaller volume than PET is more widely accepted as an injection moldable engineering plastic due to its faster crystallization rate [Pratt and Hobbs, 1976]. Hence PBT is more commonly employed in the formulation of blends also. PBT has generally been preferred over PET in the engineering plastics industry because of its superior processability, faster crystallization rate, shorter molding cycles and better properties (DTUL/impact balance) in the molded parts, particularly in the unfilled form. Nevertheless, PET is also used primarily in glass or mineral reinforced form along with nucleators and crystallization promoters. The primary motivations for blending the thermoplastic polyesters with other polymers are: (a) to improve the solvent resistance and processability of amorphous polymers such as PC, styrenics, PPE, etc., (b) to reduce the mold shrinkage of polyesters associated with their crystallization, (c) to increase the DTUL of unfilled polyesters, (d) to improve toughness. The largest volume polyester blend sold commercially is the PBT/polycarbonate blend. PET/ PC blend is also gaining commercial importance because it is similar to PBT/PC blend in properties and moldability but has some cost advantages. Unlike neat, unfilled PET that is difficult to injection mold due to its slow crystallization rate (long mold cycle times with hot molds and amorphous parts with cold molds), PET/PC blends can be molded readily using the normal hot molds (ca. 80 to 100°C) and fast mold cycles. The presence of PC helps retain high stiffness and strength in the part at the mold temperatures, to enable the demolding of distortion-free parts. The properties and applications of PET/PC blends have already been discussed under the polycarbonate blend section. Other commercial blends of PET and PBT are discussed as below.

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15.12.1 PBT/PET Blends These blends take advantage of the low cost of PET and the rapid crystallization rate of PBT. Despite their large difference in the crystallization rates PET and PBT form stable blends without the need for compatibilizating agent. This was attributed to the amorphous phase miscibility between the two components [Escala and Stein, 1979; Mondragon et al., 1989]. X-ray, DSC and IR studies indicated that the two components form separate crystalline phases and a single amorphous phase with a single Tg. Some transesterification was detected in the melt by NMR, especially at long melt residence times (• 6 min) [Mondragon et al., 1989]. The extent of transesterification under the fast extruder blending operation is however expected to be low. Commercial grades of PET/PBT blends are generally glass fiber reinforced (15-30 wt%). Compared to the glass reinforced PBT and PET, the heat distortion temperatures of the blends at 1.8 MPa are actually lower indicating the lower level of net crystallinity in the blend, an effect possibly caused either by miscibility or transesterification. The primary reason for developing these blends appears to be the improvement of surface appearance and gloss in the injection molded parts compared to those made from the individual resins. There is also some cost advantage over PBT. PBT/PET blends are used for making visible parts of both large and small appliances that need the appeal of smooth and glossy surface along with high stiffness, strength, and DTUL. There are also other electrical and automotive applications. Compared to the neat PBT molding resins, the market for the blend is still relatively small. 15.12.2 PBT/Elastomer Blends Unmodified PBT is a fairly ductile material exhibiting high elongation at break, even after crystallization. However, as to be expected of all rigid semi-crystalline polymers, molded parts of PBT show low notched Izod impact strength indicating that under conditions of stress concentration, the

resistance to unstable crack propagation is low. Impact modification of PBT was investigated through the use of several elastomeric modifiers. Commercial PBT/elastomer blends are of two types viz. (a) high impact strength type, (b) low modulus, highly flexible types. Commercial impact modified PBT grades generally contain 20 to 30 wt% of controlled particle size (< 0.3 µm), core-shell rubber modifiers [Neurey and Ott, 1981; Farnham and Goldman, 1978; Binsack, 1985]. Typical impact modifiers are: PMMA-g-SBR (MBS), PMMA-g-poly(n-BuA) (acrylic core-shell rubbers), SAN-g-PBD (high rubber ABS) or SAN-g-poly(n-BuA) (high rubber ASA). Commercial impact modified PBT grades (Table 15.31) exhibit notched Izod impact strengths in excess of 500 J/m, while retaining a good level of modulus, strength and DTUL. In these blends, the dispersion of rubber particles promotes multiple sites for crazing and localized shear yielding in the PBT matrix, thus providing mechanisms for energy dissipation during impact deformation and hence offering high resistance to crack propagation [Hourston et al., 1991]. A commercial blend of PBT with high rubber ABS as impact modifier (Pocan S1506, Bayer) has been used in Europe for automotive bumpers, mirror housings and other exterior parts [Kosoff, 1987]. The blend was reported to exhibit good heat sag resistance at 135°C and maintain high notched Izod impact strength > 700 J/m even at -29°C (Table 15.31). Substitution of ABS with ASA (SAN-grafted to crosslinked poly(n-butylacrylate) rubber particles gives a more weatherable, impact modified PBT (Ultradur KR4071, BASF). A commercial PBT blend containing ca. 25 wt% of MBS type impact modifier (Vandar® 2100, Hoechst Celanese) has been used in such exterior automotive applications as under body rivets, fuel line clips, etc. PBT/elastomer blends display a unique combination of high impact strength, dimensional stability due to their non-hygroscopic nature, excellent resistance to automotive fluids such as gasoline, oils, paint solvent, aqueous salt solutions and good heat resistance. In addition, their easy processability (low melt viscosities) lends

1099

Commercial Polymer Blends

Table 15.31. Properties of commercial PBT/elastomer blends ASTM TEST METHOD

UNITS

VALOX 357 GEC

VANDAR 2100 Celanese

TORAY PBT 5207X11 Toray

POCAN S1506 Albis Corp

PHYSICAL Density Mold Shrinkage Water Absorption, 24 hrs.

D792 D955 D570

kg/m3 % %

1,290 1.2 0.08

1,210 1.8 0.1

1,200 2.6 0.08

1,200 1.8 0.1

MECHANICAL Flexural Modulus Flexural Strength Tensile Strength at Break Elongation at Break

D790 D790 D638 D638

MPa (kpsi) MPa (kpsi) MPa (kpsi) %

2100(300) 84(12) 37(5.3) 200

1930(280) 61(8.7) 43(6.3) 150

1660(235) 54(7.7) 37(5.3) 250

1540(220)

D256

J/m(ft-lb/in) 540(10)

>800(15) 160(3)

770(14)

>900(18)

98

82 48

50

58

PROPERTY

IMPACT Izod Impact, Notched at 23°C at -40°C THERMAL Heat Deflection Temp at 0.45 MPa at 1.82 MPa

D648

38(5.4) 100

°C

themselves to the fabrication of both small and large parts. High mold shrinkage and tendency to warp are some of the drawbacks resulting from the crystallization phenomenon, which need to be addressed by proper design of the tools and the molding process. 15.12.3 PBT/Copoly(ether-ester) Elastomer Blends Commercial copoly(ether-ester) elastomers (e.g., Hytrel®, DuPont; Lomod®, GEC) are segmented block copolymers containing a polyether soft segment such as poly(tetramethylene oxide) and a hard segment that is chemically identical to polybutylene terephthalate. When the soft segment is • 50%, these block copolymers exhibit

high degree of rubbery elasticity, yet they process like thermoplastics and hence appropriately are called thermoplastic elastomers [Adams and Hoeschele, 1987]. Because the hard segments crystallize like PBT, these thermoplastic elastomers exhibit high heat resistance with Vicat softening points typically in the range of about 180°C. The hard segment crystallinity also imparts a good level of solvent resistance in these materials. The low Tg of the soft segment is responsible for the flexibility, resilience and lowtemperature toughness characteristics. Although the poly(ether-ester) block copolymers have been used in many niche applications requiring high performance thermoplastic elastomer characteristics, their high cost has been a drawback in extending to larger volume applications.

1100

M. K. Akkapeddi

To reduce the cost, these elastomers have been diluted with some PBT homopolymer. Because of the chemical similarity between the hard segment of the copoly(ether-ester) elastomers and the PBT, they form fairly compatible blends. When the hard segment content in the copoly(ether-ester) is > 80 wt%, it was found to be completely miscible with PBT, showing a single Tg, amorphous phase and co-crystallization of the PBT segments of the elastomer with PBT homopolymer. As the hard segment content was lowered to ” 60%, the blend exhibited incomplete miscibility, with two Tg’s for two amorphous phases and also two separate crystalline phases. [Runt et al., 1989]. Nevertheless, a partial miscibility was indicated due to changes in the Tg observed in DSC and dielectric relaxation spectra. The partial miscibility and low interfacial tension between the phases makes the blend very compatible. Commercial PBT/copoly(ether-ester) blends are generally richer (• 50%) in the copoly(ether-ester) elastomer content. These blends were designed for the automotive, flexible bumper fascia market (Bexloy® V, DuPont). Typical properties of these blends are shown in Table 15.32. Typically they exhibit low moduli (” 800 MPa), high elongation and toughness. The purpose of PBT in the blend is to lower the cost and improve the heat sag resistance required for paint oven capability. PBT/copoly(ether-ester) elastomer blend molded parts exhibit excellent surface finish and good paint adhesion without the need for primers. About 3 kton/y of this blend is being used for painted, flexible bumper fascias in selected luxury model cars in the USA. Because of the excellent surface esthetics of the blend, molded-in colors are also being evaluated to reduce the painting costs. Lower levels (” 20%) of copoly(ether-ester) elastomers have also been blended with PBT for high impact strength molding resin applications (Celanex®, Hoechst-Celanese). 15.12.4 Glass Reinforced, Impact Modified PET (and PBT) PET/Elastomer blends have not been commercialized in the unfilled form, due to the slow rate

