Biomedical composites
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Biomedical composites Edited by Luigi Ambrosio
Oxford
Cambridge
New Delhi
Published by Woodhead Publishing Limited, Abington Hall, Granta Park, Great Abington, Cambridge CB21 6AH, UK www.woodheadpublishing.com Woodhead Publishing India Private Limited, G-2, Vardaan House, 7/28 Ansari Road, Daryaganj, New Delhi – 110002, India www.woodheadpublishingindia.com Published in North America by CRC Press LLC, 6000 Broken Sound Parkway, NW, Suite 300, Boca Raton, FL 33487, USA First published 2010, Woodhead Publishing Limited and CRC Press LLC © 2010, Woodhead Publishing Limited The authors have asserted their moral rights. This book contains information obtained from authentic and highly regarded sources. Reprinted material is quoted with permission, and sources are indicated. Reasonable efforts have been made to publish reliable data and information, but the authors and the publishers cannot assume responsibility for the validity of all materials. Neither the authors nor the publishers, nor anyone else associated with this publication, shall be liable for any loss, damage or liability directly or indirectly caused or alleged to be caused by this book. Neither this book nor any part may be reproduced or transmitted in any form or by any means, electronic or mechanical, including photocopying, microfilming and recording, or by any information storage or retrieval system, without permission in writing from Woodhead Publishing Limited. The consent of Woodhead Publishing Limited does not extend to copying for general distribution, for promotion, for creating new works, or for resale. Specific permission must be obtained in writing from Woodhead Publishing Limited for such copying. Trademark notice: Product or corporate names may be trademarks or registered trademarks, and are used only for identification and explanation, without intent to infringe. British Library Cataloguing in Publication Data A catalogue record for this book is available from the British Library. Library of Congress Cataloging in Publication Data A catalog record for this book is available from the Library of Congress. Woodhead Publishing ISBN 978-1-84569-436-4 (book) Woodhead Publishing ISBN 978-1-84569-737-2 (e-book) CRC Press ISBN 978-1-4398-0178-9 CRC Press order number: N10042 The publishers’ policy is to use permanent paper from mills that operate a sustainable forestry policy, and which has been manufactured from pulp which is processed using acid-free and elemental chlorine-free practices. Furthermore, the publishers ensure that the text paper and cover board used have met acceptable environmental accreditation standards. Typeset by Toppan Best-set Premedia Limited Printed by TJ International Limited, Padstow, Cornwall, UK
Contents
Contributor contact details Preface
Part I Introduction to biocomposites 1
1.1 1.2 1.3 1.4 1.5 1.6 1.7 2
2.1 2.2 2.3 2.4 2.5
Natural composites: structure–property relationships in bone, cartilage, ligament and tendons M. Purbrick and L. Ambrosio, National Research Council, Italy; M. Ventre and P. Netti, University of Naples ‘Federico II’, Italy Introduction Bone Cartilage Tendons and ligaments Conclusions: implications for tissue regeneration and tissue repair Sources of further information References Design and fabrication of biocomposites L. K. Cardon and K. J. Ragaert, University College Ghent – Ghent University, Belgium; and R. P. Koster, Delft University of Technology, The Netherlands Introduction Production techniques for biocomposite parts Conventional composite processing techniques Solution-based techniques Solid free-form fabrication technologies
xiii xxi
1
3
3 4 11 14 18 19 22 25
25 27 27 30 33 v
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2.6
Influence of processing parameters on material characteristics of biocomposites Designing with biocomposites for medical applications Conclusions References
2.7 2.8 2.9 3 3.1 3.2 3.3 3.4 3.5 3.6 3.7 4 4.1 4.2 4.3 4.4 4.5 4.6
Hard tissue applications of biocomposites K. E. Tanner, University of Glasgow, UK Introduction Head and neck applications Axial skeleton applications Advantages in the use of composites for hard tissue applications Disadvantages in the use of composites for hard tissue applications Future trends References
44
Soft tissue applications of biocomposites M. Santin, University of Brighton, UK Introduction Composite nature of soft tissue extracellular matrices Composite biomaterials for soft tissue repair Use of biocomposites in clinical intervention for soft tissue repair Conclusions References
59
Part II Particular applications of biocomposites 5
5.1 5.2 5.3 5.4 5.5 5.6 5.7 5.8 5.9
38 40 40 42
Composite materials for bone repair L. Grøndahl and K. S. Jack, The University of Queensland, Australia Introduction Component selection and general design considerations Fabrication of particulate composites Fabrication of nano-composites Template-mediated formation of nano-composites Composite scaffolds Key challenges and concluding remarks Sources of further information and advice References
44 45 50 55 55 55 56
59 61 70 80 90 92
99 101
101 102 106 109 114 116 120 121 121
Contents 6
6.1 6.2 6.3 6.4 6.5 6.6 6.7 6.8 7
7.1 7.2 7.3 7.4 7.5 7.6 7.7 8 8.1 8.2 8.3 8.4 8.5 8.6 8.7 8.8 8.9 9 9.1 9.2
Composite coatings for implants and tissue engineering scaffolds M. Wang, The University of Hong Kong, Hong Kong Introduction Design of composite coatings Technologies for the surface modification of biomaterials Composite coatings for implants Composite coatings for tissue engineering scaffolds Concluding remarks Acknowledgements References
vii
127 127 130 133 148 161 167 169 169
Composite materials for spinal implants A. Gloria, R. De Santis, L. Ambrosio, National Research Council, Italy; and F. Causa, University of ‘Magna Graecia’, Italy Introduction Structure and function of the spine Materials and design of spinal implants: the state of the art Composite materials: basic concepts Polymer-based composite materials for spinal implants Conclusions and future trends References
178
Composites for dental applications S. N. Nazhat, McGill University, Canada Restorative dental composites Matrix monomers Reinforcing agents Silane coupling agents Classification of dental composites General property requirements of dental composites A brief overview of fibrous composites in dental applications Conclusion Sources of further information and advice
201
Acrylic bone cements for joint replacement S. Deb, King’s College London Dental Institute, UK Introduction Acrylic bone cement
210
178 179 181 185 187 195 197
201 202 204 205 205 207 208 209 209
210 213
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9.3 9.4 9.5 9.6 9.7 9.8
Properties of acrylic bone cements Radiopacifiers in bone cements Antibiotic-laden bone cements Composite cements Conclusion References
10
Composite materials for replacement of ligaments and tendons L. Ambrosio, A. Gloria, National Research Council, Italy; and F. Causa, University of ‘Magna Graecia’, Italy Introduction Ligaments and tendons: tissue biology and anatomy State of the art of proposed devices for replacement of ligaments and tendons Fibre-reinforced composite materials: fundamentals and technology Composite materials for tissue replacement and as scaffolds for tissue engineering Conclusions and future trends References
10.1 10.2 10.3 10.4 10.5 10.6 10.7 11 11.1 11.2 11.3 11.4 11.5 11.6 11.7 12
12.1 12.2 12.3 12.4 12.5
218 222 223 225 228 229
234
234 235 236 237 242 250 251
Injectable composites for bone repair P. Weiss and A. Fatimi, Université de Nantes, France Introduction Classifications of injectable bone substitute Stability, rheology and injectability of injectable bone substitutes Biological behaviors of injectable bone substitutes Injectable bone substitutes for bone tissue engineering Conclusion Bibliography
255
Composite materials for hip joint prostheses R. De Santis, A. Gloria and L. Ambrosio, National Research Council, Italy Introduction Properties of the hip joint Materials for hip arthroplasty Composite hip References
276
255 257 264 267 268 269 270
276 277 279 284 290
Contents 13
13.1 13.2 13.3 13.4 13.5 13.6 13.7 13.8 13.9 13.10
Harnessing the properties of fiber-reinforced composites in the design of tissue-engineered scaffolds A. T. DiBenedetto and L. Pinatti, University of Connecticut, USA Introduction Harnessing directional properties of biomaterials Morphology of load-bearing tissues The designer’s tools In-silico computational analysis Computer-aided tissue engineering Case studies Future trends Acknowledgement References
Part III Biocompatibility, mechanical behaviour and failure of biocomposites 14
14.1 14.2 14.3 14.4 14.5 14.6 14.7 15
15.1 15.2 15.3 15.4 15.5 15.6
The challenge of biocompatibility evaluation of biocomposites J. M. Anderson and G. Voskerician, Case Western Reserve University, USA Introduction Biocompatibility and the biological environment Surface effects and characterization Evaluation of biocompatibility: the relevance of employed analyses Surface characterization Future trends References Cellular response to biocomposites P. Jayakumar and L. Di Silvio, King’s College London Dental Institute, UK Introduction Skeletal regeneration and reconstruction Biocomposites Cellular response and experimental testing Conclusions References
ix
296
296 297 299 301 304 306 307 315 316 316
323
325
325 327 331 334 341 348 349 354
354 355 364 371 378 378
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16
Testing the in vivo biocompatibility of biocomposites R. Giardino, Rizzoli Orthopaedic Institute and Bologna University Medical School, Italy; M. Fini, N. Nicoli Aldini, A. Parrilli, Rizzoli Orthopaedic Institute, Italy Introduction Preclinical in vivo experimentation: ethical and legal requirements ISO 10993 and biocompatibility tests Extraction and sample preparation Irritation and sensitization test Systemic toxicity Genotoxicity, carcinogenicity, reproductive and development toxicity Hemocompatibility Tests for local effects after implantation Biocompatibility evaluation in pathological conditions Biofunctionality References
16.1 16.2 16.3 16.4 16.5 16.6 16.7 16.8 16.9 16.10 16.11 16.12 17
17.1 17.2 17.3 17.4 17.5 17.6 17.7 18
18.1 18.2 18.3 18.4
385
385 386 387 389 389 392 393 394 395 402 404 407
The mechanics of biocomposites L. Nicolais, University of Naples ‘Federico II’, Italy; and A. Gloria and L. Ambrosio, National Research Council, Italy Introduction Basic concepts in designing composite materials: lamina and laminate properties Short-fibre composites Particulate composites Polymer nanocomposites Conclusions References
411
Tribology of biocomposites S. Kanagaraj, Indian Institute of Technology, Guwahati, India; and M. S. A. Oliveira and J. A. de Oliveira Simões, University of Aveiro, Portugal Introduction Experimental consideration on tribological characterization of composites Tribology of polymer composites Conclusion
441
411 413 423 426 431 437 438
441 443 444 460
Contents
xi
18.5 18.6
Acknowledgements References
460 460
19
Fatigue behaviour of biocomposites A. Pegoretti, University of Trento, Italy Introduction Fundamentals of fatigue failure in polymer composites Fatigue behaviour of biocomposites for hard tissue applications Fatigue behaviour of biocomposites for soft tissue applications References
465
19.1 19.2 19.3 19.4 19.5
Part IV The future for biocomposites 20
20.1 20.2 20.3 20.4 20.5 20.6 20.7 20.8 21
21.1 21.2 21.3 21.4 21.5 21.6
Nanostructured biocomposites for tissue engineering scaffolds D. Meng and A. R. Boccaccini, Imperial College London, UK Introduction Processing of 2D topographies for assembly of 3D (biocomposite) structures Direct fabrication of surface nanotopographies in 3D structures Bio-nanocomposites: nanoparticles, nanotubes and nanofibres Sol–gel, direct growth and biomimetic approaches Bottom-up approaches Conclusions and future trends References Developing biocomposites as scaffolds in regenerative medicine A. Tampieri, S. Sprio, E. Landi and M. Sandri, National Research Council, Italy Introduction The new approach for developing biocomposites Bio-hybrid composites for bone-like scaffold Scaffolds with hierarchically organized structure: inspiration from nature Development of the three-layered osteochondral scaffold Future trends
465 466 473 491 495
507
509
509 513 516 520 534 536 537 538
547
547 549 550 558 562 565
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21.7 21.8
Acknowledgements References
22
Developing targeted biocomposites in tissue engineering and regenerative medicine J. A. Planell and M. Navarro, Institute for Bioengineering of Catalonia (IBEC), Spain Introduction Cell/material interactions Physical aspects of materials in tissue engineering Chemical aspects of materials in tissue engineering Specific processes in regenerative medicine Conclusions References
22.1 22.2 22.3 22.4 22.5 22.6 22.7 23
23.1 23.2 23.3 23.4 23.5 23.6 23.7
567 568
573
573 576 578 581 584 588 588
Ethical issues affecting the use of biocomposites L. Trommelmans and K. Dierickx, Centre for Biomedical Ethics and Law, Belgium Introduction Developments in biomedical composites Ethical challenges in the development of biomedical composites: risks, benefits and safety Therapy or enhancement? Implications of the application of biocomposites in therapy Conclusion References
593
Index
611
593 595 597 603 605 607 608
Contributor contact details
(* = main contact)
Chapter 1
Chapter 2
Dr Malcolm Purbrick* and Professor Luigi Ambrosio Institute of Composite and Biomedical Materials (IMCB-CNR) National Research Council Piazzale Tecchio 80 80125 Naples Italy E-mail: malc_purbrick@btinternet. com
Professor Ludwig K. Cardon* and K. J. Ragaert Faculty of Applied Engineering Sciences University College Ghent – Ghent University CPMT research group Schoonmeersstraat 52 B-9000 Gent Belgium E-mail:
[email protected]
Dr Maurizio Ventre and Professor Paolo Netti Interdisciplinary Research Centre on Biomaterials (CRIB) University of Naples ‘Federico II’ Piazzale Tecchio 80 80125 Naples Italy
Professor R. P. Koster Faculty of Industrial Design Engineering Delft University of Technology Landbergstraat 15 NL-2628 CE Delft The Netherlands
Chapter 3 Professor K. Elizabeth Tanner Departments of Mechanical and of Civil Engineering James Watt South Building University of Glasgow Glasgow G12 8QQ UK E-mail:
[email protected]
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xiv
Contributor contact details
Chapter 4
Chapter 7
Dr Matteo Santin School of Pharmacy and Biomolecular Sciences University of Brighton Cockcroft Building Lewes Road Brighton BN2 4GJ UK E-mail:
[email protected]
Dr Antonio Gloria*, Dr Roberto De Santis and Professor Luigi Ambrosio Institute of Composite and Biomedical Materials (IMCB-CNR) National Research Council Piazzale Tecchio 80 80125 Naples Italy E-mail:
[email protected]
Chapter 5 Dr Lisbeth Grøndahl* School of Chemistry and Molecular Biosciences The University of Queensland Cooper Rd Brisbane QLD 4072 Australia E-mail:
[email protected] Dr Kevin Jack Centre for Microscopy and Microanalysis The University of Queensland Level 1 AIBN (Building 75) Brisbane QLD 4072 Australia E-mail:
[email protected]
Chapter 6 Professor Min Wang Department of Mechanical Engineering The University of Hong Kong Pokfulam Road Hong Kong E-mail:
[email protected]
Dr Filippo Causa Department of Experimental and Clinical Medicine University of ‘Magna Graecia’ Biosciences Building, Level 4 University Campus ‘Salvatore Venuta’ Germaneto 88100 Catanzaro Italy E-mail:
[email protected]
Chapter 8 Dr Showan N. Nazhat Department of Mining and Materials Engineering McGill University MH Wong Building 3610 University Street Montreal Quebec Canada H3A 2B2 E-mail:
[email protected]
Contributor contact details
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Chapter 9
Chapter 11
Dr Sanjukta Deb Department of Biomaterials, Biomimetics and Biophotonics (B3) King’s College London Dental Institute Floor 17, Tower Wing Guy’s Hospital London Bridge London SE1 9RT UK E-mail:
[email protected]
Professor Pierre Weiss* and Dr Ahmed Fatimi INSERM, U791 and Université de Nantes Laboratoire d’Ingénierie OstéoArticulaire et Dentaire (LIOAD) 1 place Alexis Ricordeau 44042 Nantes Cedex 1 France E-mail:
[email protected]. fr
Chapter 10
Chapter 12
Professor Luigi Ambrosio and Dr Antonio Gloria Institute of Composite and Biomedical Materials (IMCB-CNR) National Research Council Piazzale Tecchio 80 80125 Naples Italy E-mail:
[email protected];
[email protected]
Professor Roberto De Santis*, Dr Antonio Gloria, Professor Luigi Ambrosio Institute of Composite and Biomedical Materials (IMCB-CNR) National Research Council Piazzale Tecchio 80 80125 Napoli Italy E-mail:
[email protected]
Dr Filippo Causa* Department of Experimental and Clinical Medicine University of ‘Magna Graecia’ Biosciences Building, Level 4 University Campus ‘Salvatore Venuta’ Germaneto 88100 Catanzaro Italy E-mail:
[email protected]
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Contributor contact details
Chapter 13
Chapter 15
Emeritus University Professor A. T. Di Benedetto* Institute of Materials Science University of Connecticut 97 North Eagleville Road Storrs CT 06269-3136 USA E-mail:
[email protected]
Prakash Jayakumar* and Dr Lucy Di Silvio Department of Biomaterials and Biomimetics Research King’s College London Dental Institute Floor 17, Guy’s Tower, Guy’s Campus London Bridge London SE1 9RT UK E-mail:
[email protected];
[email protected]
Dr Laura Pinatti Research Assistant, IMS Associates Program Institute Materials Science University of Connecticut 97 North Eagleville Road Storrs, CT 06269-3136 USA
Chapter 14 Professor James M. Anderson* and Dr Gabriela Voskerician Department of Pathology Case Western Reserve University 2103 Cornell Road WRB Rm 5105 Cleveland OH 44106-7288 USA E-mail:
[email protected]; gabriela.
[email protected]
Chapter 16 Professor Roberto Giardino*, Dr Milena Fini, Dr Nicolò Nicoli Aldini, Dr Annapaola Parrilli Preclinical and Surgical Studies Laboratory Rizzoli Orthopaedic Institute Codivilla Putti Research Institute via di Barbiano 1/10 40136 Bologna Italy E-mail:
[email protected] Professor Roberto Giardino Bologna University Medical School Bologna Italy
Contributor contact details
Chapter 17
Chapter 19
Professor Luigi Nicolais Department of Materials and Production Engineering University of Naples ‘Federico II’ Piazzale Tecchio 80 80125 Naples Italy E-mail:
[email protected]
Professor Alessandro Pegoretti University of Trento Department of Materials Engineering and Industrial Technologies via Mesiano 77 38123 Trento Italy E-mail: Alessandro.Pegoretti@ unitn.it
Dr Antonio Gloria* and Professor Luigi Ambrosio Institute of Composite and Biomedical Materials (IMCB-CNR) National Research Council Piazzale Tecchio 80 80125 Naples Italy E-mail:
[email protected];
[email protected]
Chapter 18 S. Kanagaraj Department of Mechanical Engineering Indian Institute of Technology Guwahati India E-mail:
[email protected] Monica S. A. Oliveira and Professor José António de Oliveira Simões* Department of Mechanical Engineering University of Aveiro 3810-193 Aveiro Portugal E-mail:
[email protected];
[email protected]
xvii
Chapter 20 Decheng Meng and Professor Aldo R. Boccaccini* Department of Materials Imperial College London Prince Consort Road London SW7 2BP UK E-mail: a.boccaccini@imperial. ac.uk
Chapter 21 Anna Tampieri*, Simone Sprio, Elena Landi and Monica Sandri Institute of Science and Technology for Ceramics – ISTEC National Research Council – CNR Via Granarolo 64 Faenza (RA) 48018 Italy E-mail:
[email protected];
[email protected];
[email protected];
[email protected]
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Contributor contact details
Chapter 22
Chapter 23
Professor Josep A Planell*, Melba Navarro Institute for Bioengineering of Catalonia (IBEC) Technical University of Catalonia CIBER-BBN C/Baldiri Reixac 10-12 Barcelona 08028 Spain E-mail:
[email protected] /
[email protected]; melba.
[email protected]
Dr Leen Trommelmans* and Professor Kris Dierickx Centre for Biomedical Ethics and Law Kapucijnenvoer 35/3 Box 7001 3000 Leuven Belgium E-mail: Helena.Trommelmans@ med.kuleuven.be; Kris.
[email protected]
To my father and mother, for their constant support throughout my life; and to my wife, Gabriella for her patience and understanding. Deep gratitude to Malcolm and Sue Purbrick for the support in book editing.
Preface
One of the major tasks of contemporary medicine is the restoration of human tissues and organs lost to diseases and trauma. This effort is in response to the demands of a health care system which incurs rising costs serving an ageing population, in a society where decreasing birth rate and increasing life expectancy are frequently not complemented by the maintenance of health and quality of life. In this endeavour, drug therapy and surgical treatments have been supported by the introduction of medical devices. Biomaterials, the functional components of medical devices, are used extensively in the treatment of disease, trauma and disability. Some of the most significant advances have taken place in the last 40 years, from the introduction of pioneering joint and heart valve replacements through to the development of so-called bioactive materials that interact with host tissues to assist healing. Anatomical structures consist of a composite of hard and soft tissues which differ drastically in composition, structure, and properties, and yet integrate and function together, effectively and harmoniously. Because of their ability to mimic the extracellular matrix structure, composite biomaterials have been developed to solve clinical cases in which non-healing conditions prevent tissue repair. Composite technology became available to the engineering community ca. 1972, and has since demonstrated great advantages in the attainment of weight savings without neglecting structural properties. This aspect has been very important for the aeronautical and automotive industries. Since then, composite materials with polymeric matrices have emerged as strong candidates to replace metals and ceramics in many contexts, recently including biomedical applications. In the early 1980s, polymer composites found use mainly in load-bearing applications such as hip joint, plates, intra-medullar nails. Their use was limited because of some negative results, owing to improper processing, and high production costs. These were encountered because the main biomedical companies had no adequate production facilities or appropriately trained personnel. The subsequent ongoing development of composite technology xxi
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Preface
within aeronautics will render composite technologies both more affordable and more transferable to biomedical applications. The use of composites in the biomedical field and the related concept of biomimicking were introduced by Professor Luigi Nicolais, University of Naples Federico II, Professor William Bonfield, then at Queen Mary & Westfield College, London, and Professor Garth W. Hastings. In the USA, several distinguished scientists worked in parallel with their European colleagues in contributing to the development of biocomposites. Notably, these included Professor Eric Baer, and Professor James M. Anderson (both at Case Western Reserve University) and Professor Harold Alexander, then at the Department of Bioengineering at New York University’s Hospital of Joint Diseases Orthopaedic Institute. Since then, tremendous advances have been made in composite materials and technology to overcome their initially perceived limitations. Many synthetic and natural polymers, both biodegradable and non-degradable, have been introduced. Biomaterials, in the form of matrix and reinforced (fibre and particle) systems were synthesized to control specific material properties (e.g., hydrophilic/hydrophobic domains, mechanical and degradation) and to modulate the biosignals through chemical and surface modification with biomolecules (e.g., peptides and amino acids), to mimic the environment of living tissue. These features could not be incorporated into monolithic materials, and so the composite approach is a unique solution for the design of biomaterials as required for tissue regeneration. Biocomposites may therefore be considered to be at the centre of any successful regenerative medicine strategy, providing many of the essential features and cues for the direction of the cells toward a functional outcome. Some of these functional outcomes include surface morphology and chemistry, mechanical properties, degradability, processability and tailorability. Access to nanotechnology has offered a completely new perspective to material scientists aiming to mimic the different types of extracellular matrices present in tissues. Techniques are now available that can reliably produce macromolecular structures of nanometer size which possess a finely controlled atomic composition and architecture. Conventional polymer chemistry, combined with novel methodologies such as electrospinning, phase separation, direct patterning and self-assembly, have been used to manufacture nanocomposites which can lead to the design of novel advanced bio-inspired materials able to mimic the different types of extracellular matrices. Nanocomposites are under intense investigation in regenerative medicine, with the objectives of changing the physical or chemical properties of biomaterials and guiding the activation of specific cellular signalling. This is a unique approach for designing multi-functional and cell-instructive materials.
Preface
xxiii
This book describes a wide range of nano-, micro- and macrostructured composite materials, both degradable and non-degradable, and related technologies for a number of applications. In approaching this task the idea was: (i) to bring together in a source book the available information; (ii) to include interdisciplinary aspects and competences; (iii) to include scientists worldwide. The authors who collaborate in this project have uniquely combined information and personal expertise in the field which may help to solicit and engender new ideas and set future challenges. An attempt has been made to include all the information needed to understand the topics and technology areas covered. Whilst this may, perhaps inevitably, not always have been successful, it is hoped that the book will be used and prove useful, not a perfect but a valuable contribution to a field that I believe has matured sufficiently to merit this publication. The book is divided into four parts. Part I covers features of the most frequently investigated natural tissues. A consideration of the relationship between their structures and mechanical properties relationship precedes an account of preparation technology and a review of the use of composites for hard and soft tissue applications. Part II is focused on the application of biocomposites. Some chapters relate to applications where commercialization is at mature stage (e.g., bone cement, dental, bone substitute) or at an early stage (e.g., hip joint), whilst other chapters are dedicated to new applications of composites that are at an embryonic or an early stage of application (e.g., tissue engineering, spine, injectable substitutes, ligament, tendons). Part III addresses the critical issue of biocompatibility in vitro and in vivo encompassing biosafety and cellular response of biocomposites. Here, a basic approach to design composites is described by including the appropriate methodology and parameters calculations for different composite structure reinforcement type (continuous fibres, short fibres, particles and nanostructures). Also described in a detailed manner are the wear and fatigue properties of various composites. Part IV is focused on the recent advances in biocomposites based on nanotechnology and bioinspired concepts leading to a multifunctional behaviours of composites for regenerative medicine. In addition, ethical issues are discussed. It is the objective of this book to deliver two types of result. First, information is given on the possibility of implementing composite technology for designing novel multifunctional biomaterials and their expected impact on applications. Second, the aim is to assist the reader to identify new routes, based upon nanotechnology, to develop the next generation of biocomposites. These biocomposites will help provide appropriate solutions to the problems of treating chronic disorders in an aging population by tailoring systems for specific patients and disease states.
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Preface
There are numerous people whose help I must acknowledge. First, I would like to thank the many authors and colleagues who have been involved in this project, directly or indirectly, for their superb contributions. My special thanks and enduring gratitude go to Professor Luigi Nicolais, who introduced me to this fascinating and challenging field, for his inspiration and his continuing support. Luigi Ambrosio
1 Natural composites: structure–property relationships in bone, cartilage, ligament and tendons M. P U R B R I C K and L. A M B R O S I O, National Research Council, Italy; M. V E N T R E and P. N E T T I, University of Naples ‘Federico II’, Italy
Abstract: The mechanical behaviour of biological tissues has been studied for decades, and the knowledge gained has been an endless source of inspiration for the design of novel composite materials. The unique mechanical properties of biological tissues arise from the exquisite multiscale hierarchical assembly of tissue macromolecules. Modern nanotechnology has shed great light on the nanometre scale features of biological tissues. What we learn from this not only provides a better understanding of the intimate relationship between the nanoand micrometre scale architecture and the macroscopic response of these tissues, but it also provides valuable insights into the mechanisms that nature uses to create these extraordinarily complex structures, starting from simple building blocks. This chapter reviews recent work on the multiscale characterization of biological tissues, ranging from macromolecules up to the whole tissue level. Key words: natural composites, tissue engineering, bone, cartilage, ligament, tendon.
1.1
Introduction
An understanding of the structures, and the structure–property relationships, for natural composites, i.e., the composite systems found in natural tissue, is fundamental to the design and realization of effective biomedical materials and tissue engineering constructs for their replacement. In this chapter, we will consider the structure–property relationships apparent in the natural composite material systems found in human bone, cartilage, tendon and ligament tissues. A similar approach could be taken with other tissues, such as skin, muscle, vascular tissue, the intervertebral disc and eye tissue. Some of these are briefly discussed elsewhere in this book. The first natural composite discussed in this chapter is bone tissue. A key feature of bone tissue is that it is organized as a hierarchical composite, whose material properties change with changes in scale. The nature of the hierarchical structure of bone will therefore be discussed before moving on 3
4
Biomedical composites
to consider the procedures used for the determination of structure– property relationships. The structure of bone tissue requires that it must be studied at each scale level, i.e., the nano-, micro- and macrostructural levels, to compile a representative understanding of its behaviour. This is essential in the determination of appropriate design parameters for composite material replacements that are generated using either synthetic processing or tissue engineering. There are three main types of cartilage in the human body, namely elastic cartilage, fibrocartilage and hyaline cartilage. This chapter will concentrate on hyaline cartilage (also called articular cartilage) whose most prominent role is as the surface covering for all the diarthrodial joints in the human body. Articular cartilage is, therefore, the most significant type to consider in the context of this book. The remaining two natural composites reviewed here, namely the tendon and the ligament, are dense connective tissues, whose main function is to transmit forces and displacements of the anatomical segments which they connect. Both tissues share several similarities in terms of microarchitecture and gross mechanical behaviour, and are therefore usually discussed in parallel, as they will be here. This chapter concludes by considering the implications of what is being learned about the structures and mechanical properties of the natural composites reviewed – bone, cartilage, tendon and ligament – for the field of tissue engineering, particularly with respect to the design, development and fabrication of constructs for tissue regeneration and tissue repair. Characteristic of natural systems, each of these tissues shows an exquisite elegance and refinement of design in meeting the demands of the particular application. For each tissue, and in each of a single tissue’s different locations in the body, the required performance is delivered through the specific refinement and adaptation of a hierarchical composite structure. This structural refinement and adaptation of the tissue must address all the diverse criteria – anatomical, mechanical, physiological, developmental, biochemical, sensing, signalling, and many others – pertaining to its particular application.
1.2
Bone
1.2.1 Structure, composition and properties: an overview Bone is a natural hybrid nanocomposite comprising a mineral component, plate-shaped hydroxyapatite particles, dispersed in an organic matrix, formed predominantly of oriented collagen. This structure gives bone its balance of stiffness, toughness, and vibrational damping properties (Fyhrie and Kimura, 1999). It also enables bone tissue to be highly versatile and adaptable in vivo. In the body, bone performs well in an extensive range
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of applications, with only small structural changes being required to adjust its properties to cope with many widely varying loading geometries. There have been many biomechanical and biomimetic studies, reviewed elsewhere (Fyhrie and Kimura, 1999), which provide the basis for the understanding of structure–property relationships in bone tissue. The structure and composition of bone, and their influence on its properties, are discussed below in further detail.
1.2.2 Organic matrix The polymeric matrix phase of bone is formed from the protein collagen. Macromolecular chains of collagen are arranged in the triple helix structure of tropocollagen, which is stabilized by hydrogen bonding between the amide groups and the osteoid water. The pitch of the tropocollagen α-helix is ca. 10 nm and its diameter is ca. 1.3 nm. The helical structure of tropocollagen stops the polyamide chains from collapsing into a random coil structure and so facilitates the orientation of collagen in biological tissues such as bone, ligaments and tendons. Further details may be found in a recent and comprehensive review (Boskey, 2005).
1.2.3 Mineral component of bone The mineral component of bone is an analogue of the naturally occurring mineral hydroxyapatite (HA). The unit cell of HA has the chemical formula Ca10(PO4)6(OH)2. However, extracted bone mineral shows a Ca : P molar ratio ranging from 1.3 : 1 to 1.9 : 1. Whilst this is owing, in part, to the contribution of organic phosphate in the bone matrix, it is also attributable to composition of the bone mineral itself. With respect to naturally occurring (geologic) HA, the bone mineral is a hydroxyl- and calcium-deficient, carbonated apatite. Further structural and functional detail (Boskey, 2005), and an account of the process of cell-mediated biomineralization (Gokhale et al., 2001) may be found elsewhere.
1.2.4 Hierarchical structure of bone: effect of scale on material properties Bone tissue is organized as a hierarchical composite, whose material properties change with changes in scale (Fig. 1.1). Bone tissue must, therefore, be studied at each scale level, i.e., the nano-, micro- and macrostructural levels, to compile a representative understanding of its behaviour. This is essential in the determination of appropriate design parameters for composite material replacements that are generated using either synthetic processing or tissue engineering. Scale considerations play a significant role in
6
Biomedical composites Collagen molecule
nm
Collagen fibres Haversian canal
256 nm
54 nm
Collagen fibril
Osteocyte lacuna
Haversian osteon
280 nm
Lamella
Hydroxyapatite microcrystal 1 nm
100 nm
1 μm Size
Cement line
Canaliculi 10 μm
200 μm
1.1 Hierarchical structure of human bone (Lakes, 1993).
understanding and modelling the behaviour of these calcified tissues. On the nanoscale, they are essentially material composites based on the interdigitation of the collagen, the most prevalent biopolymer in the body, and an apatitic mineralite component, the inorganic substance. These tissues then organize into microstructural composites to support loads, which is one of their primary functions. The macroscopic anisotropy of the properties of femoral cortical bone determine what both the Haversian microstructure and the appropriate nanostructural organization of collagen and apatite must be in order to confer the appropriate macroscopic function and behaviour. The formative sequence is that, firstly, the molecular structures of the collagen and apatite, the nanostructural level, must organize. Subsequently, these combine to form the Haversian system microscopic morphology into the appropriate fibre-like composite material necessary. This is essential to provide the bone microstructure with the required appropriate strength and stiffness. Many of the physiological functions and mechanical characteristics of bone are dependent on the size of the apatitic mineralites and their 3D spatial relationships to the collagen molecules in the organic matrix. Results from electron microscopy and low-angle x-ray and neutron diffraction provide clear evidence that the apatitic mineralites of bone are deposited almost exclusively within the collagen fibrils. In this case, the structure–function relationships depend on the amount and distribution of the solid mineral phase, which in turn affects the molecular conformation of the collagen molecules and the supramolecular structure of the fibrils. The apatitic mineralites are of very small size; their shape plays a key role both biologically and in providing the biomechanical properties which
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complement and enable the functional attributes of various bones. Atomic force microscopy (AFM) has been used to measure the three-dimensional shape and the size of bone mineralites. AFM data indicate that the average bovine bone mineralite becomes greater in length and width as the animal matures; these results are consistent with TEM data on isolated mineralites (Katz et al., 2007). Mineral crystals interdigitate and grow within channels and holes among the collagen fibrils, and so create the basic building blocks of the bone, namely the mineralized fibrils. These fibrils, however, are rarely found as isolated entities within tissues. Rather, they fuse together laterally and longitudinally, and the spatial arrangement of such arrays eventually governs the complex mechanical response of the tissue. A study of these structures (Weiner and Wagner, 1998) reported that the spatial arrangement of fibril arrays usually belongs to one of several categories, namely parallel fibrils, woven fibrils, radial arrays and plywood-like structures. Arrays of parallel fibrils are usually seen in mineralized tendons, in close proximity to bone-tendon insertion. Woven fibril structures are characteristic of immature bone tissues, and are soon replaced by ‘load-bearing’ bone types (Currey, 2002). The radial array is the most common mode of mineralized fibril assembly found in dentin, in which collagen fibrils are mostly distributed in planes parallel to the outer surface (Weiner et al., 1999). A plywood-like assembly is the basic structure of lamellar bone, which is the most common bone type in humans. Of all bone types, lamellar bone possesses probably the most sophisticated composite structure, which confers the bone tissue with unique anisotropic mechanical properties. It has been observed (Weiner et al., 1997) that lamellar bone is mainly constituted of a sequence of five sublayers of parallel fibres, and that each layer is oriented at 30 ° with respect to the previous layer. Because lamellae are composed of five layers, rather than six, and because the rotation of these layers always occurs in one direction, the lamellae are both imbalanced and asymmetric. These observations raise the question of how structural asymmetry and imbalance do not induce structural collapse caused by mechanical instability. It has been reported (Liu et al., 2000) that both the symmetry and balance are recovered at the osteonal level, where layers are wound in concentric cylinders, i.e., the orientation of the layers is reversed on opposite sides of the osteonal cylinder.
1.2.5 Structure–property relationships in bone By use of specific testing procedures, it is possible to determine the mechanical response of the tissue at different scale levels. The mechanical properties of constituents belonging to one level can differ greatly from the overall mechanical response.
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Biomedical composites
On a macroscopic level, bone tissue is assembled either in the form of cortical bone or cancellous (trabecular) bone. Usually, the outer shell of bone is compact (cortical bone), and is largely responsible for the supportive functions of bone. The bulk of the epiphysis (the rounded end of a long bone) and metaphysis (the wider portion of a long bone, adjacent to the epiphyseal plate), as well as that of irregular bones (e.g., vertebrae) is constituted by sponge-like matter (cancellous bone). The relative amounts of cancellous and cortical bone have been found to be site specific (Li and Jee, 2005). Both cortical and cancellous bone display a characteristic anisotropic mechanical response, which mainly arises from the spatial arrangement and interactions of their microconstituents. Several techniques for the assessment of the mechanical properties of cortical bone have been developed and documented. The most widely used methods include tension, compression, bending and torsion. Ultimate strength and elastic modulus are the parameters that are predominantly used to represent and identify the macroscopic response of cortical bone. The numerical values of these parameters can vary greatly, depending on the type of test used. Whole diaphyseal bones or ad hoc shaped specimens are often used in order to gather mechanical data on cortical bone tissue. The ultimate strength and elastic modulus of cortical bone evaluated in a tension test are lower than for similar specimens measured in compression. For example, the average values of tensile strength and elastic modulus of human long bones (femur and tibia) evaluated in compression are ∼200 MPa and ∼23 GPa, respectively (Reilly et al., 1974; Burstein et al., 1976; Cezayirlioglu et al., 1985), while those evaluated in tension are ∼144 MPa and ∼18 GPa respectively (Reilly et al., 1974; Burstein et al., 1976; Cezayirlioglu et al., 1985; Vincentelli and Grigorov, 1985; Rho et al., 1993). Bending tests could be performed either in three-point or four-point configuration. Such tests can provide reliable estimates of ultimate tensile strength and modulus, especially for beam-shaped specimens taken from the cortical tissue. However, bending is the method of choice for testing whole bones of small animals, for which it is impractical to cut specimens of the desired dimensions and shapes. The values of tensile strength and modulus are strongly affected by the geometry of the samples. Moreover, this type of test exhibits drawbacks, either in the execution or the analysis of the results. Material non-homogeneity, anisotropy and the irregular and variable shape of the cross-section pose some issues relating to reliability of the extrapolation of the relevant mechanical parameters. Therefore, bone specimens are usually shaped in the form of rods or beams, in order to reduce the effects of the above-mentioned parameters on the evaluation of relevant mechanical parameter. At bone tissue level, the strength and modulus of human long bone specimens are ∼184 MPa and ∼9.2 GPa, respectively (Choi et al., 1990; Lotz et al., 1991; Choi and Goldstein, 1992; Currey et al., 1997).
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The mechanical properties of cancellous bone are mainly determined by using tension, compression and bending tests. The values of both tensile strength and elastic modulus of cancellous bone are much lower than the corresponding values for cortical bone. For example the values of tensile strength and elastic modulus of human distal femur are ∼6 and ∼400 MPa, respectively (Kuhn et al., 1989; Odgaard et al., 1989). Several intrinsic factors affect the mechanical properties of both cortical and cancellous bone. Density, for example, greatly influences the mechanical parameters. As might be expected, a positive correlation between material density and mechanical properties has been observed for cortical bones. The modulus and the strength of cancellous bones are also strongly correlated with both mineral density and apparent density. However, the mechanical properties of cortical bone appear to be more affected by changes in material density than those of cancellous bone. Several techniques have been reported (Griffith and Genant, 2008) for the determination of material density, including radiography, x-ray absorptiometry, computed tomography and magnetic resonance imaging. Bone, both cortical and cancellous, exhibits anisotropic and heterogeneous (site specific) mechanical properties. This arises from the spatial arrangement and distribution of its microconstituents, basically mineral crystals and collagen fibres. Isolated specimens of the cortical bones display higher stiffness in the longitudinal direction (parallel to the loadbearing direction). In contrast, stiffness is lower in the transverse direction. In more detail, cortical bone may be considered to behave as a transversely isotropic material (Katz and Yoon, 1984). This peculiar behaviour arises from the spatial arrangement of collagen fibres and osteons, which are longitudinally oriented, i.e., parallel to the long axis of the diaphysis. The elastic modulus of cortical bone tested in the longitudinal direction is approximately 50% greater than in the transverse direction; this is shown in Table 1.1. Also, it has been observed (Bonfield and Grynpas, 1977) that the elastic modulus of bovine cortical bone varies with the testing direction, from the longitudinal orientation to the transverse orientation, according to a non-uniform trend. These results, however, cannot be simply interpreted on the basis of conventional models of fibre-reinforced composites. Strength values of cortical bone also depend on the loading direction in a similar fashion to that displayed by the elastic modulus. For example, both the tensile strength and compressive strength of a specimen loaded in the longitudinal direction are approximately 2.5 times and 1.5 times greater, respectively, than those loaded in the transverse direction (Njeh et al., 2003). Owing to its highly porous structure, the mechanical anisotropy of cancellous bone arises mainly from the spatial arrangement of its individual microstructural components, such as trabecular struts and plates.
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Biomedical composites
Table 1.1 Bone tissue elastic constants for bone tissue in various organs and sites Year1
1975
1984
1996
2003
Organ
Fm
Fm
Tb
Mn
Site
A
Method
M
Mode
C
T
E12,3 E22,3 G122,3 ν122,3 ν212,3
18.2 11.7 3.3 0.63 0.38
17.7 12.8 3.3 0.53 0.41
U
U
20.0 13.4 6.23 0.35 0.23
20.9 11.5 5.5 0.40 0.22
L
P
Md
Al
I
S
23.0 16.0 7.2 0.34 0.24
21.1 16.8 7.4 0.41 0.32
U
20.6 11.9 5.7 0.40 0.23
21.1 12.3 5.8 0.40 0.24
21.2 12.9 6.1 0.38 0.24
25.2 12.7 6.8 0.41 0.20
1
For the sake of simplicity, only four of the measured elastic constants are included here (Reilly and Burstein, 1975; Ashman et al., 1984; Rho, 1996; SchwartzDabney and Dechow, 2003). 2 Subscript 1 refers to the direction of the axis of material symmetry; subscript 2 refers to the orthogonal direction. 3 The unit of measurement for the elastic and shear moduli is GPa. Fm femur; Tb tibiae; Mn mandible; A anterior; L lateral; P posterior; Md medial; Al alveolar; I inferior border; S symphysis; C compression; T tension; M mechanical testing; U ultrasound testing.
The intrinsic heterogeneous nature of basic bone building blocks, i.e., minerals dispersed in a proteic fibril, as well as their complex spatial distribution and the presence of pores, channels and numerous interfaces, are all factors which are likely to cause bone to exhibit highly heterogeneous mechanical properties. Moreover, owing to the hierarchical structure of the tissue, mechanical heterogeneity is observed at different length scales. As we have already seen, macroscopic, i.e., tissue level, variations are observed for different anatomical locations and different loading conditions. AFMbased nanoindentation is a technique that enables the study of material and mechanical heterogeneities with a resolution of tens of nanometres, which is similar to the characteristic dimensions of the bone building blocks. Using this technique, a pattern of elasticity was elaborated on a nanometric scale for bone specimens tested in both longitudinal and transverse directions (Tai et al., 2007). Although the values of indentation moduli correlate with the macroscopic data, nanomechnical heterogeneities do not correlate directly with corresponding topographical features. Nanometric heterogeneities probably arise from an evolutionary optimization process of ductility enhancement and toughening. Furthermore, the results obtained
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suggested that nanoscale heterogeneity could be an effective means of mechanical signal amplification thus affecting cellular behaviours for the remodelling or repairing of local damage. As outlined in the previous section, many important microstructural components of bone have dimensions of only a few microns or less. The nanoindentation technique is therefore a useful tool for the assessment of mechanical properties at the osteonal, interstitial and trabecular lamellar bone levels, with a spatial resolution that can attain submicrometric dimensions. This technique allows one to determine the properties in several different directions, so gathering information relating to anisotropic behaviour, even in very small specimens such as individual trabeculae. Using this technique, it was possible to determine the mechanical response at osteonal or lamellar level. For example (Rho et al., 1993), osteons and interstitial lamellae were reported to display significantly different longitudinal stiffness (22.4 ± 1.2 and 25.7 ± 1.0 GPa, respectively.)
1.3
Cartilage
1.3.1 Overview There are three major types of cartilage in the body: elastic cartilage, fibrocartilage and hyaline cartilage. Elastic cartilage is found in the epiglottis and the eustachian tubes. Fibrocartilage exists temporarily at fracture sites, and is also permanently present in three major locations in the body, namely the intervertebral discs of the spine, as a covering of the mandibular condyle in the temporomandibular joint, and in the meniscus of the knee. This review will concentrate on the third type of tissue mentioned above, hyaline cartilage. Its most prominent role is as the surface covering for all the diarthrodial joints in the human body. Because of this, hyaline cartilage is frequently called articular cartilage, as will be the standard practice here. Hyaline cartilage also forms the growth plate which is instrumental in the growth and development of long bones during childhood. A highly significant feature of cartilage tissue at the microstructural and ultrastructural levels is the presence of water and electrolytes bound to the proteoglycans and to the collagen molecules which form the articular cartilage matrix. The fluid–solid interaction arising is a key influence on articular cartilage mechanical behaviour (Quinn and Morel, 2007; Kovach, 1996).
1.3.2 Mechanical properties of articular cartilage Articular cartilage (AC) is a heterogeneous and mechanically anisotropic tissue. Beside these complex mechanical features, AC also exhibits a time dependent (viscoelastic) behaviour and electrochemical phenomena. The
12
Biomedical composites
former arises from the fluid movement through the collagen-rich network, and from the intrinsic viscoelasticity of the collagen fibres. Electrochemical phenomena arise from interactions of the negatively charged proteoglycans (PGs) that are bound to the collagen network and the counterions dissolved in the interstitial water. In the following, the intimate relationship between AC microarchitecture and macroscopic mechanical properties under specific loading conditions will be analysed. In a physiological environment, AC has to sustain loads that are transmitted between adjacent joint surfaces. Owing to its peculiar shape and microarchitecture, it experiences locally complex and contextual stress states, such as compression, shear and tension (Mow and Guo, 2002). Compressive stresses induce a relative movement between the fluid (collagen) and solid (proteoglycan network) phases. Such movement, of course, introduces timedependent phenomena and dissipates large amounts of energy, owing to the high frictional forces involved in this mechanism. Fluid flow depends on several factors. The mesh size of the solid network, for example, is crucial. In AC, the solid phase is dense, and fluid flow is consequently hindered by the presence of highly packed collagen fibrils. In other words, high internal fluid pressures occur during compression. Moreover, mesh size does not remain constant with deformation. When the solid phase collapses under compressive strains, mesh size decreases and causes the fluid to be entrapped within the network. This, in turn, dissipates more energy and decreases permeability. The presence of negatively charged PGs also exerts an inhibiting effect on fluid flow, caused by osmotic phenomena. High compressive deformations cause a local increase in charge density, which results in enhanced osmotic pressure and constraint of fluid flow. When fluid flow vanishes, the fluid pressure is balanced by the mechanical response of the solid matrix. In this equilibrium state, the entire load is borne by the solid matrix. Accordingly, experimental tests (either confined compression or free compression) should be carried out until equilibrium is reached, in order to obtain reliable data on the mechanical behaviour of the fibrous network. For AC specimens subjected to compression, a linear relationship is observed between stress and strain up to a deformation of approximately 15–20% (Lai and Mow, 1980). Within this linear regime, swelling and recovery of initial shape is observed when the compressive load is removed. Swelling arises from the tendency of the system to attain an equilibrium between the (fixed) negatively charged PGs and the soluble counterions in the interstitial fluid. Swelling behaviour can therefore be altered by changes in the concentration of the soluble ion (Eisenberg and Grodzinsky, 1985) and by changes in the density of the PGs, which are themselves negatively charged species (Fig. 1.2). These latter two influences may be exerted individually or together. At equilibrium, with or without an external applied force, the balance between swelling and solid matrix response results in a state of pre-stress.
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Type II collagen fibril (a)
(b) Proteoglycan– hyaluronan aggragate
H2 O
–
H2O –
+ – + –
–
+
Shrinkage
–
+
– –
–
H2O
–
–
–
–
+
– –
–
–
+
Hypotonic bathing solution
–
+ +
–
–
+
(d)
–
+
– –
Hypertonic bathing solution (c)
External bathing solution
–
+
– –
–
Swelling
1.2 Swelling behaviour of articular cartilage. A Collagen (type II) fibrils interact non-covalently with matrix negatively charged proteoglycans to give an anionic nature to the collagen network. B Negative charge is balanced by counter ions. Water inflow and network swelling equilibrates the ion concentration gradient between interstitial and external solutions. C At equilibrium, osmotic pressure is balanced by matrix tension. For high cation concentration in the external solution, shown here, a modest water inflow is required, causing network shrinkage. D in contrast to C, a hypotonic external solution may cause extensive network swelling.
The underlying microstructural mechanisms that govern AC mechanical behaviour in tension are distinct from those which occur when the specimen is tested under compression. When AC is subjected to uniaxial tension, it displays the characteristic non-linear stress–strain curve, with an initial lowmodulus region that is indicative of a rearrangement of collagen fibres through the fluid viscous phase. At higher deformation, collagen fibres are aligned towards the applied load and are stretched. This reflects the stiffness of the collagen network, which depends on the physical characteristic
14
Biomedical composites 2 Tension Compression
Stress (MPa)
1.5
1
0.5
0 –0.5
–0.3
–0.1
0.1 –0.5
Strain
1.3 Tensile–compressive properties of human glenohumeral cartilage samples. The response in confined compression is in bold; unconfined tensile response is light (Huang et al., 2005).
of the fibres themselves and the interaction of the fibres with the PG reticulum. In conclusion, it is observed that the underlying mechanism by which the solid matrix at equilibrium withstands the external mechanical load, either in tension or in compression, changes dramatically. This is evident in the stretching of individual collagen fibres in tension, or collapse of the collagen network in compression. This phenomenon is manifest in the stress–strain plot of articular cartilage samples subjected to uniaxial tension or confined compression (see Fig. 1.3). The behaviour of the material is clearly bimodal, with an order of magnitude difference in tension and compression (equilibrium) moduli at the limit of zero strain.
1.4
Tendons and ligaments
Tendons and ligaments are dense connective tissues, whose main function is to transmit forces and displacements of the anatomical segments they connect. Both tissues share some similarities in terms of microarchitecture and gross mechanical behaviour, and are therefore usually discussed in parallel. Normal healthy tendons are composed of parallel arrays of collagen fibres, which are closely packed together. The fibres are mostly collagen type I (70% of the total dry weight), although there are also fibres of collagen types III and V present. The relative amount of these collagens may vary during tissue development. In particular, a larger proportion of type III and V is observed at the foetal stage. This evidence suggests that these
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Evidence: x ray EM
x ray EM
x ray
MICRO FIBRIL
SUBFIBRIL
x ray EM SEM
SEM OM
EM SEM OM
Tendon FIBRIL
Fascicle
Tropocollagen 35Å staining sites
640Å periodicity Waveform Fibroblasts or crimp structure
15Å 35Å
100–200Å
500–5000Å
50–300μ
Reticular membrane Fascicular membrane
100–500μ
Size
1.4 Hierarchical spatial arrangement of collagen in tendon.
collagens may serve as a template for the formation of mature, type I-rich, tissues (Griffith and Genant, 2008). These collagen fibres interact with other macromolecules, particularly proteoglycan, decorin and, in compressed regions of tendon, aggrecan. The reason for these interactions is not clear. There is growing evidence that certain non-collagenous molecules may have a role in regulating fibril diameter (Birk, 2001). It is assumed that, during tendon morphogenesis, fibroblasts produce the fibril precursors intracellularly. Fibrils increase in diameter in the pericellular space by accretion of extracellular collagen (Canty et al., 2004). Fibril bundles are organized to form fibres, with the elongated tenocytes closely packed between them. The collagen fibres coalesce into macroaggregates, and groups of macroaggregates are bound together by connective tissue (endotendon) to form bundles called fascicles. Finally, at the largest scale level, groups of fascicles are bound by the epitendon and peritendon to form the tendon organ. This hierarchical structure is shown in Fig. 1.4. Blood vessels may be visualized within the endotendon running parallel to collagen fibres, with occasional branching transverse anastomoses. The internal tendon bulk is thought to contain no nerve fibres, but the epi- and peritendon contain nerve endings, while Golgi tendon organs are present at the junction between tendon and muscle. It is important to note that, in the case of the ligament, bundles of collagen are present and wound helicoidally. For the anterior cruciate
16
Biomedical composites
ligament (ACL), these bundles have, by convention, been divided in two groups, namely the anteromedial and the posterolateral. These collagen bundles have great functional significance. Together and cooperatively, they enable the joint to function throughout the full range of its motion, and to resist the multidirectional loads presented to the joints. This difference in morphology from the tendon also influences the tensile mechanical response, as reported in Fig. 1.4. These aspects are a strong indication that ligament structures could not be substituted by tendon-like tissue.
1.4.1 Mechanical properties of tendons and ligaments Tendons and ligaments exert very important biomechanical functions, since they are involved in joint motion and joint stability. The knowledge of their mechanical properties is of paramount importance in several disciplines, such as orthopaedics, transplantations and tissue engineering. The assessment of the mechanical properties of both ligaments and tendons can be performed using in vivo or ex vivo tests. In vivo tests have, in principle, the advantage of taking into account loading conditions which could resemble those experienced during physical activity. However, the determination of the loads exerted on the tissue, as well as the deformations, required the use of invasive devices and could not be straightforward. Ex vivo tests, on the other hand, can be controlled and monitored more easily. Tendons and ligaments are mostly considered as tensile-bearing tissues. Therefore, uniaxial tension is a useful (in vitro) test that provides basic information on the mechanical response. However, it is not sufficient to capture the overall anisotropic behaviour of tendons and ligaments, which is a feature common to almost all soft connective tissues. In addition, more complex stress states, like compression and shear, may be experienced by the tissues in a physiological environment, and must be taken into account if a more detailed description of the mechanical response is required. A uniaxial test can be performed on isolated subunits, or on bone–ligament–bone/bone–tendon– muscle complexes. A large body of literature has been published over the past decades on the mechanical properties of, and characterization methods for, the Achilles tendon and the ACL. The Achilles tendon is often chosen as a representative model of collagen-rich tissue, owing to its very regular structure. The mechanical properties of ACL are widely documented because of its relevance in clinics since it is one of the most frequently injured ligaments of the human body. Stress–strain curves for ACL and Achilles tendon are shown in Fig. 1.5. The stress–strain relationship is not linear for either tissue, the diagrams displaying the characteristic J shape that is observed for the majority of collagen rich soft tissues. It is possible to identify some features of the stress–strain diagrams that are common for both the ACL and Achilles
Natural composites
17
50 Anterior cruciate ligament Achilles tendon
Stress (MPa)
40 30 20 10 0 0.00
0.10
0.20
0.30
0.40
Strain
1.5 Uniaxial tension stress–strain diagram for anterior cruciate ligament and Achilles tendon.
tendon. The first part of the diagram, which is usually referred to as the toe region, involves small load values. It is generally accepted that the straightening of the crimped pattern of the collagen fibres occurs during the toe region. Once straight, the fibrils begin to stretch. Larger loads are required for further elongation. Of course, not all the fibres are crimped to the same extent. Different fibre groups become straight at different elongations, i.e., the higher the elongation, the more fibres are uncrimped and therefore able to withstand loads. This process is also known as fibre recruitment, and explains the upward concavity of the stress–strain plot. Owing to the large variation in measured mechanical properties of both tendons and ligaments, it is difficult to elucidate and identify definitive features which enable discrimination between tendons and ligaments by simply comparing the mechanical response. For example, the wider toe region displayed by ACL specimens may reveal clues to the underlying microstructure. In general, ACL possesses a higher content of elastic tissue, which probably confers upon the collagen fibres a more marked crimped pattern. Also, the collagen fibres of the ligament are assembled in a less ordered arrangement than in the tendons (Silver et al., 2006). This observation is probably a consequence of the fact that most ligaments are generally subjected to more complex stress states than tendons (Netti et al., 1996). Stretching of aligned collagen fibres causes an almost linear stress–strain response. Microscopic failures of the collagenous network have been shown to occur in this linear region, which can eventually lead to rupture at sufficiently high levels of stress. Tendons and ligaments, like most biological tissues, exhibit timedependent mechanical properties. The molecular basis of the viscoelasticity of both tendons and ligaments is largely unknown. The presence of a high content of proteoglycans in the form of a hydrated matrix does, of course,
18
Biomedical composites
predispose these tissues to exhibit a viscoelastic effect, since viscous shear stresses may arise from relative gliding of collagen fibres. This occurs when the matrix acts as an elastic material which supports the fibrous elements of tendons and ligaments during small strain rate experiments (Lanir, 1978). In experiments performed on bovine ligamentum nuchae, the stress relaxation phenomenon almost vanishes after removal of collagenous components (Jenkins and Little, 1974). These data suggest that the main source of viscoelasticity arises from collagen bundles themselves, or from the interaction between collagen and the viscous matrix, or from a combination of these factors.
1.5
Conclusions: implications for tissue regeneration and tissue repair
This chapter began with our assertion that an understanding of the structures, and the structure–property relationships, for the composite systems found in natural tissue are fundamental to the design and realization of effective biomedical materials and tissue engineering (TE) constructs for their replacement. We followed this statement by considering in turn the structure–property relationships shown by natural composites in bone, cartilage, tendon and ligament tissue. After briefly examining the validity of our opening statement, this chapter will close by examining the implications of the work reviewed here, and future studies in these areas, for TE applications in orthopaedic medicine. TE aims to regenerate biological tissues in vitro by combining cells and material platforms (scaffolds) within bioreactors which provide the cellular constructs with adequate nutrients and physical stimuli. The correct balance and interplay of these three components eventually dictates the functionality of the de novo synthesized tissue. TE, as applied to orthopaedic medicine, typically focuses on the attainment of specific mechanical parameters, such as bone strength or stiffness. The ultimate function of the skeleton arises from an optimization process, whose result is a complex structure of efficient load bearing–transferring elements. The remodelling equilibrium of tissues is the key aspect of such an optimization process, with any deviation, whether pathological or traumatic, resulting in failure. The control of tissue structure, composition and macroscopic behaviour therefore requires a deep understanding of tissue morphogenesis and biomechanics. Despite the promising early results, TE constructs have, so far, been rarely implemented in clinic. This limitation mainly arises from the microstructural characteristics (histology) and macroscopic properties of the cellloaded TE construct not matching those of the native tissue. Many of the scaffolds used to date have been conventionally designed and produced. Their main aim has been to replicate macroscopic properties of the target
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19
tissue, such as shape, dimension and mechanical response, and they have rarely incorporated the nanometric features and hierarchical assemblies that are presented in natural tissues. There is, however, no doubt that, if we want to regenerate functional tissues in vitro, we necessarily have to create the correct environment for the cells to re-enact those morphogenetic events that occur in vivo during development and regeneration. This, of course, requires the assembly and choreography of the correct plethora of biochemical, mechanical, topographical cues; these must all be presented in the right place, at the right time. The achievement of this goal poses novel technological challenges, such as the development of processing technologies for producing material platforms with nanometric details. Fibre deposition modelling electrospinning and hybrid lithography are among the few technologies that can allow the production of scaffold with submicrometric features. Beside these technical issues, novel interpretative biomechanical models have to be developed in order to evaluate the potency of any tissueengineering application. As far as orthopaedics is concerned, the best way to absolutely quantify material properties of a tissue-engineered construct is through direct mechanical testing. Therefore, image analyses of tissueengineered constructs must be correlated with actual mechanical test results. The use and development of high-resolution imaging techniques, such as multiphoton excitation microscopy (Diaspro et al., 2006), fluorescence resonance energy transfer (FRET) (Lee et al., 2008), and AFM (Fisher et al., 2000) will potentially be of great benefit in the characterization of cell–receptor interactions, neo-tissue deposition and its microstructural assembly and the evolution of the mechanical properties.
1.6
Sources of further information and advice
1.6.1 Bone Ding M, Hvid I (2000), ‘Quantification of age-related changes in the structure model type and trabecular thickness of human tibial cancellous bone,’ Bone 26(3):291. Ito M, Nishida A, Koga A, Ikeda S, Shiraishi A, Uetani M, Hayashi K, Nakamura T (2002), ‘Contribution of trabecular and cortical components to the mechanical properties of bone and their regulating parameters,’ Bone 31(3):351. Laz PJ, Stowe JQ, Baldwin MA, Petrella AJ, Rullkoetter PJ (2007), ‘Incorporating uncertainty in mechanical properties for finite elementbased evaluation of bone mechanics,’ Journal of Biomechanics 40:2831. Les CM, Spence CA, Vance JL, Christopherson GT, Patel B, Turner AS, Divine GW, Fyhrie DP (2004), ‘Determinants of ovine compact bone
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viscoelastic properties: effects of architecture, mineralization, and remodelling,’ Bone 35:729. O’Neill MC, Dobson SD (2008), ‘The degree and pattern of phylogenetic signal in primate long-bone structure, Journal of Human Evolution 54:309. Otomo H, Sakai A, Ikeda S, Tanaka S, Ito M, Phipps RJ, Nakamura T (2004), ‘Regulation of mineral-to-matrix ratio of lumbar trabecular bone in ovariectomized rats treated with risedronate in combination with or without vitamin K2,’ J Bone Miner Metab 22:404. Porter D (2004), ‘Pragmatic multiscale modelling of bone as a natural hybrid nanocomposite,’ Materials Science Engineering A365:38. Rho J-Y, Kuhn-Spearing L, Zioupos P (1998), ‘Mechanical properties and the hierarchical structure of bone,’ Medical Engineering and Physics 20:92. Rho JY, Zioupos P, Currey JD, Pharr GM (2002), ‘Microstructural elasticity and regional heterogeneity in human femoral bone of various ages examined by nanoindentation,’ Journal of Biomechanics 35:189. Teo JCM, Si-Hoe KM, Keh JEL, Teoh SH (2007), ‘Correlation of cancellous bone microarchitectural parameters from microCT to CT number and bone mechanical properties,’ Materials Science Engineering C27:333. Teo JCM, Si-Hoe KM, Keh JEL, Teoh SH (2006), ‘Relationship between CT intensity, micro-architecture and mechanical properties of porcine vertebral cancellous bone,’ Clinical Biomechanics 21:235. van Lenthe GH, Huiskes R (2002), ‘How morphology predicts mechanical properties of trabecular structures depends on intra-specimen trabecular thickness variations,’ Journal of Biomechanics 35:1191.
1.6.2 Cartilage Sharma B, Elisseeff JH (2004), ‘Engineering structurally organized cartilage and bone tissues,’ Annals of Biomedical Engineering 32:148.
1.6.3 Tendon and ligament Bettinger PC, Smutz WP, Linscheid RL, Cooney WP, An K-N (2000), ‘Material properties of the trapezial and trapeziometacarpal ligaments,’ Journal of Hand Surgery 25A:1085. Boardman ND, Debski RE, Warner JJP, Taskiran E, Maddox L, Imhoff AB, Fu FH, Woo SL-Y (1996), ‘Tensile properties of the superior glenohumeral and coracohumeral ligaments,’ J Shoulder Elbow Surg 5(4):249. Buchanan CI, Marsh RL (2002), ‘Effects of exercise on the biomechanical, biochemical and structural properties of tendons,’ Comparative Biochemistry and Physiology Part A 133:1101.
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Cheng T, Gan RZ (2006), Experimental measurement and modeling analysis on mechanical properties of tensor tympani tendon, Medical Engineering and Physics 30:358. De Monte G, Arampatzis A, Stogiannari C, Karamanadis K (2006), ‘In vivo motion transmission in the inactive gastrocnemius medialis muscle– tendon unit during ankle and knee joint rotation,’ J Electromyogr Kinesiol 16:413. Fullerton GD, Amurao MR (2006), ‘Evidence that collagen and tendon have monolayer water coverage in the native state,’ Cell Biology International 30:56. Gupte CM, Bull AMJ, Thomas RdeW, Amis AA (2003), ‘A review of the function and biomechanics of the meniscofemoral ligaments,’ Arthroscopy 19(2):161. Hewitt JD, Glisson RR, Guilak F, Vail TP (2002), The mechanical properties of the human hip capsule ligaments, Journal of Arthroscopy 17(1):82. Kafka V, Jírová J, Smetana V (1995), ‘On the mechanical function of tendon,’ Clinical Biomechanics 10(1):50. Laurencin CT, Freeman JW (2005), ‘Ligament tissue engineering: an evolutionary materials science approach,’ Biomaterials 26:7530. Moglo KE, Skrazi-Adi A (2003), ‘On the coupling between anterior and posterior cruciate ligaments, and knee joint response under anterior femoral drawer in flexion: a finite element study,’ Clinical Biomechanics 18:751. Mommersteeg TJA, Blankevoort L, Huiskes R, Kooloos JGM, Kauer JMG (1996), ‘Characterization of the mechanical behaviour of human knee ligaments: a numerical-experimental approach,’ Journal of Biomechanics 29(2):151. Pena E, Calvo B, Martínez MA, Doblaré M (2007), ‘An anisotropic viscohyperelastic model for ligaments at finite strains. Formulation and computational aspects,’ International Journal of Solids and Structures 44:760. Robinson JR, Bull AMJ, Amis AA (2005), ‘Structural properties of the medial collateral ligament complex of the human knee,’ Journal of Biomechanics 38:1067. Robinson PS, Lin TW, Jawad AF, Iozzo RV, Soslowsky (2004), ‘Investigating tendon fascicle structure–function relationships in a transgenic-age mouse model using multiple regression models,’ Annals of Biomedical Engineering 32:924. Wang JH-C (2006), ‘Review: mechanobiology of tendon,’ Journal of Biomechanics 39:1563. Werner D, Kozin SH, Brozovich M, Porter ST, Junkin D, Seigler S (2003), ‘The biomechanical properties of the finger metacarpophalangeal joints to varus and valgus stress,’ Journal of Hand Surgery 28A:1044.
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1.7
References
ashman rb, cowin sc, van buskirk wc, rice jc (1984), ‘A continuous wave technique for the measurement of the elastic properties of cortical bone’, Journal of Biomechanics 17:349–361. birk de (2001), ‘Type V collagen: heterotypic type I/V collagen interactions in the regulation of fibril assembly’, Micron 32:223–237. bonfield w, grynpas md (1977), ‘Anisotropy of the Young’s modulus of bone’, Nature 270:453–454. boskey al (2005), ‘The organic and inorganic matrices’, pp 91–123 in Bone tissue engineering, Hollinger JO, Einhorn TA, Doll BA, Sfeir C (eds.), Boca Raton, CRC Press. burstein ah, reilly dt, martens m (1976), ‘Aging of bone tissue: mechanical properties’, The Journal of Bone and Joint Surgery 58:82–86. canty eg, lu y, meadows rs, shaw mk, holmes df, kadler ke (2004), ‘Coalignment of plasma membrane channels and protrusions (fibripositors) specifies the parallelism of tendon’, Journal of Cell Biology 165:553–563. cezayirlioglu h, bahniuk e, davy dt, heiple kg (1985), ‘Anisotropic yield behavior of bone under combined axial force and torque’, Journal of Biomechanics 18:61–69. choi k, goldstein sa (1992), ‘A comparison of the fatigue behavior of human trabecular and cortical bone tissue’, Journal of Biomechanics 25:1371–1381. choi k, kuhn jl, ciarelli mj, goldstein sa (1990), ‘The elastic moduli of human subchondral, trabecular, and cortical bone tissue and the size-dependency of cortical bone modulus’, Journal of Biomechanics 23:1103–1113. currey jd (2002), ‘The structure of bone tissue’, in Currey JD (ed.) Bones – structure and mechanics, Chapter 1, 3–25, Princeton, Princeton University Press. currey jd, foreman j, laketi´c i, mitchell j, pegg de, reilly gc (1997), ‘Effects of ionizing radiation on the mechanical properties of human bone’, Journal of Orthopaedic Research 15:111–117. diaspro a, bianchini p, vicidomini g, faretta m, ramoino p, usai c (2006), ‘Multiphoton excitation microscopy’, Biomedical Engineering Online, 5–36. eisenberg sr, grodzinsky aj (1985), ‘Swelling of articular cartilage and other connective tissues: electromechanochemical forces’, Journal of Orthopaedic Research 3:148–159. fisher te, marszalek pe, fernandez jm (2000), ‘Stretching single molecules into novel conformations using the atomic force microscope’, Nature Structural Biology 7:719–724. fyhrie dp, kimura jh (1999), ‘Cancellous bone biomechanics’, Journal of Biomechanics 32:1139. gokhale j, robey pg, boskey al (2001), ‘Osteoporosis’, in The Biochemistry of Bone (2nd ed.), Marcus R, Feldman D, Kelsey J, eds., San Diego, Academic Press. griffith jf, genant hk (2008), ‘Bone mass and architecture determination: state of the art’, Best Pract Res Clin Endocrinol Metab 22:737–764. huang c-y, stankiewicz a, ateshian ga, mow vc (2005), ‘Anisotropy, inhomogeneity, tension–compression nonlinearity of human glenohumeral cartilage in finite deformation’, Journal of Biomechanics 38:799–809. jenkins rb, little rw (1974), ‘A constitutive equation for parallel fibered elastic tissues’, Journal of Biomechanics 7:397–402.
