Polyolefin Blends

POLYOLEFIN BLENDS

POLYOLEFIN BLENDS Edited by Domasius Nwabunma 3M Company

Thein Kyu University of Akron

WILEY-INTERSCIENCE A JOHN WILEY & SONS, INC., PUBLICATION

Cover credit: (Third image on top right side) Reprinted from European Polymer Journal, vol. 40, Smit, G. Radonjic and D. Hlavata. Phase morphology of iPP/aPS/SEP blends, page 1439, 2004. With permission from Elsevier. Copyright ß 2008 by John Wiley & Sons, Inc. All rights reserved Published by John Wiley & Sons, Inc., Hoboken, New Jersey Published simultaneously in Canada No part of this publication may be reproduced, stored in a retrieval system, or transmitted in any form or by any means, electronic, mechanical, photocopying, recording, scanning, or otherwise, except as permitted under Section 107 or 108 of the 1976 United States Copyright Act, without either the prior written permission of the Publisher, or authorization through payment of the appropriate per-copy fee to the Copyright Clearance Center, Inc., 222 Rosewood Drive, Danvers, MA 01923, (978) 750-8400, fax (978) 750-4470, or on the web at www.copyright.com. Requests to the Publisher for permission should be addressed to the Permissions Department, John Wiley & Sons, Inc., 111 River Street, Hoboken, NJ 07030, (201) 748-6011, fax (201) 748-6008, or online at http://www.wiley.com/go/permission. Limit of Liability/Disclaimer of Warranty: While the publisher and author have used their best efforts in preparing this book, they make no representations or warranties with respect to the accuracy or completeness of the contents of this book and specifically disclaim any implied warranties of merchantability or fitness for a particular purpose. No warranty may be created or extended by sales representatives or written sales materials. The advice and strategies contained herein may not be suitable for your situation. You should consult with a professional where appropriate. Neither the publisher nor author shall be liable for any loss of profit or any other commercial damages, including but not limited to special, incidental, consequential, or other damages. For general information on our other products and services or for technical support, please contact our Customer Care Department within the United States at (800) 762-2974, outside the United States at (317) 572-3993 or fax (317) 572-4002. Wiley also publishes its books in a variety of electronic formats. Some content that appears in print may not be available in electronic formats. For more information about Wiley products, visit our web site at www.wiley.com. Wiley Bicentennial Logo: Richard J. Pacifico Library of Congress Cataloging-in-Publication Data: Nwabunma, Domasius. Polyolefin blends / edited by Domasius Nwabunma, Thein Kyu. p. cm. Includes index. ISBN 978-0-471-79058-7 (cloth) 1. Polyolefins. I. Kyu, Thein, 1948- II. Title. TP1180.P67N928 2007 668.40 234–dc22 2007021318 Printed in the United States of America 10 9 8 7 6 5 4 3 2 1

Contents

Preface

xv

Contributors

Part I

Introduction

1. Overview of Polyolefin Blends 1.1 Introduction 1.2 Olefinic Monomers 1.3 Polyolefin Homopolymers, Copolymers, and Terpolymers 1.4 Polyolefin Blends 1.5 Trends in Polyolefin Blends Nomenclature References

2. Miscibility and Characteristics of Polyolefin Blends 2.1 Introduction 2.1.1 Polyolefins 2.1.2 Blends 2.2 Polymer Blend Miscibility 2.3 Interfaces in Liquid and Polymer Mixtures 2.4 Polyolefin–Polyolefin Blends 2.4.1 Blends between Polyethylenes 2.4.2 Blends between Isotactic Polypropylene and Ethylene Propylene Copolymers 2.4.3 Blends between iPP and High Comonomer Concentration Polyethylene Copolymers 2.4.4 Blends between iPP and PB1 2.5 Binary Immiscible Blends 2.5.1 Polyolefin–Polystyrene Blends 2.5.2 Polyolefin–Polyamide Blends 2.6 Ternary Blends of Polyolefins with Other Polymers and Compatibilizing Agents 2.6.1 Surfactants and Compatibilizing Agents 2.6.2 Polyolefin–Polystyrene Blends with Compatibilizing Agents

xvii

1 3 3 4 5 7 13 16 18 27 27 27 29 30 33 36 36 38 39 40 42 43 43 44 44 45

v

vi

Contents 2.6.3 Polyolefin–Polyamide Blends with Compatibilizing Agents 2.7 Conclusions Nomenclature References

Part II

Polyolefin/Polyolefin Blends

3. Miscibility, Morphology, and Properties of Polyethylene Blends 3.1 Introduction 3.2 Structure and Properties of Polyethylenes 3.3 Applications of Polyethylene Blends 3.4 Molar Mass and Branching Distributions 3.5 Crystallization, Melting, and Branching of Polyethylenes 3.6 Miscibility and Crystallization 3.7 Theoretical Prediction of Miscibility 3.8 Rheology of Melted Polyethylene Blends 3.9 Mechanical Properties of Polyethylene Blends 3.10 Additives 3.11 Conclusions Nomenclature References

4. Miscibility and Crystallization Behavior in Binary Polyethylene Blends 4.1 4.2

Introduction Miscibility 4.2.1 Linear and Short Branched Polyethylene Blends 4.2.2 Blends of Linear and Long Branched Polyethylenes 4.2.3 Blends of Short and Long Branched Polyethylenes 4.3 Crystallization Behaviors 4.3.1 Blends of Linear and Short Branched Polyethylenes 4.3.2 Blends of Linear and Long Branched Polyethylenes 4.3.3 Blends of Short and Long Branched Polyethylenes 4.4 Conclusions Nomenclature References

5. Microscopically Viewed Structural Characteristics of Polyethylene Blends Between Deuterated and Hydrogenated Species: Cocrystallization and Phase Separation 5.1 5.2 5.3

Introduction Cocrystallization and Phase Separation of PE Blends Aggregation Structure of Chains in Lamella

48 50 50 51

57 59 59 64 65 66 68 72 74 76 77 80 80 81 82

84 84 86 86 88 89 90 90 92 93 93 94 94

97 97 98 101

Contents 5.4 Crystallization Behavior of D/H Blend Samples 5.4.1 Crystallization in the Cooling Process from the Melt 5.4.2 Isothermal Crystallization Process 5.4.3 Blending Effect on Crystallization Rate 5.5 Mixing Behavior of D and H Components 5.6 Conclusions Acknowledgments Nomenclature References

6. Thermal and Structural Characterization of Binary and Ternary Blends Based on Isotactic Polypropylene, Isotactic Poly (1-Butene) and Hydrogenated Oligo (Cyclopentadiene) 6.1 Introduction 6.2 Binary Blends 6.2.1 Blend Preparation 6.2.2 Glass Transition Temperature 6.2.3 Morphology and Spherulite Growth Rate 6.2.4 Isothermal Bulk Crystallization Kinetics 6.2.5 Temperature Dependence of the Spherulite Growth Rate and the Overall Kinetic Rate Constant 6.2.6 Melting Behavior 6.2.7 Polymorphism and Phase Transformation of Poly (1-Butene)/ Hydrogenated Oligo (Cyclopentadiene) 6.2.8 Supermolecular Structure of Isotactic Polypropylene/ Hydrogenated Oligo (Cyclopentadiene) Blends 6.3 Ternary Blends 6.3.1 Blends Preparation 6.3.2 Morphology and Spherulite Growth Rate 6.3.3 Glass Transition Temperature 6.3.4 Nonisothermal Crystallization and Melting Behavior 6.3.5 Isothermal Bulk Crystallization Kinetics of Isotactic Polypropylene Component 6.3.6 Melting Behavior of the Isotactic Polypropylene Component 6.3.7 Supermolecular Structure 6.4 Conclusions Nomenclature References

7. Morphological Phase Diagrams of Blends of Polypropylene Isomers with Poly(Ethylene–Octene) Copolymer 7.1 Introduction 7.2 Blends of sPP/POE 7.2.1 Thermal Characterization and Morphological Phase Diagrams: Undulated Lamella, Sheaf, and Spherulite 7.2.2 Growth of Single Crystals: Length, Width, and Periodicity

vii 105 106 108 113 114 117 118 118 119

121 121 123 123 123 124 125 126 131 133 136 141 141 141 143 143 145 146 147 153 154 155

157 157 159 161 165

viii

Contents 7.2.3 Phase Field Modeling for a Single-Component System: Sectorization and Ripple Formation in sPP 7.3 Blends of iPP/POE 7.3.1 Morphology Development in Relation to Phase Diagrams 7.3.2 Sectorization in iPP/POE Blends 7.3.3 Crystal Growth Dynamics in Binary Blends of iPP and aPP 7.4 Blends of ePP/POE 7.4.1 Characterization of Neat Elastomeric Polypropylene 7.4.2 Melting Transitions and Morphology Phase Diagrams of ePP/POE Blends 7.5 Conclusions Nomenclature References

8. Structure, Morphology, and Mechanical Properties of Polyolefin-Based Elastomers 8.1 Introduction 8.2 Thermoplastic Polyolefin Elastomers 8.2.1 Reactor Blends of PP, PE, and EPR: Impact Copolymer PP 8.2.2 Postreactor Blends of PP–EPR and ICP–EPR 8.3 Thermoplastic Vulcanized Elastomers 8.3.1 Dynamic Vulcanization and Morphology 8.3.2 Origin of Rubber Elasticity 8.3.3 Several Factors that Influence Mechanical Properties 8.4 Polyolefin Copolymers, Blends, and Composites 8.4.1 Polyolefin Copolymers 8.4.2 Blends and Composites 8.5 Conclusion Acknowledgment Nomenclature References

9. Morphology and Mechanical Properties in Blends of Polypropylene and Polyolefin-Based Copolymers 9.1 Introduction 9.2 Morphology and Dynamic Mechanical Properties 9.2.1 Blends with Polyethylene or Poly(butene-1) 9.2.2 Blends with Ethylene–a-Olefin Copolymers 9.2.3 Ethylene–Isotactic Propylene Copolymers 9.3 Tensile and Rheo-Optical Properties 9.3.1 Principles for Rheo-Optical Characterization 9.3.2 Blends with Ethylene–a-Olefin Copolymers 9.3.3 Blends with Novel Ethylene–Isotactic Propylene Copolymers 9.4 Solidification Process and Final Morphology 9.4.1 Morphology Formation During Crystallization 9.4.2 Structure and Properties of Injection-Molded Products

172 177 177 181 182 187 188 188 193 195 196

198 198 199 199 201 206 206 208 211 214 214 219 221 222 222 222

224 224 225 225 226 236 241 241 242 247 250 250 257

Contents 9.5 Conclusions Nomenclature References

10. Functionalization of Olefinic Polymer and Copolymer Blends in the Melt 10.1 10.2

Introduction Scope of Review 10.2.1 Free-Radical Grafting of Unsaturated Monomers to PO Chains 10.2.2 The Use of Monomers and Initiators 10.3 Functionalization of PP/PE Blends 10.3.1 Effect of Reacting Blend Formulation on Grafting Efficiency and Rheological and High Elastic Properties of Melt of Functionalized PP/PE Blends 10.3.2 Structure and Mechanical Properties of Functionalized PP/PE Blends 10.4 Functionalization of PP/EPR Blends 10.5 Functionalization Features of Blends: PE/EPR, PP(PE)/EOC, and PP(PE)/Styrene Polymer 10.6 Use of Functionalized Polyolefin Blends 10.7 Conclusion Nomenclature References

11. Deformation Behavior of b-Crystalline Phase Polypropylene and Its Rubber-Modified Blends 11.1 11.2

Introduction Deformation Characteristics 11.2.1 Static Tensile Behavior 11.2.2 Strain-Induced b ! a Phase Transition 11.2.3 Impact Behavior 11.3 Fracture Toughness 11.3.1 General Aspects 11.3.2 Mode I LEFM Approach 11.3.3 Impact Fracture Toughness 11.3.4 Essential Work of Fracture 11.4 Conclusions Nomenclature References

12. Multiphase Polypropylene Copolymer Blends 12.1

Introduction 12.1.1 Commercial Production 12.1.2 Morphology of Commercial Impact PP Copolymers

ix 264 265 265

269 269 270 270 275 284

284 291 295 297 299 300 301 302

305 305 309 310 314 323 330 330 331 339 341 347 347 348

351 351 352 352

x

Contents 12.2 12.3 12.4

Dispersive Mixing during Processing Molecular Structure of Impact PP Copolymers Coarsening in Multiphase PP Copolymer Systems 12.4.1 Background 12.4.2 Coarsening of High Impact Polypropylene 12.4.3 Coarsening of Model Blends 12.4.4 Interfacial Effects in Polypropylene Copolymer Systems 12.5 Conclusions Nomenclature References

13. Heterogeneous Materials Based on Polypropylene 13.1 13.2 13.3

Introduction The Interphase: Definition Magnitude Orders in the Interphase 13.3.1 The Dispersed Phase 13.3.2 The Matrix 13.3.3 The Interphase: Designing the Interface 13.4 Interfacial Modification of Heterogeneous Materials Based on Polypropylene 13.4.1 Composites: When the Dispersed Phase is Rigid 13.4.2 Blends: When the Dispersed Phase is Flexible 13.4.3 The Role of the Interfacial Modifiers from the Matrix Side 13.5 Interfacial Modifiers Based on Polypropylene 13.5.1 The Kinetic Approach: Basic Aspects 13.5.2 Chemical Modification of Polypropylenes by Grafting of Polar Monomers 13.6 Conclusions Acknowledgment Nomenclature References

14. Polypropylene/Ethylene–Propylene–Diene Terpolymer Blends 14.1 14.2

14.3

Introduction PP/EPDM Blends 14.2.1 Toughness and Crystallization Behaviors of PP/EPDM Blends 14.2.2 Compatibilization of PP/EPDM Blends 14.2.3 Ternary Blends and Composites from PP/EPDM Blends 14.2.4 Application of Radiation Dynamically Vulcanized PP/EPDM Blends (or Thermoplastic Vulcanizates (TPVs)) 14.3.1 Effect of Cross-linking on the Properties of PP/EPDM TPVs

357 360 360 360 365 368 370 373 376 377

379 379 380 380 382 383 383 385 387 387 388 397 397 398 407 408 408 408

411 411 412 412 414 416 417 419 420

Contents 14.3.2 Microstructure of PP/EPDM TPV 14.3.3 PP/EPDM/Ionomer TPVs 14.3.4 Mechanical and Rheological Properties 14.4 Applications of PP/EPDM Blends 14.5 Conclusions Acknowledgments Nomenclature References

15. Ethylene–Propylene–Diene Rubber/Natural Rubber Blends 15.1 15.2

Introduction Miscibility, Compatibility, and Thermodynamics of Polymer Blending 15.3 Blend Preparation 15.4 Covulcanization 15.5 Filler Distribution in NR/EPDM Blends 15.6 Morphology of NR/EPDM Blends 15.7 Compatibilization of NR/EPDM Blends 15.8 Mechanical and Viscoelastic Properties 15.8.1 Mechanical Properties 15.8.2 Dynamic Mechanical Properties 15.9 Rheological Properties 15.10 Thermal Properties 15.11 Electrical Properties 15.12 Aging Properties 15.12.1 Thermal Aging 15.12.2 Ozone Resistance 15.13 Transport Properties 15.14 Applications 15.15 Conclusions Nomenclature References

16. Phase Field Approach to Thermodynamics and Dynamics of Phase Separation and Crystallization of Polypropylene Isomers and Ethylene–Propylene–Diene Terpolymer Blends 16.1 16.2

16.3

Introduction Experimental Phase Diagrams 16.2.1 Cloud Point Phase Diagram of iPP/EPDM Blends 16.2.2 Cloud Point Phase Diagram of sPP/EPDM Blends Thermodynamic Free Energy Description of Crystalline Polymer Blends 16.3.1 Flory–Huggins Free Energy of Amorphous–Amorphous Blends

xi 423 425 428 436 437 438 438 439 441 441 442 443 444 447 448 450 452 452 456 458 460 461 462 462 463 465 466 469 469 470

473 473 475 475 476 478 478

xii

Contents 16.3.2

Extension of the FH Theory to Crystal–Amorphous Blends 16.3.3 Prediction of Phase Diagram Topologies 16.3.4 Comparison with Experimental Phase Diagrams of PP/EPDM Blends 16.4 Phase Field Modeling on Polymer Phase Transitions 16.4.1 Theory on Phase Separation Dynamics and Morphology Evolution 16.4.2 Dynamics of Crystal Growth in a Phase Separating System: iPP/EPDM Blends 16.4.3 Dynamics of Crystal Growth in a Phase Separating System: sPP/EPDM Blends 16.5 Conclusions Nomenclature References

Part III

479 481 484 486 486 488 491 496 496 497

Polyolefin/Nonpolyolefin Blends

499

17. Compatibilization and Crystallization of Blends of Polyolefins with a Semiflexible Liquid Crystalline Polymer

501

17.1

Blends of Polyolefins (High Density Polyethylene and Isotactic Polypropylene) with a Semiflexible Liquid Crystalline Polymer 17.1.1 Introduction 17.1.2 Blends of High Density Polyethylene (HDPE) with LCP 17.1.3 Blends of Isotactic Polypropylene with LCP 17.2 Crystallization Behavior of Blends of Polyolefins with a Semiflexible Liquid Crystalline Polymer 17.2.1 Isothermal and Nonisothermal Crystallization of Blends of Linear Low Density Polyethylene with a Semiflexible Liquid Crystalline Polymer 17.2.2 Crystallization Behavior of PE-g-LCP Copolymers 17.2.3 Effect of PP-g-LCP Compatibilizer on the Morphology and Crystallization of PP/LCP Blends 17.2.4 Isothermal Crystallization Kinetics of Compatibilized Blends of Polyolefins with a Semiflexible LCP 17.2.5 Crystallization and Morphology of Fibers Prepared from Compatibilized Blends of Polyethylene with a Liquid Crystalline Polymer 17.3 Conclusions Acknowledgment Nomenclature References

501 501 502 509 513

513 517 519 519

522 523 523 523 524

Contents

18. Functionalized Polyolefins and Aliphatic Polyamide Blends: Interphase Interactions, Rheology, and High Elastic Properties of Melts 18.1 18.2 18.3 18.4 18.5 18.6

Introduction Compounding and Interphase Phenomena in PA/PO Blends Rheological and High Elastic Properties of PA/PO Melts Morphology and Impact Strength of PA/g-PO Blends that Give Melts of High Viscosity and Strength New Applications of PA/g-PO Blends Conclusions

Nomenclature References 19. Plastic Deformation and Damage Mechanisms of Ternary PP/PA6/POE Polymer Blends 19.1 19.2

Introduction Materials Presentation and Experimental Methods 19.2.1 Materials Presentation 19.2.2 Morphological Study under SEM and TEM 19.2.3 Video-controlled Tensile System 19.3 Microstructure 19.4 General Mechanical Properties 19.4.1 Dynamic Mechanical Thermal Analysis 19.4.2 Toughness by Impact Loading and Yield Stress by Tension 19.5 Plastic Deformation under Uniaxial Tension 19.5.1 Definition of Volume Strain 19.5.2 True Axial Stress–Strain Relation 19.5.3 Volume Strain 19.5.4 Under Cyclic Tension 19.6 Mechanisms of Plastic Deformation and Damage 19.6.1 Damage Mechanisms in Polymer Blends 19.6.2 Influence of Damage Type 19.6.3 Damage-induced Shear Banding 19.6.4 Microscopic Observation under SEM 19.6.5 Microscopic Observation under TEM 19.6.6 Discussion 19.7 Conclusions Nomenclature References

20. Reactive Compatibilization of Binary and Ternary Blends Based on PE, PP, and PS 20.1 20.2

Introduction Friedel–Crafts Alkylation Reaction

xiii

527 527 528 534 543 550 550

551 551

556 556 558 558 559 559 561 566 566 567 569 569 571 573 575 579 579 580 581 582 589 589 596 597 597

600 600 601

xiv

Contents 20.3

Binary Blends 20.3.1 PE/PS Blends 20.3.2 PP/PS Blends 20.4 Ternary Blends: PE/PP/PS 20.5 Conclusions Nomenclature References

21. Polyolefin/Epoxy Resin Blends 21.1 21.2 21.3 21.4 21.5 21.6 21.7 21.8 21.9 21.10

Introduction Blend Preparation Miscibility Studies and Phase Diagrams Cure Kinetics Crystallization Behavior Morphology Dynamic Mechanical Properties Mechanical Properties Conductivity Studies Conclusion

Nomenclature References Index

603 603 611 615 620 620 621

623 623 624 625 629 631 637 644 647 649 656

658 659 663

Preface

Polyolefins are the most widely used commodity thermoplastics. They are of immense interest to polymer community because of their simple chemical structures and fascinating hierarchical structural organizations possible. To date, the field of polyolefins remains one of the most vibrant areas in polymer research. Polyolefin blends are a subset of polymer blends that emerged as a result of the need to meet application requirements not satisfied by synthesized neat polyolefins. In comparison to other subsets of polymer blends, polyolefin blends have distinct advantages of lower density, lower cost, processing ease, and good combination of chemical, physical, and mechanical properties. In the last several years, research and usage of polyolefin blends have increased due to new application opportunities (e.g., in medical and packaging) and the development of novel polyolefins. Although a sizable number of books on polyolefins and general polymer blends are available, only a few chapters address polyolefin blends. Currently, there is no single book that focuses exclusively on the fundamental aspects and applications of polyolefin blends. This is the primary source of motivation behind this book. The second motivation stems from the fact that new research trends in polyolefin blends such as in situ reactor blending and compatibilization/functionalization in the melt have emerged that need to be covered in a book format. This book is structured as follows: Chapter 1 serves as a guide to polyolefin blends introducing this important class of materials, why they are important, typical systems studied, issues of fundamental and applied interest, and current trends. The contributed chapters are divided into two main categories: polyolefin/polyolefin blends (Chapters 2–16) and polyolefin/nonpolyolefin blends (Chapters 17–21). Issues covered in these chapters include miscibility, phase behavior, functionalization, compatibilization, microstructure, crystallization, hierarchical morphology, and physical and mechanical properties. Most of the chapters are in the form of review articles. Some original articles are included to capture the latest development in polyolefin blends research. This book is intended to serve as a valuable reference for academic and industrial professionals performing research and development in the specific area of polyolefin blends or in the general area of polymer blends. Some review chapters include introductory materials to attract newcomers including senior undergraduate and graduate students and to serve as a reference book for professionals from other disciplines. Some knowledge of polymer chemistry, physics, and engineering, although not strictly essential, would be helpful to better appreciate the technical information of some chapters. Since this book is the first of its kind devoted solely to polyolefin blends, it is hoped that it will be sought after by a broader technical audience. xv

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Preface

The chapters in this book were contributed by highly reputed professionals from academia, industry, and government laboratories spanning several countries from various continents. All manuscripts were peer reviewed in accordance with the guidelines utilized elsewhere by top-rated polymer journals. The editors would like to thank all contributors for believing in the realization of this book and taking painstaking tasks of going through the processes of manuscript preparation, submission, review, revision, and seeking supporting documents. Finally, sincere thanks are extended to all reviewers for their invaluable help, which undoubtedly improved the quality of this book. DOMASIUS NWABUNMA THEIN KYU

Contributors

Susana Areso Capdepo´n, Department of Physics and Engineering of Polymers, Polymer Engineering Group, Institute of Science and Technology of Polymers, CSIC CL Juan de la Cierva 3, 28006 Madrid, Spain. [email protected]

Shu-Lin Bai, Centre for Advanced Composite Materials (CACM), Department of Mechanics and Engineering Science, School of Engineering, Peking University, 100871 Beijing, China. [email protected] Silvia E. Barbosa, Planta Piloto de Ingenierı´a Quı´mica, PLAPIQUI (UNS-CONICET), Camino La Carrindanga km. 7 (8000), Bahı´a Blanca, Argentina. [email protected]

Maurizio Canetti, C.N.R. Istituto per lo Studio delle Macromolecole, Via E. Bassini 15, I-20133 Milano, Italy. [email protected] Numa J. Capiati, Planta Piloto de Ingenierı´a Quı´mica, PLAPIQUI (UNS-CONICET), Camino La Carrindanga km. 7 (8000), Bahı´a Blanca, Argentina. [email protected]

Mo´nica F. Dı´az, Planta Piloto de Ingenierı´a Quı´mica, PLAPIQUI (UNS-CONICET), Camino La Carrindanga km. 7 (8000), Bahı´a Blanca, Argentina. [email protected]

Bejoy Francis, Department of Chemistry and Biochemistry, Laurentian University, 935 Ramsey Lake Road, Sudbury, Ontario, P3E 2C6, Canada. [email protected]

Soney C. George, Department of Basic Science, Amal Jyothi College of Engineering, Koovapally, Kottayam 686518, Kerala, India. [email protected]; [email protected] Christian G’Sell, Laboratoire de Physique des Mate´riaux, Ecole des Mines de Nancy, Parc de Saurupt, 54042 Nancy Cedex, France. [email protected]

Chang-Sik Ha, Department of Polymer Science and Engineering, Pusan National University, Busan 609-735, Korea. [email protected]

Jean-Marie Hiver, Laboratoire de Physique des Mate´riaux, Ecole des Mines de Nancy, Parc de Saurupt, 54042 Nancy Cedex, France. [email protected]

Benjamin S. Hsiao, Department of Chemistry, Stony Brook University, Stony Brook, NY 11794, USA. [email protected]

Boleslaw Jurkowski, Division of Plastic and Rubber Processing, Institute of Material Technology, Poznan University of Technology, Piotrowo 3, 60-950 Poznan, Poland. [email protected]

Wirunya Keawwattana, Department of Chemistry, Faculty of Science, Kasetsart University, Bangkok 10903, Thailand. [email protected]

xvii

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Contributors

Gue-Hyun Kim, Division of Applied Engineering, Dongseo University, Busan 617-716, Korea. [email protected] Il Kim, Department of Polymer Science and Engineering, Pusan National University, Busan 609-735, Korea. [email protected]

Yuri M. Krivoguz, Laboratory of Chemical Technology of Polymeric Composite Materials, V.A. Belyi Metal-Polymer Research Institute of National Academy of Sciences of Belarus, 32a Kiroc Street, 246050 Gomel, Belarus. [email protected] Thein Kyu, Department of Polymer Engineering, University of Akron, Akron, OH 44325, USA. [email protected]

Jesu´s Marı´a Garcı´a Martı´nez, Department of Physics and Engineering of Polymers, Polymer Engineering Group, Institute of Science and Technology of Polymers, CSIC CL Juan de la Cierva 3, 28006 Madrid, Spain. [email protected]

Rushikesh A. Matkar, Department of Polymer Engineering, University of Akron, Akron, OH 44325, USA. [email protected]

Liliya Minkova, Institute of Polymers, Bulgarian Academy of Sciences, Acad. G. Bonchev str. Bl.103A, 1113 Sofia, Bulgaria. [email protected]

Francis M. Mirabella, Lyondell Chemical Co., Equistar Technology Center, Cincinnati, OH 45249, USA. [email protected] Koh-Hei Nitta, Department of Chemical Engineering, Graduate School of Material Sciences, Kanazawa University, 920-1192, Japan. [email protected]

Domasius Nwabunma, 3M Company, Safety, Security, and Protection Business Services Laboratory, St. Paul, MN 55144, USA. [email protected] Emilia Pe´rez Collar, Department of Physics and Engineering of Polymers, Polymer Engineering Group, Institute of Science and Technology of Polymers, CSIC, CL Juan de la Cierva 3, 28006 Madrid, Spain. [email protected]

Stepan S. Pesetskii, Laboratory of Chemical Technology of Polymeric Composite Materials, V.A. Belyi Metal-Polymer Research Institute of National Academy of Sciences of Belarus, 32a Kirov Street, 246050 Gomel, Belarus. [email protected]

Subhendu Ray Chowdhury Department of Materials Science and Engineering, Pennsylvania State University, University Park, PA 16802, USA. [email protected]

Moonhor Ree, Department of Chemistry, Polymer Research Institute, Pohang Accelerator Laboratory, National Research Lab for Polymer Synthesis and Physics, and Center for Integrated Molecular Systems, Pohang University of Science & Technology (Postech), Pohang 790-784, Republic of Korea. [email protected]

Robert A. Shanks, School of Applied Sciences, RMIT University, GPO Box 2476V, Melbourne, Vic 3001, Australia. [email protected]

Jesu´s Taranco Gonza´lez, Department of Physics and Engineering of Polymers, Polymer Engineering Group, Institute of Science and Technology of Polymers, CSIC, CL Juan de la Cierva, 3, 28006 Madrid, Spain. [email protected]

Contributors

xix

Kohji Tashiro, Department of Future Industry-Oriented Basic Science and Materials, Toyota Technological Institute, Tempaku, Nagoya 468-8511, Japan. [email protected] Sabu Thomas, School of Chemical Sciences, Mahatma Gandhi University, Priyadarshini Hills, Kottayam 686560, Kerala, India. [email protected]

Sie C. Tjong, Department of Physics and Materials Science, City University of Hong Kong, Tat Chee Avenue, Kowloon, Hong Kong. [email protected] Shigeyuki Toki, Department of Chemistry, Stony Brook University, Stony Brook, NY 11794, USA. [email protected] Gong-Tao Wang, School of Aerospace, Mechanical & Mechatronic Engineering, The University of Sydney, Sydney, NSW 2006, Australia. [email protected] Min Wang, Centre for Advanced Composite Materials (CACM), Department of Mechanics and Engineering Science, School of Engineering, Peking University, 100871 Beijing, China. [email protected]

James L. White, Department of Polymer Engineering, The University of Akron, Akron, OH, 44325, USA. [email protected] Masayuki Yamaguchi, School of Materials Science, Japan Advanced Institute of Science and Technology, Nomi, Japan. [email protected] Jinhai Yang, Department of Polymer Engineering, The University of Akron, Akron, OH, 44325, USA. [email protected]

Part I

Introduction

Chapter

1

Overview of Polyolefin Blends Domasius Nwabunma1

1.1 INTRODUCTION Polyolefins are synthetic polymers of olefinic monomers. They are the largest polymer family by volume of production and consumption. Several million metric tons of polyolefins are produced and consumed worldwide each year, and as such they are regarded as commodity polymers. Polyolefins have enjoyed great success due to many application opportunities, relatively low cost, and wide range of properties. Polyolefins are recyclable and significant improvement in properties is available via blending and composite technologies. Polyolefins may be classified based on their monomeric unit and chain structures as ethylene-based polyolefins (contain mostly ethylene units), propylene-based polyolefins (contain mostly propylene units), higher polyolefins (contain mostly higher olefin units), and polyolefin elastomers (1). Ethylene-based polyolefins are normally produced either under low pressure conditions using transition metal catalysts resulting in predominantly linear chain structure or under high pressure conditions using oxygen or peroxide initiators resulting in predominantly branched chain structures of various densities and crystallinity levels. Propylene-based polyolefins are normally produced with transition metal catalysts resulting in linear chain structures with stereospecific arrangement of the propylene units or special stereoblock structures from a single-site catalyst. Higher polyolefins are normally produced using transition metal catalysts resulting in linear and stereospecific chain structures. Polyolefin elastomers based mainly on a combination of ethylene and propylene may be produced using metal or single-site catalysts with or without the inclusion of dienes (for cross-linking) and are mostly amorphous with high molecular weights and heterogeneous in phase structures. One

1 3M Company, Safety, Security, and Protection Business Services Laboratory, St. Paul, MN 55144, USA

Polyolefin Blends, Edited by Domasius Nwabunma and Thein Kyu Copyright # 2008 John Wiley & Sons, Inc.