of crystallization of PET. Unfilled PET and PET/ elastomer blends are not easy to injection mold under normal mold temperatures (ca. 80-100°C) in fast molding cycles. The parts tend to stick to the mold and distort. Use of cold molds allows molding of amorphous PET parts (1-2 mm) lacking heat resistance (DTUL). Upon annealing at elevated temperatures (ca. 150°C), one can develop crystallinity in PET parts, but they become brittle even in the presence of a modifier. Lack of proper adhesion between the rubber and the PET after crystallization, in general, seems to be the reason for this embrittlement. Use of reactive tougheners such as ethyleneglycidyl methacrylate copolymers [Iida, 1981] and ethylene-n-butyl methacrylate-glycidyl methacrylate terpolymer [Deyrup, 1988], leads to significantly improved toughness which is retained even after annealing [Akkapeddi and VanBuskirk, 1993]. However, at present no commercial PET/elastomer blends are offered in the unfilled form. The compositions such as those described above may be nucleated and glass filled. An ethylene copolymer rubber modified, glass filled PET (Rynite® SST) with improved notched Izod impact strength and elongation is commercially available. A glass filled PBT, impact modified with SAN-g-poly(n-BuA) core-shell rubbers (ASA) is also available commercially (Ultrablend®S, BASF). Because of the superior weatherability of ASA rubber vs. other rubbers such as ABS, MBS, etc., the PBT/ASA blend is likely to find applications in the exterior automotive applications. Mirror housings, door handles, roof racks, are typical exterior, automotive applications. Some grades of glass/mineral reinforced, impact modified PET molding resins have also been developed specifically for automotive exterior body panel applications (Bexloy®K, DuPont). This specific formulation was reported to withstand automotive on-line “E-coat” paint oven temperatures (ca. 200°C) as well as give low warpage and smooth surface finish in the molded parts. A combination of glass fiber or glass beads and/or mica is believed to be used for reinforcement. The impact modifier is more likely a reactive toughener of the ethylene-n-butylacrylate

1101

Commercial Polymer Blends

Table 15.32. Commercial PBT/polyester elastomer blends vs. PBT and polyester elastomer Blend Type

PBT/POLYESTER ELASTOMER

ASTM TEST METHOD

UNITS

PHYSICAL Density

D792

kg/m3

Mold Shrinkage Water Absorption, 24 hrs.

D955 D570

% %

MECHANICAL Flexural Modulus Flexural Strength Tensile Strength at Yield Tensile Strength at Break Elongation at Break Shore D Hardness

D790 D790 D638 D638 D638

MPa (kpsi) MPa (kpsi) MPa (kpsi) MPa (kpsi) %

D256

J/m(ft-lb/in)

PROPERTY

IMPACT Izod Impact, Notched at 23°C at -29°C THERMAL Heat Deflection Temp at 0.45 MPa at 1.82 MPa Vicat Softening Point

D648

LOMOD A-1220 GE

RITEFLEX BP H. Celanese

HYTREL 5526 Dupont

1,200 1.2 0.25

1,200

1,180

1,200

1,300

0.5

0.08

760(110)

828(120)

760(110)

211(30)

33(4.8)

25(3.6)

400 D59

200

200 D67

10(1.5) 41(5.8) 560 D55

2300(330) 84(12) 56(8) 56(8) 50 M72

NB 187(3.5)

NB 203(3.8)

NB

NB

32(0.6)

101 49 180

54 210

0.32

CELANEX 1400 H. Celanese

°C

°C

glycidyl methacrylate terpolymer type or an ionomeric type ethylene-butylacrylate-methacrylic acid terpolymer.

15.13

PBT

BEXLOY V978 Dupont

111 44 D1525

POLYESTER ELASTOMER

Specialty Polymer Blends

In the plastics industry, specialty polymers are generally considered as high performance, high priced resins whose current market volume is

176

still relatively low compared to the more widely used engineering thermoplastics and commodity thermoplastics. Specialty polymers may be of two types viz., (a) high Tg, amorphous engineering thermoplastics such as polysulfones, polyarylates, polyetherimide, polyamideimide, polyimides; (b) high melting, crystalline thermoplastics such as polyphenylene sulfide, liquid-crystalline polyesters (LCP’s), polyetherketones (PEEK, PEK). Both categories of polymers can be classified

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M. K. Akkapeddi

as high temperature polymers with long term service capability at T • 150°C and more often at T • 200°C, as defined by U.L. temperature index; i.e., temperature at which • 50% tensile strength is retained after 10,000 hours of heataging. A common structural feature of all these high temperature polymers is their essentially all aromatic backbone structures. The high temperature, specialty polymers possess certain common property advantages such as: (i) high Tg and/or Tm, (ii) inherent flame retardancy, (iii) high thermo-oxidative stability, (iv) high continuous use temperatures, (v) high mechanical rigidity and strengths, (vi) high dimensional stability (to moisture and temperature) and creep resistance. The crystalline polymers such as PPS, LCP, PEEK offer the additional advantages of high solvent resistance. Due to the inherently high cost of the specialty polymers, very few blends have been developed for commercial applications. The only driving force for the development of even the few blends of specialty polymers has been the desire to reduce the cost of the base resins by blending with lower cost engineering plastics, although this invariably results in a lower DTUL. Nevertheless, a few commercial blends of specialty polymers exist and their properties will be discussed below: 15.13.1 Polysulfone Blends Polysulfones are aromatic high Tg, amorphous polymers having a sulfone and an aromatic ether linkage in the recurring units of their backbone. Owing to this structure of their polymer backbones, they display unique resistance to hydrolysis. The two commercially significant polysulfones are (a) the polysulfone derived from bisphenol A, which is simply referred to as polysulfone (PSO; Udel®, Amoco) and (b) the polyethersulfone (PES; Victrex®, ICI). The current world-wide consumption of all polysulfones is estimated to be less than 19 kton/y about 70% of which belong to the polysulfone (PSO). The hydrolytic stability of polysulfones appears to be their unique advantage over other high heat, transparent thermoplastics

such as polycarbonate and polyestercarbonates. Hence the polysulfones have been used primarily in specialty niche applications such as steam sterilizable medical equipment, cookware, appliance parts, etc. These unique advantages of polysulfones in having transparency coupled with hydrolytic stability and heat resistance, would be lost if other polymers are blended. Hence there has been very little incentive for blends of polysulfone. In addition, the cost factor has also been a deterrent. The only commercial blends of polysulfone are those containing ABS or PET (Mindel®, Amoco), which together make up for about 1 kton/y of current market volume in the USA. 15.13.1.1 Polysulfone/ABS Blend The primary reasons for developing this blend appear to be: (a) to improve the toughness, particularly the notched Izod impact strength (b) processability (c) plateability of polysulfone. Since ABS is a low cost, tough, plateable resin, it meets the requirements. However, ABS and polysulfone lack adequate compatibility, which manifests itself in poor weld-line strength and surface appearance in the molded parts. Various types of compatibilizing additives such as phenoxy resins, styrene-maleic anhydride copolymers have been claimed to improve the weld line strength of the polysulfone/ABS blends [Robeson, 1985]. Commercial polysulfone/ABS blend (Mindel(R) A) is composed of PSO, high heat ABS (based on α-methylstyrene) a phenoxy resin compatibilizer and MBS core-shell rubber as impact modifier. The plateability and hydrolytic stability advantages of polysulfone/ABS blends have been utilized to make selected appliance, plumbing, and sterilizable equipment parts, replacing polycarbonate. Polysulfone/ABS blend has also been evaluated for plateable components of electrical and electronic parts, circuit boards and connectors. However, the lack of adequate heat resistance (DTUL lower than polysulfone) for vapor phase solderability limited its use only in conjunction with other high heat resins as laminates.