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kastelic j, galeski a, baer e (1978), ‘The multicomposite structure of tendon’, Connective Tissue Research 6:11. katz jl, yoon hs (1984), ‘The structure and anisotropic mechanical properties of bone’, IEEE Transactions on Biomedical Engineering 31:878–884. katz jl, misra a, spencer p, wang y, bumrerraj s, nomura t, eppell sj, tabib-azar m (2007), ‘Multiscale mechanics of hierarchical structure/property relationships in calcified tissues and tissue/material interfaces’, Materials Science Engineering C27:450. kovach is (1996), ‘A molecular theory of cartilage viscoelasticity’, Biophysical Chemistry 59:61. kuhn jl, goldstein sa, ciarelli mj, matthews ls (1989), ‘The limitations of canine trabecular bone as a model for human: a biomechanical study’, Journal of Biomechanics 22:95–107. lai wm, mow vc (1980), ‘Drag-induced compression of articular cartilage during a permeation experiment’, Biorheology 17:111–123. lakes rs (1993), ‘Materials with structural hierarchy’, Nature 361:511–515. lanir y (1978), ‘Structure–strength relations in mammalian tendon’, Biophysical Journal 24:541–554. lee ky, kong hj, mooney dj (2008), ‘Quantifying interactions between cell receptors and adhesion ligand-modified polymers in solution’, Macromolecular Bioscience 8:140–145. li xj, jee wss (2005), ‘Integrated bone tissue anatomy and physiology’, chapter 2, 11–56 in Deng HW, Liu YZ, Guo CY, Chen D (eds) Current topics in bone biology, World Scientific. liu d, wagner hd, weiner s (2000), ‘Bending and fracture of compact circumferential and osteonal lamellarbone of the baboon tibia’, Journal of Materials Science: Materials in Medicine 11:49–60. lotz jc, gerhart tn, hayes wc (1991), ‘Mechanical properties of metaphyseal bone in the proximal femur’, Journal of Biomechanics 24:317–329. mow vc, guo xe (2002), ‘Mechano-electrochemical properties of articular cartilage: their inhomogeneities and anisotropies’, Annual Reviews of Biomedical Engineering 4:175–209. netti pa, d’amore a, ronca d, ambrosio l, nicolais l (1996), ‘Structure–mechanical properties relationship of natural tendons and ligaments’, Journal of Materials Science Materials in Medicine 7:525–530. njeh cf, nicholson ph, rho jy (2003), ‘Mechanical testing’, in Langton CM and Njeh CF (eds), chapter 5, 125–174, The physical measurement of bone, London, Institute of Physics. odgaard a, hvid i, linde f (1989), ‘Compressive axial strain distributions in cancellous bone specimens’, Journal of Biomechanics 22:829–835. quinn tm, morel v (2007), ‘Microstructural modeling of collagen network mechanics and interactions with the proteoglycan gel in articular cartilage’, Biomech Model Mechanobiol 6:73. reilly dt, burstein ah, frankel vh (1974), ‘The elastic modulus for bone’, Journal of Biomechanics 7:271–275. reilly dt, burstein ah (1975), ‘The elastic and ultimate properties of compact bone tissue’, Journal of Biomechanics 8:393–405. rho jy (1996), ‘An ultrasonic method for measuring the elastic properties of human tibial cortical and cancellous bone’, Ultrasonics 34:777–783.
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rho jy, ashman rb, turner ch (1993), ‘Young’s modulus of trabecular and cortical bone material: ultrasonic and microtensile measurements’, Journal of Biomechanics 26:111–119. schwartz-dabney cl, dechow pc (2003), ‘Variations in cortical material properties throughout the human dentate mandible’, American Journal Physical Anthropology 120:252–277. silver fh, freeman jw, bradica g (2006), ‘Structure and function of ligaments, tendons, and joint capsule’, in Walsh WR (ed.) chapter 2, 15–48 Repair and regeneration of ligaments, tendons and joint capsule, Humana Press. tai k, dao m, suresh s, palazoglu a, ortiz c (2007), ‘Nanoscale heterogeneity promotes energy dissipation in bone’, Nature Materials 6:454–462. vincentelli r, grigorov m (1985), ‘The effect of Haversian remodeling on the tensile properties of human cortical bone’, Journal of Biomechanics 18:201–207. weiner s, arad t, sabanay i, traub w (1997), ‘Rotated plywood structure of primary lamellar bone in the rat: orientations of the collagen fibril arrays’, Bone 20:509–514. weiner s, wagner hd (1998), ‘The material bone: structure–mechanical function relations’, Annual Reviews of Material Science 28:271–298. weiner s, traub w, wagner hd (1999), ‘Lamellar bone: structure–function relations’, Journal of Structural Biology 126:241–255.
2 Design and fabrication of biocomposites L. K. C A R D O N and K. J. R AG A E RT, University College Ghent – Ghent University, Belgium; and R. P. K O S T E R, Delft University of Technology, The Netherlands
Abstract: Some basic concepts in construction of biocomposites are discussed, such as the different types of biocomposites, the composing material classes and the general families of production techniques. The conventional, solution-based and solid freeform fabrication processing technologies for biocomposites are described and their applicability and base material requirements, i.e. whether they employ solutions or the undiluted composites/compounds are explored. The technologies discussed include extrusion/injection, filament winding, compression, autoclaving, infusion, solvent casting, phase separation, electrospinning and solid freeform fabrication. Every section elaborates on the named production process itself and the resulting biocomposite construct, offering an indication of which technique is most appropriate for given product demands. The influence of the processing parameters on the composite material is evaluated, and some relevant design examples are presented. In conclusion, a comprehensive overview of the different techniques and their applicability is presented in tabular form. Key words: biocomposites, fabrication technology, scaffolds, tissue engineering.
2.1
Introduction
Biocomposite structures are either parts made of biomaterial compounds with a filler element dispersed into the matrix material, or a construct with alternating sections of different materials (Fig. 2.1). For convenience, we will refer to the former as compounds and the latter as multi-material composites in this chapter. However, similar biomaterial processing technology applies to both categories. Production techniques can roughly be divided according to the form of the base material used. On the one hand, we have the solution-based techniques, for which the biomaterial is brought into a solution to ease processing. Examples are solvent casting, particulate leaching and electrospinning. On the other hand, there are those techniques that make use of the undiluted biomaterial, processed in its pure, unaltered form. Examples are sintering, plotting and filament winding. An additional distinction can be made for the composing materials themselves, largely depending on their processing temperature and plasticity. 25
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(a)
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(b)
(c)
2.1 Schematic illustration of the different types of biocomposite constructs: (a) compounds: a matrix material is reinforced by a dispersion of filler material; (b) and (c): multi-material: alternating sections of different materials, that may be compounds themselves. (Courtesy of University College Ghent.)
Let us first consider the materials that are viscous at ambient temperatures (up to 40 °C); these include both synthetic and natural polymerbased hydrogel, composite gels, slurries and pastes. They are the only material class suitable for cell or drug inclusion during material processing, and are processed with solution-based techniques. Examples include alginate gels (natural), fibrin gels (natural), poly(vinyl alcohol) gels (synthetic). A second class comprises the thermoplastic polymer materials. These require elevated processing temperatures (up to 300 °C) in order to ‘flow’ and be shaped. Well-known examples are the degradable polyesters poly(lactic acid), poly(glycolic acid), poly-ε-caprolactone, and their copolymers. They are either processed as powders or granulates above their melt temperature, or solution-based at ambient temperatures. Another type of polymer are the thermosets: these are highly cross-linked materials, that do not weaken by heating once their cross-linking bonds have been cured (either by heating, radiation or chemically). Before curing the unprocessed material is often liquid or at least malleable. Thermosets are often used as matrix material for composites. Finally, we discern the high-temperature components; these include metals and ceramics, which require melting or sintering at temperatures above 300 °C (and sometimes over 1500 °C). When used as the strengthening filler component of a biocomposite, they may influence the required processing temperature for the matrix accordingly. In this chapter, we will discuss the various biomaterial processing techniques, their applicability to biocomposite structures, and the design requirements they entail.
Design and fabrication of biocomposites
2.2
27
Production techniques for biocomposite parts
The choice of production technique is directly affected by the used material(s) and the desired properties of the final construct. Decisive factors are: • •
• •
processing temperature of the material; (when designing multi-material parts) the desired rate of control over the distribution of the different materials, both proportional and geometrical; required overall porosity and, if applicable, fibre size of the construct; required geometrical randomness or periodicity.
We will discuss the principle and applicability of solvent casting, phase separation, electrospinning, conventional extrusion and injection, filament winding, compression, autoclaving, infusion, and solid free-form techniques like sintering and 3D plotting. This is not an exhaustive overview; some techniques that are not explored in this chapter include lyophilization, gas foaming and UV polymerization.
2.3
Conventional composite processing techniques
2.3.1 Extrusion and injection for thermoplastic materials Biocomposite compounds with a thermoplastic material as matrix may be processed by ‘conventional’ techniques, that are also applied for everyday plastic products. In these techniques, the reinforcement phase of the composite contains short fibres and/or particles. Foremost among these techniques are extrusion and injection moulding, both of which rely on the plasticity of the polymer melt. The polymer granulate (pure or compound) is dropped from a hopper system into a heated barrel, where a well-designed screw feeds the plasticized material forward. With extrusion, this feed is continuous and the material is extruded through a die into a filament, foil or profile. Polymer wires for filament winding are produced like this. With injection moulding, a certain amount of material is dosed in front of the screw, which then forcefully rams the materials into the cavity of a closed mould. When the materials has cooled, the mould is opened and the product released. Biomedical components such as screws for plates are injection moulded. Processing experience with these conventional techniques can often be applied to the smaller-scale production of thermoplastic biocomposite parts. For example, the technical variant of biodegradable poly(lactic acid) is industrially used as a material for packaging and disposable parts, and is processed by both extrusion and injection. This has given biomedical
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researchers a wide reference set for the processing requirements of the material (Koster et al. 2008).
2.3.2 Filament winding Filament winding is a conventional production technique for the creation of hollow composite structures such as pipes, pressure vessels and yacht masts. The process involves the winding of filaments, if necessary under varying amounts of stress, over a male mould or mandrel. The mandrel rotates while a carriage moves horizontally, laying down fibres in the desired pattern (Guarino et al. 2008). Once the mandrel is completely covered to the desired thickness, it can be removed, leaving the hollow final product. Filament winding is well suited to automation, where the stress on the filaments can be carefully monitored. Filaments that are wound under higher stresses normally result in a final product with higher rigidity and strength; lower stresses allow for more flexibility. The deposition angle of the filaments may also be controlled in order to influence the mechanical properties of the construct (Mollica et al. 2006). The stacking sequence in which the fibre is deposited will determine the properties of the final product. A high angle ‘hoop’ will provide crush strength, while a lower angle pattern (known as closed or helical) will provide greater tensile strength. Filament winding is currently being researched for biomedical applications with thermoplastic materials (raw or previously extruded into fibres), for example in the development of artificial tendons or vascular prostheses (Jeong et al. 2007). It is quite suitable for the manufacture of multi-material constructs, by winding different material filaments.
2.3.3 Compression Compression moulding is a conventional method in which the moulding material, generally preheated, is first placed in an open, heated mould cavity. It was first developed to manufacture composite parts for metal replacement applications and typically used to make larger flat or moderately curved parts. The mould is closed with a top force or plug member, pressure is applied to force the material into contact with all mould areas, while heat and pressure are maintained until the moulding material has cured. Compression moulding produces fewer knit lines and less fibrelength degradation than injection moulding. The process employs thermoset resins containing high levels of reinforcement fillers in a partially cured stage, either in the form of granules, powders, or preforms. Compression moulding is a high-volume, high-pressure method suitable for moulding complex, high-strength reinforcements. These materials, in general, have thermophysical properties superior to those of
Design and fabrication of biocomposites
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thermoplastics, including higher heat flux levels, thermal conductivity, impact resistance, and maintenance of mechanical properties at high temperatures (Bloom et al. 2000). Similar technologies are resin transfer moulding (RTM) and sheet moulding compound (SMC). The RTM process is suitable for structural composites with a polymer matrix; these are made by injecting a system of liquid resins into a fibrous reinforcement, which is usually preformed and pre-placed in a mould. Polymerization is carried out just after injection (Ferret et al. 1998). SMC material is also usually cut from plates to fit the surface area of the mould. Afterwards, the final part is formed by the applied pressure and heated until the curing reaction occurs. Advanced composite thermoplastics can also be compression moulded with e.g. unidirectional tapes, woven fabrics or randomly orientated fibre mats into composite structures. Paired with conventional thermoplastic processing, the advantage of compression moulding is its ability to mould large and thick walled parts. Materials may be loaded into the mould either in the form of pellets or sheet, or the mould may be loaded from a plasticizing extruder. Materials are heated above their melting points, formed and cooled. The more evenly the feed material is distributed over the mould surface, the less flow orientation occurs during the compression stage.
2.3.4 Infusion The vacuum infusion process is a technique that uses vacuum pressure to infiltrate a resin into a laminated fibre reinforcement. The fibre component is placed dry into the mould and the vacuum is applied before the resin is added. Once the complete vacuum is achieved, it will apply a suction force on the resin material and draw it into the laminate via carefully placed tubing (van Rijswijk et al. 2009). An improvement of this method is to use an additional vacuum bag to suck the excess resin out of the composite matrix. In this way, the fibre-toresin ratio will be improved and will result in a stronger final product. This will also result in a less amount of wasted resin (Thagard et al. 2003).
2.3.5 Autoclaving Another technology to process composites with a thermoset or thermoplastic matrix is autoclaving. Both heat and pressure are applied to the workpiece placed inside the autoclave. The use of elevated pressure facilitates a high fibre volume fraction and low void content for maximum structural efficiency. Typically, there are two classes of autoclave: those pressurized with steam produce parts that can withstand exposure to water, whereas those with circulating heated gas have greater flexibility and control of the heating atmosphere.
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When working with (semi)crystalline thermoplastic resins it has to be considered that the cooling rate from the consolidating temperature is a critical parameter of the production process. The cooling rate has to be tailored in such a way that an acceptable degree of crystallinity is obtained and microcracking is avoided (Fernandez et al. 2003). Processing by autoclave is far more costly than standard furnace heating and is therefore generally used only when isostatic pressure must be applied to a workpiece of comparatively complex shape. In other applications, the pressure is not required by the process but is integral with the use of steam, since steam temperature is directly related to steam pressure. Compared with autoclaving, for smaller flat parts, heated presses offer much shorter cycle times.
2.4
Solution-based techniques
2.4.1 Solvent casting Solvent casting is a simple method for the creation of uncomplicated shapes like flat sheet that is executed at room temperature. Advantages include the ease of fabrication, the lack of a need for specialized equipment and the low processing temperatures. Disadvantages include the possible retention of toxic solvent within the polymer, the limited geometries attainable and the possible denaturation of natural proteins by the solvent (Rutkokswi et al. 2002). In short, a polymer is dissolved in a solvent in a concentration that will be dependent on the required viscosity and properties of the solvent-cast film. This solution is then cast onto a surface (e.g. glass plate, or Teflon) and allowed to dry. The length of the drying phase will be dependent on the volatility of the solvent and the thickness of the film. Removal of the solvent may be enhanced by a secondary drying step under vacuum. Once the solvent has evaporated from the film, it is removed from its recipient surface and ready for use. The distribution of the composing material elements is entirely dependent on the material ratios of the biocomposite before solution, and cannot be controlled geometrically during the solvent casting process. Hence, this method is suitable for the production of compound biocomposite structures, but not for multi-material products. Solvent-cast sheets can, however, be combined with other parts to create more complex constructs (Ikada 2006).
2.4.2 Phase separation Phase separation is another solution-based technique, used for the creation of highly porous parts. For a polymer solution, the phase separation technique is based on the crystallization of the solvent (Zhang and Ma 2002);
Design and fabrication of biocomposites
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Solvent
Solution
Phase separation
Solvent separation
Polymer
2.2 Schematic representation of the phase separation process. A solution undergoes a temperature-induced (freezing) phase separation, after which the solvent component is removed. (Courtesy of Sven Vernaillen, University College Ghent.)
the different processing steps are illustrated in Fig. 2.2. The solution is prepared in a container and subsequently placed in a freezer. The drop in temperature causes the solvent to crystallize and freeze, inducing a solid– liquid phase separation. This solidified mixture is transferred to a freeze dryer, where the solvent is removed from the mixture by the applied vacuum. The result is a highly porous foam-like structure. The technique is not restricted to polymer solutions alone; it can also be applied to polymer–ceramic compounds, although this complicates the fabrication process (Zhang and Ma 1999). As with solvent casting, geometrical macro-shapes are limited to simple forms (but thicker parts may be obtained) and the dispersion of the different composite material phases is dependent on the dissolved compound and cannot be controlled geometrically. Concerning the porosity of the construct, pore size and shape in themselves can largely be controlled, but if the pores are not completely interconnected, long tortuous pathways are often created, which is disadvantageous for several biomedical applications such as release of signals or nutrient flow in the biocomposite construct (Moroni et al. 2008, Rouwkema et al. 2008).
2.4.3 Electrospinning Electrospinning, illustrated in Fig. 2.3, is a well know solution-based technology for the production of thin filaments, which uses an electrical charge to fabricate very fine fibres, typically at micro or nano scale. Electrospinning shares characteristics of both electrospraying and conventional solution dry spinning of fibres. The process is non-invasive and does not require
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Biomedical composites
Mechanical driver
Taylor cone Flow meter
Straight jet
Elongating, looping and spiralling jet
High-voltage power supply
μA 1
Onset of bending instability
μA 2 Envelope cone
V Grounded collector
2.3 Schematic representation of the electrospinning process. (Courtesy of Sven Vernaillen, University College Ghent.)
the use of coagulation chemistry or high temperatures to produce solid threads from solution. This makes the process particularly suited to the production of fibres using large and complex molecules such as biocomposites (Ramakrishna and Fujihara 2005). The standard laboratory setup for electrospinning consists of a spinneret (typically a hypodermic syringe needle) connected to a high-voltage (5 to 50 kV) direct current power supply, a syringe pump, and an electrically grounded collector plate. A polymer solution, sol–gel, particulate suspension or melt is loaded into the syringe and this liquid is extruded from the needle tip at a constant rate by a syringe pump. Alternatively, the droplet at the tip of the spinneret can be replenished by feeding from a header tank providing a constant feed pressure. This constant pressure type feed works better for lower viscosity feedstocks. For biomedical biocomposites, this technology can be used either for the production of artificial organ components, scaffolds for tissue engineering, implant material, drug delivery and wound dressing. Ultrafine electrospun fibres show clear potential for the manufacture of long fibre composite materials (Lannutti et al. 2007). Applications are limited by difficulties in making sufficient quantities of fibres to make substantial large-scale designed constructions in a reasonable time scale. For this reason medical applications requiring relatively small amounts of fibre are popular areas of applications for electrospun fibre reinforced materials. Electrospinning is hence being investigated as a source of cost-effective, easy-tomanufacture wound dressings, medical implants, and scaffolds for the production of artificial human tissues, known as tissue engineering. These biocomposite structures fulfil a purpose similar to that of the extracellular
Design and fabrication of biocomposites
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matrix in natural tissue. Biodegradable polymers such as polycaprolactone are typically used for this purpose. These fibres may then be coated with collagen to promote cell interaction, although collagen has successfully been spun directly into membranes (Matthews et al. 2002). This solutionbased fabrication technology can be used either for gels or polymer-based thermoplastic biocomposites. Multi-material composites can be achieved by spinning of different materials into consecutive layers (Vaz et al. 2005).
2.5
Solid free-form fabrication technologies
A up-and-coming group of production techniques for biocomposites consists of computer numerical controlled devices for three-dimensional (3D) production of structures. This technology is known as rapid prototyping (RP) or solid freeform fabrication (SFF). SFF is a group of technologies to manufacture objects in layer-by-layer fashion from 3D CAD files. It enables creation of objects with a higher degree of complexity than is possible with traditional manufacturing methods (Leong et al. 2003, Peltola et al. 2008). The main feature of SFF is the selective solidification or deposition of the material. We will first elaborate upon the SFF principle for selective solidification of materials, and then move on to material depositioning techniques. RP is an enhanced (automated) method for fabrication of a physical model or prototype from a 3D CAD file. Each SFF technology has an almost identical build strategy. It always starts from a 3D CAD file, preferable a 3D solid file. In general, the devices describe a 2D plane of the design via an XY numerical control system. The integrated Z-axis results in the so-called 2.5D design and production method. This file is converted into an STL file. Out of this file, the so-called ‘plateau build file’ is developed and this is converted into the final machine file, similar to CNC code files. This principle is illustrated in Fig. 2.4. Even if CAD systems all have different database structures, nowadays the industrial CAD packages define a neutral representation of the model. This standard consists of a faceted triangulated approximation of the CAD model, the so called STL file. This system was developed by the Albert Consulting Group for 3D Systems Inc (Uiterschout 1988, Jacobs 1992). It can be compared with standard triangular meshing. The top surface of the element lies on the positive side denoted by a vector. The vector is perpendicular to the shell midsurface and is directed consistent with a right-hand rule system of nodal connectivity. The more triangles are used, the better an approximation of the original parts will be reached. To avoid loss of accuracy for complex curved parts, a second neutral contour-based interface can be used (Donahue and Turner 1991, Stormer 1991). The CAD model is directly sliced into successive cross-sections known as the common layer interface format (CLI) (Vail 1995). The next step in the data processing is
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Biomedical composites 3D file Export in CAD software STL file Modifications of STL files Plateau file Input of process and material parameters Machine file Rapid prototyping machine Prototype
2.4 Illustration of the general file principle for rapid prototyping apparatus, resulting in the final prototype part. (Courtesy of University College Ghent.)
the generation of the cross-section information, comprising information about the material to be solidified (mostly the border of each cross-section and the hatching pattern) and technological information on machine process parameters. Together with the slice information, the so-called SLI file will be created (SLIce file). This file is directly used to control the RP machine, as is illustrated in Fig. 2.5 The specifics of the solidification process are related to the basic material, for which three main categories exist: polymers, metal alloys and ceramics. Ceramics are always processed as powders to produce RP parts. Resin is used for binding the ceramic particles during the creation process. Usually, a post-treatment like a firing operation is needed to solidify the binder and increase the part density and mechanical properties. When using plastics, a wider range of SFF processes are used. Some of them are related to selective polymerization of photoreactive monomer liquids like acrylate and epoxy polymers. Other processes are based on thermoplastic basic materials in powder, granulates or filament form (Jacobs 1992). These processes are known as selective laser sintering (SLS). With SLS, the material is not melted; instead, powder grains of the material are fused together by heating the materials shortly just above the melt temperature, so that only the outer layer of the powder grains melts and fuses to the adjacent grain. Alternatively, the compound contains a low-melting binder materials that will melt under the influence of the locally added energy and serve as a matrix to hold the high-melting material parts (typically ceramics or metals) together. The principle of sintering is explained in Fig. 2.6.
Design and fabrication of biocomposites Medical scanners (CT, MRI ...)
Direct slicing
CAD
IGES VDA–FS step
Convert to STL format
35
CLI file
STL file
STL editing functions
STL slicing
Hatching
CLI editing functions
CLI file
RP & T machine
SLI editing functions
STL automatic support generator
2.5 Illustration of the computer-integrated manufacturing principle for the processing of a 3D object into a prototype via SFF technology, also illustrating the need for support systems, file editing and the possibility of using medical CAD data. (Courtesy of University College Ghent.)
Thermoplastic binder
Metal particle
Before SLS process
After SLS process
2.6 Illustration of the sintering principle. Only the thermoplastic binder material (or in the case of pure materials, the outer layer of the material grains) is melted in order to fuse the particles together. (Courtesy of Sven Vernaillen, University College Ghent.)
36
Biomedical composites
SFF technologies using metal alloys nowadays are based on three main principles for selective material creation (Jacobs 1992, Wohlers 2003). The first one consists of selectively sintering the powdered metal and applying the melt on the right place to build up the product. As with polymers, this process is known as SLS. The second method is similar to ceramics and starts with powders. These powders are selectively bound together by use of different binders. This method is known as 3D printing. A third selective metal creation technology is the selective melting process. During this process, known as selective laser melting (SLM), metal powder is melted selectively at the right place. SFF by material deposition largely follows the same principle as selective solidification techniques. The difference lies in that whereas for solidification, energy (or a binder agent) is added to a bed of material where the part needs to be built, deposition techniques will ‘drop’ material onto a recipient in order to build the part in a layer-wise fashion (Landers et al. 2002). Therefore, it is often referred to as 3D plotting. Figure 2.7 shows a typical 3D plotting machine for biomaterial applications. In this setup, the temperature-controlled plotting table remains stationary, while the dispense head is mobile. The software that controls xyz-movement and dis-
2.7 BioScaffolder® machine for 3D plotting: front control panel with basic controls (software controls on adjacent computer are not shown); plotting table can be seen in the centre of the picture. The dispensing head moves over the table and deposits filaments of material. The syringe dispensing head for viscous materials is shown mounted; the auger screw dispensing head for thermoplastic polymers is at stand-by in the park area (right). (Courtesy of University College Ghent.)
Design and fabrication of biocomposites
37
pensing is based on CNC software. In more advanced systems like this one, the dispense head is interchangeable and different types of material may be processed. A relatively simple ‘syringe’ dispense head (mounted on the machine in Fig. 2.7) is usable for solution-based materials like hydrogels, pastes and slurries. A sterile syringe with the viscous material is placed in the dispense head; pressurized air is connected at the top and a disposable dispensing needle of the required gauge size is screwed in at the bottom side. While the material is pushed through the needle, the machine moves the dispense head along its calculated NC path, thus depositing filaments of the materials into the layers of the construct. Alternatively, the material is pushed down by a stepper motor-controlled plunger, which offers more precise control than the pressurized air system. Figure 2.8 shows the more sophisticated ‘auger screw’ dispense head for the processing of thermoplastic polymer materials. The material batch (either powder or granulate) is placed in the steel hopper, from where it is
2.8 Close-up of the auger screw dispensing head for 3D plotting of thermoplastic polymers, while plotting PCL material into a tube. Main elements are the material feed hopper (metallic cylinder), the feed block with heater element and the stepper motor-driven auger screw (inside block) that pushes the material through the 28 gauge needle. (Courtesy of University College Ghent.)
38
Biomedical composites
351 μm 216 μm 303 μm 203 μm 229 μm Acc.V Spot Magn Det WD 5.00 kV 4.5 80x SE 14.7 scaff
500 μm
2.9 SEM image of a PLA scaffold created by 3D plotting. (Courtesy of University College Ghent.)
pushed into the heated block by pressurized air. The materials melts inside the block and is led to a channel that contains an auger screw. This screw’s rotation is controlled by a stepper motor; it will feed the material downward towards the encapsulated needle through which the material is plotted. Filament diameters are in the order of 100 to 800 μm, depending on the mounted needle. A scanning electron microscopy (SEM) image of a scaffold section is included in Fig. 2.9. 3D plotting with interchangeable tool heads is the most able technique for the development of biocomposites with alternating material sections, the so-called multi-material constructs illustrated in Fig. 2.1. The NC software allows for complete control of the geometrical shape, orientation and relative volume of the filaments of the composing material types. An experienced operator can fine-tune the process further by adapting parameters such as feed rate, temperature and pressure.
2.6
Influence of processing parameters on material characteristics of biocomposites
Processing with biocomposites requires experience in both equipment and material characteristics. Compared with conventional polymers, most biopolymers have non-identical material characteristics (Koster et al. 2008). For example, during extrusion and injection, viscosity of conventional polymers will decrease with the processing pressure. This is called shear thinning. Biopolymers on the other hand, have an almost constant viscosity/pressure profile, illustrated in Fig. 2.10. This will influence the processing conditions.
Design and fabrication of biocomposites
39
Viscosity (poise)
1.00E+05 PS
1.00E+04
PLA 1.00E+03
1.00E+02 0.01
1
100
10 000
Shear rate (rad s–1)
2.10 For processing shear rates below 100 rad s−1, biopolymer PLA is not subject to shear thinning, a typical quality of thermoplastic polymers, where the viscosity drops with higher shear rates (higher injection pressure). (Courtesy of Nick Gereels, University College Ghent.)
Tests on technical poly(lactic acid) have proven that shrinkage of biopolymers is low. This will influence the accuracy and processing stability of the materials. For this reason, precise electrically driven control systems are suggested to increase accuracy and reproducibility of the biocomposite products. Compared with standard polymers, a small temperature and pressure processing window is available for biopolymer processing. A variation of 5 °C will immediately influence the plasticizing and sintering parameters. Altering the injection or extrusion speed results in an opposite weight change compared with conventional plastic. Specifically for 3D plotting, biopolymers are very susceptible to thermal degradation during processing (Ragaert et al. 2008). The batch-type material supply entails that a larger mass of material, which is ‘waiting to be processed’, is heated above the melt temperature for a longer period. This extensive residence time affects the polymer in such a fashion that polymer chains will break and shorten, affecting the quality of the plotted filaments and the mechanical properties of the final construct (Pietrzak and Eppley 2006). When solution-based composites like pastes and hydrogels are processed, the viscosity of the solution will greatly influence the processing parameters. In turn, the viscosity is dependent on the solution grade of the composite into the solvent, which will be determined in accordance with the required composite characteristics such as pore structure and size. A larger solvent/solute ratio will result in lower viscosity of the solution and higher porosity of the construct. A specific concern for solution-based techniques
40
Biomedical composites
is the complete removal of the organic solvent in the final step of the production process (Sachlos and Czernuszka 2003). Remnants of solvent would pose a significant problem in the human body.
2.7
Designing with biocomposites for medical applications
Currently available processing technologies offer a broad variety of innovation possibilities for both industry and research environment, especially for medical applications. A wide range of combinations between needed composite structure and available processing technologies can be selected to optimize the final product design, possibly resulting in new applications with biocomposites. Some non-restrictive examples of medical applications are reinforced prostheses, bone cements and a wide variety of scaffolds for both soft and hard tissue engineering. Concerning prostheses, besides conventional metal implants, the SFF technology can offer new opportunities for fast patientspecific parts via integration of 3D CAD design, based on medical scan data. If so required, surface coating technologies can be applied to improve the biomedical interaction (Mirza 2008). For scaffolds, multi-material scaffolds offer strong advantages. For example, one material region can promote the cell attachment and proliferation, while a second material acts as a support structure to mimic the mechanical strength of, for example, the natural bone. Appropriate biocomposites can be selected to approximate the same mechanical characteristics as natural tissue (Mikos and Temenoff 2000). The applied materials for these kind of scaffolds can be divided as biological and synthetic polymers (Burg et al. 2000). Part two of this book elaborates on some very specific examples of biocomposites for medical applications.
2.8
Conclusions
Previously mentioned technologies such as electrospinning, filament winding and SFF offer many possibilities for innovation during design and fabrication with biocomposites. To facilitate the selection of the appropriate fabrication technology and biocomposites, Table 2.1 offers a comprehensive overview of technologies and material applications. This table ranks the discussed techniques according to the previously introduced concepts such as the various biocomposite types, material classes and processing type. Additionally, a few examples of applications and their specific materials are given.
x
Phase separation
Electrospinning
1
x
x
x
x
Metal/ ceramic (high T)
3D geometry (porosity per design)
3D geometry (porosity per design)
Filament/profile 3D shape (dense) Hollow shape
Non-woven material
Porous simple geometries
Porous sheets
Final product
x
x
x
x x
x
x
x
Compound
x
x
x
x
Multimaterial
Biocomposite type
PLA, PCL fibrin gels
Ti PCL CaCO3
PLLA, PCL
PCL collagen gel ePTFE PLA
PLGA collagen gels
PLGA
Material example
Products made with one technique can be combined with other techniques in order to obtain multi-material parts.
x
x
x
SFF – 3D plotting
x
x
x x
x
Filament winding SFF – SLS/SLM
x
x
x x
(x)
x
x
Extrusion injection
x
x
Solvent casting
x
x
Technique1
Thermoplast (medium T)
Gels (ambient T)
Solution based
Undiluted material
(Matrix) material class
Material processing
Table 2.1 Overview of technologies and material applications (Courtesy of University College Ghent)
Dental implants, Scaffolds for bone tissue engineering Scaffolds for tissue engineering: bone, vascular
Artificial tendons
Assembly of scaffolds for tissue engineering Scaffolds for tissue engineering: vascular, nerve guidance Scaffolds for vascular tissue engineering Vascular prosthesis Degradable screws
Application example
42
Biomedical composites
2.9
References
bloom, p. d.; baikerikar, k. g.; otaigbe, j. and sheases, v. v. (2000). ‘Development of novel polymer/quasicrystal composite materials’. Materials Science and Engineering A 294–296, 156. burg, k. j. l.; porter, s. and kellany, j. (2000). ‘Biomaterial developments for bone tissue engineering’. Biomaterials 21(23): 2347–2359. donahue, r. and turner, r. s. (1991). ‘CAD modelling and alternative methods of information transfer for rapid prototyping systems’. 2nd International Conference on Rapid Prototyping. Dayton. fernandez, i.; blas, f. and frovel, m. (2003). ‘Autoclave forming of thermoplastic composite parts’. Journal of Materials Processing Technology, 143–144, 266–269. ferret, b.; anduze, m. and nardari, c. (1998). ‘Metal inserts in structural composite materials manufactured by RTM’. Composites Part A Applied Science and Manufacturing 29(5–6): 693–700. guarino, v., causa, f. et al. (2008). ‘Polylactic acid fibre-reinforced polycaprolactone scaffolds for bone tissue engineering’. Biomaterials 29(27): 3662–3670. ikada, y. (2006). ‘Tissue engineering: fundamentals and applications’. Interface Science & Technology, vol. 8, Elsevier. jacobs, p. (1992). Rapid prototyping and manufacturing – fundamentals of stereolithography, Society of Mechanical Engineers. jeong, s. i. et al. (2007). ‘Tissue-engineered vascular grafts composed of marine collagen and PLGA fibers using pulsatile perfusion bioreactors’. Biomaterials 28(6): 1115–1122. koster, r. et al. (2008). The bio-molding test program – a collaborative activity of industry and university. 3rd Polymers and Moulds Innovations Conference. Ghent, Belgium. landers, r. et al. (2002). ‘Fabrication of soft tissue engineering scaffolds by means of rapid prototyping techniques’. Journal of Materials Science 37(15): 3107–3116. lannutti, j. et al. (2007). ‘Electrospinning for tissue engineering scaffolds’. Materials Science and Engineering C-Biomimetic and Supramolecular Systems 27(3): 504–509. leong, k. f.; cheah, c. m. and chua, c. k. (2003). ‘Solid freeform fabrication of three-dimensional scaffolds for engineering replacement tissues and organs’. Biomaterials 24(13): 2363–2378. matthews, j. a. et al. (2002). ‘Electrospinning of collagen nanofibers’. Biomacromolecules 3(2): 232–238. mikos, a. and temenoff, j. (2000). ‘Formation of highly porous biodegradable scaffolds for tissue engineering’. Electronic Journal of Biotechnology 3(2): 114–119. mirza, j. (2008). Biomedical implant surface coatings for tissue engineering. Polymers and Moulds Innovations, Ghent, Belgium. mollica, f. et al. (2006). ‘Mechanical properties and modelling of a hydrophilic composite used as a biomaterial’. Composites Science and Technology 66(1): 92–101. moroni, l. et al. (2008). ‘Integrating novel technologies to fabricate smart scaffolds’. Journal of Biomaterials Science-Polymer Edition 19(5): 543–572. peltola, s. m. et al. (2008). ‘A review of rapid prototyping techniques for tissue engineering purposes’. Annals of Medicine 40(4): 268–280.
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pietrzak, w. s. and eppley, b. l. (2006). ‘The effect of high temperature intraoperative molding on bioabsorbable PLLA–PGA craniofacial fixation’. Journal of Craniofacial Surgery 17(5): 920–925. ragaert, k. et al. (2008). Micro extrusion for the engineering of thermoplast scaffolds. 3rd International Tissue Engineering Conference, Rhodos, Greece. ramakrishna, k. and fujihara, w. (2005). An introduction to electrospinning and nanofibers, World Scientific Publishing. rouwkema, j. et al. (2008). ‘Vascularization in tissue engineering’. Trends in Biotechnology 26(8): 434–441. rutkokswi, g. et al. (2002). Processing of polymer scaffolds: solvent casting. Methods of Tissue Engineering. A. Atala and R. Lanza, Academic Press, 681–686. sachlos, e. and czernuszka, j. (2003). ‘Making tissue engineering scaffolds work. Review on the application of solid freeform technology to the production of tissue engineering scaffolds’. European Cells and Materials Journal 5: 29–40. stormer, l. (1991). Chain tolerances and accuracy in SLA building process, PhD Thesis, University of Trondheim. thagard, j. r. et al. (2003). ‘Resin infusion between double flexible tooling: prototype development’. Composites Part A Applied Science and Manufacturing 34(9): 803–811. uiterschout, i. (1988). Interim Report on Stereolithography Project, TNO. vail, n. k., ed. (1995). New Direction – Applications for Rapid Prototyping. Proceedings of the 12th International Key Technology Laser Congress. van rijswijk, k. et al. (2009). ‘Textile fiber-reinforced anionic polyamide-6 composites. Part I: The vacuum infusion process’. Composites Part a-Applied Science and Manufacturing 40(1): 1–10. vaz, c. m. et al. (2005). ‘Design of scaffolds for blood vessel tissue engineering using a multi-layering electrospinning technique’. Acta Biomaterialia 1(5): 575–582. wohlers, t. (2003). ‘Rapid prototyping, tooling and manufacturing state of the industry’. Wohlers Report. zhang, r. and ma, p. (2002). Processing of Polymer Scaffolds: Phase Separation. Methods of Tissue Engineering. A. Atala and R. Lanza, Academic Press: 715–724. zhang, r. y. and ma, p. x. (1999). ‘Poly(alpha-hydroxyl acids) hydroxyapatite porous composites for bone-tissue engineering. I. Preparation and morphology’. Journal of Biomedical Materials Research 44(4): 446–455.
3 Hard tissue applications of biocomposites K. E. TA N N E R, University of Glasgow, UK
Abstract: Composites based on carbon reinforced epoxy resin were first used clinically in the 1970s and, although these progressed to successful clinical applications, none remained in use much beyond these initial trials. The major problems were either the inability to shape the implant to fit the patient, or the method of manufacture, along with the fact that these materials were the first generation of biomedical composites and thus bioinert. In the 1980s, the second generation, that is bioactive composites, were developed and brought into clinical trial. As surgeons have been able to shape these implants to fit their patients, the application of these materials has been more successful and, being bioactive, they have led to stronger bonds between the implant and the supporting bone, thus the implants have progressed to clinical use from their initial clinical trails. However, in the majority of cases the applications have been in low-load-bearing areas such as the head, with minimal applications in the major load-bearing parts of the skeleton. The biomedical composites currently being tested in vitro and in vivo should overcome the problems of the older generation material and thus progress to clinical use in the next few years. Key words: bioactive material, bone, biocomposites, clinical tissues, implant.
3.1
Introduction
As discussed elsewhere in this book, there are a variety of reasons to use composites in biomedical applications and a wide range of composites has been developed. In summary, the first reason for using composites is that mechanical properties can be produced that either closely match or exceed those of the natural tissue being replaced. This rationale applies to both soft tissues, such as tendons (Amis et al., 1984) and the intervertebral disc (Ambrosio et al., 1998), and in hard tissues, as with bone replacement and augmentation (Bonfield et al., 1980; 1981). This matching of the mechanical properties can include the non-linear behaviour of natural tissues and optimises the load or stress transfer across the implant–tissue interface, thus reducing the risks associated with stress shielding of the natural tissue or stress concentrations at the interface. The second reason is that the biological properties of the composite can be tailored to optimise the biological response to the presence of the implant thus a strong biological interface can be produced by the body to surround and integrate the implant 44
Hard tissue applications of biocomposites
45
(Bonfield et al., 1980; 1981). The third rationale relates solely to degradable composites, which as the implants degrade, allow the gradual transfer of load from the implant into the natural tissue (Bos et al., 1987). When we consider the specific requirements for hard tissue replacement or augmentation composites, the stiffness of the device should be similar to that of the bone and no monolithic materials have stiffnesses close to those of cortical bone (Currey, 1998). Furthermore, in most cases, the strength of the device should be higher than that of the bone it is replacing or augmenting, thus if anything fails it is the bone, which can then repair. Finally, the implant should be bioactive to bone thus encouraging bone deposition onto the device. If the device is also degradable then the degradation products should be non-toxic and it would be beneficial to release products that accelerate bone healing (Waris et al., 1994). While these may sound relatively simple requirements, when we consider the mechanical loads on a bone replacement or augmentation device, it can be seen that the requirements are high. The loads on a hip joint are 3.5 times body weight during slow walking, rising to nine times body weight during a stumble (Bergmann et al., 1983), those on the knee are 2.0 to 2.5 times body weight during normal activities (Mundermann et al., 2008) and while the peak strains on the tibial shaft are 500 με during normal walking and are increased during running to 1200 με, which assuming the Young’s modulus of bone is 20 GPa equate to 10 and 24 MPa, respectively, and furthermore include a torsional component during the push-off phase of the gait cycle (Burr et al., 1996). Thus, for a typical 80-kg person, the normal loads can be substantial, for example 2.0 kN on the knee and 2.8 kN on the hip joint, while for heavy patients the forces will be increased. Finally, the fatigue resistance needs to be good, as Wallbridge and Dowson (1982) found that the average number of load cycles applied while walking drops from 2 million per year for someone in their 20s down to 0.5 million by the time people reach their 80s, with similar numbers of load cycles predicted in the upper arm by Joyce and Unsworth (2000). Given these requirements and despite the large number of composite materials both in development and undergoing in vitro and in vivo assessment, it is probably not surprising to see that only a few composite materials have been reported as being used clinically.
3.2
Head and neck applications
3.2.1 Maxillo-facial applications The maxillo-facial region has been the first application of many materials owing to the lower mechanical requirements compared with most other skeletal applications. However, if there are problems these will be more
46
Biomedical composites
obvious owing to the more superficial implantation sites and less soft tissue covering, as reported by Böstman et al. (1990) and others when degradable polylactic acid plates were first used clinically. The first application of the hydroxyapatite (HA) high-density polyethylene (HDPE) composite HAPEXTM developed by Bonfield and colleagues (Bonfield et al., 1981; 1982; Wang et al., 1994) was as orbital floor replacement device (Downes et al., 1991; Tanner et al., 1994). Two different designs were developed, the first was a simple compression moulded disc, less than 1 mm thick and approximately 15 mm in diameter, which was used to close the base of the eye socket after fracture of the orbital floor and prevent extrusion of the soft tissues into the sinus space. The advantages of the composite were that the toughness was such that the devices should be cut intra-operatively with either a scalpel or a sharp pair of scissors. The second device was larger, being a space-filling implant for patients who had lost an eye. Here the requirement was that the implant did not change volume and bonded well to the orbital floor. The previous implants were fat pads, which could resorb with time leading to loss of volume in the orbital cavity, and either glass balls or silicone pads, which did not bond with the orbital floor and therefore could extrude from the socket. In both these groups of patients, a close bond of the HAPEXTM implant to the orbital floor could be seen on computed tomography (Fig. 3.1) and manual palpation showed a stable interface. However, for commercial reasons, these devices were never used beyond the original cohort of 18 patients.
3.1 Computed tomography (CT) scan of the head of a patient 6 months after the implantation of a HAPEXTM orbital floor implant showing partial integration of the implant, showing the lower part of the brain, two orbital sockets, the nasal cavity in the centre and the two sinus cavities (darker as they are air filled). The implant can be seen in the right orbital floor (right hand side of the figure) as an additional piece of material with similar radiographic density to bone just above the sinus (Downes et al., 1991).
Hard tissue applications of biocomposites
47
Törmälä and his colleagues at Tampere University of Technology in Finland developed a series of implants of degradable self-reinforced polylactide (SR-PLLA) rods, screws and plates (Böstmann et al., 1987; Törmälä et al., 1988; Suuronen et al., 1992; Waris et al., 1994; Ashammakhi et al., 2001), which have been commercialised. The rods and screws were manufactured by compression moulding of aligned PLLA fibres. The use of the SR-PLLA screws was extended to fractures in cancellous bone of both the femoral neck and the lateral malleous of the fibula and will be discussed later. The plates were manufactured by extruding the material into a plate, these plates were die-drawn and the thin sheets were then cut to size and moulded together to produce plates that were 0.5 mm-thick and 12 mm by 40 mm in area. These implants were initially used in craniofacial surgery (Suuronen et al., 1992; Waris et al., 1994; del Campo et al., 1994), where Suuronen et al. (1992) suggest using four of the 0.5-mm thick plates which are screwed together once each plate has been bent into the required shape. The individual plates can be bent easily, but once the four plates have been screwed together they have the required stiffness and strength for clinical applications. In the study of Waris et al. (1994), one to three plates were used (Fig. 3.2), although these patients were children between 6 months and 8 years old, so lower loads would be expected. Waris et al. suggest that ensuring a good covering of ‘well-vascularised tissue’ and not using an excessively large implant will reduce the risk of a sterile inflammatory response owing to the presence of the degrading implant. del Campo et al. (1994)
(a)
(b)
3.2 (a) One of the 0.5-mm-thick self-reinforced PLLA plates used in (b) repair of the bony defects in the skull of a patient suffering from trigonocephaly treated by cranioplasty (Waris et al., 1994).
48
Biomedical composites
used these plates to fix osteotomies in the maxilla, that is the upper jaw bone, and found good post-operative stability, without complications. Most of the reports have related to the craniofacial region (Ashammakhi et al., 2001), where they have found good responses to the implants. An advantage is that when used in growing children the gradual reduction in the mechanical properties as the implant resorbs leads to less reduction in the growth of the child’s face. They are also able to shape the devices intraoperatively by warming the implants above their glass transition temperature (Tg) of approximately 60 °C. However, it should be noted that self-reinforced PLLA plates take about 5 years to degrade compared to 6 to 12 months for self-reinforced polyglycolide (PGA) plates (Suuronen et al., 1992). Zanetti and colleagues (Zanetti et al., 2001; Zanetti and Nassif, 2003) used a composite of hydroxyapatite in polycaprolactone (Piattelli et al., 1997) which was manufactured in the form of flexible sheets 0.3 to 1.2 mm thick which were used to reconstruct both the wall of the outer ear canal in 42 patients between the ages of 14 and 64 (Zanetti et al., 2001) and to repair minor defects of the base of the skull in seven patients (Zanetti and Nassif, 2003). In the ear study, after two years, the outer ear was successfully reconstructed in 37 cases (88%) with the outer wall totally reepithelialised in 33 cases (79%). There was a re-occurrence of the clinical problem in three patients and seven patients had extruded their implants. However, compared with the results of other studies, this was considered to be a successful series with three-quarters of the patients returning to normal anatomy with no remaining infection. In their skull study, the defects treated were all less than 30 mm in diameter and after 18 to 62 months follow-up there was no extrusion or foreign body reaction. In both studies, they considered these preliminary results to be encouraging. They also comment on the ease of cutting and shaping the composite implant.
3.2.2 Aural applications The clinical success of the HAPEXTM orbital implants encouraged Smith & Nephew ENT to use HAPEXTM in middle ear implants. At the time they were producing a range of middle ear implants, such as those designed by Goldenberg (1994) and Dornhoffer (1998), with hydroxyapatite heads that contacted the tympanic membrane (ear drum) and had ultrahigh molecular weight polyethylene (UHMWPE) shafts. These shafts were cut intraoperatively to the required length to fit on the staples, the last bone of the train of three bones that transmit and amplify sound vibrations from the outer ear to the inner ear. Hydroxyapatite was used for contact with the tympanic membrane as other materials are extruded out of the ear, while the UHMWPE was used as it could be cut easily in the operating theatre to enable the implant to be tailored to fit the patient. HAPEXTM was used to
Hard tissue applications of biocomposites
49
replace the UHMWPE as there was considered to be increased possibility of the stapes bonding with an HAPEXTM shaft, increasing the long term stability of the implant, the presence of the HA particles in the polyethylene made the material easier to trim intraoperatively and finally the increased density of HAPEXTM compared with UHMWPE should increase the sound transfer through the implant shaft. Currently, these devices are still available to surgeons. Goldenberg and Driver (2000) considered the clinical success of these implants, reviewing the results for 233 patients of whom 77 had their implants in situ for more than 5 years. Overall, the hearing success rate was 56.8% with implant extrusion occurring in 5.3% and visible slippage of the implant in 7.7% of patients. They concluded that the implant provided good hearing which was stable with time and that the extrusion rate was low. Meijer et al. (2002) reviewed the histological response to 11 of these HAPEXTM implants which were retrieved 2 to 20 months after implantation owing to re-occurrence of the original clinical problems. They found a fibrous tissue layer covering all the implants (Fig. 3.3). In no cases did they find a foreign body response which they comment is in contrast to implants manufactured of Proplast® and Plastipore® used in similar applications in the middle ear.
S
H
3.3 Scanning electron microscopic appearance of a HAPEXTM middle ear prosthesis consisting of a hydroxylapatite head (H) and a hydroxylapatite – polyethylene shaft (S). The shaft has been trimmed intra-operatively to fit to the appropriate size in the middle ear. The shaft is composed of sintered hydroxylapatite and polyethylene which gives a smooth but irregular surface (Meijer et al., 2002).
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Biomedical composites
3.2.3 Dental applications As with many other activities in the field of biomaterials, dental applications lead the surgical uses of biomaterials. The major dental use of biocomposites, as tooth-filling materials to replace dental amalgam, will not be considered in this chapter. However, implantable composites have been used as posts for tooth replacement to provide support for either natural or artificial teeth or on the tongue side of teeth to support, either temporarily or permanently, a tooth where the anchorage has been weakened (Chan et al., 2006). While the Brånemark titanium tooth root replacement and similar designs are now used almost exclusively, in the past composites have been used to treat such problems. Hodosh et al. (1976) manufactured their own implants by mixing polymethylmethacrylate (PMMA) with vitreous carbon (VC). After extraction of one or more teeth, they produced moulds of the gaps in their patients’ alveolar ridges and filled these moulds with a mixture of 95% PMMA and 5% VC, presumably these are weight percentages although this is not stated, with MMA monomer and then heat polymerised. The implants were sand blasted to remove the surface layer and expose the VC. After ultrasound cleaning and sterilization, these were implanted into their patients and used to provide supports for fixed partial dentures. The authors say that they obtained rapid healing with minimal discomfort, which seems surprising given their statement that the composite ‘is a mildly irritating, strong implant material’. Their study of 15 patients also found minimal bone loss with retention of the alveolar ridge height. They do comment that the use of a material with lower modulus than bone produces better load transfer from the implant to the supporting bone.
3.3
Axial skeleton applications
3.3.1 Internal applications The first composite device to reach axial skeletal clinical application was a carbon fibre in epoxy resin composite fracture fixation plate developed by Hastings and colleagues (Hastings, 1978; Bradley et al., 1980; Ali et al., 1990) where carbon fibres were used to reinforce epoxy resin. The devices were manufactured by a combination of heat and pressure applied to 21 layers of carbon fibres in epoxy pre-preg with fibre directions along and at 45 ° to the long axis of the plates. Once the plates were manufactured, screw holes were drilled through and countersunk and the size and shape of the implant was identical to a typical metal fracture fixation plate of that time. These plates had a bending stiffness approximately one quarter that of the equivalent metal plates, yet the both fatigue limit and angulation at failure were 60%
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higher. After in vivo implantation in mice both in bone and with muscle contact to assess the biocompatibility, these plates were used in 40 forearm fractures in 29 patients (Ali et al., 1990). The application was limited to the forearm as these plates could not be bent to provide both close contact between the fracture ends and between the implant and the bone in other areas of the body, unlike metal implants. Clinically, the patients used their arms earlier than those treated with equivalent metal implants and, in 70% of the cases, the healing process was secondary healing with callus formation (Fig. 3.4) rather than primary healing with minimal callus formation, that was considered desirable at that time for internally fixed fractures. Primary healing occurs where the motion at the fracture site is less than approximately 100 μm, whereas secondary healing occurs with limited motion at the fracture site and is a substantially faster process. All the implants were removed after healing was complete. Histological analysis showed minimal response in all but six fractures, although some carbon particles were found in the soft tissue surrounding the implant. In the other cases, the responses were minor, except in two patients and one of these implants was found to be infected. As a result of the histological studies, the authors suggest that there was no reason to remove the implants in future studies. The limit on the use of these devices was their inability to be bent to fit the patient, thus limiting their use to the forearm and this constraint resulted in them being discontinued. However, thermoplastic and thermoset polymer matrices reinforcing carbon and glass fibers have been used to produce composite hip joint prostheses as discussed elsewhere in this book. More recently the self-reinforced poly-l-lactide (SR-PLLA) materials developed by Törmälä and his group (Böstmann et al., 1987; Törmälä et al.,
3.4 Clinical response to carbon-fibre-reinforced epoxy resin fracture fixation plates in the forearm showing the ‘healing by close-knit callus which was seen in 70% of cases’ (Ali et al., 1990).
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1988; Suuronen et al., 1992; Waris et al., 1994; Ashammakhi et al., 2001) have been used in internal fixation of the ankle and in femoral neck fractures. In an early study of displaced medial and lateral malleolar fractures (Böstman et al., 1987), they compared the results of treatment with cylinders of self-reinforced polylactide–glycolide (SR-PLGA) fibres with conventional metal screws. They found no differences in the anatomical and functional results with similar levels of complications in each group. In a later study also in the ankle, Joukainen et al. (2007) compared screws manufactured from self-reinforced polylactide with 70% poly-l-lactide and 30% poly-ld-lactide (SR-PLA70) with those made from 100% L PLA (SRPLLA) for the treatment of ankle fractures in 62 patients. There was minimal difference in the mechanical properties of the two devices, but the co-polymer reduced the time that the implants kept their strength in vitro from 36 to 24 weeks. They found that the patients with the co-polymer implants (SR-PLA70) had 65 days sick leave compared with 60 days for the SR-PLLA treated patients, but no other significant differences. In both groups the screw track was still visible at one year (Fig. 3.5). In a femoral neck study, Jukkala-Partio et al. (2000) treated 40 patients with subcapital fractures using three 6.3-mm-diameter SR-PLLA screws per fracture and the results were compared with 38 patients treated with three metal screws each 7.0 mm in diameter. The groups were similar in age and clinical problems and there were similar numbers of redislocations in each group, but the ability to walk and long-term range of movement were greater in the patients treated with the degradable implants. In joint replacement, Field and Rushton (Field and Rushton, 2005; Field et al., 2006) developed the Cambridge Cup to replace the horseshoe
(a)
(b)
(c)
(d)
3.5 Fracture of the lateral malleolus pre-operatively (a,b) and at the 1-year follow-up (c,d) treated with two SR-PLA70 screws. Note the syndesmotic ossification between the tibia and the fibular. The patient did not report any problems at the 1-year follow-up (Joukainen et al., 2007).
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of articular cartilage in the acetabulum in patients undergoing what would normally be a hemiarthoplasty, that is replacement of the femoral head only, after femoral neck fractures, in elderly, low-activity, osteoporotic patients. In these patients, a large femoral head articulates with the articular cartilage and with time the cartilage can wear through, but many of these patients are not well enough in themselves to be suitable for a total hip replacement. The Cambridge cup consists of 3 mm of UHMWPE supported by a 1.5-mm-thick layer of carbon fibre reinforced polybutyleneterephthalate (PBT). Implant fixation was either a 60-μm thick plasma-sprayed HA coating or PMMA bone cement supplemented with some composite fixation pins (Fig. 3.6). The concept of this design is that the implant stiffness is similar to that of the articular cartilage and subchondral bone plate structure that it is replacing. Compared with other studies using composite materials this was a much older age group of patients from 70 to 100 years with a mean age of 81.8 years, but this age range did mean that post-mortem studies have been performed, although not yet published. They found that, despite the relatively thin layer of UHMWPE, there was no significant wear, which they attributed to the flexibility of the device. The PMMA implanted prostheses migrated more than those relying on the HA coating for fixation, which would normally
3.6 Cambridge Cup showing the outer polybutyleneterephthalate (PBT) carbon-fibre composite shell with the inner UHMWPE bearing and holes for the insertion of fixation pegs (Field and Rushton, 2005).
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be considered an indicator of increased long-term risk of implant loosening (Ryd, 1985). Dual energy x-ray absorptiometry (DEXA) of the bone surrounding and supporting the implant showed that the bone density dropped in the first 6 to 12 months but then returned to the immediate post-operative levels. The post-mortem studies showed good bone contact with the HA coated implants.
3.3.2 External applications The principal applications of composites external to the body are in fracture fixation. Two major applications are in use: fibre-reinforced composites as external casts and carbon-fibre-reinforced epoxy rings used in the Ilizarov External Fixator. Plaster of Paris has many disadvantages being fairly radiopaque, mechanically brittle, of fairly high density and not being water resistant. Various orthopaedic companies have developed a wide woven fibre mesh impregnated with a settable polymer as a plaster of Paris replacement. The requirements are low radiopacity to allow the process of fracture healing to be followed radiographically without the need to remove the cast, higher failure properties, water resistance and lower density. The polymer setting initially was by heat, but the more recent formulations are light settable. After a few initial problems with the heat setting formulations in the developed world, plaster of Paris casts have been almost completely replaced with composite materials. The Ilizarov external fixator was developed by Gavriil Ilizarov in Siberia in the 1950s, but it was only in the 1980s that this external fixator became known in the West (Ilizarov, 1988). The fixator is used for fracture fixation, in limb lengthening and limb straightening. It consists of a series of thin wires, 1.5 or 1.8 mm in diameter, that pass through the bone above and below the fracture or osteotomy, these wires are then tensioned onto a series of rings that are connected to form a scaffolding system around the injured limb. For limb correction, the rings are gradually moved using the connectors that cross the osteotomy site to move the bones to their required anatomy. Initially, the entire fixator was fabricated of stainless steel, which made a heavy system with large amounts of radiopaque metal, potentially making imaging the fracture difficult. However, random chopped carbon fibre in epoxy resin and knitted Kevlar-29 in epoxy rings for the Ilizarov have been developed (Baidya et al., 2001 a and b). These rings were extensively tested both as individual structures and as parts of external fixators. They found that the composite rings were less stiff than the steel rings, but had sufficient stiffness and strength to be used clinically, while the reduced weight and radiopacity has benefits to both patients and surgeons.
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3.4
55
Advantages in the use of composites for hard tissue applications
Composites can be produced that have similar stiffness to the natural tissues they are required to replace. This modulus matching can lead to better biological responses in the form of reducing stress shielding. Furthermore, many of these newer composites contain calcium phosphates, which are released from the implant providing a supply of calcium and phosphate ions, or other biologically active moieties, that can be incorporated into the bone and accelerating the healing process. Composites implants can be tailored in the operating theatre to fit the patient, unlike most metal implants, which can only be bent but not cut, or all ceramic implants, which are difficult to cut and cannot be bent. However, the methods of shaping composites intraoperatively are different from those used with metal implants requiring heat or cutting rather than plastic deformation of the metal.
3.5
Disadvantages in the use of composites for hard tissue applications
As yet most of the composites which have been implanted have been limited to lower load-bearing applications owing to their poor strength and fracture toughness. Some of the less successful applications have been where the metal implant design has been exactly reproduced in a composite rather than starting with the implant requirements and using the properties of the composite to drive the design; i.e. the design is based on the properties of the composite to be used.
3.6
Future trends
Composites have been used in patients since the 1970s. However, the early composites were carbon containing and produced only acceptable biological responses. The newer variations of composites are bioactive producing beneficial or active biological responses. One of the major developments has been the selection of either non-degradable or degradable polymers as the basis of composites depending on the clinical application. In joint replacement and similar applications, non-degradable implants are required, whereas for fracture fixation, and even more so for tissue engineering scaffolds, degradable scaffolds will allow the patient to be treated and then forget that they ever had an artificial material implanted. It seems likely that the implant material and design will be based on the best combination for the specific clinical problem rather than treating the clinical problem with an available implant.
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It is interesting to note that there have only been a few composites that have progressed through to published clinical trials; however, considering how long it takes to progress from development of a new material to clinical use, we should expect an increasing number of composites, such as those discussed elsewhere in this book, to enter clinical use in the next five to ten years. With time, our understanding of the advantages and disadvantages of the clinical use of composites is increasing and this should accelerate the further introduction of more advanced composites whether bioactive materials or ‘smart’ materials.
3.7
References
ali m s, french t a, hastings g w, rae t, rushton n, ross e r s, wynn-jones c h (1990), ‘Carbon fibre composite bone plates development, evaluation and early clinical experience’, J Bone Joint Surg(Br), 72-B(4), 586–591. ambrosio l, de santis r, nicolais l (1998) ‘Composite hydrogels for implants’, Proc IMechE Part H: J Eng Med, 212(H2), 93–99. doi: 10.1243/0954411981533863. amis a a, campbell j r, kempson s a, miller j h (1984) ‘Comparison of the structure of neotendons induced by implantation of carbon or polyester fibres’, J Bone Joint Surg [Br], 66-B(1), 131–139. ashammakhi n, peltoniemi h, waris e, suuronen r, serlo w, kellomäki m, törmälä p, waris t (2001) ‘Developments in craniomaxillofacial surgery: use of selfreinforced bioabsorbable osteofixation devices’, Plastic Reconstructive Surg, 108(1), 167–180. baidya k p, ramakrishna s, rahman m, ritchie a (2001a) ‘Advanced textile composite ring for Ilizarov external fixator system’, Proc IMechE Part H: J Eng Med, 215(H1), 11–23. doi: 10.1243/0954411011533490. baidya k p, ramakrishna s, rahman m, ritchie a (2001b) ‘Quantitative radiographic analysis of fiber reinforced polymer composites’, J Biomater Appl, 15(3), 279–289. bergmann g, graichen f, röhlmann a (1993) ‘Hip-joint loading during walking and running, measured in 2 patients’, J Biomech, 26(8), 969–990. bonfield w, bowman j, grynpas m d (1980) ‘Composite materials for use in orthopaedics’, UK patent no. GB2085461B. bonfield w, grynpas m, tully a e, bowman j, abram j (1981) ‘Hydroxyapatite reinforced polyethylene – a mechanically compatible implant material for bonereplacement’, Biomaterials, 2(3), 185–186. bos r r m, boering g, rozema f r, leenslag j w (1987) ‘Resorbable poly-(l-lactide) plates and screws for the fixation of zygomatic fractures’, J Oral Maxillofac Surg 45(9), 751–753. böstman o, hirvensalo e, mäkinen j, rokkanen p (1990) ‘Foreign-body reactions to fracture fixation implants of biodegradable synthetic-polymers’, J Bone Jt Surg(Br), 72-B(4), 592–596. bradley j s, hastings g w, johnson-nurse c (1980) ‘Carbon fibre reinforced epoxy as a high strength, low modulus material for internal fixation plates’, Biomaterials, 1(1), 38–40.
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burr d b, milgrom c, fyhrie d, forwood m, nyska m, finestone a, hoshaw s, saiag e, simkin a (1996) ‘In vivo measurement of human tibial strains during vigorous activity’, Bone, 18(5), 405–410. chan d c n, giannini m, de goes m f (2006) ‘Provisional anterior tooth replacement using nonimpregnated fiber and fiber-reinforced composite resin materials: A clinical report’, J Prosthetic Dentistry, 95(5), 344–348, doi: 10.1016/j.prosdent. 2006.01.017. currey j d (1998) ‘Mechanical properties of vertebrate hard tissues’, Proc IMechE Part H: J Eng Med, 212-H(6), 399–411, doi 10.1243/0954411981534178. del campo a f, pohjonen t, törmälä p, waris t (1996) ‘Fixation of horizontal maxillary osteotomies with biodegradable self-reinforced absorbable polylactide plates: preliminary results’, Euro J Plast Surg, 19(1), 7–9. dornhoffer j l (1998) ‘Hearing results with the Dornhoffer ossicular replacement prostheses’, Laryngoscope, 108(4), 531–536. downes r n, vardy s, tanner k e, bonfield w (1991) ‘Hydroxyapatite–polyethylene composite in ophthalmic surgery’, In Bioceramics 4 – Proceedings of the Fourth International Symposium on Ceramics in Medicine, Edited by W. Bonfield, G.W. Hastings and K.E. Tanner, Published by Butterworth–Heinnemann Ltd, pp. 239–246. field r e, cronin m d, singh p j, burtenshaw c, rushton n (2006) ‘Bone remodeling around the Cambridge cup – a DEXA study of 50 hips over 2 years’, Acta Orthop, 77(5), 726–732. doi: 10.1080/17453670610012908. field r e, rushton n (2005) ‘Five-year clinical, radiological and postmortem results of the Cambridge Cup inpatients with displaced fractures of the neck of the femur’, J Bone Jt Surg (Br), 87-B(10), 1344–1351, doi: 10.1302/0301-620X.87B10. 16559. goldenberg r a (1994) ‘Ossiculoplasty with composite prostheses – PORP and TORP’, Otolaryngologic Clin North Am, 27(4), 727–745. goldenberg r a, daver m (2000) ‘Long-term results with hydroxylapatite middle ear implants’, Otolaryngology – Head and Neck Surgery, 122(5), 635–642. hastings g w (1978) ‘Carbon fibre composites for orthopaedic implants’, Composites, 9(3), 193–197. hodosh m, shklar g, povar m (1976) ‘Syntactic, porous polymethacrylate vitreous carbon tooth replica implants as abutments for fixed partial dentures’, J Prosthetic Dentistry, 36(6), 676–684. ilizarov g a (1988) ‘The principles of the Ilizarov method’, Bull Hosp Jt Dis, 48(1), 1–11. joukainen a, partio e k, waris p, joukainen j, kröger h, törmälä p, rokkanen p (2007) ‘Bioabsorbable screw fixation for the treatment of ankle fractures’, J Orthop Sci, 12(1), 28–34, doi: 10.1007/s00776-006-1077-y. joyce t j, unsworth a (2000) ‘The design of a finger wear simulator and preliminary results’, Proc IMechE Part H: J Eng Med, 214-H(5), 519–526, doi: 10.1243/ 0954411001535552. jukkala-partio k, partio e k, helevirta p, pohjonen t, törmälä p, rokkanen p (2000) ‘Treatment of subcapital femoral neck fractures with bioabsorbable or metallic screw fixation – a preliminary report’, Ann Chir Gynaecol, 89(1), 45–52. meijer a g w, segenhout h m, albers f w j, van de want h j l (2002) ‘Histopathology of biocompatible hydroxylapatite–polyethylene composite in ossiculoplasty’, ORL: J Oto-Rhino-Laryngology Rel Specialities, 64(3), 173–179.
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mundermann a, dyrby c o, d’lima d d, colwell c w, andriacchi t p (2008) ‘In vivo knee loading characteristics during activities of daily living as measured by an instrumented total knee replacement’, J Orthop Res, 26(9), 1167–1172, doi: 10.1002/jor.20655. piattelli a, france m, ferronato g, santello m t, martinetti r, scarano a (1997) ‘Resorption of composite polymer–hydroxyapatite membranes: a time-course study in rabbit’, Biomaterials, 18(8), 629–633. ryd l (1985) ‘Micromotion in knee arthroplasty – a Roentgen stereophotogrammetric analysis of tibial component fixation’, Acta Orthop Scand, 57(Suppl 220), 3–80. suuronen r, pohjonen t, taurio r, törmälä p, wessman l, rönkkö k, vainionpää s (1992) ‘Strength retention of self-reinforced poly-l-lactide screws and plates. An experimental study’, J Mater Sci: Mater Med, 3(6), 426–431. tanner k e, downes r n, bonfield w (1994) ‘Clinical application of hydroxyapatite reinforced polyethylene’, Brit Ceram Trans, 93(3), 104–107. törmälä p, rokkanen p, laiho j, tamminmäki m, vainionpää s (1988) ‘Material for osteosynthesis devices’, US patent 4 743 257. wallbridge n, dowson d (1982) ‘The walking activity of patients with artificial hip joints’, Eng Med, 11(3), 95–96, doi: 10.1243/EMED_JOUR_1982_011_023_02. wang m, porter d, bonfield w (1994) ‘Processing, characterization and evaluation of hydroxyapatite-reinforced polyethylene composites’. Br Ceram Trans, 93(3), 91–95. waris t, pohjonen t, törmälä p (1994) Self-reinforced absorbable polylactide (SRPLLA) plates in craniofacial surgery – a preliminary report on 14 patients’, Euro J Plast Surg, 17(5), 236–238. zanetti d, nassif n (2003) ‘Transmastoid repair of minor skull base defects with flexible hydroxyapatite sheets’, Skull Base – An Interdisciplinary Approach, 13(1), 1–11. zanetti d, nassif n, antonelli a r (2001) ‘Surgical repair of bone defects of the ear canal wall with flexible hydroxylapatite sheets: a pilot study’, Otol Neurotol, 22(6), 745–753.