3

4

Polyolefin Blends

may conclude that a given polyolefin may be a homopolymer, copolymer, or terpolymer depending on the number of monomers used in making the polyolefin, crystalline or amorphous depending on their chain conformation, configuration, and processing conditions. Today, polyolefins and polyolefin-based materials are used in many applications. These applications include transportation (automotive, aerospace), packaging, medical, consumer products (toys, appliances, etc.), electronics, cable and wire coating, thermal and acoustic insulation, and building and construction. Polyolefins can be extruded as filaments (fibers), films (cast and blown), and pipes/profiles. They can be molded into parts of various shapes. They can be foamed with physical and chemical foaming/blowing or/and can be coated onto other materials.

1.2 OLEFINIC MONOMERS The alkenes having one or more unsaturated double bonds in their structures are the monomers used to synthesize polyolefins. They have the general formula Cn H2n ; n
2. Table 1.1 shows the first 10 members of the olefinic monomers with one double bond, which are often called a-olefins. The monomers in Table 1.1 form a homologous series of hydrocarbon compounds. Thus, apart from having the same general formula, all compounds in the series have the same functional groups. Each member of the group differs from the next in the series by the CH2 group equivalent to 14 relative molecular mass units. All members of the series have similar chemical properties. The physical properties of the compounds in the series show a progressive change with increasing relative molecular mass. The first three members of the alkenes homologous series are gases at room temperature. Those containing between 5 and 15 carbon atoms are colorless liquids and the higher compounds are waxy solids at room temperature. These a-olefinic monomers may be obtained as products of the cracking of the gas– oil and naphtha fractions of petroleum distillations. They can also be obtained from synthetic organic chemistry methods.

Table 1.1

Alkene Monomers with One Double Used in the Synthesis of Polyolefins.

No. of carbon atoms (n) 2 3 4 5 6 7 8 9 10

Formula (Cn H2n ; n
2)

Name (other name)

C2H4 C3H6 C4H8 C5H10 C6H12 C7H14 C8H16 C9H18 C10H20

Ethene (ethylene) Propene (propylene) Butene-1 (butylene) Pentene-1 Hexene-1 Heptene-1 Octene-1 Nonene-1 Decene-1

Chapter 1

Overview of Polyolefin Blends

5

1.3 POLYOLEFIN HOMOPOLYMERS, COPOLYMERS, AND TERPOLYMERS Polyolefin homopolymers, copolymers, and terpolymers are foundation materials for polyolefin blends. They may be obtained via radical or ionic chain growth polymerization of alkenes using conventional free radicals (e.g., from peroxides) and organometallic complexing (Ziegler–Natta and metallocenes) catalyst systems. Polyolefin polymerization technologies and novel catalyst systems have enabled the rapid development of polyolefins with a wide range of molecular chain structures, morphologies, properties, and particle size and shape. Polyolefin homopolymers include polyethylene (PE), polypropylene (PP), polybutene-1 (PB), polymethylpentene-1 (PMP), and higher polyolefins. Table 1.2 shows the structures of commercial polyolefin homopolymers. Of these, PE and PP are the largest by amount produced yearly by the global polyolefin companies (1). PE comes in various forms differing in chain structures, crystallinity, and density levels. These are high density polyethylene (HDPE), low density polyethylene (LDPE), linear low density polyethylene (LLDPE), ultralow density polyethylene (ULDPE), and ultrahigh molecular weight polyethylene (UHMWPE). PP and higher polyolefins come in three stereospecific forms of varying densities: isotactic, syndiotactic, and atactic forms. Polyolefin copolymers involve two olefinic monomers. The process of copolymerization is normally used to control the properties of the polyolefins. Some of the consequences of copolymerization are reduced crystallinity, melting point, modulus, strength, hardness, and low temperature impact. Polyolefin copolymers are either random or block copolymers of same or different monomers and may be a single phase or heterophasic depending on the amount of comonomer, the polymerization catalyst, and the process. For polyolefin copolymer of same monomers, this can be achieved by having different segments of the copolymer with different tacticities. One can have polyolefin block copolymers of same block or of varying block lengths. One can also have polyolefin copolymers consisting of both block and random segments together in the same macromolecule. Polyolefin copolymers are usually not homogeneous in composition but are actually mixtures of copolymers of varying compositions. It is also possible with polyolefins to have block copolymers with only one monomer. These are called stereoblock copolymers and can be achieved by having sections of the polyolefin copolymer possess different tacticities. Polyolefin copolymers started with LLDPE and ethylene–propylene rubber (EPR). Today, there are polyolefin copolymers of ethylene with butene-1, hexene1, octene, cyclopentene, and norbornene and copolymers of propylene with butene-1, pentene-1, and octene-1 in addition to ethylene. There are copolymers of butene-1 with pentene-1, 3-methylbutene-1, 4-methylpentene-1, and octene in addition to its copolymers with ethylene and propylene. There are copolymers of 4-methylpentene1 with pentene-1 and hexene-1 in addition to its copolymers with butene-1 and propylene. The function of the comonomers is to reduce crystallinity, as compared to the homopolymers, resulting in copolymers that are highly elastomeric with very low

6

Polyolefin Blends

Table 1.2

Structures of Commercial Polyolefin Homopolymers.

Name (other name) Polyethylene (polyethene, polymethylene)

Chemical structure (repeat unit) CH

CH 2 n

2

Polypropylene (polypropene)

CH

2

Polybutylene (polybutene-1)

CH

n

C H2 CH3

Polyisobutylene (polyisobutene-1)

Polybutadiene

CH

CH

CH

2

CH

CH

2

CH

2 n

n

C H2

Poly-4-methylpentene-1

CH3 C H3

CH3 CH 3

Polyisoprene CH 2

C

CH

CH

2 n

glass transition temperatures, high impact strength, low modulus, low density, and are often optically transparent. The most widely used multiphase polyolefin copolymer is polypropylene impact copolymer. These copolymers are typically composed of isotactic polypropylene (iPP) and EPR. Impact polypropylene copolymers are produced by various processes, but are generally characterized by the synthesis of iPP in the first reactor and EPR in the second reactor. Therefore, these systems are typically reactor blends. Postreactor blending can be done, but the starting material is most often the reactor blend polypropylene copolymer. Polyolefin copolymers are often used for film applications or as impact modifiers. Polyolefin terpolymers contain three olefinic monomers. A well-known example is ethylene propylene diene monomer (EPDM). The diene (double bond) monomer is

Chapter 1

Overview of Polyolefin Blends

7

usually ethylidene norbornene or 1,4-hexadiene. EPDM was introduced because of the difficulty in cross-linking saturated polyolefin homopolymers and copolymers. There are also functionalized polyolefins. These are usually copolymer or terpolymer containing functional groups like epoxide, anhydride, hydroxyl, acrylate, and carboxylic acid. These functional groups are either grafted onto the polyolefin after polymerization or added directly in situ during polymerization reactions involving olefins and functional groups bearing polar monomers such as vinyl acetate, methyl acrylate, butyl acrylate, glycidyl methacrylate, and acrylic acid. Functionalized polyolefins are useful compatibilizers and impact modifiers in blends and composites containing polyolefins and nonpolyolefins. In this sense, functionalized polyolefins may be considered as additives rather than matrix materials in formulation of polyolefin blends. Commercial polyolefins often contain additives such as colorants, flame retardants, antioxidants, light stabilizers, nucleating agents, antistatic agents, lubricants (microcrystalline waxes, hydrocarbon waxes, stearic acid, and metal stearates), and so on. These additives aid the processing and fabrication of products from polyolefins. Detailed treatments about specific polyolefins, polymerization systems/ mechanism/processes, structures, properties, processing, and applications may be found in References 2–9.

1.4 POLYOLEFIN BLENDS Polymer blends (mixtures of structurally different polymers (10–19)) are of interest because synthesized polymers have not satisfied increasing application demands. Polyolefin blends are a subset of polymer blends and may be classified into two groups. The first group contains polyolefins only, which are formulated to broaden the range of structures, properties, and applications offered by polyolefins. The second group contains polyolefins and nonpolyolefins, which are formulated to mitigate some of the property drawbacks of the polyolefin or the nonpolyolefin. For a blend to be classified as a polyolefin blend, it is presumed that the polyolefin component is of significant composition in the blend. In terms of miscibility, polyolefin blends may also be classified as miscible and immiscible blends (10, 11). Polyolefin blending requires knowledge of the miscibility and crystallinity of the blend, in addition to the contributions of the components of the blend. Miscibility depends on molecular structure, blend composition, and mixing temperature. To characterize miscibility, a phase diagram is needed. Nonolefinic thermoplastic polymers that in principle may be blended with polyolefins include polyamides (nylons) such as polyamide 6, polyamide 66, polyphenylene sulfide (PPS), polyphenylene ether (PPE), and polyphenylene oxide (PPO); polyesters such as polyethylene terephthalate (PET), polybutylene terephthalate (PBT), polyethylene naphthalate (PEN), polytrimethylene terephthalate (PTT), polycarbonates, polyethers, and polyurethanes; vinyl polymers such as polystyrene (PS), polyvinyl chloride (PVC), polymethylmethacrylate (PMMA), and ethylene

8

Polyolefin Blends

vinyl acetate copolymer (EVA); block and graft copolymers (styrene–acrylonitrile copolymer, styrene–butadiene copolymer, styrene–ethylene–butadiene–styrene terpolymer, etc.); and liquid crystalline polymers (LCPs). Thermosetting resins that may be blended with polyolefins include, but are not limited to, the following: epoxies, unsaturated polyesters, phenol formaldehyde, melamine formaldehyde, urea formaldehyde, silicones, and so on. The properties that polyolefins normally contribute in blends with other polymers include high melt strength and elasticity, toughness, low viscosity for processability, low polarity, dielectric constant, and loss, and chemical resistance and moisture absorption resistance. Nonpolyolefins contribute to high modulus, heat resistance, and oxygen or solvent barrier properties. For example, barrier properties of polyolefins can be improved by blending with polymers such as ethylene vinyl alcohol and polyvinylidene chloride. Blends of polyolefins with nylons and polycarbonate (PC) allow balanced control of permeability and water retention. Polystyrene (PS) is an interesting candidate for blending with polyolefins for mechanical reasons as well as for paintability and printability. There are two classes of polyolefin blends: elastomeric polyolefin blends also called polyolefin elastomers (POE) and nonelastomeric polyolefin blends. Elastomeric polyolefin blends are a subclass of thermoplastic elastomers (TPEs). In general, TPEs are rubbery materials that are processable as thermoplastics but exhibit properties similar to those of vulcanized rubbers at usage temperatures (19). In TPEs, the rubbery components may constitute the major phase. However, TPEs include many other base resins, which are not polyolefins, such as polyurethanes, copolyamides, copolyesters, styrenics, and so on. TPEs are now the third largest synthetic elastomer in total volume produced worldwide after styrene– butadiene rubber (SBR) and butadiene rubber (BR). Two important types of elastomeric polyolefin blends are reactor-made iPP/ EPR blends and postreactor blend iPP/EPDM. The latter is called thermoplastic vulcanizates (TPVs), produced by dynamic vulcanization of blends containing a thermoplastic and an elastomer. To make iPP/EPDM TPV, the two polymers PP and EPDM are mixed with curatives, such as peroxides, phenolic resins, or sulfur with accelerators, and dynamically cured in an extruder resulting in a blend consisting of micrometer-sized elastomer particles dispersed in the PP matrix (20–24). Paraffinic oils are added in the melt mixing process for viscosity control and cost. In iPP/ EPDM TPV, the crystalline iPP resin is normally the minor phase. Recently, polyolefin plastomers have been added to the class of elastomeric polyolefin blends. Polyolefin plastomers are ultralow molecular weight linear low density polyethylenes (ULMW-LLDPE). Nonelastomeric polyolefin blends are blends of polyolefins with mostly nonpolyolefin (other thermoplastic) matrices as mentioned earlier. Polyolefin blends of commercial importance are normally made via two methods: blending in the melt either during polymerization or mechanically after the polymerization process. The first method, called in-reactor blending, involves the blending of different polyolefins (homopolymers, random, and block copolymers) in a polymerization reactor. This is enabled by the presence of multiple catalyst species

Chapter 1

Overview of Polyolefin Blends

9

in the polymerization recipe. A good example is in-reactor-made EPR/iPP blend, which is normally prepared by adding ethylene monomer to propylene monomer toward the end of propylene polymerization process. The function of EPR is to improve iPP flexibility; hence, EPR/iPP blend is often called toughened or impact PP and finds wide applications in the consumer industry. Another good example of in-reactor-made polyolefin blend is linear low density polyethylene (LLDPE), which often contains several ethylene/a-olefin copolymers that differ in ethylene contents. LLDPE and reactor-made EPR/iPP are often called thermoplastic olefins (TPOs). The second method, called postreactor blending, involves mechanical blending of a premade polyolefin with other polyolefins or nonpolyolefins in compounding extruders. A practical example of polyolefin blend made using this method is a blend of isotactic PP with cured EPDM, as described earlier. Other examples of polyolefin blends made by postreactor mechanical blending include polyolefin/polyamides (nylons), polyolefin/polyesters, polyolefin/polystyrene (PS), polyolefin/polyvinyl chloride (PVC) blends, and so on. EPR/iPP blend can also be made by postreactor mechanical blending as well as by in-reactor blending. Postreactor blending via single and twin screw compounding is still the preferred method of polyolefin blending because it is quick, easy, economical, and efficient. Table 1.3 shows a summary of specific polyolefin blends that have been studied in the literature extracted from references (25–310). These blends involve the following polyolefins: PP, PE (LLDPE, LDPE, HDPE, and UHMWPE), EPR, EPDM, and PB. Some of the blends listed in Table 1.3 are of commercial importance (6). There are many publications (journal articles and patents) on polyolefin blends in the literature. Tables 1.4 contains a summary of the journal articles (25–310) and patents published each year during the 6-year period 2000–2005. The choice of date range is arbitrary. The number of journal articles for each year was obtained from a search of electronic version of English-based polymer and polymer-related journals using the keywords polyolefin and blends. Within polyolefin keyword, the subkeywords used in the search were polyethylene (PE, LLDPE, LDPE, HDPE, UHMWPE, PE, etc.), polypropylene (PP, iPP, sPP, aPP, etc.), polybutene-1, poly-4-methylpentene-1, ethylene–diene monomer, ethylene–propylene– diene terpolymer, ethylene propylene rubber, thermoplastic olefins, natural rubber (NR), polybutadiene, polyisobutylene (PIB), polyisoprene, and polyolefin elastomer. For the polyolefin blends patent search, polymer indexing codes and manual codes were used to search for the patents in Derwent World Patent Index based on the above keywords listed in the search strategy. Table 1.4 shows an increasing trend in the number of publications. It should be noted that while generating Table 1.4, some publications may have been missed in the reference search period (2000–2005). There are several issues of interest in polyolefin blends research. They may be categorized into formulation design and processing; miscibility, structural, molecular, and property characterization; end-use properties and performance; and

10

PE/silicon rubber (76) PE/starch (85, 277, 289, 295) PE/hydrolyzed collagen (98) PE/natural rubber/PP (238, 254, 302) PE/poly(3-hydroxybutyrate) (251) PE/nitrile rubber (100, 183)

PE/polydianilinephosphazene (91)

PE/PS (47–48, 67, 79, 123, 145, 198, 219, 223) PE/PP/PA66 (71)

PE/PA6 (25, 27, 71, 88, 92, 96, 130, 140, 144, 149, 168, 204, 215, 224, 226, 272, 273) PE/PP (33, 34, 36, 47, 62, 69, 83, 93, 107, 116, 182, 196, 200, 203, 236, 237, 239, 243, 248, 255, 263, 297) PE/PP/cycloolefin copolymers (307) PE/ethylene-co-propylene-co-butene-1 (99) PP/PPE (216, 217) PP/EPDM (50, 66, 113, 121, 136, 199, 202, 207, 224, 267) PP/EPR (35, 63, 77, 118, 124, 152, 163, 195, 231, 233, 271) PP/epoxy (62, 120) PP/PA6/PS (205) PP/cyclopolyolefin (162) PP/natural rubber (64, 241, 247, 287) PP/recycle rubber(241, 247) PP/natural rubber/recycle rubber powder (249)

PE/PP/PA6 (71)

LLDPE/LDPE/wax (80)

Polyolefin blends PE/PP/EPR (35, 114) PE/PC (29, 173, 213) PE/poly(silsesquioxanes) (156) PE/ethylene acrylic elastomer (281)

Polyolefin Blends Studied in the Literature.

HDPE/UHMWPE (81) LDPE/HDPE (84, 153, 160, 221, 259, 284) LLDPE/HDPE (285) LLDPE/LDPE (72, 104, 108, 119, 128, 129, 181, 211, 214, 221, 242, 258, 261, 292, 309)

Table 1.3

PP/EPDM/epoxidized natural rubber (246) PP/epoxidized natural rubber (301) Butadiene/styrene–butadiene rubber (293) EPDM/PA66 EPDM/SBR (68) EPDM/NBR (234)

PP/EPDM/natural rubber (244–246)

PP/PC (25) iPP/sPP (40); iPP/aPP (103)

PP/PB (26, 109)

PP/PA6 (28, 54, 58, 71, 111, 170, 178, 262, 283, 286, 288, 305) PP/PA6/ethylene-co-octene (28, 31, 43, 46)

PP/PET (137) PP/PE/EPDM (49) PP/EVOH (135, 298) PP/PA66 (28)

11

PE/PBT (82, 300) PE/PS/PMMA (209) PE/wax (256)

PE/PA12 (142) PE/PTT (37) PE/PVC (52, 201) PE/starch/PCL (90) PE/ethylene-co-octene (46, 278) PE/ethylene-co-butene-1 (89, 171) PE/PA6,6 (127, 169, 273) PE/perfluoropolyether (126) PE/EVA (45, 51, 56, 78, 97, 161, 270) PE/PET (166, 167, 208, 213, 220, 274)

PP/ethylene-co-methyl acrylate (164) PP/P12 (304) PP/TPU (228) PP/SBS (74) PP/SAN (184) PP/EVA (177, 189, 191, 193, 294) PP/SEBS (74, 112, 218) PP/SEP (105, 218) PP/SEPS (105, 218) PP/PS (38, 39, 117, 155, 159, 227, 212, 229, 299, 310) PP/ethylene-co-octene (59, 87) PP/ethylene-co-butene-1 (59) PP/LCP (225, 232) Natural rubber/polystyrene (290) Natural rubber/reclaimed rubber (252) Epoxidized natural rubber/natural rubber (253)

EPDM/SAN (53, 308) EPDM/PA6 (101, 279) EPDM/polyaniline (122) EPDM/TPU (95) EPDM/PP/PA6 (262) EPDM/PA12,10 (262) EPDM/PP6,10 (266) EPDM/natural rubber (44, 65, 240, 250, 269) EPR/polydimethylsiloxane (70) EPR/PA6 (57, 60)

12

Polyolefin Blends

Table 1.4 Summary of Number of Electronic articles on Polyolefin Blends Published Between 2000 and 2005 in English Language-based Polymer and Polymer-related Journals in Comparison to the Number of Patents. Journal articles (25–310) Year Number of Articles

2000 36

2001 31

2002 42

2003 42

2004 68

2005 66

Total 285

2004 322

2005 365

Total 2349

Patents Year Number of Patents

2000 517

2001 444

2002 406

2003 295

modeling and simulation. Regarding formulation design and processing, some of the issues of interest include i. Use of block, random, and graft copolymers in compatibilization. ii. Effects of molecular structure, weight, and additives. iii. Phase behavior, miscibility, and compatibility issues. iv. Batch and continuous mixing/compounding, extrusion, and molding into films, fibers, and other articles. v. Physical blending versus reactive blending. Regarding miscibility, structural, molecular, and property characterization, some of the issues of interest include i. Morphological characterization using techniques such as scanning electron microscopy (SEM), transmission electron microscopy (TEM), atomic force microscopy (AFM), and polarized light microscopy (PLM). ii. Structural characterization using radiation scattering and diffraction techniques such as X-ray scattering (XRS), X-ray diffraction (XRD), and electron diffraction (ED). iii. Structural characterization using spectroscopic techniques such as nuclear magnetic resonance (NMR), Fourier transform infrared (FTIR), and temperature rising elution fractionation (TREF). iv. Melt rheology. v. Phase separation dynamics. vi. Thermal transitions and thermal stability. vii. Isothermal and nonisothermal crystallization behavior under quiescent and nonquiescent conditions. viii. Crystallinity and crystal structure determination by various methods such as density, differential scanning calorimetry (DSC), XRD, and ED. ix. Tacticity, copolymer sequence distribution, and comonomer composition.

Chapter 1 Overview of Polyolefin Blends

13

Regarding end-use properties and performance, some of the issues of interest include i. Mechanical (static and dynamic) behavior under tensile, shear, or compressive, and impact mode. ii. Plastic deformation. iii. Thermomechanical and thermal stability behavior. iv. Adhesion, interfacial, and interphase behavior. v. Fire resistance/flammability behavior. vi. Barrier and transport properties. vii. Optical properties (gloss, haze, and transparency). viii. Surface/tribological properties. ix. Electrical/dielectric properties. x. Aging effects (time–temperature-dependent behavior). Modeling and simulation involves theoretical analysis of processing, characterization, and performance behavior using phenomenological, atomistic, molecular dynamics, and Monte Carlo methods, among others, and comparisons with experimental results. Polyolefin blends are generally immiscible due to hydrophobic or nonpolar nature of polyolefins. The immiscibility leads to phase separation, which is responsible for the poor mechanical properties of these blends. Immiscibility of polyolefin/ nonpolyolefin blends can be mitigated by adding a proper compatibilizer. The function of a compatibilizer is to reduce interface tension (strengthen the interface between the phases) and thus improve mechanical properties of the stabilized blend (reduce the size and morphology of the phase-separated phases). The compatibilizer strengthens the interface by broadening it from a sharp change in composition and properties to a broader gradual transition interface. Thus, the chief concern in compatibilization of polyolefin/nonpolyolefin blends is phase morphology. One way to achieve compatibilization involves physical processes such as shear mixing and thermal history, which modify domain size and shape. The second way is the use of physical additives to increase attraction between molecules and phases. The third method is reactive processing, which is used to change the chemical structure of one or more of the components in the blend and thus increase their attraction to each other. Table 1.5 contains a list of compatibilizers used in the formulation of polyolefin blends. As can be seen from Table 1.5, most of the compatibilizers used in the formulation of polyolefin blends contain compounds such as maleic anhydride, acrylic and methacrylic acid, glycidyl methacrylate, and diblock and triblock copolymers involving styrene, ethylene, and butadiene.

1.5 TRENDS IN POLYOLEFIN BLENDS The first trend is the growing use of in-reactor blending technique to produce new elastomeric or toughened polyolefins for demanding applications traditionally reserved

14

Polyolefin Blends

Table 1.5

Compatibilizers Used in the Formulation of Polyolefin Blends.