Commercial Polymer Blends

15.13.1.2 Polysulfone/PET Blend Polysulfone has been blended with PET for the purposes of cost reduction and improvement in solvent stress-crack resistance. The blend was developed primarily in the glass reinforced forms for applications in electrical and electronic markets (Mindel® B, Amoco). However the DTUL, flexural strength and impact strength of the commercial polysulfone/PET blend appears to be lower than those of PET at comparable levels of glass reinforcement (Table 15.33). Lack of an adequate compatibility between PET and polysulfone and a net reduction in the crystallinity account for the lower properties. However, below the Tg of polysulfone, e.g., at T ” 150°C, the dimensional stability (modulus retention) should be somewhat better than that of glass filled PET, which loses some of its initial modulus above its Tg of ca. 75°C. Other advantages of glassfilled PSO/PET blend include lower warpage and shrinkage, better flame resistance than glass-filled PET or PBT. PSO/PET blends have also been reported to exhibit superior hydrolysis resistance relative to PET or PBT. 15.13.2 Polyarylate Blends Commercial polyarylate is an aromatic polyester of high glass transition temperature (ca. 180°C) derived from bisphenol A and a mixture of terephthalic and isophthalic acids. It is a transparent, rigid and tough thermoplastic of high heat distortion temperature (174°C). Polyarylates face competition from the more established polycarbonate and its higher heat analogues, viz. polyester carbonates; as well as the polysulfone that is superior in hydrolic stability. Although polyarylate has been commercially available for many years, its market growth has been slow due to its high cost/performance balance. Nevertheless, polyarylate’s transparency, UV/weather resistance and high heat distortion temperature properties have found specialty niche applications in automotive taillights, reflectors, etc. Polyarylates have also been used in electrical/electronic connector applications. Current world-wide production of

1103

polyarylates is estimated below 2 kton/y. Relatively few blends of polyarylate have been commercialized because by blending other polymers results in the loss of the clarity and DTUL advantages of the polyarylate. The only commercially significant blend of polyarylate appears to be a blend with PET (U 8000, Unitika; Ardel® D-240, Amoco). Polyarylate/PET blends prepared by solution or melt blending under short residence times at T ” 280°C with or without an added ester interchange inhibitor such as triphenylphosphite, are essentially phase-separated, exhibiting two glass transition temperatures, one each for a PET phase and a polyarylate-rich phase. From the observed glass transition temperatures, one can conclude that it is a partially miscible blend in which more PET dissolves in the polyarylate phase than polyarylate does in PET. The interaction parameter has been estimated to be slightly positive (χ12 ≅ 0.1) [Chung and Akkapeddi, 1993]. The miscibility between polyarylate and PET may be further driven by transesterification reaction within the melt phase, which occurs slowly at T ” 280°C but more rapidly at T • 300°C [Robeson, 1985; Eguiazabal et al., 1991]. In any case, polyarylate can readily form transparent blends when the PET content is ” 30% and the melt blending done above 300°C. Hence by adjusting the melt blending conditions, PET can be used to lower the cost and improve the chemical resistance of polyarylate while maintaining an adequate level of transparency. The heat distortion temperature is, of course, sacrificed to some extent. Except for some improved chemical resistance, processability and cost the blend does not seem to offer any compelling advantages over polyarylate and hence its applications appear to be quite limited. 15.13.3 Polyetherimide Blends Commercial polyetherimide (PEI Ultem®, GEC) is an amorphous, high performance thermoplastic with its repeat unit structure containing both the rigid aromatic imide units and the somewhat more flexible aromatic ether units.

1104

M. K. Akkapeddi

Table 15.33. Commercial polysulfone blends vs. polysulfone and PET Blend Type

PSF/ABS MINDEL A670 Amoco

PSF/PET MINDEL B322* Amoco

PSF UDEL P1700 Amoco

PET PETRA 130* AlliedSignal

% %

1,130 0.66 0.25

1,470 0.25 0.14

1,240 0.7 0.3

1,550 0.3 0.05

D790 D790 D638 D638 D638

MPa (kpsi) MPa (kpsi) MPa (kpsi) MPa (kpsi) %

2220(316) 84(12) 51(7.2) 44(6.2) 75

7030(1000) 162(23) 105(15) 105(15) 2

2740(390) 108(15.4) 72(10) 75

9130(1300) 240(34) 154(22) 154(22) 2

IMPACT Izod Impact, Notched at 23°C

D256

J/m(ft-lb/in) 380(7)

54(1)

71(1.3)

98(1.8)

THERMAL Heat Deflection Temp at 1.82 MPa

D648

149

160

173

225

3.7 0.009

3.5 0.02

ASTM TEST METHOD

UNITS

PHYSICAL Density

D792

kg/m3

Mold Shrinkage Water Absorption, 24 hrs.

D955 D570

MECHANICAL Flexural Modulus Flexural Strength Tensile Strength at Yield Tensile Strength at Break Elongation at Break

PROPERTY

ELECTRICAL Dielectric Strength D149 Dielectric Constant at 1 MHZ D150 Dissipation Factor at 1 MHZ D150

°C

KV/mm

* 30% Glass filled

Owing to the unique structure of its backbone, polyetherimide exhibits high glass transition temperature (Tg = 215°C), high mechanical strength and rigidity, yet has a good degree of ductility and melt processability. Its highly aromatic backbone structure also imparts an inherent flame retardancy and low smoke generation characteristics in the polymer. Because of these high performance characteristics, PEI finds specialty niche applications in the automotive under-the-hood, aerospace, electronic and medical equipment markets.

However, its relatively high cost has limited its market volume. Very few blends of PEI have been commercialized because blending a lower cost polymer generally compromises the performance of the base resin. Nevertheless, some PEI/PC blends have been introduced as thermoformable sheets (Ultem® 1668, GEC) for aircraft interior applications, to meet the processability and the low flammability/low smoke generation requirements (Table 15.28). Other applications in the automo-

Commercial Polymer Blends

tive and the industrial markets have also been found for this blend, the details of which have already been discussed under the polycarbonate blends section. Polyetherimide was found to be miscible with polyetheretherketone (PEEK) exhibiting a single Tg. Since PEEK is a semicrystalline polymer with a Tg of ca. 150°C, the blend should have a higher Tg (intermediate between that of PEEK and PEI) and yet possess some solvent resistance characteristics of PEEK. The blend may have some advantages in film and composite applications, but no commercial application of this blend have been developed to date. 15.13.4 Polyimide/Poly(tetrafluoroethylene) Blends High performance polyimides (Vespel®, Dupont) and polyamide-imide (Torlon®, Amoco) molding compounds are often mixed with polytetrafluoroethylene (PTFE) to fabricate a variety of lowfriction bearings, bushings and seals used in automotive, aerospace and industrial markets. The role of PTFE in these formations is simply to function as a lubricating filler. Hence these may not be considered as real blends. The PTFE particles are inert and not bonded to the polyimide matrix in the molded part. Because of this heterogeneity, higher amounts of PTFE can not be tolerated in the polyimide compounds, due to undesirable loss in tensile strength. Generally only about 10% PTFE is used, often in conjunction with additional lubricants such as graphite and/or molybdenum sulfide. 15.13.5 Polyphenylenesulfide Blends Polyphenylenesulfide (PPS) is an aromatic semicrystalline polymer with a recurring sulfide linkage in the backbone. Due to the highly aromatic nature of its structure, PPS is inherently flame retardant and also exhibits an outstanding level of chemical resistance. PPS is high melting (Tm= 285°C), yet its low melt viscosity allows easy processing and high loading of glass or mineral fillers. Since unfilled PPS is too brittle to be

1105

used by itself, PPS is generally compounded with reinforcing fillers. Glass reinforced PPS exhibits high heat distortion temperature (HDT ≅ 265°C) and high continuous use temperature (U.L. index of ca. 200°C). Owing to its unique combination of properties, PPS is experiencing a recent growth in the market interest with several producers and suppliers entering this business. Although current worldwide consumption of PPS is still relatively low (below 7 kton/y based on the neat resin) steadily improving cost and supply position of PPS may increase its usage in the future. Since most PPS is used in the glass or mineral reinforced form, blending other polymers generally has no benefits. In addition, the chemical inertness of PPS and the crystallization tendency do not promote any degree of compatibility with other polymers. The only commercial blend of PPS currently being used is the blend of PPS with PTFE. However in these formulations, PTFE is simply added as a lubricating filler. These products (Dianippon Ink, Japan) are used for making low-friction gears (for floppy disc drives), bearings, relays and other moving parts. Since a major weakness of PPS is its brittleness, some attempts have been made to improve its toughness by blending with other suitable polymers. A commercial impact modified PPS (Toray) is believed to consist of a blend of PPS with ethylene-glycidylmethacrylate polymer. A grafting reaction is expected to occur if the PPS has active end groups such as -SH or -S-Na+, which can in principle react with the epoxy group of the ethylene/GMA copolymer. Recently a PPS blend with liquid crystalline polyesters (LCP) has also been offered commercially in 40% glass filled form (Vectra® V140, Hoechst-Celanese). The liquid crystalline polyester used in this blend is a copolyester of p-hydroxybenzoic acid and 2-hydroxy-6-naphthoic acid. The melting point of this LCP and of PPS closely match, i.e., Tm = 285 to 290°C. Since there is no compatibility or reaction between the two components, LCP/PPS blend is considered to be a simple mechanical blend.

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The purpose of blending PPS seems to be simply to lower the cost of LCP without sacrificing the high heat distortion temperature and high melt flow (Table 15.34). The commercial LCP/PPS blend is designed for injection molding complex electronic parts, chip carriers, sockets and coil bobbins. It is very likely that due to its low melt viscosity, LCP still forms the continuous phase, while PPS may simply be present as a dispersed filler along with the glass fibers. Very little information about this blend has been published. There has been some commercial development activity in the blends of PPS with PPE, in which the PPE is claimed to improve the ductility of PPS, e.g., with an elongation at break of 8% in a 40% glass fiber reinforced blend [Gabriele, 1992]. The contribution of high ductility, high DTUL of 270oC and good processability (low mold shrinkage and flash), appears to make this blend to be a significant improvement over glassfilled PPS. However, no further details on the commercial usefulness of this blend have been reported to date. 15.13.6 Liquid Crystalline Polyester Blends Although liquid crystalline polyesters (LCP’s) are interesting polymers exhibiting high mechanical strength and modulus due to a high degree of selforientation, no commercial blends of LCP (except with PPS) are available. LCP blends have been

investigated by many researchers, but their high cost has precluded the successful commercialization of any such blends.