4 Soft tissue applications of biocomposites M. S A N T I N, University of Brighton, UK
Abstract: In this chapter, an overview is presented of the role of composite biomaterials as substitutes for soft tissue repair. In particular, those technological solutions where either biopolymers or synthetic biocompatible, biomimetic polymers have been used are highlighted. The potential of these biomaterials to fulfil clinical needs is underpinned by an overview of the composition and properties of the main extracellular matrix components. This overview allows the reader to critically assess the ability of the currently available composite biomaterials to perform in a clinical scenario. Examples of clinical applications are given in which composite biomaterials are fundamental to soft tissue aided repair. The clinical problems that arise when their use has been neglected, described. Key words: biocomposites, soft tissue repair, biomimetic polymers.
4.1
Introduction
The occurrence of traumatic events (accidental or surgical) in soft tissues leads to the loss of their architectural, biochemical, cellular and homeostatic properties (Hodde and Johnson, 2007). Such damage mainly occurs in the composite polymeric network in which cells reside, the extracellular matrix (ECM). The ECM is mainly composed of structural protein and polysaccharides (the glycosaminoglycans) which preserve cell viability and the delicate homeostatic balance of the tissue (see section 4.2). Intact ECM can quickly, easily, and effectively provide key extracellular components of the dermis necessary to direct the healing response in the case of tissue damage and to allow the formation of new tissue. The restoration of these properties in spontaneous healing is gradual and it takes place through a series of interconnected phases where bleeding is arrested by the formation of a clot which is gradually replaced by new mature tissue. This is progressively formed through a series of microenvironmental and biochemical changes which favour the proliferative and synthetic activity of the cells (Werner et al., 2007). Indeed, tissue repair is characterised by different phases where the clot formation is followed by an inflammatory response leading to the formation of the highly vascularised granulation tissue generally progressing into scar. With time, the scar, 59
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a tissue characterised by highly packed collagen fibrils, undergoes remodelling into healthy tissue. Together with the inflammatory response, the granulation tissue plays a fundamental role in tissue repair (Werner et al., 2007). In these two phases, important biochemical stimuli are secreted which recruit mesenchymal cells to the site of injury and induce them to proliferate and, later, to synthesise and deposit new ECM. For skin repair, epithelial cells also migrate and proliferate to cover the temporary matrix of the underlying granulation tissue. In the repairing dermis, keratinocytes and fibroblasts are mutually stimulated to synthesise growth factors and to proliferate. During this phase, fibroblasts acquire a proliferative phenotype, the myofibroblast which leads to extracellular matrix contraction, a process known to facilitate wound closure. These processes depend on a finely tuned balance between the secretion of proinflammatory cytokines and the occurrence of a transforming growth factor (TGF) β-rich microenvironment. However, tissue healing takes place spontaneously only when two fundamental conditions are fulfilled: 1. The area of damage still allows bridging of new structural components and cells sense biochemical signalling. 2. The physiological homeostasis is preserved as cell activity is not compromised by any pathological unbalance. In those cases where the area of damage is relatively large or a specific pathological conditions occurs, defect colonisation by the cells is impaired and their cross-talk through secreted biochemical signals (i.e. growth factors) fails to establish active concentration gradients (Werner et al., 2007). In certain pathological conditions such as diabetes and autoimmune diseases, tissue repair can be compromised by alterations which have not been yet fully understood (Li et al., 2008). For example, recent findings in a diabetic mouse model have highlighted the relatively lower expression of receptors for the platelet derived growth factor (PDGF-BB), an important stimulator of myofibroblast proliferation and promoter of tissue repair (Li et al., 2008). Clinicians frequently face such instances and they have to intervene to provide aid to the healing process. To this purpose, surgeons and practicioners rely upon implants which are able to repair temporarily or permanently the damaged tissue. These implants are manufactured from biomaterials which, in most cases, have been designed taking inspiration from the structural and biomechanical properties of the original healthy tissue. While trying to fulfil these characteristics, scientists and manufacturers soon realised that a close mimicking of the structure of the damaged
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soft tissue could be achieved by the development of composite biomaterials (Hodde and Johnson, 2007). This chapter provides a critical overview of the state-of-the-art in composite biomaterials for soft tissue regeneration. The available materials will be discussed taking into account their ability to simulate the tissue they are meant to replace and, for biodegradable biomaterials, their ability to simulate the composition of the temporary extracellular matrix in the different phases of tissue repair. Ultimately, the available composites will be presented on the basis of typical clinical applications and their relative performances.
4.2
Composite nature of soft tissue extracellular matrices
The biochemical composition and architecture of the ECM has developed throughout evolution to fulfil the functional needs of the connective tissue (Wiig et al., 2008). The environment surrounding the tissue cells is a dynamic milieau consisting of interstitial fluid and structural molecules of the extracellular matrix, mainly glycosaminoglycans (GAG) such as hyaluronan (HA) and proteoglycans (PGN) and proteins such as collagen and fibronectin (Fbn). The diffusion of bioactive macromolecules (e.g. cytokines and growth factors) throughout this environment depends on the physicochemical properties of the ECM structural components. More specifically, free molecules can be localised only in those spaces unoccupied by the structural components. This phenomenon is called interstitial exclusion and it plays a key role in the regulation of the plasma volume. Interstitial exclusion is generated by both structural proteins (mainly collagen) and charged GAG (Wiig et al., 2008). Collagen occupies a geometrical space and, therefore, provides a sterical contribution to the interstitial effect process. GAG contributes to molecular exclusion by electrostatic interactions through their negatively charges. This charge may need to be taken into account in the development of new ECM mimicking biomaterials to improve their ability to entrap bioactive molecules (e.g. drugs and growth factors) and to control their release upon implantation. Wiig et al. (2008) emphasise that ‘This structure-to-function relationship is always achieved by nature through the development of an ECM through composite macromolecular assembling’. Furthermore, as already mentioned in section 4.1, the ECM also plays a key and dynamic role in regulating cell functions (Daley et al., 2008). Indeed, ECM regulates cell morphology, viability, migration, proliferation and differentiation. ECM constantly undergoes remodelling mainly through a turnover of its components by cellular synthesis and enzymatic
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degradation. This process, although particularly active during the individual’s developmental age and tissue healing, takes place throughout the organism’s lifespan. Following cellular synthesis and secretion, ECM assembly is controlled by the interactions occurring between its components as well as by the tension exerted by the cells through a specific class of receptors, the integrins. Instead, degradation is controlled by proteolytic enzymatic cascades. Key to this process is stem cell self-renewal and differentiation that is itself influenced by the 3D environment within the stem cell niche (Daley et al., 2008). In summary, ECM can be described as the cell functional environment the structure of which self-organises through both dynamic macromolecular assemblying and component turnover (Ori et al., 2008). In this space, biochemical signals diffuse to promote tuned communication between cells of the same or different phenotypes. This diffusion is itself affected by the ability of the bioactive molecules to interact/penetrate throughout the ECM mesh and by the regulated assembly / disassembly of its molecular complexes. When the plethora of biochemical signalling is integrated, it becomes a sophisticated control system guiding vital biological processes such as embryonic development and organism’s homeostasis. For example, heparan sulfate, proteoglycans which are ubiquitous components of the ECM play an important role in these processes as they are able to integrate self-assembling macromolecular structures over substantial length scales. However, when altered by biochemical and cellular dysfunctions, remodelling leads to pathological conditions thus contributing to diseases. For example, matrix metalloproteinases (MMP), a family of zinc-dependent endopeptidases that degrade the extracellular matrix, play a key role in physiological tissue remodelling processes such as those occurring during pregnancy, development and tissue repair, but their activity alteration also leads to pathological conditions such as cancer and arthritis (Le et al., 2007). Interestingly, the MMP family appears to act as regulators of both activation and inhibition pathways of the inflammatory response. The ability of controlling their activity in these pathways through pharmaceutical and biomaterial strategies may lead in the future to new treatments of inflammation-related pathologies such as arthritis, cancer and chronic wounds. Fibrin, HA, Col, elastin, Fbn, and vitronectin are just a few biomacromolecules that provide a composite nature to the soft tissue ECM mesh. In this section, the structure-to-function relationship of these main biomacromolecules will be presented and discussed in the light of ECM development, healing, maturation and remodelling phases. This overview will be the platform for understanding present and future strategies underpinning the design of biomimetic composite biomaterials for soft tissue replacement and repair.
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4.2.1 Fibrin Fibrin is the main component of the blood clot. This natural hydrogel is formed through the enzymatic activation, polymerisation and crosslinking of its precursor, the fibrinogen. Fibrinogen activation and polymerisation is activated by thrombin, while its crosslinking is catalysed by a transglutaminase enzyme, the Factor XIII (Lim et al., 2008). By entrapping platelets, the fibrin mesh is able to stop bleeding. To exert its haemostatic function against the blood shear forces, fibrin needs to combine stiffness and elasticity. A recent study by atomic force microscopy and steered molecular dynamics has probed the mechanical properties of single fibrinogen molecules and fibrin protofibrils and demonstrated unique mechanical properties (Lim et al., 2008). The specific tertiary structure (folding) of these proteins appears to lead to a unique intermediate force plateau in the force–extension curve and explain its elasticity. This is generated by a stepwise unfolding of both fibrinogen coiled alpha helices and central domain. These mechanical properties change according to the variation in environmental parameters such as pH and calcium ion concentration which alter the mechanical resilience of the protein by acting on its tertiary conformation (Lim et al., 2008). Measurements of clot mechanical properties have revealed that the clots achieved maximum stiffness and minimum plasticity when specific structural parameters reached their maxima. In particular, the analysis of hydrated fibrin mesh by deconvolution microscopy has established that, after gelation, an initial fibrin network progressively evolves over time by both addition of new fibers and elongation and branching of others (Chernysh et al., 2008). Biochemical signalling molecules bind the fibres of the developing fibrin mesh (and later of the ECM) thus generating chemotactic gradients for the recruitment of stem cells as well as for the stimulation of tissue cell migration and differentiation (Nurden et al., 2008). For these reasons, platelets in a fibrin clot encourage rapid healing in several clinical scenarios such as dental implant surgery, orthopaedic surgery, skin ulcer treatment, eye and cardiac surgery. Because of its important role in tissue repair, fibrin has been the target of many material scientists and commercial products known as fibrin glues have been made available to surgeons (Zhao et al., 2008). In general, ready-to-reconstitute fibrin is provided as a kit comprising the fibrinogen freeze-dried powder together with the thrombin and Factor XIII components necessary for its activation into a polymerised and crosslinked mesh. Therefore, by varying the concentrations of the substrate as well as of the enzymes, hydrogels with different characteristics can be obtained (Zhao et al., 2008).
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4.2.2 Glycosaminoglycans and proteoglycans Recent findings have revealed the role of GAG and PGN in many biological processes such as host defense, wound healing, inflammation and fibrosis. Because of their high net negative charge, GAG and PGN (e.g. heparan sulfate) are also able to sequester growth factors and activate chemokines and cytokines as well as to confer selective permeability to the basement membranes (Yung and Tak Mao, 2007). GAG such as heparan sulfate and hyaluronan are linear carbohydrates which participate in both tissue development, remodelling and healing by regulating cell-ECM interactions and activation of chemokines, enzymes and growth factors (Taylor and Gallo, 2006). In development, sulfated PGN controls the differentiation of pancreatic endocrine cells such as the beta cells (Zertal-Zidani et al., 2007). Further insights into the mechanism of action of sulfated PGN on beta cell development could be, for example, useful in the generation of beta cells from embryonic stem cells for the treatment of diabetic subjects. Heparan sulfate also regulates enzyme/inhibitor functions, interacts with cytokines/chemokines and promotes binding and recruitment of leukocytes (Taylor and Gallo, 2006). Chondroitin sulfate/dermatan sulfate regulates growth factor activity and induces ICAM-1 expression on endothelial cells and also recruits leukocytes (Taylor and Gallo, 2006). TGF-β isoforms 1 and 2 and glial-cell-line-derived neurotrophic factor (GDNF) have been shown to bind to heparin and heparan sulfate (Rider, 2006). However, the interactions of cytokines and growth factors with these macromolecules may be very finely modulated by the exposed functional groups. Indeed, current work seems to demonstrate that the TGF-β superfamily do not share a single heparin/heparan sulfate-binding site (Rider, 2006). Furthermore, because of the close link between GAG and PGN physicochemical properties and their functions, changes of their chemical composition and structure can lead to pathological conditions (Hitchcock et al., 2008). Isomeric chondroitin sulfate glycoforms differing in position and degree of sulfation and uronic acid epimerisation play specific and distinct functional roles during development and disease onset. GAG storage of mucopolysaccharidoses leads to inflammation and apoptosis within cartilage and synovial tissue (Simonaro et al., 2008). Pharmaceutical treatments are also known to induce changes in the macromolecular structure and physiological role of both GAG and PGN (Claassen et al., 2006). For example, it has been reported that the deficiency
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of vitamin D, calcium and oestrogens in combination with biphosphonate treatment leads to an altered GAG content. As GAG and, at some extent PGN, contribute to water-storage within articular cartilages and, therefore, to its biomechanical properties, the partial loss of these biomacromolecules can severely impact on tissue functionality. Thus, the field of GAG biology provides new clues and explanations of the process of inflammation and suggests new therapeutic approaches to human disease (Na et al., 2006). Hyaluronan HA is a relatively high-molecular-weight GAG composed of repeating d-glucuronic acid, β 1–3 linked to N-acetyl-d-glucosamine β 1–4, found in body fluids and tissues, in both intra- and extracellular compartments (Bastow et al., 2008). Despite its relatively simple structure, HA acts on tissue development, homeostasis and repair in a multi-faceted manner. Through its specific interactions with CD44 cell receptors, HA promotes cell recruitment as well as to joint cavity and bone formation. In adult cartilage, HA forms macromolecular complexes with aggrecan to promote the formation of high-concentration domains and, therefore, physiologically performing compressive resilience. HA turnover is regulated by its enzymatically driven synthesis and biodegradation. Hyaluronan synthase isoforms 1, 2 and 3 and hyaluronidases finely regulate the formation and removal of this GAG during development as well as during tissue remodelling and repair. Hyaluronan also play a key role in the inflammatory response and angiogenesis (Taylor and Gallo, 2006; Slevin et al., 2007). It is known that low molecular weight HA (<2 × 106 Daltons) participates in leukocyte recruitment via interaction with CD44, while activating various inflammatory cells, such as macrophages, through CD44-dependent signalling (Taylor and Gallo, 2006). Conversely, high molecular weight HA (>2 × 106 Daltons) favours the activation of endothelial cells to form new blood vessels through its interaction with both CD44 and Receptor for HA-Mediated Motility (RHAMM, CD 168) (Taylor and Gallo, 2006; Slevin et al., 2007). As far as its use in regenerative medicine is concerned, the most important property of HA is its ability to play a role as a signal for tissue injury (Yamasaki et al., 2008) and to play an important role in tissue repair (Jiang et al., 2007). As discussed above, the simple repeating structure of HA appears to be involved in a number of important aspects of tissue injury and repair which depend on the molecular weight and location of the polymer as well as the interacting cells.
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4.2.3 Collagen Collagen is the most abundant structural protein and is the building block of most connective tissue ECM. This protein is characterised by a repetitive amino acid sequence in which hydroxyproline ensures the folding of fibrils into a triple helix. Collagen fibrils are then assembled into supramolecular structures with growing level of complexity (i.e. fibrils, fibres, bundles, fascicles, mesh) to provide the scaffold of both soft and hard tissues (Birk et al., 1997; Grant, 2007). Although the majority of collagen is of type I, several types of collagen are present in the human tissues. In this chapter, only some typical examples are reported which have been exploited or may be exploited in future technological applications. Particular attention will be given to collagen types I and II to highlight their role in soft tissues such as skin, cartilage and peripheral nerves, and to collagen type IV as the precursor of basement membrane in tissues such as skin and blood vessels. Types V and VI will also mentioned for their role in fibre assemblying, while type X will be considered to highlight its role in tissue repair.
Type I collagen Type I collagen has been studied particularly in relationship to its structural and functional role in tendons (Benjamin et al., 2008). As tendons are discussed in another chapter of the present book, type I collagen will be here discussed in the light of its presence and biological properties in other tissues. A typical example is angiogenesis that is the process leading to the development of new blood vessels from the pre-existing vasculature. This is a fundamental aspect in embryogenesis and tissue regeneration, but is also part of the development of pathologies such as tumour growth and metastasis, and neo-natal haemangioma. Likewise, a physiological induction of angiogenesis is required in tissue engineering constructs. As a consequence, research has been focusing on the biochemical and biostructural mechanisms regulating this process. Among them, type I collagen has been found to be a potent stimulator of angiogenesis both in vitro and in vivo (Twardowski, 2007). Studies have suggested that part of the type I collagen angiogenic potential resides in its ability to facilitate ligation and, possibly, clustering of endothelial cell through the plasmalemma alpha 1 beta 1/alpha 2 beta 1 integrin receptors. This interaction seems to be promoted by the amino acid sequence GFPGER of the collagen fibril. Furthermore, it has been suggested that type I collagen may contribute to cell functions by molecular mechanotransduction. Indeed, cells respond
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to mechanical stimuli at molecular scale that occur in the ECM (Peyton et al., 2007). Recent evidence show that these stimuli regulate the cytoskeletal tension thus triggering molecular switches that controls cell function. In particular, two-dimensional in vitro model systems on collagen-based substrates have shown that the substrate mechanical properties are able to affect cell morphology, migration, proliferation, and differentiation. The molecular mechanical properties of collagen also determine the macroscale biophysical characteristics of the tissue. A typical example is the vascular tissue, the mechanical properties of which depend, at large extent, on the properties of fibrillar collagens. Alongside the molecular structure, these properties are also determined by the degree of crosslinking of the type I collagen chains which are driven by the activity of enzymes such as the lysyl oxidase (Robin, 2007). It is important to observe that this enzymatically driven crosslinking optimises the collagen function and is a mechanism distinct from the non-enzymatic crosslinking occurring with aging that leads to structural changes deleterious to function. From a biomaterial viewpoint, therefore, the chemical or physical crosslinking of collagen-based biomaterials is more likely to lead to dysfunctional collagen. Type I collagen dysfunctions are also caused by other types of physicochemical changes and they are at the basis of serious pathologies (Kuo et al., 2007; Xiao et al., 2007). Electron spin resonance (ESR) indicates that type I collagen is able to scavenge hydroxyl radicals in dose- and time-dependent manner, while its denaturation product gelatine does not have this property (Xiao et al., 2007). When compared with other natural scavangers such as the glutathione, type I collagen was shown to scavenge hydroxyl radicals by a different mechanism whereby the structure of the main chain of collagen changes first, followed by carbonyl residue groups; the latter reaction leads to hydroxyl-free radicals scavenging. Therefore, when type I collagen is used as a biomaterial any physicochemical modification made during its processing may compromise its ability to scavenge free radicals. This information should be taken into account in the design of biomaterials able to control the oxidative stress induced by the inflammatory cells during the early phase of implantation. Type II collagen Together with PGN and types VI and IX collagens, type II collagen is one of the main components of the so-called pericellular matrix (PCM), the thin tissue surrounding chondrocytes in articular cartilage (Guilak et al., 2006).
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The chondrocyte/PCM complex is better known as chondron. The function of this region is not fully understood, but it has been hypothesised to play a role in controlling cell differentiation and in regulating biomechanical stimuli. Fabricated microspheres of collagen have been used to entrap hMSC by microencapsulation (Hui et al., 2008). The collagen microspheres were shown to induce hMSCs pre-differentiation into a chondrogenic phenotype in vitro. Upon implantation, these tissue-engineered cells became surrounded by cartilage-specific extracellular matrix components, among them type II collagen and aggrecan. In a similar study, to facilitate cell distribution in a tissue-engineering construct for cartilage regeneration, hMSC were encapsulated in type II collagen fibers to show that they were able to reorganise the matrix in which they were entrapped and to differentiate into spherical chondrocytes able to synthesise typical cartilage proteins (Chang et al., 2008). Type IV collagen Type IV collagen was discovered by Kefalides in 1966 (Khoshnoodi, 2008). Molecular genetics studies have shown that there are six different genes encoding this type of polypeptide chains in this type of collagen. These genes are expressed in different phases of embryonic development, providing different tissues with specific collagen IV networks. Each one of this network seems to have distinct biochemical properties. In general, type IV collagen is exclusively found in the basement membranes where it forms supramolecular networks able to control cell adhesion, migration, and differentiation (Khoshnoodi et al., 2008). Basement membranes are particularly important in tissue functionality as they are responsible for the apposition of epithelium and endothelium to the underlying skin dermal layer and blood vessel wall media, respectively (LeBleu et al., 2007). In them, type IV collagen forms a composite material with several large glycoproteins providing structural support to the tissue and functional stimuli to modulate cell activity. Matrigel is commonly used as surrogate for basement membranes in many experiments, but this is a tumour-derived material and does not contain all of the natural components of the basement membranes. For example, the interaction between the epithelial layer and the underlying mesenchyme promotes human skin morphogenesis and homeostasis of human skin. In vitro studies have shown that keratinocytes proliferate on substrates exposing type IV Collagen to form tissue with a physiological architecture (Segal et al., 2008). Conversely, cells growing on proteins not found in the basement membrane (i.e. fibronectin and type I collagen) led to tissue layers with an aberrant architecture.
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Type V and Type VI collagen Type V and VI collagen are structural proteins able to form macromolecular complexes with each other as well as with type I collagen and mucopolysaccharides (Kobayasi and Kalsmark, 2006). The propensity of these collagens to form macromolecular complexes suggests their function in ensuring the cohesion of fibrillar tissue components such as those of the dermis. In the dermis, the type VI collagen has been found entangled with type V collagen and in association with mucopolysaccharide domains. In the cartilage, the distinct mechanical properties of the PCM are defined primarily by the presence of type VI collagen (Guilak et al., 2006). The benefit deriving from a deeper understanding of the structure and function of the PCM and of the type VI collagen is, therefore, obvious in the light of both cartilage tissue engineering and cell therapy whereby the biomaterial is able to regulate finely the chondrocyte phenotype and activity.
4.2.4 Fibronectin Fibronectin is a protein present in connective tissues and in blood (known as cold fibronectin). This protein is well known to favour cell adhesion and it is particularly abundant in the basement membrane (Ruoslahti et al., 1980). The cell attachment promoting activity of fibronectin depends on the presence of specific amino acid sequences on its structure able to interact with cell integrin receptors and on its ability to form complex with other macromolecules of the ECM (e.g collagen, fibrinogen, fibrin and GAG). GAG interact with collagen/fibronectin complexes to form stable macromolecular aggregates. An amino acid sequence (Ala-Leu-Asn-Gly-Arg) of fibronectin type III domain 8-11 [FN-III (8-11)] has been found to be able to promote adhesion of human mesenchymal stem cells (Okochi et al., 2008). This sequence induces cell adhesion at levels similar to those of the Arg-Gly-Asp-Ser (RGDS) sequence. Similarly, the cross-talk between Tenascin-R (TN-R) and microglia on neural stem/progenitor cell proliferation and differentiation seems to be facilitated by fibronectin (Liao et al., 2008). Studies have shown that ECM fibronectin promotes the deposition of several ECM molecules, including collagen types I and III (Sottile et al., 2007). Furthermore, fibronectin-independent type I collagen deposition regulates the cell migratory response to fibronectin. The ability of fibronectin to bind other biomolecules is not limited to other structural elements of the ECM. Indeed, fibronectin is able to bind growth factors that are important in regulating the activity of resident cells (Clark et al., 2008).
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4.2.5 Elastin Elastin is a protein that plays a key role in providing resilience to the human skin. Elastin fibres of different lengths and distribution accumulate in the dermal compartment (Mewes et al., 2007). In addition, elastin also plays an important role in tissue remodelling and regeneration. It has been found that elastin is degraded by elastases to form fragments able to act as cytokines (Antonicelli et al., 2007). These so-called elastokines are likely to favour tissue repair by recruiting inflammatory cells, fibroblasts and smooth muscle cells and by stimulating their proliferation. Furthermore, these fragments have been shown to be inducers of angiogenesis as potent as the vascular endothelial growth factor (VEGF). However, a protracted exposure to these fragments may lead to a chronic inflammatory state. A possible role of these fragments in tumour has also been postulated as they are able to trigger excessive cell metalloproteinase activity.
4.2.6 Laminin Similarly, the type IV collagen, laminin, is one of the main components of the basement membrane separating epithelia and endothelia from the underlying stromal tissue (Merker and Barrach, 1981). This molecule has important amino acid sequences in its structure that promote cell adhesion, migration, proliferation and differentiation. This protein is also able to favour neurite outgrowth and Schwann cell development (Chernousov et al., 2008). The amino acid sequence –IKVAV – that is present in laminin has been found to promote capillary network formation by endothelial cells (Nakamura et al., 2008).
4.3
Composite biomaterials for soft tissue repair
All the biomacromolecules discussed in section 4.2 and the polymers (synthetic and natural) presented in other chapters of this book represent the building blocks or the models for the engineering of composite biomaterials for soft tissue repair. In this section, an overview of composites for soft tissue repair that are made of either biopolymers or synthetic biomimetic biomaterials or the combination of the two classes is provided. Any attempt at producing an exhaustive review of the field is frustrated by the vast number of scattered publications produced in the last two decades. Therefore, in the following sections, examples will be given which refer to the most recent publications which have approached the topic trying to fulfil the most important criteria which are the ability of the biomaterials to modulate cell behaviour and to gradually degrade upon implan-
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tation. Where possible, a particular emphasis is given to the bottom-up approach that leads from molecular assembly to new macrostructures and biofunctionalities.
4.3.1 Biopolymer-based composites for soft tissue repair: engineering methods and biological interactions The ECM components described in section 4.2 have been used as building blocks for the preparation of composite biomaterials for soft tissue repair in medical implant and tissue engineering approaches and their assembly has been controlled to mimic the ECM of the original tissue. Typically, isolated type I collagen fibrils, elastin fibres and chondroitin sulfate have been used for the preparation of collagen-elastin glycosaminoglycan scaffolds (Daamen et al., 2003). Scaffolds were prepared by varying the collagen-to-elastin ratio and by selectively introducing chondroitin sulfate and/or chemical crosslinking. In particular, carbodiimide-crosslinking created a covalent bonding of the amine groups to the carboxylic groups. Chondroitin sulfate seemed to bind to collagen scaffolds more than to collagen–elastin scaffolds and to increase the water uptake by the composite hydrogel. Collagen improved the control of the scaffold mechanical properties and porosity. Biocomposites have been fabricated from continuous sheets of elastin containing crosslinked HA (Ramamurthi and Vesely, 2005). Such composites have been proposed for the replacement of the native aortic valve cusp. After two weeks of culture, the composite biomaterial appeared to support the proliferation of neonatal rat aortic smooth muscle cells, but there was no signifcant difference in elastin secretion when compared with control tissue cells. At longer incubation times, the cells produced amounts of collagen lower than controls and a matrix layer rich in elastin was observed at the cell/biomaterial interface. The newly deposited elastin was organised into fenestrated sheets and loose fibers, a histological pattern typical of the aortic valve cusp. Similarly, the work produced by Daamen et al. (2003), the porosity of collagen–glycosaminoglycan (CG) scaffolds for tissue engineering has been modulated by freeze-drying conditions. As the conventional freeze-drying process leads to a heterogeneous porosity, the rate of freezing has been recently controlled by the the use of smaller, less warped pans. The resulting hydrogels were characterised by pores more homogeneous in size and shape. In another study, fibrillar collagen has been assembled with HA and heparin and the fibril formation kinetics has been analysed upon changes of GAG concentration (Salchert et al., 2004). This method was applied to the coating of polymeric surfaces. Electron microscopy and atomic force
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microscopy showed that the addition of GAG significantly changed both fibril size and mesh morphology. HA has also been combined with fibrin glue (Park et al., 2005). As discussed in section 4.2, the rationale for combining these two biopolymers into a composite biomaterial is that fibrin glue has been proven as a suitable carrier for cell delivery in cartilage tissue engineering and that HA is a key component of the articular cartilage. When the fibrin/HA composite was implanted in nude mice up to four weeks, cartilage-like tissue formation seemed to be faster than homogeneous fibrin glue hydrogels. Significantly higher amounts of GAG and collagen were observed along with the expression of markers typical of articular cartilage. Lyophilisation has also allowed to modulate the availability and spacial distribution of the bioligands naturally present in ECM biopolymers and to combine it with a pore size appropriate for tissue in-growth. As previously shown, collagen–glycosaminoglycan three-dimensional scaffolds did not promote tissue in-growth when the mean pore size was either lower than 20 μm or larger than 120 μm, scaffolds with a constant composition and solid volume fraction (0.005) were produced by varying the freezedrying process (O’Brien et al., 2005). Although these scaffolds were tested for their ability to support cell line osteoblasts rather than by cells relevant to soft tissue repair, it is worth observing that cell attachment and viability over 24 and 48 h were tightly dependent on the pore size. In particular, the attachment of viable cells increased linearly with the scaffold specific surface area. The uniform pore size and architecture has a significant impact not only on the biological properties of the hydrogels, but also improve the control of their mechanical properties. The effect of the pore microstructure of collagen–GAG scaffolds on their tensile, compressive and bending mechanical properties has been studied (Harley et al., 2007). The various scaffolds analysed showed stress–strain behaviour typical of low-density, open-cell foams where distinct characteristics in linear elasticity, collapse plateau and densification regimes could be observed. Scaffolds with pores of uniform diameter were found to be mechanically isotropic. Specific engineering methods have been adopted to mimic the histologically more complex architecture of certain tissues. A typical example is the mimicking of the highly organised structure of the articular cartilage. In this tissue, cells and ECM are organised in well defined zones and their overall assembly determines the tissue biomechanical functionality. To this purpose, an anisotropic pore architecture has been generated in scaffolds based on ECM components to influence the zonal organisation of chondrocytes and ECM components (Woodfield et al., 2005). A novel 3D fibre deposition technique was developed that allowed the design and manufacturing of scaffolds with a completely open-pore structure where each cavity
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was homogeneously spaced by fibres of reproducible diameter. The same technique also allowed the formation of pore-size gradients where the overall porosity and volume fraction were kept constant. In vitro experiments with chondrocytes showed that these pore-size gradients favoured anisotropic cell distribution resembling the different zones present in the superficial, middle, and lower zones of immature bovine articular cartilage. When these cell culturs were prolonged, an inhomogeneous distribution of both zonal GAG and type II collagen deposition was observed. Novel nanofibrous collagen–GAG scaffolds were obtained by electrospinning of collagen fibres blended with chondroitin sulfate in a mixed solvent of trifluoroethanol and water (Zhong et al., 2005). When collagen fibres were electrospun in the presence of 4% (by weight) chondroitin sulfate, the fibres acquired a nanoscale diameter and homogeneous structure. The increase of the chondroitin sulfate concentration to 10% (by weight) reduced the fibre diameter further, but the diameter distribution was broader, probably owing to a change in the solution properties that affected the electrospinning process. The treatment of these scaffolds by glutaraldehyde vapour led to crosslinking and, as a consequence, to a higher stability of the scaffolds to collagenase-induced degradation. The improved stability also improved the proliferation of rabbit conjunctiva fibroblasts. A two-level fractional factorial design has been used for a precise coimmobilisation of chondroitin-4-sulfate, chondroitin-6-sulfate, dermatan sulfate and heparin to chitosan membranes to produce scaffolds for cartilage tissue engineering (Chen et al., 2006). The fractional factorial design allowed the discrimínate study of the effect of each type of GAG on cell behaviour. In particular, it was shown that, changing the levels of chondroitin-4-sulfate, collagen synthesis was promoted, but not cell proliferation. Conversely, high levels of heparin and dermatan sulfate induced cell proliferation, but not the production of collagen and GAG. Moreover, GAG/chitosan membranes based on relatively low levels of chondroitin sulfate and heparin preserved the chondrocyte phenotype as assessed by cell morphology and marker gene expression. Scaffolds obtained by the combination of insoluble type I collagen fibres with soluble type II collagen fibres, insoluble elastin fibres, GAG and growth factors has allowed the tailored preparation of several bioactive biocomposites able to stimulate tissue repair (Geutjes et al., 2006). The careful purification of the used macromolecules has been considered an important, initial step. Their assembly has followed the bottom-up approach where several of the techniques described above were adopted. The collagen to elastin ratio, the freezing rate, the degree of crosslinking, and the GAG percentage were all changed and their effect was studied in vivo in a rat subcutaneous model. Histology showed that the crosslinked scaffolds preserved their structure for the period of time required to encourage the
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formation of endogenous ECM. Furthermore, when loaded with FGFb, these hydrogels significantly enhanced neovascularisation and tissue remodelling. In a similar bottom-up approach, microtextured basal lamina analogues were developed to investigate the relationship between the skin dermis– epidermis interface and the behaviour of keratinocytes (Downing et al., 2005). By a series of microfabrication techniques, patterns, negative replicates, and collagen membranes with ridges and channels of length scales similar to the basement membrane invaginations were obtained. Keratinocytes seeded for seven days formed a differentiated and stratified epidermis following the microtextured membrane morphology. Keratinocyte stratification and differentiation seemed to increase as channel depth increased and channel width decreased. However, ECM components are not the only natural polymers used for the manufacture of composite biomaterials for soft tissue repair. Indeed, biopolymers such as chitin, chitosan and alginate have also been used and engineered by different methods including chemical crosslinking and freeze-drying. In particular, chitosan has been tested as a candidate biomaterial for preparing composites for various clinical applications. For example, chitosan/gelatin complexes were fabricated by freeze-drying (Xia et al., 2004). As discussed in section 4.2, the rationale of using these two biopolymers is that chitosan is a polysaccharide potentially able to mimic the GAG and collagen structures and to exert a similar biological activity on cell adhesion, differentiation, and proliferation. Autologous chondrocytes were isolated from pig auricular cartilage and seeded onto these chitosan–gelatin scaffolds showing that elastic cartilages could be formed upon subcutaneous implantation in the porcine abdomen after 16 weeks of implantation. A detailed immunohistological analysis seemed to demonstrate that the chondrocytes were enclosed in type 11 collagen-positive chondrons similar to that of native cartilage. The presence of natural elastic fibers and GAG in the engineered cartilages was also shown. Biomechanical tests provided evidences that the extrinsic stiffness of the engineered cartilages was very similar (85%) to that of the native auricular cartilage. Chitin and chitosan were crosslinked by genipin, a natural dialdehyde, under various prefreezing temperatures to form tissue-engineering scaffolds by freeze-drying (Kuo et al., 2006). Hydroxyapatite was subsequently added to the scaffolds by a precipitation method to obtain scaffolds able to support chondrocyte attachments and growths. Relatively lower prefreezing temperatures or a higher chitin percentage led to smaller pore diameter, to greater porosity, to larger specific surface area, to higher Young’s modulus, and to a lower extensibility. Similarly, decreased prefreezing tem-
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perature and increased chitin percentages and hydroxyapatite loading improved cell proliferation as well as endogenous GAG and collagen deposition. Crosslinking of biopolymers with synthetic biomaterials has also been attempted (Santin et al., 1996) (Fig. 4.1). Crosslinked poly-(2-hydroxyethyl methacrylate) (PHEMA) hydrogels with a 70% water content have been
(a)
(b)
4.1 Scanning electron microscopy of a poly(2-hydroxyethyl methacrylate)/gelatine interpenetrating polymer network: (a) control composite biomaterial, (b) composite biomaterial showing degradation of the gelatine component after 14 days subcutaneous implantation in rats (modified from Santin M et al., 1996).
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interpenetrated with gelatine to improve the ability of the hydrogel to encourage cell adhesion in vitro and soft tissue in-growth in vivo (Santin et al., 1996). Hydrogels for chondrocyte encapsulation under physiological conditions have been developed by the combination of alkali-treated collagen with pentaerythritol poly(ethylene glycol) ether tetrasuccinimidyl glutarate (4S-PEG) as a crosslinker (Taguchi et al., 2005). Chondrocytes can be uniformly dispersed into these hydrogels where they were shown to proliferate, to secrete GAG and to express type II collagen and aggrecan over time. The crosslinking of the gel encapsulating the cells was thus shown not to affect cell viability. Therefore, these in situ setting hydrogel systems can be injected in vivo to become a valid alternative in minimally invasive surgery. Three-dimensional collagen/chitosan/GAG systems enriched by TGF-β1 have also been engineered where the growth factor has been loaded into chitosan microspheres by an emulsion crosslinking method (Lee et al., 2004). The porous scaffolds were crosslinked by 1-ethyl-3-(3-dimethylaminopropyl)carbodiimide in the presence of chondroitin sulfate. The TGF-β1 microspheres were encapsulated into the scaffold and, finally, chondrocytes were seeded in the hydrogel system and cultured for three weeks. The presence of TGF-β1 microspheres in the hydrogel clearly stimulates the chondrocytes to produce new cartilage ECM. These are typical examples of composite biomaterials for soft tissue repair obtained by the combination of ECM biopolymers and other natural polymers. The need for mimicking the specific physicochemical and bioactive properties of the natural tissues has led to more specific studies where the composite biomaterials have been prepared by considering the needs of the specific clinical applications. In view of this approach, additional examples of biocomposite biomaterials based on natural polymers will be provided in section 4.4.
4.3.2 Soft tissue repair properties of synthetic, hybrid and biomimetic material-based composites The examples of biocomposites made of ECM and other biopolymers have shown that a number of methods are now available for finely tuning their physicochemical and bioactive potential. However, the use of composite biomaterials based on natural polymers still suffers from batch-to-batch variations linked to the source of these biopolymers. For this reason, efforts have also been made to develop synthetic composite biomaterials able to mimic the ECM biochemical and structural features. Synthetic polymers such as polylactic and polyglycolic acid polymers and poly(ε-caprolactone), materials approved by the American Food and Drug Administrations, have been used in combination with each other or with
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other natural and synthetic biomaterials to provide composite biomaterials for soft tissue repair. Porous sponges based on these biomaterials have been studied for their ability to induce mouse MSC differentiation into chondrocytes (Aung et al., 2002). After one month in specific culturing conditions, the sponges appeared to form abundant GAG and collagen amounts. When implanted into nude mice, the sponges loaded with MSCs were found to induce cartilage repair. Starch-based biodegradable composites have been engineered that have shown to trigger a modest immunological reaction both in vitro and in vivo (Marques et al., 2005). Blends of corn starch have been obtained with ethylene vinyl alcohol, cellulose acetate and PCL. While the presence of cellulose acetate and PCL seemed to stimulate higher levels of immune response in an in vivo rat model, this reaction was at an acceptable level. Composite films made of poly-(d,l-lactide)/45S5 Bioglass composite films have been engineered for inducing the repair of the annulus fibrosus, the external part of the intervertebral disc (Wilda et al., 2006). Various concentrations of Bioglass were included in the polymeric matrix. These composite films appeared to be a suitable substrate for annulus cells which produced ECM where GAG and collagen were abundant. The surface of poly-(l-lactic acid) (PLLA) film by chitosan has also been performed by coupling with 1-ethyl-3-(3-dimethylaminopropyl)carbodiimide and N-hydroxysuccinimide (Cui et al., 2003). The composite improve its potential as a substrate for chondrocytes, which proliferated and increased their secretion of ECM components typical of the cartilage such as GAG and type II collagen. The possibility of finely tuning the chemistry and structure of the synthetic polymers systems has led to the synthesis and fabrication of composites at nano-scale levels. Synthetic biodegradable dermal substitutes have been fabricated from electrospun polymer scaffolds based on PLLA, PLLA and 10% (w/w) oligolactide, PLLA and rhodamine and three poly-(d,l)-lactide-co-glycolide (PLGA) random multiblock copolymers and applied to the tissue engineering of skin and oral mucosa (Blackwood et al., 2008). When tested in vitro, the PLGA 85:15 and 75:25 scaffolds were shown to favour proliferation of keratinocytes, fibroblasts and endothelial cells and the deposition of new ECM deposition. In attempting to reduce mismatch between the mechanical properties of implants and those of the surrounding tissue, nanostructured composites based on polyvinyl alcohol and bacterial cellulose nanofibres have been produced (Million et al., 2008). A novel thermal processing method under an applied strain was performed where a small amount of bacterial cellulose
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nanofibres was added to form an anisotropic nanocomposite. When tested for its mechanical properties, this nanocomposite showed stress–strain tensile properties similar to that of the porcine aorta. It has been suggested that by varying the processing conditions, nanocomposites with controlled mechanical properties can be manufactured to match closely the mechanical properties of several types of soft tissues. Synthetic nanocomposites have also been obtained by the simultaneous incorporation of highly dispersed carbon nanotubes (CNT) in polycarbonate urethan (PCU) (Khang et al., 2008). Both the nanoscale roughness of the surface and the application of an external electrical stimulation enhanced chondrocyte activity at a level higher than the control PCU. This improved cell activity was linked to higher amounts of adsorbed fibronectin. Multi-walled carbon nanotubes (MWNTs) have also been modified by a more biocompatible synthetic polymer such as the poly(2-hydroxyethyl methacrylate) [poly(HEMA) (Kumar et al., 2008)]. HEMA monomers have been grafted onto the MWNTs by free radical polymerisation. To achieve grafting, the nanotubes were previously oxidised with a mixture of nitric acid and sulfuric acid to produce carboxylic acid functionalities on their surface. The grafting of HEMA was then obtained by esterification. Although the use of MWNTs in combination with more traditional biomaterials seems to open new routes of investigation, it has to be stressed that their safety in vivo is very questionable. Indeed, carbon nanotubes may be potentially toxic to the cells when liberated from the composite. Their functionalisation with more biocompatible biomaterials (e.g. HEMA) may not be necessarily sufficient to avoid this drawback.
4.3.3 Biocompetent nanocomposites The use of synthetic nanoscale biomaterials offers new opportunities to introduce biocues which are specific to the repair mechanisms of the different types of tissues in a more controlled manner. It is believed that, in the future, biomedical commercial products based on nanostructured composite biomaterials will improve the reproducibility of both manufacturing and clinical performance. It has been shown that a collagen mimetic peptide can be synthesised by assembling peptides that include in their structure the specific amino acid sequence -Pro-Hyp-Gly- (Lee et al., 2006). These peptides are able to form a triple helix conformation that resembles native collagen fibres. Poly(ethylene glycol) (PEG) hydrogel which is known to be a non-adhesive biomaterial has been conjugated with this collagen mimetic peptide to improve the interaction of the hydrogel with cells. This bionanocomposite was synthesised by adding the peptide mimic to copolymers of poly(ethylene oxide) diacrylate. Chondrocytes were encapsulated in the composite hydro-
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gel showing after two weeks an increase in both glycosaminoglycan (87%) and collagen (103%) content higher than control PEG hydrogels. Self-assembling, macroscopic composite sacs and membranes have been produced at the interface between two aqueous solutions, one containing a megadalton polymer and the other, small self-assembling molecules bearing an opposite charge (Capito et al., 2008). The resulting structures show a highly ordered nanofibre architecture with bundles gradually aligning and reorienting at an angle of approximately 90°. It has been suggested that the formation of a diffusion barrier at the liquid–liquid interface prevents their random mixing and that the ordered membrane structure is dicated by the effect of both osmotic pressure of ions and static self-assembly. The resulting hydrogel may be exploited for soft tissue repair and other applications where a controlled biomimicking environment is required. In addition to simulating structural elements of the ECM, these bionano-composites are able to target specifically the cells through interactions of the substrate with their receptors. For example, the need for promoting vascularisation in most of the repairing tissue (except cartilage), has prompted research for the development of angiogenic bio-nanocomposites. These materials mimic the angiogenic properties of natural EMC components such as type I collagen, laminin and elastin. Although they have not always been integrated in a composite biomaterials, two examples of functional biomimetic peptides with a potential role as composite building blocks are worth mentioning. These are the elastin-like polypeptides and the ECM ligand-bearing peptides. Elastin-like polypeptides (ELPs) are composed of a pentapeptide repeat, Val-Pro-Gly-X-Gly that undergoes an inverse temperature phase transition. Although soluble in aqueous solution, these peptides aggregate when the solution temperature is raised above their glass transition temperature. Cell culture studies have shown that chondrocytes encapsulated in these hydrogels were able to preserve their phenotype thus suggesting their potential as injectable, self-setting materials for cartilage repair. Artificial ECM protein sequences have been synthesised and integrated to obtain polypeptide structure having both collagen-binding activity and active functional units able to promote capillary network formation by promoting vascular endothelial cells (Nakamura et al., 2008). A lamininderived IKVAV sequence, able to stimulate capillary network formation by the proliferation of vascular endothelial cells, was incorporated into an elastin-derived structural unit. The obtained fusion protein also presented the cell-adhesive RGD sequence and a collagen-binding domain derived from fibronectin. As a consequence, this fusion protein was shown to bind to collagen type I and promote angiogenic activity of collagen gel maybe through its proven role in promoting cell migration.
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4.4
Use of biocomposites in clinical intervention for soft tissue repair
4.4.1 Wound dressing Composite biomaterials are widely used as primary or secondary wound dressings. The general composition of these dressings consists of three layers: (i)
The semi-adherent or non-adherent layer is in direct contact with the wound and ensures that the repairing tissue does not contact other tissues. The semi-adherent or non-adherent nature of this layer allows the removal of the dressing with no interaction with the repairing tissue. Also, the hydrogel nature of this layer allows the diffusion of exudates towards the absorptive layer and the retention/release of antibiotics at the site of repair. (ii) The absorptive layer contributes to drain exudates and removes debris from the wound area. This action prevents skin maceration and bacteria growth, while maintaining a moist environment. To this purpose, hydrogels, hydrocolloids and foams have been employed. By preventing an excessive presence of exudates in the wound, this layer also contributes to the digestion of eschar and necrotic tissue by autolytic debridement. (iii) The bacterial barrier allows the diffusion of moisture vapour, while it prevents bacterial cell penetration into the absorptive layer. The most recent products have a bacterial barrier that reduces the leakage of exudates from the underlying absorptive layer thus reducing the needs for frequent dressing changes. There are many composite dressing currently available on the market. Among them, Tegaderm is a typical example of a class of dressing made of hydrocolloid adhesives, foams and alginate. These types of dressing have been used since the early 1970s and the most used formulations have an absorptive layer based either on alginate or on acrylic polymers. Viasorb is another multi-layered dressing where the absorbent layer is based on a cotton/polyester gauze. Smith and Nephew have commercialised several wound dressing types based on polyurethan films. It is claimed that in the Opsite dressing, a MVTR reactive hydrophilic polyurethan film can switch its physicochemical properties allowing water vapour to evaporate depending on the amount of exudates in the wounds. In the case of excessive exudates, the film will relax its structure allowing more evaporation.
4.4.2 Burn treatment Dermal graft substitutes are the most typical example of clinical use of composites for soft tissue regeneration. Skin substitutes have been engineered
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which are based on collagen/GAG composite 3D scaffolds. These scaffolds are commercially available as both acellular and cellular grafts and they compensate for the shortage of skin autografts (e.g. split skin grafts) in those cases where extended skin damages (e.g. extended burnt areas) occur (Ventre et al., 2009). Once epidermis is damaged there is a need both to protect the wound from microbial contamination and to prevent fluid loss. Later, the substitute has to support cell colonisation and, ultimately, the formation of a new vascularised tissue. Depending on the degree of burn depth, skin substitutes, which have been optimised for either wound cover or wound closure, can be used. Typical skin substitutes for wound covers are as follows. Biobrane (www.burnsurgery.com/Modules/skinsubstitutes/sec5.htm) This skin substitute is characterised by a bilayered structure where a nylon mesh is covered by a thin silicone membrane. The nylon filaments are functionalised by collagen-derived peptides that can improve the substitute integration with the surrounding wound bed through the adhesion of tissue cells. The silicone surface is semipermeable and prevents fluid loss and microbial contamination. The main use of Biobrane is as a temporary cover and it is removed either upon wound healing or before autografting. Transcyte (www.bu.edu/woundbiotech/bioengineered/TransCyteProdPres/ index.htm) From a biomaterial viewpoint, Transcyte bears similarity to Biobane as it is composed of a nylon mesh coated by collagen. However, this skin substitute is provided in cellular form as its mesh is seeded with allogenic fibroblasts extracted from neonatal foreskin. Before grafting, the device undergoes a freezing treatment to devitalise the cell components and preserve the newly synthesised structural proteins and growth factors. Such a treatment is performed in order to reduce the risks of immune reaction. It has to be emphasised that, since nylon is not biodegradable, both Biobrane and Transcyte cannot be strictly considered as dermal substitute. Their use is indicated as a temporary application in partial thickness burns before autograft implantation. Wound closure is achieved through the use of dermal substitutes made of composites of a complete bioresorbable nature. Among them the most used is Integra (www.integra-ls.com/products). Integra is a dermal substitute made of a bi-layer structure composed of a collagen-glycosaminoglycan sponge which is covered by a silicone membrane. The performance of this composite products have been closely examined and its clinical use is widespread. It has been demonstrated that the porous architecture of this substitute allows the migration of fibroblasts and other cell form the wound bed. After repopulation, cells start to
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remodel the scaffold by producing their own ECM and allowing the ingrowth of microvasculature. The silicone membrane temporarily mimics the epidermal layer that protects the wound from the external environment while ensuring oxygen permeation. Within three to six weeks, this silicone membrane is removed and substituted with a thin split skin autograft. For these properties, Integra is used in clinics as a wound closure device for both partial and full thickness burns. This product has gained widespread use, in particular, in the clinical treatment of third degree burn wounds and full thickness skin defects of different aetiologies. The product significantly reduces the treatment timescale and achieves final wound healing. However, the treatment of large tissue areas with this composite still leads to scar formation. To improve the clinical outcome, experimental work has been undertaken to develop a method which will allow the homogeneous integration of human keratinocytes in the scaffold. In vivo models in small size animals (athymic mice) suggest that the uniform distribution of keratinocytes allows a better wound adherence of the substitute as well as reduced wound contraction and better histological features of the healed tissue (Kremer et al., 2000). Both the dermal and epidermal layers were regenerated by the cells of human origin. Indeed, the use of cultured skin substitutes has been developed to improve the poor mechanical properties of the acellular scaffolds. It has been observed that the progressive deposition of new ECM can improve the stability of these rather soft composite scaffolds. The keratinocytes have been shown to form stratified and cornified layers separated by the dermal layer by a basement membrane (Boyce et al., 2002). Although not yet available to the surgeons, new composite scaffolds based on collagen/GAG/chitosan have been developed that can be seeded with haematopoietic progenitor cells which may improve the scaffold vascularisation without the use of growth factors (DezutterDambuyant et al., 2006). This integration was pursued with the additional seeding of keratinocytes thus providing a more complete environment where skin and vascular components were integrated to achieve a dermis equivalent.
4.4.3 Post-surgical adherence The use of meshes based on composite biomaterial has become routine practice in the surgical repair of all hernias. The use of these composite polymeric meshes has significantly reduced the hernias recurrence rates (Doctor, 2006). Composite biomaterials are widely used to prevent post-surgical adherence. They can be divided into three main classes:
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Lightweight composite meshes without barrier Absorbable barrier composite meshes Nonabsorbable barrier composite meshes.
Lightweight composite meshes without barrier VYPRO II and ULTRAPRO (Johnson and Johnson, India) are composite meshes designed to support weak tissues in cases of inguinal hernias. They are manufactured using filaments of Vicryl and Prolene or Monocryl and Prolene. The filaments undergo twisting and subsequent knitting to provide a well defined mesh architecture. The mesh is macroporous with a pore size of 4.5 mm that allows tissue in-growth and the integration into a strong 3D collagen fibrotic tissue. The integration of this mesh is also favoured by the presence of the Vicryl or Moncryl components which make this type of mesh partially resorbable. The partial resorption and tissue in-growth reduce the mesh of approximately 70% and leads to the formation of the so called ‘scar-mesh’ instead of a thinner ‘scar-plate’. These types of meshes achieve both the goal of a re-enforced tissue and an acceptable mobility of the abdominal wall. However, these meshes are not able to fulfil the current repair for minimally invasive surgery. Therefore, new meshes have been designed which can be implanted through small incisions and that can achieve tissue repair only from the abdominal wall side, while non-adhesiveness is required for the side of the mesh facing the bowel. To this purpose, absorbable and nonabsorbable barriers have been engineered. These composite meshes share the same design concept that is to present the abdominal cavity with a material that prevents adhesions with the bowel and a layer based on either polyester or polypropylene or PTFE that faces the abdominal wall and encourages tissue in-growth. Absorbable barrier composite meshes Several absorbable barrier composites are available on the market: •
•
Sepramesh biosurgical composite (Genzyme Biosurgery, Cambridge, USA) is a dual-component prosthetic biomaterial composed of macroporous polypropylene on one side, with bioresorbable, nonimmunogenic membrane of sodium hyaluronate and carboxymethyl cellulose on the other side. Seprafilm was designed to provide protection against intra-abdominal adhesion formation throughout the critical period of remesothelialisation during the first post-operative week. This absorbable barrier is able to turn into a gel in 48 h and it is able to coat the mesh for approximately 7 days and be finally cleared by the body within 28 days. Seprafilm is an
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Biomedical composites antiadhesive hydrogel that prevents adherence and reduces viscosity at the tissue interface. As a consequence, adjacent tissues can heal simultaneously, but in an independent manner. Furthermore, as the sodium hyaluronate and carboxymethyl cellulose are anionic polysaccharides, they can form a negatively-charged barrier that seems to promote the separation of healing tissues. Parietex composite and Parientene composite meshes (Sofradim, France) are composed of multifilament polyester mesh that is combined with a purified, oxidised bovine atelocollagen type I coating and then covered by an absorbable, antiadhesion film of polyethylene glycol (PEG) and glycerol. The relative inertness of PEG hydrogels and the lubrication properties of glycerol reduces the formation of tissue adherence. The collagen coating functions to promote ECM formation on the polyester mesh and to decrease the reaction of the fibrous tissue to the foreign material. The collagen, polyethylene glycol and glycerol films are resorbed within approximately three weeks of implantation. Parientene composite mesh comprises the same antiadhesive barrier but coated to polypropylene. PROCEED surgical mesh (Johnson and Johnson, India) is a sterile and multilayered mesh particularly attractive for its thin, flexible and laminate architecture that includes an oxidised regenerated cellulose fabric and Prolene mesh. This mesh is subsequently encapsulated within a polydioxanone polymer. The polypropylene mesh favours tissue ingrowth, whereas the oxidised regenerated cellulose plays a role as a bioresorbable layer that, upon implantation, separates the mesh from organ surfaces thus minimising undesired adhesions. The polydioxanone is used to bind the cellulose to the rest of the composite. Such a mesh allows tissue repair with a minimal fluid loss and reduced risk of infections.
Nonabsorbable barrier composite meshes This is a class of mesh that is also widely used in surgery. Amongst the best known ones are: •
•
Bard composix mesh (Davol Inc, Cranston, USA) is composed of a polypropylene mesh coated with a thin film of non-adhesive ePTFE. This mesh can be inserted by minimally invasive laparascopy, although this is a difficult procedure as the mesh cannot be easily compressed through the laparoscopic port. For this reason, its implantation is performed with relatively larger ports (12 mm). GORE-TEX dual mesh (W. L. Gore, USA) is a device showing two sides with distinct surface characteristics; one side is relatively smooth and
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presents micropores of 3 mm, while the other is rough and it has a porosity of approximately 22 mm. The relatively smooth surface faces the visceral organs thus reducing adhesion, while the rougher surface promotes tissue integration from the abdominal wall. This mesh is made available in two formulations which are a solid sheet and a perforated sheet. The latter seems to facilitate tissue integration. More recently, a third formulation has been made available that includes antibacterial agents such as silver and chlorhexidine into the ePTFE layer. The use of novel absorbable and nonabsorbable barriers on composite meshes to reduce the incidence of adhesion and adhesion-related complications has been evaluated in animal models and a few clinical studies have been reported. The incidence of adhesions and tenacity of adhesions were reduced for all of the barrier meshes compared with the macroporous polyester mesh. Although these new prosthetic composite biomaterials have been developed to prevent adhesion in cases of intra-abdominal placement of mesh during open and laparoscopic hernia repair, no long-term follow-up is yet available and their advantage over macroporous mesh with no antiadhesive barrier still needs to be proved. In general, the type of mesh should be chosen by the surgeon on the basis of the type of surgery. For intra-abdominal placements, meshes able to prevent bowel adhesions should be used. Another important issue is the size of the mesh. Upon inguinal hernia implantation, these meshes have to cover an area of at least 15 cm × 15 cm, whereas for repair of umbilical, ventral and incisional hernia, their size should exceed that of the defect by at least 4 cm. The defect should be exceeded by the mesh in all directions as a smaller size may lead to protrusion of the mesh into the defect and risk of relapse. The composite meshes have also to sustain the mechanical stresses related to their suturing. Usually, in the case of intra-abdominal placement of the mesh, a few strong sutures should be applied at the four corners of the mesh and subsequently fixed by anchors at a distance of 3 cm all around to prevent any bowel obstruction in between the mesh and the defect. Upon suturing, the mesh should not contract. Recently, a mesh has been developed that is made of a collagen/GAG matrix combined with polypropylene. Such a composite device was manufactured by interposing the synthetic polymeric mesh within the porous collagen/GAG matrix. This biopolymer composite coating was finally crosslinked by glutaraldehyde. When compared with non-coated meshes, at four weeks of implantation, the composite implants induced the formation of a connective tissue-like structure preventing bowel adhesions which were taking place in the non-coated synthetic mesh.
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4.4.4 Cardiovascular applications Composite biomaterials play a key role in cardiovascular applications. Composites have been engineered to fabricate vascular grafts and heart valves. In this section, typical examples of implants based on the concept of biocomposite biomaterials are provided. Vascular grafts Vascular grafts are used to replace, bypass or maintain the functionality of vascular structures, which have been altered or damaged by pathological conditions or traumas. Graft materials are either of natural source (tissue harvested either from the same patient or from a donor) or of synthetic composition. To respond to the mechanical and biocompatibility requirements, composites grafts including synthetic and/or biological components have been engineered to bypass blood circulation. The rationale underpinning these biosynthetic prostheses is to include in the final device a synthetic material able to fulfil the mechanical requirements and a biological component able to improve the implant biocompatibility. The presence of components from a natural source is particularly important to favour the integration of the implant with the surrounding tissue. In a completely synthetic graft, this is pursued mainly by an enginnering prosthesis with an appropriate porosity able to encourage tissue in-growth (Ramakrishna et al., 2001). In biocomposite devices, the presence of a biocompetent substrate such as the collagen is meant to act at the cellular level by mediating biospecific interactions between the device and the host cells. A typical example of a biosynthetic graft is the Omniflow II (Bio Nova, Melbourne, Australia, www. bionova.com), a type of graft that is indicated for the repair of peripheral arteries with a diameter of 5 to 8 mm and length of 20 to 65 cm. Omniflow II is composed of a polyester tubular mesh that is embedded into a collagen gel. A crosslinking technology seems to generate a particularly haemocompatible surface able to minimise clot formation (Koch et al., 1997). During the manufacturing process, the polyester mesh acts as a scaffold for the ordered deposition of collagen fibres providing mural strength to the composite. Heart valves Aortic valves can undergo serious damage to their structure and consequent impairment of their functionality. The replacement of damaged aortic valves with mechanical valves and Dacron tubes has become a common surgical procedure. Xenogenic materials (e.g. fixed and acellular porcine tissue) is usually used in combination with a polymeric stent supporting the valve. However, the xenogenic materials used suffer degeneration and calcification and revision surgery is soon needed to replace all the conduit (Wheatley
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and Will, 2005). For this reason, attempts have been made to obtain more stable prostheses. A typical example is the use of composite grafts in combination with a stentless valve (Urbaniski and Hacker, 2000). More specifically, stentless porcine valves were assembled with a collagen-coated woven polyester vascular prostheses. As well as the good mechanical and biocompatibility properties, this type of artificial valve has the advantage of minimising the invasiveness of possible revision surgery. In fact, the composite nature of the device allows the surgeon to remove only the deteriorated biological component leaving the tube graft in place.
4.4.5 Lower urinary tract The reconstruction of urinary tissues significantly depends on the availability of biocompatible biomaterials. Among the various biomaterials that have been proposed and tested to this purpose (Pariente et al., 2001), a study is worth mentioning that has shown the potential of composite biomaterials in a clinical investigation on a relatively large number of patients (Trabucco et al., 2004). The aim of this study was to test the performance of a composite mesh sling to support the surgical correction in cases of stress urinary incontinence. The sling was designed to prevent post-surgical complications such as chronic retention and urethral erosion which are usually associated with the implantation of synthetic polypropylene slings. Furthermore, the new composite sling design was conceived to provide the surgeon with a biomaterial easy to handle and implantable by a minimally invasive procedure. This design consisted of a tension-free sling (T Sling), a mesh composed of two different materials; a monofilament, nonabsorbable polypropylene mesh and an absorbable polydioxanone monofilament. This material was FDA and CE approved and implanted into 40 patients with urethral hypermobility by the Stamey’s procedure. A follow-up study from 15 to 70 months, showed a positive outcome (97.5%) in terms of incontinence resolution that was accompanied by no significant symptom of voiding irritation. Also, no morbidity or complications were encountered. It is believed that the mechanical properties conferred to the mesh by its composite nature as well as the better tissue integration facilitated by the bioresorbable polymer component may have led to this improvement over other noncomposite biomaterials.
4.4.6 Clinical applications lacking composite biomaterialaided soft tissue repair Cartilage Articular cartilage undergoes clinically relevant lesions which are a consequence of either traumatic events or pathological conditions, among which
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osteoarthritis and osteochondritis dissecans are the most frequent. This type of damage cannot undergo adequate self-repair, but rather results in the formation of fibro-cartilage (Oakes, 2009). This is a repair tissue where the cartilage proteoglycans are deposited together with a relatively high amount of collagenic matrix with fibrotic features. Such a tissue does not match the mechanical properties and, therefore, the function of the original healthy cartilage. In the attempt to achieve full repair of the damaged cartilage, several replacement techniques have been suggested in the last three decades. Unfortunately, none of these surgical approaches has yet provided a satisfactory clinical outcome. The most successful approach is considered to be the autologous chondrocyte implantation (ACI). This technique is based on the harvesting of chondrocytes from a different anatomical area of the same patient, their expansion in culture and their final re-implantation into the damaged site. Although bearing a great potential, the clinical outcome of this technique is limited by the tendency of chondrocyte to dedifferentiate upon culturing conditions. Biomaterials able to simulate the natural environments are therefore required in both the culture expansion step and implantation procedure. Currently, several biomaterials which have been developed up to pre-clinical stage have the potential to mimic the natural chondrocyte microenvironment and have been used as carriers for either chondrocytes or progenitor cells (Galois et al., 2004). Among them, only few are composites and, to our knowledge, none of them has yet been tested in humans. These composites are mainly based on combinations of either natural polymers such as polysaccharides and collagens or synthetic polymers and ceramic materials (e.g. PLGA/TCP) (Sherwood et al., 2002). Peripheral nerves Neurosurgeons are routinely confronted with serious cases where peripheral nerves (typically those innervating hands) are severed by traumatic events. As the nervous system physiology and structure are complex, the clinical strategies to be adopted for their repair have to rely on neural guides able to drive the healing of the different tissue components. In particular, the extension of the neural cell protrusions, known as axons, has to be guided to re-establish connection with the distal part of the severed nerve and, therefore, functionality. Although the regeneration of the axons is possible over short distances, in the case of total transection of the axon including its myelin sheath and endoneural tube (neurotmesis), the socalled Wallerian degeneration is triggered where a series of complex cellular events involving Schwann cells, macrophages, and monocytes is activated leading to the progressive degeneration of the distal part of the nerve. A rapid intervention is thus required that can stimulate the Schwann cells to proliferate and support the axonal out-growth. At the same time,
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the activation of macrophages needs to be controlled and limited to their scavenger role towards tissue debris and the migration of fibroblasts needs to be inhibited to reduce the formation of a fibrotic tissue, the neuroma that blocks the axonal repair process. Currently, the most common method for peripheral nerve repair is the autologous nerve grafting procedure where the axons of the proximal nerve stump are reconnected to their targets by the surgical implantation of a nerve that has been harvested by another anatomic site (e.g. sural nerve) of the same patient. Obviously, this technique leads to loss of donor site function, donor site morbidity, and the need for multiple surgeries in order to harvest the nerve before it is grafted. In addition, it is very difficult to harvest nerves the size of which matches that of the nerve to be repaired. As a result, the implantation of autologous nerve grafts is often inadequate. For these reasons, peripheral nerve repair has attracted the attention of many material scientists who have developed neural guides based on synthetic and natural, stable and absorbable nerve guides. These hollow tubes direct the axonal regeneration from the proximal to the distal stump, while protecting it from the invasion of fibrotic tissue. However, studies in humans using neural guides have led to contradictory results. Although not fully exploited, composite materials seem to bring some advantages over conventional silicone- or collagen-based neural guides. Indeed, composite neural guides comprised of biodegradable synthetic composite materials such as PGA and polylactide–caprolactone have been shown to have better mechanical properties and to favour a better tissue in-growth and integration. This is a significant advantage over the nonresorbable, composite neural guides made of either silicone or polytetrafluoroethylene which, while demonstrating the potential to support axonal regeneration, suffer from disadvantages such as limited or no flexibility and compression of the regenerated axons leading to chronic pain and discomfort. However, it is accepted that these composite, biodegradable biomaterials are useful only in cases presenting relatively short neural defects. For critically sized nerve lesions, engineered constructs that provide increased physical support and biological activity are required. To match these clinical requirements, composites with tailorable properties have been designed. These include neural guides made of natural (collagen, laminin, fibrin) and synthetic (polyamide, polydioxanone, polyglactin) biomaterials which are able to release neurotrophins such as fibroblast growth factor, glial growth factor, and nerve growth factor (NGF). A typical example is the neural guide made of poly(hydroxylethyl methacrylate-co-methylmethacrylate) P(HEMA-co-MMA) hydrogels which have been embedded with growth factor carriers in the form of NGF-loaded PLGA microspheres.
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The use of composite biomaterials in this field of application is considered advantageous also because of its potential to be tuned with innovative methods to fabricate porous, biodegradable conduits presenting multiple longitudinally aligned channels available to guide axonal out-growth. For example, microstructured neural guides have been fabricated by both the coated mandrel and mandrel adhesion techniques which permit introduction into the guide of various numbers of channels of different structure, organisation and diameter. To this purpose, the composite nerve guides were obtained from the combination of PCL and porous collagen-based beads (CultiSphers) (Bender et al., 2003). The incorporation of the collagen beads improved neuronal cell adhesion and viability, while promoting neurite extension. These biocompatibility properties seem to be accompanied by suitable mechanical properties. However, it has to be outlined that, as with the clinical intervention for cartilage repair, these composite neural guides have not yet made an impact on clinical practice. Finally, it has to be considered that the tailored design of new composite neural guide may also offer a future advantage in the repair of damage to the central nervous system that has not been explored in this section.
4.5
Conclusions
The concept of regenerative medicine is underpinned by the availability of new biomaterials able to participate actively in the tissue repair process and promote its physiological pathway. To this purpose, it is a natural consequence that the sights of material scientists, industry and clinicians worldwide have been set on the development of new biomaterials able to mimic both the structure and the biochemical properties of the tissues to be repaired. In particular, in the last two decades, the research activity has been focusing on the development of biomaterials able to mimic the ECM that is the environment where cells exert their functions. From a material science viewpoint, the ECM is a composite material where different structural elements are combined to form a unique unit able to regulate tissue functions. Therefore, a large proportion of the research effort aiming at repairing soft tissues has been based on using biopolymers of proteic or polysaccharide composition which are naturally present in the ECM and which present to the cells specific bioligands able to play a role as substrate for their adhesion and to regulate their activity. These studies have been paralleled by others based on the use of synthetic biomaterials which may still be able to provide the cells with a suitable environment for growth, while offering industrial advantages such as reduced production costs and higher batch-to-batch reproducibility.