Compatibilizer

Polyolefin blend

PP-g-maleic anhydride Ethylene-co-acrylic acid Ethylene-co-glycidyl methacrylate Styrene-b-ethylene-co-propylene-g-maleic anhydride PE-g-maleic anhydride Styrene-co-ethylene-co-butadiene-costyrene-g-maleic anhydride Sodium-neutralized ethyleneco-methacrylic acid PE-g-glycidyl methacrylate Styrene-co-ethylene-co-butadiene-costyrene-g-glycidyl methacrylate Styrene-co-ethylene-co-butadiene-co-styrene Styrene-co-ethylene-co-butadiene-costyrene-g-maleic anhydride PE-g-maleic anhydride EPR-g-maleic anhydride PE-g-maleic anhydride PP-g-maleic anhydride Styrene-co-ethylene-co-butadiene-co-styrene Styrene-co-butadiene Styrene-co-ethylene-co-butadiene-co-styrene Styrene-co-butadiene-co-styrene Liquid polybutadiene þ diakyl peroxides Chlorinated PE Ethylene-co-acrylic acid Bismaleimide-g-PE PE-g-maleic anhydride PE-g-epoxidized natural rubber PE-g-diethyl maleate Sodium neutralized ethylene-co-acrylic acid PE-g-glycidyl methacrylate PE-g-maleic anhydride Ethylene-co-glycidyl methacrylate Styrene-co-ethylene-co-butadiene-co-styrene Styrene-co-ethylene-co-butadiene-co-styrene PP-g-3-isopropenyl-a,a-dimethylbenzene isocyanate PP-g-succinic anhydride and PP-g-fluorescein PP-g-maleic anhydride Ethylene-co-butyl acrylate-g-maleic anhydride Styrene-co-ethylene-co-butadiene-co-styreneg-glycidyl methacrylate

PE/PA6 PE/PA6 PE/PA6 PE/PA6

(27, 88) (27, 92, 140, 204, 215, 224) (27, 144) (130)

PE/PA6 (130) PE/PA6 (130, 168) PE/PA6 (92) PE/PA6 (144) PE/PA6 (168) PE/PP/PA6,6 (71) PE/PP/PA6,6 (71) PE/PA12 (142) PE/PP (47) PE/PP (71, 239) PE/PP (71, 239) PE/PS (47) PE/PS (48, 115) PE/PS (79) PE/PS (145) PE/PS (223) PE/PVC (52) PE/hydrolyzed collagen (98) PE/hydrolyzed collagen (98) PE/acrylonitrile-co-butadiene (100) PE/scrap rubber powder (165) PE/PET (166) PE/LCP (188) PE/PET (220) PE/ethylene-co-acrylic acid (281) PE/plasticized tapioca starch (289) PP/PA6,6 (28) PP/PA6 (28) PP/PA6 (54) PP/PA6 (111) PP/PA6 (58, 170, 178, 262, 283, 286, 289) PP/PA6 (170) PP/PA6 (178)

Chapter 1 Overview of Polyolefin Blends Table 1.5 (Continued) Compatibilizer

Polyolefin blend

PP-g-bismaleimide PP-g-acrylic acid EPDM-g-maleic anhydride Ethylene-co-octene-g-maleic anhydride PP-g-maleic anhydride Styrene/AlCl3-catalyzed friedel–crafts alkylation PP-g-PS Styrene-co-ethylene-co-butadiene-co-styrene Styrene-co-ethylene-co-butadiene-costyrene þ ionomer resin Ionomer Zn2þ Styrene-co-ethylene-co-propylene-g-glycidyl methacrylate Styrene-co-ethylene-co-butadiene-costyrene-g-glycidyl methacrylate PP-g-monomethyl itaconate PP-g-dimethyl itaconate PP-g-maleic anhydride PP-g-acrylic acid PP-g-monomethyl itaconate PP-g-primary amine PP-g-secondary amine N,N-m-Phenylene bismaleimide N,N-m-Phenylene bismaleimide Polyoctenamer Polystyrene-modified natural rubber PP-g-maleic anhydride Hydroxylated EVA Ethylene-co-butyl acrylate-co-glycidyl methacrylate EPDM-g-maleic anhydride EPDM-g-styrene-co-acrylonitrile Ethylene-co-octene-g-maleic anhydride Mercapto-modified EVA EPDM-g-maleic anhydride EPDM-g-maleic anhydride EPDM-g-vinyltriethoxysilane EPDM-g-vinyloxyaminosilane EPR-g-maleic anhydride EPR-g-MAH þ epoxy-modified lignosulfonate Styrene-co-isoprene

PP/PA6 (262) PP/PA6 (262) PP/PA6 (279) PP/PA6 (286) PP/epoxy (62, 120) PP/PS (21, 117) PP/PS (155, 299) PP/PS (299) PP/PS/PA6 (205) PP/EVOH (135) PP/PET (137) PP/PET (137) PP/EPDM (139) PP/EPDM (139) PP/EPDM/PA6 (262) PP/EPDM/PA6 (262) PP/EPR (163) PP/TPU (228) PP/TPU (228) PP/natural rubber/PE (238) PP/EPDM/natural rubber (244–245) PP/PE/natural rubber (254) PP/natural rubber (288) PP/EVA (294) PP/EVA (294) PTT/PE (37) EPDM/natural rubber (44, 240, 250) EPDM/SAN (53) EPDM/TPU (91) EPDM/NBR (234) EPDM/PA6,10 (266) EPDM/PA12,10 (266) EPDM/PE (275) EPDM/PE (296) EPR/PA6 (57, 60) EPR/lignin (172) Natural rubber/PS (290)

15

16

Polyolefin Blends

for more expensive engineering thermoplastics. There is also a growing interest in the use of elastomer-rich thermoplastic olefin blends that contain 60–75% rubber and 40– 25% polypropylene for automotive interior applications because they provide the required amount of lower and upper service temperature performance. There is also growing use of beta nucleator functional additive in polyolefin blends, for example, PP/ EPR blend. A beta nucleating agent leads to blend with improved impact and ductility. The second growing trend is the impact modification of polyolefin blends using styrenic block copolymers, which are known to be clear, strong, have low glass transition, compatible with PP, form interpenetrating polymer networks, and very efficient in contrast to maleic anhydride-grafted polyolefins. The third trend is the growing interest in functionalization of polyolefin blends in their melt by means of reactive extrusion. Particular attention has been paid to blended systems PP/PE, PP/EPR, PE/ethylene–octene copolymer (EOC), PP/EOC, PE/PS, and PP/PS functionalized in melt by reactive extrusion. The major field for application of functionalized polyolefin blends is compatibilization of blends of condensation polymers, where they can be used in place of homopolyolefins. The fourth trend is spurred by environmental sustainability concerns and the need for increased recyclability and reuse of polyolefin blends. In this regard, there is increasing replacement of PVC by polyolefin–polyolefin blends. There is also an increase in recyclability of EPDM rubber vulcanizates since EPDM is the fastest growing elastomer among synthetic rubber and the most used of nontire rubbers. Also, cryogenically ground rubber tires are being used as fillers for polyolefin blends such as LLDPE/HDPE. The fifth trend is the use of beta nucleator as opposed to alpha nucleator in PP containing blends such as PP/EPR. The use of beta nucleating agent results in blend with improved impact and ductility The sixth trend is the growing interest in foams, nonwoven, and elastic materials based on polyolefin blends The seventh trend is the increasing use of novel processing methods. For example, there is growing use of supercritical fluids (e.g., supercritical carbon dioxide and nitrogen gases) to foam polyolefin blends for density reduction. There is use of ultrasound to, for example, devulcanize cross-linked rubber. There is use of solid-state shear mechanical processing to break the polyolefin blend material into submicron particles to make environment friendly (water-based) polyolefin dispersions. There is use of electrospinning technique to make polyolefin fibers and in particular nanofibers. In conclusion, the importance of polyolefin blends has increased dramatically in the last couple of decades and this is sure to continue as new polyolefins are made and as new applications are sought for these materials.

NOMENCLATURE AFM BR

Atomic force microscopy Butyl rubber

Chapter 1 Overview of Polyolefin Blends

DSC ED EOC EPDM EPR EVA EVOH FTIR HDPE LCP LDPE LLDPE NBR NMR NR PA PA12 PA12,10 PA6 PA6,6 PA6,10 PB PBT PC PCL PE PEN PET PIB PLM PMMA PMP PP iPP sPP aPP POE PPE PPO PPS PS PTT PVC SAN SB

Differential scanning calorimetry Electron diffraction Ethylene octene copolymer Ethylene propylene diene monomer Ethylene propylene rubber Ethylene-co-vinyl acetate Ethylene-co-vinyl alcohol Fourier transform infrared High density polyethylene Liquid crystal polymer Low density polyethylene Linear low density polyethylene Nitrile butadiene rubber Nuclear magnetic resonance Natural rubber Polyamide Polyamide 12 Polyamide 12,10 Polyamide 6 Polyamide 6,6 Polyamide 6,10 Polybutene-1 Polybutylene terephthalate Polycarbonate Polycaprolactone Polyethylene Polyethylene naphthalate Polyethylene terephthalate Polyisobutylene Polarized light microscopy Polymethylmethacrylate Polymethylpentene-1 Polypropylene Isotactic polypropylene Syndiotactic polypropylene Atactic polypropylene Polyolefin elastomer Poly(2,6-dimethyl-1,4-phenylene ether) Polyphenylene oxide Polyphenylene sulfide Polystyrene Polytrimethylene terephthalate Polyvinyl chloride Styrene-co-acrylonitrile Styrene-co-butadiene

17

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Polyolefin Blends

SBR SBS SEM SEBS SEP SEPS TEM TPE TPO TPU TPV TREF UHMWPE ULDPE ULMW-LLDPE XRD XRS

Styrene butadiene rubber Styrene-co-butadiene-co-styrene Scanning electron microcopy Styrene-co-ethylene-co-butadiene-co-styrene Styrene-co-ethylene-co-propylene Styrene-co-ethylene-co-propylene-co-styrene Transmission electron microscopy Thermoplastic elastomer Thermoplastic olefin Thermoplastic polyurethane Thermoplastic vulcanizate Temperature rising elution fractionation Ultrahigh molecular weight polyethylene Ultralow density polyethylene Low molecular weight linear low density polyethylenes X-ray diffraction X-ray scattering

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Polyolefin Blends

205. M. A. Debolt and R. E. Robertson, Polym. Eng. Sci., 44, 1800 (2004). 206. M. H. Ha and B. Y. Kim, Polym. Eng. Sci., 44, 1858 (2004). 207. W. Feng and A. I. Isayev, Polym. Eng. Sci., 44, 2019 (2004). 208. Z.-M. Li, B.-H. Xie, R. Huang, X.-P. Fang, and M.-B. Yang, Polym. Eng. Sci., 44, 2165 (2004). 209. K. P. Tchomakov, B. D. Favis, M. A. Huneault, F. Champagne, and F. Tofan, Polym. Eng. Sci., 44, 749 (2004). 210. J.-T. Yeh, S.-S. Huang, and H.-Y. Chen, Polym. Eng. Sci., 45, 25 (2005). 211. Y. Fang, P. J. Carreau, P. G. Lafleur, and S. Ymmel, Polym. Eng. Sci., 45, 343 (2005). 212. V. Thirtha, R. Lehman, and T. Nosker, Polym. Eng. Sci., 45, 1187 (2005). 213. H.-S. Xu, Z.-M. Li, S. Y. Yang, J.-L. Pan, W. Yang, and M.-B. Yang, Polym. Eng. Sci., 45, 1231 (2005). 214. Y. Fang, P. J. Carreau, and P. G. Lafleur, Polym. Eng. Sci., 45, 1254 (2005). 215. F. P. La Mantia, L. Canfora, and N. Tzankova Dintcheva, Polym. Eng. Sci., 45, 1297 (2005). 216. S. S. Morye, Polym. Eng. Sci., 45, 1369 (2005). 217. S. S. Morye, Polym. Eng. Sci., 45, 1377 (2005). 218. Y. Matsuda, M. Hara, T. Mano, K. Okamoto, and M. Ishikawa, Polym. Eng. Sci., 45, 1630 (2005). 219. C. Lu, X. Yu, and S. Guo, Polym. Eng. Sci., 45, 1666 (2005). 220. F. Pazzagli and M. Pracella, Macromol. Symp., 149, 225 (2000). 221. J. Murı´n, J. Uhrin, and I. Choda´k, Macromol. Symp., 170, 115 (2001). 222. B. L. Lopez, C. Gartner, and M. Hess, Macromol. Symp., 174, 277 (2001). 223. D. Hlavata´, Z. Krulisˇ, Z. Hora´k, F. Lednicky´, and J. Hroma´dkova´, Macromol. Symp., 176, 93 (2001). 224. M. R. Arroyo and M. L. Bell, Macromol. Symp., 170, 181 (2001). 225. S. Bualek-Limcharoen, S. Saengsuwan, T. Amornsakchai, and B. Wanno, Macromol. Symp., 170, 189 (2001). 226. F. P. la Mantia, R. Scaffaro, A. Valenza, A. Marchetti, and S. Filippi, Macromol. Symp., 198, 173 (2003). 227. J. Pointeck, P. Po¨tschke, N. Proske, H. Zhao, H. Malz, D. Beyerlein, U. Schulze, and B. Voit, Macromol. Symp., 198, 209 (2003). 228. Q.-W. Lu, C. W. Macosko, and J. Horrion, Macromol. Symp., 198, 221 (2003). 229. Y. Roiter, V. Samaryk, S. Varvarenko, N. Nosova, I. Tarnavchyk, J. Pointeck, P. Po¨tschke, and S. Voronov, Macromol. Symp., 210, 209 (2004). 230. T. Trongsatikul, D. Aht-Ong, and W. Chinsirikul, Macromol. Symp., 216, 265 (2004). 231. H. H. Kausch, Macromol. Symp., 225, 165 (2005). 232. G. Guerrica-Echevarrı´a, J. I. Eguiazaba´l, and J. Naza´bal, J. Polym. Sci. B. Polym. Phys., 38, 1090 (2000). 233. F. Mighri, M. A. Huneault, A. Ajji, G. H. Ko, and F. Watanabe, J. Appl. Polym. Sci., 82, 2113 (2001). 234. M. G. Oliveira and B. G. Soares, J. Appl. Polym. Sci., 91, 1404 (2004). 235. J. Luettmer-Strathmann, J. Chem. Phys., 123, 014910 (2005). 236. D. Heine, D. T. Wu, J. G. Curro, and G. S. Grest, J. Chem. Phys., 118, 914 (2003). 237. R. Strapasson, S. C. Amico, M. F. R. Pereira, and T. H. D. Sydenstricker, Polym. Test., 24, 468 (2005). 238. A. Hassan, M. U. Wahit, and C. Y. Chee, Polym. Test., 22, 281 (2003). 239. C. Li, Y. Zhang, and Y. Zhang, Polym. Test., 22, 191 (2003). 240. S. H. El-Sabbagh, Polym. Test., 22, 93 (2003). 241. H. Ismail and Suryadiansyah, Polym. Test., 21, 389 (2002). 242. J.-Z. Liang, Polym. Test., 21, 69 (2002). 243. C. M. Tai, R. K. Y. Li, and C. N. Ng, Polym. Test., 19, 143 (2000). 244. Halimatuddahliana, H. Ismail, and H. Md. Akil, Polym. Plast. Technol. Eng., 44, 1429 (2005).

Chapter 1 Overview of Polyolefin Blends

25

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Polyolefin Blends

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Chapter

2

Miscibility and Characteristics of Polyolefin Blends James L. White1 and Jinhai Yang1

2.1 INTRODUCTION It is our purpose here to summarize the characteristics of blends of polyolefins. We consider both blends of polyolefins with other polyolefins and with nonpolyolefins including polyamides and polystyrenes. In the case of interpolyolefin blends, our primary concern is with miscibility. In the case of blends of polyolefins with nonpolyolefins, the blends are all immiscible and our concern is with phase morphology. We also consider three component blend systems where the third component is a surfactant or compatibilizing agent, which collects at the interface.

2.1.1 Polyolefins Olefins are a group of unsaturated hydrocarbons of the general formula CnH2n. The polymers of olefins are known as polyolefins. Commercial polyolefins mainly include homopolymers of the olefin monomers: ethylene, propylene, butene-1, isobutene, and 4-methylpentene-1. These homopolymers have structure units as shown in Fig. 2.1, where the asterisk indicates asymmetric carbon atoms. Thus, polypropylene (PP), poly (butene-1) (PB1), and poly(4-methylpentene-1) (P4MP1) have different tactic forms. The most important commercial polyolefins are polyethylene, polyisobutene, and the isotactic forms, that is, iPP, iPB1, and iP4MP1. Polyisobutene was first polymerized by the IG Farbenindustries (BASF) in the late 1920s. Polyethylene was first polymerized by ICI in the late 1930s in a branched form (1). Linear polyethylene

1

Department of Polymer Engineering, The University of Akron, Akron, OH 44325, USA

Polyolefin Blends, Edited by Domasius Nwabunma and Thein Kyu Copyright # 2008 John Wiley & Sons, Inc.

27

28

Polyolefin Blends

CH2 CH2

PE

n

*

CH2 CH n CH3

PP

CH3 CH2 C n CH3

PIB

*

CH2 CH n CH2 CH3 PB1

*

CH2 CH n CH2 CH CH3 CH3 P4MP1

Figure 2.1 Structure units of polyolefins.

was developed by Ziegler and coworkers (2) and by Phillips Petroleum (3) in the 1950s. The other isotactic polyolefins were polymerized by Natta et al. (4,5) using modified Ziegler catalysts in the 1950s. It is a very interesting question why only these five types of polyolefins are commercialized, considering that there are many more types of olefin monomers. This is largely associated with their crystalline melting temperatures. Polyethylene, polypropylene, and polybutene-1 have reasonable melting temperatures, in particular, 135, 165, and 120
C, respectively. However, the melting temperatures decrease greatly with increase of the pendant group length in the isotactic polymers, for example, polypropylene (165
C), polybutene-1 (120
C), polypentene-1 (70
C), polyhexylene-1 (55
C), polyheptylene-1 (40
C), and polyoctylene-1 (38
C) (6,7). Introducing methyl groups onto the side group raises the melting temperatures, for example, poly(3-methylbutene-1) (300
C), poly(4-methylpentene-1) (235
C) (7). Polyisobutylene, which crystallizes with difficulty has become a significant synthetic rubber. In addition to the homopolymers, there are various polyolefin copolymers. Linear low density polyethylene (LLDPE) is probably the largest. This involves ethylene copolymerized variously with butene-1, hexene-1, octene-1, and 4methylpentene-1, and so on. Here the second comonomer content is around 5 mol%. In the 1990s, Exxon (8) and Dow Chemical (9,10) began to make polyolefin copolymers using new generation of Ziegler–Natta catalysts called metallocene catalysts. Polymers were also made with high comonomer contents and the density was as low as 0.86 g cc1 (compared to 0.83 g cc1 for amorphous polyolefins). It should be noted that these are random copolymers. Recently, Dow Chemical (11) made polyolefin block copolymers using ‘‘chain shuttling polymerization.’’ These block copolymers display melting temperature around 40
C higher than those of random copolymers of equivalent density. Ethylene–propylene copolymers, which usually contain roughly 50–75 mol% of ethylene, are also important polyolefin copolymers. These comonomers are elastomers and are defined as ethylene– propylene copolymer (EPM) or ethylene–propylene rubber. Terpolymers are commercially more important than EPM and are produced by polymerizing ethylene and propylene together with a diolefin, which are called ethylene–propylene diene terpolymer (EPDM). Since the late 1990s, another group of polyolefins, cyclopolyolefins, were developed (12,13). They are homo- or copolymers formed by the polymerization

Chapter 2 Miscibility and Characteristics of Polyolefin Blends

29

Figure 2.2 Norbornene form (I).

of norbornene derivatives of form (I) shown in Fig. 2.2. The bridged ring structure of norbornene does not readily fit into crystalline lattices, which together with comonomers make the resins completely amorphous with high transparency and high glass transition temperature, for example, 200
C with 60 mol% of norbornene (14). They are used in competition with polymers such as polycarbonate and polymethylmethacrylate for applications such as optical discs and lenses.

2.1.2 Blends Though more and more polymers have been synthesized since 1930s, however, they could not satisfy the fast growing millions of applications especially in electric, architecture, and automobile after 1960s. Thus, both blends and composites are compounded greatly to compensate the gap. So far, compounding is the most fast, economic, and efficient material development method for new applications. At present, about 36% of the synthetic resins are used in blends and about 39% in compounds with particle and fibers (15). Polymer blends date to the nineteenth century. In 1846, Parkes patented the first polymer blends of cis-1,4 polyisoprene (natural rubber) and trans-1,4 polyisoprene (gutta percha) (16). The modern era of polymer blending is thought to begin with the development of high impact polystyrene (HIPS) (17) and poly(acrylonitrile–butadiene–styrene) terpolymer (ABS) (18–21) in the 1950s and polyphenylene ether (PPE)/ polystyrene (PS) blends by General Electric (22) in the 1960s. Polymer blends can be divided into two groups: miscible and immiscible blends. Miscible blends are homogeneous and stable. Their properties tend to be intermediate. However, they are relatively few. Most polymer blends are immiscible. Their properties are strongly affected by their phase morphologies, which are decided by their viscosity, interfacial tension, and processing methods. In this review we will describe polyolefin blends. Many of these blends involve polar polymers with polyolefins.

30

Polyolefin Blends

These blends are immiscible and their interfaces are unstable. Special interfacial treatments are required to make them suitable as materials of commerce. A second group of blends are those in which the components are all polyolefins. These blends will be miscible or nearly so. Also in this paper the general questions of blend miscibility and interfacial characteristics will be treated.

2.2 POLYMER BLEND MISCIBILITY The question of polymer blend miscibility derives out of the nineteenth century development of thermodynamics and studies in the same periodic of binary mixtures of low molecular weight liquids. From a thermodynamic viewpoint, Gibbs (23) formulated the stability of multiphase systems in terms of the quantity G defined by (in modern notation (24)) G ¼ U þ PV  TS ¼ H  TS

ð2:1Þ

where U is the internal energy, P is the pressure, V is the volume, T is the absolute temperature, and S is the entropy. According to Clausius (25), equilibrium was associated with the maximization of Suniv , the entropy of the universe. Gibbs showed that this was equivalent to G for a system at constant temperature and pressure to reach a minimum. Thus, the formation of a solution requires DG < 0

ð2:2Þ

The quantity, DG, known today as the Gibbs free energy may be expressed DG ¼ DH  TDS

ð2:3Þ

There is an extensive literature on modeling DH, the enthalpy or heat of solution, and DS, the entropy of solution. The simplest situation involves isotropic molecules with no energetic interactions. We may consider an ensemble of systems of equal numbers of component molecules and of energy (microcanonical ensemble). The entropy of mixing involves Boltzmann’s formulation of entropy as (24,26) S ¼ kb ln V

ð2:4Þ

where kb is the Boltzmann constant and V is the number of configuration. Considering N1 the number of isotropic molecules of molecule ‘‘1’’ and N2 the number of isotropic molecules of molecule ‘‘2’’ randomly distributed on a lattice, this leads to V¼

ðN1 þ N2 Þ! N1 !N2 !

ð2:5Þ

On the basis of Stirling’s approximation, as long as N is large the log of factorial, N!, is given by ln N! ¼ N ln N  N

ð2:6Þ

Chapter 2 Miscibility and Characteristics of Polyolefin Blends

31

The entropy of mixing is then DSm ¼ kb ðN1 þ N2 Þ½x1 ln x1 þ x2 ln x2 

ð2:7Þ

where x1 ¼

N1 N2 ; x2 ¼ N1 þ N2 N1 þ N2

ð2:8Þ

The second problem, which needed to be resolved was the heat of solution, DH. For solutions with very similar molecules such as benzene and toluene, or hexane and octane, the heat of mixing is near zero. For various pairs of molecules, for example, acetone and chloroform, significant heat is evolved (DH < 0) and the solution process is said to be exothermic. For other pairs of molecules, heat is absorbed (DH > 0) and the solution process is called endothermic. As early as 1906, van Laar (27) modeled the heat of mixing of a binary solutions using van der Waals equation of state. The views of van Laar were subsequently expanded by van Laar and Lorenz (28), Hildebrand (29,30), and Scatchard (31). The latter authors considered a mixtures of two fluids each with isotropic molecules with interaction energies, c11 and c22 . The interaction energy between the two fluid molecules is c12 . Scatchard (31) obtained the form DUm ¼ Um  U1 x1  U2 x2 ¼ Vx1 x2 ðc11 þ c22  2c12 Þ where V is the mixture volume. He further presumed pffiffiffiffiffiffiffiffiffiffiffiffi c12 ¼ c11 c22

ð2:9Þ

ð2:10Þ

which leads to 1=2

1=2

DUm ¼ Vx1 x2 ðc11  c22 Þ2 Hildebrand (32) defined ‘‘solubility parameters’’ of liquids as   DUi 1=2 1=2 ; i ¼ 1; 2 di ¼ cii ¼ Vi

ð2:11Þ

ð2:12Þ

Thus, Equation 2.11 is equivalent to DUm ¼ Vx1 x2 ðd1  d2 Þ2

ð2:13Þ

Hildebrand (32) also determined experimental values for various liquids. Typical values are listed in Table 2.1 taken from Hildebrand and Scott (33). A major problem with Equation 2.13 is that it leads to the heat of mixing being positive, that is, always endothermic, absorbing heat. Many solutions are exothermic, that is, heat is evolved when the solution is formed. There have been various efforts to model exothermic heats of mixing. These are generally dated to Dolezalek (34) who considered dissolution to represent processes similar to chemical reactions. The modeling of the formation of solutions that we have described assumes that there is no change in volume in mixing. Theories of liquids that consider cell models

32

Polyolefin Blends Table 2.1 Solubility Parameters of Organic Liquids at 25
C. Solubility parameter, (J cm3)1/2

Substance n-Hexane n-Octane n-Hexadecane Cyclohexane Benzene Ethyl benzene Naphthalene Fluorinated hexane (C6F14) Fluorinated octane (C8F18) Fluorinated cyclohexane (C6F12) Fluorinated benzene (C6F6) Dimethyl ether Carbon disulfide Chloroform

7.30 7.55 8.0 8.20 9.15 8.80 9.9 5.6 5.7 6.0 8.1 8.8 10.0 9.3

Adapted from Reference 33.

volume changes begin with Lennard-Jones and Devonshire (35) in 1937. This approach was extended to solutions by Prigogine and Garikian (36) and LonguetHiggins (37) in the 1950s. These developments up to the mid 1950s are described by Prigogine in his book ‘‘The Molecular Theory of Solutions’’ (38). Volume change in mixing can affect the change in DG and induce phase separations which are not predicted from Equations 2.7 and 2.13. Early investigations of the thermodynamics of polymer solutions indicated that solutions exhibited large deviations from ideality. Calorimetric measurement showed that heats of solution were small indicating that these deviations were due to the entropy of mixing. Subsequently, Meyer (39) proposed that the low entropy of mixing was associated with the reduced number of configurations available to polymer chains compared with low molecular weight molecules. More quantitative formulations were subsequently developed by Huggins (40) and notably Flory (41–43). The entropy of mixing DSm was shown to have the form DSm ¼ kb ðN1 þ N2 Þ½x1 ln f1 þ x2 ln f2 

ð2:14Þ

where f1 and f2 are the volume fractions of solvent and polymer. The heat of mixing DHm in polymer solutions that has been expressed by Flory (43) is DHm ¼ xkb TN1 f2

ð2:15Þ

where x characterizes the interaction energy per solution molecule divided by kb T. According to Flory x may be positive or negative. The free energy of solution, DGm , may be expressed DGm ¼ kb TðN1 ln f1 þ N2 ln f2 þ xN1 f2 Þ

ð2:16Þ

Chapter 2 Miscibility and Characteristics of Polyolefin Blends

33

This formulation was subsequently extended to polymer blends by Scott (44) in 1949 as follows:    V f1 f2 ð2:17Þ ln f1 þ ln f2 þ xf1 f2 DGm ¼ kb T M2 V s M1 where M1 and M2 are the numbers of segments of the two polymers, V is the total volume of mixture, and Vs is the volume of segments, which usually are regarded as the volume of repeating units of the polymers. The formulation of Scott (44) does not present the range of phenomena occurring in polymer blends. Various binary blends exhibit lower critical solution temperatures (LCST) where phase separations occur at lower temperature. Other blends exhibit upper critical solution temperatures (UCST) where miscible blends exhibit phase separations at higher temperatures (45). It was shown by McMaster (46) that volume changes occurred in mixing.