15.14

Thermoset Blend Systems

Thermoset systems are, by definition, threedimensional polymer networks formed by thermal radiation or chemically induced polymerization and crosslinking (curing) reactions of multifunctional monomers or prepolymers. Since a thermosetting resin solidifies upon curing and cannot be remelted or reprocessed, it is a necessary to mix all the required ingredients including the reinforcing fillers or fibers, modifiers, stabilizers and additives during the initial monomeric or prepolymer stage before being fabricating into the final shape by curing. The mixing is facilitated by the low viscosity of the system. Among the most widely used commercial thermosets are: (a) unsaturated polyesters, (b) phenolics, (c) epoxies, (d) vinyl esters, (e) polyurethanes, (d) polyimides, etc. Many commercial thermosets are quite often used as complex mixtures of several co-reacting monomers and prepolymers, specifically formulated to suit a given end-use application. For example, in coating and adhesive applications, often mixtures of different epoxies, differing in chemical structure and/or molecular weights, are used as required. Even two different thermoset-

Table 15.34. Typical properties of LCP/PPS blend (Vectra® V140, Hoechst-Celanese) (40 wt% glass fiber reinforced) Density Water absorption, 24 hr Flexural Modulus Flexural Strength Tensile Strength Tensile Elongation DTUL Dielectric Constant Dielectric Strength

(kg/m3) (%) (MPa) (MPa) (MPa) (%) (°C, 1.8 MPa) (at 1kHz) (kV/mm)

1,670 0.02 16,550 248 165 1.4 265 3.7 23.6

Commercial Polymer Blends

ting monomers may be mixed such as in the case of single-package epoxy-phenolic molding compounds [Fry et al., 1985]. However, even these hybrid thermoset systems are not usually considered as blends. Since the molecular weights of the epoxy and novolaks are low and during curing they become integral parts of the polymer network through co-reaction, the system may be called a co-reacted thermoset. Although it is sometimes hard to distinguish some thermoset blends from co-reacting thermosets, thermoset blends can be ideally classified as follows: A) Thermoset-thermoset blends and interpenetrating networks (IPN’s) B) Thermoset-thermoplastic blends and semiinterpenetrating networks (Semi-IPN) C) Rubber modified thermosets An interpenetrating network (IPN) is defined as a combination of two polymer networks, at least one of which is in the presence of the other. The distinction between interpenetrating network and blend may often be based on morphology and the degree of phase separation. If there is a sufficient degree of molecular interaction, the phase separation tendency is suppressed and a true molecularly or morphologically uniform interpenetrating network can be achieved. Very few systems can form truly homogeneous blend networks and, in reality, some microheterogeneity is invariably observed. Compatible IPN’s generally form < 5 nm size domains and incompatible IPN’s or blends form domain size > 30 nm. 15.14.1 Thermoset/Thermoset Blends As previously mentioned, several commercial hybrid thermosets are known to be co-reacting thermosets, i.e. when the mixture of two different thermosetting monomers or prepolymers is cured, there is a simultaneous graft or co-reaction between the components along with the crosslinking reactions. These systems may therefore be considered as co-polymerizing thermosets and not as true blends. Examples of such systems are: phenolic novolak/epoxy resin (or epoxy novolaks);

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melamine prepolymer/epoxy resin; carboxyl terminated unsaturated polyester/epoxy thermosets. The graft or copolymer network reaction involves the reaction between the phenolic-OH or an amine or a carboxyl and the epoxide group. Some examples of what may be considered as actual thermoset/thermoset blends are the epoxy and bismaleimide thermosets modified with aromatic dicyanate esters such as bisphenol A dicyanate. Aromatic cyanates crosslink by cyclotrimerization to form a network of triazine ethers. Crosslinked cyanate ester systems typically exhibit higher glass transition temperatures (Tg • 250°C), lower moisture absorption and lower dielectric constants than conventional epoxy thermosets. Hence by mixing dicyanates or their corresponding prepolymers with epoxy resins, and then co-curing, the desired level of property improvements can be achieved. Such hybrid thermosets have been used in printed circuit board manufacture. Similar hybrid thermosets consisting of bismaleimide resins mixed with bisphenol A dicyanate ester have been commercially available (BT resin, Mitsubishi gas chemical). The presence of crosslinked triazine ether network in the matrix of bismaleimide crosslinked network is believed to improve the toughness, reduce the moisture sensitivity and improve the dielectric properties of the bismaleimide thermoset without a sacrifice in heat resistance. In a majority of cases, the thermoset/thermoset blends are actually formulated by the fabricator or the end user during the fabrication and processing of such materials as composite prepregs, printed circuit boards, laminates and adhesives. The formulations and compositions are often kept proprietary and are designed to meet their own individual requirements. 15.14.2 Thermoset/Thermoplastic Blends 15.14.2.1 Low-profile/low-shrinkage Additives for Unsaturated Polyesters Unsaturated polyesters account for nearly 75% of the reinforced thermosetting composite volume

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of usage [Atkins, 1978]. Unsaturated polyesters are derived from the polycondensation of glycols such as propylene glycol and maleic anhydride with some phthalic anhydride added to control the degree of unsaturation. This unsaturated polyester prepolymer is further mixed with styrene monomer (ca. 40%) and glass fiber reinforcements and then polymerized via a radical initiator and heat to the final thermoset network. Reinforced unsaturated polyesters are generally processed as sheet molding compound (SMC) or bulk molding compound (BMC). In the SMC process, the resin, fillers, glass fibers, catalyst and other additives are mixed together and cast into a sheet sandwiched by polyethylene film. The sheet is then cured by compression molding in a matched-die mold to the required part shapes such as automotive exterior body panels. Although parts fabricated from reinforced polyester thermosets exhibit excellent rigidity, strength and toughness, the high shrinkage associated with the polymerization of the matrix causes sink marks and depressions in the surface where ribs or bosses are located. Hence for parts requiring close tolerance and surface appearance such as automotive body panels, SMC based on unsaturated polyester could not be used for a long time, until it was found that addition of some thermoplastics to the formulation helped reduce the shrinkage and produce smooth class “A” surface. Polyvinylacetate (PVAc) and vinyl acetateacrylic copolymers (VAc-A), thermoplastic polyurethanes, polyethylene, polystyrene and polycaprolactone are some of the candidates for low-profile shrinkage additives to SMC and BMC. PVAc and VAc copolymers are the most widely used thermoplastic additives. Typically a low-profile SMC recipe contains about 15% unsaturated polyester resin, 8% thermoplastic additive, 50% calcium carbonate and 27% glass fiber. It is believed that during the curing the thermoplastic forms a separate phase, which counteracts the curing shrinkage in the matrix initially by thermal expansion and subsequently by void formation. Almost all of the SMC used in the automotive industry is based on low profile, lowshrinkage additives.

Thermoplastic additives in SMC and BMC accounted for nearly 8 kton/y consumption in the USA [Skeist, 1992]. A primary requirement for the polymer additive is that it must be amorphous with a low to moderate Tg and fairly soluble or dispersible in the resin matrix initially, but capable of phase separation during the polymerization. 15.14.2.2 PPE/epoxy Blends Poly(2,6-dimethyl 1,4-phenylene ether) (PPE) is a high Tg (215°C) ductile polymer with a low dielectric constant and extremely low moisture absorption. Epoxy thermosets exhibit a very good combination of useful properties such as good adhesion, low-shrinkage, high electrical resistivity and good thermal properties. The brittleness and the moisture sensitivity of epoxies can be improved by blending a thermoplastic additive such as PPE. In addition, since PPE has a high Tg, the thermal properties are not sacrificed, but improved. Furthermore, the dielectric constant is lowered. Such blends have been used in glass reinforced composite form for laminates and printed circuit boards. In a typical formulation of epoxy resin (Epon 825, Shell) cured with aluminum alkoxide, incorporation of 30% PPE increased the elongation at break from <2% to >17%. The heat distortion temperature increased from 160 to 195°C. The dissolution of PPE in epoxy formulationraised its viscosity at 200°C from 0.2 to 4 Pas [Anonym., 1991]. 15.14.2.3 Thermoplastic/silicone Semi-IPNs Although the dynamically vulcanized blends such as EPDM/PP (Santoprene®) and NBR/PP (Geolast®) have sometimes been referred in the literature as semi-IPNs, we considered them as blends of crosslinked elastomer dispersions in a thermoplastic matrix and as such treated them under the elastomer blends. There is yet another class of thermoplastic/thermoset blend system in which a minor amount of the crosslinkable monomer(s) is allowed to polymerize in the thermoplastic matrix forming a loose network.