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More recently, the use of composite biomaterials comprising either natural or synthetic polymers or their hybrid combinations has been extended with the advent of nanostructured biomaterials in which the composite nature is reflected at nanoscale level to finely tune the cell microenvironment and to modulate more precisely the process of tissue repair. As a consequence, composite biomaterials have become a widespread solution to clinical cases involving damage to soft tissues. From the relatively more superficial and simple application of wound dressing to the need of tissue integration by hernia meshes and urological slings, composite biomaterials play a key role in advanced clinical treatments which are now routinely performed worldwide. Moreover, each year, composite biomaterials save the life of many patients suffering of vascular complications. The advanced technology applied to the manufacturing of vascular grafts has been explored elsewhere in this book, but still deserves a mention in this chapter about composite biomaterials for soft tissue repair. However, it is surprising to note that many clinical scenarios where the patient’s life is threatened or its quality is significantly compromised, composite biomaterials still have not developed into products. Cartilage and peripheral nerve repair have been described as examples where the promising research progress has not been matched by industrial interest. It is evident that, if framed in the context of the finely tailored physicochemical and biochemical properties of the natural ECM (see section 4.2), the full potential of composite biomaterials in soft tissue repair applications has not yet been fully explored. For example, although many studies have been performed to study the characteristics of bound water in hydrogels, very little is known about water structuring in composite biomaterials and, moreover, how this structuring can be exploited to favour controlled diffusion of growth factors and other important metabolites upon implantation. Also, biodegradable composite biomaterials rely on the spontaneous hydrolysis of the polymeric chains. The synthesis of new biomaterials able to degrade upon cell enzymatic activity has been considered in monolytic hydrogels and, to our knowledge, not in composite biomaterials. Finally, the design of composite biomaterials has so far been pursued to match the tissue mechanical properties. Little has been done to create composite biomaterials which are able to tune their physicochemical properties to those of the evolving ECM upon healing. The development of new nanostructured biomaterials may lead to a significant progress towards these goals. Indeed, the ability of achieving biomaterials with a composite nanostructure is likely to favour the creation of more controlled bulk and surface properties as well as of highly controllable microenvironments able to expose docking sites to both growth factors and cell receptors and to respond finely to cell activity. This is likely to be a significant advantages for future tissue engineering products.
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Finally, careful considerations will need to be made to assess the real clinical benefit of implanting composite biomaterials. This assessment should take into account both the significant improvement of the clinical outcome and the likelihood of the higher costs of production that these materials may bring.
4.6
References
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5 Composite materials for bone repair L. G R Ø N DA H L and K. S. JAC K, The University of Queensland, Australia
Abstract: Component selection and general design considerations are described for the production of inorganic/polymer composites for use in bone applications, with a focus on biodegradable systems. A review is presented of particulate composites (including nano-composites) with emphasis on the achievement of well-dispersed and well-bonded reinforcing phases, approaches to the in situ formation of nanocomposite materials, and the fabrication of 3D porous, composite scaffolds. Finally, some of the important challenges associated with the production of composite materials for bone repair are highlighted. Key words: biocomposites, bioactivity, mechanical properties, nanoparticles, tissue engineering, scaffolds.
5.1
Introduction
There exists a vast range of composite materials intended for use in bone applications. These differ with respect to the components chosen and to the way in which these components are assembled and processed. The focus in this chapter will be on biocomposites consisting of an inorganic dispersed phase in a polymer matrix. The design features important for biomaterials that interface with bone are similar to those developed for cartilage, tendon and ligament, but biocomposites for these applications are described elsewhere in this book. This chapter will not include organic/inorganic fibre composites, polymer–polymer composites, multiphase ceramics, materials with composite coatings or cell-material composites. The reader is referred to other chapters of this book as well as to the additional reading material listed at the end of this chapter for further details on these types of materials. This chapter starts with a discussion of component selection and general design considerations for the production of inorganic/polymer composites. Following this, fabrication of particulate composites including the pioneering work of William Bonfield on particulate hydroxyapatite (HAP)/high density polyethylene (HDPE) composite materials will be presented. In the next section, the use of nano-sized inorganic particles will be discussed and the issue of achieving a dispersed well-bonded reinforcing phase in nano-composites will be addressed. Approaches to in situ formation of 101
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nano-composite materials will then be illustrated by recent work on chitosan and collagen composites. A section on the production of composite 3D porous scaffolds will follow and the chapter will conclude with a description of key challenges and concluding remarks.
5.2
Component selection and general design considerations
In selecting the components for a composite material the key requirements that exist for any biomaterial must be fulfilled. Thus, the components must be bio-acceptable, non-immunogenic, non-toxic, and sterilizable. In addition, the cost of production, the scalability of the fabrication process and the ability to fabricate accurately sized and shaped devices are also of importance when choosing materials and processes. Finally, the polymer in the composite must have thermal properties that allows for processing and mechanical properties that approach those of the native tissue which it is to replace, since mechanical compatibility between the biocomposite and bone will encourage bone growth in and around the implant (Kohane and Langer, 2008; Bonfield, 1987). Some of the most extensively studied polymers are listed in Table 5.1 together with some of their properties and key advantages/disadvantages for use in bone applications. These polymers can be subdivided into those that are naturally occurring in nature and those that are prepared synthetically. It is generally accepted that natural polymers including polysaccharides (e.g. chitin and chitosan) and proteins (e.g. collagen) possess better bio-compatibility, but their relatively poor mechanical properties and very rapid degradation rates limit their range of applications. The poly-(β-hydroxyalkanoates), however, are naturally occurring polymers displaying good biocompatibility, suitable mechanical properties and slow degradation rates thus making them potential materials for bone applications (Chen and Wu, 2005). Synthetic polymers have the advantage that they can be tailored to give a wide range of properties and more predictable uniformity, and they are generally free from the problems of immunogenicity, which can be a concern for naturally derived materials, e.g. bacterially produced poly(hydroxybutyrate) and poly(hydroxybutyrate-cohydroxyvalerate) [PHB(V)] (Wu et al., 2006). However, it should be noted that the synthetically prepared polymers may contain low levels of chemical impurities (e.g. trace amounts of catalyst or monomer) which may lead to alternative challenges in terms of both bio- and regulatory acceptability. Additionally, the polymers listed in Table 5.1 can be categorised into biostable (i.e. those which do not degrade in vivo), and biodegradable types. Whereas the use of bio-stable polymers avoid the issue of matching the implant degradation rate with tissue regeneration, the biodegradable ones have the potential to produce an implant that with time is substituted by
E = 0.4 GPa, UTS = 10 Mpa
Bio-polymer, variable2
Bio-polymer, variable2 Bio-polymer, variable2
PCL
Collagen
Chitin Chitosan
Variable2 Variable2
Bio-degradable Degradation time (6–12 months) Bio-degradable, degradation rate faster than either homopolymer Bio-erodible, slow degradation rate (>24 months) Variable2
Bio-degradable >24 months (PLLA)
Osteoconductive; human recombinant possible Insoluble in water Soluble in acidic solution
Semicrystalline, high tissue compatibility
Tuneable degradation rate (1–12 months)
Highly crystalline, Relatively short in vivo stability
Bio-inert, ductile; allowing for high volume % of inorganic filler to be incorporated High strength but brittle, slow degradation time Less crystalline than PHB; can be modified by 3HV content, improved ductility compared with PHB, slow degradation time Semi crystalline, slow degradation time
Advantages/disadvantages
(Abraham et al., 2008, Cen et al., 2008) (Rinaudo, 2006) (Rinaudo, 2006)
(Murugan and Ramakrishna, 2005)
(Rezwan et al., 2006)
(Rezwan et al., 2006)
(Murugan and Ramakrishna, 2005, Rezwan et al., 2006)
(Murugan and Ramakrishna, 2005) (Misra et al., 2006)
(Murugan and Ramakrishna, 2005)
References
HDPE: High density polyethylene; PSU: polysulphone, PHB: poly(hydroxy butyrate), PHBV: poly(hydroxy butyrate-co- hydroxy valerate), PLA: poly(lactic acid), PGA: poly(glycolic acid), PLGA: poly(lactic-co-glycolic acid), PCL: poly(caprolactone); E: Young’s modulus; UTS: ultimate tensile strength. 2 Highly dependent on biological source and type of crosslinking used during processing.
1
E = 1.4–2.8 GPa, UTS = 40–55 MPa (depending on composition)
PLGA
PGA
PLA
(PLLA) or amorphous (PDLLA), E = 2.7 GPa (PLLA), 1.9 GPa (PDLLA), UTS = 50 MPa (PLLA), 29 MPa (PDLLA).
Bio-erodible
E = 2.5 GPa, UTS = 36 MPa, highly crystalline E = 2.5–0.5 GPa, UTS = 70–20 MPa (Depending on 3HV content)
PHB Bio-erodible
Bio-stable
E = 0.88 GPa, UTS = 35 MPa
HDPE
PHBV
Degradation rates
Mechanical properties
Polymer1
Table 5.1 A selection of polymers reported for biocomposite fabrication
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the living tissue. The class of synthetic biodegradable polymers most widely studied for bone applications are the polyesters poly(lactic acid) (PLA), poly(glycolic acid) (PGA), poly(lactic-co-glycolic acid) (PLGA) and poly(εcaprolactone) (PCL) all of which have their own advantages and disadvantages when used in bone applications (see Table 5.1). Biodegradable polymers may degrade either via a bulk- or a surface-eroding mechanism which affects the long term stability of the material and it is generally accepted that surface eroding polymers are preferred for use as scaffolds (Rezwan et al., 2006). An important consideration during the design of materials, therefore, is knowledge of both the mechanism and rate of degradation of the selected polymer. Since most of the polymers possess neither the required strength to match the mechanical properties of bone nor the required bioactivity, an inorganic component is typically introduced to improve the overall mechanical stiffness and strength, as well as the bioactivity of the resulting composite material. Therefore, the inorganic materials commonly studied for biocomposite fabrication, and which are the focus of this chapter, have relatively large moduli and strengths as well as proven bioactivities. A number of physical characteristics of the inorganic material affect their bioactivity and biostability and therefore the resulting composites. These characteristics include density, phase purity and crystallinity, as well as the overall size and shape of crystallites and/or primary particles (Vallet-Regi, 2001). In Table 5.2 the most commonly used inorganic materials for bone applications are listed together with their key advantages for use in bone biocomposites. It is worth noting that while the calcium phosphates used in composites for bone repair are of the biodegradable type, the micrometre-sized bioglasses and glass ceramics are biostable. Glasses and glass ceramics with a range of specific compositions from the Na2O–CaO–P2O5–SiO2 system have the capacity to bond to bone and, hence, are termed bioactive. The glasses are amorphous single-phase materials and promote the rapid bonding of bone to the surface of the implant whereas the bioactive glass ceramics (BGC) are multi-phase materials, with relatively good strength and toughness. However, the bonding of the native bone to the surface does not occur as rapidly as for the glasses (Vallet-Regi, 2001). The calcium phosphate minerals such as HAP and β-tricalcium phosphate (β-TCP) are obvious choices as resorbable ceramics in bone repair since HAP dominates the mineral phase in bone. Fundamental studies and clinical applications have demonstrated that calcium phosphate biomaterials are biocompatible and osteoconductive. (Vallet-Regi, 2001; Bohner, 2000; Langstaff et al., 1999). When implanted in vivo, these materials are non-toxic and bond directly to bone without any intervening connective tissue layer. In some composite materials the inorganic component can also provide a buffering capacity when used in conjunction with biodegradable polymers that produce acidic degradation
Composite materials for bone repair
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Table 5.2 Inorganic materials reported for biocomposite fabrication (Adapted from Rezwan et al., 2006; Vallet–Regi, 2001) Material
Key properties
Key advantages
HAP
Osteoconductive, bioactive, E = 80–110 GPa, UTS = 100 GPa
β-TCP
Bioactive, bio-resorbable
Bioglass® (BG)
Bioactive, binds to both hard and soft tissue, amorphous, non-resorbable, E = 500 MPa, UTS = 35 GPa Bioactive, amorphous with crystalline nuclei, non-resorbable, E = 215 MPa, UTS = 118 GPa
Easily processed into nano-sized crystals, small crystals (<10 μm) are bio-resorbable Faster degradation rate than HAP More bioactive than HAP
A-W glass ceramic
More bioactive than HAP
products (e.g. PLGA) thereby avoiding locally increased acidity and subsequent adverse immune response. Whereas it has long been known how to produce and stabilised HAP on the nano-scale (i.e. with at least one dimension <100 nm), only few recent reports on the fabrication of nanosized bioglasses exist (Brunner et al., 2006). It is evident that there exist a number of component choices when designing a biocomposite and it clearly depends on the exact intended application (i.e. load-bearing or not, bio-stable or degradable) as to the selections of components. One common goal in biocomposite fabrication is to achieve reinforcement of a relatively low modulus polymer with a high-modulus, high-strength reinforcing inorganic material utilising the plastic flow of the polymeric material under stress to transfer the load to the reinforcing phase. This can result in a composite with a higher strength and modulus than the polymer and better toughness and processability than the pure inorganic material. The ability to achieve this goal is highly dependent on the relative volume fraction of each of the phases that can be obtained whilst ensuring that the inorganic phase is well-dispersed in (Fig. 5.1A) and well-bonded to the polymer matrix. The interfacial bond strength thus has to be sufficient for the load to be transferred from the matrix to the reinforcing phase if the composite is to be stronger than the polymeric matrix material. In addition to the importance of creating a large interfacial bonding strength, it is also greatly beneficial to have a large interfacial area between the phases. In general, a large interfacial area provides greater transfer of load from the matrix to the reinforcing phase which leads to an improvement in mechanical properties. Moreover, this area is dependent on the shape and size of the reinforcing particle and hence to optimise this aspect of
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Biomedical composites (a)
(b)
(c)
(d)
A
B
Surface
Bulk
5.1 Composite material where grey background illustrates the polymer matrix and the light grey lines signify reinforcing crystals: A top view (a) poorly dispersed crystals in the polymer matrix; (b) well dispersed crystals and thus enhanced mechanical properties. B Side view. (c) poor surface presentation; (d) good surface presentation and thus improved bioactivity.
composite design much effort has involved the inclusion of nano-sized particles. Finally, in addition to improving the strength of the composite material, high bioactivity is necessary and this can be achieved through a high surface presentation of the bioactive inorganic component (Fig. 5.1B). Unfortunately, this is an aspect that is often overlooked as will become evident in this chapter.
5.3
Fabrication of particulate composites
The use of a bioactive inorganic filler in composite materials for bone biomaterials was pioneered by Bonfield et al. in the early 1980s (Bonfield et al., 1981) with the production of particulate HAP/HDPE composites. Owing to the ductile nature of HDPE, composites with up to 45 vol% (73 wt%), HAP could be produced using the fabrication process of compounding, powdering and compression moulding (Wang et al., 1994). Uniform distribution of the HAP particles was observed on the surface of high vol% HAP composites and the HAP particle size was unchanged from that of the HAP particles in the powder used. The lack of agglomeration of the HAP particles was attributed to the large shear forces generated during the compounding process (Wang and Bonfield, 1996). The Young’s modulus and tensile strength of the composite materials increased with
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increasing filler content reaching values of 5.5 GPa and 19 MPa, respectively, for a 45 vol% HAP/HDPE composite (Wang et al., 1998). In vitro osteoblast studies revealed attachment of the cells to the HAP islands followed by spreading across the implant surface (Huang et al., 1997). In vivo studies showed an excellent bone–biomaterial interface attributed to both the bioactive nature of the implant and the mechanical matching of the implant with bone (Bonfield and Luklinske, 1991). Implants produced from HAP/HDPE are now used as middle ear implants, under the trade name HAPEXTM (Swain et al., 1999). Another example of using a bioactive filler for reinforcement of a biostable polymer is the Bioglass®/HDPE composites. The inclusion of Bioglass® rather than HAP in HDPE composites was done to improve the overall bioactivity. Using similar processing methods as for the HAP/HDPE composites it was possible to include up to 50 vol% Bioglass®. It was reported that composite materials with greater than 30 vol% Bioglass® possessed mechanical properties similar to cancellous bone and similar in magnitude to the HAPEXTM analogue (Juhasz et al., 2004). Simulated body fluid (SBF) studies showed that the Bioglass® composites have higher bioactivity than the HAP composites, as expected (Rea et al., 2004). Additionally, composites of HAP with biostable polymers such as poly(1,4phenylene ether ether sulfone) (PSU) (Wang, 2003; Chlopek et al., 2006), polypropylene (Liu and Wang, 2007), high-impact polystyrene (HIPS) (Gong et al., 2004), and poly(ether ether ketone) (Abu Baker et al., 2003) have also been reported. The polymers are chosen for investigation because of their superior mechanical properties (stiffness, strength or toughness) compared with HDPE, and in the case of PSU its stability with respect to oxidation, hydrolysis and radiation, which are typical processes used in the sterilisation of medical devices. It is reported that the incorporation of HAP leads to increases in the mechanical properties of the resulting composites, provided that sufficient dispersion of filler particles can be achieved. Indeed, in the work of Gong et al. (2004), the authors have shown that the degree of dispersion and the interfacial adhesion of micrometre-sized HAP particles in a HIPS matrix could be enhanced through an in situ polymerisation of styrene at the surface of the HAP particles. Furthermore, Chlopek et al. (2006) have investigated the temporal variation of mechanical properties of a 15 wt% HAP/PSU composite under in vitro conditions via the measurement of a series of creep curves at various applied stresses. They found that the HAP/PSU composite showed long-term stability only when the applied stress was ⱳ20% of the initial (static) strength of the material, suggesting stringent limitations on the applicable sites of implantation of these materials. These findings are also more general in that they highlight the need for the measurement of dynamic mechanical properties in vitro, as opposed to static measurements, when assessing the performance of
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biomaterials. When comparing the overall performance of the biostable composites, however, there appears to be no definitive conclusions in the literature as to the general superiority of any of these composites over the others, including HPDE. The use of biodegradable polymers in the fabrication of particulate reinforced composites has predominantly involved the use of the polyhydroxyalkanoates, e.g. PHB and PHBV (Chen and Wu, 2005) as well as synthetic polyesters. However, it is not within the scope of this chapter to comprehensively cover particulate reinforced composites of both classes of polymers in detail and the reader is referred to the reading list in section 5.8. The sections on nano-composites and composite scaffolds will, however, include studies on a much broader range of polymers (sections 5.4 and 5.6). HAP particle reinforcement of PHB and PHBV produced by dry mixing and melt processing resulted in composites displaying compressive strength of 62 MPa (30 wt% HAP) (Galego et al., 2000). In vitro testing in SBF demonstrated high bioactivity as evident from formation of a bone-like apatite layer on the material surface with the bioactivity increasing with filler content (10–30 vol% HAP) (Ni and Wang, 2002). Coskun et al. (2005) studied the inclusion of HAP rods (2–4 × 20–30 μm) as filler in PHB and PHBV composites produced by blending and injection moulding. SEM images indicated the HAP rods were highly embedded and homogeneously distributed within the polymer matrx. However, based on the mechanical properties, it was concluded that the composite did not possess a strongly bonded interface. Indeed, the Young’s modulus decreased with inclusion of filler (15 wt%) whereas the tensile strength increased for PHB but decreased for PHBV composites. In vivo implantation into bone tissue of particulate-reinforced HAP/PHBV or β-TCP/PHBV composites have shown good bone integration and formation of a layer of bone-like apatite at the implant interface after implantation (Luklinska and Bonfield, 1997; Luklinska and Schluckwerder, 2003). Notably, none of these studies compared the bioactivity with that of the PHBV polymer alone. However, a separate in vivo study on HAP/PHB particulate composites failed to detect improved bioactivity when compared with PHB (Doyle et al., 1991) and this was attributed to the lack of surface presentation of HAP. Composite Bioglass®/PHB films (5 and 20 wt%) were produced by solvent casting from a chloroform suspension of 5 μm Bioglass® particles which were dispersed by sonication (Misra et al., 2007). The presence of Bioglass® reduced the crystallinity of the PHB polymer. In addition, it was proposed that poor wetting of the particles caused poor interfacial bonding which, in turn, resulted in lowered Young’s modulus for the composites compared with the pure polymer. Although the presence of Bioglass® particles on the surface was only evaluated using microscopy, it was concluded that there were few particles on the surface and that the particles were
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present mainly in the bulk of the film. This explained the relatively low bioactivity as evaluated in vitro by immersion in SBF. However, owing to the lack of a comparative study on a pure PHB control being included in their work, it is not clear if the filler improved the bioactivity to any extent.
5.4
Fabrication of nano-composites
In micrometre composite fabrication, agglomeration of particles can to a large extent be avoided by using blending, solvent processing, extrusion and/or sonication. However, these methods are generally insufficient when working with nano-sized crystals. Indeed, one of the main challenges in nanotechnology is to achieve good dispersion of the individual nano-sized (primary) particles which owing to their high surface area have a strong tendency to agglomerate. There are a few examples of using the concept of dispersion of the nanosized HAP (nHAP) particles in an organic solvent which can also be used to dissolve the polymer of choice. An important consideration here is the colloidal stability of the nHAP particles in the solvent which have led to the choice of DMF (Hao et al., 2002; Deng et al., 2001) which offers relatively good stability of nHAP particles (Rai et al., 2008a). Co-precipitation of nHAP particles and a polymer from a combined suspension allows production of a nano-composite powder which can be further processed. For such composites of PCL or PLA an increase in tensile modulus was observed with increasing nHAP content, but no change in the tensile strength was found (Hao et al., 2002; Deng et al., 2001). Since the particles in these composites appeared to be well dispersed when studied by microscopy, the authors suggested that the observed properties can be attributed to poor interfacial bonding between the hydrophilic nHAP particles and the hydrophobic polymers. However, recent work by Schaefer and Justice (2007) has shown using ultra-small-angle scattering techniques that in many apparently ‘well-dispersed’ materials determined by microscopy, there exists micrometre-scale agglomeration of the nano-sized particles which, in turn, reduces the overall efficiency of the filler particles to provide reinforcement of the composite. In the production of nHAP/polymer composites, a number of different approaches which involve innovative chemistry, as illustrated in Fig. 5.2, have been employed not only to create well-dispersed particles but also to enhance the interfacial bonding between the particles and the matrix polymer. These approaches involve surface modification of the nHAP particles before composite fabrication and include adsorption of surfactants or polyelectrolytes as well as surface grafting techniques. The successful use of surfactant adsorption to nHAP was demonstrated by Kim (2007) who used oleic acid to both increase the colloidal stability of nHAP in a chloroform
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5.2 A schematic representation of the surface modification of HAP crystals by (a) surfactant adsorption, (b) polyelectrolyte adsorption, (c) chemical grafting (not drawn to scale).
suspension and to mediate the interaction between the hydrophilic nHAP crystals and the hydrophobic PCL matrix. An increase in both tensile stress and elastic modulus was observed for composites prepared using oleic acid compared with those prepared using non-stabilised nHAP. The latter composites displayed an increase in the modulus but not strength, similar to that observed for the DMF dispersed systems. Kim (2007) also investigated the biological response and found that the cell proliferation was enhanced on the nano-composite produced using oleic acid. However, it could not be concluded whether this was due to the different surface presentation of nHAP, differences in surface roughness of the films or to the nano-size of the reinforcing particles. Adsorption of polyelectrolytes such as poly(acrylic acid) (PAA) and heparin to nHAP particles has been well studied (Misra 1996; Rees et al., 2002) and the resulting enhancement in colloidal stability well documented (Rai et al., 2008b; Noohom et al., 2009; Misra 1996). The use of PAA-coated nHAP in the preparation of PHBV composite powders which were subsequently processed into discs via compression moulding or films via solvent casting (Fig. 5.3) showed an enhanced surface presentation and a concomitant improved cell response compared with composites produced using micrometre-sized HAP particles (Cool et al., 2007). In addition, the mechanical properties were found to be enhanced for particle loadings of up to ca. 15 wt%. At higher loadings the composites were found to contain significant large-scale agglomeration of particles and they could not be processed into structures suitable for mechanical testing (Noohom et al., 2009). It was also shown that the level of mechanical improvements observed was well below the theoretical values predicted and this is most likely the result of the presence of micrometre-scale agglomeration as discussed by Schaeffer and Justice (2007). The choice of heparin for stabilising nHAP particles (Rai et al., 2008a) was based on its dual effect in enhancing colloidal stability and
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5.3 10% w/w HAP/PHBV solvent cast films produced using PAA stabilised HAP nano crystals. TEM images (a) HAP clusters are visibly scattered throughout the composite (indicated by arrow) but (b) at higher magnification, individual HAP particles with spindle-like morphology are visible. SEM images (c) solvent cast PHBV film and (d) 10% HAP/PHBV composite. Micron-sized HAP agglomerates are visible on the surface of the HAP/PHBV composite films. The dimensions of the scale bars are: (a) 5 μm, (b) 200 nm and (c) and (d) 10 μm (images supplied by Dr C K A Wu, The University of Queensland).
providing biological stimulus, i.e. it was used as a model for bone-specific heparin sulfates (Nurcombe et al., 2004). Solvent cast films produced from co-precipitated composite powder displayed enhanced wettability and thus nHAP surface presentation (yielding an advancing contact angle of 35° for 30 wt% composite) compared with an insignificant change for composites prepared using non-stabilised particles (Rai et al., 2008a). In addition, the mechanical properties were much improved for the heparin-stabilised composite films yielding a four-fold increase in both elastic modulus and tensile strength for the 30 w/w% composite. Again, non-stabilised composites displayed similar mechanical properties to the pure polymer. The enhancements seen were to a large extent attributed to the heparin-stabilised nHAP particles being well dispersed in the polymer matrix with minimal agglomeration observed by electron microscopy.
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Surface grafting of polymers onto nHAP particles has been achieved by predominantly two methods, e.g. plasma polymerisation (Nichols et al., 2007) and chemical reaction to the surface hydroxyl groups (Lee et al., 2006; Hong et al., 2004; Lui et al., 1998; Qiu et al., 2005; Liu et al., 2008a). In the first method, nHAP particles were modified by acrylic acid plasma polymerisation to introduce tethered PAA chains. By manipulating the plasma parameters (i.e. power and gas pressure) two types of coatings resulted: one being degradable and possessing a high percentage of carboxylic acid groups which yielded a highly hydrophilic coat (advancing contact angle of 15°), the other which was stable and less hydrophilic (advancing contact angle of 57°). Both types of particles were introduced into solvent cast PLGA films (3 wt%) with mixing achieved using ultrasonication. It was found that both types of coatings improved the modulus and yield strength of the composite films (compared with non-coated nHAP composites) and, in addition, the use of particles with degradable coatings yielded significantly better results (Nichols et al., 2007). Chemical reactions with the nHAP surface hydroxyl groups have been achieved using the two chemical approaches of grafting to and grafting from. In the former, a pre-formed polymer chain is attached to surface hydroxyl groups, in the latter grafting is initiated from the surface hydroxyl groups. The grafting from approach in this context was pioneered by Lui et al. (1998) with the fabrication of PEG-grafted nHAP particles. More recently, Qiu et al. (2005) used the grafting to approach to attach PLLA via a condensation reaction between the oligomer carboxylic acid end group and the nHAP surface hydroxyl groups. These particles formed a stable dispersion in chloroform and were subsequently used in composite formation (5–30 wt% filler) using mechanical mixing. It was found that the tensile modulus increased with increasing filler % over the range studied, while the tensile strength increased with the incorporation of up to 15% filler and then decreased to reach values of the pure polymer at 30% filler. Both modulus and strength were vastly improved compared with non-grafted nHAP composites and this was attributed to the improved distribution of the particles within the polymer matrix on the micrometre scale. This was elegantly supported using EDX to measure Ca elemental distribution maps of the fracture surfaces. The grafting from approach for the surface modification of nHAP has been more extensively used. Some of the advantages of this technique are that it offers higher graft density and an easier rate to purification as compared with the grafting to method. Hong et al. (2004) grafted PLA onto nHAP and subsequently introduced these particles in PLLA composites (1–30 wt%). Increased particle dispersion as compared with non-grafted nHAP was observed by TEM and, although better tensile strength was found for composites made from grafted nHAP, no difference in tensile modulus between stabilised and non-stabilised particle composites were
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observed. Improved grafting from methods were developed by Lee et al. (2006) in which linker molecules (e.g. lactic acid and polyethylene glycol) were introduced before grafting. This was done to decrease steric effects and increase the nucleophilicity of the hydroxyl groups and it allowed for more efficient grafting (especially for the PEG-modified nHAP particles). In this study, grafting from PCL by ring opening polymerisation was performed. nHAP/PCL composites with 10 wt% filler showed improved tensile modulus as well as tensile strength when using grafted nHAP as compared with the non-grafted system. A subsequent study evaluated the in vitro biocompatibility of the composites using protein adsorption and fibroblasts cell culture (Lee et al., 2007). Improved biological activity was observed for composites produced using grafted nHAP compared with the nongrafted analogue, but the concentration of nHAP particles on the surface was not evaluated and, therefore, it is difficult to say what caused this improvement. The recently developed ability to fabricate Bioglass® on the nano-scale has resulted in studies using these nano-fillers in biocomposites for bone tissue engineering applications. Misra et al. (2008) compared the inclusion of micro- and nano-sized Bioglass® in PHB composite films prepared by solvent casting. It was found that the micrometre reinforced composites (10–30 wt% filler) displayed a reduction in Young’s modulus compared with the pure polymer, while the introduction of the nano-sized filler improved the Young’s modulus. However, the improvement was greatest for the low filler content (10 wt%) and decreased with increasing amount of filler (up to 30 wt%). The nano-sized filler was found to be present at the material surface (assessed by SEM and contact angle measurement) and this resulted in higher protein adsorption and bioactivity (assessed by SBF) compared with pure polymer and composites with micrometre-sized filler. In a study by Liu et al. (2008a), the grafting to approach was used to modify nano-sized Bioglass® particles with PLLA. These particles were subsequently introduced into PLLA discs prepared by hot pressing. Increase in tensile modulus with filler content (4–20 wt%) was observed for both grafted and non-grafted Bioglass® composites, but only grafted Bioglass® composites showed an improvement in tensile strength and only up to a filler content of 6 wt%. As can be seen from the examples given in this section, the use of intelligent manipulation of the nano-sized particles can lead to improvements in mechanical properties of nano-composite materials. However, only the approach of heparin adsorption was successful in improving both strength and modulus of the composite up to high (30 w/w%) filler content. Most of the remaining strategies were effective to a filler content of 15% or less. There is, however, no clear evidence that it is owing to their nano-scale size that these reinforcing particles improve the biological response. In all cases, it can indeed equally well be attributed to secondary effects such as either
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surface presentation of the filler or processing effects such as altered surface roughness. These secondary effects are brought about by the nano-size of the filler particles, and thus it can be concluded that the use of nano-sized particles is generally highly beneficial.
5.5
Template-mediated formation of nano-composites
An attractive method of nano-composite fabrication is the use of biopolymers (e.g. collagen or chitosan) to act as nucleation sites for HAP growth and thus to use the so-called bottom-up approach. This can be achieved either by immersing a biopolymer construct into a solution containing calcium and phosphate ions such as simulated body fluid, or it can be achieved by allowing HAP nucleation and growth on dissolved biopolymer chains or fibrils (Murugan and Ramakrishna, 2005). Both methods allow for nHAP particles to form at discrete locations which to some extent circumvent the issue of nHAP agglomeration. In the following text, examples of the latter method will be described in further detail. Several groups have successfully grown nHAP crystals in solutions containing macromolecules and/or polymers aiming to mimic the ultrastructure of bone. It was thus shown that the growth of nHAP crystals takes on directional aggregation similar to that found in bone in the presence of chondroitin sulfate (Jiang et al., 2005; Rhee and Tanaka, 2000; Rhee and Tanaka, 2001). The biopolymer chitosan has also been used to produce nHAP composites. In one such study a chitosan/phosphoric acid solution was added to a Ca(OH)2 suspension (Yamaguchi et al., 2001) into which the chitosan would precipitate and simultaneously nucleate and grow nHAP. Composites which showed discrete nHAP particles (30 mm × 10 nm) by TEM yielded increasing Young’s modulus values with increasing nHAP filler content, however, the strain at failure decreased with filler content. An alternative method was developed by Rusu et al. (2005) in which a more controlled stepwise co-precipitation approach was used. This allowed for size control of the nHAP particles formed. Recently, Kithvar et al. (2009) used an in situ precipitation method to prepare composite suspensions containing nHAP and formaldehyde-treated chitosan. The method involved transforming an amorphous calcium phosphate phase into nHAP (30–35 nm) in a chitosan solution. Solvent cast films of the nano-composites showed significant increases in both the Young’s modulus and ultimate tensile strength with increased weight fraction of nHAP reaching values of 17.0 GPa and 208 MPa, respectively, for films containing 70 wt% nHAP (Fig. 5.4). In an attempt to closely mimic the structure of natural bone, nHAP/collagen composites have been studied extensively and only a small sample of
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5.4 SEM image of a 70% (w/w) HAP/chitosan composite produced by co-precipitation (image provided by Dr P Kithva, The University of Queensland).
recent work will be included here. Roveri et al. (2003) performed direct nucleation of nHAP crystals onto dissolved collagen fibers and achieved composites with nHAP crystals aligned with their c-axis along the collagen fibres. Another study by Kikuchi et al. (2001) used a similar approach and obtained self-assembled fibre bundles which were 20 μm in length and composed of collagen fibrils 300 nm in length which had nano-sized plate shaped crystals of 50 nm incorporated. This composite was tested for its mechanical properties using three-point bending and found to have a Young’s modulus of 2.5 GPa when prepared under optimum conditions of pH 9 and 40 °C. In vivo performance was investigated in a 20 mm bone defect of a beagle tibia using an Ilizarow frame stabilised fracture model. It was found that bone remodelled within 8 weeks and it was proposed that this composite is a suitable replacement for autologous bone grafts. This type of composite has the advantages of being produced from two osteoconductive constituents giving rise to excellent biological performance. However, in general the mechanical properties are not appropriate for load bearing in vivo applications. Various methods have been used to cross-link the collagen biomaterials to improve their mechanical properties and prolong their degradation time, including physical treatments (e.g. dehydration or UV irradiation) and chemical treatments (e.g. carbodiimide or glutaraldehyde) (Wahl and Czernuszka, 2006). Owing to the potential toxicity of chemically cross-linked biomaterials, physical cross-linking is preferred. However, the composites prepared still do not possess the required
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mechanical integrity and this currently limits the use of nHAP/collagen composites in load-bearing bone-graft applications.
5.6
Composite scaffolds
The fabrication of three dimensional (3D) porous structures with controlled, tunable, architectures remains an important challenge in the field of tissue engineering and has been the subject of many reviews; see e.g. (Hutmacher et al., 2007; Kohane and Langer, 2008; Langer, 1995; Rezwan et al., 2006; Yunos et al., 2008). In general, such scaffolds are engineered to enhance the rate of tissue repair by providing a large surface area for promoting vascularisation and by integration of the regenerating tissue within the scaffold. This is often coupled with the delivery of, for example, stem cells or growth factors. Furthermore, scaffolds for tissue engineering applications are designed such that they will be fully degraded once the native tissue has reformed. For applications in bone repair, a number of key parameters, in addition to the general requirements for bulk composite materials discussed in section 5.2, need to be addressed when choosing both the materials for and the method of fabrication of the scaffolds. As discussed below, the resulting scaffolds should have optimal porosity as well as an optimal distribution of pore sizes and shapes and their mechanical properties, and degradation rates need to be considered. The ability to form dense 3D networks of interconnected pores of a suitable size to allow for vascularisation and the penetration of osteogenic cells is a key requirement. It has been proposed that pore sizes of >300 μm and maximal porosities of up to 90% are best suited for successful regeneration of bone (Karageorgiou and Kaplan, 2005). Recently, however, Linnes et al. (2007) have proposed that an optimal pore size in order to avoid a prolonged pro-inflammatory and to increase angiogenesis in vivo is ca. 35 μm. Moreover, it has been suggested that this optimal pore size may be general to a range of types of tissue scaffold, including hard tissues. However, the findings of the in vivo investigations of scaffolds with well defined pore sizes tested in a bone fracture site are yet to be published. In addition, the desire for high porosities must be tempered by the need to maintain suitable mechanical properties within the scaffold throughout its lifetime (Karageorgiou and Kaplan, 2005; Bonfield, 2006). The maximum porosity that can be used is, therefore, highly dependent on the mechanical properties of materials used, their degradation properties and their intended implantation site. Scaffolds with porosities of 70–95% are typically reported. The development of mechanical properties which are sufficient to be at least self-supporting over the intended lifetime of a scaffold and in some cases matching those of native bone, e.g. in the case of load-bearing implants,
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is another key challenge. This requirement is more restrictive on the choice of materials when compared with the bulk materials as the stiffness and strength are greatly reduced owing to porosity and in addition must be maintained in vivo (i.e. in wet conditions). Moreover, for these degradable scaffolds the rates of degradation of the scaffold must be balanced with the rates of tissue regeneration (and any decline in mechanical properties) to maintain the integrity of the scaffold throughout its lifetime. In addition, the mechanism of degradation of the polymer (i.e., surface or bulk eroding) may be dependent on the thickness of the material. Hence, a polymer that degrades by a surface erosion mechanism in a thick specimen may degrade via bulk erosion when fabricated as a thin specimen (Burkersroda et al., 2002). Given the typically large surface areas and thin walls engineered into tissue scaffolds, the issue of degradation rate and mechanism need to be determined for the scaffolds and cannot be simply extrapolated from the behaviour of bulk polymers. There exists a range of methods for the fabrication of porous polymeric and composite scaffolds which include solvent casting and salt/ particulate leaching (Hou et al., 2003), microsphere templating (Linnes et al., 2007), photo-lithography (Bryant et al., 2007), sintering (Gross and Rogriquez-Lorenzo, 2004), thermally induced phase separation (TIPS) (Cao et al., 2006), rapid prototyping (Russias et al., 2007), and solid free form fabrication (Hutmacher and Cool, 2007). It is beyond the scope of this chapter to describe in detail the methods and their relative merits; instead the reader is directed toward currently available reviews e.g. (Rezwan et al., 2006). It should be noted here, however, that, although all methods allow for some control over the pore size, scaffolds prepared by salt/particulate leaching always have high porosity which greatly affects their mechanical properties, whereas, for example, solid free-form fabrication allows for much lower porosity which produces scaffolds that are, in general, more suitable for bone tissue engineering. Moreover, techniques such as rapid prototyping and microsphere templating allow for the detailed control of the distribution of pore sizes and shapes and can therefore be used to investigate the effect of these parameters on the scaffold performance. The choices of component materials for composite scaffolds are primarily governed by the same design considerations discussed in section 5.2 and, hence, the vast majority of composites scaffolds reported for bone repair contain HAP, β-TCP or bioactive glasses as the reinforcing phase. Composites of these inorganic particles with the biodegradable polyesters (PCL, PLA, PLGA) have been most extensively investigated (Charles-Harris et al., 2008; Hong et al., 2008; Hutmacher et al., 2007; Kim et al., 2006a and 2006b; Lei et al., 2007; Ma et al., 2001; Minamiguchi et al., 2008; Russias et al., 2007; Verma, 2006; Wei and Ma, 2004). In addition, the fabrication
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and characterisation (physical, in vitro and in vivo) of composite scaffolds of PHAs have also received much attention in the literature (Jack et al., 2009; Li and Chang, 2004 and 2005; Sultana and Wang, 2008; Wang, 2006; Wang et al., 2004 and 2005). Composites of calcium phosphates with natural polymers such as collagen, polysaccharides and silk fibres have also been reported (Du et al., 1998; Jiang et al., 2008; Liu et al., 2008b). In the vast majority of these works the authors have demonstrated that the incorporation of inorganic particles within the polymer matrix leads to significant reinforcement of the composite scaffold. For example Wei and Ma (2004) have shown that the incorporation of HAP particles led to an increase in the compressive modulus from ca. 4 to 8 MPa when 30–50 wt% of HAP was added to a PLLA scaffold fabricated by TIPS. However, only a small difference in the modulus of the composites was shown when comparing the incorporation of 50 wt% nano-sized HAP with micro-sized HAP. In recent work we demonstrated that the addition of as little as 2 wt% of nano-sized HAP particles to a porous scaffold of PHBV produced by the TIPS method resulted in a significant increase in both the compressive modulus (1.8 to 4.5 MPa) and strength (0.8 to 2.5 MPa) (Jack et al., 2009). A three-fold increase in compressive modulus upon addition of 20 wt% of nano-sized HAP to PHBV TIPS scaffolds was also reported by Sultana and Wang (2007). In addition, in the work of Jack et al. (2009) the compressive modulus continued to show an increase to ca. 10 MPa after immersion of the composite scaffold in SBF for 2 weeks whereas no increase was observed in a scaffold of pure PHBV during this time (Fig. 5.5). The addition of BCG to PLLA to produce composite scaffolds has also been investigated. For example, Hong et al. (2008) have fabricated scaffolds with a range of compositions of BCG/PLLA (up to 30 wt% of BCG) by TIPS. They have shown that the addition of up to 30 wt% of BCG led to an increase in the compressive modulus (from ca. 5.5 to 8 MPa) and strength (from ca 0.28 to 0.33 MPa), for up to 10 wt% of BCG at least. Moreover, in all of these works, the addition of the reinforcing particles was found to have little
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effect on the degree of porosity of the scaffolds, which were in the range of 88 to 95%. In general, for the highly porous biodegradable composite scaffolds reviewed in this work, the maximum obtainable values for the elastic moduli were of the order of a few to a few tens of MPa. One exception is the work of Russias et al. (2007) in which a highly ordered scaffolds (75% porosity) containing 70 wt% HAP in PCL and PLA were prepared by a printing method and the resulting Young’s modulus was reported to be as high as 150 MPa when measured parallel to the printing plane. In general, these values are less than that of typical cancellous bone (50–500 MPa) and orders of magnitude less than that of cortical bone (7–30 GPa), but are large enough to provided suitable mechanical integrity to the scaffolds to allow them to be self-supporting. Arguably of greater importance given that the current values of mechanical properties reported are significantly less than that observed in bone, is that the incorporation of bioactive fillers has been shown to provide significant advantages to the bioactivity of the scaffolds both in vivo and in vitro. HAP and BCG particles have extensively been shown to enhance the rate of mineralisation onto the surface of the scaffold from SBF (and in vivo) when incorporated into a range of polymeric matrices which potentially provides materials with a higher degree of osteogenic potential in addition to enhanced mechanical properties (Hong et al., 2008; Jack et al., 2009; Kim et al., 2006a and 2006b; Wei and Ma, 2004). It is also widely reported that the incorporation of bioactive fillers results in materials that show higher levels of cell attachment, differentiation, proliferation and penetration into the scaffold (Charles-Harris et al., 2008; Kim 2006a and 2006b; Ma et al., 2001; Minamiguchi et al., 2008; Wang et al., 2005) and to the enhanced formation of bone in vivo. In addition, it has also been demonstrated that the presentation of the HAP at the surface of the scaffold is of key importance to the full realisation of these improvements (Jack et al., 2009; Wang et al., 2005). Finally, the incorporation of biodegradable filler particles can provide a method for modifying the rates of degradation of the scaffold and can also act as a buffer against undesirable degradation products (Li and Chang, 2004 and 2005). Although the question of efficacy of nano- versus micro-sized particles has also been discussed for bone scaffolds, most typically when comparing HAP-filled polymers, see e.g. (Wei and Ma, 2004; Kim and Fisher, 2007), there appears to be an absence of systematic studies of the effect of the filler particle size on the mechanical properties or bioactivity. Although it is difficult to draw definitive conclusions with regard to the question of the enhancement of mechanical properties, in general, it appears that any extra reinforcing effects of nano-sized particles may be most significant at lower particle concentrations compared with higher loadings in agreement with those observed in the bulk composite
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materials. In these bulk composites it has been shown that the stabilisation of primary particles (i.e. prevention of agglomeration), particularly at higher particle content, is of vital importance to maintaining effective reinforcement (Nohoom et al., 2009; Rai et al., 2008a; Schaefer and Justice, 2007). Interestingly, it is noted that although there has been much research into methods, characterisation and understanding of the stabilisation of nanoparticles in bulk composites, there appears to be few reports of these phenomena in porous composite scaffolds. Such investigations will no doubt become more abundant in the future. In the case of bioactivity, however, it is generally accepted that the increased surface area to mass ratio of the nanoparticles in combination with any higher degree of surface presentation results in higher levels of bioactivity, e.g. protein adsorption, in vitro mineralisation and cell proliferation and differentiation (Kim and Fisher, 2007; Kim et al., 2006a and 2006b; Sato and Webster, 2004; Wei and Ma, 2004).
5.7
Key challenges and concluding remarks
There remain a number of key challenges in the fabrication of biodegradable and bioactive composites for bone regeneration that have not yet been achieved. One of the major challenges is the ability to control and tune the rate of degradation of the materials. That is the composites must be designed to degrade gradually in vivo while being replaced by newly formed tissue, thereby providing a means of load transfer from the implant to the new bone during degradation. Moreover, since the appropriate rate of degradation is not only determined by the implant (i.e. components, fabrication method, overall size and porosity) but also on the disease state of the patient, clearly the details of the degradation process of the implant must be able to be accurately predicted and also be able to be tailored for a given circumstance. Clearly an important challenge in the fabrication of nano-composites is the ability to obtain as completed as possible dispersion of the primary particles and to maintain this dispersion throughout the lifetime of the composite, if the full benefits of the nano-scale reinforcing particles are to be realised. More knowledge and understanding of methods for achieving these higher degrees of dispersion are still required even in bulk materials. Moreover, only limited attempts have been made to incorporate nano-sized particles into 3D porous scaffolds and, thus far, none of the more sophisticated stabilising agents used in production of solid films or discs have been transferred into 3D scaffolds. A greater understanding of the interplay between the methods of stabilisation and scaffold fabrication will need to be developed and the effects of stabilisation methods on the in vitro and in vivo response of the scaffold will need to be investigated. Finally, in
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order to produce better osteoconductive and osteoinductive 3D scaffolds a greater understanding and new materials which are fabricated by template-mediated formation of nano-composites may prove crucial to the achievement of these goals.
5.8
Sources of further information and advice
As pointed out throughout this chapter, not all aspects of biocomposites for bone regeneration could be extensively covered. The reader is therefore referred to the following key review papers and books for further details. Murugan, B; Ramakrishna, S (2007) ‘Nanoengineering biomimetic bonebuilding blocks’ in Topics in applied physics: molecular building blocks for nanotechnology – from diamonoids to nanoscale materials and applications, Ed. G A Mansoori, F G Thomas, A Lahsen, G Zhang. Springer. Seal, B L; Otero, T C; Panitch, A (2001) ‘Polymeric biomaterials for tissue and organ regeneration’ Materials Science and Engineering R, 34, 147–230. Burg, K J L; Porter, S; Kellam, J F (2000) ‘Biomaterial developments for bone tissue engineering’ Biomaterials, 21, 2347–2359. Rahaman, M N; Mao, J J (2005) ‘Stem cell-based composite tissue constructs for regenerative medicine’ Biotechnology and Bioengineering, 91, 261–284.
5.9
References
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6 Composite coatings for implants and tissue engineering scaffolds M. WA N G, The University of Hong Kong, Hong Kong
Abstract: Most currently used implant materials in orthopaedics and dentistry lack osteoconductivity. Over the past two decades, various surface modification techniques have been investigated and developed to improve the osteoconductivity of metallic, polymeric and ceramic biomaterials. To overcome various problems and to fulfil different purposes, bioactive composite coatings are formed on metallic biomaterials using some of these techniques. The composite coatings thus produced have been shown to possess better mechanical and/or biological properties than single-component, monolayer coatings. Since it emerged two decades ago, tissue engineering holds great promises for the regeneration of human body tissues for the diseased or damaged body parts. Some particular coating techniques such as biomimetic deposition can be employed to form bioactive composite coatings on biodegradable polymer tissue engineering scaffolds so that the scaffolds can provide an enhanced substrate for cell migration, adhesion and proliferation and eventually, tissue formation. The composite coatings strategy has been increasingly adopted by biomaterials researchers in the development of materials for human tissue repair and regeneration. Key words: composite coatings, bioceramics, osteoconductivity, implants, tissue engineering, scaffolds.
6.1
Introduction
Aided by advances in biological and medical sciences, materials science and engineering has triumphed in recent decades by providing various implant materials for human tissue repair for millions of patients all over the world. Nowadays, materials including metals, polymers, ceramics and composites are used clinically for implants and medical devices in many medical and dental areas (Black and Hastings, 1998; Ratner et al., 2004). Even though they have served their roles relatively well, some of the currently used implant materials that were not originally developed for medical applications have shortcomings in biological environments. One example of these shortcomings is the bioinertness of metallic biomaterials such as traditional orthopaedic metals of stainless steel and Co–Cr alloys, which results in the formation of fibrous tissue on implants of these materials in the human body (Ratner et al., 2004). Increasingly, Ti, Ti–6Al–4V alloy and NiTi shape 127
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memory alloy (SMA) are used for implants in orthopaedics, dentistry and cardiovascular surgery, etc owing to their unique properties (Brunette et al., 2001; Yahia, 2000). Nevertheless, the bioinertness of these metals and, in the case of NiTi SMA, the release to surrounding tissues of Ni ions, which are cytotoxic (Plant et al., 2005), during long-term implantation have limited their medical applications. Ceramic biomaterials were not seriously considered as viable biomaterials until the 1970s. Over the past 30 years, bioceramics, which include ceramics, glasses and glass–ceramics that are used for human hard tissue repair, have been intensively investigated (Winter et al., 1982; Oonishi et al., 1989; Prado and Zavaglia, 2009). Bioceramics such as bioinert alumina and toughened zirconia ceramics are now used in implants for hip joints (Wang, 2002). For over two decades since the early 1980s, extensive research has been conducted on bioactive bioceramics such as hydroxyapatite [HA, Ca10(PO4)6(OH)2], Bioglass® and A-W glass–ceramic (de Groot, 1983; Yamamuro et al., 1990; Hench and Wilson, 1993) because these materials can form chemical bonding with bone after implantation. (In this chapter, ‘bioactive’ is used interchangeably with ‘osteoconductive’.) However, most bioactive bioceramics including HA are weak ceramics and hence cannot be used on their own for load-bearing applications in human bodies (Wang, 2002). Through detailed studies of bioceramics such as Bioglass® and A-W glass–ceramic, it was established that the prerequisite for a material to be bioactive is its ability to form (or induce to form) apatite on its surface after coming into contact with bone in the body (Hench, 1991; Kokubo, 1992). The implant made of a bioactive material is bonded to bone through this intermediate apatite layer. Many investigations have shown that this apatite is similar in composition and structure to the apatite in bone. Therefore, since the 1980s, various techniques have been investigated by many groups around the world to form an apatite layer on metals such as Ti in order to make them osteoconductive. Some of these surface apatite-forming techniques for metals can also be applied to polymers (Wu et al., 2006a). Furthermore, investigations have been made into coating bioinert bioceramics such as alumina with an apatite layer in order to achieve bioactivity (Jiang and Shi, 1998). While plasma spraying has become the standard industrial technique for producing bioactive hip prostheses (Lacefield, 1993), low-temperature coating processes (‘low temperature’ here means temperatures below, or well below, 100 °C, which include room temperature and human body temperature), such as the biomimetic deposition of an apatite layer on metal substrates, are attracting increasing attention in the biomaterials community owing to their various advantages. For biomaterials such as NiTi SMA, bioactive coatings formed on their surfaces can serve two purposes: (1) to promote bone formation in vivo,
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and (2) to act as a physical barrier to prevent (or minimize) Ni ion release from the implant to surrounding tissues. In recent years, there have been increasing research activities on the surface modifications of NiTi SMAs for biomedical applications. Tissue engineering, since the definition and use of the term in 1988, (Skalak and Fox, 1988), has attracted great attention in science, engineering, medicine, and society in general. As a new and multidisciplinary endeavour, tissue engineering offers the potential of (a) eliminating re-operations by using biological substitutes, (b) using biological substitutes to solve problems of implant rejection, transmission of diseases associated with xenografts, and shortage in organ donations, (c) providing long-term solutions in tissue repair or treatment of diseases, and (d) potentially offering treatments for medical conditions that are currently untreatable such as fulminant hepatic failure. The discipline has made rapid advances over the past 20 years owing to more knowledge being gained in biology and medicine, the advances in physical sciences and technologies, and more willingness for collaboration and actual deeper collaborations among biologists, clinicians, engineers and scientists (McIntire, 2003). Synthetic tissue engineering scaffolds are considered an important part of a successful tissue engineering strategy. The scaffolds are designed to provide a structural framework as well as a microenvironment for the seeded cells and to facilitate the formation of new tissues. In bone tissue engineering, it is desirable that the tissue engineering scaffolds are osteoconductive. Currently, the scaffolds are mainly made of biodegradable polymers (Lanza et al., 2000; Ratner et al., 2004) but these polymers such as poly(l-lactic acid) (PLLA) and poly(ε-caprolactone) (PCL) are non-osteoconductive. For fabricating polymeric tissue engineering scaffolds, a range of techniques have been developed and can be used (Atala and Lanza, 2002). Nature produces the best designs and makes the best materials. Human body tissues are natural composite materials (actually, they are natural ‘nanocomposites’). Bone, for example, is a composite material with a hierarchical structure. It has two levels of composite structure: first, the bone apatite-reinforced collagen forming individual lamella at the nanometre to micrometre level; second, the osteon-reinforced interstitial bone at the micrometre to millimetre level. Using bone as the template, Bonfield and his co-workers in the early 1980s pioneered the research on bioactive bioceramic–polymer composites, which mimic the bone apatite–collagen structure and hence are sometimes called ‘bone analogue biomaterials’, for hard tissue replacement and regeneration (Bonfield et al., 1981). Since then, a variety of bioactive bioceramic–polymer composites of various characteristics have been investigated and developed in order to meet the various clinical requirements (Wang, 2003). With the obvious advantages of overcoming the shortcomings of constituent materials, having synergistic effects
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from the combination of different materials and tailoring composite properties through careful composite design, over the past two decades, the composite approach in developing new biomaterials has gained acceptance in biomaterials R & D. The composite strategy can also be employed for the surface modification of metal implants and tissue engineering scaffolds by forming composite coatings on their surfaces, thus greatly enhancing their osteoconductivity and other properties. Composite coatings offer distinctive advantages over single-component, monolayer coatings. Most coating techniques are adaptable for forming composite coatings on metal substrates and a few methods can be used to produce bioactive coatings on polymer-based tissue engineering scaffolds. This chapter presents our rationale, strategies and efforts in developing various osteoconductive composite coatings using various technologies for metal implants and bone tissue engineering scaffolds.
6.2
Design of composite coatings
For the surface modification of biomaterials, there are two major categories of designs for composite coatings: 1. a single-layer coating consisting of a matrix and a secondary phase which is distributed in the coating matrix; and 2. a coating having a layered structure, with the layers being different from each other and fulfilling different functions. They can be further classified into three designs, as shown in Fig. 6.1. (For simplicity, in this chapter only two-constituent composite coatings are discussed. Composite coatings containing more than two constituents can be made using some of the coating technologies discussed in this chapter and may be useful for some particular medical applications.) In a type I composite coating, the secondary phase (represented by dots in the diagrams of Fig. 6.1) is uniformly distributed in the matrix of the composite coating. This secondary phase can be a bioactive material (or useful biomolecules), with the coating matrix phase providing the support. It can also be a toughening material for the original single-component, bioactive coating. In a type II composite coating, the secondary phase has a continuous concentration gradient in the coating, thus forming a functionally graded coating (FGC), which can accomplish different goals for different coatings (viz. FGCs). For instance, when using plasma spraying to form a bioactive coating, an FGC can be designed to reduce the residual thermal stress of plasma sprayed coating; or, an FGC can be designed for an enhanced bonding between the coating and metal substrate. [This concept of FGC comes from the fabrication and use of functionally graded materials (FGMs) in industries such as aerospace and electrical power generation (Koizumi and
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Coating
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Substrate
(a)
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(b) L4 Coating
Coating
L3 L2 L1
C1
Substrate
(c)
C2
Substrate
(d)
6.1 Designs of composite coatings for biomaterials: (a) type I design, (b) type II design, (c) type III design, (d) type III (simple) design. (A higher density of dots in the coating indicates a higher density of the secondary phase in the composite coating.)
Niino, 1995; Kaysser and Ilschner, 1995). As in an FGC, an FGM has a gradient compositional change from the surface to the interior of the material.] However, the type II composite coating design is an idealized situation. In reality, type III composite coatings are fabricated and used instead of type II composite coatings owing to technical difficulties in producing type II composite coatings via most coating-forming techniques. The type III composite coating is a layered FGC, which is relatively easy to produce for most coating technologies. It consists of discrete layers, with the adjacent layers having similar composite compositions and, in each layer, the secondary phase is uniformly distributed in the coating matrix (only within this layer of the overall, layered composite coating). Here, the requirement of similar composition between adjacent layers is important because sufficiently dissimilar layers may cause delamination within the layered structure of the composite coating. (On the other hand, using the delamination strategy, an FGC of the layered structure may be able to reduce the effective
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driving force for coating fracture and arrest advancing cracks between the layers. This is a general concept for designing tough engineering materials and is beyond the scope of this chapter.) An example of the compositional arrangement for a type III composite coating can be as follows: (i) (ii) (iii) (iv)
Layer Layer Layer Layer
1 2 3 4
(L1) (L2) (L3) (L4)
containing containing containing containing
100% of material A; 75% of material A and 25% of material B; 50% of material A and 50% of material B; 25% of material A and 75% of material B.
with material B (or A) being uniformly distributed in material A (or B) in each layer. In its simplest form, the type III composite coating can be reduced to a composite coating consisting of two layers of different materials (i.e., type III (simple) composite coating), as shown in Fig. 6.1. In this two-layer structure, the bottom layer (i.e., the C1 layer) can act as an intermediate layer bonding the top, bioactive layer (i.e., the C2 layer) to the substrate, thus improving the original poor (or unsatisfactory) adhesion of the bioactive material (which makes the original monolayer coating of C2) to the substrate. This type of two-layer bioactive composite coating is analogous to some thermal barrier coatings (TBCs) in aeronautical applications which consist of a ‘bond coat’ (the C1 layer) and a ‘top coat’ (the C2 layer) where ‘C2’ fulfils the function of isolating heat from the metal substrate (Patterson et al., 2008). In another situation, the C1 layer can be a bioactive layer which induces the in vitro (or in vivo) formation of apatite and the C2 layer is the bioactive apatite layer formed in vitro (or in vivo). For bioactive coatings (including composite coatings) formed on metal substrates, it is important to have good adhesion to the substrates and this adhesion should be stable throughout the service life of the metal implants. The adhesion can be achieved through two mechanisms: 1. mechanical bonding: this is achieved through mechanical interlocking where the coating penetrates into surface irregularities or pores in the substrate surface; 2. chemical bonding: the coating (or the ‘bond coat’ in the composite coating) forms chemical bonds with the substrate. Mechanical interlocking between the coating and substrate relies on the surface roughness or porosity of the substrate. Grit-blasting is a common method to roughen the metal surface and the surface roughness created depends on the grit (Al2O3, SiC, or other hard particles) and grit size for grit-blasting, blasting pressure, blasting time, blasting angle, etc. Etching by strong acids such as HCl can also be used for making the metal surface rough. It is essential that before a coating is manufactured on a metal substrate, the surface of the metal is free from grease and also any contaminant including grits used for grit-blasting the surface.
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Chemical bonding of the coating to substrate is desired for coatings produced by most techniques to ensure strong adhesion, but it is avoided for coatings made by a few other techniques such as plasma spraying. (If chemical bonding occurs for a plasma sprayed coating, the chemical compound(s) formed at the coating-substrate interface may undermine the long-term adhesion of the coating.) It is sometimes desirable that the interface between the coating and substrate is a diffused one, i.e., there is no distinct (or discernible) boundary that separates the coating from substrate after coating formation, and that the chemical composition changes continuously from the substrate to coating, which may provide a very strong bonding for the coating. A number of coating techniques are capable of achieving the diffused interface for various coatings. Composite coatings for polymer-based tissue engineering scaffolds can use type I and type III (simple) coating designs shown in Fig. 6.1, with the substrate now being the struts of scaffolds. In a type I composite coating, the matrix of the coating can be an osteoconductive bioceramic and the secondary phase a polymer or biomolecules. Even though the polymer scaffold is intended to degrade during the tissue regeneration process and hence a strong adhesion of the coating to pore walls of the scaffold is not required (and not achievable) for the entire service time of the scaffold, sufficient bonding strength is nonetheless desired. In a type III (simple) coating, again the C1 layer can be a bonding layer on which the C2 layer of specially required tissue engineering functionality (e.g., a layer of biomolecules suitable for achieving enhanced cell attachment) is attached.
6.3
Technologies for the surface modification of biomaterials
The general requirements for bioactive bioceramic coatings for metal implants and their fabrication technologies include: 1. no alteration of the structure and composition of the bioceramic coating formed (if changes take place for the bioceramic during coating formation, the coating formed may not be bioactive), 2. good adhesion between the coating and implant, 3. no strength reduction of the metal implant after coating formation, and 4. if possible, (a) low cost for producing the coating, and (b) the technique being suitable for mass production. Only some coating technologies can meet most, if not all, of these requirements and the actual utilization of a particular coating technique sometimes is dependent on the compromise that can be made between the technology itself and other non-technological factors.
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Plasma spraying emerged in the 1980s as a feasible way to produce bioceramic coatings on metal substrates (de Groot et al., 1987) and has become industry’s choice for coating hip prostheses with a so-called ‘HA coating’ for cementless hip replacement (Lacefield, 1993). However, there are inherent problems such as decomposition of the feedstock HA powder for using the plasma spraying technique to form ‘HA coatings’ (Wang, 2004). Various groups have therefore sought alternative ways to form bioactive bioceramic coatings on metal implants. Techniques such as dipping-and-sintering (Aksakal and Hanyaloglu, 2008; Carsten et al., 1995), blast coating (Ishikawa et al., 1997; Yeo et al., 2008), electrochemical deposition (Kawashita et al., 2008; Monma, 1993), radiofrequency magnetron sputtering (Nakamura et al., 2007; Wolke et al., 1997), excimer laser deposition (Mroz et al., 2009; Singh et al., 1994), pulsed laser deposition (Tanaskovic et al., 2007; Wang et al., 1997), ion beam assisted deposition (Blalock et al., 2008; Wang et al., 2001a), plasma immersion ion implantation (Liu et al., 2005), and biomimetic deposition (de Groot et al., 2000; Kokubo, 1998; Sun and Wang, 2008a) were investigated for fabricating bioactive bioceramic (including glass) coatings on metal substrates. ‘Line-of-sight’ is a common problem for most of the aforementioned coating techniques, which form coatings of uniform thickness only on implants of simple shapes.
6.3.1 Plasma spraying Plasma spraying is one of the thermal spraying processes that utilize a high energy heat source to melt and to accelerate fine particles onto a prepared surface (Pawlowski, 2008). Upon impact, these molten particles (‘droplets’) cool down and solidify instantly by heat transfer to the underlying substrate and therefore form, by accumulation, a coating consisting of lamellae. Thermal spray processes have been used for many years to deposit layered coatings for various purposes such as wear resistance, thermal barrier, biocompatibility and bioactivity. The main thermal spraying techniques include flame spraying, plasma spraying, high velocity oxygen fuel (HVOF) spraying, vacuum plasma spraying, arc metallization, and detonation gun spraying. Among them, flame spraying, HVOF spraying and plasma spraying have been often investigated for fabricating bioactive bioceramic coatings on metal substrates. For plasma spraying, the creation and utilization of plasma as the hightemperature source are realized in the plasma torch (also called ‘plasma gun’) (Heimann, 1996; Pawlowski, 2008). As shown schematically in Fig. 6.2, the plasma torch consists of a cone-shaped thoriated cathode (tungsten) and a cylindrical anode (copper). Gases (for forming plasma) flow through the annular space between the two electrodes and an arc is initiated by
Composite coatings for implants and tissue engineering scaffolds Anode of plasma torch Water cooling chamber
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Powder feeder Metal substrate Plasma flame
Cathode of plasma torch Plasma arc Plasma gas
Plasma sprayed coating
6.2 Schematic diagram of the plasma spraying technique for forming coatings.
high-frequency discharge. The stream of gas which flows between the two electrodes stretches the arc, so that in its course from one electrode to the other, the arc loops out of the nozzle of the torch as a plasma flame. Gases such as Ar, He, H2 and N2 are mainly used as the arc gases (plasma-forming gases). The temperatures in a plasma flame are normally 10 000–15 000 °C. Therefore, in principle, almost any metal or ceramic including refractory metals or oxides can be melted and deposited to form a coating by plasma spraying. During plasma spraying, feedstock materials (in the form of powder, rod, or wire) are introduced into the plasma flame to be melted. The melted materials arrive on the surface of the target (the metal substrate) after having been sufficiently heated and accelerated by the plasma jet. When the droplets impact the target, they are flattened and spread out on the substrate surface and form a coating through successive impingement (Fig. 6.2). The velocity and temperature of powder particles or droplets are directly related to the plasma gas type. Most powders used for plasma spraying have particle sizes of 10–90 μm in diameter. Powders with this narrow size distribution are preferred in order to achieve uniform heating and acceleration of a single component powder. A constant powder feeding rate is a prime condition for achieving uniform coating thickness and good quality coatings. The powders to be injected into the plasma must also possess good flow properties, which is again associated with their morphology (preferred shape: spherical) and
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SHA coating
Substrate
6.3 SEM micrograph showing the layered structure of a plasmasprayed ‘HA’ coating.
size (Ye et al., 1999). The plasma sprayed coating is formed by the build-up of successive layers of particle droplets flattened upon impact and hence the coatings display layered structures. Figure 6.3 is a scanning electron microscopic (SEM) image of the cross-section of a dense ‘HA’ coating produced by plasma spraying using flame spheroidised HA (SHA) particles. The orientation of lamellae in the coating was mostly parallel to the Ti–6Al– 4V substrate. Good formation of the coating was evident by discontinuous lamellar lines in the coating. Owing to the large number of parameters involved in the plasma spraying process and the complex relationships that exist among these parameters, precise control and optimization of the plasma spraying deposition process is a tedious and expensive operation. The main advantages of using plasma spraying to form ‘HA’ coatings include good adhesion between the ‘HA’ coating and substrate and reproducible results; in addition, plasma spraying is an industrial process, enabling mass production. However, owing to the nature of the plasma spraying process, there are severe drawbacks: (a) bioceramic decomposition at plasma spraying temperatures; (b) high thermal residual stresses in the coating; and (c) the presence of defects such as unmelted particles, voids and cracks, if spraying parameters are not optimized. Nevertheless, extensive studies have been made by various groups on using the plasma spraying technology to form bioceramic coatings and there is a large body of literature on plasma spraying of monolayer bioceramic coatings (Hench and Wilson, 1993; Yamamuro et al., 1990). The thickness of plasma sprayed HA
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coatings normally ranges from 50 to 200 μm. It is feasible to use plasma spraying to produce multi-layered FGCs on metal implants.
6.3.2 Spraying-and-sintering To overcome the problems encountered in plasma sprayed bioceramic coatings, composite bioceramic coatings containing HA (or other bioactive bioceramics) can be produced using the spraying-and-sintering technique (Kumar and Wang, 1999, 2000, 2002a, 2002b; Kumar et al., 2001). The spraying-and-sintering technique is similar to the dipping-and-sintering method. Instead of dipping a metal sample in the bioceramic suspension or slurry, for spraying-and-sintering, the metal substrate is sprayed (using an airbrush) with a layer (or layers) of bioceramic particles that have been suspended in distilled water (or a solvent) in the storage bottle of a spraying gun. After spraying to form the powder coat on the metal substrate in the atmospheric condition, the coating is formed by sintering the powder coat in a high-temperature furnace (either conventional or vacuum furnace). The thickness of coatings (controlled by the number of passes during the spraying process) is 100–250 μm for FGCs. The spraying-and-sintering technique is a very simple and effective method for producing composite coatings containing HA or HA FGCs. It has several advantages over dipping-and-sintering as it eliminates the optimization of properties such as concentration and viscosity of bioceramic suspensions (or slurries) for dip-coating and makes it easy to form multilayered composite coatings in a reasonable production time. The FGCs on Ti–6Al–4V can be sintered at 900 °C, which is much lower than the plasma spraying temperature. To use even lower sintering temperatures and to eliminate voids in coatings, a HA–glass FGC can be designed, as shown in Fig. 6.4, and fabricated. In this design, a biocompatible (and preferably but not necessarily bioactive) glass is used as a bond coat for achieving enhanced adhesion between the top coat of HA and the metal substrate. The glass flows at the sintering temperature, and this helps to eliminate voids in the coating formed. It is ideal if the coefficient of thermal expansion (CTE) of the glass is compatible with that of the metal substrate in order to avoid crack formation in the coating or delamination of the coating from the metal substrate.
6.3.3 Ion beam assisted deposition To improve properties such as surface hardness and corrosion resistance of metals, ceramics and polymers for various engineering applications, since the mid-1970s, many surface modification techniques based on ion implantation, such as ion beam deposition (IBD), ion beam assisted deposition
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HA
HA + glass
Glass Ti–6Al–4V substrate
6.4 Schematic diagram showing the design (cross-sectional view) of an HA – glass functionally graded coating to be fabricated using the spraying-and-sintering technique.