2.3 INTERFACES IN LIQUID AND POLYMER MIXTURES The concept of surface tensions between air and liquids seems to date back to Thomas Young (47). Measurements of surface tensions of liquids (48) and its relationship to waves on liquids (49) and the breakup of capillary jets and viscous filaments (50,51) occur throughout the nineteenth century. In thermodynamics point of view, the surface tension is the increase of Gibbs free energy as a result of creating a unit surface area. The surfaces of liquids can be easily understood as an elastic film that inherently tends to shrink to decrease its surface. A summary of surface tensions of various low molecular weight liquids is contained in Table 2.2. It can be seen that the highest values are for molten metals. For organic liquids, the highest values are for polar compounds. Aromatic and aliphatic compounds are similar and the lowest values are for fluorinated compounds. It was believed that the intermolecular interactions play similar role in determining cohesive energy density and surface tension. The surface tension was suggested to be estimated by the cohesive energy density based on the following equation (53)   Vð298KÞ 4 2=3 ð2:18Þ gðTÞ ¼ 0:75  ecoh  VðTÞ where g is the surface tension, ecoh is the cohesive energy density, V is the molar volume, and T is the absolute temperature. However, Equation 2.18 cannot be used to evaluate the polymers containing polar units such as amide, urea, and hydroxyl groups, which form strong hydrogen bonds because the effect of hydrogen bonding on surface tension is generally smaller than its effect on cohesive energy density (54). For high molecular weight materials, the surface tensions of these materials has been found to vary with their molecular weight roughly as (55) g ¼ g1 

Ke M 2=3

ð2:19Þ

34

Polyolefin Blends Table 2.2 Surface Tensions of Low Molecular Weight Compounds. Surface tension, dyn cm1

Material Iron Gold Silver Tin Lead Mercury Water (H2O) Ethylene glycol Epsilon-caprolactam Phthalic acid Lactic acid Styrene Acrylic acid Methyl methacrylate Benzene Cyclohexane Hexafluorobenzene n-Octane n-hexane Octafluoro-2-butene Decafluorobutane

Temperature,
C

1962 1083 882 540 442 318 74.5 50.7 47.1 42.8 39.5 32.6 29.0 26.1 28.8 25.2 22.6 21.5 18.5 12.7 11.8

1200 1200 1200 455 455 455 20 20 20 20 20 20 20 20 20 20 20 20 20 20 20

Adapted from Reference 52.

where g 1 is the surface tension at infinite molecular weight, Ke is a constant, and M is the molecular weight. The surface tension of polymers seems to vary with temperature approximately as (56) g ¼ g 0 ð1  T=Tc Þ11=9

ð2:20Þ

where g 0 is the surface tension at T ¼ 0K and Tc is the critical temperature whose value is around 1000 K for most polymers. The temperature coefficient of surface tension is given by differentiation of Equation 2.20 to be dg=dT ¼ ð11=9Þðg 0 =Tc Þð1  T=Tc Þ2=9

ð2:21Þ

Thus, the coefficient is practically a constant at ordinary temperature, that is, T  Tc ¼ 1000 K. The surface tensions of some polymers are listed in Table 2.3. It came to be realized that interfaces between low molecular weight liquids in immiscible liquid mixtures possess an interfacial tension. Experimental measurements of interfacial tension between low molecular weight liquids were first made in the nineteenth century (59,60). Various experimental techniques have been developed to measure interfacial tensions. A popular method has been the pendant drop measurement, which was originally developed by Andreas et al. (61)

Chapter 2 Miscibility and Characteristics of Polyolefin Blends

35

Table 2.3 Surface Tension (g) of Polymers at Different Temperature in dyn cm1 .

Polymer Linear polyethylene Atactic polypropylene Isotactic polypropylene Isotactic polypropylene

Molecular weight g mol1 Mw ¼ 67; 000 — — Mw ¼ 267; 000

g g (20
C) (140
C)

—

25.5

Polystyrene Polystyrene Poly(ethylene terephthalate)

Mn ¼ 44; 000 Mn ¼ 9290 Mn ¼ 25; 000

40.7 44.6

28.8 22.7 23.1 19.7 (210
C) 19.2 (180
C) 32.1 31.5 36.7

Poly(ethylene oxide) Poly(methyl methacrylate) Poly(e-caprolactam)

M ¼ 6000 Mn ¼ 3000 —

42.9 41.1

33.8 32.0

Poly(hexamethylene adipamide) Bisphenol-A-polycarbonate Poly(dimethyl siloxane) Polytetrafluoroethylene

Mn ¼ 19; 000

46.5

— Mn ¼ 75; 000 M ¼ 1088

42.9 20.9 21.5

30.2 (270
C) 35.7 14.3 13.7

Polybutene-1

35.7 29.4 30.1 30.0

g (180
C)

References

26.5 20.4 20.8 19.1 (220
C) 17.5 (220
C) 29.2 28.8 34.2

(55, 56) (55, 56) (55, 56) (57)

30.7 28.9 36.1 (265
C) 28.3 (300
C) 33.3 12.3 11.1

(55, 56) (55, 56) (55, 56)

(57) (55, 56) (58) (55, 56)

(55, 56) (55, 56) (55, 56) (55, 56)

Adapted from References 55–58.

for measuring surface tension. Another method involves the breakup of molten polymer threads. Interfacial tensions of some polymer pairs are listed in Table 2.4. Measurements of interfacial tensions of polymer melts were reviewed by Wu (55), Koberstein (65), and Demarquette (66). The measurements usually need long equilibrium time because of the high viscosities of polymer melts. The measurements can be divided into two groups: static methods in which interfacial tension is calculated based on the equilibrium profile of the drops and dynamic methods that study the evolution of fiber or drop profiles with time. Static methods include pendant drop method, sessile drop method, and rotating drop method. Dynamic methods include breaking thread method, imbedded fiber method, and deformed drop retraction method. Interfacial tension plays a very important effect in controlling the phase morphology of immiscible polymer blends. In mixing of polymer blends, the minor phase is deformed into long fibrils or thin films by shear or elongation flows, and then is broken into small particles. At the same time, the dispersed particles can collide and merge together, which is called coalescence process. The final morphology is decided by the balance between breakdown and

36

Polyolefin Blends

Table 2.4

Interfacial Tensions of Molten Polymer Blends in dyn cm1 .

Polymer A

Polymer B

Polyethylene Polyethylene Polyethylene

Polystyrene Polysulfone Poly(phenylene sulfone) Poly(ethylene terephthalate) Polycarbonate Polyamide 6 Polyethylene Polystyrene Polycarbonate Polyamide 6

Polyethylene Polyethylene Polyethylene Polypropylene Polypropylene Polypropylene Polypropylene

Interfacial tension, dyn cm1

Temperature,
C

References

4.0 7.0 7.9

290 290 290

(62) (62) (62)

9.4

290

(62)

12.5 12.8 1.63 4.5 8 12.2

290 290 290 290 290 290

(62) (62) (63) (64) (64) (64)

Adapted from References 62–64.

coalescence. Large interfacial tension resists this process in the following three ways. First, larger interfacial tension indicates lower interfacial adhesion, which prevents passing shear stress from one phase to another phase. Second, the interface with large tension is like an elastic film as we mentioned earlier. Inherently, it tends to shrink the deformed phases into spherical shape. Thus, it resists both the shear or elongation deformation of the dispersed phase. Third, large interfacial tension increases the tendency to coalescence. The low interaction between the two phases leads the matrix molecules easy to be excluded when two dispersed particles collide together. Also, large interfacial tension accelerates the merging process of the collided particles.

2.4 POLYOLEFIN–POLYOLEFIN BLENDS There is a history of investigations of blends of polyolefins. Many of these blends were not produced on purpose but were the results of incompletely understood polymerization processes. Examples of these are the various studies of elastomeric polypropylenes, which are mixtures of polypropylenes of varying tacticity levels (67–71). These materials often have interesting technological properties. It is of more concern to consider the results of specially prepared blends of known characterized polymers.

2.4.1 Blends between Polyethylenes Hundreds of polyethylenes are commercially available, and their blends have been used to provide desired range of properties. The blends have high level of mis-

Chapter 2 Miscibility and Characteristics of Polyolefin Blends

37

cibility. As mentioned earlier, LLDPE is a copolymer of ethylene with other olefins, for example, butene, hexene, and octene. Thus, it includes many short branches. LDPE is synthesized by free radical polymerization at high pressure. It includes both short and long chains. Thus, except molecular weight, both branch type and branch content are the two most important molecular structures of polyethylenes. Although the blends of polyethylenes are nearly miscible, their miscibility has been studied extensively. Three methods have mainly been used to study the miscibility. Hill et al. (72–75) mainly used differential scanning calorimetry(DSC) combined with transmission electron microscope (TEM). The melts of the blends are quenched into acetone at freezing point. If DSC shows only one melting peak and there is only one group of crystal morphology based on TEM, the blend is thought miscible in the melt. If DSC shows double melting peak and there are two groups of crystal structures, the blend is thought to be separated in the melt. Hussein et al. (76– 79) used rheology method. They measured both the dynamic viscosity and viscosity at zero shear rate of the blends. If the two parameters of the blends show log-additive linear relationship based on blend content, the blends are considered miscible. If they deviate from the linearity, the blends are thought immiscible. Alamo et al. (80–82) used small-angle neutron scattering (SANS). One component of the blends is deuterated. It was concluded that deuterating can increase the interaction parameter. In some cases, it is enough to lead to phase separation. Tashiro et al. (83–86) argued that they found phase-segregated structures between polyethylene and deuterated polyethylene. These studies suggest the following conclusions regarding the miscibility of polyethylene blends. (1) The branch content is the most important factor to control the miscibility. The miscibility decreases with increase of branch content. (2) The branch type also affects the miscibility. The miscibility decreases with increase of branch length. For example, LLDPE with 4.4 mol% of butene is miscible with high density polyethylene (HDPE) at all compositions (82). However, LLDPE with 4 mol% of hexene is immiscible with HDPE when the HDPE fraction is less than 60 wt% (87), and LLDPE with 2.1 mol% of octene is claimed immiscible with HDPE when HDPE fraction is less than 50 wt% (73). It is hard to deduce the effects of long-branch chains in LDPE. (3) Decreasing molecular weight increases the miscibility. For example, in a study of Hill et al. (72), HDPE with molecular weight 105 g mol1 shows phase separation with the LDPE with 5.2 mol% of branches when the HDPE fraction is lower than 50 wt%. However, HDPE with molecular weight 2553 g mol1 is miscible with LDPE at all HDPE composition. (4) Most of the blends possessed an UCST. Thus, the blends become more miscible with the increase of melt temperatures. (5) Usually, the blends become more miscible with the increase of linear polyethylene content. (6) In summary, the miscibility between polyethylenes is dependent on molecular structures (molecular weight, branch content, and type), temperature, and composition. Thus, phase diagrams based on the composition and the temperature are necessary to characterize the miscibility of specific polyethylene blends.

38

Polyolefin Blends

2.4.2 Blends between Isotactic Polypropylene and Ethylene Propylene Copolymers There have been extensive applications of isotactic polypropylene (iPP)/EPM blends. These were used to produce rubber toughened polypropylene blends and subsequently polyolefin thermoplastic elastomers (88,89). Most commercial EPMs contain more than 50 mol% of ethylene, and these are elastomers. The solubility parameter of EPM should be intermediate to those of polyethylene and polypropylene dependent on ethylene content. Thus, it is often used to compatibilize PE/PP blends (90,91). IPP–EPM blends were generally considered to be immiscible (92–95). Chen et al. (96) have reported miscibility and a LCST in blends of isotactic polypropylene and an ethylene-propylene terpolymer (EPDM). Seki et al. (97) reported that iPP was miscible with EPM with both 19 and 47 mol% of ethylene based on both transmission electron microscopy (TEM) and SANS. However, the molecular weight of the iPP was below 10; 000 g mol1. More recently both Nitta et al. (98) and Kamdar et al. (99) reported that the miscibility between iPP and EPM, whose molecular weight is higher than 100; 000 g mol1 , is strongly dependent on ethylene content in EPM. The miscibility increases with decrease of ethylene content. IPP was miscible with EPM with ethylene content lower than around 17 mol% based on both dynamic mechanical analysis (DMA) and TEM. The EPM domain sizes decrease with decrease of ethylene content as shown in Fig. 2.3.

Figure 2.3 TEM micrographs for the iPP/ethylene–propylene copolymer (80/20) blend films. The propylene contents by mol% in the copolymers are (a) 89.3; (b) 84.3; (c) 76.5; (d) 52.5. (From Reference 98 with permission from Elsevier.)

Chapter 2 Miscibility and Characteristics of Polyolefin Blends

39

2.4.3 Blends between iPP and High Comonomer Concentration Polyethylene Copolymers Since about 1990, ethylene copolymers with substantially high level of olefin comonomers (e.g., butene, hexene, and octene) have become available (100–104). Compared with most commercial EPM, they are available in a pellet form that facilitates continuous processing with twin screw extruders. These copolymers are becoming widely used to substitute for EPM to toughen iPP (105–107). The miscibility between these copolymers and iPP is also strongly dependent on the comonomer content. In a 1996 paper, Yamaguchi et al. (108) of Tosoh studied the miscibility of isotactic polypropylene with ethylene–hexene-1 (EHR) and ethylene–butene-1 (EBR) copolymers. It was found that iPP was miscible with ethylene–hexene-1 copolymers with more than 50 mol% of hexene based on both TEM and DMA (as shown in Figs. 2.4 and 2.5). However, though DMA indicated miscibility between iPP and ethylene– butene-1 copolymers with 56 and 62 mol% of butene-1 (as shown in Fig. 2.4), TEM suggested that there were tiny particles in the blends (as shown in Fig. 2.5). Thomann et al. (109) showed that ethylene–butene-1 copolymers are miscible with iPP when the butene-1 content was higher than 78 mol%. Yamaguchi et al. subsequently investigated the mechanical properties of blends of polypropylene with ethylene–hexene copolymers (110,111). Interestingly

Figure 2.4 Variation of the mechanical loss modulus (E00 ) with temperature for the iPP (O), the copolymers: EHR57 (with 57 mol% of hexene-1); EBR56 (with 56 mol% of butene-1); EBR62 (with 62 mol% of butene-1) (D), their blends (.). (From Reference 108 with permission from John Wiley & Sons, Inc.)

40

Polyolefin Blends

Figure 2.5 TEM micrographs of thin sections of series of blends with 50wt% of iPP. EHR57 has 57 mol% of hexene-1; EBR56 has 56 mol% of butene-1; EBR62 has 62 mol% of butene-1. (From Reference 108 with permission from John Wiley & Sons, Inc.)

Yamaguchi and Miyata (112) found a lack of miscibility of these copolymers with syndiotactic polypropylene (as shown in Fig. 2.6).

2.4.4 Blends between iPP and PB1 Based on the research of Thomann et al. (109), iPP is miscible with ethylene–butene copolymer with 78 mol% of butene. It is an interesting question whether iPP is miscible with PB1. For iPP/PB1 (30/70 by weight) blends made by melt mixing, Cham et al. (113) reported that liquid–liquid demixing happened at 180, 220, and 250
C. The demixing

Chapter 2 Miscibility and Characteristics of Polyolefin Blends

41

Figure 2.6 Transmission electron micrographs for (a) iPP/EHR and (b) sPP/EHR. Hexene-1 content is 57.1 mol%. The samples were stained with RuO4. (From Reference 112 with permission from American Chemical Society.)

process was observed by optical microscope, which indicates that the refractive index difference between iPP and PB1 is large enough for microscopy observation. Bartczak et al. (114) also reported that PB1 particles were found to disperse in the iPP spherulites when the blends were isothermally crystallized at 125
C. These suggest that the blends are immiscible at least up to 250
C. Cham et al. (113) found some evidence of partial miscibility. The demixed blends were isothermally crystallized at 145
C. It was found that the crystals were nucleated in the iPP domains, then, they gradually extended to all the domains. Surprisingly they continued to grow into the continuous PB1 phase with low density. When the crystals growing in PB1 phase reach some iPP domains, they can nucleate the crystal grow in those domain. Since at 145
C PB1 cannot crystallize, the crystals growing in PB1 phase must be PP crystals. This indicates that in the continuous PB1 phase there are iPP molecules, indicating that they are partially miscible. The authors also showed that the iPP crystals growing in PB1 phase have lower melting temperatures than those growing in iPP domains. It was reported that iPP/PB1 blends showed only one tand peak in a DMA test (113,115,116). This was usually taken as a proof for the miscibility between the amorphous iPP and PB1. However, a phase separation was reported at 250
C by Cham et al. (113). They suggested an explanation in their paper. First, the difference of Tg of iPP and PB1 is not large, around 25
C. Second, the separated phases are not pure iPP or PB1 phases, but iPP-rich and PB1-rich phases. Thus, the difference of the Tg of the two phases should be less than 25
C. The small Tg difference may make the two tand peaks merge together. Here, we also want to mention another point based on the study of Hsu and Geil (116), where blends with iPP content lower than 30 wt% show only one peak, but, those with higher iPP content show two peaks (as shown in

42

Polyolefin Blends

Figure 2.7 DMA curves of iPP/PB1 blends. (From Reference 116 with permission from John Wiley & Sons, Inc.)

Fig. 2.7). First, the tand peak of PB1 is much larger than that of iPP. Thus, the iPP peak in the blends with low iPP content is easy to hide, but the PB1 peak in the blends with low PB1 content cannot. It was reported by Hsu et al. (116) and Lee et al. (117) that iPP can nucleate the crystallization of PB1 to increase its crystallization temperature, and PB1, especially in the blends with high content of PB1 more than 70 wt%, can obviously retard the crystallization of iPP. On the basis of the publications available so far, we can make the following conclusions: (1) iPP and PB1 are partially miscible and phase separation occurs at reported temperatures up to 250
C. Melt mixed blends should have a two-phase morphology. (2) The blends made by precipitation from dilute solution show homogeneous mixing. But they are in a metastable state and tend to demix at high temperature.

2.5 BINARY IMMISCIBLE BLENDS There is a large literature involving binary blends of polyolefins, especially polyethylene and isotactic polypropylene, with other polymers. Among the polymers included in these studies are polystyrenes and polyamides. Polyolefins contribute to increase toughness, processability, chemical resistance, and moisture absorption resistance to these polymers. The other polymers contribute high modulus, heat resistance, and oxygen or solvent barrier properties to polyolefins. However, these blends are all immiscible. Useful polymer pairs generally require a compatibilizing agent as described in Section 2.6.

Chapter 2 Miscibility and Characteristics of Polyolefin Blends

43

Figure 2.8 Scanning electron microscopy micrographs of PS/PP (70/30) blend prepared at 190
C.

2.5.1 Polyolefin–Polystyrene Blends Polyolefin–polystyrene blends have long been studied. They are immiscible showing two-phase morphologies. The blends showed poor mechanical properties, especially elongation at break and impact strength, much lower than those predicted based on an additive rule (118). Their fracture surfaces were observed by electron microscope (119,120). As shown in Fig. 2.8, the dispersed phase is easy to be pulled from the matrix and leaves very smooth surface, indicating low interfacial adhesion. Because of the high interfacial tension, the morphology of the blends is not stable. Coalescence readily occurs in the molten state. As suggested by Macosko et al. (121), in melt mixing of immiscible polymer blends, the disperse phase is first stretched into threads and then breaks into droplets, which can coalesce together into larger droplets. The balance of these processes determines the final dispersed particle sizes. With increase of disperse phase fraction (usually more than 5 wt%), the coalescence speed increases and the dispersed phase sizes increase (121–123).

2.5.2 Polyolefin–Polyamide Blends Polyolefin–polyamide melt blends are striking in not only the lack of miscibility, but also the large interfacial tensions between the two melt phases. Investigations of these phenomena in our laboratories (118,124–126) have made numerous studies of these polymer blend systems and found that their phase morphology are quite unstable and trend to coalesce especially under quiescent or low deformation rate conditions. Similar to polyolefin–polystyrene blends, they also show weak interfacial adhesion (118,124,127) (as shown in Fig. 2.9). The mechanical properties of the

44

Polyolefin Blends

Figure 2.9 Scanning electron microscopy micrographs of PA12/PP (70/30) blend prepared at 190
C.

blends are also much lower than those predicted based on the additive rule (125). The disperse phase particles increase in size with increase of concentration (126,128).

2.6 TERNARY BLENDS OF POLYOLEFINS WITH OTHER POLYMERS AND COMPATIBILIZING AGENTS There is a long history of the use of surfactants with a mixture of low molecular weight molecules. These were used as soaps in the preparation of emulsion and other applications. This science and technology were highly developed in the 1920s. The use of high molecular weight surfactants dates back to the 1950s and 1960s and was only studied scientifically in the 1970s.

2.6.1 Surfactants and Compatibilizing Agents The term surfactant has the same meaning with the term surface active agent. Actually it is a blend of ‘‘surface active agent.’’ It has been extensively accepted to substitute the term surface active agent since the early 1960s. Surfactants are usually organic amphipathic compounds with both hydrophobic and hydrophilic groups. They can lower the surface tension of liquids to allow them to spread easier. Their earlier name ‘‘surface active agent’’ was given because of this characteristic. Also they can lower the interfacial tension between two liquids. This characteristic leads them to be applied extensively to make emulsions. Soap is the earliest surfactant. The earliest soap was made by boiling fats together with ashes of plant, which dated as early as 2800 B.C. in ancient Babylon. In 1823, Michel Eugene Chevreul, a French chemist, worked out the structure of fats

Chapter 2 Miscibility and Characteristics of Polyolefin Blends

45

as compounds of organic acids and glycerol. He showed that alkalis converted fats into soaps. He isolated stearic acid and various other organic fatty acids. He sought to commercially develop stearic acid. This was the beginning of the scientific studies of modern soaps and surfactants. From then on more and more surfactants were developed and were used in more applications, for example, detergents, fabric softeners, emulsifiers, paints, and so on. Though surfactants were mainly used to stabilize the immiscible small molecular mixtures, they were also reported in more recent to be used to make immiscible polymer blends more compatible (129–132). They could suppress the coalescence of dispersed particles to decrease their sizes, but could not increase the interfacial adhesion to increase the mechanical properties. The term ‘‘compatibilizing’’ was defined by Gaylord in 1966 (133) as ‘‘rendering a mixture of two or more polymeric materials permanently miscible so as to form a homogeneous composition that has useful plastic properties and which does not separate into its component parts.’’ Actually, in his definition the term ‘‘miscible’’ was not appropriate and another term ‘‘uniform’’ would be better. A few years later, he defined another term ‘‘compatibilization’’ as ‘‘absence of separation or stratification of the components of the polymeric alloy during the expected useful lifetime of the product’’ (134). Different from surfactants, substance considered as compatibilizing agents are today considered not only to decrease the interfacial tension, but also to enhance mechanical behavior presumably by increasing the interfacial adhesion. They must be high molecular weight compounds or form high molecular weight compounds in the blends. They stay on the interface with the part of their molecular chain in one component of the blends and the other part of chain in the other component. Though it was usually thought that the most efficient compatibilizing agents are block copolymers of the two polymers in the blends, however, there are only few block copolymers commercially available. Thus, the most applicable way is to use some commercially available polymers whose surface tensions are between those of the two blend components. For blends between polyolefins and some polar polymers, for example, polyamide, polycarbonate, polyester, reactive group grafted polyolefins are usually used, for example, maleic anhydride grafted polyethylene for polyamide/PE blends. In mixing, the maleic anhydride can react with the amine groups on polyamide to produce polyamide grafted polyolefins, which have good compatibilizing effect. Table 2.5 indicates that compatibilizing agents can decrease the interfacial tensions of some polymer melt pairs greatly.

2.6.2 Polyolefin–Polystyrene Blends with Compatibilizing Agents 2.6.2.1 Polyethylene–Polystyrene Blends with Compatibilizing Agents In case of the first studies using polymer compatibilizing agents, Anderson et al. of Dow Chemical in 1960 (135) irradiated polyethylene with g rays and

46

Polyolefin Blends

Table 2.5 Interfacial Tensions between Polymer Melt Pairs Including Compatibilizing Agents. MAH–PP: Maleic Anhydride Grafted Polypropylene; SEBS: Hydrogenated Triblock Copolymer of Styrene and Butadiene; MAH-g-SEBS: Maleic Anhydride Grafted SEBS; PEMA–Zn: Poly(ethylene-co-methacrylic Acid) Ionomer Neutralized by Zinc. Binary system PE/PA 6

PE/PS

Compatibilizing agent None MAH-g-PP MAH-g-SEBS PEMA-Zn None SEBS

Temperature,
C

Interfacial tension, dyn cm

230 230 230 230 180 180

12.5 2 1.5 1.8 5.8 1.1

Adapted from Reference 118 by Chen & White.

subsequently treated it with styrene to make styrene grafted polyethylene, which increased the tensile strength of PS/LDPE blends. Locke and Paul in 1973 (136) also reported that styrene-grafted polyethylene could improve the tensile strength, modulus, and elongation of LDPE/PS blends and greatly decreased the dispersed particle sizes. Kishimoto et al. in 1970 (137) masticated HDPE/ PS blends with the presence of a free radical initiator. They suggested that the formed block graft copolymers improved the properties of the blends. Carrick in 1970 (138) prepared polyethylene grafted polystyrene copolymers based on Friedel–Crafts (F-C) alkylation reaction of polystyrene. The aluminum chloride (AlCl3) was put into the solution of PE/PS blends in cyclohexane to initiate the reaction. After a period the solution was precipitated in methyl–ethyl ketone (MEK) to extract polystyrene. In the following years, Heikens et al. (139–142) used polyethylene-graft-polystyrene prepared as suggested by Carrick to compatibilize LDPE/PS blends and found that they showed good effect to increase interfacial adhesion and impact strength. In 1980, Sjoerdsma et al. (143) reported that tapered butadiene–styrene block copolymer (PS-(PS-PB)random-PB) could increase the interfacial adhesion of LDPE/ PS blends. In the followed years, hydrogenated butadiene–styrene copolymers (SEBS or SEB) was used as compatibilizing agents for HDPE/PS blends by Lindsey et al. (144), Fayt et al. (145), and for LDPE/PS blends by Fayt et al. (146–149). Fayt et al. (147,148) also reported that tapered hydrogenated butadiene–styrene copolymers (PS-(PS-PB)random-PB) showed better compatibilizing effect than pure hydrogenated butadiene–styrene diblock copolymer in decreasing dispersed phase sizes and increasing the elongation and tensile strength of LDPE/PS blends. Today these butadiene–styrene copolymers, for example, SB, SBS, SEB, SEBS, and so on, have been the major group of compatibilizers for PE/PS blends (150–157), presumably because of their large commercial availability. With 5–10 wt% of these compatibilizers in PE/PS blends, the dispersed phase sizes are decreased, and the elongation at break and impact strength are increased.

Chapter 2 Miscibility and Characteristics of Polyolefin Blends

47

Gao et al. (158) showed that LDPE-g-PS prepared based on Friedel-Crafts alkylation reaction could increase the impact strength, elongation at break and tensile strength of LDPE/PS blends. Wang et al. (159) reported that adding dicumyl peroxide (DCP) into HDPE/PS/SBS blends can obviously increase their impact and tensile strength. The mixing order affected the blend properties very much. However, DCP did not show such effect with HDPE/PS blends. 2.6.2.2 Polypropylene–Polystyrene Blends with Compatibilizing Agents The most commonly used compatibilizers for PP/PS blends are also di- or triblock copolymers of styrene and butadiene (SB and SBS) and their hydrogenated products (SEB and SEBS) (160–164). They form dispersed phases in both pure PP and PS. In PP/PS blends, they locate at the interface to connect both PP and PS phase together. Thus, the interfacial tension is decreased and the dispersed phase sizes are greatly decreased. Polystyrene-b-poly (ethylene-co-propylene) block copolymers were also used to compatibilize PP/PS blends (165,166). They showed similar effect to block copolymers of styrene and butadiene. Fig. 2.10 indicates that they locate in the blend interface. Graft and block copolymers of propylene and styrene have been developed to compatabilize PP/PS blends. Del Giudice et al. (167) and Xu and Lin (168) have synthesized PP-b-PS. Kim et al. (169) and Li et al. (170) first polymerized propylene together with some functional monomers, then polymerized styrene from these monomers units to form polystyrene branches. Diaz et al. (171,172) grafted PP chains onto PS chains based on F-C alkylation reaction when mixing PP/PS blends in the presence of AlCl3 catalyst and styrene. All these copolymers help form very

Figure 2.10 Transmission electron micrograph of iPP/aPS 70/30 blends compatibilized with poly (styrene-b-ethylene-co-propylene) (SEP) and stained with RuO4: 10 wt% of SEP. (From Reference 165 with permission from Elsevier.)

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strong interfacial adhesion to decrease the dispersed phase size and increase the tensile strength.