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Commercial Polymer Blends

Examples of such systems are silicone semiIPNs in thermoplastics that have been recently commercialized (Rimplast®, Petrarch, div. of Hüls) [Anonym., 1983]. The silicone semi-IPNs consist of mixing a hydride-containing silicone prepolymer and a vinyl functionalized silicone polymer into a thermoplastic matrix such as PA, PBT, thermoplastic polyurethane (TPU) or styrene-ethylene/butylenestyrene (S-EB-S) block copolymer elastomer. The two silicone prepolymers co-react in the thermoplastic matrix during melt extrusion and injection molding to form a partially crosslinked network within the thermoplastic matrix. The crosslinking reaction may be catalyzed by a small amount of suitable organometallic catalysts. The blends typically contain ca. 5-20% silicone. Injection molded or extruded parts are further heat-treated to complete the curing reactions. There is, of course, a significant level of phase separation. In the thermoplastic molding compounds such as glass-filled PA and PBT, addition of the silicone semi-IPN in small amounts (ca. 5-10%) is reported to reduce the mold shrinkage, improve mold release, and increase wear and friction resistance. Polyamide-silicone blends have already been discussed under PA blends section. Elastomeric silicone IPN with TPU and S-EB-S thermoplastic elastomer matrices have found some medical applications [Carew and Deisher, 1989]. The silicone contributes to the excellent release characteristics and to the bio-compatibility. Typical applications include medical tubing, catheters, implants, diaphragms, seals, gaskets, etc.

The commercial volume for silicone IPNs is, however, still very small. 15.14.3 Rubber Toughened Thermosets Most commercial rigid thermosets of high Tg exhibit brittleness and low tensile elongation because of the inherent nature of crosslinked network structures. Addition of rubbery dispersions into the thermoset matrix should improve the ductility and impact strength of the matrix by promoting the absorption of strain energy through multiple crazing and shear deformation sites in the matrix. However, dispersion of high molecular weight rubbers into the monomeric or prepolymer mixtures of the thermosetting resin matrix is usually difficult due to a viscosity mismatch and a lack of solubility or compatibility. Hence a low molecular weight, reactive elastomer is normally used for impact modification of thermosets. The low molecular weight of the rubbery prepolymer aids its easy dissolution or dispersability in the thermosetting resin. The reactive functionality couples the rubber covalently to the growing polymer network during the curing reaction. Hence the rubber toughened thermosets may also be considered as co-reacted thermosets and not true blends. Rubber toughened epoxy resins are the well known examples of impact modified thermosets utilizing reactive rubbery prepolymers. Epoxy resins can be toughened or flexibilized by any one of the following types of oligomeric reactive elastomers:

Oligomeric Elastomer

Reactive functionality

Polypropylene glycol diglycidyl ether [Riew, 1976] Polyaminoamides (condensation products of polyamines and ‘dimer’ acids) [McAdams, 1985] Liquid polysulfides [McAdams, 1985] Aliphatic polyesters [Drake, 1983] Liquids, butadiene-acrylonitrile copolymers (Hycar® ATBN or CTBN, B.F. Goodrich) [McGerry, 1968; Drake, 1975]

Epoxide end groups Amine groups Thiol groups Carboxyl and -OH groups Amine or carboxyl end groups

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These oligomeric reactive rubbers co-react with the epoxy resins through their corresponding reactive end groups thus incorporating rubbery blocks into the crosslinked epoxy network. For impact modification usually 10 wt% of the reactive rubber is used. For flexibilizing the thermoset higher levels (up to 50 %) are needed. The type and the amount of the oligomeric rubber used depends upon the degree of toughness and flexibility required in the product. The rubbery segments must phase separate after curing into discrete domains for effective impact modification without sacrificing the glass transition temperature or heat resistance of the matrix. Generally, 1-5 µm size rubber particles promote craze formation while shear deformation is promoted by rubber particles of < 0.5 µm. Systems possessing both small and large particles, i.e., bimodal distribution, provide maximum toughness [Riew, 1976]. Elastomer modified epoxy resins are primarily used in composites, structural adhesives and electronics applications.

15.15

Conclusions

Polymer blends will continue to play a significant role in the commercial development and application of plastic materials because of their desirable overall cost/performance balance and their ability to fill the performance gaps between existing single resins. In many cases, polymer blends have provided unique combination of properties and processability advantages not available in single resins. While further growth in commercial polymer blends will undoubtedly be strongly driven by specific market needs and applications, there are also several technological factors which can drive and sustain this growth. The following are some of the key technology factors that are likely to influence the future growth in commercial polymer blends: 1. Use of low cost recycled polymer feedstock of acceptable quality. 2. New compatibilization/reactive alloying strategies for immiscible polymer blends which fulfill the desired economic, safety and productiv-

ity targets. 3. Improved technology for obtaining reproducible and stable morphologies under commercial extrusion and molding conditions. 4. Better understanding of the correlations between rheology, morphology and mechanical properties to help optimize polymer blend design. 5. Development of efficient toughening technology (impact strength, fracture toughness and ductile/brittle transition temperature). 6. Improving the long-term service life and performance of polymer blends (thermal aging/ embrittlement resistance, creep and fatigue resistance, weatherability, etc.). 7. Developing cost-effective processing (compounding and post-fabrication) technology. 8. Improving the recyclability and reprocessability aspects of polymer blends, particularly with respect to the retention of properties after multiple processing histories, to increase the efficiency of regrind use. 9. Development of cost-effective technology for polymer blends that can continue to bridge the performance gaps between the commodity, engineering and specialty polymers.

15.16

References

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CHAPTER 16

ROLE OF POLYMER BLENDS’ TECHNOLOGY IN POLYMER RECYCLING

L. A. Utracki

National Research, Council Canada, Industrial Materials, Institute, Boucherville, QC, Canada

16.1

Abstract

There are three principal methods of plastic waste recycling: polymer or mechanical recycling, feedstock recycling, and energy recovery. The most desirable is the former, hence this Chapter will focus on the mechanical recycling. The process usually involves: segregation, cleaning or washing, shredding and compounding. During the melt compounding stage, the material is reformulated to fit performance criteria of the foreseen application. It also undergoes devolatilization, stabilization, compatibilization, alloying, filtering and pelletizing. The ABC of the plastics recycling technology is: alloying, blending, and compounding. Accordingly, the following text contains three principal parts describing: (i) the principles of polymer blending, (ii) use of polymer blend technology in plastics recycling, and (iii) polymers’ and blends’ recycling. The first part provides an outline of the principal aspects of the thermodynamics (miscibility and interfacial properties), compatibilization (either by addition of a compatibilizer, by reactive or mechanical compatibilization), and flow behavior of multicomponent, polymeric systems (including the fundamental principles of the morphology modeling). The fundamentals are followed by a summary of the polymer alloying, blending, and compounding methods, as well as it outlines the principles of morphology modeling. Because of the importance of re-stabilization of recycled polymers, these problems are also discussed. The following part focuses on the plastics recycling aspects, while the last discuses recycling of commingled post-consumer polymeric waste, that of polymer blends, it lists blends designed for repeated recycling. The chapter ends with an information on the biodegradable resins and blends.

16.2

Introduction

In 1996, the world production of thermoplastics reached 131 million ton/year. Europe generated 33.1 wt% of it, North America — 33.0%; South America — 3.6%; Asia L.A. Utracki (Ed.), Polymer Blends Handbook, 1117-1165. © 2003 Kluwer Academic Publishers. Printed in the Netherlands.

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— 28.5%, and Africa — 1.8%. As shown in Figure 16.1, by the year 2000, this amount is projected to increase by another 20 million ton/year. The annual consumption per capita varies from 180 kg in Belgium to few hundred grams in Pakistan. Similarly, the growth rate of plastics production varies from 3-4% in the developed countries to 13-19% in the developing ones. A simple algebra shows that for the global population of 6 billion and the annual average consumption of 150 kg/person the saturation level may require 900 million tons. This volume is sufficient to form a band all around the Equator, one kilometer wide and 22 meters tall (ca. as tall as a seven story building). All this mass must be somehow dissipated — only a small portion will degrade during the service and storage periods. The majority should be recovered, re-used, and recycled.

annum. For the plastics industry the average rate is predicted as AGR = 4.4 [BCC, 1997]. The increase is mainly due to rapid development of the plastics consumption in the developing countries. When this large captive market is taken into account, the leveling of the plastics production is expected to exceed one billion ton/year [Utracki, 1989, 1997]. At the end of its service-time most of this mass becomes a solid waste. Even if plastics constitute but a small fraction of the total waste (4-6%), such an amount represents large economic value. Many countries have adopted policy known as 3R’s: re-use, repair, and recycle [Banerji, 1997].

Figure 16.2. The world consumption of plastic will grow from 121 million tons in 1995 to 151 million tons in 2000, or by 4.4% per annum [BCC, 1997].

Figure 16.1. World production of man-made plastics, excluding rubbers and fibers, is expected to grow from 30 thousand tons in 1900 to 151 million tons in 2000, i.e., by a factor of 5033.