(IBAD) and ion beam mixing (IBM), and techniques based on plasmaassisted ion implantation, such as plasma source ion implantation (PSII) and plasma immersion ion implantation (PIII), have been developed and are now widely used. Ion beam sputtering deposition (IBSD) was investigated as a method for producing bioactive bioceramic coatings on metal implants because it produced thin coatings with high density and good adhesion (Ong and Lucas, 1994; Ong et al., 1992). In this process, ionized argon gas was used to sputter atoms from a bioceramic target. The sputtered atoms built up on the Ti substrate that was placed in the path of the sputtered material. Both argon ion beam and nitrogen ion beam can be used for the ion-beam bombardment of the bioceramic target. By using a nitrogen ion beam to bombard the bioceramic target to form bioactive calcium phosphate (Ca–P) coatings, it is possible to further enhance the bone bonding ability of the coatings. If used as a beam to homogenize the coating being formed, nitrogen ions are able to penetrate the whole coating layer and react with the Ti substrate to form TixNy, which can improve the bonding strength of the coating. An ion beam sputtering / ion beam mixing deposition (IBS–IBMD) technique was therefore investigated to produce thin bioactive Ca–P or glass coatings on Ti substrate (Chen et al., 1999; Wang et al., 2000, 2001a, 2001b, 2001c, 2001d, 2002a). An IBS–IBMD system used in the studies is shown schematically in Fig. 6.5. For producing Ca–P coatings on Ti substrate, before IBS–IBMD, pure Ti plates (20 × 20 × 1 mm in dimension) were mechanically polished and ultrasonically cleaned with acetone and alcohol. HA powder was cold pressed into discs and sintered (at 1150 °C for 3 h). These HA discs were to be used as the ion-beam sputtering target. (The sputtering target disc
Composite coatings for implants and tissue engineering scaffolds Vacuum: 2.8~3.7 × 10–4 Pa
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Sputtered atoms 4
2
Sputtering target disc
Ion beam
Ion beam
Coating Metal substrate 3
1
6.5 Schematic diagram showing the ion beam sputtering / ion beam mixing deposition (IBS–IBMD) technique: 1 low-energy ion source for sputtering; 2 high-energy ion source for mixing; 3 rotating sample holder; 4 sputtering target holder.
could also be made from a bioactive glass when bioactive glass coating was fabricated through IBS–IBMD.) The IBS–IBMD system consisted mainly of a Kaufman ion source, a Freeman ion source, a target holder and a rotating sample holder (Fig. 6.5). The deposition chamber was evacuated to a base pressure of 2.8∼3.7 × 10−4 Pa. Before Ca–P coating deposition, etching of the substrates with 800 eV and 40 mA cm−2 argon ions needed to be performed for 30 min to clean the surface of Ti substrate. The energetic ion beam was produced by ionizing high-purity argon gas (99.999% pure). After cleaning, the stage was rotated so that the Ti plates were placed in the path of sputtered atoms from the target disc. Monolayer Ca–P coatings could be sputter-deposited by an Ar+ beam with 1200 eV and 40 mA cm−2 for 60 min. For comparative studies, different combinations of two ion beams could be used for sputtering and mixing, respectively, during IBS– IBMD. Ar+ beam sputtering- and Ar+ beam mixing-deposited Ca–P coatings were produced first by sputtering the HA target with a 1400 eV and 40 mA cm−2 Ar+ beam for 1.5 h and then by using the second Ar+ beam with 60 keV to homogenize the coatings. The dosage of the Ar+ mixing beam could be 2.0 × 1016 ions cm−2. For Ar+ beam sputtering- and N+ beam mixingdeposited Ca–P coatings, after sputtering the HA target with a 1400 eV and 40 mA cm−2 Ar+ beam for 1.5 h, a N+ beam with 60 keV was used to homogenize the coating, and its dosage could also be 2.0 × 1016 ions cm−2. Obviously, mechanical and biological properties of IBS–IBMD coatings (bioactive Ca–P or glass) were influenced by coating conditions such as ion beam combination and ion beam parameters. The as-deposited coatings
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(either Ca–P or glass) were amorphous and very thin (∼500 nm for Ca–P or glass coatings), which are very different from coatings produced by plasma spraying or spraying-and-sintering. Cell culture experiments showed that bioactive monolayer Ca–P or glass coatings fabricated by IBS–IBMD promoted the proliferation of osteoblastic cells. Through modifications of the coating process, IBS–IBMD can be used to fabricate thin bioactive FGCs on metal substrates.
6.3.4 Biomimetic deposition In biomaterials R&D, it is through in vivo animal experiments that important information about the bioactivity of a biomaterial is gained. However, there are a variety of factors which can affect the bioactivity of a biomaterial and sometimes it is difficult, if possible at all, to distinguish and/or evaluate the effects of individual factors on bioactivity in the complex situation of in vivo experiments. Therefore, simple in vitro methods have been constantly sought for investigating the bioactivity of materials. Kokubo and his coworkers started the trend of using a simulated body fluid (SBF) for assessing the bioactivity of materials (Kokubo and Takadama, 2006; Kokubo et al., 1990), using their own A-W glass-ceramic for initial investigations. Research groups in other parts of the world adopted the similar strategy for bioactivity assessment, even though their test fluids used in the experiments could be different (de Groot et al., 2000; Hench, 1991). Using the SBF that was introduced by Kokubo and his co-workers (Ohtsuki et al., 1991) or other test fluids, we studied the in vitro bioactivity of various biomaterials (Chen and Wang, 2002; Huang et al., 1997a, 1997b; Ni and Wang, 2002; Wang and Ni, 2004; Wang et al., 1998, 2001, 2002; Weng and Wang, 2001; Weng et al., 2001, 2002, 2005). (Researchers need to be aware, though, that in the biomaterials R&D community, there are arguments about the validity of using in vitro bioactivity results to ‘project’ in vivo bioactivity of the material.) Looking in the opposite direction of in vitro bioactivity studies, SBF can be used to form a bioactive apatite coating on a material if the material is bioactive and, thus, ‘biomimetic deposition’ came into being in the 1990s through the pioneering work conducted by groups led, respectively, by Kokubo, de Groot and a few others (de Groot et al., 2000; Kokubo, 1998). As schematically shown in Fig. 6.6, biomimetic deposition is a simple process. A bioactive material (in this case, a surface pre-treated metal sample) is immersed in the SBF (or one of its variants) contained in a water bath which has a reciprocal motion (a shaker bath). The SBF in the container (a plastic bottle for example) is maintained at human body temperature (36.5 °C). Depending on the bioactivity of the material (for example, a Ti surface treated by different methods can exhibit different degrees of bioactivity) and the strength of the solution for immersion [Classic SBF is
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Lid or bottle cap Metal sample
Simulated body fluid
Water
pH 7.4, 36.5 °C
Shaker bath
6.6 Schematic diagram showing the biomimetic deposition technique.
the original SBF introduced by Kokubo’s group (Ohtsuki et al., 1991), which is an improvement over an earlier one (Kokubo et al., 1990), that has ion concentrations nearly equal to those of human blood plasma; 2SBF is a high-strength SBF whose ion concentrations are twice those of human blood plasma; 5SBF has ion concentrations five times those of human blood plasma; and so on.], an apatite layer can form on the material surface within days during immersion. Investigations have shown that this apatite is similar in composition and structure to apatite in bone and it is thus termed ‘bonelike apatite’ (BLA) (Kokubo, 1998). Some studies have suggested that the coatings biomimetically formed in SBF and similar solutions are not apatite but octacalcium phosphate [OCP, Ca8(HPO4)2(PO4)4·5H2O] (Lu and Leng, 2004, 2005). In this chapter, the traditionally adopted term ‘bone-like apatite’ is used for the biomimetically deposited Ca–P coatings. A few metals including Ti and its alloys can induce the formation of bone-like apatite on their surfaces after they have undergone suitable chemical and thermal treatments. Surface modification of implantable metals through surface apatite formation via biomimetic deposition appears to be a good approach for improving the usability of these metals in the medical field. Biomimetic deposition takes place at human body temperature and does not require harsh reaction conditions such as high (or low) pressure, or high temperature, thus avoiding the major disadvantages of other surface apatite-forming techniques such as those discussed in previous sections of this chapter. Biomimetic deposition is a ‘low-temperature’ process, ‘low temperature’ here being comparative to those in plasma spraying, spraying-and-sintering, and other techniques that involve the use of high temperatures ranging from several hundred degrees Celsius to over ten thousand degrees Celsius.
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Mechanisms and kinetics of biomimetic apatite formation and growth on Ti surfaces have been studied using a variety of analytical techniques (Wang and Wang, 2001a, 2001b, 2002, 2003; Wang et al., 2002b, 2003, 2004). It was found that in vitro apatite formation depended on the solution and the procedure that were used for biomimetic deposition. Electrochemical impedance spectroscopy (EIS) can be a useful tool for monitoring the formation and growth of apatite on pretreated metals. In one series of studies (Wu et al., 2006b, 2006c, 2006d, 2007), a low-temperature process involving the use of H2O2 solution for metal surface oxidation was investigated and developed for the pretreatment of Ti substrate before biomimetic deposition of the apatite layer in SBF, thus totally eliminating the deleterious thermal treatment of metal substrates which is required in the ‘biomimetic processes’ used by other research groups. Hot water ageing of the titania gel formed through H2O2-oxidation was used in order for the gel to crystallize and become bioactive. The crystallized synthetic TiO2 layer on Ti subsequently induced the formation of apatite in SBF. Figure 6.7 exhibits the surface morphology of Ti substrate after surface oxidation through H2O2 treatment and after biomimetic apatite deposition in SBF. Our more recent studies have focused on establishing and using the low-temperature process (i.e., hydrogen peroxide oxidation and hot water ageing with subsequent biomimetic apatite deposition in 5SBF) for the surface modification of NiTi SMA (Sun and Wang, 2008a, 2008b, 2008c, 2008d, 2008e, 2008f).
6.3.5 Other surface modification techniques for implants As stated earlier, coating techniques such as dipping-and-sintering, blast coating, magnetron sputtering, and laser deposition have also been investigated by various groups for the surface modification of metal implants. Electrochemical deposition is another attractive method, which is a simple and low-cost technique and is also a low-temperature process that is capable of fabricating an apatite coating on metal substrates within a few hours. It was thus investigated for forming apatite coatings on Ti and Ti alloys (Kawashita et al., 2008; LeGeros et al., 2008). When appropriate deposition parameters are used, it can produce apatite coatings which are homogeneous in chemical composition and microstructure. Furthermore, electrochemical deposition of bioactive coatings is not only suitable for metal implants with complicated shapes and porous structures, but it can also provide the possibility of incorporating proteins and/or useful species (such as collagen) in coatings. Therefore, electrochemical deposition is a promising technique for the surface modification of metallic biomaterials. Our recent work has focused on using this technique to fabricate bioactive apatite coatings and apatite-containing composite coatings on NiTi SMA (Sun and Wang, 2009; Sun et al., 2009a). Figure 6.8 shows the nano-
Composite coatings for implants and tissue engineering scaffolds
(a)
(b)
Mag = 4.02 K X
2 μm
407
EHT = 5.00 kV WD = 6 mm
20 KV
Signal A = InLens Photo No. = 330
10 μm X2,000
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Date: 25 Aug 2005 Time: 15:02:14
23 mm
6.7 Surface morphology of a Ti plate: (a) after H2O2-oxidation treatment; (b) after biomimetic apatite deposition in SBF.
particulate morphology of electrochemically deposited monolayer apatite coating on NiTi SMA substrate.
6.3.6 Coating techniques for tissue engineering scaffolds Good tissue engineering scaffolds are a pillar of the successful tissue engineering approach in treating tissue loss and organ failure of human bodies. For bone tissue engineering, major issues of tissue engineering scaffolds include the use of appropriate materials for scaffolds, control of porosity
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Mag = 100.00 K X
100 nm
EHT = 4.00 kV Photo No. = 5408
Signal A = InLens Date: 25 Mar 2008
WD = 3 mm Time: 15:46:02
6.8 Surface morphology of an apatite coating on NiTi SMA fabricated through electrochemical deposition.
and pore characteristics of scaffolds, mechanical strength of scaffolds, scaffold degradation properties, and bioactivity (i.e., osteoconductivity, or osteoinductivity if possible) of scaffolds. Biomaterials experience with biodegradable polymers in the pre-tissue engineering time (before 1987) has heavily guided researchers towards using biopolymers such as PLLA and poly(lactic acid-co-glycolic acid) (PLGA), which are accepted by the medical profession and approved by the US Food and Drug Administration (FDA), for constructing tissue engineering scaffolds. In recent years, other polymers have also been developed for tissue engineering applications (Lanza et al., 2000; Ratner et al., 2004). There are many reports on scaffolds made of these polymers. Compared with the strengths of metals and ceramics for medical applications, the strengths of non-porous biodegradable polymers are already very low. With the introduction of pores in the polymers to form tissue engineering scaffolds, the strengths of porous polymer structures are further decreased (Gibson and Ashby, 1997). On the other hand, polymers such as PLLA, PLGA, PCL, polyhydroxybutyrate (PHB) and poly(hydroxybutyrate-co-hydroxyvalerate) (PHBV) are considered to be non-osteoconductive. In order to find better scaffolds for bone tissue engineering, a composite strategy can be adopted (Wang, 2003). Polymerbased scaffolds containing bioactive bioceramics can be produced in which the bioceramics can serve two purposes: (a) making the scaffolds osteoconductive, and (b) reinforcing the scaffolds. With this composite strategy, there are two main routes for the manufacture of polymer-based osteoconductive scaffolds for bone tissue engineering:
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1. producing scaffolds that incorporate bioceramic particles using a variety of techniques including porogen leaching (Weng et al., 2002), emulsion freezing / freeze-drying (Sultana and Wang, 2006, 2007, 2008a, 2008b), electrospinning (Tong and Wang, 2006, 2007a, 2007b), and selective laser sintering (SLS) (Duan et al., 2008a, 2008b, 2008c; Zhou et al., 2007, 2008a, 2008b), and 2. coating the pore surface of a polymer scaffold with a thin layer of apatite or apatite-containing composite layer through biomimetic processes (Chen et al., 2005a, 2005b, 2005c, 2006a, 2006b, 2007, 2008a, 2008b). These two routes have respective advantages and disadvantages and the latter can be used to make the readily fabricated non-bioactive polymer scaffolds osteoconductive. Incorporating bioceramic particles (microparticles previously, and nanoparticles in recent years) into biodegradable polymers to form bioactive scaffolds can be successful, but there is an upper limit to the amount of particles that can be incorporated. It has been shown that bone-like apatite could form in vitro on bioactive composite scaffolds (Weng et al., 2002), indicating osteoconductivity of these scaffolds. As indicated above, this route of making bioactive bone tissue engineering scaffolds is being explored, with several fabrication techniques such as SLS, electrospinning and emulsion freezing / freeze-drying being currently investigated actively and with bioactive nano-sized particles (HA and Ca–P) being incorporated in the scaffolds. Even for these nanocomposite scaffolds, surface modification is sometimes performed for scaffold improvements. As PHBV is a hydrophobic biodegradable polymer, for improving surface wettability and hence achieving enhanced cell attachment, the surface of pores in HA/ PHBV nanocomposite scaffolds can be coated with a thin layer of collagen using the dip-coating method (Sultana and Wang, 2009). In dip-coating, composite scaffolds were immersed in a dilute collagen solution for 1 h. After being removed from the collagen solution, the scaffolds were air dried. It was shown that for the collagen-coated composite scaffolds, osteoblastic cells became attached to the pore walls, where they were anchored well by discrete filopodia. In a separate study, to achieve a balance of hydrophilicity/hydrophobicity for nanocomposite scaffolds, which is important for protein adsorption and subsequent cell behaviour, the approach of physical entrapment of hydrophilic gelatin on the struts of Ca–P/PHBV scaffolds made by SLS was adopted (Duan and Wang, 2008). In this simple process, gelatin was dissolved in a miscible mixture of 2,2,2-trifluoroethanol (TFE) and water, which is a solvent (TFE) and a nonsolvent (water) for PHBV, respectively. The sintered Ca–P/PHBV scaffolds were immersed in the gelatin solution. After the immersion treatment, the surface modified Ca–P/PHBV composite scaffolds were rinsed to remove non-entrapped
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gelatin and dried at room temperature. During the gelatin entrapment process, the PHBV matrix of scaffolds swelled but did not dissolve in the gelatin solution. Gelatin molecules in the solution then diffused into the swollen surface of PHBV and were entangled with PHBV molecules on the scaffold strut surface. After the scaffold was removed from the gelatin solution and immersed in water, which is a non-solvent for PHBV, the surface rapidly contracted and gelatin molecules on the polymer surface were entrapped and immobilized. These gelatin molecules were only planted in the PHBV area of the strut surface of Ca–P/PHBV composite scaffolds, leaving originally exposed Ca–P particles on the struts still uncovered and hence not affecting the osteoconductivity of the composite scaffolds. The contact angle measurements showed that the physical entrapment of gelatin improved the hydrophilicity of PHBV and hence sintered composite scaffolds. Using non-osteoconductive biodegradable polymer scaffolds, biomimetic apatite deposition on pore surfaces of the scaffolds provides an alternative route for fabricating bioactive composite scaffolds. In the classic biomimetic process to form an apatite layer on metal or ceramic surfaces, normal-strength SBF is commonly used (see section 6.3.4.) and it usually takes 1–4 weeks to form the apatite layer. However, as polymeric scaffold materials such as PLLA are easily hydrolysed in water, a much shorter coating time must be used in the biomimetic deposition process in order to avoid hydrolysis of the biodegradable polymer scaffolds. Therefore, an accelerated biomimetic deposition process which employs a higher-strength SBF (2SBF, 5SBF, etc.) can be used. Using accelerated biomimetic deposition, an apatite layer can form on the surfaces of poly(glycolic acid) (PGA) fibre meshes and PLLA scaffolds within 24 h (Chen et al., 2005a, 2005b, 2005c). The biomimetic deposition of coatings in the ‘static’ condition involved the same experimental setup as that shown in Fig. 6.6. In this condition, even though the shaker water bath makes reciprocal motions, the coating solution (e.g., 5SBF) in the container (e.g., a capped bottle) may not flow through the porous scaffolds thoroughly. Nevertheless, bone-like apatite can be formed on pore surfaces of the PLLA scaffolds (Fig. 6.9a). There are concerns that in the static biomimetic deposition process, less apatite is formed in the interior of a scaffold than the peripheral region because of easier ion exchange between the peripheral region and the surrounding coating solution. The formation of a uniform spatial distribution of apatite coating in the scaffold (interior and surface areas) may be achieved by using dynamic biomimetic deposition, as shown in Fig. 6.10. With dynamic biomimetic deposition, a circulation is established, in which the coating solution flows directly through the polymer scaffolds (Chen et al., 2007). Therefore, high transport of ions is achieved between the internal regions and flowing
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6.9 Bone-like apatite coating formed on pore surfaces in the interior of PLLA scaffolds through accelerated biomimetic deposition: (a) apatite coating formed through static biomimetic deposition; (b) apatite coating formed through dynamic biomimetic deposition.
coating solution, and consequently more apatite particles are formed on the internal pore walls in the scaffold. Figure 6.9b shows the apatite layer formed on the surface of a pore in the interior of a PLLA scaffold through dynamic biomimetic deposition. With modification(s) of the process, composite coatings can be fabricated on polymer scaffolds, thus enhancing their biological performance.
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6.10 A schematic diagram showing the dynamic biomimetic deposition technique for bone tissue engineering scaffolds.
6.4
Composite coatings for implants
The coating technologies described and discussed in previous sections have been used to make single-component, monolayer apatite (or glass) coatings on metal substrates or on polymer tissue engineering scaffolds. With appropriate modification(s) of the coating technique and/or using suitable mixture(s) of powders as feedstock material(s) (or coating solutions containing a second or even a third component), composite coatings of various compositions and microstructures, as schematically represented in Fig. 6.1, can be formed by these technologies for obtaining improved/enhanced mechanical and/or biological properties for implants or scaffolds. This section presents and discusses some of the composite coatings formed so far in our research.
6.4.1 Plasma sprayed coatings Our work on plasma spraying of bioactive Ca–P coatings has concentrated on obtaining toughened coatings with high bonding strength with Ti substrate (Wang et al., 1999; Yang et al., 1998). As HA is a brittle and relatively weak ceramic, some tough but bioinert bioceramics such as TiO2, Al2O3 and ZrO2 can be used to form HA-based bioactive ceramic composite coatings for improved mechanical performance. Other calcium phosphates such as tricalcium phosphate (TCP), Ca3(PO4)2, can also be considered for toughened bioactive coatings. Additionally, owing to its biodegradability, TCP in the HA-based coatings can gradually dissolve after implantation so that enhanced osseointegration is promoted. For achieving good adhesion
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between the coating and metal substrate and improved coating properties (mechanical and/or biological), a three-layer FGC structure was employed, which is exemplified by the HA–TCP composite coating as shown in Fig. 6.11. The layer which bonds with the substrate (layer 1: the HA layer in this case) should be a mechanical support for the whole coating system. In the meantime, the layer on the surface of the FGC (layer 3: the TCP layer in this case) should be bioactive (and sometimes biodegradable) in order to possess good osteoconductivity. The layer which is located between the aforementioned two layers (layer 2: the 50% TCP + 50% HA composite layer in this case) may have a compromised mechanical/biological performance. In producing such a layered composite coating, pure HA and TCP powders were used as feedstock materials for plasma spraying to produce layer 1 and layer 3 of the FGC, respectively. For layer 2 of the FGC, HA and TCP powders were mixed well using a planetary mill to produce an HA and TCP particle mixture (termed ‘composite powder’). This composite powder was used for plasma spraying to make the 50% TCP + 50% HA composite layer. During the plasma spraying operation for making HA–
α-TCP
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6.11 Coating structure (cross-sectional view) of a three-layer HA–TCP FGC produced by plasma spraying.
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TCP FGC, the layers of FGC were formed in sequence (layer 1, layer 2 and, finally, layer 3) using different feedstock materials for plasma spraying, with no time lapses between plasma spraying of adjacent layers. In a similar manner, a layered HA–TiO2 FGC was fabricated: (1) layer 1, pure TiO2; (2) layer 2, 50% TiO2 + 50% HA composite; and (3) layer 3, pure HA. Both the HA–TCP and HA–TiO2 FGCs follow type III coating design (Fig. 6.1), with the top (layer 3) and bottom (layer 1) layers being single-phase (singlecomponent) layers. Micro- and nanoindentation techniques can be used to assess mechanical property changes across the cross-section of FGCs (Wang et al., 1999, 2004). Figure 6.12 shows the variation of microhardness in a plasma sprayed HA– TiO2 FGC. The first layer, which was composed of TiO2, exhibited a high microhardness value (around 940 VHN). Microhardness decreased dramatically to about 400 VHN in the second layer which had 50 wt.% of HA. The microhardness value of the third layer (i.e. the HA layer) was similar to the value obtained for plasma sprayed monolayer HA coatings. The microhardness in the transitional area between adjacent layers was relatively low. These low values may be attributed to the relatively high (tensile) residential stress at these locations. A similar trend in microhardness changes was established for the HA–TCP FGC. When cracks were induced at the tips of indentations and radiated from the tips in bioceramic FGCs under microindentation loads, indentation fracture toughness (KIc) could be calculated using well established theories. KIc of the TCP layer in HA– TCP FGC was about 1.41 MPa m−1/2, which is higher than that of pure, monolayer HA coatings (0.5–1 MPa m−1/2, depending on the particle size of
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6.12 Variation of microhardness on the cross-section of a plasma sprayed HA–TiO2 FGC.
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feedstock HA powders). Furthermore, KIc of the HA/TCP composite layer of HA–TCP FGC remained at 1.28 MPa m−1/2, whereas KIc of the HA layer of HA–TCP FGC decreased to 1.05 MPa m−1/2. For the TiO2 layer of HA– TiO2 FGC, microcracks could not be induced under the indentation load of 2 kgf. Therefore, it could be inferred that fracture toughness of the TiO2 layer in the FGC was far higher than that of pure HA coatings. KIc of the HA/TiO2 composite layer of HA–TiO2 FGC which contained 50 wt% of TiO2 was 1.76 MPa m−1/2, which was still much higher than 0.47 MPa m−1/2 that was obtained from the HA layer of HA–TiO2 FGC. These results indicated that both TCP and TiO2 were effective in increasing microhardness and fracture toughness of ‘HA coatings’, with TiO2 offering a greater toughening efficiency. These studies have demonstrated the effectiveness of adopting the composite strategy in improving mechanical properties of well established bioceramic coatings in clinical use.
6.4.2 Coatings produced by spraying-and-sintering For the particular combination of HA or HA-related materials as (part of) the coating and Ti or Ti alloys as the substrate, the spraying-and-sintering technique can only be used to fabricate composite coatings, not HA or other bioceramic monolayer coatings, on Ti-based metals for the following reasons: (1) the sintering temperature for the coating should be around 900 °C or lower to avoid causing structure damage to and mechanical property reductions of the metal substrate and to avoid the formation of deleterious titanium compounds at the coating-substrate interface; and (2) the sintering temperature cannot be much lower than 1000 °C otherwise a dense and well-bonded coating based on one of the pure bioactive Ca–P (e.g., HA or TCP) cannot be achieved. (Dense and sufficiently strong HA bars or rods are obtained by sintering at 950–1300 °C, with 1360 °C being the HA decomposition temperature which should not be exceeded.) Even though through spraying-and-sintering, HA–TCP FGC (Kumar et al., 2001), HA–TiO2 FGC (Kumar and Wang, 2002a) and HA–ZrO2 FGC (Kumar and Wang, 1999) were made by sintering them at 900 °C and these FGCs were studied, the HA–glass FGC appeared to be the only system that had achieved good coating integrity and compact coating structure (Kumar and Wang, 1999, 2000, 2002b), possessing the potential for clinical applications. The thickness of the FGCs formed was 100–250 μm. Figure 6.4 shows the idealized microstructure for HA–glass FGC with a continuous HA concentration gradient in the coating. In practice, the HA–glass FGC was fabricated using four layers to form the composite coating (type III
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coating design, as shown in Fig. 6.1): (1) layer 1: 100% glass; (2) layer 2: 75% glass + 25% HA; (3) layer 3: 50% glass + 50% HA; and (4) layer 4: 25% glass + 75% HA. The glass powder for fabricating HA–glass FGC in our investigations was produced in-house. Its composition (by weight) was 67.7% SiO2, 10.4% B2O3, 5.2% Al2O3, 8.3% Na2O, 4.2% K2O, 2.1% Li2O, 1.05% ZrO2 and 1.05% TiO2. To make stock materials for spraying in forming powder coats which eventually became layers of the composite coating after sintering, HA powder of various weight percentages such as 25, 50 and 75% was mixed with glass powder of corresponding weight percentages of 75, 50 and 25%, respectively, to form composite powders. Composite powders of different compositions, for instance, 25% HA + 75% glass, were mixed well using a planetary mill. They were subsequently washed in acetone and dried and were ready for spray-coating. The Ti–6Al–4V substrate was grit-blasted with Al2O3 grit and then cleaned in acetone before being used for coatings. The surface roughness (Ra) of grit-blasted Ti–6Al–4V plates was controlled to be 3–5 μm. The HA–glass FGC was produced by spraying composite powders layer by layer using an airbrush with an oil-less diaphragm pump under controlled pressure at room temperature. Sintering of powder coats on Ti–6Al–4V plates at 900 °C for 5 min was performed to ensure good adhesion of the bottom layer (glass layer) to the metal substrate and also between composite layers. HA-glass FGC was thus uniformly formed on the Ti–6Al–4V substrate, which had a dense structure (Fig. 6.13). X-ray diffraction (XRD) analysis of the HA–glass FGC (and also other types of
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6.13 Coating structure (cross-sectional view) of an HA-glass FGC produced by spraying-and-sintering.
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FGCs such as HA–TCP FGC) revealed that no HA decomposition was encountered with the spraying-and-sintering technique. For HA–glass FGC, Fourier transform infrared spectroscopy (FTIR) confirmed the presence of the hydroxyl group (at 3564–3572 cm−1) and the phosphate group (at 962– 1092 cm−1), indicating that the HA in the FGC retained functional groups for its bioactivity. Nanoindentation was also used to evaluate mechanical properties of FGCs. Figure 6.14 displays a load-displacement curve obtained from a nanoindentation made on the cross-section of an HA-glass FGC. Using a maximum indentation load of 300 mN, the hardness and elastic modulus of the HA-glass FGC were determined to be 5.99 GPa and 27.3 GPa, respectively. When nanoindentation was applied to the HA–TCP FGC which was made following the same procedure and using the same processing parameters, the hardness and elastic modulus of the HA–TCP FGC were 0.68 GPa and 15.4 GPa, respectively, indicating that at 900 °C, the HA–TCP FGC was not well sintered and hence had lower mechanical properties than the HA–glass FGC. Using Ti–6Al–4V as the substrate, the HA–TCP FGC (and other FGCs without glass) could not be improved as the sintering temperature could not be raised.
6.4.3 Coatings produced by ion beam assisted deposition In the normal IBS–IBMD process (Fig. 6.5), the sputtering target disc is made of only one material such as HA or bioactive glass and hence only single-component, monolayer coatings are formed on the metal substrate. However, if a series of sputtering target discs are made of a composite
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6.14 Typical nanoindentation load–displacement curve obtained from the cross-section of a HA–glass FGC produced by spraying-andsintering.
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(consisting of material A and material B) of different compositions and are sequentially changed for the sputtering–deposition process, an FGC, which is type III coating design as shown in Fig. 6.1, can form on the metal substrate. Therefore, using this concept, an investigation was conducted to form a calcium phosphate–Ti FGC through IBS–IBMD (Wang et al., 2001e). This FGC should combine a strong adhesive strength to the Ti substrate with good bioactivity at the coating surface (or near-surface). In the investigation, hydroxyl poly calcium sodium phosphate (HPPA) with the chemical formula of (NaCanPmOk)M·H2O was used as the bioactive component. HPPA powder discs and discs of HPPA powder mixed with pure Ti powder in different proportions were made through cold pressing. These discs were used as the ion beam sputtering targets. They had Ti contents of 100, 90, 70, 50, 30, 10 and 0% and the sequential change of sputtering target discs during IBS–IBMD was also of this order, forming, first of all, a Ti layer on the Ti substrate. For fabricating the HPPA–Ti FGC on Ti substrate, the IBS–IBMD system and operation were the same as described in section 6.3.3, but with the sequential change of sputtering target discs in the coating forming process. Two types of HPPA–Ti FGCs were produced: (1) using ion beam sputtering only for coating deposition without ion beam mixing (the IBS only process); (2) using ion beam sputtering and ion beam mixing for coating deposition (the full IBS–IBMD process). When IBS–IBMD was employed for making the FGC, an Ar+ beam (1200 eV and 40 mA cm−2, 30 min) for sputtering each target disc and another Ar+ beam (60 keV, 1.5 × 1015 ions cm−2 dosage) for mixing (i.e., homogenizing within each coating layer of the FGC) were used. For investigating the distribution of elements in the deposited thin coatings, X-ray photoelectron spectroscopic (XPS) analysis was conducted. Depth profiling was accomplished by sputter-etching the coating and obtaining chemical analysis data from the sputter-etched area. Using Ar+ ions, a differentially pumped ion gun (equivalent to a sputtering rate of about 10 nm min−1) with a potential of 4 kV was used to sputter-etch the samples. The analysis of the sputtered surfaces was performed until the surface of the Ti substrate was reached. XRD analysis of FGCs indicated that they were amorphous. FTIR analysis of FGCs confirmed the presence of phosphate group (at 1092 cm−1) but no distinct absorption bands of hydroxyl group were observed. New absorption bands at 1436 and 939 cm−1 were present for the carbonate group which was brought about by the deposition process. Figure 6.15 exhibits the surface morphology of HPPA–Ti FGC produced by IBS–IBMD. The coatings had a smooth surface even though they were very thin. The depth profiles of Ca, P and Ti of the HPPA–Ti FGC produced by IBS–IBMD are
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6.15 Surface morphology of HPPA–Ti FGC produced by ion beam sputtering / ion beam mixing deposition (IBS–IBMD).
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6.16 Depth profiles of Ca, P and Ti of the HPPA–Ti FGC produced by IBS–IBMD.
shown in Fig. 6.16. It can be seen that from the surface to the interior of the FGC, the concentration of Ti increased, while concentrations of calcium and phosphorus decreased. In addition, compared with FGC produced by IBS only, the composition changes in FGC produced by IBS–IBMD were more gradual. Owing to the IBMD process, complete mixing of Ti and calcium phosphate could be achieved. Consequently, there was elemental Ti on the surface of the FGC produced by IBS–IBMD, while no elemental Ti was detected on the surface of the FGC produced by IBS only. In FGCs produced by IBS–IBMD, because of the compositional gradient, the residual stress caused by the difference in coefficient of thermal expansion was relaxed to some extent. Therefore, when a post-deposition heat treatment
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was used to crystallize the Ca–P in the FGC, the occurrence of coating layer peeling-off in FGCs was minimal. In another study, it was found that under certain sputtering/mixing conditions, using a single HA sputtering target disc, an FGC could also be created on Ti substrate by the IBS–IBMD technique (Wang et al., 2001f). But this FGC is not as desirable as and similar to the HPPA–Ti FGC produced by IBS–IBMD.
6.4.4 Coatings formed by biomimetic deposition As has been pointed out in section 6.3.4, the low-temperature route for biomimetic apatite deposition, which involves H2O2-oxidation and hot water aging with subsequent biomimetic apatite deposition in SBF, was well established for Ti-based metals. Well-crystallized anatase thin films with excellent in vitro bioactivity could be produced on Ti surfaces by soaking the Ti substrate in 30 mass% H2O2 solution at 80 °C for 2–72 h, followed by a hot water aging at 80 °C for 72 h. During hot water ageing, the amorphous titania gel produced by H2O2-oxidation hydrolysed and re-precipitated back to the Ti substrate to form anatase nanocrystals. Ti treated by this low-temperature chemical modification of the surface could induce apatite formation in SBF within 24 h. The apatite layer formed was dense and uniform. The trace elements in the titania layer were found to affect greatly the apatite formation on Ti in SBF. The Cl incorporated in the titania layer did not hinder apatite formation but F did. Using this low-temperature method, a two-layer composite coating (type III (simple) coating design, as shown in Fig. 6.1), which consisted of the crystallized TiO2 bottom layer and the apatite surface layer, could be fabricated on Ti substrate (Wu et al., 2006c, 2007). Using 5SBF instead of SBF for forming the apatite surface layer more quickly, the ‘accelerated’ low-temperature biomimetic deposition method has been applied to NiTi SMA (Sun and Wang, 2008a, 2008b, 2008c, 2008d, 2008e, 2008f). Before H2O2-oxidation treatment, energy dispersive X-ray spectroscopic (EDX) analysis revealed the distinctive presence of Ni and Ti elements in the surface of NiTi SMA samples that were polished and cleaned via the established procedure, whereas the content of elemental O was too low for O to be detected, indicating the surface of untreated samples mainly consisted of homogeneously distributed Ti and Ni with a near 1 : 1 atomic ratio. After H2O2-oxidation treatment, EDX results showed a significantly reduced amount of Ni on as-oxidized surfaces of NiTi SMA (Fig. 6.17), which improves the biocompatibility of the material. (The cracks exhibited in Fig. 6.17a were artefacts caused by drying up of the titania layer under SEM.) The surface of H2O2-oxidized samples was mainly composed of Ti and O elements. The Ni element (only 8.2 at.%, down from 50.8 at.% in the original NiTi SMA surface) mostly existed in the ‘cracked’ area (Fig. 6.17b), indicating that the sample surface had a
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Spectrum 1
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6.17 Surface analysis of an H2O2-oxidized NiTi SMA sample: (a) an SEM image of the surface; (b) Ni element mapping of the surface by EDX.
relatively Ni-free layer. Potentiodynamic polarization tests were conducted for untreated and H2O2-oxidized NiTi SMA samples. As can be seen from potentiodynamic polarization curves, the corrosion potential (Ecorr) of the as-oxidized NiTi SMA sample (−142 mV) was significantly higher than that of the untreated sample (−473 mV), revealing that the thermodynamic driving force which causes corrosion was remarkably stronger for the
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untreated sample. In addition, a large shift of the breakdown potential (Eb) was recorded from 428 mV for the untreated sample to 1010 mV for the as-oxidized NiTi SMA sample. Furthermore, the current density (Icorr) of the as-oxidized NiTi sample at Eb was much lower than that of the untreated sample. After potentiodynamic polarization tests (which are ‘chemical stability tests’ in the current context) in SBF, deep and irregular cavities were visible on the surface of the untreated sample, whereas no significant pitting or holes was observed on the surface of the as-oxidized NiTi sample, indicating that the as-oxidized NiTi SMA sample had better chemical stability and hence biocompatibility. The results obtained from both chemical stability tests and SEM examination suggest that the titania layer fabricated by H2O2-oxidation and subsequent hot water ageing was much more stable chemically and did not break down easily in the simulated body environment. This oxide layer was homogeneous and compact, had a thickness of about 4 μm and did not contain much Ni. On its own, it should be able to act as a physical barrier to prevent (or retard) the release of Ni ions from the NiTi SMA implant. This physical barrier should be stable in the human body environment for sufficient time after implantation. The synthetic TiO2 layer (the anatase phase of TiO2 after crystallization by the hot water ageing treatment) also possessed the ability to induce in vitro formation of bone-like apatite, which forms the top layer of a two-layer composite coating on NiTi SMA substrate. By immersing asoxidized NiTi SMA in the 5SBF solution, apatite gradually formed on as-oxidized NiTi SMA and the thickness of the apatite layer grew with increasing immersion time. Figure 6.18 displays XRD patterns for NiTi SMA at different stages of composite coating formation. As can be seen from Fig. 6.18b, apart from diffraction peaks arising from the NiTi SMA substrate, a broad peak at ∼25.3 ° corresponding to poorly crystallized anatase was present for the surface of as-oxidized NiTi SMA samples. Meanwhile, the NiTi SMA peak from the substrate at ∼42 ° was reduced in intensity after the H2O2-oxidation and hot water ageing treatment owing to the formation of the titania layer on substrate surface, indicating the formation of a relatively thick TiO2 layer that reduced the NiTi SMA peak. After soaking in 5SBF for 4 h, apart from the diffraction peaks from NiTi SMA and anatase, broad peaks associated with low crystallinity bone-like apatite at 2θ = 26 ° and 32 ° were detected for the NiTi SMA sample (Fig. 6.18c). The intensity of the apatite peak at 2θ = 32 ° increased gradually as the immersion time increased to 24 h (Fig. 6.18d). With increasing immersion time, diffraction peaks arising from the NiTi SMA substrate decreased in intensity and finally disappeared from XRD patterns, indicating that the apatite layer formed on the as-oxidized NiTi SMA substrate became thicker, resulting in the suppression of X-ray diffraction of the NiTi SMA substrate.
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6.18 XRD patterns of NiTi SMA samples: (a) untreated NiTi SMA sample; (b) as-oxidized NiTi SMA sample, (c) as-oxidized NiTi SMA sample after soaking in 5SBF for 4 h; (d) as-oxidized NiTi SMA sample after soaking in 5SBF for 24 h.
Figure 6.19 exhibits the cross-sectional morphology of an apatite/TiO2 composite coating formed on NiTi SMA, which clearly reveals that the uniform and compact composite coating firmly adhered to the substrate, indicating good adhesion strength. For the composite coatings formed on NiTi SMA, there was no obvious interface between the outer apatite layer and the inner TiO2 layer, which was probably because the apatite and TiO2 layers bonded together tightly. Compared with the thickness of the apatite/ TiO2 composite coating formed on the Ti substrate (about 28.5 μm), the composite coating formed on the NiTi SMA substrate was thinner (about 17.0 μm) after the same soaking time (24 h) in 5SBF, suggesting that apatite formation on the Ti substrate was much faster than that on the NiTi SMA substrate. The adhesion strength of apatite/TiO2 composite coatings, as was determined through scratch tests, was higher for Ti substrate than that of corresponding NiTi SMA substrate. But NiTi SMA samples having the composite coating exhibited better wear properties than Ti samples also having the composite coating. The apatite/TiO2 composite coating on NiTi SMA should not only promote bone formation in vivo but also serve as a physical barrier to prevent Ni ion release from the implant to surrounding body tissues.
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C
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Date: 18 Jul 2007 Time: 14:28:15
6.19 Cross-sectional view of the apatite/TiO2 composite coating formed on NiTi SMA substrate through accelerated biomimetic deposition. (C: apatite/TiO2 composite coating.)
6.4.5 Coatings fabricated by other techniques Electrochemical deposition was investigated to form an apatite/collagen composite coating (type I coating design in Fig. 6.1) on NiTi SMA at room temperature (25 °C) in 2SBF (Sun et al., 2009a). The coating was composed of spherical apatite particles (0.5–1.5 μm in size) and fibrous collagen. Apatite clusters seemed to have nucleated and grown on collagen fibrils, forming the apatite/collagen composite coating. The Ca : P ratio of the composite coating, as determined by EDX, was about 1.38, which is slightly higher than that of OCP (with a Ca : P ratio of 1.33) and much lower than that of HA (with a Ca : P ratio of 1.67). XRD and FTIR results indicated that apatite in the composite coating was amorphous, calcium-deficient, and carbonated. This composite coating was shown to have improved the wettability of NiTi SMA. Plasma immersion ion implantation and deposition (PIIID) is a better technique than plasma immersion ion implantation (PIII) for fabricating thin, dense, tough, adherent and biocompatible (and maybe also bioactive) coatings on NiTi SMA alloys for their wider applications in the medical field. It was used to make (Ti–O)–Ti composite coatings (type III coating design in Fig. 6.1) on NiTi SMA recently (Sun et al., 2009b). The fabrication of (Ti–O)–Ti composite coatings was performed using a multi-purpose plasma immersion ion facility. The vacuum chamber of this facility was evacuated to a base pressure of 5.0 × 10−3 Pa. To remove any residual pol-
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lutant or surface oxide, sputtering of the sample surface using an argon plasma was performed before the PIIID process. Afterwards, a Ti transition layer was firstly fabricated on NiTi SMA samples, which increased the bonding strength of the Ti–O coating to substrate, using a pulsed cathodic arc plasma source. Subsequently, O2 plasma was generated in the vacuum chamber through radio-frequency (RF) glow discharge and the Ti–O layer formed on the Ti layer. Two-layered (Ti–O)–Ti composite coatings were thus fabricated on the surface of NiTi SMA. XRD patterns of PIIID-treated NiTi SMA samples did not show any peaks that could be attributed to TiO2 and only displayed Ti peaks and relatively small peaks arising from the B2 and B19′ phases of NiTi SMA substrate, indicating that the Ti transition layer consisted mainly of poorly crystallized Ti and that the Ti–O outer layer was amorphous. Compared with untreated NiTi SMA, peaks from the NiTi SMA substrate in XRD patterns of PIIID-treated samples were suppressed, suggesting the effectiveness of (Ti–O)–Ti composite coating. SEM examination showed that the surface of composite coatings was smooth and no coating delamination was observed. The cross-sectional morphology of PIIID-treated samples indicated that (Ti–O)–Ti composite coatings were uniform and compact and had a thickness of about 0.5 μm. No obvious boundary existed between the Ti layer and Ti–O layer. EDX element mapping of PIIID-treated samples revealed that Ni was depleted from the surface owing to composite coating formation on PIIID-treated samples. Pin-on-disc test results indicated that the wear resistance of PIIID-treated samples was improved owing to the protective composite coating. Therefore, the layered (Ti–O)–Ti composite coatings fabricated by PIIID effectively improved both mechanical properties and biocompatibility of NiTi SMA.
6.5
Composite coatings for tissue engineering scaffolds
As has been pointed out in section 6.3.6, single-component, monolayer coatings can be made on some tissue-engineering scaffolds for improving their performance (biological and/or mechanical).With appropriate modification(s) or enhancement of the coating techniques, scaffolds can be coated with composite coatings for mainly improving their biological performance.
6.5.1 Tissue engineering materials and scaffold manufacture Various materials have been investigated or used for making porous scaffolds for tissue engineering applications (Lanza et al., 2000; Ratner et al., 2004). For bone tissue engineering, osteoconductivity of the scaffold is one of the major considerations. It is also highly desirable that, if a composite
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scaffold is used for bone tissue engineering, the composite scaffold should be totally biodegradable. (Some of currently investigated composite scaffolds are not totally biodegradable because in any one of these composite systems, one of the constituent materials of the composite does not exhibit degradation in vitro or in vivo.) The scaffolds can be made using different techniques and in most situations for fabricating useful and usable scaffolds, the tissue engineering material and the scaffold production technology are considered together. Although in recent years opinions have surfaced that polymers used for tissue engineering scaffolds do not necessarily need to be degradable in certain applications, most research in the field has followed the generally accepted notion of tissue engineering by using biodegradable polymers. Traditionally, polyesters such as PLLA and PLGA are materials of choice for tissue engineering. However, other synthetic biodegradable polymers such as poly(propylene fumarate) (PPF) have also been investigated for making tissue engineering scaffolds. Natural, biodegradable polymers such as PHBV, chitosan and collagen are also used in tissue engineering, having advantages over synthetic polymers because their degradation products are part of the body or body fluids. But natural polymers have their own problems for tissue engineering applications. The degradation rate of these polymers (natural or synthetic), which is affected by various factors but predominantly by the average molecular weight of the polymer, controls the degradation rate of scaffold. Recognizing the importance of controlling the degradation rate and other properties of scaffolds, polymer blends have also been of interest for constructing tissue engineering scaffolds (Sultana and Wang, 2008b). Having previously been overlooked for tissue engineering applications, some bioactive and biodegradable bioceramics have now been investigated for constructing tissue engineering scaffolds (Pereira et al., 2005; Shah et al., 2005). Recent investigations have also looked into the possibility of using biodegradable metals for scaffolds (Witte et al., 2007). As non-porous bioceramics or metals are much stronger and stiffer than non-porous polymers, bioceramic or metal scaffolds possess higher strength and stiffness than polymer scaffolds at the same porosity level. However, it should be noted that the strength of a material decreases exponentially with an increasing degree of porosity and highly porous scaffolds (sometimes with the porosity level above 90%) are required for tissue engineering applications. Therefore, whether it is made of a polymer, a bioceramic or a metal, the scaffold on its own is not destined for load-bearing applications. Another attractive group of materials for tissue engineering are hydrogels. Polymeric hydrogels are cross-linked macromolecular networks formed by hydrophilic polymers swollen in water or biological fluids. They offer unique properties that can be used for tissue engineering of various human tissues (Nicodemus and Bryant, 2008). Hydrogels of both natural polymers
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and synthetic polymers have been investigated for tissue engineering applications. The composite approach has been adopted by many researchers around the world in developing new biomaterials. Increasingly, composite scaffolds for tissue engineering have been investigated and used. In the premier international biomaterials conferences, special symposia have been organized for biomedical composites (7th World Biomaterials Congress, 2004) and composite tissue engineering scaffolds (8th World Biomaterials Congress, 2008). For bone tissue engineering, as pointed out in section 6.3.6, bioceramic–polymer composite scaffolds, which take the advantage of osteoconductivity of the bioceramics, are promising structures for bone cell attachment and proliferation and tissue formation and hence are actively investigated by many research groups. Polymer-based tissue engineering scaffolds are fabricated using different techniques (Atala and Lanza, 2002). Commonly used scaffold-making methods such as porogen leaching, emulsion freezing / freeze-drying, and fibre bonding are conventional technologies in chemical engineering and other industries. These techniques produce scaffolds with, mostly, irregular pores although the pore size and sometimes also pore shape can be controlled through the careful selection of scaffold manufacture parameters. Rapid prototyping (RP) technologies (including SLS), which are widely used in the traditional manufacturing industries, have attracted great attention in recent years for constructing tissue engineering scaffolds (Leong et al., 2003). One of the appealing aspects of using RP technologies is that the pore size, pore shape and porosity of a scaffold can be designed and the scaffold of the specific architecture can then be made according to the design (and clinical requirements). This, when properly used, gives particular advantages to some tissue engineering strategies.
6.5.2 Biomimetically deposited coatings for scaffolds As has been stated in section 6.3.6, there are two routes for making osteoconductive bioceramic–polymer composite scaffolds. In the biomimetic deposition route, a bioactive apatite coating can be formed on pore surfaces of polymer scaffolds (Chen et al., 2005b, 2005c, 2007). It was shown that this apatite coating promoted osteoblastic cell (Saos-2 cells) attachment and significantly enhanced cell proliferation and differentiation (Chen et al., 2006b). Natural bone is a composite of organic collagen fibers and inorganic apatite. A composite coating consisting of natural collagen and bioactive apatite appears attractive for the surface modification of polymer scaffolds. It is possible to fabricate an apatite/collagen coating having a layered structure, as the type III (simple) coating design in Fig. 6.1, on polymer scaffolds.
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However, an apatite/collagen coating of the type I composite coating design (Fig. 6.1) is better in several aspects in terms of coating structural integrity and biological performance. Using experimental setups as shown schematically in Fig. 6.6 and 6.10, when the coating solution was changed from pure 5SBF to 5SBFC (5SBFC is the 5SBF solution containing dissolved collagen), apatite/collagen composite coatings could be built up on pore surfaces of PLLA scaffolds in the biomimetic deposition process (Chen et al., 2005a, 2006a). Figure 6.20 shows an SEM micrograph of the pore surface of a PLLA scaffold which had been coated with an apatite/collagen coating through biomimetic deposition. Collagen fibrils and submicron size (200–600 nm) apatite particles could be seen in the composite coating. The collagen fibrils were interlapped randomly with each other. XRD patterns of the composite coating revealed a characteristic apatite peak at 2θ = 32 °. The Ca : P atomic ratio of apatite in the composite coating, as was determined by EDX, was 1.17. In the FTIR spectrum for the coating formed in 5SBF, strong absorptions at 1031 and 563 cm−1 owing to the ν3 and ν4 vibration of PO43−, respectively, were observed. The absorption band around 1645 cm−1 arising from the ν3 absorption of CO32− in the coating was present. (It is a common knowledge in the bioceramics field that apatite or an apatite coating synthesized in the atmospheric condition is carbonated apatite.) The spectrum also exhibited the characteristic broad band for OH− stretching at 3570 cm−1. In the FTIR spectrum for the coating formed in 5SBFC, besides the absorption peaks arising from PO43−, CO32−, and OH− functional groups, amide
EHT = 20.00 kV 2 μm
Mag = 10.00 K X WD = 9 mm Photo No. = 2242 Detector = SE1
6.20 Apatite/collagen composite coating formed on a PLLA scaffold through biomimetic deposition within 24 h.
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peaks at 1655 cm−1 and 1550 cm−1 were also present. These results showed that the co-precipitation of apatite and collagen in the biomimetic deposition process using the 5SBFC solution allowed the formation of a submicrometre apatite/collagen composite coating on PLLA scaffolds. Cell culture studies indicated much enhanced biological properties of PLLA scaffolds with the apatite/collagen composite coating (Chen et al., 2006b). Figure 6.21 shows Saos-2 cell morphology on an apatite/collagen composite coated PLLA scaffold after 2-day cell culture. Cells spread on the scaffold and exhibited filopodia. The MTT assay was used as a measure of relative cell viability. After Saos-2 cells were cultured in scaffolds for 8 days, the cell viability was evaluated using MTT assay. (Experiments were run in triplicate per condition. All results were expressed as mean ± standard deviation. The single factor analysis of variance technique was used to assess the statistical significance of results. Differences with p < 0.05 were considered to be significant.) Figure 6.22 shows the absorbance of formazan produced by Saos-2 cells on PLLA scaffolds with and without a coating after 8-day culture. Compared with the control (pure PLLA scaffold), higher absorbance was obtained when the scaffolds were coated with apatite or apatite/collagen with p < 0.05 or p < 0.01. Compared with the apatite coating, the viability of Saos-2 cells cultured on apatite/collagen composite coating was significantly higher (p < 0.05). Alkaline phosphatase activity (ALP) was measured after 8-day cell culture to assess the differentiated osteogenic activity of the cell constructs. The ALP activity of Saos-2 cells cultured on PLLA scaffolds with an apatite/collagen composite coating was significantly higher than that on the PLLA scaffold control (p < 0.05)
10 μm
20 μm
6.21 Saos-2 cells after 2-day culture on PLLA scaffolds coated with an apatite/collagen composite coating.
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Cell viability
4
a*
b*
3
2
1
0
PLLA scaffold control Apatite coating
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6.22 MTT assay results: formosan absorbance expressed as a measure of cell viability of Saos-2 cells seeded in PLLA scaffolds with and without coating. (*p < 0.05, **p < 0.01. (a) PLLA scaffold with apatite coating compared with pure PLLA scaffold; (b) PLLA scaffold with apatite/collagen composite coating compared with PLLA scaffold with apatite coating; (c) PLLA scaffold with apatite/collagen composite coating compared with pure PLLA scaffold).
and also than that on PLLA scaffolds with apatite coating (p < 0.05). These results indicate the importance of surface composition of tissue engineering scaffolds in modulating cell adhesion, proliferation and differentiation of osteogenic cells. The apatite/collagen composite coating was similar to natural bone. It offers a biomimetic microenvironment for osteoblastic cell attachment. PLLA scaffolds with an apatite/collagen composite coating are more suitable for osteoblast-like cell adhesion and proliferation than pure PLLA scaffolds or PLLA scaffolds with an apatite coating. A recent study (Chen et al., 2008b) showed that spreading of human mesenchymal stem cells (hMSC) on the apatite/collagen composite coating was better than on apatite coating or PLLA surface. Integrin expression and focal adhesion on the apatite/collagen composite coating was higher than on the apatite coating or PLLA surface. These results further indicate the usefulness and role of apatite/collagen composite coating for polymer scaffolds, using either osteoblastic cell or hMSC as the cell for bone tissue engineering. For the surface modification of bone tissue engineering scaffolds, signal recognition ligands and sequence may be incorporated into the coating to mediate the response of anchorage-dependent cells such as osteoblasts and osteoclasts. Matrix proteins such as laminin, fibronectin, and bone morphogenetic proteins (BMPs) play important roles in enhancing osteoblastic functions. Using the accelerated biomimetic deposition process, these proteins can be contained in the 5SBF and hence protein-containing apatite
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coatings or apatite/collagen coatings can be made on bone tissue engineering scaffolds.
6.5.3 Surface modification of scaffolds using other techniques The surface modification of tissue engineering scaffolds can be achieved by the incorporation of adhesive peptides such as Arg–Gly–Asp (RGD) and covalent immobilization of bioactive proteins such as growth factors. In bone tissue engineering, effective and sustained delivery of growth factors such as bFGF, VEGF and BMP-2 at the target site are of significant importance. Heparin, a sulfated polysaccharide, is known to have the binding affinity with a number of growth factors and is thus capable of blocking the degradation of protein and prolonging the release time. Based on the previous work of physical entrapment of hydrophilic gelatin on the struts of Ca–P/PHBV scaffolds made by SLS (Duan and Wang, 2008), heparin was immobilized via covalent conjugation onto gelatin-modified Ca–P/PHBV scaffolds (Duan and Wang, 2009). This composite coating structure follows the type III (simple) coating design shown in Fig. 6.1: struts of the composite scaffolds being the substrate, gelatin being the C1 layer of the composite coating and serving as the ‘bond coat’, and heparin being the C2 layer of the composite coating. In the coating process, the gelatin-modified Ca–P/ PHBV nanocomposite scaffolds were pre-wetted through a treatment and then soaked in an activated heparin solution for 4 h, washed with phosphatebuffered saline (PBS) and dried overnight at room temperature. Heparin was thus immobilized on gelatin-modified scaffolds by conjugating carboxylic groups in heparin to amine groups in gelatin via EDC/NHS chemistry. The content of conjugated heparin was analysed using the toluidine blue colorimetric method. The amount of immobilized heparin, as was determined by the toluidine blue assay, was 41.8 ± 0.4 μg/scaffold (using two-layer Ca-P/PHBV scaffolds for experiments, with each scaffold weighing ∼70 mg). To study the stability of immobilized heparin, scaffolds with surface immobilized heparin were immersed in PBS for up to 14 days. It was shown that the amount of heparin immobilized on scaffolds decreased gradually from 41.8 ± 0.4 to 22.3 ± 0.7 μg/scaffold after 14 day immersion in PBS. There are other techniques that are useful and can be used for the surface modification of tissue engineering scaffolds (Agrawal et al., 2006; Ma et al., 2007).
6.6
Concluding remarks
The first generation of modern biomaterials encompasses biocompatible engineering materials which have extended their applications into the
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medical field. Owing to their distinctive characteristics, useful properties and proven record, they have remained in service as implant materials and will continue to be used in the medical field in the foreseeable future. ‘Designer biomaterials’ such as bioactive bioceramic–polymer composites are specially developed materials according to clinical requirements and are used to solve various medical problems. Tissue engineering holds the promises of tackling problems of tissue loss and organ failure by developing biological substitutes that can fit into and enhance the body’s own repair process in the injured state (any implantation in the body is considered to be ‘intentional’ injury by the body) or facilitate the restoration of tissue/ organ functions. Forming suitable tissue engineering scaffolds thus plays a pivotal role in successful tissue engineering strategies. The surface modification of existing, accepted metallic biomaterials improves their biological performance and is an important area in biomaterials R&D. The surface modification of tissue engineering scaffolds can significantly enhance their properties (biological and/or mechanical) and therefore broaden their applications. Using composite coatings for the surface modification of metal implants and tissue engineering scaffolds has distinctive advantages and will be increasingly investigated and adopted in the future. As can be seen from our work as well as research performed by other research groups on the surface modification of metallic biomaterials, bioactive bioceramic composite coatings can fulfil; the functions of: 1. promoting and enhancing bone tissue formation on metal implants; and 2. protecting metal implants from corrosion in the human body environment and preventing (or minimizing) the release of cytotoxic ions from metal implants to surrounding tissues, while maintaining the strength and stiffness, which are required for loadbearing applications, of metals used for the implants. For tissue engineering scaffolds, osteoconductive composite coatings can enhance cell adhesion, proliferation, and differentiation and subsequently, bone tissue formation. There is already a sufficiently large body of knowledge about ‘structural’ coatings (including composite coatings presented and discussed in this chapter and also those investigated by others), ‘structural’ coatings here meaning mainly bioceramic coatings. One of the future research emphases in surface modification in the medical field will be in the direction of achieving composite coatings with more biofunctions or being more biofunctional. For bioinert metal implants, there are quite a few surface modification techniques that can be used to fabricate bioactive single-component, monolayer bioceramic coatings or composite coatings on their surfaces in order to make them osteoconductive. Composite coatings need to be carefully designed, taking into account the coating materials and coating fabrication
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technology. The selection of a particular surface modification technique for making the coatings depends on a number of factors. Low-temperature processes for coating fabrication have obvious advantages over processes that involve the use of high temperature and/or high (or low) pressure. Biomimetic deposition appears to be increasingly used in biomaterials development. The biomimetic deposition of apatite coating or composite coating on polymeric tissue engineering scaffolds provides a simple means of rendering the scaffolds osteoconductive. The biomimetic processes can also be used to add or enhance other functions for tissue engineering scaffolds. There is no doubt that new and practical technologies, especially techniques in the nanotechnology field, will emerge and these can be utilized for the surface modification of biomaterials and tissue engineering scaffolds.
6.7
Acknowledgements
I thank Professor Luigi Ambrosio, the Editor of this book, for inviting me to contribute a chapter on composite coatings. Professor Ambrosio’s invitation made me pause and reflect on our research on the surface modification of various biomaterials that has been conducted over the past ten years in Singapore and Hong Kong and also made me ponder on what we can do and will be doing in this area in the future. I would like to thank the research students, research associates and post-doctoral researchers for conducting the research work with me on surface modification of metallic biomaterials and tissue engineering scaffolds in Nanyang Technological University (NTU), Singapore, The Hong Kong Polytechnic University (PolyU) and The University of Hong Kong (HKU), Hong Kong. I wish to thank my collaborators in these universities and also other universities for their useful discussions and support. Assistance provided by technical staff in various laboratories of the universities is much appreciated. Research funding from funding agencies (MOE in Singapore, and RGC in Hong Kong) and universities (NTU, PolyU and HKU) is gratefully acknowledged. I thank my research students and staff, past and present, for providing me with the figures that I have used for this chapter. The credit goes to: H.-W. Tong for Fig. 6.2; X.-Y. Yang for Figs. 6.3, 6.11 and 6.12; J.-M. Wu for Fig. 6.7; T. Sun for Figs. 6.8, 6.17, 6.18 and 6.19; Y. Chen for Figs. 6.9, 6.20, 6.21 and 6.22; R. R. Kumar for Figs. 6.13 and 6.14; and C.-X. Wang for Figs. 6.15 and 6.16.
6.8
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7 Composite materials for spinal implants A. G L O R I A, R. D E S A N T I S, L. A M B R O S I O, National Research Council, Italy; and F. C AU S A, University of ‘Magna Graecia’, Italy
Abstract: In this chapter, the state of the art of spinal implants such as interbody spacers and intervertebral disc (IVD) prostheses made of conventional materials, is described and the role of composite biomaterials for spine applications highlighted. In particular, the possibility of designing multifunctional devices with tailored mechanical properties is emphasized. A biomimetic approach to developing an IVD prosthesis with appropriate biological, transport and mechanical properties, is also presented. Finally, in the field of composite biomaterials for spine applications, future challenges and strategies are discussed. Key words: polymer composites, spinal implants.
7.1
Introduction
The spine is a co-operative system of elements and its unique function is to provide trunk flexibility and protection of the spinal cord and nerve roots that pass through the spinal canal and foramen. During daily activities (sitting, standing, jogging, walking, lifting a weight . . .) the loads experienced by intervertebral discs and other spinal elements can be several times higher than the body weight, and, more importantly, repetitive and fluctuating. However, spinal diseases are a serious medical problem which affects many people in the world. Almost all of the spinal components (e.g. ligaments, intervertebral discs, vertebral bodies, facets and laminae) may be dissected at surgery, depending upon the pathology or deformity (Martz et al., 1997). In order to restore spinal stability and function, medical practice today uses a large number of man-made devices designed to restore, replace or improve the function of degenerated tissues or organs, and to correct abnormalities. Intervertebral discs (IVDs) give flexibility to the spine and enable the body to twist and bend into a wide range of postures. Each IVD consists of a soft centre, called the ‘nucleus pulposus’ surrounded by an outer wall, the ‘annulus fibrosus’. The annulus presents a layered structure where each layer is reinforced by a regular pattern of collagen fibres. IVD degeneration usually involves dehydration of the nucleus often accompanied by small 178
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tears in the annulus. This is a common reason for the back pain which affects many people in the world. The pain may become a chronic disabling condition (Hukins, 2005; Tsantrizos et al., 2005). Some chronic back-pain patients may benefit from spinal fusion, a surgical approach in which adjacent vertebrae are fused together. This is usually realized through the use of a man-made device or a bone graft that promotes bone ingrowth and eventual fusion of the vertebral bodies. However, this approach provides reduction of spine flexibility and biomechanics alteration. Accordingly, there is an increasing attention on the restoration of flexibility through the replacement of the damaged IVD with an artificial one. Combinations of several mechanical, chemical and physical properties are usually required to design multifunctional materials for spine applications. Metals, polymers, ceramics and composites are used to design spinal implants, since these materials can span a range of properties to meet the desired needs. However, in contrast to conventional materials (metals, ceramics and polymers), fibre-reinforced polymers have gained an important role in the development of prosthetic devices with improved and tailored mechanical properties. In this context, this chapter provides an overview of the state of the art in materials for designing spinal implants. Fibre-reinforced composite materials will be discussed taking into account their ability to match the mechanical properties of natural interacting structures, and to reproduce the structure and behaviour of the tissue that they are meant to replace, as in the case of IVD prostheses.
7.2
Structure and function of the spine
The spine plays two distinct and apparently conflicting roles. First, it must provide mobility to the trunk and transfer loads from the head and trunk to the pelvis, representing a strong, yet mobile axis onto which the appendicular skeleton is applied. Second, it must protect the spinal cord and the roots of delicate nerves connecting the brain to the periphery. This proper blending of mobility, stability and structural integrity is crucial to fulfil these goals simultaneously and the dual function is due to a linked structure consisting of 33 vertebrae superimposed on one another (Rothman and Simeone, 1992). The spine can be divided into four distinct regions: cervical, thoracic, lumbar and sacral. In particular, the human spinal column is a complex structure made up of 24 individual vertebrae plus the sacrum; the sacral coccygeal region is formed by nine fused vertebrae, and articulates the left and right ilia at the sacroiliac joints to form the pelvis. Motion is allowed into three planes: flexion–extension, axial rotation and lateral bending.
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Adjacent vertebrae are linked by three articulations defined as the ‘triple joint complex’, which consists of the cube-shaped vertebral bodies sandwiching an intervertebral disc (IVD) and matching facet joint posteriorly. The main function of the disco-vertebral joint is to transmit compressive loads while providing flexibility. Like the hip and knee, the facet joints are characterized by a smooth cartilage surface, lubricating joint fluid and a covering capsule. They allow small degrees of flexion and extension, limit rotation, and protect the IVD from translational shear stresses (Bao et al., 1996). The principal biomechanical function of the vertebral body is to support the compressive loads of the spine resulting from body weight and muscle forces. The vertebrae are separated by IVDs and are united by artificial capsules and ligaments. The IVD is a composite structure made up of a nucleus pulposus surrounded by a lamellar fibrous structure, the annulus fibrosus. The annulus fibrosus is a multilayered structure with each layer characterized by well-organized collagen fibres embedded in a proteoglycan– water gel, running in opposite directions in adjacent layers (Markolf and Morris, 1984; Cassidy et al., 1989). The nucleus is a semifluid mass enclosed within the annulus and the two endplates. It is mainly composed of water and proteoglycans, which form a gel-like matrix. The type and orientation of collagen fibres in the IVD have an important influence on how the load is distributed. In the disc there is a gradation of collagen type and orientation from annulus to nucleus. In the concentric lamellae the orientation of annulus fibres vary from 62° at the periphery to 45° in the vicinity of the nucleus, with respect to the spinal axis, thus imparting a structurally graded architecture to the IVD (Cassidy et al., 1989). The IVD is covered on the upper and lower surfaces by a thin layer of cartilaginous endplates, which contain micropores that allow the exchange of water, nutrients and products of metabolism (Bao et al., 1996). The main role of IVD is to act as a shock absorber for the spine. Many spine-related disorders have been identified over the years and often one disorder has a cascading effect on the other. Several reasons including IVD ageing, birth deformities, metastasis, and mechanical loads caused by sports and work lead to spine disorders (Ramakrishna et al., 2001). The two most common spinal surgery approaches, discectomy and fusion, are far from the ideal treatment for disc degenerative disease (Bao et al., 2000). Although discectomy provides a reasonably good short-term result in relieving low back pain, it causes reduction of the disc height in almost all patients and also increases the instability of the treated disc. Accordingly, the consequences of these changes in anatomy and structural stability are twofold. On the other hand, spinal fusion means surgical immobilization of joint between two vertebrae and different techniques are usually employed.
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The fusion approach mainly consists of using particular devices such as baskets, cages and threaded inserts, made of metals or bioceramics. These devices are designed and manufactured so that tissues can grow into them, thus ensuring rigid anchoring of prostheses to the bone. However, by eliminating motion, fusion also significantly alters the normal biomechanics of the spine. Even though single-level IVD fusion may not impair the patient’s normal function and activity as seriously as knee or hip fusion, it increases stress and strain on the discs at adjacent levels (Bao et al., 2000). Therefore, the ideal solution would seem to be an artificial IVD.
7.3
Materials and design of spinal implants: the state of the art
It is well known that conventional materials (i.e. metals, polymers and ceramics) have been considered to make spinal implants over the past years, since these materials have a range of properties sufficient to meet the desired needs. The materials used and the various designs of spinal devices clearly affect their performance.