2.6.3 Polyolefin–Polyamide Blends with Compatibilizing Agents To increase the affinity between polyolefins and polyamides, some polar groups, such as maleic acid or its anhydride, can be grafted onto polyolefin chains. In mixing the polar groups react with polyamides to form polyamide-grafted polyolefins, which are the real compatibilizing agents. The studies about polar group grafted polyolefins were stimulated by trial to increase the dyeability of polyolefin fibers or wettability of polyolefin films, which was earlier than their application to compatibilize polyolefin– polyamide blends. In the 1960s, Nowak and Jones (173–175) of Dow Chemical grafted reactive reagents such as acrylic acid onto melt polyethylene in a screw extruder. In the 1960–1970s, unsaturated acid group grafted polyolefins were seen to be used as compatibilizing agents for polyolefin/polyamide blends first in patents (176–181), then in publications by Ide and Hasegawa (182), Braun and Eisenlohr (183). Since the late 1980s, compatibilizing agent effects on polyolefin/polyamide blends were extensively studied. 2.6.3.1 Polyethylene–Polyamide Blends with Compatibilizing Agents The most important group of compatibilizers for polyethylene/polyamide blends is the free radical initiated graft copolymers of polyethylene, SEBS, EVA, EPDM, and polypropylene with unsaturated acid or its anhydride. They include graft copolymers of polyethylene with maleic acid or its anhydride, glycidyl methacrylate, and ricinoloxazoline maleinate; graft copolymers of SEBS with glycidyl methacrylate, maleic anhydride, and ricinoloxazoline maleinate; graft copolymers of EVA with mercaptoacetic acid and maleic anhydride; graft copolymer of EPDM with maleic anhydride; graft copolymer of polypropylene with maleic anhydride. The acid or anhydride groups of the grafted copolymers can react with the -NH2 groups in polyamide to form polyolefin-graft-polyamide, the real compatibilizers. These graft copolymers have been called precursor compatibilizers. An extensive study comparing various graft copolymers was given by Chen and White (118). Because the free radical initiated graft reaction can also lead to the cross-linking of polyethylene, copolymers of ethylene and with acrylic acid (184,185), glycidyl methacrylate (184,186), methacrylic acid and 10-undecenoic acid (187–189) were synthesized to compatibilize polyethylene/polyamide blends. The poly (ethyleneco-methacrylic acid) ionomers neutralized by sodium (184) and zinc (45,118,190– 192) has also used as compatibilizers. High energy irradiation, used to modify the surface of fibers or films at beginning, was also used to compatibilize the polyethylene/polyamide blends (193–196).

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2.6.3.2 Polypropylene–Polyamide Blends with Compatibilizing Agents The most commonly used compatibilizers for the blends are graft copolymers of polypropylene with maleic anhydride (PP-g-MAH) (197–203), glycidyl methacrylate (PP-g-GMA) (204,205), and acrylic acid (PP-g-AA) (198,206). They can increase the tensile strength and impact strength of the blends. Favis et al. (207– 210) used an ionomer, a copolymer of ethylene and a mixture of methacrylic acid, zinc methacrylate, and isobutylacrylate to compatibilize PP/PA blends and found it effective to decrease the dispersed particles and improve mechanical properties. To further increase the toughness of polyamides and polyamide-polypropylene blends, elastomers grafted with maleic anhydride were used, for example, SEBS-g-MAH (203,211–214), poly(ethylene-co-octene) grafted maleic anhydride (POE-g-MAH) (202,215), EPR-g-MAH (214) and EVA-g-MAH (216). For example, Fig. 2.11 shows the toughening effect of EPR-g-MAH and SEBS-g-MAH on PP/polyamide 6 blends. For PP-rich blends, the notched impact strength of the toughened blends is usually less than 20 kJ m2 , however, for polyamide-rich blends around 100 kJ m2 notched impact strength was obtained. For polypropylene/polyamide/grafted rubber ternary blends, TEM was proved to be a good tool to study their phase morphology. Without grafting, EPR (217) or SEBS (218) formed dispersed particles together with the minor component of the blends separately; however, after grafting with maleic anhydride they could encapsulate the dispersed minor component of the blends (214,217–221).

Figure 2.11 Izod impact strength of polyamide 6/polypropylene blends modified with 20% EPR-gMA or SEBS-g-MA. The composition shown on the abscissa is the percentage of polypropylene in the blend on a rubber-free basis. The dashed lines correspond to the properties of unmodified binary blends. (From Reference 214 with permission from Elsevier.)

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2.7 CONCLUSIONS The miscibility between polyolefins is dependent on molecular structure, blend composition, and mixing temperature. Phase map is needed to characterize their miscibility. The binary blends between polyolefins and polar polymers are immiscible and have large interfacial tension. They show weak interfacial adhesion, poor mechanical properties, and trend to coalesce in the melt state. Compatibilizing agents are needed in these binary blends, which can decrease interfacial tension, prevent phase coalescence, and improve properties largely. Compatibilizing agents can be grafted or biblock copolymers of the two blend components, or can be prepared by functionalizing one blend component, which can graft onto the other component to form grafted copolymers through the functional groups in mixing. To be effective, these compatibilizing agents must stay at phase interface. Both their molecular structures and melt mixing procedure can affect their diffusion behavior to the phase interface.

NOMENCLATURE G U DUm P V Vs T Tc S DS DSm Suniv H DHm kb V N N1 N2 x1 x2 c11 c22 c12 d d1

Gibbs free energy Internal energy Internal energy change in mixing Pressure Volume Volume of segments of polymers Absolute temperature Critical temperature Entropy Entropy change Entropy change of mixing Entropy of the universe Enthalpy Enthalpy change in mixing Boltzmann constant Number of configuration Number of molecule Number of molecule 1 Number of molecule 2 Mole fraction of molecule 1 Mole fraction of molecule 2 Interaction energy between molecule 1 Interaction energy between molecule 2 Interaction energy between molecules 1 and 2 Solubility parameter Solubility parameter of molecule 1

Chapter 2 Miscibility and Characteristics of Polyolefin Blends

d2 f1 f2 x M1 M2 g g0 g1 ecoh M Ke

51

Solubility parameter of molecule 2 Volume fraction of molecule 1 Volume fraction of molecule 2 Interaction parameter Number of segments of molecule 1 Number of segments of molecule 2 Surface tension Surface tension at zero absolute temperature Surface tension of molecules with infinite molecular weight Cohesive energy density Molecular weight Molecular weight–surface tension constant

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Part II

Polyolefin/Polyolefin Blends

Chapter

3

Miscibility, Morphology, and Properties of Polyethylene Blends Robert A. Shanks1

3.1 INTRODUCTION Polyethylenes (PE) are a group of polymers having the same chemical structure, but differing in molecular architecture. They have subambient glass transition temperatures so the amorphous phase is rubbery. Several additional subambient relaxation temperatures have been detected depending on the branching characteristics. Molecular flexibility contributes yield, drawing, and toughness. Strength and moderate thermal resistance are derived from crystallinity. Weak intermolecular forces means that interactions and crystallinity result from symmetry and close packing. Branching restricts symmetry and close packing. Branch points are excluded from the crystals. Polyethylenes have many applications and through their price and performance they rank as commodity polymers. They are used in films, sheets, fibers, pipes, and many types of moldings. Polyethylene products are prepared by extrusion, injection molding, thermoforming, and rotational casting. A particular polyethylene must be chosen for the molding process, product properties, and longer term product performance. Even with the large range of polyethylenes available, it is difficult to choose a polyethylene with the best properties. Blending is usually the solution to property optimization. More polyethylenes are manufactured as blends than they are as the pure polymers. Most polyethylene blends are with other polyethylenes, though some blends are formed with ethylene copolymers with polar monomers such as vinyl acetate (VAc), methyl acrylate (MA), butyl acrylate (BA), glycidyl methacrylate

1

School of Applied Sciences, RMIT University, GPO Box 2476V, Melbourne, Vic 3001, Australia

Polyolefin Blends, Edited by Domasius Nwabunma and Thein Kyu Copyright # 2008 John Wiley & Sons, Inc.

59

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Polyolefin Blends

(GMA), acrylic acid (AA), or grafted with maleic anhydride (MAn). Hydrocarbon waxes, sometimes highly branched, are other compounds that may be blended with polyethylenes. Polyethylene in its ideal form is a linear polymer (LPE) that should correctly be called poly(methylene), Table 3.1. LPE can reach high crystallinity (Xc > 0:90) due to its compact symmetrical molecules. It has low strength dispersive molecular forces. The melting temperature (Tm ) of LPE is variable but usually about 135
C. The equilibrium melting temperature of LPE has been measured many times and is about 145
C. LPE is insoluble at ambient temperatures, but it can be dissolved in hydrocarbon, alkylaromatics, and halogenated hydrocarbons at temperatures approaching its melting temperature. Two crystalline morphologies have been observed; the chain folded form is most common and forms spontaneously from

Table 3.1

Properties of Linear Polyethylene.

Property

Value

Comment

Monomer

 CH2 CH2  

Polymerization

Ziegler–Natta, supported metal oxides such as Philips, Unipol; and metallocene 120 to 60
C 137
C

Symmetrical and compact for close packing, dispersive forces; no other intermolecular interactions, polarity, or bulky steric groups Stereospecific chain formation required; weak interaction/complexation with any initiators/catalysts; metal coordination complexes required

Glass transition (Tg ) Melting temperature (Tm )

Crystallinity

0.60–0.95

Molar mass

Wide range, 103–107 g mol1

Density

0.90–0.96 g cm3

Stability

No chemical groups to absorb ultraviolet light and no reactive tertiary hydrogens

An elastomer or viscous liquid at ambient folded chain crystals; strength properties derived from crystallinity; low thermal resistance. Extended chain crystals can melt at about 145
C Properties dependent on crystallinity, with amorphous phase providing flow, flexibility, and toughness. Kinetically stable folded chain crystals; thermodynamically stable extended chain crystals can be formed by gel drawing Molar mass must be high since intermolecular forces are weak; all values are available from waxes to ultrahigh molar mass. Dependant on crystallinity; often HDPE contains some butene comonomer to avoid brittleness The ideal polyethylene structure should be very stable; however, some unsaturated end groups, impurity oxygen containing groups, and catalyst residues accelerate degradation

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61

Figure 3.1 Density, crystallinity, and polyethylene type scale.

the melt, being kinetically favored. The chain extended form is thermodynamically favored, but can be formed by only flow elongation of molecules, typically by gel spinning. Significant characteristics of LPE are molar mass (M) and molar mass distribution. Molar mass and morphology are the only variables available to provide required properties with LPE. LPE is usually called high density polyethylene (HDPE) because of the high density arising from high crystallinity. Commercial HDPE often contains a small proportion of comonomer, typically 1-butene, to decrease its brittleness. The comonomer will lower the density, crystallinity, and melting temperature; Tm may be in the range 130–135
C. A schematic showing polyethylene type, density, melting temperature, and spherulitic crystal ranges is shown in Fig. 3.1. Branching of polyethylenes provides the second dimension, after molar mass, with which to control properties, Tables 3.2 and 3.3. Branching of polyethylenes is a complex topic; in this review it will be treated starting with the ideal copolymer structures formed with the new metallocene catalysts. Branched polyethylenes (BPE) provide increased toughness though decreased modulus and strength compared with LPE. Branches are obtained by copolymerization with 1-alkenes, such as Table 3.2 Properties of Typical Branched Polyethylenes. Properties

VLDPE-B VLDPE-O

Comonomer Butene Catalyst type M MFI (dg min1) 27.0 0.901 Density (g cm3) Mw 58,000 Mw/Mn 2.65 Comonomer 6.3 Content (%mol) Tm (
C) 92.7 Tc(
C) 76.6

VLDPE-O(ZN)

LDPE-highM LDPE-lowM

Octene S 1.0 0.908 96,700 2.86 2.4

Octene ZN 1.0 0.912 120,000 3.80 4.2

P 7.0 0.919 474,000 23.3

P 22.0 0.918 89,000 4.4

105.4 90.3

123.0 100.5

106.5 87.8

103.9 87.1

Data were obtained from data sheets published by the manufacturer. ZN ¼ Ziegler–Natta catalyst, S ¼ Constrained geometry single-site catalyst, M ¼ Metallocene, and P ¼ Peroxide. Crystallisation (Tc) and melting (Tm) temperature were obtained using DSC at scanning rates of 10
C min1.

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Polyolefin Blends

Table 3.3

Types of Polyethylenes and Their Properties.

Polyethylene

Structure

Comment

HDPE or linear polyethylene, poly(methylene) LDPE

Linear carbon chain

See Table 3.1

Random branching with either short (average C4) or long (many C) branches; dependant on autoclave or tubular reactor type

LLDPE

Short branches introduced by a comonomer, usually butene, hexene, or octane; with many others; comonomer amount and structure limited with ZN catalysts

mLLDPE

Metallocene-catalyzed PE with more uniform short branch distribution; increased monomer content and structure available Increased comonomer content giving lower density and lower Tm. ZN and metallocene type available Highest comonomer contents, lowest Tm ; crystallization can continue below ambient; can be dissolved at ambient

Many short branches limit crystallinity and lamella thickness (reduce Tm ); long branches provide nonNewtonian rheology and melt strength Short branch content is controlled by comonomer amount, distribution of comonomer not controlled. Slurry (or solution) and gas phase variations. Bimodal crystal distributions and possibly two liquid phases Structure and property combinations highly specialized; new applications possible with a wide range of molar mass and branching Soft PE that can be blended with stronger PE to improve toughness, heat sealing, and elasticity Elastomeric properties that require cross-linking if used alone. Often used in blends

VLDPE

ULDPE

1-butene, 1-hexene, and 1-octene. Polar polyethylenes are formed by copolymerization with VAc, MA, BA, GMA, or AA. Grafting of MAN can be performed in a solution reaction or in an extruder. The properties of BPE are determined by the proportion of branches, or more specifically the methylene sequence length (MSL) between branches, in addition to molar mass and its distribution. When the comonomers are unevenly distributed along polyethylene molecules, MSL will be a distribution or this can be expressed and a comonomer distribution. A two-dimensional contour chart must define the architecture of BPE, with molar mass distribution and comonomer distribution axes. The comonomer distribution can be intermolecular, such that some molecules have mainly linear structure while others have many branches. The comonomer distribution may be intramolecular, where the density of comonomer units changes along the length of a molecule, so that segments

Chapter 3 Miscibility, Morphology, and Properties of Polyethylene Blends

63

Figure 3.2 PE branch distributions, intermolecular branch distribution (top and middle), and intramolecular branch distribution (bottom).

of a molecule may be mainly linear, while other segments have many comonomers or branches (1). Inter- and intramolecular branch distributions are illustrated in Fig. 3.2. The morphology is independent of branch length, that is, comonomer structure, since the branches are excluded from crystals. BPE polymerized by Ziegler–Natta, Philips, Unipol, or similar multisite catalysts show increased dispersion of comonomers. The distributions are typically bimodal. Gas phase and solution polymerizations differ in distribution, with solution or slurry prepared BPE generally showing broader distributions. Bimodal distributions of branches result in bimodal distributions of lamella thickness and hence melting temperatures. Bimodal BPE behave like blends though there is poor control over the structure and composition of each component. The amount of braches that can be included is limited to low amounts compared with single-site-catalyzed BPE where highly branched elastomeric BPE can be produced. There is controversy as to whether the components are miscible in the liquid phase and if they become immiscible at what difference in branch composition and at what temperature. The phase diagram is proposed to be of an upper critical solution temperature type. The BPE described are prepared by copolymerization and only short branches are present. PBE with both short and long branches are prepared from radical polymerization in the traditional high pressure process. These PE are called low density polyethylene (LDPE). A molecular model of a conformation of a typical LDPE is shown in Fig. 3.3. LDPE with different proportions of short and long branches are formed using autoclave or tubular reactors. There are many short branches with average

Figure 3.3 Molecular model of typical LDPE with 1000 carbons, 3 long branches, and 30 short branches.

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Polyolefin Blends

length of four carbons formed by intramolecular chain transfer by the growing chain. There are relatively few long branches, about 10% of the short branches, formed by intermolecular chain transfer by a growing chain to a terminated chain. The special properties of LDPE compared with other PBE are brought about by the long branches. Long branches contribute to melt rheology during processing by providing increased shear thinning and melt strength. The branch distribution in LDPE is random because of the mechanism of formation, and the overall distribution demonstrated by melting is narrow compared with most multisite-catalyzed BPE. Techniques are now available for polymerization with single-site catalysts that allow long branches to be included by copolymerization with alkene-terminated BPE formed during the same polymerization reaction. Alternatively, the long branches can be added to any PE composition by blending with LDPE.

3.2 STRUCTURE AND PROPERTIES OF POLYETHYLENES The important factors (in order of most to least) in controlling the properties of polyethylenes are as follows: Proportion of branches (i.e., the composition of the comonomer). Each comonomer unit in the polymer is unable to crystallize, so this breaks the ability of the polyethylene chain to take part in crystallization. The distribution of the comonomer units in the chain (are they distributed evenly or are they clustered within some molecules or within parts of some molecules). This is determined by the type of catalyst—metallocenes give relatively even distribution of branches. Ziegler–Natta catalysts give broad distribution. Bimodal distribution of branches, similar to the above two, except that the melting shows two endotherms. Usually a broad lower temperature melting and a sharp higher temperature melting, the latter for crystals of molecules with few branches. This is caused by uneven reactivity giving two phases in the reactor. Type of branches—1-butene comonomer gives ethyl branches, 1-hexene comonomer gives butyl branches, and 1-octene comonomer gives hexyl branches. The length of branches has a small influence, but the number of branches is more important since the branches do not take part in crystallization. Molar mass (inversely proportional to melt flow index, MFI) has a small influence. Generally in the chain folded crystal structure, the molar mass will not be important. It will mean that a particular molecule has more folds in the crystal. Melting is determined by the crystal thickness. Crystal thickness is determined by crystallization conditions and the presence of branches that sterically prevent crystallization of some segments. Often molecules with different numbers of branches will be immiscible in the liquid state (melt). So a polymer with a broad branch distribution may have two

Chapter 3 Miscibility, Morphology, and Properties of Polyethylene Blends

65

liquid phases and so crystallize with two crystal populations. This is another manifestation of the above three, within the one polymer. Blends of polyethylenes with much different branch contents will also show two immiscible liquid phases.

3.3 APPLICATIONS OF POLYETHYLENE BLENDS Rheological control is a primary reason for the preparation of PE blends. The PE chosen to provide optimum properties for a particular application may not have desired processing characteristics, Table 3.4. Ideally shear thinning is preferred so that high molar mass PE required for a product can be processed under moderate conditions. Even a small amount of BPE with significant shear thinning can impart shear-thinning properties when used in a blend. Branching increases the elasticity, elongation at break, and toughness of PE films, though modulus, yield stress, and break stress are decreased. Heat sealing of films is enhanced by BPE since branching reduces the melting temperature. PE films are usually produced by the blown film extrusion method. The alternative method is sheet extrusion followed by passing through chilled rollers. Blown film is more economic; biaxial orientation is introduced by the draw-off and blow ratios. Bubble stability is critical to the blown film process. Bubble stability is provided by the melt strength and rheological characteristics of the PE. Long

Table 3.4 Some Examples of Typical Polyethylene Blends and Their Applications. Polymer 1

Polymer 2

HDPE

LDPE

HDPE

LLDPE

LLDPE mLLDPE UHMWPE

HDPE

Application

Toughened HDPE with improved processing, especially melt strength; HDPE and LDPE are immiscible LLDPE Pipe grade HDPE required rheology for (or LDPE) extrusion; rigidity and impact resistance for pipes LDPE Typical film combination for increased melt strength, elasticity combined with LLDPE strength. These are the most widely used PE blends mLLDPE or Packaging films with improved melting and mVLDPE flow character for heat seal LDPE Introduction of long branches to assist processing into specialized metallocene PE HDPE Lower molar mass (HDPE) to moderate the properties of ultrahigh molar mass PE for processing and flexibility HDPE Bimodal molar mass distributions provide (‘‘reactor blends’’) unique properties that cannot be achieved by mixing the two PE grades

66

Polyolefin Blends

branches and high molar mass increase melt strength and so LDPE is preferred for blown film manufacture, but other properties may be required of the film after formation. The solution to optimization of production and application properties is to use blends of PE to contribute their unique characteristics. Such an important blend use for films is LLDPE–LDPE. Better-defined properties can be obtained using particular BPE formed by single-site catalysts. These BPE are more expensive, so a blend with lower cost PE will often provide the desired contribution to properties. Often combinations of properties are required and these are more applicable to a PE blend. A combination of high melt strength and shear-thinning rheology requires long branches, with high film strength from low branch content, to allow thin films to be made. Heat sealable films require a low temperature melting component, together with a strength-providing component. PEs are being increasingly used for cable coverings, either for large cables or for flexible appliance coverings. Many competing requirements are expected of these cable coverings, most specified by regulations or standards. The inherent properties of the PE are often supplemented by peroxide or grafted silane-induced cross-linking. Metallocene PE have suitable properties for cable covering, but blends are needed to satisfy all specifications. Plasticised poly(vinyl chloride) has been used so a PE covering will often need to possess similar properties to be considered as a substitute. Extruded pipes need to be rigid yet tough so a significant component is likely to be HDPE, yet toughness must be included and flexibility to enable coiling. Toughness of HDPE can be improved by inclusion of some comonomer, though a blend with a PE plastomer or elastomer will produce dispersed phase elastomer that is know to provide efficient toughness. Other extruded profiles for building components will require similar properties. PE moldings such as garbage bins need a combination of stiffness, strength, and toughness. In addition high melt flow is needed to completely fill the mold before solidification. HDPE is the base polymer, but other blended polymer such as a PE elastomer will be required to increase the toughness while retaining strength and dimensional stability, particularly when the bin is used to carry a load. PE is used for large rotationally molded tanks. The tanks must be strong, stiff, and tough as for the other molded products. Rotomolding is a special molding technique where flow of the molten polymer is required under a small centrifugal or gravitational force and the walls must be void free. The slow molding process usually does not produce stresses within the walls. Low molar mass could provide suitable melt flow, but the tank would have insufficient strength so the PE must be a blend to obtain the individual distinct properties. See Table 3.5 for a summary of processing methods, the PE and product requirements.

3.4 MOLAR MASS AND BRANCHING DISTRIBUTIONS Metallocene catalysts provide PE with control of molar mass and a narrower molar mass distribution than other catalysts. These PE will have precise control of

Chapter 3 Miscibility, Morphology, and Properties of Polyethylene Blends

67

Table 3.5 The Main Processing Methods for Polyethylenes and Their Requirements. Processing method

Polyethylene

Comment

Extrusion

All types

Cast-film extrusion

LLDPE, HDPE

Laminate-film extrusion

LLDPE, LDPE

Coextrusion

VLDPE–mPE

Blown-film extrusion

VLDPE–LDPE blends

Injection molding

HDPE

Vacuum forming

HDPE, VLDPE

Rotational casting

LDPE, LLDPE

Powder coating

LDPE, LLDPE

Extrusion is the main process for compounding with additives, fillers, pigments, blended polymers, and then forming pellets for subsequent processing. Further processing may form various profiles such as pipes, beading, wire, and cable coating Film is extruded through a slit-die then passed through chilled rollers to crystallize and smooth the surface; orientation by cold drawing PE can be extruded onto other substrates, paper, and paperboard mainly; though other polymer films and metal foils are used Two or more polymer films are extruded together for combination properties not available by blending; strength and heat-seal layers The tubular extrudate in expanded into a thinner film; melt strength to resist bursting is needed Molding of articles of many sizes, from rubbish bins to small fittings requires melt high flow and lubrication Partially melted sheet is rapidly shaped with high extensional flows and critical edge strengths Melting, flow, and particle coalescence are required; particularly flow under little more than gravity, though log times are involved Powder must adhere to hot objects, then flow under adsorption and surface tension forces

properties associated with molar mass, but this may not be effective in providing a broad range of properties. A broad molar mass distribution will combine good processing with sufficient strength, but with less precise control. Multisite catalysts give a broad molar mass distribution, but the molar mass can be controlled to give a range of grades for each polymer. Radical polymerization to manufacture LDPE gives the least control of molar mass and distributions or offer very broad due to the many long branches.

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Polyolefin Blends

Bimodal molar mass distributions have become of increasing interest due to their combination of processability and mechanical properties. Combinations of toughness, tensile strength, and ultimate strain can be found in bimodal PE. Bimodal HDPE can retain modulus and yield strength, while providing some flexibility and toughness. The alternative of blending with a BPE often decreases modulus and tensile strength too much for applications such as extruded pipe. Resistance to creep is an important property retained in bimodal HDPE. The bimodal PE are usually not blends of differing molar mass, they are created in a polymerization reactor. Bimodal LLDPE are available where the high molar mass component contributes to tensile strength and creep resistance required in films, while the low molar mass facilitates processing and flexibility of the film. Branch distribution is the orthogonal distribution to molar mass distribution. Branching determines crystallinity and melting temperature range, so the mechanical properties are more dependent on branch distribution than molar mass distribution. As described above, catalyst technology has mainly revolutionized the copolymerization reaction. The single-site or metallocene catalysts produce a more uniform structure, in both molar mass and branch distribution. These catalysts are less selective of monomer, and the branch distribution becomes more statistical compared with multisite catalysts such as Zeigler–Natta. The distribution of short branches approaches that of LDPE, except that there are normally no long branches. Long branches can be introduced through single-site copolymerization. This is done by maximizing termination by disproportionation so that unsaturation is formed at chain ends. The end unsaturated chains then become macromonomers for further copolymerization with ethylene and the respective 1-alkene (2). The long branches are desired for their contribution to shear thinning and melt strength during processing. These rheological properties assist blown film production. Many LLDPE are compositionally heterogeneous, since they contain molecules with few branches and molecules that are highly branched. The more highly branched fraction can be separated by solvent extraction. Small-angle neutron scattering (SANS) has confirmed the prediction that structurally disperse LLDPE display liquid–liquid phase separation, even if the overall branch content is low. These LLDPE behave like polyethylene blends even though they are the result of a single polymerization. PE with high levels of 1-alkenes can be produced using single-site catalysts. The near statistical highly branched PEs are elastomeric. They have similar compositions to EP and EPDM, but they are formed from comonomers such as 1-butene, and their comonomer distribution is more uniform. They have processing characteristics similar to thermoplastics, as opposed to rubbers.

3.5 CRYSTALLIZATION, MELTING, AND BRANCHING OF POLYETHYLENES The melting temperature of PE varies over a broad range with the equilibrium melting temperature as the upper limit for LPE. The melting temperature varies with lamella

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thickness that in turn varies with crystallization conditions. Slow crystallization will give larger lamella thickness. Slow crystallization will occur at higher isothermal conditions or upon slow cooling. As the crystallization temperature (Tc ) approaches the melting temperature, the melting temperature approaches the equilibrium melting temperature (Tm
). Branching produces a new limiting temperature called the copolymer melting temperature (Tmc ). Tmc still depends on crystallisation conditions and represents the limiting case for Tc ¼ Tmc ; however, Tmc depends on the mole fraction of branches, or more specifically the noncrystallizable component. Branches suppress lamella morphology due to their causing variable crystallizable sequences (3). Branches are excluded from the lamella, thus the lamella thickness is limited by the length of ethylene segments, the methylene sequence length (MSL). Methyl and sometimes ethyl branches can be included in lamella if crystallization is fast; the resulting crystals are less well formed and the melting temperature is further reduced. Linear polyethylene (LPE or HDPE) has been blended with LLDPE with varying branch distributions. The critical branch content for phase separation was lower for mPE, where the branches were more evenly distributed, than for Ziegler– Natta LLDPE, where the branches are heterogeneously distributed. Where cocrystallization occurred, the unit cell of the crystals showed marked expansion (4). Crystallization conditions were shown to be important for the cocrystallization of blends of linear and branched polyethylenes, low isothermal temperatures promote cocrystallization as did quench cooling (5). Figure 3.4 shows the crystallization (on cooling at 10
C min1) followed by melting (on heating at 10
C min1) of ZN–octene–LLDPE. This slurry or solutionpolymerized LLDPE shows considerable molecular or branching heterogeneity according to the broad bimodal melting temperature range. There is a sharp higher

Figure 3.4 Crystallization and melting of ZN–octene–LLDPE after cooling and then heating at 10
C min1 using a Perkin-Elmer Pyris 1 DSC.