Recycling is considered to be either a reduction of waste, alleviation of disposal problems, or means for the recovery of material and/or energy. This chapter focuses on the material recovery as the most pertinent. 16.2.1

The world consumption of major polymers is shown in Figure 16.2. As the annual growth rate (AGR in %) indicates, it is expected that during the last five years of the 20th century, production of ABS will increase about 4 % per annum, that of LDPE (including LLDPE’s), HDPE, PS and PVC 5 %, PP about 7 %, and PET nearly 8% per

Sources of Plastics Waste

Plastics waste comes from two main sources: processing plants (ca. 10%) and post-use scrap (90%). The former is nearly totally reprocessed (i.e., • 96%) by the processors using the plastics (or mechanical) recycling methods. Recycling of the later material very much depends on the

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Role of Polymer Blends’ Technology in Polymer Recycling

geographical location. For example, in 1995 in Western Europe 70% went into landfill, 17% was burned for energy, 9% was mechanically recycled, 1% was thermolyzed and the rest was burned [Anonym., 1997]. By contrast, in India 91% was recycled by mechanical reprocessing and 9% goes into landfill [Mathur, 1997]. The following text will discuss mechanical recycling of mainly the post-use plastics waste. The single-component plastics and mixed plastics having well-defined composition and known history are best suited for the mechanical recycling. These recyclates can be used to substitute the virgin resins, i.e., they can be reprocessed to have the same functional properties as the comparable virgin materials. Structural heterogeneity’s and residual impurities, e.g., different pigments, fillers, flame retardants may be present even in carefully sorted and cleaned recyclates. From the point of view of economy, post-use plastic recycling makes sense only where an efficient system of collecting and sorting the waste is well designed and implemented. The main difficulties of recycling are related to their collection, segregation and decontamination. From the technical point of view, the variability of composition and the level of contamination present in the waste material pose the greatest problems. For these reasons, it is convenient to group the commingled plastics waste into four categories [Leidner, 1981]: ∗ The easiest to recycle is the in-house generated, processing scrap. At present, blending (up to

40 wt%) with virgin resins reuses virtually all these materials. ∗ To the second category belong clean, industrial scraps. These materials are collected in a full range of chemical homogeneity, from a single resin (e.g., off-spec) to commingled (e.g., PE and PVC from wire-&-cable plants). The former type is usually upgraded by ABC, or downcycled. The second category resins are alloyed either in the molten or solid state. To this category belong materials that enter the close-loop applications, e.g., bottle crates, trash bins, automotive parts (such as bumpers or battery cases), PET bottles, and from variety of other origins, including the industrial scrap. ∗ Waste of similar plastic materials from individual sector of the plastics applications, viz. packaging plastics comprising different polyolefins, such as polyethylenes (HDPE, LDPE, LLDPE), polypropylene (PP) and/or copolymers, PS and EPS from fast-food chains, PET and PET -copolymers from textile mills, rigid and plasticized PVC building waste (sidings and piping), etc. ∗ The most difficult is the post-consumer waste, PCW, e.g., from households, consisting of a variety of composition, viz. PO, PVC, PS, polycondensates (viz. PA, PC, PET), etc. As shown in Table 16.1, composition of PCW varies from one country to another. The plastics that belong to this category, first must be separated then cleaned. Alternatively, they may be processed, e.g., into a variety of plastics woods..

Table 16.1.

Composition of post-consumer thermoplastics waste in different regions or countries (in wt%)

Polymer

W. Europe

HDPE LDPE PP PS PVC PET Others

22 32 13 13 10 5 5

Germany 21 40 13 11 11 1 3

Italy

USA

19 36 18 8 5 12 2

27 30 11 16 4 11 1

Japan 16 26 18 21 8 5 6

India 12 20 31 5 31 1 -

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Table 16.2.

Uses of recyclable plastics and their content in the total amount of produced waste by sector

No.

Sector

1 2 3 4 5 6

Packaging Building Electrical/electronics Automotive Agriculture Others (e.g., household)

Use of plastics as wt% of total use 41 20 12 7 4 16

The plastics use in the five principal sectors (as wt% of the total consumption of plastics) is listed in Table 16.2. The relative magnitude of the plastics waste is also indicated as a percentage of the total waste generated by the sector. Even in the packaging industry that consumes 41 wt% of all plastics, these materials constitute less than one weight percent of the total waste. The largest contribution of plastics to the total waste is in the electrical and electronic industry that consumes 12 wt% of total plastics, generating nearly 13 wt% of recyclable plastics in their waste streams. This waste originates mainly in two sources: housings for the electronic instruments and wire and other parts insulation. 16.2.2

The Methods of Plastics Recycling

Over the years, several plastics’ disposal methods have been developed [Leidner, 1981]: 1. Polymer recycling, also known as mechanical recycling. The method is the most desirable, and it should be used whenever it is ecologically, commercially, and technically feasible. 2. Feeedstock recycling, also known as thermolysis of rubbers and plastics. The method is used to recycle soiled, mixed rubbers and plastics into monomers, feedstock gasses, oils, and condensates, from which new polymers can be synthesized. It should be used only when and where the first method cannot be carried out. 3. Energy recovery, also known as incineration of organic waste. This is the last recourse, to

Plastics waste as wt% of all wastes 0.87 0.33 12.70 7.00 0.11 7.88

be used when and where neither of the two previous methods is viable. 4. Use of soiled, commingled plastics waste in blast furnace as carbon donors. 5. Wasteful and polluting burning. Recyclability should not be an afterthought, but one of the performance criteria of plastics products. When designing a new part, the material selection should consider the total cost, i.e., that of resins, of forming, assembling, decorating, customer satisfaction, esthetics, service life-span, ease of disposal, and recycling. Thus recyclability, as well as the recommended strategy for recycling, should be considered at the beginning of the process, not at the end. Several strategies of the plastics recycling have been developed. From the point of view of the recycled material performance, these can be divided into three categories: (1) downcycling, (2) recovery of properties, and (3) upgrading. In many cases the strategy depends more on the material properties than on the recycler’s desires. Thus, for example, General Electric Company proposed that its engineering resins should be recycled for the sequential use in progressively less demanding applications — the strategy that became known as downcycling. The company demonstrated that PPE, PC, PEST resins and their blends recovered from automobiles after 10 years of weathering, can successfully be reused as valuable engineering materials, e.g., for household goods, toys, etc.

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Role of Polymer Blends’ Technology in Polymer Recycling

Recovery of properties may be based on specific polymers and their blends (viz. syndiotactic polystyrene, polyesteretherimide, some blends of polyphenyleneether) that have been shown to maintain high performance for at least 5-10 reextrusions. It has been also shown that properly stabilized resins (e.g., polyolefins) are able to maintain performance for at least five re-extrusions [Herbst et al., 1997]. The upgrading strategy comprises several elements of the polymer blends technology, viz. alloying (i.e., compatibilization and/or impact modification), blending to the desired morphology, and compounding with other additives (e.g., stabilizers and fillers) — the three ABC elements of plastics recycling. For example, to compensate for the hydrolytic decrease of polyamides or polyesters molecular weight, the reactive coupling (e.g., using bi-functional oxirane compounds) can be used to repair the molecular damage. The methods used to recycle plastics in Western Europe (in 1993, by sector) are listed in Table 16.3 [Radusch, 1997]. Thus, at this late date over 3/4 of European plastics waste were buried, with only 6% mechanically recycled. Note, that at the same time, 80% of plastics were recycled in India, 16% in South America, 12% in Japan, 10% in China and the same amount in the USA [Ratra, 1997]. The cost and benefit analysis for different plastics recycling strategies in Germany are presented in Table 16.4 [Bauermeister et al., 1994]. The

Table 16.3.

Plastics waste disposal in Western Europe — 1993

Sector Waste Packaging Building Electrical Automotive Agriculture TOTAL

mechanical recycling of plastics is clearly the most profitable. The difference between the total cost and that to recycler is due to special subsidy per ton of recycled resin, paid by the Duales System Deutschland (DSD). However, in other countries where the subsidy does not exist, the relative cost-to-benefit ration is expected to be similar. Synthetic polymers can conveniently be divided into three groups: commodity, engineering and specialty. Their relative annual production can be approximated by the ratio 100:10:1. Thus, it is logical that the main emphasis has been for recycling of the commodity resins (see Figure 16.3). It has been also estimated that polymer blends constitute about 36% of all thermoplastic resins on the market. As it will be evident from the following text, blending technology plays a dominant role in the plastics recycling. This technology offers numerous advantages for the resin manufacturer as well as for the designing engineer: 1. Better processability, product uniformity, low scrap thus process economy. 2. Plant flexibility – quick change of formulation. 3. Product tailored to customer needs. 4. Blending offers a unique method for economic development of materials having a wide range of performance characteristics. 5. The technology can be adapted for filler and reinforcement incorporation.

Recycling by method (kton) Polymer recycling

Burning for energy

609 15 3 39 70 915 (5.6%)

246 11 21 26 13 2,425 (15.0%)

Incineration 60 1 8 4 9 592 (3.6%)

Landfill 1620 753 486 773 635 12,279 (75.7%)

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Table 16.4.