7.3.1 Interbody spacers Several interbody fusion techniques have been used for many years and, as previously described, they are based on the concept of removing all or most of the IVD and stabilizing the operated segment. Many types of cages have been developed, including ring-like, rectangular or tapered, and hollow threaded cylindrical cages (Blumenthal and Ohnmeiss, 2003; McAfee, 1999; Weiner and Fraser, 1998). Non-threaded cages do not compromise the endplate integrity, but, in contrast to threaded systems, the geometry may be inadequate to obtain a good surface match to the endplate (Steffen et al., 2000; Gloria et al., 2008). Many types of cage implants have been inserted by different approaches. Bagby developed smooth stainless-steel cylinders with multiple holes drilled throughout the walls, described as ‘baskets’. These devices were implanted into intervertebral holes in the necks of horses leading to chronic cervical spine disorder. Moreover, drilling and hammering before insertion of the device caused problems and vertebral fractures. Therefore, some workers concluded that constructs with smooth walls and insertion with use of hammering were not appropriate for a human application (Ray, 1997; Gloria et al., 2008). The BAKTM interbody fusion system is an example of cage structure, which is a hollow, threaded and porous titanium–aluminium–vanadium (Ti–6Al–4V) shell into which bone grafts can be packed. In order to better contain the bone grafts the ends of the lumbar device are capped with an
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ultrahigh molecular weight polyethylene (UHMWPE). This cylindrical device can be implanted either centrally or bilaterally between the vertebral bodies (Martz et al., 1997). Waisbrod (1988) developed a cobalt–chrome metal sponge with a porous structure similar to that of cancellous bone, but as with all metallic porous structures there is a substantial increase in the surface area that can lead to problems with increased release of metal ions. A study performed by Merritt and Brown (1985) suggested that patients with cobalt–chrome alloys and stainless-steel prostheses may experience increased pain, loosening of the implant and bone resorption. Moreover, inflammatory responses and infections cause lower pH levels at the implant site, which can increase fretting corrosion. Consequently, titanium mesh blocks were developed as a substitute for autogenous bone grafting. They are quite compressible and characterized by a mechanical compliance which is close to that of cancellous bone. For this reason they should reduce stress shielding effects, but because of their claimed flexibility there may be problems of collapse (Leong et al., 1994; Martz et al., 1997; Gloria et al., 2008). McAffee et al. (1986) proposed a poly(methylmethacrylate) (PMMA) bone cement for fixation of the spine, eliminating the need for external bracing. However, many complications at the bone–cement interface (i.e. failure) resulted in loss of fixation. Yamamuro et al. (1990) developed a ceramic prosthesis for replacing a lumbar vertebra in a sheep. This glassceramic implant composed of apatite and wollastonite has high biocompatibility and apatite–wollastonite glass–ceramic spacers, in place of the bone graft, have been used in conjunction with laminoplasty in the cervical zone with some success. A variety of clinical procedures have been proposed in order to achieve successful vertebral fusion. Even though it is very difficult to determine if a segment has achieved a bony union, many complications have been reported with cage use, similar to those documented for traditional interbody fusion using bone grafts (Blumenthal and Ohnmeiss, 2003). However, initial biomechanical studies have suggested that threaded lumbar intervertebral cages are stable enough to be used as stand-alone devices (Brodke et al., 1997; Zdeblick and Phillips, 2003).
7.3.2 Intervertebral disc prostheses Fusion alters the biomechanics of adjacent vertebral levels, which may be the reason for disc degeneration, symptomatic hypertrophic facet arthropathy, spinal stenosis and osteophyte formation, which have all occurred at levels adjacent to a fusion site. There are a number of other drawbacks related to fusion, including loss of spinal mobility, graft collapse resulting in suboptimal sagittal balance,
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autograft harvest site pain, and alteration of muscular synergy. Even in the cervical spine, where the results seem to be better than in the lumbar region, many studies have reported that the rate of disc degeneration in levels adjacent to a cervical fusion is approximately 3% per year. Other works have demonstrated that arthrodesis increases local motion in adjacent segments as well as their contribution to global cervical motion; this may play a crucial role in the acceleration of adjacent-level disc degeneration (Traynelis, 2002). A great number of IVD prostheses have been developed, ranging from a ball bearing prosthesis to sandwiches of porous coated metal endplates and elastomeric cushions (Martz et al., 1997). Ideally, artificial discs should have similar functions and properties as the natural structure (Ramakrishna et al., 2001). Consequently, material properties to be considered in designing prostheses are biocompatibility, endurance and resistance to long-term compressive creep. Thus, metals, polymers, ceramics, and combinations of metallic and non-metallic materials have been considered over the past years. The concept of a disc prosthesis was first proposed in 1956 in a French patent by van Steenbrugghe (1956), but it was not until 17 years later that Urbaniak et al. (1973) reported the first prototype of an IVD prosthesis, which was implanted into a chimpanzees. Since then, many other IVD prosthesis concepts have been proposed (Bao and Yuan, 2000). One of the first attempts to carry out disc arthroplasty was undertaken by Nachemson 40 years ago. Fernstrom proposed to reconstruct IVDs by implanting stainless-steel balls in the disc space (Traynelis, 2002). Component materials can be divided into metals, non-metals, and metals in combination with non-metals. The main advantage of using metals alone for realizing an IVD prosthesis is the high fatigue strength compared with a non-metal design. Because patients with back pain are on average approximately 40 years old, the device should have at least 40 years of fatigue life (Bao and Yuan, 2000). Considering that metals are usually much stiffer than the natural IVD they are used to replace, they must be specially designed to reduce stiffness and provide appropriate flexibility. The most widely tested all-metal disc prosthesis consists of two Ti–6Al– 4V springs pocketed between two forged or hot isostatically pressed cobalt– chromium–molybdenum alloy plates with a posterior hinge allowing flexion and extension (Hedman et al., 1991; Hellier et al., 1992). In designing this implant, great efforts were made to meet various design concepts and criteria, focusing the attention on biocompatibility, dynamics, kinematics and long-term resistance. The springs were chosen to reproduce the stiffness of a natural IVD. Vertically projecting lugs are positioned at the front and side of the plate members, through which screws can be placed for fixation. Fatigue tests were performed on the individual components, which
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underwent one hundred million cycles without failure. In a sheep model, it was shown that fibrous tissue does not grow between the hinges or around the coils in the short term. However, such ingrowth would be expected to interfere with the implant mechanics and, because of the potential risk associated with its use, it has yet to be applied in humans. The principle advantage of using a metal IVD prosthesis is its high fatigue strength, whereas the main benefit of a non-metal design is its mechanical similarity to the natural disc. Considering, its lower elastic modulus, a nonmetal design may more closely replicate IVD kinematics. The concept of an IVD prosthesis with a non-metal design was introduced by van Steenbrugghe (1956). The implant consists of a multicomponent device encompassing intermediate cushion inlayers with a plastic body of varying shapes, but no experimental tests performed on this device have been reported. The most widely tested device in the non-metal category is that designed by Lee et al. (1990). This design consists of a soft central elastomeric core, which reproduces the function of the nucleus, reinforcing fibre sheets with specific alternating fibre orientation in six to fifteen laminae embedded in a second elastomer, which emulates the function of the annulus, and two stiff plates. This device is able to reproduce both the compressive modulus and the compressive–torsional stiffness of the natural disc, but the lack of adequate implant and vertebra fixation is thought to be the great obstacle to its clinical use (Bao and Yuan, 2000). In order to take advantage of both metal and non-metal materials, overcoming the drawbacks related to the use of either of them alone, many researchers have combined both types of materials in their designs, thus realizing a metal–polymer–metal sandwich design. A metal plate is used to anchor the device to the vertebral bodies through spikes, tabs with screws, or porous coating for ingrowth, whereas the polymer provides the typical flexibility of an IVD. Although many IVD prostheses have been proposed over the years, only three have been tested extensively and used clinically. In particular, the LINK SB Charité prosthesis has provided the largest and longest clinical trial of all existing artificial discs. This device was developed by Buttner-Janz and colleagues (1988), and it has been characterized by several designs and manufacturing modifications from the first to the third and current generation. The LINK SB Charité design features two metal endplates with spikes or teeth that allows them to be anchored without cement to the vertebrae, and an interposed polymer core. In particular, the Link SB Charité III, now in its third generation, is the most widely implanted total-disc replacement (Traynelis, 2002). Overall, more than 2000 devices have been implanted worldwide. The Charité III consists of a UHMWPE spacer, which is available in different sizes, surrounded by a radiopaque ring for radiographic localization. This core spacer interfaces with two separate endplates con-
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structed using cast cobalt–chromium–molybdenum alloy. The endplates are also coated with titanium and hydroxyapatite to promote bone bonding. The prosthesis is available in different sizes and its unique design allows it to provide unconstrained kinematic motion across the implanted spinal segment device. Since 1984, over 2000 prostheses of all three generations have been implanted in Europe and the most detailed clinical data for this device were reported by Griffith et al. (1994). The ProDisc was developed by Thierry Marnay. This prosthesis consists of two cobalt–chromium–molybdenum alloy endplates with vertical wings. The metal endplates are coated with pure titanium to improve osteointegration. A monoconvex polyethylene core is inserted in the caudal endplate. Because of the monoconvex configuration of the polyethylene core, this prosthesis can be inserted with considerably less segmental distraction than that required for the Charité III device. Once the polyethylene core is firmly anchored to the caudal endplate, there are only two movable parts, which results in a kinematic behaviour that is best described as semi-constrained. The first Acroflex prosthesis consisted of a hexene-based polyolefin rubber core vulcanized to two titanium endplates; in contrast to the Charité III and ProDisc prostheses, transmission of motion is possible if there is good osteointegration of the endplates. The second generation Acroflex-100 consists of an HP-100 silicone elastomer core bonded to two titanium endplates.
7.4
Composite materials: basic concepts
Clinical experience clearly demonstrates that not all the materials commonly used in the engineering applications are suitable for biomedical devices. The various materials used in biomedical applications may be divided into: a) metals, b) ceramics, c) polymers and d) composites that are made from various combinations of a), b) and c). Each type of material shows peculiar aspects that are particularly suitable for specific applications. Metals are characterized by high strength, ductility and resistance to wear, but usually they are too stiff compared with natural tissues, have higher density, and are subject to corrosion and release of metal ions. Polymers are too flexible and too weak to satisfy the mechanical requirements for specific applications, such as implants in orthopaedic surgery. Ceramics are highly biocompatible and corrosion resistant and have high compression strength, but they also show brittleness, low fracture strength and lack of resilience. Accordingly, polymer-based composite materials provide an alternative choice to overcome many shortcomings of the above mentioned conventional materials (Ramakrishna et al., 2001). In order to evidence the advantages in the use of polymer composite biomaterials, it is possible to consider
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the mechanical compatibility of conventional materials with host tissues. Based on the structural or mechanical compatibility with tissues, metals or ceramics are considered for hard tissue applications, whilst polymers are used for soft tissue applications. However, the elastic moduli of metals and ceramics are at least 10–20 times higher than those of hard tissues (Ramakrishna et al., 2001). For instance, the mismatch of stiffness between the bone and metallic or ceramic implants negatively affects the load transfer mechanism at the bone/implant interface. In the load sharing between the bone and implant, the amount of stress carried by each of them is directly related to their stiffness. Consequently, bone is insufficiently loaded compared with the implant. This phenomenon is called ‘stress shielding’ or ‘stress protection’. Thus, the stress shielding affects the bone remodelling and healing process leading to increased bone porosity. Stress shielding can be reduced by matching the stiffness of the implant with that of the host tissues, producing desired tissue remodelling. In this respect, the use of polymers is interesting, but their low elastic modulus associated with low strength strongly limits their biomedical applications (Ramakrishna et al., 2001). Many advantages have emerged in using composite materials in the preparation of biomedical devices, since they may exhibit low elastic modulus and high strength at the same time. In particular, fibre-reinforced polymeric composites have attracted much attention as a new class of materials for manufacturing prostheses with improved and tailored mechanical properties (Ambrosio et al., 1996, 1998; De Santis et al. 2004; Gershon et al., 1990; Gloria et al., 2007). By controlling the volume fractions and the local and global arrangement of the reinforcement phase, the properties of an implant can be suitably varied and tailored in order to match those of the host tissues (De Santis et al., 2000; Gershon et al., 1990; Gloria et al., 2008; Mallick, 1988). All of this suggests that polymer-based composite materials offer a greater potential in terms of mechanical properties than homogeneous monolithic materials. Moreover, human tissues are composite materials with anisotropic properties depending on the roles, amount and structural arrangement of various components (e.g. collagen, elastin, hydroxyapatite) of the tissues. These similarities have led to the development of polymer composite biomaterials (Ramakrishna et al., 2001). When compared to metals and ceramics, polymer-based composite biomaterials clearly present great advantages such as the absence of the corrosion and fatigue failure of metal alloys, the release of metal ions, and the low fracture toughness typical of ceramic materials (Ramakrishna et al., 2001). Additionally, metallic implants complicate post-operative assessment with x-rays, computed tomography (CT) and magnetic resonance imaging (MRI) through reflection and artefacts. On the other hand, polymer com-
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posite biomaterials are fully compatible with the modern diagnostic methods such as CT and MRI (Ramakrishna et al., 2001). As previously described, the mechanical properties of an implant play an important role as they affect tissues healing and remodelling, and the importance of designing devices that mimic replaced tissue is evident (Gershon et al., 1990). Mechanical behaviour clearly varies with the type of tissue, and anisotropy represents an important characteristic common to most biological tissues. Another important feature of most natural soft tissues such as tendons, ligaments and IVDs is non-Hookean behaviour; hence, their stress–strain curve is non-linear displaying a marked toe or J-shaped region (De Santis et al., 2004; Gershon et al., 1990; Netti et al., 1996). Through the use of composite laminate design a wide selection range for mechanical properties, with a given material combination determined according to the biomedical requirements, may be provided (Gershon et al., 1990). With regard to the manufacturing process, filament winding represents one of the most interesting technologies which are commonly used to realize multifunctional biomedical devices with tailored mechanical properties. Filament winding is a technology used to manufacture axially symmetric fibre-reinforced structures by helically winding fibres impregnated with resin or reactive solution onto a suitable mandrel; the combination of the translation of the transverse carriage and the rotation of the mandrel, as well as the diameter of the systems that increases during winding, allow the winding angle to be varied thus modulating the mechanical properties (Iannace et al., 1995).
7.5
Polymer-based composite materials for spinal implants
Differently from vertebral bodies, metal devices are homogeneous isotropic materials and their design focuses only on the geometry (van Rietbergen et al., 1993). Metal interbody fusion devices are typically very strong and stiff, but they are characterized by an excessive stiffness which leads to stress shielding and implant subsidence. As consequence of the stress shielding, the bone remodelling is driven toward the bone resorption (Gloria et al., 2008; Martz et al., 1997). The solution involves composite interbody fusion devices, which are material-structure designs that can offer a wide range of possibilities in implant design. Analogously, current IVD prostheses do not reproduce the unique mechanical and complex viscoelastic behaviour of the natural structure. Thus, an approach which involves composite materials science and technology could be considered in order to design alternative IVD prostheses with suitable biological, mechanical and transport properties.
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7.5.1 Composite interbody fusion devices In recent years, there has been great interest in fibre-reinforced polymers materials for orthopaedic prostheses designed in order to avoid the bone loss observed around metallic prostheses owing to stress shielding effects. These composite materials allow a powerful tailoring of the prostheses stiffness, thus improving the stress transfer between the prostheses and bone, therefore reducing the stress shielding and stress concentration effects (De Santis et al., 2000; van Rietbergen et al., 1993). Thermoset resins reinforced with carbon fibres were the first choice for composite prostheses, but the toxic effect of unreacted monomers has suggested the use of thermoplastic polymers as matrix for implants (De Santis et al., 2000; Gloria et al., 2008; Merolli et al., 1999; Peluso et al., 1991a, 1991b, 1994). In this context, polymers under investigation as matrix for composite implants mainly include poly(ether ether ketone) (PEEK) and poly(ether imide) (PEI) (De Santis et al., 2000; Gloria et al., 2008; Merolli et al., 1999). Generally, poly(aryletherketones) (PAEKs) have been increasingly employed as biomaterials for orthopedic, trauma, and spinal implants. PAEK is a relatively new family of high temperature thermoplastic polymers, consisting of an aromatic backbone molecular chain, interconnected by ketone and ether functional groups. Two PAEK polymers used in orthopedic and spinal applications are PEEK and poly(aryl ether ketone ether ketone ketone) (PEKEKK). Owing to their chemical structure, polyaromatic ketones show stability at high temperatures (exceeding 300 °C), resistance to chemical and radiation damage, compatibility with many reinforcing agents (e.g. glass and carbon fibres), and a strength greater than that of many metals. Although polyaromatic polymers can exhibit an elastic modulus of about 3–4 GPa, the modulus can be tailored to closely match cortical bone (18 GPa) or titanium alloy (110 GPa) by reinforcing them with carbon fibres (Kurtz and Devine, 2007). By the late 1990s, PEEK was considered as the leading high-performance thermoplastic candidate for replacing metal implant components, especially in the orthopaedic field and, in April 1998, PEEK was offered commercially as a biomaterial for implants (Invibio, Ltd., Thornton-Cleveleys, UK). AcroMed (Cleveland, OH, now DePuy Spine, Raynham, MA) suggested PAEKs for cages realization. Initially, both chopped and continuous carbon fibre composites were investigated as cages (Kurtz and Devine, 2007). Brantigan et al. (1991) and Ciappetta et al. (1997) developed composite cages for lumbar fusion, consisting of PEEK and poly(sulphone) (PS) reinforced with carbon fibres, trying to match the elastic modulus of the bone. Although many works describe the use of PEEK reinforced with carbon fibres for spine implants, several studies also involve the use of neat PEEK for both
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cervical and lumbar spinal cages. Since the use of neat PEEK for spine is a relatively recent development, the published works are generally limited to in vitro biomechanical studies or short-term outcomes in animal studies or human clinical trials. Recent studies with PEEK cages have also suggested ways to improve or accelerate fusion performance by combining the devices with the use of hydroxyapatite. However, among the polymers used as matrix for composite implants, the thermoplastic amorphous PEI has also been selected because of its high mechanical properties, marginal water sorption, a high service temperature. PEI reinforced with carbon and glass fibres has already been investigated through in vivo (Merolli et al., 1999) and in vitro tests (De Santis et al., 2000). It has been shown that these PEIbased composites can be suitable for clinical applications. Gloria et al. (2008) designed and realized a novel threaded PEI reinforced with continuous carbon fibres interbody fusion device, which can provide high strength and its mechanical properties can be tailored to match the bone mechanical behaviour. This PEI fibre-reinforced cage was manufactured using filament winding and compression moulding technologies. Then it was coated with hydroxyapatite (Fig. 7.1). It is well known that bone is a natural composite which mainly consists of particles of hydroxyapatite in collagen and hence a coating of hydroxyapatite as a bioactive phase should promote a bony union, providing a long term chemical bonding between the composite device and vertebral bone (Bonfield et al., 1998).
7.1 Image of PEI fibre-reinforced cage realized using filament winding and compression moulding technologies (Gloria et al., 2008).
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As for the evaluation of mechanical properties of this PEI fibrereinforced cage, compression tests were performed according to the practice ASTM F 2077, which is intended to provide a basis for the mechanical comparison among past, present and future non-biological intervertebral body fusion device assemblies. Results from compression tests have evidenced a suitable mechanical behaviour. The intradiscal height shall be generally determined from vertebral body and disc morphometric data at the considered level of application along the spinal column (cervical, thoracic and lumbar) and many heights are suggested by ASTM F 2077, spanning from 4 to 10 mm. Mechanical properties of the PEI reinforced with continuous carbon fibres interbody fusion device were also assessed through a porcine model and results were compared with the properties of natural IVDs and the properties of IVDs prosthesized with a commercial titanium cage (Inter Fix, SOFAMOR Danek). In particular, local compression tests have led to some understanding of the effects of the interbody device inside natural segments, by comparing the results with those previously obtained from natural IVDs. Similarly to the segments prosthesized with titanium cage, the prosthesized IVDs obtained by inserting the fibre-reinforced PEI-based cage have locally presented a mean strain value which is about one or two orders of magnitude smaller than the natural L4–L5; the effect is the reduction of movements between the two adjacent vertebrae. Unfortunately, all of the mechanical tests were performed in vitro but the loads applied to the devices may differ from the complex loading seen in vivo, and therefore the results from these tests may not directly predict in vivo performance. In any case, the results obtained could be used to compare mechanical performance of different intervertebral body fusion device assemblies (Gloria et al., 2008).
7.5.2 Composite intervertebral disc prostheses Current IVD prostheses available on the market frequently undergo failure, mainly owing to wear or mismatch between mechanical proprieties of the device implanted and the natural tissue (Gloria et al., 2007; Shikinami et al., 2004). Although Link SB Charité III (Buttner-Janz, 1992) and Acroflex (Serhan, 1999) use materials with high biocompatibility with the surrounding tissues and an unrestricted biological mobility, they have not overcome these interface failure problems (Shikinami et al., 2004). Another drawback of current IVD prostheses is also related to the effect of the available sizes on the mechanical behaviour of the spine. In fact, current prosthetic implants are characterized by standard sizes and usually consist of a polymer core interposed between two metallic plates. For instance, the ProDisc prosthesis is available with 2 mm height variations, but 3D non-linear element models
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suggest that a difference of 2 mm may significantly affect the spinal curvature in that region (Rohlmann et al., 2005). Moreover, among the several IVD prostheses available none of them have the hydrophilics characteristic and hence the real properties of the natural structure. Biological materials are dynamic, complex and multifunctional, characteristics which are difficult to achieve in purely synthetic systems. It was previously believed that artificial biomaterials had to be designed to provide a high strength associated with a high modulus of elasticity at low strain levels. An attempt was made to achieve this combination of properties through the use of metals, ceramics and plastics with a relatively high strength as biomaterials. However, in contrast to most artificial materials, soft biological tissues are characterized by a large amount of strain before failure, and they are flexible and tough, showing a high strength. Consequently, it is impossible to combine all of these features by using materials with a single structural arrangement (Gloria et al., 2007; Shikinami et al., 2004). On the other hand, stiff and hard materials such as metals are often considered for manufacturing the superior and inferior interface endplates. However, the engineering alloys widely used in orthopaedic surgery (stainless steel, cobalt–chromium and titanium alloys) possess high elastic moduli compared with those of vertebral bone and a stiff implant may cause bone resorption (Bonfield et al., 1998). To satisfy all the requirements for designing an alternative intervertebral disc prosthesis, the use of poly(2-hydroxyethylmethacrylate) (PHEMA) hydrogels was considered. These hydrogels have already been used in a wide variety of biomedical applications because of their biocompatibility, high permeability and high hydrophilicity (Hoffman, 2002; Netti et al., 1993; Peppas et al., 2000). However, the mechanical properties of these materials in the hydrated state are not suitable for biomedical applications where high mechanical strength is required (Ambrosio et al., 1996, 1998; Netti et al., 1993). The mechanical properties of polymer hydrogels have been improved by the incorporation of a hydrophobic component, such as poly(caprolactone) (PCL), and polymeric fibres (Ambrosio et al., 1996, 1998; Davis et al., 1991; De Santis et al. 2004). The inclusion of hydroxyapatite and/or calcium phosphate would be beneficial, as these materials are bioactive and can stiffen polymers, as is required for the realization of the endplates (Giordano et al., 2006). The principle of mimicking the natural structure of IVD has led to design a fibre-reinforced composite hydrogel, which is able to match the mechanical properties of natural IVDs and the surrounding tissues (Ambrosio et al., 1996, 1998; De Santis et al. 2004). PHEMA-based networks composite hydrogels reinforced with poly(ethyleneterephthalate) (PET) fibres have been designed for potential
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use as IVD substitutes (Ambrosio et al., 1996, 1998). These devices were manufactured using filament winding and moulding technologies. The mechanical behaviour of PHEMA/PCL semi-interpenetrating polymer networks (s-IPNs) composite hydrogels reinforced with PET fibres was investigated (Ambrosio et al., 1996, 1998). As PCL is a biodegradable polymer, and its degradation leave voids in the network, thus undermining mechanical properties, poly(methylmethacrylate) (PMMA), a biostable polymer, was considered as a replacement for PCL to improve the mechanical behaviour of the hydrophilic composite structures. Accordingly, PHEMA/PMMA s-IPNs composite hydrogels reinforced with PET fibres were also designed and characterized (Gloria et al., 2007). This proposed approach enabled a nucleus/annulus synthetic system for IVD prosthesis to be designed and prepared that was characterized by a softer and more hydrophilic PHEMA-based s-IPNs and a stiffer and less hydrophilic outer fibrous part, which reproduces the annulus (Fig. 7.2) (Ambrosio et al., 1996, 1998; De Santis et al., 2004; Gloria et al., 2007). Although various forces act on the motion of the spine segment, it may be assumed that, in normal daily life, the loads transferred through the IVD are mainly compressive (Gloria et al., 2007; White and Panjabi, 1990). Thus, static and dynamic compressive tests were carried out on the designed composite structures. All of the dynamic tests were performed on swollen
Fibre-reinforced hydrogel as artificial annulus fibrosus Hydrogel-based artificial ancleus Hydroxyapatite-reinforced polyethylene composite endplates
(a)
(b)
(c)
7.2 (a) Schematic representation of a composite biomimetic total IVD prosthesis, which comprises a pair of hydroxyapatite-reinforced polyethylene composite endplates, a hydrogel-based nucleus and a peripheral structure made of fibre-reinforced hydrogel (Ambrosio et al., 2007); (b) image of PHEMA/PMMA 80/20 (w/w) s-IPN composite hydrogel reinforced with PET fibres as nucleus/annulus substitute (anterior view); (c) multicomponent IVD prosthesis consisting of swollen PHEMA/PMMA 80/20 (w/w) s-IPN composite hydrogel reinforced with PET fibres as nucleus/annulus substitute, and two hydroxyapatite-reinforced polyethylene composite endplates in order to anchor the device to the vertebral bodies (Ambrosio et al., 2007).
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14 12 Stress (MPa)
10 8 6 4 2 0
0
0.05
0.1
0.15
0.2
Strain (mm mm–1)
7.3 Typical stress–strain curve of swollen PHEMA/PMMA 80/20 (w/w) s-IPN composite hydrogel reinforced with PET fibres having a winding angle from 45° to 65°, compressed up to a load level of 17.0 kN without breaking (Gloria et al., 2007).
fibre-reinforced devices in physiological solutions at 37 °C by a servohydraulic MTS Bionix 858 Test System. The compressive J-shaped stress– strain curve (Fig. 7.3) obtained for the swollen PHEMA-based composite hydrogels reinforced with PET fibres is typical of soft biological tissues, such as articular fibrocartilage and IVDs, and the initial upward concavity (toe region) suggests a relatively high flexibility at low strain levels. Despite the high flexibility, with attendant low modulus and hence high compliance, high compressive strengths can be achieved, as shown by the compressive stress–strain curve (Gloria et al., 2007). The toe region is mainly caused by the matrix properties and realignment of fibres, which straighten their crimped waveform and reorient themselves in the transverse direction. During the loading process the influence of the fibres sharply increases and the linear region is related to the fibres straightening. Previous studies on canine IVDs have demonstrated that mechanical properties are strongly dependent on the spine location. In particular, the modulus in the linear region increases by a factor of 3 over the length of the spinal column, spanning from 32.0 MPa at C2–C3 level to 115.0 MPa at L6–L7 one, in terms of mean values. Analogously, the maximum stress varies from 8 to 19 MPa (Cassidy et al., 1990a). By varying the composition of the hydrogel-based matrix, the winding angle and the amount of the PET fibres, it is possible to modulate the hydrophilicity and the mechanical properties of the composite structure. For instance, compression tests performed on swollen PHEMA/PCL s-IPNs composite hydrogels reinforced with PET fibres have highlighted that the
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Table 7.1 Compressive properties of swollen composite PHEMA/PCL s-IPNs reinforced with 40% by volume of PET fibres having a winding angle from 45° to 65°, and composite PHEMA/PCL 70/30 s-IPN reinforced with 50% by volume of PET fibres having a winding angle from 45° to 90°. Modulus (E), maximum stress (σmax) and maximum strain (εmax) reported as mean value ± standard deviation (Ambrosio et al., 1996) Sample PHEMA/PCL 45°–65° PHEMA/PCL 45°–65° PHEMA/PCL 45°–65° PHEMA/PCL 45°–90°
E (MPa)
σmax (MPa)
εmax (mm/mm)
90/10; winding angle
30.17 ± 2.95
12.00 ± 0.60
0.52 ± 0.02
80/20; winding angle
60.82 ± 4.82
14.30 ± 0.39
0.48 ± 0.03
70/30; winding angle
73.30 ± 6.30
17.20 ± 1.25
0.40 ± 0.01
70/30; winding angle
129.70 ± 6.80
20.30 ± 0.90
0.32 ± 0.02
mechanical properties are a function of concentration PCL content in the s-IPN and the winding angle (Table 7.1) (Ambrosio et al., 1996, 1998). As shown in Table 7.1, a decrease in the maximum strain associated with increases in the compressive modulus and maximum stress may be observed with increasing PCL content. In particular, the compressive modulus increases from 30 to 73 MPa and the maximum stress from 12 to 17 MPa. Considering the values of compressive modulus and maximum stress reported for canine IVDs, PHEMA/PCL 90/10, 80/20 and 70/30 (w/w) s-IPNs composite hydrogels reinforced with 40% by volume of PET fibres having winding angle from 45° to 65° would present optimum compressive properties for an IVD in the cervical and thoracic regions of the spinal column. PHEMA/PCL 70/30 s-IPN reinforced with 50% by volume of PET fibres with a winding angle varying from 45° to 90° shows compressive modulus and maximum stress of 129 and 20 MPa, respectively. However, as with soft biological tissues, IVDs exhibit viscoelastic behaviours such as stress–relaxation and creep, which show the ability of the natural structures to attenuate the stress concentration when they are strained, and to limit rapid deformations when they are subjected to high stresses (Fung, 1993; Gloria et al., 2007). Results from creep tests performed on these PHEMA-based composite hydrogels show that creep is rapid at the outset, gradually levelling off over time. Comparing creep responses of the hydrophilic composite structure (Ambrosio et al. 1998; Gloria et al., 2007) to those of natural IVDs (Cassidy et al., 1990b), the fibre-reinforced hydrogel has shown a flatter creep curve and hence a higher dimensional stability, resulting in smaller strain values at creep. The differences observed may be probably ascribed to the higher
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water content of natural IVDs and to the mode of testing, which does not consider the flow of water through the surfaces, as would occur between two adjacent vertebrae. Obviously, the evaluation of compressive modulus, strength and creep behaviour is important in the first step of IVD prostheses design and the knowledge of viscoelastic properties and long-term behaviour are crucial to define design parameters. As the non-linear stress–strain curves in compression, the viscoelasticity of the PHEMA-based fibre-reinforced hydrogel cannot be interpreted through the simple theory of linear viscoelasticity. Accordingly, compressive dynamic–mechanical tests were also performed at an increasing level of mean (or static) load ranging from 200 to 2200 N with a dynamic amplitude of 100 N. The frequency was scanned over a wide range. Creep-fatigue tests were carried out on swollen PHEMA/PMMA s-IPNs reinforced with PET fibres up to ten million cycles at 2 Hz. The sinusoidal compressive load ranged from 200 to 2200 N, corresponding the load on human lumbar IVDs while lying supine and bending position, respectively, for an individual of 100 kg bodyweight. Creep recovery steps of 8 h (sleeping time) at 200 N (supine position) were also performed after different fatigue cycles (Gloria et al., 2007). The results from fatigue tests have highlighted a high long-term performance, since the fibre-reinforced hydrogels underwent ten million cycles without failure, corresponding to the recommended minimum conditions (Hedman et al., 1991; Shikinami et al., 2004). This is also supported by the fact that values of compressive storage modulus do not drop with increasing the number of cycles, indicating that no structural failure occurred. Moreover, the ability of the composite device to recover from the strain has been evidenced by the creep–recovery curves, which have shown a high mechanical stability after fatigue test (Gloria et al., 2007).
7.6
Conclusions and future trends
In order to avoid the mechanical mismatch which exists between traditional spinal devices and natural tissues, the design of new high-performance and multifunctional materials has to involve polymer-based composite materials. In contrast to commercial metal interbody fusion devices, composite cages may show mechanical properties closer to those of natural interacting structures. Thus, through a suitable composite cage design it is possible to control stress–strain distributions and hence the mechanical signals to bone, thus avoiding stress-shielding and stress-concentration effects and also the corrosion and release of metal ions that are typical of the metallic implants. Recently an approach which involves fibre-reinforced bioresorbable fusion cages has also been proposed. They were manufactured through compounding, braiding and compression moulding technologies. In
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particular, composites containing varying amounts of β-tricalcium phosphate (β-TCP) embedded in a poly(lactide) (PLA70) matrix with and without poly(lactide) (PLA96) fibre reinforcement were investigated, and the feasibility of using these composites for spinal fusion implants was considered (Huttunen et al., 2006). Even if the results obtained from this study are interesting, further mechanical tests as well as in vivo studies are needed before clinical trials. As for IVD prostheses, to overcome the drawbacks related to current devices, the research has been focused on the basic concept of mimicking the architecture of natural IVDs. A biomimetic approach has been adopted to design a fibre-reinforced hydrogel, which is able to match the mechanical properties of natural IVD and the surrounding tissues. The compressive stress–strain behaviour of natural IVDs has been reproduced by the hydrophilic composite structures, while dynamic–mechanical measurements have suggested that the complex viscoelastic behaviour of the disc tissues can be emulated by selecting the suitable matrix and an opportune design of the composite structure. Creep–fatigue tests have shown endurance, resistance to long-term compressive creep and a high dimension stability (Gloria et al., 2007). In this context, the main feature of the proposed hydrogel-based fibre-reinforced structures is that it is possible to tailor the mechanical properties, in order to optimize its characteristics at several locations along the spinal column, and to obtain the correct implant height and size by eventually integrating computer tomography, computer numerical control machining, filament winding and moulding technologies. Since polymer-based composite materials are different from conventional ones, testing methods that have been used to assess the properties of implants made of homogeneous and isotropic materials may be not appropriate for testing multifunctional composite prostheses. Accordingly, there is an increasing need for further improvements and standardization of testing methods, which could contribute to the design of highperformance devices. However, much of the recent research in the area of the IVD has focused on the biologic restoration of the IVD using growth factors, gene therapy, tissue engineering, and cells (Sakai, 2008). These approaches still have several limitations. In this field, problems are strongly related to the complex structure of IVD consisting of three tissues (nucleus, annulus and cartilage endplate) that differ histologically, chemically and physiologically. In general, IVD is populated by chondrocytes in the endplates, vacuolated notochordal cells and/or chondrocyte-like cells in the nucleus pulposus, and fibroblast-like cells in the annulus fibrosus (Leung et al., 2006). Many research groups are investigating tissue-engineered replacements for the nucleus or the whole disc using various cell sources (Hukins, 2005; Leung et al., 2006; Revell et al., 2007; Sakai, 2008 ).
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The successful implantation of tissue-engineered discs into mice may be considered an important step in the right direction (Hukins, 2005). Mizuno et al. (2004) engineered both the annulus and the nucleus tissues and tried to reproduce the anatomic shape of IVDs. In particular, a mesh of bioresorbable polyesters (polyglycolic acid coated with polylactic acid) was used to make an annulus scaffold seeded with cells from sheep annulus. Inside this annulus Mizuno et al. (2004) placed a hydrogel (sodium alginate mixed with calcium sulphate) seeded with cells from sheep nucleus. This scaffold was then implanted in a mouse and the results obtained were encouraging. Moreover, current findings have highlighted the potential of allogenic and autogenic mesenchymal stem cells to arrest IVD degeneration or to promote partial regeneration in several animals models (Leung et al., 2006). In conclusion, even though much progress has been made in the field of IVD regeneration, there are many questions and challenges that need to be addressed; thus, future studies are required not only to optimize methodology and efficacy, but also to better understand the microenvironment, phenotype and differentiation of IVD cells.
7.7
References
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cassidy jj, hiltner a, baer e (1989), ‘Hierarchical structure of the intervertebral disc’, Connective Tissue Res, 23, 75–88. cassidy jj, hiltner a, baer e (1990), ‘The response of the hierarchical structure of the intervertebral disc to uniaxial compression’, J Mater Sci: Mater Med, 1, 69–80. cassidy jj, silverstein ms, hiltner a, baer e (1990), ‘A water transport model for the creep response of the intervertebral disc’, J Mater Sci: Mater Med, 1, 81–89. ciappetta p, boriani s, fava gp (1997), ‘A carbon fiber reinforced polymer cage for vertebral body replacement: a technical note’, Neurosurg, 41(5), 1203–1206. DOI: 10.1016/S0303-8467(97)82078-1. davis pa, huang sj, ambrosio l, nicolais l, ronca d (1991), ‘A biodegradable composite artificial tendon’, J Mater Sci: Mater Med, 3, 359–364. DOI: 10.1007/ BF00705368. de santis r, ambrosio l, nicolais l (2000), ‘Polymer based composite hip prostheses’, J Inorg Biochem, 79, 97–102. DOI:10.1016/S0162-0134(99)00228-7. de santis r, sarracino f, mollica f, netti pa, ambrosio l, nicolais l (2004), ‘Continuous fibre reinforced polymers as connective tissue replacement’, Comp Sci Tech, 64, 861–871. DOI:10.1016/j.compscitech.2003.09.008. fung yc (1993), ‘Biomechanics: mechanical properties of living tissues’, New York, Springer. gershon b, cohn d, marom g (1990), ‘Utilization of composite laminate theory in the design of synthetic soft tissues for biomedical prostheses’, Biomaterials, 11, 548–552. giordano c, sanginario v, ambrosio l, di silvio l, santin m (2006), ‘Chemicalphysical characterization and in vitro preliminary biological assessment of hyaluronic acid benzyl ester – hydroxyapatite composite’, J Biomater Appl, 20, 237–252. gloria a, causa f, de santis r, netti pa, ambrosio l (2007), ‘Dynamic–mechanical properties of a novel composite intervertebral disc prosthesis’, J Mater Sci: Mater Med, 18, 2159–2165. DOI: 10.1007/s10856-007-3003-z. gloria a, manto l, de santis r, ambrosio l (2008), ‘Biomechanical behavior of a novel composite intervertebral body fusion device’, J Appl Biomater Biomech, 6(3), 163–169. griffith sl, shelokov ap, büttner-janz k, lemaire jp, zeegers ws (1994), ‘A multicenter retrospective study of the clinical results of the LINK SB Charité intervertebral prosthesis. The initial European experience’, Spine, 19, 1842–1849. hedman tp, kostuik jp, fernie gr, hellier wg (1991), ‘Design of an intervertebral disk prosthesis’, Spine, 16, 256–260. hellier wg, hedman tp, kostuik jp (1992), ‘Wear studies for development an intervertebral disc prosthesis’, Spine, 17(Supp l6), S86–S96. hoffman as (2002), ‘Hydrogels for biomedical applications’, Adv Drug Delivery Rev, 43, 3–12. hukins dwl (2005), ‘Tissue engineering a live disc’, Nature Materials, 4, 881–882. huttunen m, ashammakhi n, tormala p, kellomaki m (2006), ‘Fibre reinforced bioresorbable composites for spinal surgery’, Acta Biomater, 2, 575–587. iannace s, sabatini g, ambrosio l, nicolais l (1995), ‘Mechanical behaviour of composite artificial tendons and ligaments’, Biomaterials, 16, 675–680. kurtz sm, devine jn (2007), ‘PEEK biomaterials in trauma, orthopedic, and spinal implants’, Biomaterials, 28, 4845–4869.
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ray cd (1997), ‘Threaded titanium cages for lumbar interbody fusion’, Spine, 22, 667–680. revell pa, damien e, di silvio l, gurav n, longinotti c, ambrosio l (2007), ‘Tissue engineered intervertebral disc repair in the pig using injectable polymers’, J Mater Sci: Mater Med, 18, 303–308. DOI: 10.1007/s10856-006-0693-6. rohlmann a, zander t, bergmann g (2005), ‘Effect of total disc replacement with ProDisc on intersegmental rotation of the lumbar spine’, Spine, 30, 738–743. rothman rh, simeone fa (1992), The spine. 3rd Ed, Philadelphia, WB Saunders Company. sakai d (2008), ‘Future perspectives of cell-based therapy for intervertebral disc disease’, Eur Spine J, 17(Supp l4), S452–S458. serhan h (1999), Spinal disc prosthesis, World Patent 99/20209. shikinami y, kotani y, cunningham bw, abumi k, kaneda k (2004), ‘A biomimetic artificial disc with improved mechanical properties compared to biological intervertebral discs’, Adv Funct Mater, 14, 1039–1046. steffen t, tsantrizos a, fruth i, aebi m (2000), ‘Cages: designs and concepts’, Eur Spine J, 9, S89–S94. traynelis vc (2002), ‘Spinal arthroplasty’, Neurosurg Focus, 13(2), Article 10, 1–7. tsantrizos a, ito k, aebi m, steffen t (2005), ‘Internal strains in healthy and degenerated lumbar intervertebral discs’, Spine, 30, 2129–2137. urbaniak jr, bright ds, hopkins je (1973), ‘Replacement of intervertebral discs in chimpanzees by silicone–Dacron implants: a preliminary report’, J Biomed Mater Res Symp, 4, 165–186. van rietbergen b, huiskes r, weinans h, sumner dr, turner tm, galante jo (1993), ‘The mechanism of bone remodeling and resorption around press-fitted THA stems’, J Biomech, 26, 369–382. van steenbrugghe mh (1956), ‘Improvements in joint prosthesis’, French Patent 1,122,634. waisbrod h (1988), ‘Treatment of metastatic disease of the spine with anterior resection and stabilization by means of a new cancellous metal construct. A preliminary report’, Arch Orthop Trauma Surg, 107, 222–225. wang a, lin r, stark c, dumbleton jh (1999), ‘Suitability and limitations of carbon fiber reinforced PEEK composites as bearing surfaces for total joint replacements’, Wear, 225, 724–727. weiner bk, fraser rd (1998), ‘Spine update lumbar interbody cages’, Spine, 23, 634–640. white aa, panjabi mm (1990), Clinical biomechanics of the spine. 2nd Ed, Philadelphia, JB Lippincott Company. yamamuro t, shikata j, okumura h, kitsugi t, kakutani y, matsui t, kokubo t (1990), ‘Replacement of the lumbar vertebrae of sheep with ceramic prostheses’, J Bone Joint Surg, 72B, 889–893. zdeblick ta, phillips fm (2003), ‘Interbody cage devices’, Spine, 28(15S), S2–S7.
8 Composites for dental applications S. N. N A Z H AT, McGill University, Canada
Abstract: Composites, both particulate and fibrous, have been used in dentistry for the past few decades in a number of areas such as restorative, prosthetic, endodontic, and orthodontic dental application. They offer the advantages of aesthetics and mechanical properties that match those of dentine. This chapter will focus on the light polymerisable methacrylate restorative dental composites, and will touch upon the prefabricated fibrous composites for post and core applications. Key words: dental composites, polymers, methacrylate, prosthetics.
8.1
Restorative dental composites
Dental resin composites have been used clinically for approximately half a century in the restoration of teeth, mainly in the anterior region as well as small defects in the posterior region. Although these composites tend to have a shorter clinical service life span when compared to amalgam based restorations, they are becoming increasingly popular, mainly due to the demand for aesthetic filling, i.e. fillers that match tooth colour. As in the case of bone, dentin and enamel are examples of natural composite systems; combining collagen and hydroxyapatite crystals at different volume fractions, and the use of composite systems that match some of the aesthetic, physical and mechanical properties of these tissues is logical. As in all composite systems the properties of the restorative dental composites depend on the properties of the constituents, i.e. the organic based matrix and the inorganic filler, their respective volume fractions, and the quality of the interface separating them. This chapter focuses on light polymerisable methacrylate-based restorative dental composites. These are initially provided to the dental clinician as soft single-paste systems, which then harden at the site of application under a visible light activated polymerisation process carried out at ambient conditions. In the 1960s, Bowen (1962) pioneered these composite systems by developing a methacrylate resin based on bisphenol A glycidyl methacrylate (Bis-GMA) and reinforced with silanated fillers. This composition revolutionised the use of these materials as dental composites. After experimenting with the reinforcement of epoxy resins he switched to methacrylate-based systems as epoxies were slow to polymerise. This high 201
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molar mass Bis-GMA monomer allowed for polymerisation to take place without the significant shrinkage experienced by the lower molar mass monomers such as methyl methacrylate. Currently, there are three main components that make up the composite systems that are used in restorative dentistry: 1. A continuous phase, which comprises the polymer-based matrix. Before polymerisation or curing, this is dominated by the monomer systems. It also contains polymerisation initiators, which trigger the hardening process as well as stabilisers that increase the shelf life of the unpolymerised paste. 2. A discontinuous phase, which comprises an inorganic filler to impart an increase in mechanical and physical properties. These can be based on quartz, silica glasses, or fused silica. 3. A coupling agent, which comprises a silane coupling agent that chemically bonds the filler to the matrix. Although the advancement in the field has increased the clinical lifetime of these composites, they are less successful in posterior than in anterior regions owing to greater mastication stresses that are imposed in the former. Moreover, they generally still suffer from two main problems that ultimately end in their premature failure: chemically, they undergo shrinkage owing to polymerisation that could lead to marginal leakage, which is the collection of oral fluids and bacteria in between the filling material and the tooth causing decay, and mechanically, they tend to have poor resistance to wear or abrasion. Along with focusing on improving these two drawbacks, research into dental composites have also been aimed at improving on biocompatibility (toxicity) as well as the development of drug releasing composites (e.g. fluoride and chlorhexidine) that prevent secondary caries and the degeneration of the mineral phase of teeth.
8.2
Matrix monomers
Monomer selection in restorative dental composites influences a number of properties such as the viscosity of paste, shrinkage owing to polymerisation, the coefficient of thermal expansion, water absorption, and the mechanical properties of polymerised composite. Early developments of resins based on methyl methacrylate suffered drawbacks such as high shrinkage, high coefficient of thermal expansion and the frequent occurrence of secondary caries. Nowadays, the main methacrylate-based monomer systems are made of aliphatic and aromatic dimethacrylate of Bis-GMA, urethane dimethacrylate (UDMA) triethylene glycol dimethacrylates (TEGDMA). Some of the properties of these monomers are given in Table 8.1. The difunctional Bis-GMA monomer has a high molar mass, which
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Table 8.1 Types of methacrylate-based monomers used in the matrix of restorative dental composites (from Moszner N, Salz U, 2001). Adapted from Progress in Polymer Science, vol. 26, no. 4, Moszner N, Salz U, New developments of polymeric dental composites, 535–576, copyright (2008).
Monomer
Molar mass (g mol−1)
Viscosity (Pa s)
Density (g cm−3)
Density of polymer (g cm−3)
ΔS (%)
Bis-GMA UDMA TEGDMA
512 470 286
500–800 5–10 0.1
1.151 1.110 1.072
1.226 1.190 1.250
−6.1 −6.7 −14.3
allows for lower overall shrinkage owing to polymerisation. However, because of its high molar mass, Bis-GMA also suffers from high viscosity and therefore other viscosity-reducing dimethacrylate monomers (e.g. TEGDMA) are used in combination with Bis-GMA. This is important as it will allow for easy mixing, manipulation and application of the paste at ambient temperature before polymerisation. Before applications, the monomer components are kept in a single-paste system which includes an initiating system consisting of a photoinitiator molecule, usually camphoroquinone and an amine activator, usually N,Ndimethyl-p-toludine. Polymerisation is carried out by free radical polymerisation that is activated by visible light. When exposed to light (λ = 400 to 510 nm) polymerisation takes place at increments of up to 2 mm thick samples. Each polymerisation step is usually about 40 s. The process of polymerisation is the formation of chains containing large numbers of similar repeat units, namely the monomers. These are covalently linked to each other along the backbone as well as in between the chains. This is initiated by the breakup of the carbon–carbon double bonds (C=C) in the starting monomers to create larger molecules that are chemically linked to each other in a three-dimensional network i.e., crosslinked. The percentage degree of conversion measures the consumption of C=C bonds, and the greater the conversion percentage the greater the extent of polymerisation. As well as being used as diluents, the dimethacrylates are used to allow for crosslinking to take place between the chains, which increases the resistance to solvent attack. However, these diluents can only be used in smaller amounts relative to Bis-GMA as they will increase the shrinkage caused by polymerisation as given in Table 8.1. Shrinkage occurs because of the reduction of interchain distances which can be up to 50%. The magnitude of shrinkage caused by polymerisation is dependent on the molar mass of the monomer, and it decreases with an increase in molar mass. The prevention of shrinkage owing to polymerisation is important not only in preventing secondary caries, but also in the reduction of the presence of stress concentrations between the dental composite and the tooth cavity.
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Table 8.2 Property requirements of the monomers and polymers and their effects on the clinical performance of restorative dental composites (from Moszner N, Salz U, 2001). Adapted from Progress in Polymer Science, vol. 26, no. 4, Moszner N, Salz U, New developments of polymeric dental composites, 535–576, copyright (2008). Monomer/polymer requirements
Effects on clinical performance
Low shrinkage caused by polymerisation
Easier processing and no marginal leakage and stress concentrations Shorter polymerisation times Sufficient physical and mechanical properties Longer term durability Lower failure rate
Higher rate of polymerisation Cross-linking capacity Glass transition temperature above 60°C Greater resistance to oral conditions e.g., temperature fluctuations, continuous dynamic loading, fluid content Higher light and coloration stability of the polymer Low toxicity
Longer term aesthetics Minimum toxicological risk to the patient
The volume shrinkage owing to polymerisation can be calculated from equation (1) below: S=
ρ p − ρm 100 ρp
(8.1)
where ΔS is the percentage shrinkage caused by polymerisation, and ρm and ρp are the densities of the monomer and polymer, respectively. The shrinkage is dependent on the number of C=C in the monomer relative to the molar mass, amount of filler and the volume of composite. The successful clinical performance of the polymer-based matrix of restorative dental composites is dependent on a number of factors which are listed in Table 8.2.
8.3
Reinforcing agents
As in all composite systems, reinforcing agents are used in restorative dental composites to improve a number of chemical, mechanical and physical properties such as shrinkage caused by polymerisation, stiffness, strength, wear resistance as well as water sorption and coefficient of thermal expansion. For this purpose, a number of filler particle types have been used which include quartz, silica based glasses, and colloidal silica. The larger sized fillers are usually produced by the grinding or milling of quartz or glasses to produce particles in the range 0.1 to 100 μm. Quartz fillers were used in particular during the early development of dental restor-
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ative composites. Owing to the hardness of quartz, they are difficult to grind down and the composites tend to have less polished surfaces which cause abrasion of the other surfaces. Silica based glasses are normally doped by heavy metal oxides such as barium, strontium and zirconium oxides. These doping oxides impart radiopacity that is they absorb or reflect X-rays, which is an important requirement in the diagnosis of secondary caries. Colloidal silica particles of approximately 40 nm size can be obtained either through precipitation or pyrolysis. Owing to their small particle size they have a large surface area ranging between 50 and 300 m2 g−1. This large surface area increases the viscosity of the paste before polymerisation owing to the increase in polar bonds between filler and monomers.
8.4
Silane coupling agents
In composite systems, the efficacy of reinforcement is significantly affected by the quality and nature of the interface between filler and of the matrix. Therefore, adhesion promoters are often used to wet the fillers, and to prevent agglomerate as well as void formation. Silane coupling agents are also used as they impart a chemical bond between the filler and matrix, which ultimately is the best form of bonding in order to achieve increased strength as it allows for stress to be transferred from matrix to filler. The primary function of a coupling agent is to provide a strong chemical link between the methacrylate polymer matrix and the oxide groups on the glass filler surface. One of the most common adhesion promoters is γ-methacryloxypropyltrimethoxysilane. This organosilane agent reacts with both the methacrylate polymer and filler’s silanol groups by forming a siloxane bond. When hydrated, the methacrylate groups covalently attach to the monomer while silanol groups attach to the particle surface. The silane coating also leads to a strong water-resistant bond thus reducing the hydrolytic degradation of this interface. In the absence of a silane coupling agent, the filler-matrix interface deteriorates rapidly in the presence of water which diffuses through the matrix and attacks through hydration. This results in the large reduction of the strength of the composites. Water also acts as a plasticiser in resins. Organosilanes are, however, less efficient in fillers not containing silica.
8.5
Classifications of dental composites
The properties of composites are dependent on a number of factors including the properties of the constituent matrix and fillers, the volume or weight fraction of these constituents, the distribution of the discontinuous phase (filler particles) and the quality of the interface separating the interface.
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Table 8.3 Classification and some properties of restorative dental composites (from Anusavice, KJ, 1996). Reprinted from Phillips’ Science of Dental Materials, Anusavice KJ, Restorative Resins, 273–299, copyright (1996). Property1
Composite classification
φ (μm)
Unfilled resin Traditional Small particle Microfilled Hybrid
– 8–12 0.04–0.4 1–5 0.6–1.0
Vf
E (GPa)
– 0.60–0.65 0.65–0.77 0.20–0.35 0.60–0.65
2.4 8–15 15–20 3–6 7–12
σc (MPa)
σt (MPa)
CTE (10−6/°C)
Water sorption (%)
70 250–300 350–400 250–350 300–350
24 50–65 75–90 30–50 70–90
92.8 25–35 19–26 5–60 30–40
1.7 0.5–0.7 0.5–0.6 1.4–1.7 0.5–0.7
1 φ is the filler particle size, Vf is the filler volume fraction, E is the composite elastic modulus, σc and σt are the composite compressive and tensile strengths respectively, and CTE is the composite coefficient of thermal expansion.
Restorative dental composites are classified in terms of the size of the filler reinforcing agent. Table 8.3 lists the classification of the composite according to the range of filler particle size and provides some of the mechanical and physical properties. These fillers affect the properties of the composites where filler content, viscosity and shrinkage caused by polymerisation as well as the physical and mechanical properties are essentially linked. Depending on the type of filler i.e., classification, these composites often contain fillers ranging from 0.3 to 0.7 volume fraction. In some restorative dental composites a combination of filler types are used to improve the packing, which are known as hybrid composites. The properties of the various restorative dental composite classifications are summarised as follows: 1. Traditional composites have the largest particle size and distribution. These quartz filler incorporating composites were used in the initial phase, however, these tend to have rough surfaces. 2. Small particle-filled composites are characterised by a high filler volume fraction which is the result of a combination of particle size and distribution allowing for increased filler packing. The majority of fillers are silica-based glass particles as well as small amounts of colloidal silica. Owing to the increased filler volume fraction, this classification of restorative dental composites demonstrates the best properties when compared with the other classifications. These composites also tend to have smoother surfaces than traditional composites. 3. Microfilled composites incorporate colloidal silica as the filler. As in the case of nanocomposites, these nanosized particles suffer from agglom-
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eration. Furthermore, owing to the large surface area, they cause an increase in viscosity of the paste mixture before polymerisation, even at low inclusions. Therefore, to produce composites with higher filler volume fraction, pre-polymerised composites particles of silane-treated colloidal silica are first prepared and then mixed into another paste of monomer and colloidal silica. This whole complex is then subsequently polymerised to form the final composite. This process also succeeds in reducing the shrinkage caused by polymerisation. However, owing to the relatively lower filler volume fraction, these composites have lower mechanical and physical properties than other classes of composites. However, they have the smoothest surface finish compared with the other classifications. 4. Hybrid composites combine two types of reinforcing agents to increase the volume fraction. A significant portion of the filler consist of glass particles of smaller particle sizes than those found in the small particle filler composites and up to 0.2 wt fraction of colloidal silica. Hybrid composites offer intermediate properties when compared with the other classifications, and also have smooth surface finish.
8.6
General property requirements of dental composites
As discussed above, the incorporation of silanated fillers into the polymer matrix results not only in the decrease in the shrinkage owing to polymerisation, but also in an increase in the mechanical properties of elastic modulus and strength. As the mouth is a highly hydrated environment, the water absorption behaviour of the composites is an important factor to consider. Water absorption normally takes place in the polymer matrix and is dependent on factors such the presence of hydrophilic groups or inclusions. If it occurs at high levels, it can lead to a plasticising effect on polymer, which ultimately lowers the glass transition temperature and mechanical properties. Other consequences of the high levels of water sorption include the hydrolysis led breakdown and loss of bond at the matrix filler interface which leads to a high rate of wear. Water absorption also leads to the volumetric swelling of the composite causing cracks as well as the leaching of low molar mass components such as un-reacted monomers, which can be toxic. Water absorption can be controlled by the crosslinking density in the matrix. Other physical properties include the coefficient of thermal expansion (CTE). The fillers tend to have CTE properties similar to that of teeth which are much lower than those of the matrix. Therefore a reduction in CTE of the composites can be achieved by increasing the filler volume fraction. Also there are a number of problems that are related to the biocompatibility
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of restorative dental composites, which include the toxicity of the materials used and marginal leakage. As the polymerised materials are inert, toxicity is only a concern if there are leachable components such as unreacted monomers or low molar mass species. The leaching out of these may cause allergic reactions within patients. Marginal leakage is a major concern, resulting in secondary caries, the major source of the dental filling replacement. This occurs as a result of shrinkage caused by polymerisation leaving a gap between the filling composite and the cavity thus allowing for bacterial ingrowth. For this reason, recent developments of dental composites have also been concerned with preventing plaque accumulation, such as developing fluoride ion releasing systems. These ions are intended to reduce the acid-induced degeneration of the mineral phase of the tooth through the formation of fluorohydroxyapatite. Although fluoride ions can be released from the matrix, more often they are incorporated into the glass filler through barium or strontium fluorosilicate glasses. Other antimicrobial agents such as chlorhexidine have also been incorporated into the matrix though physical entrapment.
8.7
A brief overview of fibrous composites in dental applications
Fibrous composites have been used in dental applications for up to 30 years. Their applications have ranged from aesthetic orthodontic archwires and brackets, to posts in post/core applications in the restoration of endodontically treated teeth. They were originally developed to replace the various metallic materials such as stainless steel, cobalt–chrome alloys, titanium alloys, nickel titanium, and gold. Endodontic treatment is required for teeth that have undergone trauma and severe carries. More recent research has paid attention to the development of fibre-reinforced polymer composites as posts. These are characterised as being more aesthetic, and have superior mechanical properties to metallic based posts. The high modulus of the metallic posts offer disadvantages in adjustment, and may cause the fracture of tooth root. Composites are designed to reduce the modulus mismatch between the dentine and post, i.e., they have similar modulus values to dentine. Other mechanical properties of significance are compressive strength and fatigue behaviour. A further issue that needs to be considered is corrosion resistance, for example stainless steel which contains nickel, has been associated with nickel allergy. The coefficient of thermal expansion of the composites should be also similar to dentine. The materials considered in composite posts are carbon or silica fibre reinforced epoxy resin. The fibres are in the range of 10 μm in diameter, which may be longitudinal, braided, or woven. The dimensions of the posts
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range from less than 1 mm up to 2 mm in diameter, and 20 mm in length and their shape may be parallel or tapered. These composite posts are designed to bond well to resin-based composite core materials. In case of failure, they can also be easily removed compared with metallic posts.
8.8
Conclusion
Both fibrous and particulate composites have been used in dental applications. These application are also set to grow with time. While the main drive for the use of composites has been aesthetics, they also offer other advantages over traditional metal-based materials owing to their chemical, mechanical and physical properties.
8.9
Sources of further information and advice
Anusavice KJ Chapter 12 in Restorative Resins in Phillips’ Science of Dental Materials, WB Saunders Company. 1996, pp 273–299. Bowen RL. Dental filling material comprising vinyl–silane treated fused silica and a binder consisting of the reaction of bisphenol and glycidyl methacrylate. US Patent 3,006,112, 1962. Cheung W. A review of the management of endodontically treated teeth: post, core and the final restoration. The Journal of the American Dental Association 2005, 136, 611–619. Drummond JL, Degradation, fatigue and failure of resin dental composite materials. Journal of Dental Research 2008, 87, 710–719. Fujihara K, Teo K, Gopal R, Ramakrishna S, Foong KWC, Chew CL. Fibrous composite materials in dentistry and orthopaedics: review and application. Composites Science and Technology 2004, 64, 775–788. Moszner N, Salz U. New developments of polymeric dental composites. Progress in Polymer Science 2001, 26, 535–576. Peutzfeldt A. Resin composites in dentistry: the monomer systems. European Journal of Oral Sciences 1997, 105, 97–116.
9 Acrylic bone cements for joint replacement S. D E B, King’s College London Dental Institute, UK
Abstract: The physical and mechanical properties of poly(methyl methacrylate) cement have been extensively studied and a number of variables are known to affect the characteristics of a given property, the influential ones being the composition and porosity. The incidence of infection, bone resorption and the fact that in areas of bone filled with acrylic cement there is no prospect of bony regeneration, has led to research in numerous directions. In this chapter, existing data on acrylic bone cements are reviewed and a brief overview of PMMA composite cements is presented. The future direction of research and the need for standardisation of test methods and long- term clinical studies are highlighted. Key words: bone cement, poly(methyl methacrylate), bone regeneration.
9.1
Introduction
Disease conditions, such as osteoarthritis, osteoporosis or rheumatoid arthritis cause degeneration of the joint. The most common reason for having a hip or knee replaced is osteoarthritis, a degenerative joint disease causing breakdown of the cartilage in the joints. Total joint replacement has been widely adopted to treat these debilitating illnesses, as it is often the only course of treatment to alleviate pain and improve the quality of life. Given the increase in population of the elderly and associated disease and trauma, the number of joint replacements worldwide is expected to rise. Joint replacement surgery has become one of the most common operations in orthopaedic surgery with good prognosis. The success and widespread use of joint replacements in the management of arthritic conditions and trauma has made a significant impact in modern healthcare. The successful outcomes of total hip/knee joint replacements have enabled people to live more active lives and have a better quality of life. Over the past five decades, advances in the design, construction, technique and implantation of artificial hip/knee joints have resulted in a high degree of success.
9.1.1 Hip and knee joint The hip has a ball and socket joint that is subjected to a wide range of movement. Under normal conditions, the ball moves smoothly in the socket 210
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with a lining of articular cartilage. Conditions that commonly lead to a hip joint replacement are osteoarthritis, rheumatoid arthritis, fracture of the neck of the femur and or trauma that leads to irreversible damage. In a total hip replacement an artificial hip joint is surgically implanted. There are two components of a hip joint: (i) the acetabulum or the cup that replaces the hip socket and is generally made out of ultra high molecular weight polyethylene or highly crosslinked polyethylene, metals and ceramics or a combination thereof (Fig. 9.1); and (ii) a metal or ceramic head attached to a metal stem that is inserted into the femur to provide stability to the prosthesis. The knee joint is more complex; in a healthy knee, the ligaments hold the bones in place thus allowing the joint to function properly. In a normal knee, four ligaments help hold the bones in place so that the joint works properly. When a knee becomes arthritic, these ligaments can become scarred or damaged. During knee replacement surgery, some of these ligaments, as well as the joint surfaces, are substituted or replaced by the new artificial prostheses. The knee joint is subjected to not only flexing and extending but
Acetabular cup Femoral component
(a)
(b)
9.1 (a) Ball and socket joint; (b) artificial hip joint. (Fig. 9.1b reprinted with kind permission of Japan Medical Materials Corporation.)
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also a certain degree of rotating. Knee joints also come in two forms, an uncemented and cemented prosthesis. The prosthesis consists of two parts and is implanted into the ends of the femur, the tibia and the patella with or without bone cement. Metal implants are generally used on the femur end whereas polymers such as polyethylene are used on the tibia and patella surface, but combinations such as metal–metal, ceramic on ceramic and ceramic on plastics are also used.
9.1.2 Joint replacement Joint replacements may be cemented, uncemented and, in some cases, a combination of the two depending on the surgical requirements. Cemented fixation of the hip allows the patient to be mobilised in the very early stages post-operatively and this results in more rapid rehabilitation. The cemented prosthesis is highly dependent on the stable interface between the prosthesis and the cement and the mechanical bond between cement and bone, thus a strong bond at the bone–cement–prostheses interface is important. Cemented surgery involves use of self-curing bone cements that assist in anchoring the prosthetic components to contiguous bone whereas uncemented replacements use prosthetic components that are placed by a press fit technique that allows bone to grow into the porous surface of the implant, leading to a biological fixation. The bone cement has often being implicated in the loosening of the prosthetic components. Crack or fatigue fractures have been noted in the cement present in the femoral component and this has been suggested to be one of the causes of failure. Wear particles, also generated during the function of the joint have been proven to contribute to loosening via the process of osteolysis. However, it is difficult to attribute one particular reason for the failure of joint replacement as the causes are multifactorial and, despite cement failure being recognised as one such factor, cemented surgery is recommended for the older patient and patients with rheumatoid arthritis or poor bone density. Cementless surgery was introduced in the 1980s as a means to facilitate a direct bonding between the implant and the bone that would eliminate the use of the cement. These implants are either textured or coated to create a surface topography that would promote bone growth onto the surface of the implant. The designs of the prosthesis that are used for cementless surgery are different, especially in size. The procedure also entails the precise preparation of the femur to allow a good contact of the implant with the bone. Although cementless implants obviate the use of bone cement the procedure requires longer healing time as the stability of the prosthesis is dependent on bone ingrowth. Total joint replacements need permanent implants that are required to provide function over a period of time. The long-term performance of each
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component is one of the determinants of the success of the surgical procedure. However, wear particulate debris, implant loosening, and failure of cement are some of the factors that necessitate further surgery. Nevertheless, the success of a joint replacement lasting approximately 20 years is around 80%, thus it is has a considerable role in alleviating pain and improving the quality of life for a large number of the population. Poly(methylmethacrylate) PMMA bone cement, clinically the most widely used bone cement for anchoring joint prosthesis, is discussed in detail in the following sections. The PMMA bone cement is essentially a mixture of the prepolymer of PMMA with a newly formed matrix of selfcuring PMMA and it can either be alluded to as an interpenetrating network or even a composite owing to the presence of inorganic salts such as barium sulphate or zirconia that renders the cement radio-opaque for clinical radiological detection. Attempts to enhance the mechanical, physical and biological properties of acrylic bone cements to improve their clinical performance has been addressed by numerous researchers (Deb 1999) and the inclusion of fibres and or particulate fillers has been successful in some cases, although the handling characteristics of the cement are compromised. Improvements such as the inclusion of safer new activators with high-molecular-weight monomers have also been reported (Rojo 2008; Deb 2008; Vazquez 2002). However, with the need to enhance cement properties there has been a substantial amount of research dedicated to enhancing the mechanical properties of acrylic cements and composite technology has also been applied to this field. A short discussion related to the development of composites is presented in section 9.6, but the clinical use of such cements is not common and only a selected few cements have undergone clinical trials.
9.2
Acrylic bone cement
Poly(methylmethacrylate) PMMA, self-curing acrylic cements were first used by Charnley in the 1960s (Charnley 1960) to anchor the hip prosthesis that led to the clinical success and wide use of this surgical procedure. The placement of an artificial hip or knee implant is essentially to restore the function of the damaged joint. Thus, the design, material properties, shape and surgical technique determine the load transfer characteristic, which is an important determinant of the long-term success of the replacement surgery as the bone remodels and responds directly to the loading pattern. The PMMA bone cement is placed between the bone and the implant and the bone–cement–implant construct enables the cement to transfer body weight and service loads more effectively from the prosthesis to the bone. It is important to reflect on the properties of the hip implant itself, which is usually a metal made from Co–Cr alloys, stainless steel and other metals
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that usually have a much higher modulus than bone, thus stress shielding can lead to bone resorption and eventually the loosening of the implant. Self-curing PMMA cements have been widely used in dentistry and orthopaedic surgery as filling agents and for fixation of joint prosthesis with a fair degree of success. The wide array of the properties of the two-part self-polymerising bone cement makes it a very attractive biomaterial for medical applications. The main function of the bone cement is to transfer body weight and service loads from prosthesis to bone and immediate immobilisation of the prosthesis. PMMA bone cement provides adequate fixation of the femoral stem for up to 10 to 15 years, but beyond that time, failure is nearly unavoidable, even using the latest methods of femoral canal preparation, cement mixing and delivery. There are many possible reasons for the failure of the prosthesis, including sepsis or mechanical failure of the cement mantle. Acrylic bone cement is currently the most widely used biomaterial for anchoring cemented arthroplasties to contiguous bone and performs admirably owing to its range of properties; it is, however, beset with some well known disadvantages (Lewis 1997; Deb 1999). The polymerisation reaction is exothermic and generates temperatures of 66 to 120 °C. The high rise in temperature is often a cause of necrosis and impairment of blood circulation, and it is one of the reasons for formation of fibrous tissue around the bone–cement interface. Methylmethacrylate (MMA) can cause hypotensive effects that may induce adverse systemic effects if it enters the blood stream. Unreacted monomer resulting from incomplete polymerisation may also leach into the surrounding tissues, leading to chemical necrosis. Furthermore, PMMA is a brittle material with low fracture toughness, poor fatigue life and a mismatch of modulus between the cement and the bone, factors that may all contribute to mechanical failure.
9.2.1 Composition of PMMA bone cement Poly(methylmethacrylate) (PMMA) bone cement is a two-component system comprising of a solid and a liquid phase and the typical components are presented in Table 9.1. Commercial bone cement formulations mainly vary in their pre-polymer powder composition (solid phase) with either pure PMMA, methylmethacrylate–methylacrylate (MMA-MA) or MMA– ethylacrylate (EA) copolymers, or polystyrene/PMMA copolymers to name a few examples (Demian and Mc Dermott 1998; Kühn 2000). Setting of bone cement The solid powder phase when mixed with the liquid phase allows the monomer methylmethacrylate to wet the beads of the pre-polymerised
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Table 9.1 The typical constituents of an acrylic bone cement Powder component
Liquid component
Poly(methylmethacrylate) (PMMA) 89% w/w Benzoylperoxide (BPO) 0.75% w/w Radiopacifiers such as zirconia or barium sulfate
Methylmethacrylate (MMA) 97% v/v N,N-Dimethyl-p-toluidine (DMT) 2.7% v/v Hydroquinone 750 ppm Colouring additives such as chlorophyll
PMMA, which also allows the initiator, benzoyl peroxide to be released from the PMMA beads leading to the onset of the reaction. The initial stages of mixing first yield a sandy mixture which progresses on to a doughlike mixture, at which point the cement needs to be placed into the cavity so that the flowability of the paste is not lost. The free radical addition polymerisation occurs in three steps as elucidated below: •
Initiation: the initiator benzoyl peroxide (BPO) decomposes to form free radicals in presence of a tertiary amine, usually N,N-dimethyl-4toluidine (DMPT) incorporated within the liquid phase. C6 H 5 COO ⎯ OOCC6 H 5 benzoyl peroxide
•
•
→
2 ( C6 H 5 COO )i benzoyl free radical
Propagation: chain propagation occurs as the monomer units add on, the reaction thus proceeds, leading to the formation of the polymer chain (Fig. 9.2). Termination: the reaction ceases either via the combination or disproportionation of free radicals.