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Polyolefin Blends

Figure 3.5 DSC melting curves for HDPE, LDPE, and their 1:1 blend.

melting double peak at 125
C, where the double peak is attributed to melting– recrystallization–melting during the DSC scan. Crystallization similarly shows a sharp higher temperature exotherm followed by a broad large exotherm (not both exotherms and endotherms are positive since the heat flow data have been converted to apparent specific heat by baseline corrections and calibration). DSC melting scans of blends of polyethylenes reveal changes in all aspects of crystallinity as shown by the melting peaks, though some peaks of the original components are retained in modified form. This indicates different crystal populations that may arise from liquid–liquid immiscibility or fractionation of a singlephase liquid during crystallization. Figure 3.5 shows melting of HDPE and LDPE and their blend. The HDPE in the blend melts similarly to the pure HDPE but at a slightly lower temperature. The melting of LDPE in the blend occurs at the same temperatures as the pure LDPE, but the main peak at 110
C is much reduced indicating that LDPE is miscible with the HDPE in the amorphous phase and that crystallinity of the blend is much reduced. Figure 3.6 shows the melting of a blend of LDPE–VLDPE and the melting of the pure components. The individual peak melting temperatures are retained in the blend, but the distribution of crystals melting under each peak has changed. This indicates miscibility of the components with similar branching distributions to provide a new branching distribution in the liquid phase, which may exist as a phase separated liquid phase. A small low temperature endotherm shifted from about 90
C in pure LDPE, to 80
C in the blend. The LDPE–ULDPE blend shown in Fig. 3.7 provides the greatest difference between the PE in these examples. ULDPE melts between 30 and 75
C because it is highly branched; such PE are called thermoplastic polyolefin elastomers (TPO). ULDPE and LDPE in the blend seem to mainly melt independently, though with overall decreased crystallinity. The main peak shift is the LDPE lower melting temperature peak shifting from 90
C to about 80
C. The melting endotherms of LDPE and

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Figure 3.6 DSC melting curves for LDPE, VLDPE, and their 1:1 blend.

ULDPE do not overlap, so it can be assumed that the branching distributions do not overlap making it unlikely that any components from each polymer would be dissolved in the other. Within a BPE with a distribution of branches, there will be a distribution of lamella thickness. This will result in a broad melting range. BPE with a bimodal distribution of branches will have a bimodal distribution of lamella thickness and a corresponding melting temperature range. When two polyethylenes are blended, assuming they are miscible, they will cocrystallize only where they have common MSL. Some molecular segments in each BPE will crystallize independently of

Figure 3.7 DSC melting curves for LDPE, ULDPE, and their 1:1 blend.

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Polyolefin Blends

the other BPE in the blend. Note that it is the segments that enter lamella, not whole molecules. Where there are intramolecular branch distributions, then segments of the one molecular can join different lamella as appropriate to their MSL. Thermal fractionation techniques have been used to separate crystals with differing lamella thickness formed from molecular segments with differing methylene sequence lengths. Stepwise isothermal cooling (SC) has been used to allow each fraction to crystallize according to the temperature appropriate to each methylene sequence length, where longer sequences can crystallize at higher temperatures (6–8). Another technique that provides better resolution of the fractions is successive self-nucleation and annealing (SSA) (9). SSA has been used to fractionate ethylene–hexene copolymers to distinguish the contribution of catalyst on the branch distribution and structure (10).

3.6 MISCIBILITY AND CRYSTALLIZATION Polyethylenes have a broad distribution of melting temperatures resulting from varying lamella thickness. The lamella thickness is determined by the fold length, which is limited by MSL. At each crystallization temperatures, only those molecules with sufficient MSL can crystallize. If a blend is homogeneous, then crystallization will proceed, which is dependent on MSL and temperature, irrespective of the component from where the molecules originated. The overall distribution of lamella thickness will be a combination of the lamella thicknesses if each component were crystallized separately. Difference in the distribution will depend on the conditions of crystallization, such as rate of cooling. This situation will provide a broadened and probably bimodal melting endotherm. In this case phase separation does not occur, but fractionation according to MSL occurs for the combined blend components. If the PEs in a blend are immiscible at the temperature of crystallization, then each phase must crystallize independently and there will normally be a disperse phase and a continuous phase, except for nearly equal volume fractions where the phases could be cocontinuous. Though the phases crystallize independently, they will not crystallize at the same temperatures, the crystallization temperatures will depend on branch density. This situation will provide a bimodal melting endotherm and a double crystalline morphology. Note that the two crystalline phases will not have the composition of the original PE. The composition of the phases will be according to the tie-lines of the phase diagram (Fig. 3.8). As the blend is cooled and the mutual solubility changes, the composition of the two phases will change. Descriptively, the more linear PE will dissolve the more linear molecules from the more branched PE of the blend. Conversely, the more branched PE will dissolve the more branched molecules from the less branched PE. In this case phase separation occurs before crystallization and continues during crystallization. The phase equilibria correspond to a typical U-shaped upper critical solution temperature diagram. The competition between liquid–liquid phase separation and crystallization from homogeneous melt has been observed for

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Figure 3.8 Temperature–composition phase diagram for an immiscible polyethylene blend with differing branching.

blends of ethylene copolymers with hexene and butene, using time resolved SAXS and WAXS, complemented by optical and atomic force microscopies (11). Similar conclusions were made concerning blends of ethylene copolymers of butene, hexene, and octane (12–14). In both miscible and immiscible blends, the resulting crystalline morphology will be complex and there will be a broad melting temperature range. Properties will be significantly modified by the blend morphology and events such as impact; elongation or creep conditions will differ. Only a small amount of a modifying PE may be necessary to provide significant changes. If miscibility is required then the two polyethylenes must be similar, but different enough to change morphology in the blend. Blends of LLDPE with LDPE with only small differences in branch composition (3% and 5%, respectively) were found to be immiscible in the liquid (15). Blends of HDPE with long-chain branched polyethylenes (HBPE) prepared from metallocene catalysts have been studied by DSC and their crystal structures interpreted in terms of phase behavior. The HBPE contained long-chain branches and short branches formed form octane comonomer. HBPE with 7.5–12.0% octane exhibited phase separation, whereas HBPE with 2% octane were found to be miscible with HDPE over the whole composition range. Long branches were few and did not contribute to the immiscibility (16). Crystallization and melting of ternary blends containing mPE, LLDPE, and LDPE have been studied where the mPE and LLDPE varied in content and had the same melting temperature, Tm ¼ 122
C and a fixed content of LDPE, Tm ¼ 114
C. Crystallinity increased with mPE content, and it was proposed that mixed crystals may be formed since no separated melting peaks were observed for the component PE (17). Blends of LPE and several different poly(ethylene-cohexene)s were shown,

74

Polyolefin Blends

by DSC measurements of melting, to be immiscible even with only 1.8 mol % hexene (18). SSA has been used to determine the miscibility of linear and branched polyethylenes and the results showed that only those PE fractions that were similar in branch content and distribution of short branches were miscible in the melt. The SSA thermal fractionation helped to distinguish miscibility effects from cocrystallization effects (19). Phase diagrams have been proposed for blends of polyethylenes where the first component is linear, or less branched, and the second component is more branched. The method involves quench cooling each blend composition from the melt at various temperatures, so that there is insufficient time for liquid–liquid phase separation. Separated blends are considered to be separated at the particular temperature of the melt prior to quenching. TEM and DSC are used to characterize the quenched blends. The phase diagrams exhibit UCST behavior, depending on the difference in branching content of the component polyethylenes. Branch length has been found not to be important since it is the branch points that are excluded from the crystals (20,21).

3.7 THEORETICAL PREDICTION OF MISCIBILITY The miscibility of polyolefins in the melt is difficult to predict. Their miscibility or immiscibility is in contrast to the chemical similarity of their structural groups: methyl, methylene, and methyne. Polyolefin melts have similar structures, including density–temperature relationships and optical characteristics such as the refractive index, making most characterization techniques inappropriate for the detection of phase separation. The general rule of like dissolves like being the determinant of miscibility is inadequate for polyolefin miscibility. Some blends of polyethylenes are considered to phase-separate the above crystallization temperatures of the components according to an upper critical solution temperature spinodal phase diagram (Fig. 3.8). Small-angle X-ray scattering and hot-stage optical microscopy have not provided definitive evidence of phase separation. Crystallization of polyethylene from solution has been shown to follow selfnucleation to form gels where the network links are formed by the crystals. Crystallization occurs at a temperature-dependent critical gelation concentration. The lamella may join edge-to-edge when near coplanar (22). Crystallization from miscible polyethylene blends is expected to be consistent with solution gelation studies. Two processes are predicted to occur on solidification of polyolefin blends. Crystallization from a homogeneous melt of the blend where the composition of the crystalline phase will initially contain molecules with the longest methylene sequence length irrespective of which component from which the molecules originated. Alternatively liquid–liquid separation may occur first, followed by independent crystallization of the two liquid phases, and in this case the composition of the liquid phases will be determined by the temperature and associated tie-lines of the phase diagram for the blend. Analogous phase separation mechanisms have been observed for polyethylene–polypropylene blends (23–25).

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Studies, where one of the blended polyethylenes was deuterated, have revealed immiscibility of polyolefins with SANS, by exploiting the difference in coherent scattering lengths between proton and deuterium nuclei to provide suitable enhanced scattering contrast. SANS techniques have been applied to polyolefin blends to investigate the thermodynamic interactions relating to phase behavior and purecomponent pressure–volume-temperature (PVT) relationships. Qualitative predictions of the miscibility of model homopolymer blends of PE and deuterated PE have been achieved by simulation. Optical microscopy, differential scanning calorimetry, rheological measurements, and scanning electron microscopy have been applied. Molecular architecture and the effects of molar mass have been studied in relation to phase separation and PVT data (26). Thermodynamic interactions in polyolefin blends have been shown to originate from induced- dipole forces but they depend upon the component structures. SANS has been used to determine interactions of blends, including solubility data, light scattering to determine phase boundaries, and PVT measurements to establish cohesive energies. Isotope and microstructure interactions have been shown to contribute to the Flory–Huggins interaction parameter in mixtures of deuterated and protonated PE. Lack of a single dominant thermodynamic parameter determining phase behavior means polyolefin blends are influenced by small variations in molecular architecture. Binary phase behavior was correlated with statistical segment length asymmetry and its dependence on the Flory–Huggins interaction parameter. Additional contributions from temperature, composition, and molar mass have been established, so Flory–Huggins theory of polymer blends and their thermodynamics has been extended to include them. Miscibility loop behavior has been predicted and described using a temperature-dependent interaction parameter, leading to an UCST phase diagram. The predictions have been attributed to nonrandom packing effects where constituents have sufficient asymmetry in the polymerization indices (27). Molecular packing may have a dominant influence on polyolefin miscibility, so the accurate prediction of polyolefin melt structures would enable understanding of reasons for the difficulty of PE miscibility prediction to be obtained. In the melt, molecules of a homopolymer polyolefin, such as LPE, generally exist in random conformations, but some order is present because of weak intramolecular and intermolecular interactions, such as dispersive forces. Molecular conformation is determined by the configuration of branches, which influence the packing of the PE chains, determining the melt structure, density, and, hence, the miscibility of PE in the melt. The physical properties of the melt are thus determined by conformation. The branch structure, length, frequency, and distribution determine the most favorable conformations. The conformations allow or limit interactions to the extent that PE with differing branch structures may be immiscible. The thermodynamic stability of a single liquid phase has been theoretically predicted by computation of pure component and blend parameters dependant upon entropic considerations (28). Microphase separation has been ascribed to a considerable increase in unfavorable noncombinatorial entropy, using a Flory–Huggins equation to predict the free energy of mixing (29).

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Polyolefin Blends

3.8 RHEOLOGY OF MELTED POLYETHYLENE BLENDS Linear polyethylene rheology is dependant on the critical molar mass for entanglements (940 g mol1 (30), which is low due to lack of chain stiffening groups or side chains. Chain entanglements and branch entanglements provide non-Newtonian rheology for higher molar mass PE. Entanglements of branches increase the melt strength of PE with long branches, especially LDPE, and also some newer mPE that are polymerized to contain long branches in addition to the short branches from a comonomer. Short branches do not interact or entangle sufficiently to provide melt strength. PE with differing branch content have been shown to form immiscible liquid–liquid mixtures. The rheology of immiscible PE blends deviate from additivity expected for a homogeneous liquid. For example, HDPE and LLDPE are miscible when the branch content is low (<4 branches/100Cs). However for >8 branches/ 100 Cs, phase separation takes place. The rheology of melted polyethylene blends is of importance in processing (31). Processing is often by extrusion where the main product is film, either from the blown film process or cast film. The molecular structure and hence meltingtemperature profile determines screw design. Usually PE blends melt over a broad temperature range, so a long compression zone is required. Orientation of blown film requires high melt strength that is enhanced by long branches while subsequent film strength is dependent on crystallinity requiring long methylene sequence lengths and so few branches. The optimum structure for processing and strength of PE films is best achieved by blending different PE. Injection molding requires high melt flow and hence low molar mass, while the mechanical properties will require high molar mass; the optimum properties may require a blend of PE with differing molar mass, though bimodal molar mass distribution. Foams are prepared by gas injection molding, and successful foams require high melt strength. Rotational casting, powder coatings, heat sealing all require melt flow under relatively low shear, so non-Newtonian characteristics may not have the opportunity to become apparent. Rheology can be modified by processing aids and lubricants and some of these are low molar mass PE, such as PE waxes and microcrystalline waxes. Rheology has been used to interpret the miscibility of polyethylene blends in lieu of direct measurement of phase separation. The sensitivity of various models has been shown suitable for distinguishing homogeneous melts from liquid–liquid phase separated melts (32,33). Such phase separated melts may be of advantage in extrusion processing of polyethylene blends; particularly in blown film formation where melt strength and rheological characteristics are critical to bubble stability, film uniformity, orientation, and clarity. Metallocene poly(ethylene-cobutene)s were blended with LDPE and the viscosities were increased at low frequencies though shear thinning was more evident. The relaxations were retarded in the presence of LDPE and the effects were more noticeable for the metallocene copolymers with higher butene content (34). Metallocene HDPE and metallocene LLDPE blends were confirmed by rheological analysis to be miscible, contrary to blends of the same metallocene HDPE with LDPE (35).

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3.9 MECHANICAL PROPERTIES OF POLYETHYLENE BLENDS PE blends are compatible; they can be miscible, cocrystallize, or crystallize with separate crystallite distributions. These morphologies effect modulus, yield and any double yield phenomena, strain hardening, elasticity and recovery, creep, creep recovery, stress relaxation, tear strength. These properties are all measured using static stress or strain. Typical tensile stress–strain curves for LLDPE, LDPE, and a blend of LDPE–EP rubber are shown in Fig. 3.9. LDPE shows a sharp yield, plastic flow with some strain hardening with increasing strain. The LDPE–EP blend is a thermoplastic vulcanizate, though it still yields, it extends with constant stress. LLDPE shows a sharp yield followed by plastic flow, until extensive strain hardening occurs over the strain range 4 to 13. Strain hardening is a problem for films where rapid stretching to high strain is required, such as pallet wrap, stretch wrap, and agricultural films. The film exhibits decreased elasticity during strain hardening, as crystallinity is oriented in the direction of flow and the crystalline morphology changes. Figure 3.10 shows atomic force microscopy of the Ziegler–Natta-catalyzed LLDPE containing 5% hexene, shown in the stress–strain curve (Fig. 3.9). The original crystals become elongated and orientated during straining with loss of film elasticity, increase in surface roughness, and loss of film transparency. Blending with LDPE or metallocene VLDPE will reduce the strain hardening. Polyethylenes undergo interlaminar deformation during tensile elongation. This can give rise to a double yield phenomenon. At the onset of the first tensile yield, chain slip and lamella rotation occur and this process is reversible. The second tensile yield is irreversible and coinsides with lamella fragmentation (36). These mechanically induced morphological changes and the observation of any double yield phenomena are dependent on several structural factors that can be controlled by

Figure 3.9 Tensile stress–strain curves for LLDPE, LDPE, and LDPE–EP blends.

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Polyolefin Blends

Figure 3.10 Atomic force microscopy of ZN–hexene–LLDPE before strain (a) and after strains of 3 (b) and 6 (c) (10 mm  10 mm).

polyethylene blending. Homogeneous or heterogeneous slips have been identified as responsible for different macroscopic deformation regimes. Homogeneous slip, arising from dislocations within the crystals, has a higher thermal activation and predominates as temperature increases or strain rate decrease. Heterogeneous slip, operating through defective crystal block boundaries, proceeds with lower strain hardening (37). Dynamic mechanical properties exhibit side chain or branch motions; short main chain segment motions, main chain segmental motions, recrystallization, and melting. These transitions are observed as inflections in the storage modulus curve with temperature, peaks in either the loss modulus or damping factor (tan(d)) curves. Figure 3.11 shows the dynamic mechanical spectroscopy (DMS) of a ZN–VLDPE at 1 Hz in tensile mode. The glass transition temperature (maxima of the loss modulus

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Figure 3.11 Dynamic mechanical spectroscopy of ZN–VLDPE in tensile mode using frequency of 1 Hz.

peak, or the damping factor peak, as often used) is at about 40
C. The transitions of individual components of a blend are often detected independently. The activation energy for each of the transitions can be determined using multifrequency data and correlation with the Arrhenius equation. Williams–Landel–Ferry (WLF) analysis of the multifrequency data is used to construct time (frequency)–temperature master curves to obtain plateau or terminal viscoelastic properties. These properties reveal the contributions of differing branching levels or distributions within a blend. This is illustrated in Fig. 3.12 that shows a LDPE–EP rubber blend 1, 2, 5, 10, and 20 Hz in

Figure 3.12 Dynamic mechanical spectroscopy of LDPE–EP rubber blend in tensile mode using synthetic frequency multiplexing of 1, 2, 5, 10, and 20 Hz.

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Polyolefin Blends

tensile mode measured using synthetic frequency multiplexing. The glass transition temperature (50
C) is shifted to higher temperatures with higher frequencies. A second, lower temperature transition is shown in the range 70 to 90
C. Metallocene poly(ethylene-cooctene)s showed changes in the b-transition region that were related to motions of segments located at interfacial regions (38).

3.10 ADDITIVES Many additives are included in PE and their blends. These include colorants, stabilisers, ultraviolet protection, flame retardants such as antimony oxide and decabromodiphenyl ether, or halogen free flame retardants such as magnesium and aluminum hydroxides. Prodegradents are an alternative to stabilizers, such as prooxidants that are currently being used in some packaging films to reduce environmental impact. Surface properties such as tack, antiblocking, corona discharge treatment for printing, and its subsequent reversion over time are modified by additives. Additives for processing include internal and external lubricants. Microcrystalline waxes, hydrocarbon waxes, stearic acid, and metal stearates such as calcium, magnesium, and zinc stearates are common lubricants. Often blends with lower melting temperature polyethylenes, particularly those with long branches are preferred for processing enhancement. Elastomeric or plastomeric polyethylenes often require cross-linking, so radical initiators such as cumyl peroxide are added, or alkoxyvinylsilanes grafted for subsequent water-induced cross-linking.

3.11 CONCLUSIONS The polyethylenes are a large and diverse family of chemically similar polymers; increased understanding is required to learn how to combine them to achieve blends with a custom range of properties. Single-site (metallocene)-catalyzed PE have brought a new life cycle to PE developments and applications. These new PE have initially found applications in blends with existing PE whose characteristics are well known but require processing and performance improvement. Most PE blends are with other PE to provide combinations of molar mass, short branching, and long branching. Other blends are with PE copolymers with polar or functionalized comonomers so that increased surface polarity and functionalization can be included. Production and cost pressures increase the performance requirements for PE. PE can be enhanced to replace other polymers as well as expanding into areas occupied by other materials. The similarity of the chemistry and lack of distinguishing features for many characterization techniques make investigation of the structure–property– performance of PE blends a difficult and often controversial task. Critical requirements for PE are rapid processing, thinner yet stronger films, ease of direct printing or adherence of labels, rapid heat sealing, high strain rate at high strain for stretch wrap, plastic memory for shrink wrap, cross-linkability to resist creep and stress relaxation, cross-linking to resist elevated temperatures when the

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low melting temperature of PE would otherwise be a constrain, increased elasticity covering the range from traditional elastomers to stiff thermoplastics. PE compositions need performance compatible with newer high speed processing equipment and high speed packaging and finishing machinery. The new PE and their blends are meeting these needs and opening new opportunities where they easily meet or exceed expectations. All of these properties and structural diversity are now being derived from the polymer with the simplest chemical structure.

NOMENCLATURE Tg Tm Tm
Tmc Tc PE LPE mPE M HDPE BPE LLDPE VLDPE ULDPE LDPE TPO TPE VAc MA BA GMA AA MAn MSL SANS PVT UCST SAXS WAXS SSA SC HBPE DMS ZN

Glass transition temperature Melting temperature Equilibrium melting temperature Copolymer melting temperature Crystallization temperature Polyethylene Linear polyethylene Metallocene polyethylene Molar mass High density polyethylene Branched polyethylene Linear low density polyethylene Very low density polyethylene Ultralow density polyethylene Low density polyethylene Thermoplastic polyolefin Thermoplastic polyolefin elastomer Vinyl acetate Methyl acrylate Butyl acrylate Glycidyl methacrylate Acrylic acid Maleic anhydride Methylene sequence length Small-angle neutron scattering Pressure–volume–temperature Upper critical solution temperature Small-angle X-ray scattering Wide-angle X-ray scattering Successive self-nucleation and annealing Stepwise isothermal cooling Long-chain branched polyethylenes Dynamic mechanical sprectroscopy Ziegler–Natta catalyst

82 EP TPO

Polyolefin Blends

Ethylene–propylene copolymer elastomer Thermoplastic polyolefin elastomer

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Chapter

4

Miscibility and Crystallization Behavior in Binary Polyethylene Blends Moonhor Ree1

4.1 INTRODUCTION Polyethylene (PE) is currently the most widely used commercial polymer throughout the world, and the industrial market for this polymer is still growing due to the variety of applications that are based around its superior properties, such as high chemical and mechanical resistances, easy processability, and low specific gravity, as well as low manufacturing costs. So far, several types of PE materials have been developed, each with different levels of mass density, branching content, and molecular weight. These PE materials include linear high density PE (HDPE), linear ultrahigh molecular weight PE (UHMWPE), linear low density PE (LLDPE), and low density PE (LDPE) (Fig. 4.1) (1–5). These polymers are usually synthesized by the conventional polymerization of the ethylene monomer, or with one or more a-olefin comonomers (1-butene, 1-hexene, 1-octene, etc.) with the aid of Ziegler–Natta catalysts (1–5). Advances in the catalysts used for the polymerization of olefins have allowed manufacturers to expand their product line to accommodate a wide range of densities and branching structures. In particular, new high performance PEs are mainly produced via the use of metallocene catalysts (i.e., single-site catalysts), which have been shown to have superior properties compared with conventional PEs (6,7). The single-site nature of the catalysts allows for the formation of polymers

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Department of Chemistry, Polymer Research Institute, Pohang Accelerator Laboratory, National Research Lab for Polymer Synthesis and Physics, and Center for Integrated Molecular Systems, Pohang University of Science & Technology (Postech), Pohang 790-784, Republic of Korea Polyolefin Blends, Edited by Domasius Nwabunma and Thein Kyu Copyright # 2008 John Wiley & Sons, Inc.

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Figure 4.1 Schematic molecular structures of LDPE, LLDPE, HDPE, and UHMWPE.

with a narrower molecular weight distribution and a more uniform distribution of comonomers, as compared with conventional PEs (6,7). HDPE gives rise to stiffer materials with good tensile properties but poor impact and tear resistance (1,5). Moreover, UHMWPE has excellent mechanical properties but very poor melt properties, causing severe processing difficulties (4,5). LLDPE, which is synthesized by copolymerization of ethylene with one or more a-olefin comonomers, has short chain branching and presents excellent mechanical properties, such as impact and tear resistance, as well as high tensile strength (1–5). However, due to a narrow molar mass distribution, LLDPE exhibits an elevation in the melt viscosity and a lowering of the melt strength, causing relatively poor processability (1–5). In general, metallocene-catalyzed LLDPE is synthesized with a much narrower molecular weight distribution, compared to that of Ziegler–Nattacatalyzed LLDPE, and exhibits very poor melt properties that cause processing difficulties (6–15). On the contrary, LDPE is a polymer composed of both long and short chain branches and is well known to have good melt properties and therefore efficient processability (16). In general, the processing behavior of PEs mainly depends on parameters such as the weight-average molecular weight and distribution, comonomer content, and the long-chain-branching content and distribution (17–22). In particular, the longchain-branching content is known to have a large impact on polymer processability; PE polymers with long-chain-branching demonstrate excellent processing properties (18,23).

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These PE polymers offer a broad spectrum of structures, properties, and applications. However, the blending of different types of PEs (HDPE, LDPE, LLDPE, and UHMWPE) has attracted growing interest because of the potential for obtaining low cost materials with improved mechanical properties and better processabilities, as compared to those of the pure constituents (1–34). Nowadays, 70% of PEs in the market are blends (24). The processability and properties of PE blends are dependent on the melt miscibility. Moreover, the properties are also dependent on the morphological structure of the blend, which is basically a combination of the crystallization behavior and melt miscibility. Therefore, the miscibility and crystallization behavior of PE blends have been prevalent research topics over the last two decades.