L. A. Utracki

Cost-to-benefit analysis of plastics recycling in Germany (DM/ton) [Bauermeister et al., 1994]

Operation

Neat plastics recycling

Mixed plastics recycling

Hydrogenation

Collection Separation Treatment Process Total cost Recycler cost Selling $ Cost-selling $ Comment

700 800±100 550±150 480±120 2530 ±550 925±575 1015±485 0±700 BEST

700 700 partial 850±350 2250±450 490 350±50 -500±400 OK

900 700 500 500 2660±300 650±150 240±30 -820±130 OK

Neat Mixed plastics plastics thermolysis thermolysis 900 700 500 335 2435±300 500 290±30 -510±100 OK

6. Blending is essential for incorporating recycles with virgin resins. 7. Blending technology can: ∗ up-grade recyclates’ performance, ∗ be used for reactive molecular repairs, ∗ can extend the use of high performance polymers, etc. 8. Blends are inherently recyclable by virtue of morphology regeneration. 9. Most modern blends (see patents) are formulated for multiple recyclability without a loss of the performance characteristics. 10. Blends can be processed using standard equipment. By definition, polymer blends are mixtures of at least two macromolecular species, polymers and/or copolymers. For this reason any mixture of polymers should be treated as a blend. Most blends of chemically different polymers are immiscible. However, it has been documented that blends of chemically identical polymers (viz. LLDPE’s) can also be immiscible when their macromolecules sufficiently differ in chain configuration. The polymer/polymer miscibility is limited to: mixtures of homologous fractions, rare cases of polymeric pairs with strong specific interactions, or systems with the so-called “miscibility

900 700 200 300 2085±250 650 70±20 -415±70 POOR

Gassification

Incineration

Burning

Landfill

900 700 450 450 2500±300 650±150 120±20 -750±150 OK

300-450 500±200 850±150 675±175 75 -300±50 POOR

300-450 350 750±50 500 18±12 -370±50 POOR

375 450±250 790±210 500 -200 WORST

window”. The miscible blends have been studied not to produce single-pase commercial systems, but rather to develop proper strategy of compatibilization. For more details on the thermodynamics and miscibility see Chapter 2. Thermodynamics of Polymer Blends and Appendix 2. Miscible Polymer Blends. Evidently, to recycle mixtures of chemically different polymers they should be treated as blends, hence compatibilized, dispersed to the suitable morphology and stabilized. However, it is noteworthy that also mixtures of chemically identical polymers that were subjected to degradation, should be considered blends. Even if these would be miscible, adequate homogenization is required to ascertain the desired performance. To obtain well performing material out of a mixture of immiscible polymers two principal operations must be performed: compatibilization and compounding. Fundamental aspects of these will be discussed in Part 16.3 The principles of polymer blending. However, interested reader is advised to search detailed information is specific chapters of this Handbook, viz. on compatibilization in Chapters 4 and 5, on flow, morphology and compounding in Chapters 7, 8 and 9, respectively. The immiscible blends must be compatibilized. However, in special cases the compatibilization

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Role of Polymer Blends’ Technology in Polymer Recycling

may not be required, especially if the polymers are chemically similar. Thus, when either the mixture contains only a small amount of the second polymer, e.g., φ2 < 0.1, or the composition is near the phase inversion, so the phase cocontinuity can be generated, the compatibilization is not as important. For the other compositions of immiscible polymer blends compatibilization is a necessity. The aim of compounding is to ascertain optimum morphology (for a specific application). To accomplish this task, nowadays compounding and reactive compatibilization are carried out in compounders as well designed, integrated steps of a single process. Flow field that exists during compounding affects the size and shape of individual domains, imposes a gradient of concentration and/or morphology across the part, and it may affect the miscibility. Before discussing the fundamentals of blending, it seems appropriate to complete this Introduction with brief discussion of three, widely different areas of plastics recycling: Commercial recycled resins, Plastics recycling in the automotive industry, and Polymeric post-consumer waste recycling. 16.2.3

Plastics Recycling — Examples

16.2.3.1 Commercial Recycled Resins Recycled resins (all commodity and some engineering ones) are commercially available. A comparative list of prices of the virgin and recycled resins is shown in Table 16.5. The prices of recycled plastics are quite competitive. Evidently, while the quoted prices of commodity resins are reasonably close to these of virgin materials, the situation vis-à-vis condensation polymer is atrocious. The only way to make profit on the latter materials is to upgrade them by the recyclers ABC and supply compounds with well-defined and reproducible set of properties to a specific market niche. As it will be discussed later, the additiontype polymers can be reused after restabilization and compounding, whereas the condensation-type ones may require “molecular repair” to bring up

the molecular weight to the level that existed prior to the hydrolytic degradation. Table 16.5.

Resin ABS PC HDPE LLDPE LDPE PET-bottle PP PS-crystal PS-high heat PVC-flakes

List price of selected resins in US $/ton [Plast. News, Dec. 1997] Virgin

Recycled

1958-2002 3256-3586 836-990 924-1032 1012-1210 1188-1254 660-792 1012-1034 1034-1056 762-770

792-902 1122-1232 792-902 638-748 594-704 814-924 550-638 770-880 792-902 286-462

Several recycled trade-name resins that are also on the market (see Table 16.6). As the data in Figure 16.3 demonstrate, their performance is slightly (by about 10%) reduced [Huckestein, 1995]. In spite of the price advantage, the recycled resins still constitute only a small fraction of the thermoplastics’ consumption — it varies from 16.5% for PET, 5.0% for HDPE, 2.2% for LDPE, and to a fraction of one percent of PS, PVC, and other resins. Numerous products are commercially available from the recyclates: ∗ injection molded from recycled PC multipallet systems (GE) ∗ agricultural films comprising 30 wt% recycled PE (Dow) ∗ fleece jackets, sweaters and other knitwear from recycled PVC/PET (Reprise Plast.) ∗ PP automotive battery cases; PP pallets, PP carpets ∗ PP/EPDM automotive bumpers (since the early 1990’s, bumper recycling has been commercial in Europe) ∗ HDPE/LDPE low-pressure pipes, bottle crates ∗ Styropor – EPS for concrete and other industrial applications (BASF)

1124

Table 16.6.

L. A. Utracki

Available, recycled trade-name resins [Huckestein, 1995]

Trade name

Composition

Manufacturer

Aurum™ Bayblend™ Dylite™ EPS Naxell™ Orgablend™ Retain™ Ultramid™RC

Fully recyclable PI Blends of ABS & recycled EPS with 25% recycles Recycled PC Recycled PA/PP alloys Recycled PS/PO blends Recycled PA’s + GF

Mitsui Toatsu PC Bayer Arco Chem. Co. MRC Polymers Elf Atochem Dow Plastics BASF

∗ Commingled plastics pallets, plastic wood, noise protection walls, etc. 16.2.3.2 Automotive Industry

Figure 16.3. Comparison of properties of selected virgin and recycled resins from BASF.

∗ ∗ ∗ ∗ ∗ ∗

ABS electrical and electronic housings, automotive parts Remo-Tec are moldings of recycled PU foams to desired shape (Hennecke) PVC bottles, roofing sheets, blister packagings, window profiles PBT/PC automotive bumpers PA containers, hub cups, electrical and electronic housings Recycled PET is used for an inner layer of new, 5-layer Coca-Cola bottles. The outer layers are virgin PET, separated from the inner by two thin layers of PA-m XD6

The automotive industry started to use plastics in 1946. Since then, its content steadily increases (see Figure 16.4). Today, the plastics for automotive applications are dominated by polymer alloys, blends, and composites. The reason for it is the need for automotive parts to show a wide range of performance characteristics that are virtually impossible to meet using a single-component polymer. Thus, for example, in Saturn, front fenders and rear quarter panels are from PA/PPE, door outer skins are from PC/ABS, the bumper fascias are TPO, etc.

Figure 16.4. Increase of polymeric content in North American family cars: 1978-2000.

Role of Polymer Blends’ Technology in Polymer Recycling

The automotive industry is strongly committed to recycling. For example, the new cars are designed for easy dismantling of the plastic body parts for a closed-cycle recycling. The industry favors: (1) recyclable resins, (2) reduction of the variety of resins, and (3) resins that can easily be compatibilized and recycled. Subaru was the first manufacturer to offer fully recyclable PP/TPE bumpers for all car models. To recycle, the bumpers are chopped into sections and placed between rollers rotating at different speeds. This removes 99% of paints and coatings. The sections are then re-stabilized and granulated, ready for re-use. Similarly, Nissan aims for 100% bumper-to-bumper recycling of PPE/PA. Fiat has a policy of 100% recycling its TPO bumpers into automotive parts that do not demand similar level of performance, viz. air ducts, liners, etc. The PP battery cases are being recycled as fender liners. Ford adopted a different strategy. The company recycles the PC/PBT Xenoy™ alloy bumpers (after removing 99.7% of paint) into tail light assemblies. Alternatively, by adding 25 wt% of the recycled material to virgin resin it uses it to mold bumpers, guide brackets, etc. In 1996, the company used 16.5 kton, or 17%, of postconsumer recycled resin. The recyclates content is to grow steadily to 50% in the year 2000 [Grande, 1996]. 16.2.3.3 Polymeric Post-consumer Waste (PCW) Recycling The most difficult for recycling is the commingled post-consumer waste (PCW), especially from the municipal waste stream. Complete PCW recycling plants are available from several equipment manufacturers at a cost of about $3-5 million for 1 ton/hr capacity. However, most reprocessors tend to assemble the reprocessing lines from individual units, frequently modifying the equipment to better suite individual conditions. The number of manufacturers for separate machines, viz. sorters, separators, grinders, washers, shredders, aggregators, classifiers, dryers, extruders, and others, runs into thousands.