Curing parameters The mixing of the two phases of the cement triggers a chain of events and the stages as observed by the surgeon is explained in terms of the following steps of the polymerisation process: 1
2
On mixing the powder and the liquid, the monomer MMA begins to dissolve the beads of PMMA thus releasing the benzoyl peroxide and a sandy texture is obtained. The mixture progresses to a smooth consistency losing its sandy texture (dough time) with the progressive dissolution of the beads and the polymerisation process is initiated. The dough time is defined as the time after mixing from which the cement progresses to a smooth consistency and does not adhere to surgical gloves and it is imperative that the
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2(C6H5COO)•
Benzoyl peroxide
Benzoyl free radical CH3
R• +
CH3
CH 2 =C-COOCH3
Free radical CH3
R-CH 2-C •-COOCH3
Methyl methacrylate CH3
R-CH2 -C• -COOCH3 + CH 2 =C-COOCH3
CH3 R-CH 2 -C-COOCH 3 CH 2 C•COOCH3 CH 3
9.2 Sequence of the MMA polymerisation reaction.
3
4
cement is placed in the cavity between the dough time and working time, since the ability of the cement to flow decreases after the working time, impeding the interdigitation of the cement with bone. As the reaction proceeds, the viscosity of the mix gradually rises and heat is generated (propagation). The polymerisation is a highly exothermic reaction and, in addition, the gel or Trommsdorff–Norrish effect contributes to the temperature rise, a characteristic of radical polymerisation of acrylates when the termination rate is dramatically decreased owing to the high viscosity of the medium. Finally, the viscosity rises very rapidly and the cement mass sets when the maximum exotherm temperature is reached (setting time).
The polymerisation of the bone cement is predominantly based on the bulk polymerisation of MMA, which leads to a biphasic network of prepolymerised PMMA particles dispersed into a newly polymerised MMA matrix. The reaction follows a typical free radical polymerisation mechanism as shown in Fig. 9.2 (Bauer 1966). During the propagation step where growing chains are formed, not only does the viscosity rise but heat is also liberated as the conversion of a mole of MMA releases around 13 kcal of energy (Kuhn 2000). The chemical composition of the cement, the working environment, and the mixing and dispensing methods are factors that influence the curing parameters and properties of the cement. Although most commercial formulations have very similar powder compositions, the variables are mainly
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Table 9.2 Typical curing parameters of some commercial bone cements
Cement
Dough time1 (min)
Working time1 (min)
Setting time1 (min)
Peak temperature (°C)
Palacos® CMW1 Simplex P
2.0 2.2 3.0
9.0 8.5 9.0
11.0 11.0 10.0
73.0 88.0 90.0
1
Recorded ambient temperature (21–22°C).
in the differing copolymers used, the molecular weight of the pre-polymers, the bead size, the ratio of activator: initiator, and the radiopaque agents used. Although the setting properties are primarily associated with the composition of the cement, many other parameters have been found to affect the setting properties of bone cements, an important factor being the environmental temperature (Lewis 1997; Deb 1999). Higher environmental temperatures decrease both working and setting time, which can influence the handling of the cement. There are also reports that even the specific brand of surgical gloves has a significant effect on dough time measurements (He 2003). The other important environmental factor is the presence of oxygen since it inhibits the polymerisation reaction and the concentration of oxygen has a significant effect on the setting time. The influence of oxygen mainly occurs at the earlier stages of the bone cement setting process and delays the rise in temperature by lowering the free radical concentration. As the oxygen concentration increases, the setting time increases, but it does not affect the dough time (He 2003a). Table 9.2 presents typical values of curing parameters of some commercial cements. The recommended powder to liquid ratio (P : L) for most cement formulations is 2 : 1 and any changes to the P : L ratio results in the alteration of the viscosity of the cement mix and the curing parameters. A higher powder to liquid (P : L) ratio results in a shorter setting time, owing to the relative increase in the amount of the initiator that is present in the solid phase, whereas a lower P : L ratio favours the wetting stage as more liquid MMA is able to dissolve the PMMA beads thus improving the handling (Lautenschlager 1984); however, optimisation is important as it can affect the net shrinkage. Residual monomer and shrinkage The polymerisation of MMA is accompanied by shrinkage and presence of small amounts of unreacted monomer. The average low temperature during the course of the polymerisation and the rigidity of the macromolecules inhibit conversion of the monomer to the polymer. The other factor that
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leaves uncured monomer is the high viscosity of the mix during the polymerisation that restricts diffusion of free radicals, thus adding to total residual monomer content. The presence of unreacted monomer can be detrimental to the mechanical properties of the cement, either because it is leached from the cement matrix or it acts as a plasticiser. Shrinkage in bone cements is another parameter that has been related to loosening of the prostheses. Polymerisation of acrylic bone cement is accompanied by a change in density whilst the liquid monomer, with a density of 0.937 g ml−1, converts to a polymer of higher density (1.18 g ml−1). Thus, it results in volume shrinkage, which is of clinical concern as implant loosening may follow post placement. Pure MMA polymerisation exhibits approximately a 21% shrinkage; however, as the orthopaedic bone cement is presented as a pre-polymerised PMMA mix with only one third of it containing the monomer, the shrinkage decreases to about 7%. Higher concentrations of polymer powder decrease net shrinkage but compromise the clinical handling and mechanical properties. However, the net shrinkage is lower than expected owing to the inherent porosity that arises during the mixing stage through the incorporation of air bubbles or because of evaporation of the monomer at very high temperatures. Shrinkage has also been associated with poor inferior mechanical properties and Orr et al. (2003) have confirmed the appearance of cracks around the stem of a hip prosthesis and related it to shrinkage with the residual stresses sufficient to cause crack initiation before functional loading. Thermal shrinkage has also been implicated in contributing towards shrinkage as residual stresses develop that may be sufficient to initiate cracks at the interface between the implant and cement (Jasty 1991; Ahmed 1982). The net volume shrinkage is greater in vacuum mixed cements and can be as high as 7–8% compared with hand mixed cements that shrink between 1 and 3% as a consequence of the loss of the air bubbles during the application of vacuum.
9.3
Properties of acrylic bone cements
9.3.1 Porosity Porosity is generally considered to have an adverse effect on the survival of cemented total hip arthroplasties (Dunne 2003; Lewis 1999; Murphy 2002) and has been shown to affect the mechanical properties of bone cements [Dunne 2003; James 1992; Murphy 2000]. During the mixing of bone cement components there is a possibility of incorporation of air bubbles and in addition, the steep increase in temperature accompanied with the reaction can also cause some of the volatile MMA to evaporate generating voids in the matrix. The relatively low fracture toughness of PMMA-based bone cements has been associated with porosity because
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pores are marked as stress concentration points and crack initiators sites; increased porosity may also contribute to the presence of microcracks (McCormack 1996; Murphy 2002; Topoleski 1993). Pores in cement mantles have been clearly identified to affect the fatigue properties of cements. A study by Topoleski et al. on retrieved bone cement samples reported that the pores acted as nucleation sites for microcracks, thus accelerating the crack propagation within the cement mantle. However, pores can also be beneficial since they lead to an increase in the damage zone in front of the crack, which blunts the crack thus decreasing the rate of crack propagation (Topoleski 1990, 1993). From laboratory studies, it has been established that porosity is detrimental to the fatigue performance of acrylic bone cements with pore size and pore distribution being important parameters. Hoey and Taylor (2009) reported the effect of porosity on fatigue strength by comparing a handmixed and vacuum-mixed cement and used the theory of critical distances to conclude that failure initiated from large pores in vacuum-mixed cements and from cluster of pores close to each other in the hand mixed cements. The size of the pore was also important as it acted as stress concentration points, resulting in crack initiation. Porosity is not all detrimental in bone cements as it has some beneficial effects, especially in the release of antibiotics (Baker 1988, Belt 2000) and compensates for shrinkage to an extent.
9.3.2 Mechanical properties PMMA bone cement is used extensively as a load-bearing prosthetic component in total hip replacement surgery. The bone cement is placed between the stiff metallic implant (E∼214 GPa) and bone, which has a significantly lower modulus (E∼20 GPa), thus it functions as a compliant buffer and allows for stress transfer at the bone–cement interface. The mechanical properties of the hardened bone cement are very important in terms of clinical success, as during the in vivo application the cement has to withstand compressive and shear stresses. The tests used to characterise the mechanical properties of bone cements, range from quasi-static to in vivo fatigue load conditions. The FDA requires the pre-clinical mechanical testing of the cement and ASTM provide guidelines for standard testing protocols. However, a survey of the literature indicates that there is a wide variation in test methods and specimen geometries being employed, thus making it difficult to compare the mechanical properties of different cements. Stress–strain characteristics of poly(methyl methacrylate) bone cement Acrylic bone cements are viscoelastic in nature and thus properties such as tensile, compressive, flexural and shear properties depend on the strain rate.
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PMMA bone cement is relatively weak in tension but resistant in terms of compression. The tensile strength and elastic modulus values reported lie in the range of 22–50 MPa and 1.7–3.2 GPa, respectively. Compressive strength values of 64–103 MPa, which are related to 50–70% strength of the cortical bone, and elastic modulus values were between 2.1–3.4 GPa, the same order of magnitude as those for cortical bone (Table 9.1). The wide range of the values reported in literature can be attributed to the variations in composition of polymer and monomers, particle size, morphology, molecular weight, powder: liquid ratio, differences in amount and type of radiopacifiers, sterilisation techniques and testing regimes. The determination of mechanical properties such as flexure and shear, as well as the analysis of toughness to fracture and dynamic properties such as fatigue and impact, gives good information about the dynamic mechanical behaviour of these materials. An active implant in vivo is subjected to cyclical loading and typically can experience up to 1 million cycles per year during normal activities such as walking. The cement is considered the weak link in cemented hip implants and thus minimising cement fracture is important. Properties such as fatigue, creep and stress–relaxation describe the long term properties of cements and significantly influence the transfer of loads. The apparent fracture toughness K1C is used to denote the resistance of a porous material to crack propagation. The fracture toughness of bone cements range from 1.00 to 2.32 MPa m−1/2 and exhibit an inverse relationship with porosity and viscosity (Murphy 2002; Topoleski 1993; Vila 1999). The variability in the reported data stems from the fact that different specimen configurations, sizes, applied loads, loading rate, storage conditions, mixing methods and sterilisation techniques have been used to conduct the tests (Lewis 1999). Thus, there exists a need for a systematic study of cements with controlled standard parameters. Dynamic fatigue is more widely recognised as a cause of fracture and the usual approach followed in representing the fatigue behaviour is S–N graphs, where stress or strain data is plotted against the number of cycles to failure. Fatigue failure and fatigue crack growth have been related to one of the primary causes of the mechanism of failure in bone cements (Deb 2008; Krause 1988; Kuhn 2000; Lewis 2003). A review of the literature on fatigue of bone cements, with some typical values shown in Table 9.3 (Lewis 2003), concluded that there was a paucity of comparable data and this has been clearly identified as the lacunae that exists in the characterisation of the fatigue properties of cements. A myriad of variables influence fatigue life of cements, making it difficult to state conclusively the factors that decisively influence fatigue properties, but they include intrinsic properties such as composition that includes the type of inclusions such as the radiopacifier, activator–initiator, any reinforc-
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Table 9.3 Fatigue: mean number of stress cycles to fracture, the influence of mixing technique and inclusion of antibiotic
Cement Palacos R Simplex P Simplex P Simplex P + 2.4 g of tobramycin
Mixing method
Stress
Frequency
Hand mixing Hand mixing Vacuum mixing Hand mixing
Tension– compression Tension– compression Tensioncompression Tensioncompression
2 Max strain of 0.005 2 Max strain of 0.005 2 Max strain of 0.005 2 Max strain of 0.005
Results Nm Mean number of stress cycles to failure1 23 000 18 000 32 000 11 000
1
For all directly moulded cylindrical specimens (Lewis 2003).
ing agent, antibiotics, molecular weight of the prepolymer, and the viscosity of the mix, and also extrinsic factors such as environmental temperature, mixing methods, specimen configuration, test conditions such as the magnitude of applied stress, frequency of loading, humidity, testing in air or wet conditions and the conditioning of the specimen before testing. Ageing cements in physiological fluids is commonly carried out before testing the mechanical properties of cements with a view to mimic in vivo conditions. The uptake of fluids has been shown to adversely affect the mechanical properties resulting in a lowering of the stiffness of the cements. Although ageing cements in physiological fluids have been shown to adversely affect cements, there are only a few studies that have compared explanted cements and correlated the mechanical properties with in vitro testing conditions. However, a few studies that exist on explanted cements have failed to make a direct correlation of time with fracture toughness (Hughes 2003; Ries 2006), which may be attributed to the myriad of parameters that influence the fracture toughness of cements such as viscosity, composition (Lewis 2000), sterilisation methods (Chapiro 1962; Graham 2000; Harper 1997; Lewis 1998) and mixing methods (Lewis 2000a). The bone/cement/prosthesis interface is an important construct in the hip joint and transfer of stresses and the micro-mechanical properties are expected to influence the long-term performance of the cemented joint replacement surgery. There have been numerous studies to characterise the mechanical properties of the acrylic bone cement but relatively few studies address the properties at the interface (Arola 2006; Funk 1998). The strength of the cement/bone interface has been reported in the past albeit using overly simplified models. The micro-mechanical behaviour of the composite construct was reportedly tested by Mann et al. (2008) using non-destructive
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fully reversible tension–compression loads whilst monitoring the local micromotion of the interface that showed that the interface was more compliant than the bone or the cement. The interdigitation of the cement with the bone is an important aspect, but the roughness of the surface of the bone vary and the presence of blood, fats and fluids are also expected to affect the levels of interdigitation. Janssen et al. (2008) used a micromechanical modelling approach to study the effect of frictional, morphological and material properties on the mechanical response using a finite element approach. Although the acrylic cement does not function as a ‘glue’ or adhesive by virtue of its nature, the study by Janssen supports the fact that frictional phenomena or mechanical interlocking at the interface is important and the interfacial gaps that existed were detrimental thus suggesting that pressurisation of the cement had a clinical advantage. Attempts at directly measuring the transient and residual radial stresses generated at the stem–cement interface during the polymerisation of the cement suggest that the bone cement shrinks towards the stem and the temperature gradient between the stem and the cement can affect the direction of shrinkage in the cement (Race 2002; Wang 2004).
9.4
Radiopacifiers in bone cements
Radiological detectability of the bone cement is vital for post-operative follow up as the prosthetic components are inherently radiopaque, whereas the cement is not. PMMA-based bone cements are intrinsically radiolucent and are usually rendered radiopaque by the addition of heavy metal salts such as barium sulfate or zirconia. Among the commercial acrylic bone cements used, all contain either BaSO4 or ZrO2 particles as the radiopacifier (Kuhn 2000; Lewis 1997). The general consensus on the inclusion of radiopacifier particles is that it causes deleterious effects because it is essentially incompatible with the PMMA polymer matrix (Deb 2002). The particulate additives barium sulfate and zirconia have been linked to bone resorption (Sabokbar 1997) increased wear rates owing to particulate debris and degradation of some of the mechanical properties (Bhambri 1995). The extent of the effect on the properties is related to the amount, size and morphology of the inorganic particulates included. The addition of radiopacifiers in PMMA cements can adversely affect both the mechanical and physical properties. Barium sulfate has been shown to increase compressive strength and Young’s modulus while decreasing ultimate tensile strength (Ginebra 2002; Vazquez 1997). Zirconium dioxide, on the other hand does not alter the tensile or compressive properties (Ginebra 2002) and has also been reported to enhance the fracture toughness and fatigue crack propagation resistance significantly; however, the quantity, particle and shape also have an important role.
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Of more concern is that the radiopacifiers included in the cements have been implicated in bone resorption. The wear particles generated have a role in the pathogenesis of implant loosening and have been linked to the presence of tissue macrophages in the fibrous pseudomembrane. These cells are believed to interact with wear particles and release inflammatory cytokines detected in the pseudomembrane in retrieval studies. TNF-α, IL-1β and IL-6 have all been found in such tissues and shown to be released by macrophages in response to particulate debris. Radiopacifying agents that are compatible with the polymer matrix are alternative routes for achieving radiopacity in organic polymers. Mechanical properties of acrylic bone cements have been reported to improve by substitution of the radiopaque agent with monomers such as 2,5-di-iodo-8-quinolyl methacrylate, triphenylbismuth (Deb 2001; 2002) and copolymers of methacrylate/methacrylic monomers containing tri-iodobenzoate functional groups (Ginebra 1999). PMMA cements are also being increasingly used in vertebroplasty (Deb 2008; Kaufmann 2006) and strontium oxide as a radiopaque agent has been investigated by a number of researchers (Cheung 2005; Wong 2004a, 2004b; Zhao 2004). The studies so far conclude that strontium oxide shows promise as a radiopacifying agent and can be applied for PMMA cements for both arthroplasty and or vertebroplasty.
9.5
Antibiotic-laden bone cements
Infection continues to be one of the major complications following joint replacement surgery, particularly in relation to the occurrence of ‘super bugs’ second only to the fatal effects of pulmonary embolism. Joint replacement sites can acquire infections in many ways, from peri-operative primary seeding or contamination owing to non-sterile implant materials during or after the surgical procedure to haematogenous spreading, where the primary infection originates at a distant site and is transferred to the implant site. Currently used surgical solutions include, debridement and irrigation and methods involving excision arthroplasties which are associated with patient morbidity. Antibiotic therapies, either by systemic administration or prophylactically using antibiotic laden bone cements (ALBC) are other modalities of treatment that have contributed to the growing success of joint replacement surgery. Acrylic bone cements containing antibiotics release them gradually over time in such a way that the local levels of antibiotics exceed the minimal inhibitory concentration of the pathogens of interest. The advantage of this site-specific delivery is that the amount of antibiotics released at a local level is usually much higher than that achieved with parenteral therapy whilst allowing a lower systemic absorption. The use of ALBC reduces the incidence of infection especially if used in combination with oral administration of antibiotics (Espehaug 2002); however, the
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clinical efficacy of antibiotic-releasing bone cements is not universally accepted and the long-term exposure to low doses from antibiotic-releasing bone cements in patients may be a cause for developing antibiotic resistance (Belt 2001). One of the most widely used antibiotics in bone cements is gentamicin that has been in clinical use over 50 years and exhibits a concentration-dependent antibacterial activity (Lacy 1998). Furthermore, gentamicin in the form of gentamicin sulfate has a number of properties that make it particularly suitable for inclusion in acrylic bone cement formulations, such as wide-spectrum antibacterial activity, solublility in aqueous media, low allergenicity and thermal stability. Whilst the addition of antibiotics such as gentamicin or tobramycin is known to reduce the risk of infection (Bucholz 1981, Hanssen 1994; Scott 1999), some reports present conflicting results (Bourne 2004; Hendriks 2004; Wahlig 1980). The Norwegian Register included 22,170 total hip replacements in a study (Engesaetar 2003) that reported that antibiotic-containing cements reduced the rate of infection from 0.7 to 0.4% (p = 0.001) when comparing groups that received 24-h intravenous antibiotics prophylactically or received both 24-h intravenous treatment, along with an ALBC. Antibiotic elution from the cement matrix occurs by surface diffusion, bulk diffusion and through pores and defects from the cement mantle. The mode of mixing is of significance with particle size and uniformity of distribution governing the elution kinetics of the antibiotic. There is a rising trend of antibiotics being added manually to the cement by the surgeon and often the antibiotic is selected to suit the antibacterial activity. This practice not only has an effect on the elution of the antibiotic but also affects the mechanical performance of the cement. Although the addition of antibiotics does not seem to affect the compressive or tensile properties, few studies report the effects of inclusion of high amounts of antibiotics. It is important to review the properties of the cement that is mixed with a high amount of antibiotic as often surgeons tend to add larger amounts of antibiotic especially in cases of revision, when used as articulating spacers and when combination of antibiotics have been used (Hoffman 2005; Koo et al. 2001; Langlais 2006). He et al. (2003) investigated the addition of a range of concentrations of gentamicin in Palacos R from 2.4 to 13.0 wt%, to recommend that a maximum of 6.5 wt% of the antibiotic would maintain the compressive strength of the cements. However, cements containing higher quantities of antibiotics have been shown to exhibit a gradual decrease in their mechanical properties post immersion in body fluids. One study reported that tripling the dose of Flucloxacillin in Simplex P from 2 to 6 g caused a marked decrease in the compressive strength when tested after 4 weeks immersion in phosphate buffered saline with evidence of increased porosity and clumping of the radiopacifiers (Pelletier 2008). Such effects can lead to the catastrophic failure of cements in vivo, a risk that may be associated
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with the high amounts of antibiotics in the cement. The fact that the major volume of the antibiotic is eluted from the cement in the early stages can lead to a large concentration of voids that can cause either catastrophic failure or a more gradual degradation of the cement over time. These effects on the bulk properties are not evident in the short-term tests simply because the matrix of the cement remains undistributed whilst the antibiotic is dispersed within it. Mixing techniques such as vacuum mixing and the use of closed proprietary systems is also expected to influence the presence of voids within the cement and the release kinetics of antibiotics from a cement mantle. Vacuum mixing in general, decreases the net amount of pores within a newly formed cement mantle and, although some studies claim that vacuum mixing has little or no effect on the amount and rate of drug release, this is debatable since the release of drug is predominantly through voids and defects rather than diffusion. This observation may be related to the determination techniques used to assess drug release or the differences in surface geometries utilised. As vacuum mixing decreases the net amount of pores in a cement matrix, it is difficult to conceive that antibiotic elution remains unaffected in a denser matrix. It is the opinion of the author that no single mechanism occurs in the elution of drugs from cement matrices, but rather a combination of diffusion and loss of molecules through surface defects and bulk irregularities in the cement matrix. The amount of drug present in the cement, the composition of the acrylic cement, the presence of hydrophobic additives, the blending method of the drug within the uncured cement, the surface area of the cement mantle and porosity are factors that will influence the elution of drugs. In general, it has been reported that elution is greatest in the early phase and that only a part of the drug impregnated is released, which may suggest that any complications caused by the prolonged presence of antibiotics in the systemic circulation will be minimal; however, using clinically close models and conducting long-term recall studies will reveal a truer picture.
9.6
Composite cements
PMMA cements are classified as bioinert materials, which provide a structural support to the implant but form a barrier for direct healing owing to the presence of fibroblastic cells at the interface. The PMMA cement has often been implicated in loosening of the prosthesis and it is thought that a direct bone-bonding ability of the cement would enhance the fixation. Thus, to confer bone-bonding ability, the main approach so far has been the addition of calcium phosphates such as hydroxyapatite (HA) bioactive glasses and tricalcium phosphates in the acrylic PMMA cements to yield composite cements (Cho 2005; Dalby 1999; Mousa 1999; Shinzato 2000).
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The inclusion of particulates such as HA, A-W glass ceramics can additionally improve the mechanical properties of the cement, but it is important to bear in mind that PMMA is a brittle material and inclusion of particulate fillers is expected to further embrittle the matrix. Furthermore, in order to achieve an improvement in mechanical properties, the interfacial adhesion between the filler and the matrix is an important parameter and the use of coupling agents is an option, but the effect on bioactivity needs to be ascertained. The bone–cement adhesion may contribute significantly to the long-term performance of the surgical procedure. Potentially bioactive PMMA cements have been reported containing more than 20 wt% of wollastonite ceramics in PMMA cements with increasing bioactivity with higher contents of the ceramic. These cements were reported to form a homogeneous and thick apatite layer after 21 days of immersion in SBF (Mousa 1999; Rentería-Zamarrón 2008). Composite cements with PMMA as the organic matrix containing different fillers such as apatite–wollastanite glass ceramic (AW– GC), sintered HA and bioactive glass bead fillers of MgO–CaO–SiO2–P2O5– CaF2 have been formulated (Shinzato 2000) and compared for in vitro mechanical properties and in vivo osteoconductivity. Filler contents as high as 70% by weight were used in the study and the bending strength measurements reported showed significantly higher flexural strength for the glasscontaining cement. The glass-containing cement also exhibited higher osteoconductivity than the cements containing AW–GC or HA, making it a promising alternative. However, important data on the rheology of the cements and curing kinetics are not well described, which are essential parameters for adequate clinical handling. Another study reported the inclusion of glass spheres in PMMA cements (Vallo 2000) and reported that 50 wt% glass particles could be added with significant increases in flexural modulus and fracture toughness. The improvement in the fracture toughness was attributed to the toughening caused by particle reinforcement. In this study, the PMMA powder was replaced by the glass component, which would effectively decrease the viscosity of the mix as the monomer content would increase and, as expected, the cements yielded higher residual monomers. The shrinkage of these cements are not reported, but it is expected to be higher owing to the higher monomer content that may only be compensated to an extent by the presence of the filler. Although, the cement was reported to have better handling properties the high concentration of residual monomers is expected to be detrimental owing to the resulting higher fluid uptake. A commercial acrylic bone cement modified by the incorporation of different weight fractions of polycrystalline HA, and the influence of the filler proportion on the flow characteristics and the mechanical behaviour of the resultant composite was reported. The study highlighted the fact that there were limits to the amount of filler that could be
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incorporated in a PMMA cement without compromising the flowability of the cements. Vallo et al. (1999) reported that up to 15 wt% HA could be added with enhancements in flexural modulus and fracture toughness, but higher concentrations limited the workability. The cellular interaction of a PMMA bone cement containing 17.5% by weight of hydroxyapatite was reported by Dalby et al. (1999). The in vitro biological interaction at the cellular level was studied using primary human osteoblast-like cells with cell proliferation and growth being assessed by measurement of total cellular DNA and tritiated thymidine ([3H]-TdR) incorporation. The results showed that HA/PMMA was a better substrate for human osteoblast-like cells than PMMA and it showed increased cell proliferation and alkaline phosphatase activity. The cells were also reported to attach preferentially onto the HA particles present on the composite surface. PMMA-based composites having a range of concentrations (0, 30, 40 and 50 wt%) of a glass filler, namely Ca3(PO4)2–SiO2–MgO were recently reported by Lopes et al. (2009). The polymer used was a copolymer of poly(methylmethacrylate)–co–(ethylhexylacrylate) (PMMA–co–EHA) matrix filled with a glass that was reportedly shown to exhibit in vitro bioactivity via the formation of a hydroxyapatite-like layer on interaction with simulated body fluid (Lopes 2008). The effect of filler proportion (0, 30, 40 and 50 wt%) on the bending properties showed that a maximum flexural strength of 29 MPa coupled with an elastic modulus of 1.1 GPa was obtained at an intermediate filler concentration (30 wt%). Hydroxyapatite, chitosan and PMMA composite have also been proposed as alternative bioactive bone cements to PMMA (Kim 2004). The composites show higher water uptake, weight loss and porosity with lower compressive strength and modulus than PMMA cements. A higher porosity with higher water uptake is expected to lead to plasticisation of the matrix with time and the inferior mechanical properties may lead to failure. The porosity may allow more bone ingrowth but the effect on the fatigue properties needs thorough examination before clinical use. The exothermic temperature of the composite cement was lower than PMMA cements with a higher intrusion capacity, which is an advantage. In vitro and in vivo tests showed that the chitosan composites were better than PMMA cements and have been proposed as a potential replacement for PMMA cements. Other researchers have explored the inclusion of titania in PMMA (Goto 2005) and reported that cements containing 50 to 60% silanised nano-sized titania have similar compressive strengths to PMMA cements and exhibit direct apposition to bone with significantly higher osteconduction in vivo as experimented in rat tibiae. Multiwall carbon nanotubes, owing to their large surface to volume ratio, have also been investigated for the augmentation of PMMA cements and the data has shown that an enhanced fatigue life can be achieved (Marrs 2006).
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Another approach has been the design of bioactive bone cements based on organic–inorganic hybrids via the modification of the organic matrix. (Miyazaki 2008). PMMA bone cements were modified by the incorporation of calcium silicate gels from methacroyloxypropyl trimethoxysilane and other calcium salts. The alkoxysilane with the calcium salts renders the PMMA bioactive and it becomes capable of interacting with simulated body fluids to deposit apatite, thus enhancing the bone bonding ability. It is claimed that these cements maintain excellent handling properties whilst exhibiting bone bonding ability. PMMA is the principal component of an acrylic bone cement that has been used clinically over 50 years with success. However, the drawbacks of these cements are well known and have thus warranted improvements. A variety of materials have been added to augment the mechanical properties of these cements but none of them have been so far deemed successful enough to replace the PMMA acrylic bone cement. The reinforcement of PMMA includes the addition of polymer to metals fibres and particulate fillers that have been explored for augmentation of the cement in a bid to improve fracture toughness and fatigue characteristics, mechanical properties and enhance bioactivity via the inclusion of bioactive fillers. However, there have been problems with the modifications attempted thus far such as an increase in viscosity, compromised flow and clinical handling, debonding of the matrix, alteration of the setting kinetics, fluid uptake and plasticisation of the matrix. Thus, so far none of these modifications have led to an alternative to the PMMA acrylic bone cement. An additional cause has also been the reticence of surgeons to replace the use of the PMMA cement that has an acceptable and established clinical performance, which also limits the industry in exploiting the potential of some of the augmented cements.
9.7
Conclusion
PMMA is the most widely used acrylic biomaterial for anchoring prosthesis to the contiguous bone in cemented arthroplasties. The clinical performance of these arthroplasties is related to a myriad of variables, both exogenous and endogenous; however, the quality of the bone cement is of paramount importance. Porosity has been deemed to be one singular property that has been considered to affect a large number of mechanical properties, but a comparison of reported data is inconclusive owing to the variability in the composition, mixing, dispensation method, specimen preparation and conditions of storage. It is recognised that the performance of the cement is related to both material and structural behaviour. In addition complex load patterns at the cement–bone–implant interface exist and an understanding of the stress distribution in the various sections is important especially at
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the implant–cement interface, as it has been implicated in the loosening of the prosthesis. Finite element analysis, three-dimensional modelling and the determination of the stress distribution in a model mimicking in vivo conditions will help elucidate the exact requirements for improved cements. With the increasing need for incorporating specific antibiotics or combination of antibiotics, systematic studies on elution kinetics and mixing methods are still required to provide a more general trend on the effect of antibiotics in cements.
9.8
References
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10 Composite materials for replacement of ligaments and tendons L. A M B R O S I O, A. G L O R I A, National Research Council, Italy; and F. C AU S A, University of ‘Magna Graecia’, Italy
Abstract: In this chapter, the role of composite biomaterials as a unique type of material able to reproduce the complexity of the hierarchical structure and the peculiar mechanical properties of ligaments and tendons is discussed. The fundamentals of fibre-reinforced composite materials are reviewed, highlighting critical aspects in designing devices for ligaments and tendons replacement. The various approaches for replacement and regeneration are explored, with the aim of benefiting from composite materials science. Prospective and future challenges for composite materials for ligaments and tendons replacement and regeneration are outlined. Key words: fibre-reinforced composite materials, filament winding technology, mechanical behaviour, viscoelastic properties.
10.1
Introduction
The healing of damaged ligaments and tendons is a controversial issue. The use of artificial prostheses to repair and stabilize knee ligaments is preferred to autologous replacement both to avoid sacrificing the patient’s own tissues as well as to significantly decrease rehabilitation time. Over the past years, a variety of artificial devices have been proposed and used in clinical practice. These artificial ligaments are composed of threads of biocompatible polymers woven in different directions to form highly interlaced macroscopic structures. However, there is no device to-date that sufficiently matches the mechanical behaviour of natural tissue to be considered for long-term applications. Amongst the drawbacks related to the current synthetic replacements for ligaments and tendons are wear, degradation of the materials, chronic inflammation and low fatigue strength. Clinical and biomechanical investigations have been increased to analyze the high failure rate and clinical problems owing to the use of prosthetic ligaments and to design alternative structures. The challenge remains to duplicate the complex structure and mechanical properties of the natural ligaments and tendons, by using the benefits of composite materials science. Moreover, even though great improvements have been obtained in skin regeneration, 234
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the regeneration of diseased tissues, such as ligaments and tendons, is far from ideal. In order to better understand the parameters involved in the design of alternative prostheses and scaffolds, a description of the biology, structure, mechanical properties and materials used is needed.
10.2
Ligaments and tendons: tissue biology and anatomy
Tendons and ligaments are dense connective tissues constituted by a protein phase (collagen and elastin) and a polysaccharide phase (proteoglycans). The overall mechanical properties are determined by the relative amount of the two phases, geometrical factors and the conformation and orientation of the individual constituents. The main function of tendons is to transfer to bones the force resulting from contraction of muscles, thus generating the movement of the articulations. Ligaments stabilize the joints, preventing abnormal movements (O’Connor and Zavatsky, 1963). The need to design and manufacture artificial prostheses to repair or replace damaged ligaments has become important in recent years because of the increase of tendons and ligaments injuries (e.g. anterior cruciate ligament). A careful analysis of the structure and the inner organization explains the observed mechanical behaviour of these tissues (Arnoczky et al., 1963). In fact, the design of artificial prostheses for tendons and ligaments cannot be separated from a profound understanding of the structure and the functionality of the organs that need repair or replacement (Iannace et al., 1995). In tendons, the collagen fibres are aligned along the axial direction, conferring anisotropy to the material such that the material is stiffer in the direction of the fibres. In ligaments, on the other hand, collagen fibres have a wave-like organization, thus contributing to the ligament properties only at higher values of deformation, that is after their alignment in the loading direction (Amiel et al., 1984). In particular, the anterior cruciate ligament (ACL) is a dense and highly organized cable-like tissue composed of collagens (types I, III and V), elastin, proteoglycans, water and cells. Ligaments have a hierarchical structure with different levels of organization, including collagen molecules, fibrils, fibril bundles and fascicles, that run parallel to the long axis of the tissue. The collagen molecule, a glycine-rich triple helix, is the main constituent of ligaments. Collagen molecules assemble sequentially into microfibrils, subfibrils and fibrils (20–150 nm in diameter) before forming fibres (1–20 μm in diameter) that mutually crosslink to make up a further subfascicular unit (100–250 μm in diameter). These subfascicular units are surrounded by a loose band of connective tissue known as the endotenon. The collagen fibrils in ligaments naturally display a periodic change in direction called
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crimp pattern. In the ACL, this crimp pattern repeats every 45–60 μm (Silver, 1994). The ligament is surrounded by a sheath of vascularized epiligament (Silver, 1994). An additional level of structure exists in the ACL, the collagenous network being twisted by approximately 180 ° between the femoral and tibial attachment sites. The ACL also contains anteromedial and posterolateral bands. The entire continuum of fascicles is surrounded by the paratenon, a connective tissue cover similar to, but much thicker than, the epitenon (Amiel et al., 1990). Cellular components, mainly fibroblasts, not only synthesize fibrillar collagen but also enzymatically break down and remove old collagen as part of a renewal process. The human ACL has an average length of 27–32 mm, and a cross-sectional area of 44.4–57.5 mm2 (Vunjak-Novakovic et al., 2004).
10.3
State of the art of proposed devices for replacement of ligaments and tendons
Approximately 200 000 Americans required ligament reconstructive surgery in 2002, with the total expenditure exceeding US$5 billion. This was further augmented by costs associated with loss of productivity, healthcare and social benefits. Reconstruction of the ACL is the most significant procedure in this category (Vunjak-Novakovic et al., 2004; Pennisi, 2002). At approximately one case per 3000 Americans (in 1999), the incidence of ACL injuries is high and has increased over the years (Fu et al., 1999). As it occurs in the case of bone, three options have been tried for the repair or replacement of damaged ligaments using biological substitutes, namely autografts (patellar tendon with bony attachments or two of four hamstring tendons harvested from the patient at the time of surgery) (Fu et al., 1999; Weitzel et al., 2002), allografts (frozen ligaments with bony attachments) (Noyes et al., 1984) and xenografts (bovine tissue cross-linked with glutaraldehyde) (Woo and Buckwalter, 1988). In particular, autografts have produced the most satisfactory long-term results and are referred to as the ‘gold standard’. Donor site morbidity remains the limiting factor of patellar tendon grafts as it is often associated with pain, muscle atrophy and tendonitis, resulting in prolonged rehabilitation periods, which cause an immunological foreign-body response that hinders tissue remodelling (Laurencin et al., 1999; Yahia et al., 1997). Many efforts have been made to use synthetic materials in ligaments replacements. Several non-degradable synthetic materials have been proposed for ACL repair, including polyethylene terephthalate (StrykerDacron and Leeds-Keio ligaments), polypropylene (Kennedy Ligament Augmentation Device) and poly(tetrafluoroethylene) (Gore-Tex). Some of these synthetic grafts (Gore-Tex, Stryker, Kennedy Ligament Augmentation Device) have been conditionally approved by the FDA for testing and
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augmentation but are not recommended for primary ACL repair because they are unable to duplicate the mechanical strength and surface properties of the complex ACL (Friedman et al., 1985, Yahia et al., 1997). The Stryker-Dacron ligament prosthesis is made of Dacron tapes wrapped in a Dacron sleeve (Laurencin et al., 1999; McCarthy et al., 1993). The LeedsKeio ligament is a woven porous made up of polyethylene terephthalate and attached to woven tapes. It is designed to allow the ingrowth of new tissue (Fujikawa, 1998; Silver, 1994). The Kennedy Ligament Augmentation Device consists of a polypropylene cylindrical prosthesis with a diamondbraided construction. It is designed to be implanted in conjunction with biological grafts such as the patella tendon (Silver, 1994). The Gore-Tex prosthesis is composed of an expanded poly (tetrafluoroethylene fibre) that is wound into loops, which are joined together to form a braid (Bolton and Bruchman, 1985; Olson et al., 1988). Although these synthetic devices initially supply the function of the ligaments that they replace or protect the ligament that they augment, they fail over time because they cannot reproduce the mechanical behaviour of the ligament. Repeated elongation of these devices leads to permanent deformation at the points of stress. Contact with sharp edges of the bone tunnel causes abrasions that weaken the implant and create debris that can cause synovitis in the joint. Woven prostheses face the additional problems of axial splitting, low tissue infiltration, low extensibility and abrasive wear. Eventually, these implants fail due to fragmentation, stress shielding of new tissue, fatigue, creep and production of wear debris (Laurencin et al., 1999; Laurencin and Freeman, 2005; Smith et al., 1993; Woo et al., 1994). Clearly, alternative ACL replacement and reconstruction methods would be advantageous.
10.4
Fibre-reinforced composite materials: fundamentals and technology
Over the past years fibre-reinforced polymeric composites have attracted much attention as a new type of material for designing biomedical devices, such as those for ligaments and tendons. Even if the biocompatibility of a prosthesis represents a necessary condition, the healing and remodelling of a natural tissue and the successful implantation of a device are strongly affected by the stress field induced on the tissue (Gershon et al., 1990). Thus, the mechanical properties of the implant play an important role and vary with the characteristics of the replaced tissue. The necessity of matching the mechanical behaviour of the natural tissue has driven the research toward fibrous composite materials and their anisotropy, which is an important feature of most biological tissues. Another important aspect shown by fibrous soft biological tissues, such as ligaments and tendons, is related to their non-Hookean response under physiological
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loads and to the J-shaped stress–strain curve, which they exhibit (Gershon et al., 1990). A great advantage of using composite design in the preparation of soft tissue devices is that they provide a wide selection range for the compliance and eventually the ability to control the chosen value. Basically, the compliance along the x axis 1/Ex of an orthotropic lamina with its principal axes oriented at angle θ with respect to the coordinate axes (Fig. 10.1) is given by the following equation (Gershon et al., 1990; Hull, 1981): cos4 θ sin 4 θ 1 ⎛ 1 1 2ν = + + − LT ⎞ sin 2 2θ 4 ⎝ GLT Ex EL ET EL ⎠
(10.1)
where Ex represents the Young’s modulus in the x direction, EL, ET, GLT and nLT are the Young’s moduli, the shear modulus and Poisson’s ratio, respectively, of the composite lamina with L and T denoting the principal material axes, longitudinal and transverse. Thus, compared with a homogeneous isotropic material, a composite structure may offer a much higher level of design sophistication using a choice of material combinations with a wide selection range of constituent volume proportions. If a choice of relative constituent properties and concentrations determines the basic properties of a lamina, the orientation angle θ is also important since it represents another factor to the design options as expressed in equation (10.1). Many implications of equation (10.1) with advanced composite materials reinforced with high modulus fibres are usually reported in composite material textbooks (Argawal and Broutman, 1980; Mallick, 1988). Clearly, y T L
q x
10.1 Schematic representation of an orthotropic lamina with its principal axes (longitudinal L and transverse T) oriented at angle θ with respect to the x and y axes.
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at θ = 0 °, Ex is equal to EL, whereas at θ = 90 °, Ex is equal to ET. For 0 ° < θ < 90 °, the trend of Ex strongly depends on the elastic properties of the lamina and may be determined by investigating the first and second derivatives of equation (10.1). At intermediate values of θ the Young’s modulus Ex can either be >EL or <ET, depending on the value of the combination of EL, ET, GLT and nLT (Gershon et al., 1990). It may be possible to express the conditions for a maximum and minimum for the function Ex = f(θ). For a given set of elastic properties, it is possible to obtain the location of the maximum/minimum by equating the first derivative of equation (10.1) to zero. All of these considerations suggest that the possibilities of obtaining a minimum or a maximum for Ex (which means a maximum or a minimum for the compliance, respectively) have significant implications for the design options of soft composite materials beyond the inherent merits of composite design (Gershon et al. 1990). From a materials science point of view, most of new developments in designing high performance and multifunctional materials are obtained by combining materials and technologies, such as filament winding, which are already known. Accordingly, the design of soft fibre-reinforced composite using filament winding techniques involves many parameters and considerations.
10.4.1 Principles of soft composite design Filament winding is a technology for producing axially symmetric composite structures by helically winding fibres on a suitable mandrel through the combination of the translation and rotation of the transverse carriage and the mandrel, respectively. During the winding of the fibres, the diameter of the system increases and the winding angle changes (Iannace et al., 1995). A schematic representation of the variation of the winding angle as a function of the increasing radius of the composite cylinder is reported in Fig. 10.2, whilst Fig. 10.3 shows a unit element deforming under an applied load. The distribution of the angle in the entire system can be calculated from the speed of the transverse carriage and the rotation of the mandrel through the following expressions (Fig. 10.2): AC = ν t CB = ωπ r1t CB′ = ωπ r2 t tan θ1 = (ωπ r1 ) ν tan θ 2 tan θ1 = r2 r1
(10.2)
where n is the speed of the transverse carriage, ω the rotational speed of the mandrel, t the time, and θ the angle of the helix around the mandrel (Iannace et al., 1995).
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Fibres
B
B′ pr1
q1
pr2
q2
A
C
10.2 Scheme of the soft fibre-reinforced structure and variation of the angle as a function of the increasing radius during the winding.
ds1 A′ q2 A q1
B
B′
C
10.3 Triangular unit element as a part of an orthotropic lamina deforming under the applied load.
As for the modelling of this complicated system, it is useful to consider the composite structure as an ideal system where the winding angle is constant and equal to the mean of the initial and final angles. The overall composite structure is therefore modelled by a lamina with the fibres oriented as ±θ (Iannace et al., 1995).
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If a small stress dσ1 is applied on A in the direction of AC, the system changes its geometrical configuration (Fig. 10.3) as a consequence of the soft matrix. The ‘fibre’ AB becomes A′B′ and its length Lf is assumed to be constant. If another stress dσ2 is applied on A′ in the same direction as dσ1, the point A′ will be on A″, aligned on the segment A–A′, and so on until the winding angle reaches a critical value depending on Φf and θ. It has been assumed that this critical angle can be calculated on the basis of the maximum deformation that the system can support owing to the percentage of matrix present in the composite (Iannace et al., 1995). The unit element represented in Fig. 10.3 can be considered as part of an orthotropic lamina that can be continuously deformed if the matrix itself allows this deformation. The decrease in the area of the triangle element results in a closer movement together of the fibres. This is accompanied by a simultaneous compression of the matrix. The maximum ideal deformation is therefore characterized by:
α ⋅ φm =
Ai − Ac Ai
0 ≤α ≤1
(10.3)
when α = 1. Ac represents the area of the unit element ABC shown in Fig. 10.3, when the system reaches its critical configuration, Ai is the same area in the initial geometrical configuration and Φm is the volumetric fraction of the matrix. The empirical parameter α, which has to be evaluated by fitting the experimental results with the theoretical predictions, takes into consideration the limited deformability of the matrix and has to be considered as a function of the matrix rigidity. In fact, if α = 0, the matrix does not allow the deformation process shown in Fig. 10.3 and the mechanical response of the system does not present the characteristic ‘toe’ region but rather the typical stress–strain behaviour of rigid composites. The final geometrical configuration that the system reaches is a function of the initial condition θi, the volumetric fraction of matrix and the parameter α:
θc =
1 arcsin [(1 − αφm ) sin ( 2θ i )] 2
(10.4)
The deformation of the overall composite structure can be written as (Fig. 10.3): dε i =
cos θ i +1 −1 cos θ i
(10.5)
In the equation (10.5) the deformation of the fibres is neglected and their length Lt is considered as constant (Iannace et al., 1995).
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At each step i, Young’s modulus Ei may be evaluated through the classical lamination theory of composite laminates (Gershon et al., 1990; Mallick, 1988) according to equation (10.1), where Ei is referred to the direction of the applied stress, EL, ET, GT and nLT are Young’s moduli, the shear modulus and the Poisson ratio, respectively, of the composite lamina, whilst L and T denote the principal material axes, as previously discussed. The elastic properties of the lamina along the principal axes were calculated with the following expressions (Iannace et al., 1995; Jones, 1975; Mallick, 1988). EL = Ef φf + Emφm Ef Em ET = Ef φm + Emφf Gf Gm GLT = G f φm + Gmφf
(10.6)
At each step i the total strain εTi and the total stress σTi applied on the system will be:
ε Ti =
i cos θ i − 1 = ∑ dε j cos θ1 j =1 i
σ Ti = ∑ dσ j
(10.7)
j =1
The last two expressions allow the stress–strain curve to be drawn by knowing the initial condition of the system (i = 1) and the properties of the materials. However, these equations must be solved through the parameter θi because of the impossibility of explicitly defining a function σ = f (ε). When the system reaches the critical configuration characterized by an angle equal to θC, the elastic modulus of the composite becomes constant and equal to EC (Iannace et al., 1995). In the light of what has been said, such an approach represents a powerful tool to design fibre-reinforced composites as devices for tendons and ligaments replacement, enabling materials scientists to improve their mechanical performance.
10.5
Composite materials for tissue replacement and as scaffolds for tissue engineering
The design of artificial tendons and ligaments prostheses must simulate the natural structure providing a similar mechanical behaviour (Ambrosio et al., 1998; De Santis et al., 2004; Iannace et al., 1995). Adequate mechanical properties are important for the implant function. It is extremely difficult to find materials with appropriate biological and mechanical properties. A variety of materials and structures have been proposed for clinical use.
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However, most of them do not perform adequately and cannot be considered for long term applications (Ambrosio et al., 1994; Amis and Kempson, 1999; Durselen et al., 1996; Duval and Chaput, 1997; Turner and Thomas, 1990). A composite structure (fibre/matrix) can be suitable to realize an artificial tendon or ligament that mimics the behaviour of natural ones. Thus, a great effort has been devoted to develop several fibre-reinforced composite structures based on polymer matrix reinforced with fibres as ligaments and tendons substitutes, benefiting from composite materials theory and technologies (De Santis et al., 2004). Work on the replacement of natural tendons with composite prostheses has been reported. Absorbable polymer–carbon systems were investigated by Parson et al. (1983). Stol et al. (1985) described results obtained from poly(2-hydroxyethyl methacrylate) (PHEMA)/collagen composite tendons. Studies on a composite tendon, consisting of PHEMA reinforced with poly(ethylene terephthalate) (PET) fibres, were also reported (Kolarik et al., 1981; Migliaresi and Nicolais, 1980). The development of a biodegradable composite artificial tendon prosthesis, composed of a water-swollen PHEMA/poly(caprolactone) (PCL) blend hydrogel matrix reinforced with poly(glycolic acid) fibres was presented by Davis et al. (1991). The prosthesis was composed of fibres wound in a hydrogel-based matrix and oriented in such a way as to obtain the desired mechanical behaviour. Consequently, several developments have been reported in the design of soft composite structures capable of emulating the complex behaviour of natural tendons and ligaments. The effect of the microscopic structure on the macroscopic mechanical properties of a fibre-reinforced soft composite was investigated. A composite material consisting of a soft hydrated matrix (PHEMA) reinforced with a hydrophobic fibre phase [poly(ethylene terephthalate) (PET)] wound into a helix was manufactured and then characterized. The composite was made by reinforcing a soft hydrogel matrix with PET fibres wound into a helix using a filament winding machine. The soft matrix was obtained by radical polymerization starting from a prepolymer solution composed of the monomer 2-hydroxyethyl methacrylate (HEMA), ethylene glycol as plasticizer (molecular weight 300), ethylene dimethacrylate (EDMA) as crosslinking agent, and 2,2-azobisisobutyronitrile (AIBN) as initiator. The bundled fibres were wound at one of two fixed angles onto a polyethylene tubing support. To enhance fibre–matrix adhesion, the fibres were coated with the prepolymer (or monomeric) matrix solution before winding. The composite was then assembled by inserting the PET fibre tube in a cylindrical Teflon mould and pouring the prepolymer solution over it. The polymerization reaction was carried out at 90 °C for 1 h. Finally,
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(a)
(b)
(c)
(d)
10.4 (a) Scheme of a filament wound structure; (b) x-ray radiograph of a fibre-reinforced structure obtained through filament winding technique; (c) photograph showing two hollow composite structures; (d) composite structure in the form of a tube having two ends to insert the ligament prosthesis into the bone tunnel.
the polyethylene support was removed to obtain a hollow composite cylinder. These composite prostheses can be represented by continuous fibres helically wound to form a hollow composite cylinder (Fig. 10.4). As described in 10.4.1, the system may be simplified in a unit element which can be considered as part of an orthotropic lamina and its elastic properties are given by the composite laminate theory (Iannace et al., 1995). Results from static tensile tests performed on the composite structures at 37 °C in physiological solution are reported in Fig. 10.5. Figure 10.5 clearly shows tensile stress–strain curves of the composite prostheses realized with two different winding angles (19 ° and 35 °) of the fibres compared with the behaviour of natural tendons and ligaments. Static stress–strain curves of composite prostheses with two different winding angles (19 ° and 35 °) show that both curves present the initial toe
Composite materials for replacement of ligaments and tendons 70
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ACL composite prosthesis ACL Tendon composite prosthesis Tendon
60
Stress (MPa)
50 40 30 20 10 0 0
0.1
0.2
0.3
0.4
0.5
0.6
0.7
0.8
Strain (mm mm–1)
10.5 Tendon fibre-reinforced prosthesis with winding angle of 19 ° (䊏) and ACL fibre-reinforced prosthesis with winding angle of 35 ° (䊉) compared with the stress–strain behaviour of Achilles’ tendon (䊐) and ACL (䊊).
regions (low modulus region) followed by a progressive increase of the modulus up to the linear region that endured the failure. The higher winding angle resulted in a more extensive toe region. It is notable that the fibre composition in the two composites were identical and, therefore, the difference between these two curves is merely the result of the structural organization of the PET fibres derived from the different winding angles (Ambrosio et al., 1998a, 1998b). The mechanical behaviour of the composite prostheses is compared in Fig. 10.5 with the stress–strain response of a rabbit anterior cruciate ligament (ACL) whose data were taken from previous works (Iannace et al., 1995; Netti et al., 1996). The trend of the curves is similar to that shown by the natural system. The composite material having a winding angle equal to 35 ° displays the closest behaviour. The low-modulus region is related to the properties of the matrix (PHEMA) and its interaction with PET fibres, while the high-modulus region and final properties are mainly the result of the properties of the fibres. Accordingly, the contribution of the fibres is related to the winding angle (Ambrosio et al., 1998b). However, since natural ligaments and tendons present a viscoelastic behaviour, a proper design of artificial prostheses should take into consideration time-dependent contributions (Ambrosio
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Biomedical composites 250 Dynamic amplitude: 5 N Storage modulus E′ (MPa)
200
1 Hz 10 Hz
150
100
50
0
0
20
40
60
80
Static load (N)
10.6 Results from tensile dynamic-mechanical measurements performed on fibre-reinforced ligament prosthesis: storage modulus E′ versus static load at two different frequencies, 1 Hz (䊊) and 10 Hz (䊐).
et al., 1998a, 1998b; Netti et al., 1996). Thus, a viscoelastic analysis was also carried out on PHEMA/PET composite structures at 37 °C in physiological solution, by imposing a sinusoidal load and monitoring the response. The dynamic behaviour can be characterized by just two parameters, the storage or elastic modulus E′ and the loss or viscous modulus E″. Storage and loss moduli are correlated with the energy stored in the material and to the dissipated energy, respectively. These parameters provide an indication of the viscoelastic behaviour. Owing to the non-linear characteristic stress–strain curves, the viscoelasticity of such composite prostheses and natural structures cannot be interpreted by the simple theory of linear viscoelasticity (Ambrosio et al., 1998b; Netti et al., 1996). Consequently, tensile dynamic–mechanical tests were performed in a load control mode at an increasing level of static (or mean) load, ranging from 10 to 90 N with a dynamic amplitude of 5 N, in a wide range of frequencies (Ambrosio et al., 1998b). Viscoelastic analysis shows that the storage modulus E′ increases with frequency and static load (Fig. 10.6). In particular, for natural tissues, the trend of E′ with frequency and static load reflect the non-linear behaviour of the composite structure (Netti et al., 1996). The modulus increases until a constant value is reached, according to the stress–strain curves obtained from static tensile tests (Ambrosio et al., 1998b).
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16
Loss modulus E″ (MPa)
14 12 10 Dynamic amplitude: 5 N
8 6
1 Hz 10 Hz
4 2 0
0
20
40
60
80
Static load (N)
10.7 Results from tensile dynamic-mechanical measurements performed on fibre-reinforced ligament prosthesis: loss modulus E″ versus static load at two different frequencies, 1 Hz (䊊) and 10 Hz (䊐).
Figure 10.7 highlights the dependency of the loss modulus E″ versus static load at two different frequencies (1 and 10 Hz). When a small static load is employed, the loss modulus increases with the frequency and this behaviour progressively changes increasing the static load. In fact, when higher values of the static load are reached the loss modulus begins to decrease with frequency. The variation of the viscoelastic response observed during the increase of the applied load results from a modification of the contribution that matrix and fibres make towards the overall dynamic–mechanical response (Ambrosio et al., 1998b). At low static load levels, the increase of loss modulus with frequency is typical of rubbery materials and this behaviour can be related to the rubbery hydrogel matrix. In this condition, the PET fibres make a poor contribution to the overall viscoelastic behaviour, but this contribution becomes more and more significant by increasing the static load. Moreover, the decrease of loss modulus with frequency, shown at higher static loads, is characteristic of glassy materials and this suggests that in these conditions the contribution of the fibres is relevant (Ambrosio et al., 1998a, 1998b). The composite prosthesis is therefore subjected to gradual transition from a rubbery-like to a glassy-like behaviour during the loading process and this allows the complex viscoelastic behaviour observed in natural ligaments and tendons to be reproduced (Netti et al., 1996).
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Basing on these principles, composite structures made up of polyurethan matrix (HydroThaneTM) reinforced with PET fibres were also designed and realized by filament winding to model the morphology and mechanical properties of natural ligaments and tendons (De Santis et al., 2004). The mechanical behaviour of HydroThaneTM reinforced with PET fibres was then investigated (De Santis et al., 2004). In conclusion, the stress–strain behaviour of natural ligaments and tendons can be reproduced by hydrophilic composite structures. The tensile response of such fibre-reinforced prostheses may be modulated by varying the winding angle of the fibres, which determines the extent of the toe region and the mechanical response in the linear region. The behaviour of these soft composite prostheses can be controlled by a suitable design of the hydrophilic fibre-reinforced structures realized using the filament winding technique; the control of the winding angle of the fibres thus allows a wide range of mechanical properties to be obtained (Ambrosio et al., 1998b). Furthermore, dynamic–mechanical measurements suggest that the viscoelastic properties of natural ligaments and tendons can be reproduced by selecting an appropriate matrix and an opportune geometrical design of the soft composite prostheses (Ambrosio et al., 1998b). However, there is also a great interest in tissue-engineered solutions to ligaments and tendons injuries. Tissue engineering may be defined as the application of biological, chemical and engineering principles toward the repair, restoration or regeneration of living tissue using biomaterials, cells and factors alone or in combinations (Cooper et al., 2005). Ideal ligament and tendon scaffolds should possess suitable mechanical strength. They should be biodegradable, porous, biocompatible and able to promote the formation of ligamentous tissue. Several studies have already proposed potential ACL scaffolds using collagen, silk, biodegradable polymers and composite materials (Altman et al., 2002; Cooper et al., 2005; Dunn et al., 1995; Jackson et al., 1994, 1996). The architecture of the tissue-engineered scaffold is an important design parameter that can modulate biological response and long-term clinical success of the scaffold; for example, it has been reported that a minimum pore diameter of 150 mm is suggested for bone and 200–250 mm for soft tissue ingrowth (Konikoff et al., 1974; von Recum, 1986; Yahia et al., 1997). Overall scaffold porosity can modulate the functionality and gross cellular response to the implant. The presence of pore interconnectivity extending through an implant increases the overall surface area for cell attachment, which in turn can enhance the regenerative properties of the implant by allowing tissue in-growth into the interior of the matrix (Cooper et al., 2005). The FDA approved the use of the poly-(α-hydroxy ester)s [polylactic acid (PLA), polyglycolic acid (PGA) and copolymers, polylactide-co-glycolide (PLAGA)] for several clinical applications and they were investigated for
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use in tissue engineering (Cooper et al., 2005; Wise et al., 1995; Laurencin et al., 1999; Morgan and Yarmush, 1999). The growing interest on the use of biodegradable materials is because these materials do not elicit a permanent foreign body reaction, since they are gradually reabsorbed and replaced by natural tissue. In the long term, fatigue properties of the material may be less of a concern as the scaffold is eventually replaced by natural tissue. Therefore, PLAGA fibres, owing to their well-documented biocompatibility, biodegradability, and extended clinical use as sutures and fixations devices, were chosen for study as part of a tissue-engineered scaffold. Cooper et al. (2005) developed a novel PLAGA scaffold based on a 3D fibrous hierarchical design, utilizing custom braiding techniques, which permits controlled fabrication of substrates with a desired pore diameter, porosity, mechanical properties and geometry. Such a scaffold would provide the newly regenerating tissue with a temporary site for cell attachment, proliferation, and mechanical stability. This 3D braided PLAGA fibre scaffold was made up of three regions: femoral bone tunnel attachment site, ligament region, and tibial tunnel attachment site. The attachment sites for the bone tunnels showed a lower porosity and smaller pore size then the ligament region. This pre-designed heterogeneity in the grafts was aimed to promote integration of the graft with bone tissue and resist the abrasive forces within the bone tunnels. The advantages of this system compared with other systems were controlled porosity throughout the scaffold, which is lacking in most ACL artificial implants. The 3D braiding system allowed for custom production of scaffolds with mechanical properties similar to those of natural ACL tissue in order to overcome problems of stress shielding during tissue in-growth. Moreover, the intertwining of the fibres within the 3D braid prevents total catastrophic failure of the scaffold owing to a small rupture. Three-dimensional braiding is defined as a system where three or more braiding yarns are used to form an integral braided structure, with a network of continuous filament and yarn bundles with fibrous architecture oriented in various directions. Three-dimensional braiding systems can produce thin and thick structures in a wide variety of shapes through the selection of yarn bundle size (Cooper et al., 2005; Ko and Pastore, 1985; Ko, 1987; Ko et al., 1988, 1989). The results of this study demonstrated that processing parameters such as braiding angle can be manipulated in order to increase or decrease porosity and mode pore diameter. This is critical to the development of tissue-engineered ligaments because there is an optimal pore size that must be created in order to promote tissue in-growth. These looked similar to what would be expected of natural ligament tissue. When the same number of yarns was used for the rectangular and circular braids the circular braid geometry showed a significant increase in maximum tensile load (Cooper et al., 2005).
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In addition to scaffold architecture and degradability, cell source and cellular response play also an important role in ACL tissue engineering. In this study, the primary criteria for cell selection was based on whether the alternative cell source could reproduce or mimic the response of native ACL cells when exposed to the designed replacement scaffold. In their work, Cooper et al. (2005) also performed an in vitro assessment of scaffold biocompatibility, where cell attachment, growth, and long-term matrix elaboration by primary ACL cells were compared with those of a murine fibroblast line. The primary criteria for cell selection were based on whether the alternative cell source could reproduce or mimic the response of native ACL cells when exposed to the designed replacement scaffold. A composite scaffold for ACL tissue engineering made up of a bioactive porous matrix reinforced with biodegradable fibres was also proposed (Guarino et al., 2007). A hyaluronan-based material, HYAFF11®, which is a benzyl ester of hyaluronic acid, was chosen as matrix and PLLA fibres were used as reinforcement. These fibres were helically wound by using the filament winding technique. A porosity range from 100 to 500 μm was obtained through the solvent casting/salt leaching technique. It is well known that hyaluronan derivatives are bioactive, effectively reproducing the extracellular matrix domain. In particular, the benzyl ester of hyaluronic acid showed a good biocompatibility and the capability to promote cell adhesion and proliferation (Campoccia et al., 1998). Preliminary studies about 3T3 mouse fibroblasts adhesion and proliferation on the proposed scaffolds provided interesting results. (Guarino et al., 2007).
10.6
Conclusions and future trends
It was previously believed that artificial biomaterials had to be designed to provide a high strength associated with a high modulus of elasticity at low strain levels. It was attempted to achieve this combination of properties through the use of metals, ceramics and plastics with a relatively high strength as biomaterials (Gloria et al., 2007; Shikinami et al., 2004). However, in contrast to most artificial materials, soft biological tissues are characterized by a large amount of strain before failure, and they are flexible and tough, showing a high strength (Shikinami et al., 2004; Gloria et al., 2007). Accordingly, the design of new high performance prostheses has been proposed by using composite materials science and technology. The adopted approach has evidenced the ability to design soft fibre-reinforced prostheses, PHEMA/PET and HydroThaneTM/PET, which are able to match the mechanical properties of natural ligaments and tendons, taking into account their strength and their compliance.
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Although the need for high quality implants will continue to drive the research toward the design of multifunctional fibre-reinforced prostheses, an alternative method for the treatment of ligaments and tendons injuries involves the use of tissue engineering approach. As for the tissue engineering approach, the fibre-reinforced HYAFF11®/ PLLA scaffolds (Guarino et al., 2007) previously described combine the features of the biomaterials in second and third generations, simultaneously possessing bioactivity, degradability and eventually the capability to incorporate biomolecules (Causa et al., 2006). Ge et al. (2006) proposed a ‘dual’ composite structure, which, it is claimed, could better meet the basic requirements for a scaffold. The two ends of the proposed dual structure should allow osteogenesis and be integrated with host bone after implantation, while the middle part of it should host tissue ingrowth and, in time, attendant functionality. In cross-section, it could be composed of multiple layers, each individual layer fulfilling different functions, as well as degrading at different rates (Ge et al., 2006). The outer layer of such a structure should block inflammation cytokines and other macromolecules from knee joints, while allowing free exchange of nutrient ions. The middle layer should provide a good microenvironment for tissue ingrowth and developing functionality. In due course, growth factors stored in the middle layer would be released to promote faster tissue ingrowth (Ge et al., 2006). The core of the composite structure would comprise intact multiple layers, imparting the mechanical strength necessary for ACL reconstruction, possibly enhanced by a fibre-reinforced matrix. The overall mechanical behaviour of the structure would remain stable during degradation and always match that of the ACL (Ge et al., 2006). The growth factors stored within the core of the composite structure, principally to promote blood supply and tissue functionality, would ideally be released at a later stage and at a stable rate. However, tissue engineering techniques still present several limitations such as those related to the role of mature tissues, which are difficult to restore, and to the fact that flexible fibrocartilagineous tissues are subjected to complex dynamic loading processes for long periods in vivo (Shikinami et al., 2004).
10.7
References
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ambrosio l, de santis r, iannace s, netti pa, nicolais l (1998), ‘Viscoelastic behavior of composite ligament prostheses’, J Biomed Mater Res, 42, 6–12. ambrosio l, de santis r, nicolais l (1998), ‘Composite hydrogels for implants’, Proc Inst Mech Eng, 212, Part H, 93–99. amiel d, billings e jr, harwood fl (1990), ‘Collagenase activity in anterior cruciate ligament: protective role of the synovial sheath’, J App Physiol, 69, 902–906. amiel d, frank c, harwood f, fronek j, akeson w (1984), ‘Tendons and ligaments: a morphological and biochemical comparison’, Orthop J. Res, 1, 257–265. amis aa, kempson sa (1999), ‘Failure mechanisms of polyester fiber anterior cruciate ligament implants: a human retrieval and laboratory study’, J Biomed Mater Res, 48, 534–539. argawal bd, broutman lj (1980), Analysis and performance of fiber composites, John Wiley & Sons, USA. arnoczky sp, matyas jr, buckwalter ja, amiel d (1963). In Jackson DW. The anterior cruciate ligament, New York, Raven Press. bolton cw, bruchman wc (1985), ‘The GORE-TEX expanded polytetrafluoroethylene prosthetic ligament. An in vitro and in vivo evaluation’, Clin Orthop, 196, 202–213. causa f, netti pa, ambrosio l, ciapetti g, baldini n, pagani s, martini d, giunti a (2006), ‘Poly-ε-caprolactone/hydroxyapatite composites for bone regeneration: in vitro characterization and human osteoblast response’, J Biomed Mater Res, 76A, 151–162. davis pa, huang sj, ambrosio l, nicolais l, ronca d (1992), ‘A biodegradable composite artificial tendon’, J Mater Sci: Mater Med, 3, 359–364. de santis r, sarracino f, mollica f, netti pa, ambrosio l, nicolais l (2004), ‘Continuous fibre reinforced polymers as connective tissue replacement’, Compos Sci Technol, 64, 861–871. dunn mg, liesch jb, tiku ml, zawadsky jp (1995), ‘Development of fibroblastseeded ligament analogs for ACL reconstruction’, J Biomed Mater Res, 29, 1363–1371. durselen l, claes l, ignatius a, rubenacker s (1996), ‘Comparative animal study of three ligament prostheses for the replacement of the anterior cruciate and medial collateral ligament’, Biomaterials, 17, 977–982. duval n, chaput da (1997), ‘Classification of prosthetic ligament failures’, In L’Hocine Yahia, Ligaments and ligamentoplasties, Berlin, Springer-Verlag. friedman mj, sherman oh, fox jm, (1985), ‘Autogenic anterior cruciate ligaments (ACL) anterior reconstruction of the knee’, Clin Orthop, 196, 9–14. fu fh, bennett ch, lattermann c, ma cb (1999), ‘Current trends in anterior cruciate ligament reconstruction, Part 1: biology and biomechanics of reconstruction,’ Am J Sports Med, 27, 821–830. fujikawa k (1988), ‘Clinical study of anterior cruciate ligament reconstruction with the Leeds-Keio artificial ligament’, In Friedman MJ and Ferkel RD, Prosthetic ligament reconstruction of the knee, Philadelphia,W. B. Sanders Company. ge z, yang f, goh jc, ramakrishna s, lee eh (2006), ‘Biomaterials and scaffolds for ligament tissue engineering’, J Biomed Mater Res, 77A, 639–652. gershon b, cohn d, marom g (1990), ‘Utilization of composite laminate theory in the design of synthetic soft tissues for biomedical prostheses’, Biomaterials, 11, 548–552.