4.2 MISCIBILITY 4.2.1 Linear and Short Branched Polyethylene Blends Ree (5) first studied quantitatively the miscibility of two HDPE/LLDPE blend systems using the small-angle neutron scattering (SANS) technique: (i) blends of a HDPE-d (237,000 weight-average molecular weight Mw , 2.06 polydispersity PDI, and 98% deuteration) and a LLDPE-B (114,000 Mw , 4.50 PDI, and 18 ethyl branches per 1000 backbone carbons); (ii) blends of a HDPE-d (109,000 Mw , 1.80 PDI, and 98% deuteration) and a LLDPE-O (96,000 Mw , 4.50 PDI, and 15 hexyl branches per 1000 backbone carbons). For these blends in the melt, the Flory–Huggins interaction parameter x was determined to be small, with positive values of the order 10 4, but did not exceed the critical x values, that is, the upper limit of the stability of a single miscible phase. Moreover, these blends revealed no SANS cloud point over the temperature range 140–300 C, indicating that the HDPE blend systems with LLDPE-B and LLDPE-O are both miscible in the melt. Nicholson et al. (35) extended the SANS analysis approach to accommodate blends of HDPEs with hydrogenated and deuterated polybutadienes (PB, a LLDPE-B and PB-d, a LLDPE-B-d) with different amounts of ethyl branches ranging from 18 to 106 per 1000 backbone carbons: HDPE (111,830 Mw and 1.23 PDI), HDPE-d (116,580 Mw and 1.27 PDI), LLDPE-B-d(1) (102,660 Mw , <1.05 PDI, and 18 ethyl branches), LLDPE-B-d(2) (22,620 Mw , <1.05 PDI, and 22 ethyl branches), LLDPE-B(1) (161,650 Mw , <1.05 PDI, and 39 ethyl branches), LLDPE-B(2) (101,230 Mw , <1.05 PDI, and 65 ethyl branches), LLDPE-B(3) (76,850 Mw , <1.05 PDI, and 88 ethyl branches), and LLDPE-B(4) (179,670 Mw , <1.05 PDI, and 106 ethyl branches). Alamo et al. (36) and Agamalian et al. (37) also carried out SANS measurements on the binary blends prepared from hydrogenated and deuterated HDPEs (94,000–201,000 Mw and 1.03–3.4 PDI) and hydrogenated and deuterated PBs (LLDPE-Bs: 21–106 branches, 78,000– 123,000 Mw , and <1.07 PDI). Both groups found that the melt miscibility of the HDPE/LLDPE-B blend is dependent on the extent of ethyl branches; in the blend, the LLDPE-B component containing more than 60–80 branches per 1000 backbone carbons induced phase separation (35–37). The miscible blends in the melt

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were found to have x values on the order of 10 4 (35–37). The effect of the ethyl branches on the melt miscibility of the HDPE/LLDPE-B blends was further confirmed by molecular dynamics simulations (38–41). A similar branching effect on the miscibility was observed in the HDPE/ LLDPE-O blends, which were prepared from HDPE (49,400 Mw and 3.60 PDI) and LLDPE-O polymers (69,200–104,000 Mw and 1.8–3.40 PDI) with 2–87 hexyl branches per 1000 backbone carbons (42). It was observed that the critical branch number in the LLDPE-O component capable of causing immiscibility in the HDPE/ LLDPE-O blend was 50 branches per 1000 backbone carbons, as determined using inverse gas chromatography, rather than the SANS technique. The miscibility of HDPE/LLDPE blends was also studied using differential scanning calorimetry (DSC), dynamic mechanical thermal analysis (DMTA), rheological analysis, excimer fluorescence analysis, and transmission electron microscopy (TEM). Lee et al. (43) studied blends of HDPE (121,000 Mw and 16.8 PDI) with LLDPEs (LLDPE-B: 89,300 Mw and 3.8 PDI; LLDPE-O: 93,100 Mw and 3.6 PDI) having 15–16 branches per 1000 backbone carbons using DSC and DMTA, and found the blends to be miscible. Lee and Denn (44) conducted rheological measurements on the blends of HDPE (40,000 Mw and 3.1 PDI) with LLDPE-O polymers (68,300–134,000 Mw , 3.2–3.8 PDI, and 10–16 hexyl branches per 1000 backbone carbons), and found that the HDPE/LLDPE-O blends are miscible in the melt over the whole composition range. Zhao et al. (45) attempted to examine the miscibility of HDPE/LLDPE blends using excimer fluorescence analysis; the LLDPE component in these blends was labeled with a chromophore. They found that the LLDPE component is miscible with the blended HDPE. It is noteworthy that in such studies using chromophores, the effect of the labeled chromophore on the blend miscibility should be considered. Tabtiang et al. (46) used rheological analysis to study the observed immiscibility of blended HDPE (45,000 Mw ) with LLDPE-O having a very high content of branches, typically 125 hexyl branches per 1000 backbone carbons. Furthermore, Hill et al. (47, 48) reported on the use of DSC and TEM analyses to study the partial miscibility of quenched HDPE/LLDPE blends (quenched from the melt). The authors studied the miscibility of blends of HDPEs (50,000–100,000 Mw and 2.8– 3.5 PDI) with LLDPE-B polymers (10,000–160,000 Mw , 2.0–5.0 PDI, and 0.5–22 ethyl branches per 1000 backbone carbons) as well as LLDPE-O polymers (37,000– 130,000 Mw , 2.0–2.2 PDI, and 10.5–59 hexyl branches per 1000 backbone carbons). They found that the blends with LLDPE polymers having 10–40 branches per 1000 backbone carbons exhibited a closed loop phase diagram at low HDPE concentrations. The obtained phase diagrams are quite different from those obtained from the SANS results discussed above. In fact, in the DSC and TEM study, Hill et al. assumed that the phase structure of the blends in the melt is retained in the blend samples quenched from the melt, and if any crystallization is involved in such the quenched blend samples, the structure of the formed crystals is dependent on the phase structure. It was determined that the quenched blend samples do not have zero crystallinity. On the contrary, for the above SANS studies it is noted that the used LLDPE-B-d samples were prepared from deuterations of the polybutadienes, which

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are different from LLDPE-B samples synthesized by the copolymerization of ethylene and 1-butene. As reviewed above, the miscibility of the HDPE/LLDPE blends seems to depend on the molecular weight, polydispersity, composition, and temperature, as well as on the branch content and length of the components. However, the HDPE and LLDPE samples used in the blend studies were limited to certain ranges of molecular weight, polydispersity, and branch content. Further, there was a distinct lack of information regarding the phase diagram. Therefore, more quantitative analysis is still needed in order to completely understand the miscibilities and phase diagrams of HDPE/ LLDPE blends.

4.2.2 Blends of Linear and Long Branched Polyethylenes Ree (5) first studied in detail the miscibility of HDPE/LDPE blends using the SANS technique: HDPE (198,500 Mw and 2.28 PDI), HDPE-d (237,000Mw , 2.06 PDI, and 98% deuteration) and LDPE (286,000 Mw , 15.98 PDI, and 27.6 branches per 1000 backbone carbons: 26 short and 1.6 long branches). For the blends in the melt, x was determined to have small, positive values on the order of 10 4, which did not exceed the critical x values, that is, the upper limit of the stability of the single miscible phase. Moreover, the blends did not show a SANS cloud point over the temperature range 140–300 C. These SANS results indicated that the HDPE/LDPE blends are miscible in the melt. Alamo et al. (49) extended the SANS analysis to blends of HDPEs (101,000–201,000 Mw and 2.9–3.6 PDI) with LDPEs (110,000–212,000 Mw , 5.1–17.2 PDI, and 12–16 short branches and 1.5–2.8 long branches per 1000 backbone carbons), and found that the HDPE/LDPE blends in the melt state form a homogeneous liquid mixture throughout the whole composition. In addition to the SANS technique, other analytical methods were employed to investigate the miscibility of HDPE/LDPE blends. Lee et al. (43) investigated the miscibility of HDPE/LDPE blends using DMTA: HDPE (121,000 Mw and 16.8 PDI) and LDPE (73,000–98,000 Mw and 8.7–9.2 PDI; 32–34 branches per 1000 backbone carbons). They found that HDPE/LDPE blends are miscible in the amorphous phase. Lee and Denn (44) conducted rheological measurements on the blends of HDPE (40,000 Mw and 3.1 PDI) and LDPE (388,000 Mw and 10 PDI) at 160 C; no detailed information on the branches in the LDPE was given. These results suggest that the HDPE/LDPE blends are partially miscible in the melt, depending on the composition. Barham et al. (50) studied the miscibility of HDPE/LDPE blends using DSC and TEM: HDPE of 98,000 Mw and LDPE of 208,100 Mw (16 short and 10 long branches per 1000 backbone carbons). They observed that the HDPE/LDPE blends are partially miscible in the melt depending on the composition. Hill et al. (51–54) conducted DSC, rheology, and TEM measurements on blends of HDPEs (2550–2,000,000 Mw and 1.4–12.0 PDI) and LDPE polymers (90,000–208,100 Mw and 8.0–8.5 PDI; 10–16 short and 1–16 long branches per 1000 backbone carbons), and constructed phase diagrams, indicating that the blends are partially miscible

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depending on the composition, molecular weight, and branch content; phase separation took place in the LDPE-rich region and the size of the phase separation region in the phase diagram decreased with decreasing molecular weight of HDPE. As pointed out in the earlier section for HDPE/LLDPE blends, the validity of the techniques of Hill et al. is still controversial because the authors assumed that quenching a sample from the melt state to the solid state does not change the morphology of the blends in the melt. In other work in this field, Fan et al. (40) investigated the miscibility of HDPE/LDPE blends using molecular dynamics simulations and suggested that the critical branch content of LDPE required to induce immiscibility with HDPE is 25 branches per 1000 backbone carbons. As discussed above, there is still no consensus on the miscibility of the HDPE/ LDPE blends. Therefore, this subject needs more detailed investigation.

4.2.3 Blends of Short and Long Branched Polyethylenes Kyu et al. (2) and Ree et al. (3,5) studied blends of LLDPE-B (114,000 Mw , 4.50 PDI, and 18 ethyl branches per 1000 backbone carbons) and LDPE (286,000 Mw , 15.98 PDI, and 26 short and 1.6 long branches per 1000 backbone carbons) using in situ small-angle light scattering (SALS) and DSC and found that these blends are miscible across the whole composition range. Similar miscibility results were reported for LDPE blends with LLDPE-B and LLDPE-O polymers by other research groups. The DSC and DMTA analysis of Lee et al. (43) confirmed that the blends of LLDPEs (LLDPE-B: 89,300 Mw , 3.8 PDI, and 15–16 branches per 1000 backbone carbons; LLDPE-O: 93,100 Mw , 3.6 PDI, and 15–16 branches per 1000 backbone carbons) and LDPEs (73,000–98,000 Mw and 8.7–9.2 PDI; 32–34 branches per 1000 backbone carbons) are miscible at all compositions. Chen et al. (55, 56) obtained miscible blends from LDPEs (89,000–474,000 Mw and 4.4–23.3 PDI) with LLDPEB (58,000 Mw , 2.65 PDI, and 31.5 ethyl branches per 1000 backbone carbons) and LLDPE-O polymers (96,700–474,000 Mw , 2.86–3.80 PDI, and 12.0–64.5 hexyl branches per 1000 backbone carbons). In these blends, miscibility was achieved for all compositions, regardless of the branch lengths and contents considered. The steady-state rheological, thermal, and mechanical properties of the binary blends consisting of LLDPE-H (690,000 Mw , 1.86 PDI, and 18.6 butyl branches per 1000 backbone carbons) and LDPE (660,000 Mw , 9.43 PDI, and 9 short and 5 long branches per 1000 backbone carbons) were investigated by Yamaguchi et al. (29, 30). According to the rheological measurements, LLDPE is miscible with LDPE in the melt state. Silveira and Choi (57) also studied the miscibility of blends of LDPE with six different LLDPE polymers using inverse gas chromatography; the molecular characteristics of the used PE polymers were not given. These blends were found to be miscible regardless of the molecular characteristics of the LLDPEs used. Lee and Denn (44), however, reported somewhat different results on LLDPE/ LDPE blends. They investigated blends of LDPE (388,440 Mw , 10.0 PDI, and 32 branches per 1000 backbone carbons) with LLDPE polymers (68,300–134,000 Mw and 3.2–3.8 PDI; 10–16 short branches per 1000 backbone carbons) at 160 C using

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rheological analysis and found that the blends are partially miscible in the melt depending on the composition. Muller et al. (58) studied the miscibility of LDPE (110,000 Mw ) and LLDPE-O (130,000 Mw ) blends by means of DSC. These blend samples too were found to be partially miscible, depending on the composition. Using rheological analysis, Hameed and Hussein (33) also investigated the miscibility of blends of LDPE (100,000 Mw , 4.14 PDI, and 7.8 branches per 1000 backbone carbons), LLDPE-B (110,000 Mw , 1.78 PDI, and 20.4 ethyl branches per 1000 backbone carbons), and two LLDPE-H (67,000 Mw , 1.85 PDI, and 21.1 butyl branches per 1000 backbone carbons; 108,000 Mw , 1.83 PDI, and 18.0 butyl branches per 1000 backbone carbons) and found that the blends were partially miscible. Interestingly, the blends of LDPE with LLDPE-H (67,000 Mw ) were almost miscible throughout the whole composition range. In contrast, the blends of LDPE with LLDPE-H (108,000 Mw ) were partially miscible. The miscibility of the blend was nonsymmetric with respect to composition; the addition of a small amount of LLDPE-H (108,000 Mw ) to LDPE was more likely to cause immiscibility than the addition of a small amount of LDPE to the LLDPE-H polymer. However, increasing the branch length of LLDPE from ethyl to butyl showed no effect on the miscibility with LDPE. Hill et al. (59) examined the miscibility of LLDPE-O (40,000 Mw , 4.2 PDI, and 15 hexyl branches per 1000 backbone carbons) and LDPE (112,000 Mw , 12.0 PDI, and 15 short and 10 long branches per 1000 backbone carbons) using DSC and TEM and found that the blend produces a complex phase diagram. The realization of the phase diagram indicates that the blend is partially miscible, depending on the composition and temperature window. However, the phase diagram was constructed from the DSC and TEM data measured for the quenched samples. Thus, such phase diagrams still need to be verified by direct measurements. Contrary to the reports discussed above, Liu et al. (60) reported rheological results indicating that the LLDPE-B (45,700 Mw , 2.1 PDI, and 13.5 ethyl branches per 1000 backbone carbons) and LDPE (102,000 Mw , 6.8 PDI, and 27.5 branches per 1000 backbone carbons) blend is an immiscible system. As reviewed above, there is still no consensus on the miscibility of LLDPE/ LDPE blends, or the effects of short and long branches on the miscibility. Furthermore, the lack of detailed SANS analysis on LLDPE/LDPE blends is surprising, although the miscibility specifications of these blends are highly demanded, particularly in relation to their various applications in the polymer industry.

4.3 CRYSTALLIZATION BEHAVIORS 4.3.1 Blends of Linear and Short Branched Polyethylenes Datta and Birley (61) first reported on the crystallization behavior of HDPE/LLDPE blends using DSC and wide-angle X-ray diffraction (WAXD). In the DSC measurements, the blends revealed only one endothermic peak for both fast and slow cooling conditions, indicating that the blend components were cocrystallized. The X-ray

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analysis confirmed that the unit cell dimensions of the crystallized HDPE/LLDPE blends are identical to those of the HDPE. Here, it is noted that details of the molecular characteristics of the LLDPE were not given. Similar cocrystallization was evident in the DSC measurements for blends of HDPEs (29,000 Mw and 4.08 PDI; 121,000 Mw and 16.8 PDI) with an LLDPE-O (93,100 Mw , 3.6 PDI, and 15 hexyl branches per 1000 backbone carbons) (43) and LLDPE-Bs (64,000 Mw , 4.57 PDI, and 12.5 ethyl branches per 1000 backbone carbons; 89,300 Mw , 3.8 PDI, and 16 ethyl branches per 1000 backbone carbons) (25, 43). The cocrystallization was also observed in HDPE blends with LLDPE-O having a high number of hexyl branches (e.g., 32) (62). Hu et al. (1) studied in detail HDPE/LLDPE blend systems using DSC, WAXD, small-angle X-ray scattering (SAXS), small-angle light scattering (SALS), and Raman longitudinal-acoustic-mode spectroscopy (LAM). In their study, the blend samples were prepared from HDPE (160,000 Mw and 7.1 PDI) and LLDPE-B (114,000 Mw , 4.50 PDI, and 18 ethyl branches per 1000 backbone carbons). Their DSC results are similar to those of Datta and Birley (61). The WAXD, SAXS, and LAM analyses of Hu et al. found that the branches are excluded from the polyethylene unit cell, resulting in thinner lamellae. This finding was confirmed by molecular dynamics simulations (63). Tashiro et al. (64, 65) investigated the crystallization behavior of blends of HDPEs (80,000–126,000 Mw and 5.3– 5.7 PDI) with LLDPE-B polymers (61,000–75,000 Mw , 2.0–3.1 PDI, and 17 and 41 ethyl branches per 1000 backbone carbons) using DSC and infrared spectroscopy. They found that even under slow cooling conditions, the LLDPE-B polymer containing 17 branches cocrystallizes with the HDPEs over the whole composition range. In contrast, the LLDPE-B polymer with 41 branches was found to partially cocrystallize with HDPEs, depending on the composition. Rana (66, 67) studied in detail the cocrystallization kinetics of a HDPE/LLDPE blend system using DSC and SALS; in that study, a HDPE and an LLDPE-O having 30 hexyl branches per 1000 backbone carbons were used, but their molecular weight and polydispersity were not documented. It was found that, on the basis of Avrami exponent, both the HDPE and LLDPE-O polymers form cocrystallites. The incorporation of LLDPE segments into the HDPE crystallites progressively dilutes the properties of HDPE in the blend. The variation in crystallization half-time (t1/2) with increasing LLDPE-O content shows three distinct stages: t1=2 increases slowly in the region of 0–40 wt% LLDPE-O (HDPE-rich blend), remains constant in the 30–70 wt% LLDPE-O containing region (middle region of blend composition), and increases sharply with further increases of the LLDPE-O content. These variations of t1=2 qualitatively account for the change in crystallinity percentage, the Avrami exponent, and crystallite-size distribution. These observations were further supported by SALS measurements. The effect of branch distribution of LLDPE on the crystallization of HDPE/ LLDPE blends has also been investigated. The DSC and scanning electron microscopy (SEM) studies of Lee et al. (68) found that a Ziegler–Natta-catalyzed LLDPEB having a relatively heterogeneous branch distribution (179,000 Mw , 4.2 PDI, and 16 ethyl branches per 1000 backbone carbons) cocrystallizes with HDPE (147,000 Mw and 3.7 PDI) in a wide range of blend compositions, and at the same time,

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separate crystallization occurs to a small extent. A metallocene-catalyzed LLDPE-B polymer with a relatively homogeneous branch distribution (170,000 Mw , 2.8 PDI, and 17.5 ethyl branches per 1000 backbone carbons) was also found to cocrystallize with HDPE in a wide range of blend compositions. On the contrary, the metallocenecatalyzed LLDPE-B polymer undergoes separate crystallization to a relatively larger extent compared to that of the Ziegler–Natta-catalyzed LLDPE-B (68). From the results discussed above, it is concluded that the crystallization behavior of HDPE/LLDPE blends depends on the number, length, and distribution of branches in the LLDPE component. In general, it was found that increases in the number and degree of branch distribution in the LLDPE component reduce the tendency of cocrystallization in the LLDPE blend with HDPE.

4.3.2 Blends of Linear and Long Branched Polyethylenes Lee et al. (43) studied the crystallization behavior in blends of HDPE (121,000 Mw and 16.8 PDI) and LDPE (73,000–98,000 Mw and 8.7–9.2 PDI; 32–34 branches per 1000 backbone carbons) and reported that the HDPE and LDPE components cocrystallize in the blend. In contrast, in DSC and WAXD studies on HDPE/LDPE blends cooled both rapidly and slowly, Datta and Birley (61) found that both the rapidly and slowly cooled blend samples exhibited two distinct endothermic peaks in the DSC measurements, indicating that the HDPE and LDPE components in the blends crystallize separately. However, no data were given on the molecular characteristics of LDPE. Similar separate crystallization behavior was also reported for HDPE/LDPE blends by Ree (5) and Song et al. (69, 70). In these studies, HDPE (160,000 Mw and 7.1 PDI) and LDPE (286,000 Mw , 15.98 PDI, and 26 short and 1.6 long branches per 1000 backbone carbons) were employed. In the DSC measurements, the blends revealed two peaks characteristic of both exotherms and endotherms, and indicative of the formation of separate crystals. This separate crystallization phenomenon was also found at the crystallite and lamellar levels using WAXD, SAXS, and LAM. The SAXS spectra of the blends were interpreted with respect to the segregation phenomena of interfibrillar scales during crystallization of the blend components. The SALS studies were performed in order to consider the morphology and crystallization behavior of the blends at the super structural level. Light scattering results demonstrated that the blend samples were predominantly volume-filled as a result of the primary crystallization of the HDPE component, whereby the LDPE component crystallized as a secondary process within the spherulites of the high density component. In addition, the SALS intensities from the blend samples of various compositions were interpreted on the basis of the spherulite radius, its internal disorder, and the volume fraction crystallinity. The crystallization process of the LDPE component in the blends was investigated by means of DSC. In this study, it was evident that the degree of crystallinity changes linearly with logarithmic time during isothermal crystallization of the LDPE component.

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From the results discussed above, it is concluded that the crystallization behavior of HDPE/LDPE blends depends on the number, length, and distribution of branches in the LDPE component. In general, most HDPE/LDPE blends reveal a strong tendency to undergo separate crystallization. Cocrystallization is rarely observed in HDPE/LDPE blends.

4.3.3 Blends of Short and Long Branched Polyethylenes Kyu et al. (2) and Ree et al. (3, 5) studied the crystallization behavior of blends of LLDPE-B (114,000 Mw , 4.50 PDI, and 18 ethyl branches per 1000 backbone carbons) and LDPE (286,000 Mw , 15.98 PDI, and 26 short and 1.6 long branches per 1000 backbone carbons) using DSC, WAXD, SAXS, SALS, and LAM. They found that the LLDPE-B and LDPE components crystallize separately. The SAXS and SALS results indicate that when the LLDPE/LDPE blends are cooled from the molten state, the LLDPE-B crystallizes first and forms volume-filling spherulites, which is closely followed by crystallization of the LDPE within the spherulites formed by LLDPE-B. Similar separate crystallization behavior was reported by Lee et al. (43) for blends of LLDPE-B (89,300 Mw , 3.8 PDI, and 15–16 branches per 1000 backbone carbons) with LDPEs (73,000–98,000 Mw and 8.7–9.2 PDI; 32–34 branches per 1000 backbone carbons). Separate crystallization was also reported for blends of LDPEs with LLDPEs having butyl and hexyl branches. Using DSC analysis, Drummond et al. (71) observed that separate crystallization of the individual blend components takes place in the blends of LLDPE-H (146,000 Mw , 3.1 PDI, and 19.9 butyl branches per 1000 backbone carbons) and LDPE (78,700 Mw , 31.0 PDI, and 20.1 branches per 1000 backbone carbons). The DSC studies of Lee et al. (43) revealed that for blends of LLDPE-O (93,100 Mw , 3.6 PDI, and 15–16 branches per 1000 backbone carbons) with LDPEs (73,000–98,000 Mw and 8.7–9.2 PDI; 32–34 branches per 1000 backbone carbons), the blend components also undergo separate crystallization. Hill et al. (69) reported similar separate crystallization behavior in blends of LLDPE-O (40,000 Mw , 4.2 PDI, and 15 hexyl branches per 1000 backbone carbons) and LDPE (112,000 Mw , 12.0 PDI, and 15 short and 10 long branches per 1000 backbone carbons). As reviewed above, LLDPE/LDPE blends tend to undergo separate crystallization of the individual components, depending on the number, length, and distribution of branches in the two blend components.

4.4 CONCLUSIONS The miscibility and crystallization behavior of the three binary PE blends, HDPE/ LLDPE, HDPE/LDPE, and LLDPE/LDPE, were reviewed. In general, differences in the number, length, and distribution of branches in the PE blend components are the major factors governing their miscibility and crystallization phenomenon. In particular, the content and distribution of branches significantly affect both the miscibility and crystallization. Moreover, the presence of a few long chain branches, as well as

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the observed differences in the molecular weight and molecular weight distribution, all affect the miscibility and crystallization. It is noted, however, that studies on both the miscibility and crystallization behavior of PE blends are quite limited due to experimental difficulties, in particular those caused by the chemical repeat units of linear and branched PEs having the same molecular nature. Therefore, the development of more quantitative analytical techniques is still needed in order to fully understand the miscibility and crystallization behavior of PE blends.

NOMENCLATURE DSC DMTA HDPE HDPE-d LDPE LLDPE LLDPE-B LLDPE-B (hydrogenated PB) LLDPE-B-d LLDPE-H LLDPE-O PB PE PDI SANS TEM UHMWPE Mw x

Differential scanning calorimetry Dynamic mechanical thermal analysis High density linear polyethylene Deuterated high density linear polyethylene Low density branched polyethylene Linear low density polyethylene Linear low density polyethylene prepared by the copolymerization of ethylene and 1-butene Linear low density polyethylene prepared by the hydrogenation of polybutadiene Deuterated LLDPE-B Linear low density polyethylene prepared by the copolymerization of ethylene and 1-hexene Linear low density polyethylene prepared by the copolymerization of ethylene and 1-octene Polybutadiene Polyethylene Polydispersity index Small-angle neutron scattering Transmission electron microscopy Ultrahigh molecular weight polyethylene Weight-average molecular weight Flory–Huggins interaction parameter

REFERENCES 1. S. R. Hu, T. Kyu, and R. S. Stein, J. Polym. Sci. B: Polym. Phys., 25, 71 (1987). 2. T. Kyu, S. R. Hu, and R. S. Stein, J. Polym. Sci. B: Polym. Phys., 25, 89 (1987). 3. M. Ree, T. Kyu, and R. S. Stein, J. Polym. Sci. B: Polym. Phys., 25, 105 (1987). 4. P. Vadhar and T. Kyu, Polym. Eng. Sci., 27, 202 (1987). 5. M. Ree, Ph.D. Thesis, University of Massachusetts, Amherst, 1987. 6. G. M. Benedikt and B. L. Goodall (eds.), Metallocene-Catalyzed Polymers: Materials, Properties, Processing and Markets, William Andrew, New York, 1998.

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49. R. G. Alamo, J. D. Londono, L. Mandelkern, F. C. Stehling, and G. D. Wignall, Macromolecules, 27, 411 (1994). 50. P. J. Barham, M. H. Hill, A. Keller, and C. C. A. Rosney, J. Mater. Sci. Lett., 7, 1271 (1988). 51. M. J. Hill, P. J. Barham, A. Keller, and C. C. A. Rosney, Polymer, 32, 1384 (1991). 52. M. J. Hill and P. J. Barham, Polymer, 33, 4099 (1992). 53. M. J. Hill, P. J. Barham, and A. Keller, Polymer, 33, 2530 (1992). 54. M. J. Hill, Polymer, 35, 1991 (1994). 55. F. Chen, R. Shanks, G. Amarasinghe, J. Appl. Polym. Sci., 81, 2227 (2001). 56. F. Chen, R. Shanks, G. Amarasinghe, J. Appl. Polym. Sci., 95, 1549 (2005). 57. M. D. L. V. Silveira and P. Choi, Annu. Technol. Conf. Soc. Plast. Eng., 58(2), 2454 (2000). 58. A. J. Muller, V. Balsamo, F. Da Silva, C. M. Rosales, and A. E. Saez, Polym. Eng. Sci., 34, 1455 (1994). 59. M. J. Hill and C.C. Puig, J. Appl. Polym. Sci., 65, 1921 (1997). 60. C. Liu, J. Wang, and J. He, Polymer, 43, 3811 (2002). 61. N. K. Datta and A. W. Birley, Plast. Rubber Proc. Appl., 2, 237 (1982). 62. S. J. Mahajan, B. L. Deopura, and Y. Wang, J. Appl. Polym. Sci., 60, 1517 (1996) 63. M. Doran and P. Choi, J. Chem. Phys., 115, 2827 (2001). 64. K. Tashiro, M. Izuchi, M. Kobayashi, and R. S. Stein, Macromolecules, 27, 1221 (1994). 65. K. Tashiro, R. S. Stein, and S. L. Hsu, Macromolecules, 25, 1801 (1992). 66. S. K. Rana, J. Appl. Polym. Sci., 61, 951 (1996). 67. S. K. Rana, J. Appl. Polym. Sci., 69, 2599 (1998). 68. S. Y. Lee, J. Y. Jho, and W. Huh, J. Ind. Eng. Chem., 4, 258 (1998). 69. H. H. Song, R. S. Stein, D.-Q. Wu, M. Ree, J. C. Philips, A. LeGrand, and B. Chu, Macronolecules, 21, 1180 (1988). 70. H. H. Song, D.-Q. Wu, B. Chu, M. Satkowski, M. Ree, R. S. Stein, and J. C. Philips, Macronolecules, 23, 2380 (1990). 71. K. M. Drummond, J. L. Hopewell, and R. A. Shanks, J. Appl. Polym. Sci., 78, 1009 (2000).

Chapter

5

Microscopically Viewed Structural Characteristics of Polyethylene Blends Between Deuterated and Hydrogenated Species: Cocrystallization and Phase Separation Kohji Tashiro1

5.1 INTRODUCTION Much attention has been paid in these decades to polymer blends from scientific and practical points of view. Polyethylene (PE) blends between the species with different degrees of branching contents have also been widely investigated (1–23). In these studies, it is very important to trace the aggregation state of each component separately and at a molecular level. However, it had been difficult to distinguish the behaviors of each component separately because these components have almost the same chemical structure of carbon and hydrogen atoms. One possible and useful idea is to use a deuterated PE sample as one component. The deuterated species can be separated from the hydrogenated species by utilizing the difference in the vibrational frequency (24–26) and the neutron scattering cross-sectional area (27–31). Therefore, the infrared and Raman spectroscopies

1 Department of Future Industry-Oriented Basic Science and Materials, Toyota Technological Institute, Tempaku, Nagoya 468-8511, Japan

Polyolefin Blends, Edited by Domasius Nwabunma and Thein Kyu Copyright # 2008 John Wiley & Sons, Inc.