1125

The method of recycling depends on the type of the waste stream and the degree of cleanliness. Relatively clean PCW can be processed by methods used for industrial scrap. Numerous techniques have been developed for reclaiming the most readily identifiable constituents of such mixed waste streams as soft drink bottles (PET) and milk containers (HDPE). After separation of them, the remaining part of PCW (known as tailings) is made of PE’s (50-75 wt%), styrenics (25 wt%) and other plastics. These can be cleaned, comminuted, melted, upgraded by blending, and melt-formed. They can also be reinforced, foamed, or stretch-formed into a variety of products for specific applications. Alternatively, PCW tailings can be used to produce plastic wood or plastic lumber, known to have excellent long-term outdoor performance. The experience indicates that it is easier to recycle unfilled than filled polymers. In all cases re-stabilization of the resin is required. Some polymers (e.g., PI, PPS, POM, and LCP) are less susceptible to degradation than others. By contrast, the condensation polymers, viz. PA, and PEST, are difficult to reprocess without loss of performance. In most cases, the engineering resins and polymer blends require upgrading. The mechanical recycling of PCW has been accomplished either in the molten or solid state. Melt recycling of washed, dried and granulated commingled polymers, in essence, follows the same procedure as that described for compounding of polymer blends (see Chapter 7. Compounding Polymer Blends). However, these resins must be carefully re-stabilized using stabilizer systems that will not be damaged by the presence of other additives. Some commingled polymers may also require addition of odors suppressants. When compounding polymers at high stresses, the free radical generated by the mechano-chemical means may provide adequate compatibilization. For example, an extruder was developed with the residence time, τ > 7 seconds [Patfoort, 1976]. The machine generates such high shear stresses that extruded, immiscible polymers (e.g., recycled PS and PE) show good mechanical properties. Several other machines of this type have been

1126

L. A. Utracki

developed since. They all involve intensive mechanical mixing, that produces enough chain scission. Recombination of the free radicals in situ generates sufficient concentration of copolymer, to compatibilize the system. The generated under high stress, non-equilibrium morphology is then locked by quenching. Best performance has been observed for systems with co-continuous morphology. Uni- or bi-axial stretching in solid state (using methods similar to these of the steel industry) leads to significant enhancement of properties, especially when the PCW contains mostly PO or PEST. In another description of recycling PCW into plastics with good and reproducible set of performance characteristics, first the waste plastics were shredded, metals and heavy elements were removed, then it was washed, paper was removed, dried and grounded. The composition was verified by FTIR, the additives and stabilizers were added. The material was mechanically compatibilized by compounding it in a high speed mechanical shearing machine, with the screw speed 30x higher than that of a TSE. The mass was kinetically heated to 230°C, in-situ generating the compatibilizing copolymers. The compound was fed to a melt extruder, degassed, filtered, and pelletized. A blend comprising 63.7 wt% HDPE, 8.75 wt% PVC, 5 wt% PET, 8,75 wt% PP, 8.75 wt% PS and 5 wt% ABS had higher than HDPE tensile strength at yield (but lower at break), higher Young modulus, the same flexural modulus, and lower elongation at break. The ABS/PVC scrap showed similar performance to PP [La Mantia, 1996; La Mantia et al., 1996; Dubrulle D’Orhcel, 1996]. There are several methods of PCW recycling in the solid state. These are characterized by intensive mechanical action at low temperature, generating free radicals that lead to co-reaction between the macromolecules. Created morphology can be stabilized either by the chemical or physical means, viz. crystallization. 1. Ball-milling in a shaker ball-mill at acceleration 12.3 g, a frequency 29 Hz and T = -150°C. The final powder had the particle size of about 2 mm. Placed under vacuum, it was consolidated under pressure at a temperature

between 3 and 100°C higher than the melting point, Tm. The solid state alloying resulted in homogenous materials, with high hardness and tensile strength [Shaw, 1993]. 2. Commingled resins’ performance was upgraded without going through the melting and compatibilization stage by solid-state sheet rolling. 3. The solid-state shear extrusion (SSSE) utilizes Berstorff co-rotating twin-screw extruder. The first zone is slightly heated, while the kneading zone was modified for intensive cooling to 15-60°C. Within this zone, the particles are repeatedly broken and reformed by impact, compression, attrition, and shear. The process changes PCW into powder with the particle size, 20 ” d ” 2000 µm. The powder was used directly for injection molding. Thus pulverized PO’s (viz. HDPE/LLDPE, HDPE/PP, etc.) showed about 20% improved impact strength, and about 10% decreased elongation at break in comparison to melt blended systems. The moldings had good surface appearance and did not delaminate upon breaking [Khait, 1994; 1995; 1996]. The process can easily be adopted for incorporation of upgrading additives: compatibilizers, impact modifiers, stabilizers, pigment, dyes, etc.

16.3

The Principles of Polymer Blending

16.3.1

Thermodynamics

Miscibility of polymer blends is usually presented as an isobaric phase diagram with the lower critical solution temperature, Tc = LCST (see Figure 16.5). The binodals indicate limits of miscibility, while the spinodals the boundary between the meta-stable and spinodal regions. Thus, polymer blends are usually miscible at lower temperature, outside the binodal and are immiscible at higher temperature inside the spinodal region. As discussed in Chapter 2 of this Handbook, there is a significant difference in the rate of phase separation and the generated morphology when a single-phase blend is “quenched” into either the

Role of Polymer Blends’ Technology in Polymer Recycling

meta-stable or spinodal regions. The mechanism that take place within the metastable region is called “nucleation and growth (NG)” and as the name indicates it resembles crystallization — NG leads to polydispersed drops. Since phase separation within the spinodal region generates phase co-continuity (by the mechanism of “spinodal decomposition, SD),” these blends may not need compatibilization. Furthermore, the SD blends most efficiently combine the advantageous properties of both polymers, e.g., high modulus of one component with high elongation at break of another.

the interphase in binary blends is characterized by low entanglement density, low viscosity and poor interphasial adhesion. These interphase drawbacks can be alleviated by compatibilization. The interphase thickness depends on the miscibility of the polymeric component as well as on the compatibilization. For uncompatibilized binary, strongly immiscible systems, the interphase thickness ∆l = 2 nm. The thickest interphase has been observed for reactively compatibilized polymer alloys: ∆l = 65 nm. For most blends, the interphase thickness is in between these two limits. The importance on the interphase can be appreciated noting that its volume will be the same as that of the dispersed phase when the drop diameter (without interphase) is about 500 nm. It is noteworthy that in most commercial polymer alloys the drop diameter is about five times smaller, making the importance of the interphase much greater. 16.3.2

Figure 16.5. Binary phase diagram at constant pressure, with the lower critical solution temperature,Tc = LCST. The solid and broken lines indicate binodal and spinodal curves, respectively. The single phase, two meta-stable regions, and a spinodal region are shown. Majority of polymer blends (whose miscibility depends on specific interactions) shows this type of behavior.

For proper understanding of the immiscible polymer blends it is important to take into account the interphase. In binary blends, the interphase thickness, is inversely proportional to the interfacial tension coefficient, thus, poorer the miscibility, larger the interfacial tension coefficient and smaller the interphase thickness. Owing to the thermodynamic forces the polymeric chain-ends concentrate at the interface and the low molecular weight components diffuse to it as well. Thus,

1127

Compatibilization of Polymers

Only in rare cases (mixtures with φ2 < 0.1, or compositions near the phase inversion) the compatibilization may not be necessary. However, even in these cases compatibilization has beneficial effects on blends’ performance. All the other immiscible blends should be compatibilized. The compatibilization must: 1. Reduce the interfacial tension to facilitate dispersion, 2. Stabilize the generated morphology against modification during the subsequent processing steps, and 3. Enhance adhesion between the polymers’ domains, facilitating the stress transfer, hence improving the mechanical properties of the product. The methods of compatibilization are discussed in details in Chapter 4. Interphase and compatibilization by addition of a compatibilizer, and in Chapter 5. Reactive Compatibilization of this Handbook. The efficiency of added compatibilizer is determined by its preferential location at the interface. At the thermodynamic equilibrium, this is accompanied by its dissolution in all phases. Increasing the compatibilizer concentration and/or its molecular weight, first micelles then meso-phases may be formed. Compatibilizer must be designed by

1128

L. A. Utracki

taking the thermodynamic and kinetic parameters into account. Three compatibilization methods are recognized: 1. Addition of a compatibilizer: either a small amount (0.5 to 4 wt%) of precisely tailored copolymer, 15 to 35 wt% of a multi-purpose compatibilizer-cum-impact modifier, or a cosolvent (viz. Phenoxy™ in PBT/PMMA blends). It is useful to distinguish copolymers having inherent morphology, capable to impose specific structure (usually phase co-continuity!) on the compatibilized blend, e.g., SEBS in PP/PC. The inherent morphology is related to the regularity of the molecular structure, thus expected only for nearly monodispersed anionic block copolymers. The other copolymers have irregular morphology and, depending on composition and flow, lead to either disperse or co-continuous blend structures. 2. Reactive compatibilization, that generates the desired quantity of either block or graft copolymers that form chemical bonds across the interface. The reactive compatibilization may take place in the molten or solid state. 3. Physical compatibilization that generates fine, non-equilibrium morphology, and locks it by nucleated crystallization. Th