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and the neutron scattering method can be used effectively by utilizing these characteristic differences between the deuterated and hydrogenated species. In the utilization of these two isotopic components, we need to notice the fact that the deuterated species does not behave in perfectly the same manner as the hydrogenated species. In contrast to the case of low molecular weight compounds, the chemical potential of the fully deuterated PE species is appreciably different from that of the hydrogenated PE species because the aggregation of many CD2 groups results in the large isotope effect even if the difference in chemical potential between single CD2 and CH2 groups is negligibly small, that is, the polymeric effect (32). Therefore, the deuterated high density PE [DHDPE, –(CD2CD2)n–], for example, is predicted quite reasonably to behave in a different manner from the normal high density PE [HDPE, –(CH2CH2)n–], giving the detectable difference in the melting and crystallization temperatures. This situation is seriously related to the phenomena of phase separation and cocrystallization between DHDPE and HDPE species. In fact, the blend sample between DHDPE and HDPE was reported to show the phase segregation when the blend sample was slowly cooled from the melt (27–31). The phase segregation or the self-aggregation of D (or H) chains becomes serious in some problems. For example, the phase segregation problem of DHDPE/HDPE blend sample is serious in the research theme of chain-folding mechanism where the clarification of the trace of one isolated chain is important (30). We prepare the blend sample of H chains with very low concentration of D chains, where the D chains must not aggregate together but should be isolated from each other to avoid the optical interference between the D chains in the small-angle neutron scattering experiment. Different from the case of HDPE, the chemical potential of linear low density PE (LLDPE) can be modified by controlling the degree of side group branching and the miscibility between D and H chains may be enhanced. In fact, the blend of DHDPE and LLDPE was found to be almost perfectly miscible and can cocrystallize together in a common crystal lattice. In this chapter, we will describe the structure and crystallization behavior of PE blends between DHDPE and LLDPE with various degrees of ethyl branching on the basis of experimental data of infrared spectra, wide-angle (WAXD), and small-angle X-ray scatterings (SAXS), neutron scattering, thermal analysis, and so on so that the chain aggregation state can be figured out concretely at the molecular level (33–47).

5.2 COCRYSTALLIZATION AND PHASE SEPARATION OF PE BLENDS Table 5.1 shows the PE samples used in a series of our research. The blend samples were prepared between DHDPE and LLDPE with different degrees of ethyl branches. The PE components were mixed together in boiling p-xylene solutions at the concentration of about 2 wt% followed by rapid precipitation into methanol at room temperature. Figure 5.1a shows the DSC thermograms measured in the cooling process from the melt for a series of DHDPE/LLDPE(2) blend samples. The crystallization temperature is originally different by about 15
C between the pure D and H

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Table 5.1 Characterization of PE Samples. Mw DHDPE HDPE LLDPE(2) LLDPE(3)

80,000 126,000 75,000 61,000

Mn 14,000 24,000 37,000 20,000

Ethyl branching per 1000 carbon atoms 2–3 1 17 41

components. The blend sample shows a single exothermic peak. The peak temperature changes continuously in between the original two points depending on the D/H content. The infrared spectra, WAXD, and SAXS measured in the cooling process from the melt supported the DSC data (33–36). For example, Figure 5.1b shows the temperature dependence of infrared band intensity, in which the crystalline bands of

Figure 5.1 (a) DSC thermograms and (b) the temperature dependence of infrared absorbance estimated for the D and H infrared crystalline bands in the cooling process from the melt measured for a series of DHDPE/LLDPE(2) blend samples. The starting temperature of crystallization is slightly different between these two curves probably because of the difference in the cooling rate, the monitoring point of temperature, and so on. But essential behavior is the same.

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the two components start to appear simultaneously near the temperature region corresponding to the standing-up point of the DSC curve shown in Fig. 5.1a. Besides, as will be mentioned in a later section, the splitting width of infrared bands, due to the correlation coupling between the neighboring chain stems, changes continuously depending on the D/H relative content. All these experimental data allow us to apply the idea of cocrystallization to the DHDPE/LLDPE(2) blend system even when the crystallization occurs at slow cooling rate. As already pointed out, the D and H chains in the DHDPE/HDPE blends crystallize separately when the cooling rate is slow, although the degree of phase segregation depends on the blend ratio. At this point, therefore, the discovery of perfectly cocrystallizable pair of two different PE components is quite important, which can be utilized for the structural study of a single D chain embedded in a matrix of H chains without any concern about the D/H segregation problem (46). Different from the DHDPE/LLDPE(2) case, the blend of DHDPE with LLDPE(3) with higher degree of ethyl branchings shows the phase separation phenomenon. The DSC thermogram consists of the two peaks corresponding to the meltings of individual components (33–36). In the crystallization from the melt, the infrared bands and the WAXD peaks characteristic of each component are observed to appear at the different temperatures corresponding to their crystallization. In this way, the cocrystallization and phase separation are dependent on the degree of branching of the LLDPE component. This difference in the crystallization behavior affects sensitively the morphology or the chain aggregation state. Table 5.2 shows the summary of the various structural parameters estimated for these two typical blend samples (37). For example, the spherulite size (as estimated from the small-angle light scattering data) and the long period between the neighboring lamellae in the spherulite (as estimated from the SAXS data) are different depending on the blend. As for both of the spherulite radius and the long period, the DHDPE/ LLDPE(2) blend shows the intermediate size between those of the original two components. But, for the DHDPE/LLDPE(3) blend, the spherulite radius is Table 5.2

Morphology of PE Samples.

Spherulite radius, mm ˚ Long period, A ˚ Unit cell, A a b ˚ Lattice size, A (110) (200) a

LLDPE(3)

DHDPE/ LLDPE(3)

8.6 187

2.4 195

8.7 277

7.48 4.96

7.46 4.96

7.63 5.03

7.45a 4.96a

184 134

225 176

117 100

218a 190a

DHDPE

LLDPE(2)

5.3 231

9.3 158

7.42 4.93 272 216

DHDPE/ LLDPE(2)

The apparent values deduced from the overlapped profiles between DHDPE and LLDPE(3) components.

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(b) DHDPE/LLDPE(3) blend

D H

Figure 5.2 Stacked lamellar models of polyethylene blend samples.

remarkably larger than that of the original two components and the long period is also very long compared with the original ones. In the DHDPE/LLDPE(2) case, both the D and H chains coexist to form the common lamellae. In the DHDPE/LLDPE(3) case, on the contrary, the D and H species form their own lamellae individually, which are mixed together to form the stacked lamellar structure in longer period. In Fig. 5.2 are illustrated schematically the lamellar stacking modes in these samples.

5.3 AGGREGATION STRUCTURE OF CHAINS IN LAMELLA As pointed out above, the blend of DHDPE and LLDPE(2) components show almost perfect cocrystallization phenomenon. Combining the various experimental data, we can discuss the chain aggregation state in the crystal lattice for each component individually. Figure 5.3a and b shows the infrared spectra taken at room temperature for a series of DHDPE/LLDPE(2) blend samples. The infrared spectra of pure components, that is, DHDPE and LLDPE(2) show the doublet bands due to the correlation splitting. The vibrational modes, in which the two chains arrayed along the (110) direction in the orthorhombic unit cell vibrate with the phase difference of 0 and p, are optically active and give the bands at the different frequency positions. The band splitting width is dependent on the strength of interactions between the neighboring chain stems. If the D chain stems invade the (110) array of the H chain stems and disturb the interaction between the H chain stems (H    H) as illustrated in Fig. 5.4, then the vibrational coupling between the neighboring H chain stems is cut and the band splitting does not occur anymore (24–26). In other words, depending on the spatial arrangement of D and H chain stems in the crystalline lamella, the band splitting width is changed. As shown in Fig. 5.3c and d such a remarkable change of band splitting width cannot be seen in the case of DHDPE/LLDPE(3) blend samples. The systematic behavior of infrared spectra observed for DHDPE/LLDPE(2) blend samples can be interpreted reasonably as follows. Using the simply coupled

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Figure 5.3 Infrared spectra measured at room temperature for a series of (a and b) DHDPE/ LLDPE(2) blend samples and (c and d) DHDPE/LLDPE(3) blend samples. Right: the CD2 rocking bands. Left: the CH2 rocking bands.

oscillator model, we can derive the simple equation concerning the band splitting width (46, 48, 49). As shown in Fig. 5.4c, an array of H chain stems along the (110) direction is assumed to be cut by an invasion of D stems. The band splitting width Dn is given by the following equation. Dn ¼ Dn0 cos½p=ðN þ 1Þ

ð5:1Þ

where Dn0 is a splitting width of the infinitely repeated array of oscillators (H chain stems) and N is the effective number of stems in the (110) direction. If we assume that the statistically random arrangement of D and H chain stems and the distribution are

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Figure 5.4 Chain-packing modes of orthorhombic polyethylene crystal. (a) The crystal lattice consists of all the same species. (b) A single chain stem (CH2) is surrounded by CD2 stems. (c) An illustration of CH2 and CD2 stem array along the (110) direction. (From Reference 46 with permission from the Society of Polymer Science, Japan.)

not affected by the states of the neighboring chains, then the probability to have a sequence of D(H)pD is expressed by f ðpÞ ¼ ð1  XÞ2 X p1

ð5:2Þ

where X is the fraction of the H stems and the probabilities of pairs of D–H (H–D) and H–H sequences are given by 1  X and X, respectively. Then the averaged size hNi of the D(H)pD clusters is calculated as hNi ¼ Spf ðpÞ=Sf ðpÞ ¼ 1=ð1  XÞ

ð5:3Þ

where the summation is over p ¼ 1 to infinity. By substituting the hNi into Equation 5.1, we can get the band splitting width Dn as Dn ¼ Dn0 cos½pð1  XÞ=ð2  XÞ

ð5:4Þ

In a similar way, we can predict the fraction of singlet H band component for the sequence of DHD. In Equation 5.2, f ðp ¼ 1Þ is the probability of getting a single band and f ðp > 1Þ is the probability of getting doublet bands. Therefore, the singlet

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Figure 5.5 Dependence of (a) the splitting width and (b) the single band component fraction estimated for the CH2 rocking bands on the H chain fraction in the DHDPE/LLDPE(2) blends and in the blend of n-C30H62/n-C30D62. (From Reference 46 with permission from the Society of Polymer Science, Japan.)

fraction or f ðp ¼ 1Þ=½ f ðp ¼ 1Þ þ f ðp > 1Þ is given as 1  X by assuming the equal molar extinction coefficients for single and doublet bands. As shown in Fig. 5.5, where the data on the D/H blends of n-alkane (n-C30H62 and n-C30D62) are also given, these equations are found to satisfy the observed infrared data quite well. The wide-angle neutron scattering profiles measured for a series of DHDPE/ LLDPE(2) blend samples can be reproduced also reasonably by assuming the random distribution of the D and H chain stems in the lamella. In Fig. 5.6 are compared the observed and calculated wide-angle diffraction profiles measured for DHDPE/LLDPE(2) 50/50 blend sample. The neutron scattering profiles were calculated using a software Cerius2 for the model in which 20 D and 20 H chains were randomly packed in the crystal cell and each chain consisted of 20 methylene units. As seen in Fig. 5.6, the relatively intense low angle scattering is characteristic of the D and H blend lattice. In the case of DHDPE/LLDPE (3) blend, this low angle scattering becomes intense when the sample is melted, but it is quite weak when the sample is solidified because of the segregation of two components.

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Figure 5.6 Wide-angle neutron scattering profiles measured at the various temperatures for (c) DHDPE/LLDPE(2) and (d) DHDPE/LLDPE(3) blend samples in comparison with the profiles of (a) the amorphous state and (b) the crystalline state calculated for the pure DHDPE and the blend of D and H chain components. In (c), the low angle broad scattering is detected at any temperature due to the homogeneous mixing of D and H species in the same crystallite state as well as in the melt state. In (d), the low angle scattering can be seen only in the molten state, while it becomes lower when the sample crystallizes into the separated phases of D and H chains. These observations are consistent with the simulated results of (a) and (b).

From all these data analyses, we can definitely say that the D and H chain stems are distributed statistically randomly in the crystalline lamellae of the D/H cocrystallized blend. This conclusion is quite important in relation with the chain-folding problem, a controversial research theme that had been discussed for a long time (30). The random distribution of the D and H chain stems naturally supports the idea that the D and H chains reenter randomly into and out of the crystalline lamellae as shown in Fig. 5.7. The regular adjacent reentry model is impossible to apply at all as for as the melt-crystallized sample is concerned.

5.4 CRYSTALLIZATION BEHAVIOR OF D/H BLEND SAMPLES In order to investigate the cocrystallization and phase segregation behaviors of the above-mentioned PE blend samples, we performed the experiments for both the isothermal and nonisothermal crystallizations. In the nonisothermal crystallization, the temperature is changed gradually and the WAXD, SAXS, or infrared spectra are collected as a function of temperature. In the isothermal crystallization, the

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Figure 5.7 Illustration of the chain-folding models. (a) A regular adjacent reentry model, (b) a random reentry model, and (c) a cluster model consisting of the mixed structures of (a) and (b). (From Reference 46 with permission from the Society of Polymer Science, Japan).

temperature is changed rapidly from the molten state to the crystallization temperature and is kept constant for a long time, during which the WAXD, SAXS, or infrared spectra are measured as a function of time.

5.4.1 Crystallization in the Cooling Process from the Melt Figure 5.8 shows the infrared spectral change of DHDPE/LLDPE(2) blend of 75/25 wt % measured during the nonisothermal crystallization or in the cooling process from the melt (36). Generally, the infrared CH2 and CD2 bending bands show approximately the singlet and the doublet profiles, respectively. The band profiles were separated into the singlet and doublet components by a deconvolution method. As shown in Fig. 5.9a, when the crystallization started, the singlet component of the H bands appeared at first, just when the D bands showed both the doublet and singlet components at almost the same ratio. That is to say, the H chain stems in the crystal lattice are isolated from each other by being surrounded by D chain stems at the early stage of cocrystallization. As the temperature deceased furthermore, the H and D

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Figure 5.8 Temperature dependence of infrared spectra measured for DHDPE/LLDPE(2) 75/25 blend sample during the cooling process from the melt. In these profiles, the contribution of the amorphous band was subtracted. (From Reference 36 with permission from the American Chemical Society.)

chain stems gathered randomly around the initially created nuclei. In the case of DHDPE/LLDPE(3) blend, on the contrary, the situation was a little different. As mentioned above, the D and H species crystallize separately in the different lamellae. But, more strictly speaking, some parts of the H species are trapped in the D crystallites. As shown in Fig. 5.9b, at the early stage of crystallization, the singlet band of H chain stems was detected in the crystallization temperature region of the D species. When the temperature was lowered down to 90
C, where the main part of LLDPE(3) started to crystallize, both the doublet and singlet band components were observed to increase in intensity. That is to say, some portion of H chains was induced to cocrystallize with the surrounding D chains. Below 90
C, the H chain stems crystallized separately from the growing CD2 crystals, as illustrated in Fig. 5.9b. This observation may be due to more or less inhomogeneous distribution of side group branchings along the skeletal chain; the parts with lower degree of branching can be cocrystallized with the D chains likely as the case of DHDPE/LLDPE(2) blend, but the parts with higher branchings are separated from the D chains.

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Figure 5.9 Schematic illustration of aggregation behavior of D (o) and H (.) chain stems in the crystallite during the crystallization process from the melt. (a) DHDPE/LLDPE(2) and (b) DHDPE/ LLDPE(3).

5.4.2 Isothermal Crystallization Process Before the description of crystallization behavior of DHDPE/LLDPE(2) blend system, it is valuable to review the structural evolution process of LLDPE(2) itself during the isothermal crystallization process from the melt (44–46). The timedependent measurement of infrared spectra was performed during the temperature jump from 160 to 105
C, for example, where the degree of supercooling DT, evaluated as a difference between the actually-set crystallization temperature and the starting point of crystallization, was about 4
C. The data were interpreted on the basis of the information about the infrared bands characteristic of such various types of molecular conformation as all-trans, trans-gauche, and so on. After the temperature jump, the infrared bands (1368 cm1, etc.) characteristic of the so-called conformationally disordered trans-form consisting of long trans-segment with invasion of some gauche bonds were found to generate as shown in Fig. 5.10, where the time ¼ 0 s was defined as the time at which the temperature reached just the predetermined crystallization temperature. After that, these infrared bands started to decrease in intensity and the infrared bands intrinsic to the orthorhombic crystal form or the long trans-zigzag bands started to increase. Therefore, it is considered that the conformationally disordered but not perfectly random chain stems are generated at first from the melt. After that, they are regularized to the longer

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Figure 5.10 Time dependence of the integrated intensity estimated for the infrared bands characteristic of the partially disordered trans-form (1368 cm1) and the regular trans-zigzag form (728 cm1) in the isothermal crystallization process from the melt. The time 0 s was defined as the time when the temperature reached just the predetermined crystallization point (Tc ). (From Reference 44 with permission from the Society of Polymer Science, Japan.)

trans-zigzag stems in the orthorhombic crystal. As shown in Fig. 5.11, we performed the time-resolved WAXD measurement during the isothermal crystallization by using a synchrotron radiation source. In the early crystallization process, the amorphous halo peak was observed and shifted to higher scattering angle with the passage of time (47). This indicates the closer packing of the random coils in the melt. This is almost detected at the same time when the disordered trans-zigzag bands were observed. It might be speculated that the partially regularized trans-zigzag segments are created in the amorphous coils and these segments gather gradually to form a

Figure 5.11 Time-dependent wide-angle X-ray scattering pattern taken for LLDPE(2) in the isothermal crystallization process.

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Figure 5.12 Time dependence of small-angle X-ray scattering patterns measured for (a) LLDPE(2) and (b) DHDPE in the isothermal crystallization process from the melt. It should be noticed that the intensity exchange (at (b)) can be observed between the firstly appeared peak and the secondly appeared peak in the crystallization process.

closer packing of chains in the amorphous region. Although the X-ray diffraction peaks characteristic of the hexagonal lattice was not detected clearly before the appearance of the peaks of the orthorhombic cell, there should exist more or less ordered structure consisting of partially ordered trans-segments, which might be similar to the nematic liquid–crystalline state. This structure changes finally to the orthorhombic lattice. The time-resolved SAXS measurements were made in parallel (45). As shown in Fig. 5.12a, immediately after the temperature jump (DT ¼ 4
C), the invariant Q, which is approximately proportional to the total crystallinity, could be detected at almost the same time when the partially disordered trans-bands (Fig. 5.13a) appeared. After that the SAXS peak corresponding to the stacked ˚ long period was detected. The lamellar thickness lamellar structure with about 800 A also increased gradually. At around 150 s after the jump, another SAXS peak ˚ long period was detected as a shoulder and increased corresponding to the 400 A in intensity, while the original SAXS peak decreased. This SAXS profile change was detected more clearly in the isothermal crystallization experiment of DHDPE as shown in Fig. 5.12b. The similar phenomenon was observed for the isothermal crystallization of polyoxymethylene (POM) (50–52). The amorphous region sandwiched between the neighboring lamellae crystallizes to a new lamella, resulting in ˚ long period. This type of stacked lamellar the observation of the structure of 400 A structure formation was proposed for the LLDPE/LDPE blends (2–6, 11, 12). The intensity exchange between these two SAXS peaks in Fig. 5.12 is not perfect but occurs only partly, indicating that the secondary crystallization occurs only locally at

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Figure 5.13 Comparison of vibrational spectroscopic data with SAXS data collected for (a) LLDPE (2) and (b) DHDPE samples in the isothermal crystallization process from the melt. Q: invariant, hdi: lamellar thickness, Li and I(Li): the long period and the corresponding SAXS peak intensity of lamellar ˚ period and L2 for 400 A ˚ period. stacking structure: L1 for 800 A

the various positions of original lamellae, as reported for the isothermal crystallization of POM (50–52). In parallel to these studies, we evaluated the radius of gyration Rg for D and H chains in the DHDPE/LLDPE(2) blend samples by measuring the small-angle neutron scattering at various temperatures (46). As shown in Fig. 5.14, Rg takes essentially the same value in both the melt and the solid state. In other words, the spatial size of a chain is not significantly affected even when drastic environmental changes occur in the crystallization process from the melt to the stacked lamellar structure. By combining the data shown in Fig. 5.13 (for both LLDPE(2) and DHDPE) and the information on Rg , we can describe the structural change of a chain during the crystallization process as illustrated in Fig. 5.15. The D and H chains are mixed homogeneously in the molten state. Once the crystallization starts to occur, the conformational ordering starts to occur and the partially disordered trans-segments are generated in the random coil, the details of which will be described in a later section. These partially disordered trans-chain segments are gathered together to form the lamella with Rg remaining almost unchanged. The D and H chains are randomly mixed and aggregate together with the random folding mode. As a result, the D and H chain stems are randomly mixed in the crystalline lattice. The lamellae

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Figure 5.14 Temperature dependence of the radius of gyration (Rg ) estimated for the D (solid symbol) and H (open symbol) chains of the DHDPE/LLDPE(2) blend samples with the different D/H ratios. (From Reference 46 with permission from the Society of Polymer Science, Japan.)

Figure 5.15 Model of structural evolution process in the isothermal crystallization of polyethylene chains. The random chain segments change to the locally regularized trans-form. The stacked lamellar ˚ . The structure changes further to the stacked lamellar structure is formed with the long period 800 A ˚ period by inserting the new lamella into the original lamellae. It is noted that the structure of 400 A radius of gyration is kept unchanged during this process. (From Reference 46 with permission from the Society of Polymer Science, Japan.)

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˚ long period. With the passage of time, the lamellar are stacked with about 800 A ˚ long period changes to the final state of 400 A ˚ long period stacking structure of 800 A through the secondary crystallization process in the amorphous region sandwiched between the originally formed lamellae. The Rg is kept almost the same even after an occurrence of the secondary crystallization as seen from the SANS data (Fig. 5.14). On the basis of the accumulated knowledge on the isothermal crystallization behavior of LLDPE (and DHDPE) sample, we investigated the crystallization of the DHDPE/LLDPE(2) blend sample through the time-resolved infrared and SAXS measurements. Similarly to LLDPE(2), the infrared bands characteristic of partially disordered trans-segments appeared at first, suggesting the formation of closer packing structure of chain in the amorphous region as discussed in the previous paragraphs. Then the infrared bands of long trans-segments started to appear to show the regularization as the orthorhombic lattice. It was noticed that the doublet bands of the DHDPE component and the singlet of LLDPE(2) component were observed to appear at first and simultaneously. After that, both the doublet and singlet bands increased in intensity for both the D and H species. This observation is consistent with that in the nonisothermal crystallization of DHDPE/LLDPE(3) blend as described in the previous section. The H chain stems are surrounded by the D stems in a small nucleus at the early stage of crystallization, then the D and H chain stems gather together randomly around these small nuclei to produce larger crystallites (Fig. 5.9).

5.4.3 Blending Effect on Crystallization Rate In the preceding section, the structural evolution process was described at the molecular level as well as the lamellar level. In the experiment about the structural evolution in the crystallization process, we noticed the difference in crystallization rate between various kinds of PE samples. Figure 5.16 shows the increase in the crystallization sensitive infrared band intensity estimated for DHDPE, HDPE,

Figure 5.16 Time dependence of infrared band intensity measured for various PE samples during the isothermal crystallization from the melt. The degree of supercooling was about 3
C. (From Reference 38 with permission from the American Chemical Society.)

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Figure 5.17 Comparison of crystallization rate between the pure components and their blend sample in DHDPE/LLDPE(2) system in the isothermal crystallization process from the melt. (From Reference 38 with permission from the American Chemical Society.)

LLDPE(2), and LLDPE(3) during the isothermal crystallization at almost the same degree of supercooling. The increasing rate of the band intensity is higher for PE sample with lower degree of branching: HDPE > DHDPE > LLDPE(2)  LDPE(3). Figure 5.17 compares the crystallization rate between the D and H components in the DHDPE/LLDPE(2) blend and DHDPE. The D and H components in the blend, which are named D (blend) and H (blend), respectively, in this figure, are found to crystallize at almost the same rate, and this crystallization rate is much higher than that of the pure DHDPE.

5.5 MIXING BEHAVIOR OF D AND H COMPONENTS As already mentioned, a pair of DHDPE and LLDPE(2) cocrystallizes, while that of DHDPE and LLDPE(3) separates when cooled from the melt. In both the cases, however, the SANS experiment indicates that the D and H species are mixed homogeneously in the molten state (46). In what mechanism do the segregation and cocrystallization occur from such a homogeneously mixed molten state? For example, it might be reasonable to speculate that the D and H chains in the melt experience the diffusional motion at different rates and the chains of similar species gather to crystallize into their own lamellae, resulting in the phase segregation. In the cocrystallization phenomenon, the D and H chains diffuse at almost the same rate in the melt and are stabilized into the state of coexistence in the same lamella. (Of course, these two species are needed to be compatible to each other to be stabilized in the common lattice.) Therefore, we need to investigate the interpenetration behavior of the D and H chains in the molten state. We investigated the diffusion of D and H chains in the melt to understand the aggregation process of these two kinds of PE chains to form the crystalline lamellae (42). In the actual experiment, as shown in Fig. 5.18, the H and D films were contacted at the edges and the interpenetration process of D and H chains was observed by measuring the change in the infrared band

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Figure 5.18 Schematic illustration of the sample used for the IR microscopic observation at the interfacial boundary of DHDPE and LLDPE(2). (a) The cover slit for infrared spectral measurement was set at the open box position. (b) The sample was shifted from the contact position to measure the infrared spectra at the position x.

intensity of the H and D chains at the interfacial part as functions of time and position. The infrared spectral measurement at the interfacial part was made using an infrared microscope. Figure 5.19 shows the time dependence of infrared band intensity of D and H components measured for a pair of DHDPE and LLDPE(2) films at the interfacial part of the H film side. The H band decreased in intensity and the D band increased

Figure 5.19 (Left) Infrared spectral change during the interdiffusion procedure at the interfacial parts of DHDPE/LLDPE(2) films. H-side: profile changes of D and H bands measured at the LLDPE(2) side. D-side: profile changes of D and H bands measured at the DHDPE side. (Right) Time dependence of integrated intensity estimated for the D and H components at the H side.

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with time. From these curves, the diffusion constant G was evaluated as 5:0 109 cm2 s1 pon the basis of Fick’s equation of diffusion; cðx; tÞ ¼ ð1=2Þc0 erfc ½x=ð2 ðGtÞÞ where a system of two infinitely long plates connected at one edge is assumed to solve the equation of diffusion @cðx; tÞ=@t ¼ Gð@ 2 cðx; tÞ=@x2 Þ. The cðx; tÞ is the concentration at a position x and at a time t, c0 is the initial concentration, and erfc is a complementary error function: p Rt erfcðtÞ ¼ 1  ð2= pÞ 0 expðy2 Þdy. Figure 5.20 compares the infrared spectra

Figure 5.20 (a) Infrared spectra measured at the various positions for a pair of contacted DHDPE and LLDPE (2) films. (b) Infrared spectra measured for a pair of DHDPE and LLDPE (3) films. The samples were quenched after being kept at 160
C for 10 min. The measured positions are indicated by figures with micrometer unit (refer to Fig. 5.18).

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Figure 5.21 Schematic illustration of diffusion process of the D and H chain components passing through the interface during the melt. (a) DHDPE/LLDPE(2) and (b) DHDPE/LLDPE(3). When the sample was quenched from the melt, the chains were crystallized with the relative positions remaining the same.

measured for DHDPE/LLDPE(2) and DHDPE/LLDPE(3) systems at the various positions apart from the interface for the samples quenched from the molten state. In the case of DHDPE/LLDPE(2)