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gloria a, causa f, de santis r, netti pa, ambrosio l (2007), ‘Dynamic-mechanical properties of a novel composite intervertebral disc prosthesis’, J Mater Sci: Mater Med, 18, 2159–2165. guarino v, causa f, ambrosio l (2007), ‘Bioactive scaffolds for bone and ligament tissue’, Expert Rev Med Devices, 4(3), 405–418. guidoin mf, marois y, bejui j, poddevin n, king mw, guidoin r (2000), ‘Analysis of retrieved polymer fiber based replacements for the ACL’, Biomaterials, 21, 2461–2474. hull d (1981), An introduction to composite materials, Cambridge, Cambridge University Press. iannace s, sabatini g, ambrosio l, nicolais l (1995), ‘Mechanical behaviour of composite artificial tendons and ligaments’, Biomaterials, 16, 675–680. jackson dw, heinrich jt, simon tm (1994), ‘Biologic and synthetic implants to replace the anterior cruciate ligament’, Arthroscopy, 10, 442–452. jackson dw, simon tm, lowery w, gendler e (1996), ‘Biologic remodeling after anterior cruciate ligament reconstruction using a collagen matrix derived from demineralized bone. An experimental study in the goat model’, Am J Sports Med, 24, 405–414. jones rm (1975), Mechanics of composite materials, New York, MacGraw-Hill. ko fk, pastore cm, head aa (1989), Handbook of industrial braiding, Covington, KY: Atkins and Pearce, Inc. ko fk, pastore cm (1985), ‘Structure and properties of an integrated 3-D fabric for structural composites’, Am Soc Test Mater STP, 864, 428–439. ko fk, soebroto hb, lei c (1988), ‘3-D net shaped composites by the 2-step braiding process’ SAMPE J, 33, 912–932. ko fk (1987), ‘Braiding’, In Reinhart TJ, Engineering materials handbook: composites, vol. 1, ASM International, Metals Park, Ohio, 519–528. kolarik j, migliaresi c, stol m, nicolais l (1981), ‘Mechanical properties of model synthetic tendons’, J Biomed Mater Res, 15, 147–157. konikoff jj, billings w, nelson lj, hunter jm (1974), ‘Development of a single stage active tendon prosthesis. I. Distal end attachment’, J Bone Joint Surg Am, 56, 848. laurencin ct, ambrosio ama, borden md, cooper ja (1999), ‘Tissue engineering: orthopedic applications’, In Yarmush ML, Diller KR, Toner M, Annual review of biomedical engineering, Palo Alto, CA: Annual Reviews, 19–46. laurencin ct, freeman jw (2005), ‘Ligament tissue engineering: an evolutionary materials science approach’, Biomaterials, 26, 7530–7536. laurencin ct, ko fk, borden md, cooper ja, li wj, attawia m (1999), ‘Fiber based tissue engineered scaffolds for musculoskeletal applications, in vitro cellular response’, In Neenan T, Marcolongo M, Valentini RF, Biomedical materials: drug delivery, implants and tissue engineering, Symposium Held November 30– December 1, 1998. Boston, MA, USA: Materials Research Society. mallick pk (1988), Fiber-reinforced composites. New York, Marcel Dekker. mccarthy dm, tolin bs, schwendeman l, friedman mj, woo sl (1993), The anterior cruciate ligament: current and future concepts, New York, Raven Press, 343–356. migliaresi c, nicolais l (1980), ‘Composite materials for biomedical applications’, Int J Artif Organs, 3, 114–118. morgan jr, yarmush ml (1999), Tissue engineering methods and protocols, New Jersey, Humana Press Inc.
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netti pa, d’amore a, ronca d, ambrosio l, nicolais l (1996), ‘Structure–mechanical properties relationship of natural tendons and ligaments’, J Mater Sci: Mater Med, 7, 525–530. noyes fr, butler dl, grood es, zernicke rf, hefzy ms (1984), ‘Biomechanical analysis of human ligament grafts used in knee ligament repairs and reconstructions’, J Bone Joint Surg, 66, 344–352. o’connor jj, zavatsky a (1963), ‘The anterior cruciate ligament’, In Jackson DW, New York, Raven Press. olson ej, kang jd, fu fh, georgescu hi, mason gc, evans ch (1988), ‘The biochemical and histological effects of artificial ligament wear particles: in vitro and in vivo studies’, Am J Sports Med, 16, 558–570. parson jr, alexander h, weiss ab (1983), In Szycher M, Biocompatible polymers, metals and composites, Lancaster, PA: Technomic, 873–905. pennisi e (2002), ‘Tending tender tendons’, Science, 295, 1011. shikinami y, kotani y, cunningham bw, abumi k, kaneda k (2004), ‘A biomimetic artificial disc with improved mechanical properties compared to biological intervertebral discs’, Adv Funct Mater, 14, 1039–1046. silver fh (1994), Biomaterials, medical devices and tissue engineering: an integrated approach, Chapman & Hall, London, UK, 92. smith ba, livesay ga, woo sl (1993), ‘Biology and biomechanics of the anterior cruciate ligament’, Clin Sports Med, 12, 637–670. stol m, tolar m, adam m (1985), ‘Poly(2-hydroxyethyl methacrylate)–collagen composites which promote muscle cell differentiation in vitro’, Biomaterials, 6, 193–197. turner ig, thomas np (1990), ‘Comparative analysis of four types of synthetic anterior cruciate ligament replacement in the goat: in vivo histological and mechanical findings’, Biomaterials, 11, 312–329. von recum af (1986), Handbook of biomaterials evaluation: scientific, technical and clinical testing of implant materials, New York, Macmillan. vunjak-novakovic g, altman g, horan r, keplan dl (2004), ‘Tissue engineering of ligaments’, Ann Rev Biomed Eng, 6, 131–156. weitzel pp, richmond jc, altman ga, calabro t, kaplan dl (2002), ‘Future direction of the treatment of ACL ruptures’, Orthop Clin North Am, 33, 653–661. wise dl, trantolo dj, altobelli de, yaszemski mj, gresser jd, schwartz er (1995), Encyclopedic handbook of biomaterials and bioengineering–Part A: materials, New York, NY: Marcel Dekker. wise dl, trantolo dj, altobelli de, yaszemski mj, gresser jd, schwartz er (1995), Encyclopedic handbook of biomaterials and bioengineering–Part B: applications, New York, NY: Marcel Dekker. woo sly, an kn, arnoczky sp, wayne js, fithian dc, myers bs (1994), ‘Anatomy, biology, and biomechanics of tendon, ligament and ligament and meniscus’, In Simon SR, Orthopaedic basic science, USA, American Academy of Orthopaedic Surgeons, IL, 45–87. woo sly, buckwalter ja (1988), Injury and repair of the musculoskeletal soft tissues. American Academy of Orthopaedic Surgeons, IL, USA, 45–101 yahia l, hagemeister n, drouin g, rivard ch, rhalmi s, newman n (1997), Ligaments and ligamentoplasties, Berlin, Springer-Verlag.
11 Injectable composites for bone repair P. W E I S S and A. FAT I M I, Université de Nantes, France
Abstract: New percutaneous techniques using injectable biomaterials have been developed in bone surgery. The challenge is to associate the bioactivity, osteoconduction and fast substitution of the calcium phosphate biomaterials with rheological properties to use them in minimally invasive surgery. The best way to achieve this objective is to associate different compounds in an injectable composite with properties appropriate to the medical indication. To understand the various strategies involved in the research laboratory, we propose an easy classification based on the physical presentation of the medical devices proposed by the market. Key words: biomaterials, calcium phosphate, composite, injectable bone substitutes.
11.1
Introduction
Currently, the most commonly used injectable bone cement in orthopedics is poly(methyl methacrylate) (PMMA). Not only is this cement not degradable but its high curing temperatures can also cause necrosis of the surrounding tissue. For more than 25 years, calcium sulfate (Coetzee, 1980) and calcium phosphate (CaP) biomaterials have been used (De Groot, 1980) in various clinical applications such as for filling bone defects (Daculsi et al., 1990b), bone augmentation in spinal arthrodesis (Cavagna et al., 1999; Passuti et al., 1989), periodontal treatment (LeGeros, 1988), and as coatings on metal implants (Delecrin et al., 1994). CaP biomaterials are considered to be biocompatible and osteoconductive, achieving coalescence with bone tissue (Daculsi et al., 1989, 1990a). More recently, new percutaneous techniques using injectable biomaterials have been developed in spinal and orthopedic surgery (Taylor et al., 2007; Verlaan et al., 2006). The acrylic cementation of vertebrae and filling of bone cysts were the first documented applications (Malawer and Dunham, 1991). Although acrylic cements fulfill the requirements of injectability, filling complex-shaped bone defects and very firmly setting in situ, these materials nevertheless lack osteoconductivity and degradability. The challenge is to associate the bioactivity, osteoconduction and fast substitution of the calcium phosphate biomaterials with rheological properties to use them in minimal invasive surgery. The best way to achieve this objective is to associate different compounds. 255
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There are at least two materials for injectable composite substitutes. The first is the bioactive calcium phosphate in a particulate state and the second is a more or less viscous liquid, for the injectability of bioactive material. A third component could be used to increase chemical–physical properties or biological behavior. As a result of their chemistry, several substitutes can harden in the body just after injection whereas other remains viscous until bone invasion solidifies them. In all cases, the material is a multiphase composite material, but the interface changes with the biological substitution process.
11.1.1 Background to injectable bone substitutes Injectable CaP biomaterials should associate efficient bone colonization on implantation with non-invasive surgical techniques. Two types of injectable bone substitutes (IBS), are being developed in laboratories: calcium phosphate cement and calcium phosphate suspensions. The concept of apatitic calcium phosphate cement (CPC) was first introduced by LeGeros in 1982 (LeGeros et al., 1982). The first patent on hydraulic CPC cement (self-setting or self-hardening) was obtained by Brown and Chow in 1988 (Brown and Chow, 1987). In the 1990s, considerable efforts were made to develop injectable bone substitutes. Self-setting calcium phosphate cements (CPC) were the first IBS developed for percutaneous applications (Drissens et al., 1994). The calcium phosphate cements were first used as bone substitutes to repair cranio-facial defects. The histological observations on an animal model (Costantino et al., 1991; Friedman et al., 1991) showed a bone-implant interface with proliferation of osseous cells in the volume of the implant after several weeks, and then a slow reduction in the biomaterial with a new bone formation after several months. These hydraulic cements, however, are not ready-to-use, requiring extemporaneous mixing with various in situ setting times. Furthermore, most CPC wash out when they come into contact with body fluids before setting (Takechi et al., 1998). Additionally, once hardened, CPC produce a dense material with irregular microporosity and slowly degrade in vivo (Carey et al., 2005; Constantz et al., 1995), whereas numerous studies have shown that interconnected macropores are needed to facilitate bone in-growth (Daculsi and Passuti, 1990). A second type of injectable bone substitute, consisting of CaP ceramic granules suspended in a vicous water-soluble polymer carrier phase, has been developed.
11.1.2 Why do we need injectable bone substitutes? With classical surgery we can use calcium phosphate ceramic blocks, however, surgeons need to open the body and to remove bone to enable
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the placement of these blocks. The classical surgical procedure is invasive and may cause a high level of inflammation which can delay the healing process. To decrease patient immobilization times, which is more comfortable and less expensive, surgical technique development is directed towards minimally invasive procedures. IBS formulations in particular are attractive since they can be used for the purpose of minimally invasive surgical procedures and can be molded to exactly fill irregular bone defects.
11.1.3 Objectives of injectable bone substitutes The main objective of IBS is to enhance wound healing and bone in-growth when it is not possible to anticipate self-healing and bone regeneration. This process occcurs through its osteoconductive and bioactive properties. First of all, the material has to be biocompatible and free of any cytoxicity or immunogenicity just after implantation and over a long period of time because of its degradation products. Following bone loss, the principal indications are in critical size bone cavity. The material has to be injectable and sterilisable and has to disappear and be replaced by new functional bone during the bone remodeling process within a 10-year period and for the whole body.
11.2
Classifications of injectable bone substitute
As for all classifications, it is not easy to compile a classification of the entire injectable bioactive composite for bone repair. The first classification is by the chemistry of the compound used to produce these materials. As with the definition of composites, all these materials are associations of different materials and it could be difficult to characterize all of them. The only easy way to differentiate them is the chemistry of the mineral phase, such as calcium sulfate, a calcium phosphate such as brushite [CaHPO4·2H2O], apatite [Ca5(PO4)OH], hydroxyapatite (HA), beta tricalcium phosphate [βTCP, β-Ca3(PO4)2] and a biphasic calcium phosphate that is an association of HA and βTCP. These calcium phosphates can be the result of a cementing process which involves the end product of an acid–base reaction in water. There are two main types of hydraulic calcium phosphate cements (CPC), depending on the end product of the setting reaction: brushite and apatite CPC. Calcium phosphate granules or fillers can be synthetic-like sintered ceramics after grinding or fillers of biological origin like bovine or porcine bone. The second classification is by mechanical properties, which can decipher two types of IBS: the first is the hardening IBS and the second is nonhardening and is a calcium phosphate suspension. The third classification is by the presentation, ready-to-use or not.
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Class II: dense calcium phosphate matrix associated with particulate porogens
11.1 Schematic representation of IBS physical presentation: for Class 1, cells spread on the osteoconductive particular granules; for Class II, cells progress inside the material by the interconnected pores from the surface.
Whether the porosity is interconnected or not remains to be answered as does whether we use cells with IBS for bone tissue engineering. One way to present IBS is via its physical presentation to the host bone and cells (Fig. 11.1). This classification is very simple and can explain the kinetics of bone in-growth which is at the origin of the biological and mechanical properties, and is more relevant than the initial mechanical properties. The first physical presentation, Class I, is particulate calcium phosphate in a rapidly resorbable matrix which disappears quickly, allowing rapid osteoconduction from the wall of the bone cavity to the center. The fillers have slow resorption kinetics and act as the scaffold for the bone in-growth. The main advantage of this presentation is the fully interconnected properties (Weiss et al., 2003) of the composite after the matrix dissolution and/or degradation. Class I IBS comprise (i) non-hardening and (ii) hardening materials. For the hardening materials, the matrix can be organic or mineral. The disadvantage is the lack of initial mechanical properties. The second physical presentation, Class II, is a dense calcium phosphate matrix associated with particulate porogens. These porogens will disappear after dissolution and/or degradation that gives controlled porosity to the IBS to increase the kinetics of bone in-growth and substitution. The greatest advantage of this presentation is the initial mechanical properties but the disadvantage is a slow kinetics of bone in-growth depending on the macroporosity interconnection.
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11.2.1 Class I: particulate scaffold in a matrix Non-hardening IBS: suspensions The development of a composite bioactive material was carried out in accordance with the good results reported with both calcium phosphate ceramics and CPC involving formation of calcified tissues (Weiss et al., 1994). The most attractive feature of this injectable biomaterial is the mineral component with various HA/βTCP ratio making it possible to control its dissolution and precipitation kinetics and, subsequently, the bioactivity of the bone substitute (Daculsi et al., 1989; LeGeros, 1991, 2002; Nery et al., 1992). This material is an injectable ceramic consisting of biphasic calcium phosphate as the bioactive material, in a matrix of hydroxypropylmethylcellulose (Grimandi et al., 1998; Weiss et al., 1994) as the carrier. The degradation of the cellulosic polymer and the stability of this composite material were evaluated (Bohic et al., 2001; Weiss et al., 1997, 1998, 1999). The biocompatibility of cellulose and its derivatives has been documented (Barbie et al., 1990; Grimandi et al., 1998). This injectable calcium phosphate ceramics suspension (ICPCS) is ready-to-use and osteoconductive, but lacks initial mechanical strength. Bone in-growths develop very rapidly owing to the material’s interconnected macroporosity (Fig. 11.2 and 11.3A) (Gauthier et al., 1999b; Weiss et al., 1999). A similar concept was developed by Chazono et al. (2004) using βTCP granules as the bioactive fillers and sodium hyaluronate as the carrier matrix. The new bone formation inside the implanted bone area was similar to the previous suspension: 10% at two weeks and 20–30% between three and four weeks. These comparative results seem to favor the role of the physical presentation of the biomaterial compared with the chemical composition when the composition and the formulation are similar. Another way of making an injectable calcium phosphate suspension is to use natural inorganic bovine-derived hydroxyapatite matrix (ABM) combined with a synthetic cell-binding peptide on its surface as the bioactive fillers in a viscous phase, which is a water solution of glycerol and a ionic cellulosic ether, sodium carboxymethylcellulose (Barros et al., 2006). The water suspension of nano hydroxyapatite cannot be considered to be a composite but it is an injectable nano calcium phosphate suspension with an indication similar to the other composite. It has a viscous, fluid-like consistence and can be injected into a bone defect. A pig animal model has shown that complete resorption of the hydroxyapatite nanoparticles occurred after 12 weeks with 22% bone in-growth after three weeks (Thorwarth et al., 2005). The results were similar to autogenous bone after 8 weeks of implantation.
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11.2 Reconstructed microtomographic images of the IBS before implantation (a) and 3 weeks after implantation in femoral defects (b). New bone appears in gray, BCP ceramic in white, and soft tissues in black (bars: 100 μm). (a) General aspect of the composite IBS before implantation. The presence of the polymer confers on IBS its rheological properties and manages interconnected intergranular spaces. (b) Image of bone ingrowth in a femoral defect filled with IBS containing 40–80 μm BCP particles after 8 weeks of implantation. New bone formation formed an interconnected, newly formed bone network that developed extensively in the intergranular spaces.
Calcium phosphate ceramics with a hardening matrix Ceramics in resorbable organic matrix A self-hardening injectable calcium phosphate ceramics suspension (SHICPCS) is a composite made of BCP granules, the bioactive scaffold, in a pH-sensitive self-hardening hydrogel made of silated hydroxypropylmethylcellulose (Si-HPMC) (Weiss et al., 1995) as the matrix. The granules of BCP are associated with silated hydrogel (Si-HPMC). The guiding principle of Si-HPMC is its hydrophilic and liquid property (it is viscous before mixing with the calcium phosphate load and injection) and its pHcontrolled process of reticulation. The silated hydrogel/calcium phosphate composite involved self-reticulation obtained by the pH change as catalyst and with exothermic effect (Bourges et al., 2002; Fellah et al., 2006; Turczyn et al., 2000). Hoffmann et al. (2007) have devised a novel biomaterial consisting of the biodegradable compounds hydroxyapatite, chitosan and starch.
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100 μm
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11.3 Backscattered scanning electron micrographs of new bone formation in femoral defects filled with IBS after implantation. New bone appears in dark gray, BCP particules in white calcium phosphate cement in gray and soft tissues in black: (a) Class I a: non-hardening suspension of BCP granules in HPMC water solution showing bone in-growth inside the implant 3 weeks after implantation. Centripetal new bone formation occurred both in close contact with the surface of the BCP particles and in intergranular spaces. (b) calcium phosphate cement (CPC) image of the interface between cement and newly formed bone 3 weeks after implantation. New bone formation occurred in close contact with the surface of the cement without any fibrous interface but bone did not progress inside the CPC. (c) Class I b: hardening suspension of ceramics (BCP) in resorbable mineral matrix MCPC (Daculsi G, Khairoun I., 2007) 6 weeks after implantation showing the same bone in-growth processes as Class Ia with remaining BCP and piece of CPC. (d) Class II b: porous hardening matrices (I. Khairoun, Graftys HBS®, Graftys SAS, Aix en Provence, France) 8 weeks after implantation showing bone in-growth inside calcium phosphate matrix.
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Hydroxyapatite particles were modified in a layer-to-layer reaction with oxidized starch and deacetylated chitosan. In this manner, two kinds of particles with different surfaces were generated, one featuring deacetylated chitosan, the other one exhibiting oxidized starch at the outermost layer. Mixed together, these two kinds of particles form a homogeneous powder, which can be transformed into a paste by adding water. The paste particles bind to each other in an aqueous medium, and there is also the possibility of adjusting the paste viscosity according to the surgeon’s needs, rendering it a promising injectable filling material for bone defects in orthopedic surgery. An association of BCP bioceramics and fibrin sealants was also used to produce bioactive hardening materials (Le Nihouannen et al., 2007a, 2007b). As is the case for all IBS, this association is easy to handle, and allows filling bone cavities that are available for bone apposition. In addition, the improvement of bioceramic performance is achieved by the addition of bioactive factors included in fibrin sealant. Another strategy in this IBS category is the use of bioactive fillers like calcium phosphate or bioactive glass in a resorbable dry polymer matrix like polylactic polymers (Aho et al., 2004). The mechanical properties of these composite seem to be appropriate for bone tissue reconstruction (data not published) but the kinetics of bone substitution is very slow because the percentage of new bone in-growth into the composite was 6–8% at 23 weeks compared with the same percentage with ICPCS in the same animal model (Gauthier et al., 1999b).
Ceramics in resorbable mineral matrix The composite matrix can be a mineral self-setting CPC. To allow cells invasion and substitution, the matrix has to be quickly removed by rapid dissolution or bioresorbtion without any local toxicity owing to the level of liberated ions. Khairoun and co-workers have developed calcium phosphate ceramics such as BCP as the bioactive fillers in a highly resorbable CPC (Fig. 11.3C) (Daculsi et al., 2007; Julien et al., 2007). The end product and resorbable mineral matrix is calcium-deficient apatite. For bioactive fillers, the calcium phosphate can be βTCP granule powder. In addition, the end product resorbable mineral matrix is brushite (Apelt et al., 2004). The bioactive fillers can also be hydroxyapatite granule powder and resorbable mineral matrix calcium sulfate (Nilsson et al., 2003). The powder phase and liquid phase of the product are mixed in an apparatus combining mixture and injection.
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11.2.2 Class II: Porous hardening matrix Cements The thick consistency of CPC makes them unfit for injection (Khairoun et al., 1998) and various formulations have been proposed to improve their handling. For example, the addition of substances such as glycerin (Chohayeb et al., 1987), silicone gel (White and Goodis, 1991), polyethylene glycol, liquid paraffin, glycerol (Sugawara et al., 1992; Takagi et al., 2003), cellulosic (Cherng et al., 1997; Takagi et al., 2003; Yoshikawa et al., 1994) compound can increase the rheological properties and control the setting time. Other components include dispersants, binders, plasticizers, drugs that can be incorporated to modify their biological properties and their injectability. Some of the first class of IBS (ICPCS) are ready for use without any preparation and so decrease the risk of contamination during the surgical procedure. For purposes of hardening CPC, Carey and co-workers have developed a premixed calcium phosphate composite to present a ready-touse injectable and hardening CPC (Carey et al., 2005). The premixed CPC is CPC powder + non aqueous liquid + gelling agent + hardening accelerator. Although calcium phosphate cements are hardening materials with compressive mechanical strength, they are likely to shatter with catastrophic fractures. These poor cementing properties have severely restricted use to uniquely non-load-bearing applications. The use of the cement is limited to the reconstruction of non-stress-bearing bone, and clinical usage has been limited by brittleness. CPC provide dense biomaterials with irregular microporosity (Constantz et al., 1995) whereas macroporosity is known to be an essential factor for homogeneous and early bone colonization (Daculsi and Passuti, 1990; Gauthier et al., 1998; van Blitterswijk et al., 1986). Calcium phosphate cement is a massive material without macroporosity that allows good osteoconductivity on its surface, but acts as a barrier to colonization inside the implant (Fig. 11.3B) (Weiss et al., 2003). Moreover, the surface (circumference and tiny peripheral fissures) accessible for cell action and the resorption process is minimal. As a result, cement degradation is limited and substitution is slow. Macroporosity is known to be an essential factor for the homogeneous bone colonization of CaP biomaterials. To overcome these limitations, researchers devised new composites of calcium phosphate cement with fillers which can dissolve quickly in water and allow defined macroporosity. These porogens can be sugar or polysaccharides such as chitosan, manitol (Takagi and Chow, 2001; Xu et al., 2001), cellulosic ether (Fig. 11.3D) (Khairoun et al., 2006) or polylactic granules (Link et al., 2006). Because of the brittle properties of the cements, VicrylTM fibers have been used by Xu and co-workers to improve the mechanical properties and resorbability (Ethicon, Somerville, NJ, USA) (Xu et al., 2008).
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11.3
Stability, rheology and injectability of injectable bone substitutes
11.3.1 Stability One of the most important properties is the stability of the biomaterial and each of its phases before sterilization. The method of sterilization, however, must not interfere with the bioactivity of the material or alter its chemical composition which could, in turn, affect its biocompatibility or degradation properties. Such sterilization must not interfere with the physical properties of each phase and/or with the final formulations such as stability, rheology and injectability. IBS, especially injectable calcium phosphate ceramic suspensions (ICPCS), consist of two phases (i.e., calcium phosphate particles and polymer solution) and are considered as ‘ready-to-use’ biomaterials. This generation of injectable biomaterials poses a number of problems of stability during storage. Phase separation between particles and polymer solution leads to an inhomogeneous suspension. The word stability is ubiquitous in material science but means different things to different people. In a concentrate suspension it is possible to identify three different physical processes that may cause instability in a particulate system, namely a change in particle size owing to Ostwald ripening, particle aggregation and particle sedimentation. Although these processes may occur simultaneously they are here treated separately in order to promote an understanding of the physical forces involved. The engineering sciences have long been interested in models describing the settling (or sedimentation) of particle ensembles in viscous fluids. Zeta potential analysis demonstrates how the particle–surface interaction energy is very sensitive to changes in particle size. Specifically, the key attributes of these profiles, namely, the height of the repulsive energy barrier and the depth of the secondary energy well, both increase in magnitude with the increasing particle diameter. On the other hand, the presence of polymer leads to a slight decrease in the absolute zeta potential of the calcium phosphate (CaP) particles at fixed particle size. Adsorbed layers of polymer on the surface can enhance the stability of suspensions by an effect known as ‘steric stabilization’. Various approaches have been proposed to describe the hindered settling process. Most of them consider the hindrances as if dependent on the volume fraction of solids. The relative settling velocity can be calculated using Stokes’s law in suspensions with a minimal volume fraction of solids in which the particles can be considered as completely isolated (Stokes, 1845): uSt = d2(ρs − ρf)g/18μf where uSt is Stokes’s velocity, ρs and ρf are the densities of the solid particles and liquid, respectively, d is the diameter of
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the particle, g is the gravitational acceleration and μf is the viscosity of the suspending liquid. This law is considered as the fitting model to describe the settling process for ‘laminar’ flow. Although Stokes’s law is a simple expression for the relative settling particle velocity in viscous fluids; it is valid only for minimal volume fraction of solids in which the particles are not affected by each other and the container walls (Fatimi et al., 2008). In practice, settling velocity is affected by many factors. One important factor is the concentration of particles. In concentrated suspensions, the settling velocity is influenced by the interparticle forces which normally reduce the settling velocity with respect to Stokes’s velocity. This process is commonly referred to as hindered settling. Hindered settling is found to be a function of solid concentration.
11.3.2 Rheology The rheology of CaP suspensions is used to predict the injectability of this generation of biomaterial. In 1965, one of the first studies related to CaP suspension rheology investigated the viscosity of dicalcium phosphate suspensions (Bujake, 1965). Results demonstrated appreciable shear thinning behavior and suggested significant particle–particle interaction in these sus. pensions. Over most regions of the shear rate (γ ), the empirical power-law .n equation τ = Kγ was proposed to describe flow curve of CaP suspensions, where τ is the shear stress, K is the consistency factor and n is the flow index. Rao and Kannan examined the yield stress and viscosity of hydroxyapatite suspensions (Rao and Kannan, 2001). For all suspensions, the researchers observed a yield stress and a shear-thinning followed by shear-thickening behavior. Generally, shear thickening appears to occur at high particle loading (Knowles et al., 2000). Friberg et al. (2001) measured the viscosity of βTCP suspension by varying the liquid-to-powder ratio (LPR), employing powders of two medium particle sizes, and adding three different modifiers. More recently, Baroud et al. (2005) have studied the rheological properties of concentrated aqueous βTCP suspensions. This study has reported measurements of the yield stress and the viscosity as a function of LPR and milling time of the powder. The LPR clearly affected the rheological properties of CaP suspensions. Increasing LPR results in a more dilute solution with less particle–particle interaction, and hence lower viscosity and yield stress. The effect of milling time was significant, viscosity and yield stress increased as a function of the milling time (Bujake, 1965; Knowles et al., 2000). Liu et al. (2006) studied rheological properties of concentrated aqueous injectable CaP cement. Their investigations showed that CaP cement presented viscoplasticity and thixotropy. Results of this study confirmed the dependence of the technological parameters such as LPR, temperature and particles size on the rheological behavior of CPC (Liu et al., 2006).
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11.3.3 Injectability Injectability is often addressed as one of the most significant properties of biomaterials to be used in minimally invasive surgery. Several works have been reported but none of them provide any link between the injectability of the biomaterials and their rheological properties. The injectability of CaP biomaterials has been investigated by several groups, but most studies were conducted on setting CaP cements. Khairoun et al. (1998) measured the injectability of setting cement pastes by measuring the percentage of paste that could be extruded from a syringe fitted with a needle. The injectability of CaP cement through a syringe is evaluated by the necessary pressure for injection. This pressure can be separated into two components: (i) the pressure required to overcome the yield stress and to trigger the flow of the suspension, and (ii) the pressure required to maintain the flow of a suspension. This pressure is mainly determined by the apparent viscosity of the suspension. Several methods for altering the injectability of CaP cement are available. The injectability of CaP suspension in general can be altered by: • • •
variation of the LPR; modification of the particle size distribution; addition of various polymer solutions.
Bohner and Baroud (2005) have already proposed a theoretical approach to CaP cement extrusion and have studied the effect of geometry and formulation of the paste on injectability. However, their model was limited by considering the flow through a porous medium of a Newtonian fluid. Based on these remarks, more recently, we have proposed a model for injection which takes into account the rheological behavior of the injectable CaP biomaterial and the geometrical parameters of the injection system. In this study, a modeling of the injection of viscoelastic fluids is developed and applied to non-setting CaP biomaterials. The modeling takes into consideration the shear-thinning properties of the sample as well as the detailed injection conditions (syringe and needle radii, needle length and injection speed). This model was first tested for model fluids (a Newtonian and a shear thinning fluid). The case of CaP biomaterials was studied by comparing theoretical values of the extrusion pressure with experimental data obtained under about different conditions (extrusion and CaP concentration). The main advantage of this approach is that a rheological characterization of the considered fluids allows a prediction of their injectability in various conditions. The comparison between experimental and predicted values is based on the use of a power law behavior (Bujake, 1965) of the flowing materials under the injection conditions.
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11.4
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Biological behaviors of injectable bone substitutes
11.4.1 Cells behavior The size of bone cells is about 10 to 50 μm depending on their origin or spread. All calcium phosphate injectable composites are made of calcium phosphate particulates in a viscous matrix. The original size of the bioactive fillers at the time of injection or after biodegradation should influence cell behaviors such as proliferation, apoptosis and differentiation. Cellular reactions to implanted biomaterials are influenced by the shape and size of powder granules. In a study of calcium phosphate ceramic from 10 to 400 μm (Malard et al., 1999), bone-marrow cultures showed a significantly higher rate of multinucleated giant cells in the 10–20 μm range. In the 200–400-mm range, bone in-growth was significantly higher than the 80–100 μm range and controls. Inflammatory reactions were poor in marrow cell cultures and in histological retrievals. In another study performed by Pioletti et al. (2000), the functions of osteoblasts in the presence of βTCP, brushite and cement particles were quantified. Two particle sizes were prepared. The first size corresponded to the critical diameter range 1–10 μm and the second size had a diameter of more than 10 μm. CPC particles could adversely affect the osteoblast functions. A decrease in viability, proliferation and production of extracellular matrix was measured. A dose effect was also observed. A ratio of 50 CPC particles per osteoblast could be considered as the maximum number of particles supported by an osteoblast. The smaller particles produced stronger negative effects on osteoblast functions than the larger ones. Future CPC development should minimize the generation of particles of less than 10 μm. On the other hand, nano particles produced the best effect on the promotion of cell growth and inhibition of cell apoptosis when their size decreases from 80 to 20 nm (Shi et al.).
11.4.2 Animal models Although CPC showed good direct contact with newly formed bone, it was slightly degraded 3 weeks after implantation (Khairoun et al., 2002). The cells are unable to penetrate the CPC biomaterial and to produce a complete mineralized tissue substitution (Gauthier et al., 2003) whereas bone in-growth with IBS becomes greater at the expense of the biomaterial, as porosity and interconnection increase (Weiss et al., 2003). Moreover, the injectable ceramics favor an earlier and more extensive osteogenesis than large BCP ceramic granules and provide an osseous architecture with a progressive but early improvement in mechanical properties (Gauthier
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et al., 2001). Originally, this injectable BCP/polymer composite material or IBS was successfully used on dog models on a vertebral disk site (Passuti et al., 1995) and to fill periodontal pockets or bone alveolar sockets (Boix et al., 2004, Boix et al., 2006; Gauthier et al., 1999a). These in vivo studies showed that Class I IBS such as calcium phosphate suspension supported extensive bone colonization (Gauthier et al., 1999b; 2003; Weiss et al., 2003) and the newly formed bone was in perfect continuity with the trabecular host bone structure (Gauthier et al., 2005) with higher mechanical strength. A study using particles of different sizes demonstrated that smaller grains result in faster bone formation, with 50% of the original BCP resorbed after 2 weeks in the small (40–80 μm) grain composite (Gauthier et al., 1999b). However, the major drawback of this system is that it has no significant initial mechanical properties. If bone in-growth is very rapid, this problem may eventually be overcome, but the lack of sufficient initial mechanical characteristics can also lead to difficulties in maintenance of the composite within the defect during surgery. This possible drawback in such clinical situations justifies the development of hardening material with the main drawback of slow cell invasion and bone substitution (Weiss et al., 2003). New bioactive injectable composites seek to overcome this problem by including resorbable matrix or porogens (Link et al., 2006; Takagi and Chow, 2001; Xu et al., 2008). Another approach is the use of polymer constructs (Temenoff and Mikos, 2000) but they can be less biocompatible than calcium phosphate materials because of their chemical hardening process.
11.5
Injectable bone substitutes for bone tissue engineering
Tissue engineering is ‘an interdisciplinary field that applies the principles of engineering and life sciences to the development of biological substitutes that restore, maintain, or improve tissue function or a whole organ’ (Langer and Vacanti, 1993). The Class I particulate scaffold in a matrix is the best way to carry adherent living cells to a specific area for bone apposition or in-growth in a large defect. Trojani et al. (2006) used Si-HPMC/BCP biomaterial as an injectable composite which was mixed extemporaneously with undifferentiated bone marrow stromal cells (BMSCs) prepared from C57BL/6 mice, injected in subcutaneous and intramuscular sites and retrieved 4 and 8 weeks after implantation. Dissection of the implants revealed a hard consistency and the absence of a fibrous capsule reflecting a good integration into the host tissues. Histological analysis showed mineralized woven bone in the granule inter-space with numerous active osteoclasts attached to the particles as assessed by the presence of multinucleated cells positively stained for TRAP activity and for the a3 subunit of the V-ATPase. Small vessels were
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homogeneously distributed throughout the implants. Similar results were obtained in SC and IM sites and no bone formation was observed in the control groups when cell-free and particle-free transplants were injected. These results indicate that this injectable biphasic calcium phosphate– hydrogel composite mixed with undifferentiated BMSCs is a new promising osteoinductive bone substitute. Elabd et al. (2007) demonstrated that human multipotent adipose-derived stem cells, hMADS, injected subcutaneously into nude mice, in the presence of Si-HPMC/BCP biomaterial, are able to induce the formation of a highly vascularized mineralized woven bone displaying numerous osteoblasts, osteocytes and osteoclasts, whereas cell-free/HIBS leads to empty implants. Furthermore, immunostaining of 8-week-old implants with human specific anti-osteocalcin antibody showed that the implanted hMADS cells had differentiated into osteoblasts and had promoted bone formation in nude mice. It appears that a low proportion of the cells present in the 8-week-old implants are positive for human antigens. The remaining cells are mouse osteoblasts and/or osteocytes. Thus, it is tempting to postulate that hMADS cells once committed to the osteogenic lineage acquire the ability to cause the invasion of the implant by host cells via an unknown endocrine/ paracrine mechanism which then co-participates in tissue formation. These results are in agreement with previous observations showing that when quail marrow cells were subcutaneously implanted with a biomaterial into nude mice, osteogenesis was observed as a two step phenomenon, i.e., a first step in which donor cells are largely responsible for the observed osteogenesis, and a second one in which host cells predominate. The Class II porous hardening matrix can be used for injectable bone tissue engineering but cells have to be protected during the mechanical mixture of the different component of the CPC matrix. For this purpose, Xu et al. (2008) used alginate beads to encapsulate cells during mixing and cement setting. Kneser et al. (2005) proposed fibrin gel. Both confirmed viability of cells inside the biological composite constructs with lack of bone formation in the last proposition.
11.6
Conclusion
Although the transplantation of autogenous, vital bone remains the method of choice for the treatment of bony defects, alloplastic materials are often preferred because they are readily available and their use avoids a donor site defect. Synthetics materials are more abundant and safer in order to circumvent any biological contamination risk. Several strategies can be invoked to achieve injectable bioactive bone substitutes. The main difficulty is forging a compromise between fast bone in-growth inside the materials and the first mechanical properties that are antagonistic. Choosing the right
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strategy can therefore have consequences for the clinical indications of each type of bioactive composite.
11.7
Bibliography
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pioletti, d. p., takei, h., lin, t., van landuyt, p., ma, q. j., kwon, s. y. & sung, k. l. (2000) The effects of calcium phosphate cement particles on osteoblast functions. Biomaterials, 21, 1103–1114. rao, r. r. & kannan, t. s. (2001) Dispersion and slip casting of hydroxyapatite. J Am Ceram Soc, 84, 1710–1716. shi, z., huang, x., cai, y., tang, r. & yang, d. Size effect of hydroxyapatite nanoparticles on proliferation and apoptosis of osteoblast-like cells. Acta Biomater, 5(1), 338–345. stokes, g. g. (1845) On the theories of internal friction of the fluids in motion. Trans Cambridge Philos Soc, 8, 287–319. sugawara, a., nishiyama, m., kusama, k., moro, i., nishimura, s., kudo, i., chow, l. c. & takagi, s. (1992) Histopathological reactions of calcium phosphate cement. Dent Mater J, 11, 11–16. takagi, s. & chow, l. c. (2001) Formation of macropores in calcium phosphate cement implants. J Mater Sci Mater Med, 12, 135–139. takagi, s., chow, l. c., hirayama, s. & sugawara, a. (2003) Premixed calciumphosphate cement pastes. J Biomed Mater Res B Appl Biomater, 67, 689–696. takechi, m., miyamoto, y., ishikawa, k., nagayama, m., kon, m., asaoka, k. & suzuki, k. (1998) Effects of added antibiotics on the basic properties of anti-washout-type fast-setting calcium phosphate cement. J Biomed Mater Res, 39, 308–316. taylor, r. s., fritzell, p. & taylor, r. j. (2007) Balloon kyphoplasty in the management of vertebral compression fractures: an updated systematic review and metaanalysis. Eur Spine J. temenoff, j. s. & mikos, a. g. (2000) Injectable biodegradable materials for orthopedic tissue engineering. Biomaterials, 21, 2405–2412. thorwarth, m., schultze-mosgau, s., kessler, p., wiltfang, j. & schlegel, k. a. (2005) Bone regeneration in osseous defects using a resorbable nanoparticular hydroxyapatite. J Oral Maxillofac Surg, 63, 1626–1633. trojani, c., boukhechba, f., scimeca, j.-c., vandenbos, f., michiels, j.-f., daculsi, g., boileau, p., weiss, p., carle, g. f. & rochet, n. (2006) Ectopic bone formation using an injectable biphasic calcium phosphate/Si-HPMC hydrogel composite loaded with undifferentiated bone marrow stromal cells. Biomaterials, 27, 3256–3264. turczyn, r., weiss, p., lapkowski, m. & daculsi, g. (2000) In situ self hardening bioactive composite for bone and dental surgery. J Biomater Sci Polym Ed, 11, 217–223. van blitterswijk, c. a., grote, j. j., koerten, h. k. & kuijpers, w. (1986) The biological performance of calcium phosphate ceramics in an infected implantation site. III: biological performance of beta-whitlockite in the noninfected and infected rat middle ear. J Biomed Mater Res, 20, 1197–1217. verlaan, j. j., oner, f. c. & dhert, w. j. (2006) Anterior spinal column augmentation with injectable bone cements. Biomaterials, 27, 290–301. weiss, p., bohic, s., lapkowski, m. & daculsi, g. (1998) Application of FT-IR microspectroscopy to the study of an injectable composite for bone and dental surgery. J Biomed Mater Res, 41, 167–170. weiss, p., daculsi, g., delecrin, j., grimandi, g. & passuti, n. (1994) CNRS Patent: Biomaterial composition – preparation proceeding. WO 95/21634. France. weiss, p., gauthier, o., bouler, j. m., grimandi, g. & daculsi, g. (1999) Injectable bone substitute using a hydrophilic polymer. Bone, 25, 67S–70S.
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weiss, p., lapkowski, m., daculsi, g. & dupraz, a. (1995) Brevet: composition pour biomatériau, procédé de préparation n 95-09-582 du 7/8/95. WO 97/05911. weiss, p., lapkowski, m., legeros, r. z., bouler, j. m., jean, a. & daculsi, g. (1997) Fourier-transform infrared spectroscopy study of an organic-mineral composite for bone and dental substitute materials. J Mater Sci Mater Med, 8, 621–629. weiss, p., obadia, l., magne, d., bourges, x., rau, c., weitkamp, t., khairoun, i., bouler, j. m., chappard, d., gauthier, o. & daculsi, g. (2003) Synchrotron x-ray microtomography (on a micron scale) provides three-dimensional imaging representation of bone ingrowth in calcium phosphate biomaterials. Biomaterials, 24, 4591–4601. white, j. m. & goodis, h. (1991) In vitro evaluation of an hydroxyapatite root canal system filling material. J Endod, 17, 561–566. xu, h. h., quinn, j. b., takagi, s., chow, l. c. & eichmiller, f. c. (2001) Strong and macroporous calcium phosphate cement: Effects of porosity and fiber reinforcement on mechanical properties. J Biomed Mater Res, 57, 457–466. xu, h. h. k., weir, m. d. & simon, c. g. (2008) Injectable and strong nano-apatite scaffolds for cell/growth factor delivery and bone regeneration. Dent Mater, 24, 1212–1222. yoshikawa, m., toda, t., oonishi, h. & al., e. (1994) Osteocompatibility and biocompatibility of tetracalcium phosphate cement. Bioceramics, 7, 187–192.
12 Composite materials for hip joint prostheses R. D E S A N T I S, A. G L O R I A and L. A M B R O S I O, National Research Council, Italy
Abstract: Composite hip prostheses have gained an increasingly important role in the development of prosthetic devices. These materials can be engineered more accurately than monolithic structures. Biomechanical properties of tissues composing the hip joint and the concept of composite material and anisotropy are introduced. The state of the art is reviewed for prosthetic components comprising monolithic and composite materials for hip arthroplasty. The final section shows the composite approach to develop the hip stem, including technology, modelling and testing. Key words: fibre-reinforced composite materials, hip prosthesis.
12.1
Introduction
Hip fracture represents a major health problem facing elderly people. Hip prosthesis is the greatest challenge of biomaterial design to meet the needs of joint arthroplasty. The loads experienced by the lower limbs are large; in the femur they can be several factors higher than the body weight. Unusual combinations of mechanical, chemical and physical properties are required by hip prostheses. High strength and toughness, long life resistance, tailored stiffness, resistance to impact, abrasion and corrosion, adequate transparency to electromagnetic waves for diagnostic purposes, are the combination of properties relevant for tissue substitute design. Composite hip prosthesis, with use of continuous-fibre-reinforced polymers, has gained an increasingly important role in the development of prosthetic devices. Composite materials can be engineered more accurately than monolithic structures (single-phase materials, such as metals), thus allowing the development of more effective tissue substitutes. Basic biomechanical properties of tissues composing the hip joint and the concept of composite material and anisotropy are introduced in the first section, ‘Properties of the hip joint’. The following section, ‘Materials for hip arthroplasty’, reports the state of the art of prosthetic components of the hip prostheses. This is divided into three subsections, each of which contains two paragraphs: the former reports prosthetic components made of monolithic materials, the latter illustrates the efforts in designing com276
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posite materials for hip arthroplasty. The final section, ‘Composite hip joint’, shows the composite approach to develop the hip stem. This section is divided into three subsections dealing with technology, modelling and testing.
12.2
Properties of the hip joint
The hip joint is classified as a ball and socket joint. It is formed by the articulation of the femoral head and the acetabulum through a synovial joint, while ligaments and muscles stabilise and provide propulsion. The main constituents of femoral bone are the extracellular matrix (i.e. collagen, apatite and water) and osteocytes which control and adapt the structure performance (You et al., 2004). The soft and ductile collagen fibrils (about 100 nm in diameter) are reinforced by stiff and brittle platelet apatite crystals at a subnanostructural level (Rho et al., 1998), these fibres are bonded by less than 1% by weight of a non-fibrillar organic matrix (Currey, 2001; Fantner et al., 2005). The elastic properties of collagen and hydroxyapatite are about 1.5 and 110 GPa, respectively. These constituents are conveniently organised allowing a wide range of mechanical properties to be covered (De Santis et al., 2007). The hierarchical structure of bone shows a composite organisation at a multi-scale level. A composite is a material constituted by two (or more) phases: a continuous phase and a dispersed phase. Cortical or compact bone represents the outer shell of the femur. Osteons or the Haversian systems are the basic elements distinguished at a microscopic level. An osteon is generally an arrangement of concentric lamellae, each constituted of mineralised collagen fibres oriented according to predefinite patterns. At a higher scale level, osteons are organised and oriented according to preferential directions (i.e. the longitudinal direction in the shaft of the femur). It is not surprising that the concept of engineering reinforced composite (Nicolais, 1975) stands for cortical bone at a multi-scale level (Lenz and Nackenhorst, 2004; Cowin, 1999). This bone has been initially considered as a transversely isotropic material. Using piezoelectric transducers and sensors the anisotropy in the Young’s modulus has been observed (Bonfield and Grynpas, 1977). Table 12.1 shows the elastic constants measured through mechanical testing. Table 12.2 reports the yield and failure properties of human compact bone from the femur in the longitudinal and transverse direction. Trabecular bone is a highly anisotropic tissue; the elastic modulus depends on the region and direction of loading (Deligianni et al., 1991; Van Rietbergen et al., 2003). The elastic modulus of macroscopical trabecular specimens reaches values up to 1 GPa (Ciarelli et al., 2000). The elastic modulus anisotropy, averaged in the whole proximal femur, suggests values of 0.13, 0.06 and 0.05 GPa in the axial, sagittal and coronal direction, respectively (Augat
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Table 12.1 Elastic constants: Direction 1 is the longitudinal direction of the femur, Ei and Gij are the elastic moduli and nij is the Poisson’s ratio (Ashman et al., 1984; Dabestani and Bonfield 1988; Reilly and Burstein 1975) Year
1975
1984
1987
E1 [GPa] E2 [GPa] G12 [GPa] n12 n21
17.0 11.5
20.0 13.4 6.23 0.35 0.23
22.5 13.4 6.23 0.32 0.23
0.46 0.31
Table 12.2 Anisotropy in the yield and failure properties of human compact bone from the femur (Currey and Butler, 1975; Jepsen and Davy, 1997; Reilly and Burstain 1975) Mechanical test
Tension
Direction
Longitudinal Transverse Longitudinal Transverse Longitudinal
Yield strain (%) Yield stress (MPa) Failure strain (%) Failure stress (MPa)
0.73
Compression
Bending
1.41
1.41
107.9
Torsion
1.3 55.8
3.8
0.7
133
53
5.2 208
67
208
74.1
et al., 1998; Majumdar et al., 1998). Using specimens oriented in the main direction of the trabecular network from the femoral neck, this bone shows a modulus of 3.23 and 2.7 GPa in compression and tension, respectively, whereas the greater trochanter indicates an elastic modulus of about 0.6 GPa in both compression and tension loading mode (Morgan and Keaveny, 2001). The mechanical load which acts on the femoral bone depends on the forces transferred at the hip and knee joints. The line of action of these forces hinges upon the configuration of each joint (e.g. the orientation of the cup or pelvis toward the head or the femoral axis) and upon the forces and reactions provided by muscles and ligaments involved in the specific attitude. The hip joint bears and transmits very high loads to bone amplified by the offset between the loading axis and the femoral axis. A bending
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Table 12.3 Resultant force on the hip joint as a function of activity for a body weight of 70 kg Activity
Slow walking
Fast walking
Car exitting
Running
Slope ascent
Force (N)
2000
3200
3700
3700
4900
moment on the neck of the femur can be clearly recognised (Singleton and LeVeau, 1975). The bone in the neck region is prone to fracture in the elderly, mainly owing to the degenerative effects of osteoporosis. Apart from the supine and the two legs standing postures, the force on the hip joint is always a multiple of the body weight (Inman, 1947). It is a factor of about 3 to 5 during walking, increasing during attitudes such us exiting a car and running to more than a factor of 7 during a fast slope ascent or heel strike (McLeish and Charnley, 1970; Morrison, 1970; Hof, 2001). The magnitude of this force is important in assessing the strength requirements of the hip prostheses (Akay and Aslan, 1996; Paul, 1997; Yildiz et al., 1998). Table 12.3 provides the magnitude of this load considering a body weight of 70 kg. The performance of the hip joint depends on the optimised combination of articular cartilage, a load-bearing connective tissue covering the bones involved in the joint, and synovial fluid, a nutrient fluid secreted within the joint area, which lubricates the joint to reduce friction and prevents articular cartilage erosion (Fung, 1993; Netti and Ambrosio, 2002; Long et al., 1998; Mensitieri et al., 1996). The femur–pelvis interface is completed by a intracapsular ligamentum, which attaches directly from the head of the femur to the acetabulum, and the fibre-reinforced capsule, which covers the femoral head and neck. In the capsule, a circular band of fibres forms a collar around the femoral neck, while longitudinal fibres travel along the femoral neck and carry blood vessels (Villar and Santor, 2006).
12.3
Materials for hip arthroplasty
Total hip arthroplasty is the most common surgery approach to restore the biomechanics of a damaged hip. The technique introduced in the 1960s by John Charnley represents a milestone in orthopaedic hip surgery. Figure 12.1 shows the configuration of a healthy hip joint and the main components of a total hip replacement. The prosthesis consists of a stemmed femoral component in conjunction with an acetabular cup. The neck and the head of the femoral component may form a unique device with the stem, or they can be separate modules joined through tapered surfaces. Both the femoral and acetabular
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Head Stem
Epiphysis
Cement Hip prosthesis
Femur
Diaphysis
12.1 Representation of healthy hip joint (left) and the main components of a total hip replacement (right).
components can be cemented or non-cemented. Cemented prosthesis are especially indicated when the hosting tissue does not meet general requirements of conformity and healthy structure (Bourne, 1996). Partially cemented hip prostheses are another solution which combines the features of cemented and cementless fixation (Claes et al., 2000).
12.3.1 Composite bone cements Bone cements based on self-polymerising poly(methyl methacrylate) (PMMA) represent the main synthetic biomaterials for anchoring the prosthetic components to bones. These cements consist of a solid powder phase made of PMMA or related copolymer and a liquid monomer phase (MMA) (Kuhn, 2000). After mixing the solid and liquid phases, the polymerisation takes place via a free radical reaction and the kinetics are regulated by the concentrations of the initiator and the activator (Dunne and Orr, 2001). The cement layer has the main task of resisting and transferring the loads between the natural and synthetic coupled materials. Bone cements also provide a mechanical buffer between bone and hip prosthetic components, reducing stress concentrations and absorbing mechanical shocks (De Santis et al., 2003; McCormack and Prendergast, 1996; Ronca and Guida, 1993; Van Rietbergen et al., 1993). In order to improve properties of bone cements the composite strategy has been investigated. The main limitation of this composite design arises from a technological point of view. Unlike other prosthetic devices, cements are processed into the human body. Therefore, a combination of rheological properties need to be satisfied before cement setting. As a consequence, the
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only class of composite material which can be investigated as bone cements is the particle or short-fibre-reinforced polymer (Nicolais et al., 1979). Using this approach, the introduction of a bioactive phase (e.g. hydroxyapatite) in the PMMA matrix is suggested in order to enhance the quality of the bone-cement interface (Dalby et al., 2001) whereas the use of a reinforced PMMA composite is proposed in order to improve the mechanical bulk properties of the cement (Gilbert et al., 1995). Multiwall carbon nanotubes (Marrs et al., 2007) have also been added at weights of 2 and 5% in order to improve fatigue performance of acrylic based bone cements, while SiO2 nanocomposites (Carotenuto et al., 1996) and phase-change materials (De Santis et al., 2006) have been incorporated in PMMA matrix in order to optimise properties and to control the thermal profile during the exothermal setting of the material.
12.3.2 Composite acetabular cups Prosthetic cups can be screwed, press fitted or cemented to the acetabulum. Ultra-high-molecular-weight polyethylene (UHMWPE) represents the material most used for the acetabular cup. This polymeric choice has been mainly motivated by tribological properties required at the joint between the acetabular cup and the prosthetic femoral head. The main drawback of this material is aseptic loosening related to wear debris originating from the polyethylene cup (Zimmerman et al., 2002) and creep deformations as the glass transition temperature of this polymer is lower than body temperature. In an effort to reduce polyethylene wear, alternative bearing kinematics joints, such as metal on metal and ceramics, are being explored (Galante, 1998; Ebied and Journeaux, 2002). Carbon-fibre-reinforced poly(ether ether ketone) (PEEK) and Kevlar fibre-reinforced polyethylene represent the composite strategy adopted to increase strength, wear and fatigue properties of acetabular cups. It is recognised that the presence of a second phase, such as fibres, determines a loss of the mobility of the polymeric amorphous chains and, consequently, a decrease of the rate of the free volume relaxation of PEEK (D’Amore et al., 1993). A critical issue of Kevlar-reinforced polymers is the debonding owing to the mismatch in thermoelastic properties between fibre and matrix during cool down from processing temperatures (Deteresa and Nicolais, 1988). However, the wear resistance of carbon-fibrereinforced PEEK against metals and ceramic is higher than UHMWPE (Wang et al., 1999). On the other hand, using the compression moulding technology, polyethylene cups reinforced with a weight of 10 and 30% of Kevlar fibres showed and elastic modulus closer to bone (3.7 and 6.0 GPa, respectively) and a frictional behaviour close to high-density polyethylene (HDPE) (Chowdhury et al., 2004). Carbon fibre reinforcement of
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resins for acetabular cups has been used in the form of long and continuous fibres as well as short and chopped fibres. The resulting composites against ceramic heads showed wear rates lower than UHMWPE (Howling et al., 2004).
12.3.3 Composite hip stem The prosthetic stem is the hip component which is joined to the femur. The stem fits into the medullary canal passing through the epiphysis of the femur. Thus, the stem mainly substitutes the cancellous bone, occupies part of the medullary canal and replaces part of the femur head tissues. It is not surprising that no matter which synthetic material is used, this prosthesis always alters the biomechanical stress field of the natural femur. However, the use of stem fitting into a canal is still the only approach to restore other human joints (e.g. the knee) and also dental crowns. Traditionally, metals such as high-strength stainless steel, forged cobalt–chromium alloys and titanium aluminium vanadium alloy have been used as materials for stem manufacturing (Galante, 1998). The choice of metallic materials mainly relies on the fatigue strength of metals. The material composing the hip stem has to deal with bone properties. The strength of bone can be higher than 100 MPa (Table 12.2). Compared with bone which is characterised by an autorepairing capability, synthetic materials undergo progressive and localised structural damage as a consequence of cyclic loading. Therefore, the fatigue resistance of the synthetic stem has to be higher than the strength of bone. Among single-phase materials, only metals satisfy this requirement. However, these materials are much stiffer than bone. Consequently, the large difference in stiffness between bone and metal causes the stress shielding effect leading to bone atrophy and aseptic loosening in the long-term period. Calcar bone resorption is always a consequence of MHP designs, the loss of proximal bone is estimated to be around 40–60% over five years. The higher the stiffness of the prosthesis, the greater the bone remodelling, hence the greater the bone loss owing to the stress-shielding effect (Van Rietbergen et al., 1993). Moreover, the stress-shielded bone does not heal completely and is susceptible to refracture after the removal of the metallic implant (Codran 1969). Another drawback of metals is allergic reaction, which may promote an escalation of uncontrolled corrosion phenomena (Tonino, 1976). More specific, ion release may cause local adverse tissue reactions as well as allogenic responses (Lamovec J., 1988; Martin A., 1988). Reinforced polymers or composites have gained an increasingly important role in the development of new stem materials because they can be engineered more accurately than monolithic structures (single-phase materials), thus allowing the development of more effective tissue substi-
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tutes. However, the unusual combination of properties required for the material design and the high fatigue strength, strongly reduces the number of suitable composites. As materials science and technology stands today, the only possible way to satisfy the biomechanical needs and extend material property combinations relies on the development of a polymeric composite reinforced by continuous carbon fibres. The biomechanics of bone growth, absorption and fracture healing are related to material properties, structural properties and bonding characteristics of the implant. Metal stems are homogeneous isotropic materials, thus metal hip stem design focuses on the geometry (Van Rietbergen et al., 1993), whereas composite hip prostheses (CHP) are material–structure designs, providing many new option and possibilities in implant design (Evans and Gregson, 1998). Fibre-reinforced composite materials can offer strength similar to that of metals, and also more flexibility than metals. Mechanical properties of a CHP can be tailored to meet the bone mechanical behaviour (Bobyn et al., 1992). Carbon fibre reinforcing thermoset resins such as epoxies were the first choice for composite prostheses (Ambrosio et al., 1987). Toxicity of eventual unreacted monomers suggests the use of thermoplastic polymers as matrix for implant applications (Petillo et al., 1994). Polymers under investigation as matrices for composite implants include poly(sulfone) (PS), poly(ether ether ketone) (PEEK) and poly(ether imide) (PEI) (Akay and Aslan, 1995, 1996; Evans and Gregson, 1998; Shirandami, 1990; Yildiz et al., 1998a, 1998b). Material–structural designs may differ, although a common challenge is a stem more flexible than those of MHP in order to improve proximal stress transfer (Chang et al., 1990; Svesnesson et al., 1977; Wilke et al., 1994). These engineering polymers are characterised by excellent mechanical properties, thermal stability, marginal water absorption and relatively easy processing. In addition, their high solvent and thermal resistance allows the production of sterilizable medical devices. Moreover, the selected materials have demonstrated, at the same time, both positive and negative properties for the particular applications. PEEK has excellent mechanical stability but requires critical processing conditions owing to its temperature-sensitive semi-crystalline structure. Polysulfone has shown poorer mechanical properties following saturation in Ringer’s solution. PEI has proved to be an excellent substrate for cell spreading and growth, eliciting no cytotoxic response or haemolysis, coupled with both easy processability and resistance to sterilization (γ rays and autoclave) (Peluso et al., 1994). PEI has been considered as a matrix for composites reinforced with drop-off plies of glass and carbon fibres for designing a composite hip joint prostheses with adequate stability and mechanical properties (De Santis et al., 2000).
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12.4
Composite hip
As already discussed, the biomechanical properties combination required for composite hip stem limits composite design to continuous carbon fibre reinforcement. The continuous fibre reinforcement is particularly indicated as mechanical tailoring involves anisotropy in order to meet hosting tissue requirements (Table 12.1). The elastic modulus of a unidirectional fibrereinforced lamina is directly proportional to the amount of reinforcement (Nicolais, 1975). The elastic modulus of particles or chopped fibres composites varies according to the rule of mixtures and to the shape of the reinforcement (Nicolais and Nicodemo, 1974). Consequently, using a fixed amount of reinforcement, the elastic modulus of the continuous fibre composite is higher than the particle or short-fibre-reinforced composite.
12.4.1 Stem technologies Several technologies involving continuous fibre reinforcement have been investigated to manufacture composite hip prostheses. The compression moulding of mat laminae, the braiding of fibres and the filament winding approaches represent the main efforts to achieve this aim. These technologies are illustrated in Fig. 12.2. Compression moulding is a closed mould process. The equipment is characterised by tool surfaces consisting of two hardened, ground, polished and often chromium-plated steel surfaces composing the external profile of the hip stem. The matched metal dies are mounted in a hydraulic press. Mould heating is achieved using electrical heaters, steam lines, or hot oil flowing through internal cavities. The heating curing time depends on the thermoset or thermoplastic which is used (order of a few minutes). This technology provides a rapid and repeatable process for making thermoset and thermoplastic composites parts. Thermoplastic composites, based on continuous Compression moulding +q
Filament winding
Braiding
–q
V V w
w
12.2 Basic technologies used for the prototyping of continuous-fibrereinforced hip prostheses.
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fibre reinforcement in the form of mat, are first heated to the processing temperature in an infra-red or hot air oven (Ambrosio et al., 1996). The charge is then transferred onto the cooler tool and the compression between the male and female moulds forces the resin and fibres to fill the mould cavity. In other instances, the plastic resin (which may be pre-heated) is loaded into the mould through one or more canals. When the moulds are forced one on the other, the applied pressure and heat forces the resin to fill the cavity of the mould. Moulding pressures (order of hundreds of bar) are maintained as the matrix cools and solidifies. Once the cross-linking is completed the composite part is extracted from the mould and the flash is trimmed. Using compression moulding of thermoplastic fibre-reinforced composites, graphite epoxy and carbon-fibre-reinforced PEEK (Yildiz et al., 1998) and drop-off plies of carbon and glass fibre reinforced PEI (De Santis et al., 2000) were processed to manufacture prototypes of tailored hip stem prostheses. Filament winding is an open mould process for composite materials suitable for producing axial symmetric composite structures. The fibre reinforcement is helically wound on a mandrel. By combining the translation and the rotatory movements of the transverse carriage and the mandrel, respectively, it is possible to manufacture cylindrical composite devices with tailored mechanical properties. The winding angle and the content of fibres allow fine tailoring of the elastic properties of the composite. The reinforcement may be fed through a resin bath and the winding process is strongly dependent on the resin characteristics since the matrix viscosity controls the friction between the tensioned fibres and resin. When the designed thickness of layers is reached the polymerisation is carried out (e.g. by heating or IR radiation) and the mandrel may be removed. The high degree of fibre orientation and the external position of the reinforcement produce a high stiffness composite in the form of hollow cylinders (i.e. shafts). Composite-assisted manufacturing (CAM) fulfil automatic procedures and more complex geometry may be realised using more than one degree of freedom of the mandrel and carriage. This technique is suitable for replicating dense (Ambrosio et al., 1998) and hard (De Santis et al., 2004) connective tissues analogues. On the other hand, the braiding process is useful to make near-netshaped devices. The system consists of a mandrel stock unit, a mandrel supply unit and a braider unit. The mandrel has the shape of the core of the composite hip stem and is supplied to the braider. The mandrel goes through the multiple braider units and yarns are intertwined from the braided layer of the composite. Several layers of fabric can be braided over each other to produce the required thickness and desired mechanical properties. Several patents of composite hip stem deal with the filament winding or the braiding approaches. A block of composite reinforced with helically wound carbon fibres is machined to form the composite stem (Ainsworth et al., 1990). Carbon fibres are filament wound around a box mandrel, four composite hips being cut from the corners of the reinforced box (Davidson
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et al., 1995). Braided fibres on a metal stem core has also been patented (Hamilton et al., 1994), manufactured (Simões and Marques, 2005) and investigated (Simões and Marques, 2005). The braiding approach has also been applied onto a polymeric mandrel core using carbon fibres, and the resulting composite hip stem provides a fine tailoring of the modulus (Trentacosta et al., 1995). Composite stem hip prostheses have also been realised using the combination of braiding and injection moulding technology. The wound carbon fibres obtained through the braiding approach are deposited into a mould having the negative shape of the stem prosthesis, and polyamide 12 is injected (Campbell et al., 2008). Figure 12.3 shows two continuous fibre-reinforced hip prostheses realised using the filament winding and the compression moulding approaches.
12.4.2 Modelling Composite materials are investigated according to their stacking sequence. Symmetric laminates are commonly realised to avoid the bending–inplane coupling, which in non-symmetric ones causes an undesirable warping owing to inplane loads and temperature change. A laminate in which all of the plies have a counter-ply with an opposite sign is defined as balanced; it is said to be unbalanced if the number of positive and negative orientations is not the same. In other words, the laminate is balanced when there are an equal number of plies with positive and negative orientations as depicted in Figure 12.2 (Nicolais, 1975; Yildiz et al., 1998a).
12.3 Two continuous-fibre-reinforced hip prostheses realised using the filament winding (left) and the compression moulding (right) approaches.
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By tailoring the stiffness of the prosthesis both along its length and through its thickness, it is possible to change the pattern of load transfer between the prosthesis and the bone. Accordingly, an appropriate ply orientation could be selected for the design of a composite implant which could produce more favourable stresses and deformations for the femoral bone. A finite element model (FEM) combined with a mathematical description of adaptive bone remodelling show the high performances of CHP in terms of mechanical stability and tissue conservation (Apicella et al., 1994; Kuiper and Huiskes, 1997; Mihalko et al., 1992; Van Rietbergen et al., 1993). In many studies relating to stresses and strain, bone has been assumed to be an orthotropic material described through nine independent constants and linear elasticity constitutive equations have been used in order to treat it adequately (Apicella et al., 1994). Mechanical properties of titanium alloy, composite fibre-reinforced laminae, natural cortical and trabecular bone are summarized in Table 12.4. FEM has clearly highlighted the advantage of using a composite material for the stem of the hip prostheses (Srinivasan et al., 2000). In particular, a decreasing rigidity from the head to the tip region has provided a stress transfer from the prostheses to bone more uniform than metal stems (Apicella et al., 1994). Compared with metals, the mechanical properties of composites vary from ply to ply, position to position, depending upon the type of reinforcement, fibre orientation, fibre volume fraction and laminae sequence (Yildiz et al., 1998a). Yildiz (1998) realized a 3D finite element analysis for analysing composite hip prostheses in the femur by proposing a ply drop-off composite element which could contain multidirectional fibres and ply dropoffs owing to the variation of the cross-section of the stem from the proximal region down to the distal one. A CHP made up of approximately 200
Table 12.4 Elastic properties of femoral bone, titanium alloy and fibrereinforced composites used in the finite element model, expressed in terms of elastic moduli (Ei, Gij) and Poisson’s ratio nij Elastic properties
Cortical bone
Trabecular bone
Titanium alloy
Carbon fibre reinforced PEI
Glass fibre reinforced PEI
E1 (GPa) E2 (GPa) E3 (GPa) G12 (GPa) G13 (GPa) G23 (GPa) n12 n13 n23
11.5 11.5 17 3.6 3.28 3.28 0.58 0.31 0.31
2 2 5 2 2 2 0.45 0.45 0.45
123 123 123 47 47 47 0.32 0.32 0.32
139 16 16 4.8 4.8 3.3 0.24 0.26 0.26
43 14 14 4.3 4.3 3.2 0.24 0.20 0.20
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unidirectional T300-976 graphite–epoxy composite plies was manufactured and then tested in bending using a displacement control mode (Yildiz et al., 1998a).
12.4.3 In vitro testing Experimental testing of composite hip prosthesis is necessary to assess the strength and to validate or calibrate models and technological processes. The proper knowledge of the hip biomechanics is essential to improve orthopaedic treatments designed to solve the specific pathology and disorder. The numerical and experimental approaches, used together and in parallel, represent a very powerful tool for optimising the mechanical function of a synthetic device (Huiskes et al., 1981). The long term stability of implanted prostheses is strongly related to the stress transfer between implants and bone. The fatigue strength of hip prosthesis is measured according to ISO 7206-1 and ISO 7206-3 or ASTM F1612-95 standards. Figure 12.4 shows the set up of a fatigue test on a hip prosthesis. These standards are also suitable for testing metallic hip implants. The behaviour of a composite hip prosthesis has to be assessed in conjunction with the surrounding femoral bone. Also the fatigue behaviour of the cement mantle is important where cemented hip prostheses are considered (Colombi, 2002). Synthetic models of long bones such as the femur are com-
12.4 Set up of a fatigue test on a hip prosthesis.
Composite materials for hip joint prostheses
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12.5 Fatigue behaviour of the glass and carbon fibre reinforced PEI hip prosthesis. Vertical displacement of the femoral head within a time window of 3 s after 1, 5, 7 and 10 million cycles.
mercially available. They are appropriate for certain types of biomechanical tests (Cristofolini et al., 1996). They have also been used to predict proximal bone stress-shielding in total hip reconstruction (McNamara et al., 1997). These models are made of glass fibre reinforced epoxy and polyurethan foam (which mimic the cortical and the spongy bone respectively). Synthetic models reproducing the anisotropy of bone tissue can also be obtained through the filament winding approach (De Santis et al., 2004). Experimental and numerical testing of composite hip prostheses has been carried out on glass and carbon fibre reinforced PEI (De Santis et al., 1998), on the AS4-PEEK graphite/thermoplastic unidirectional composite implant (Yildiz et al., 1998). A carbon fibre reinforced PEEK prosthesis has also been investigated in conjunction with a synthetic glass reinforced femur (Akai and Aslan 1996). This approach is also preferred for metal hip prostheses (Britton and Prandergast, 2005). Hip simulators provide an experimental complex loading of the hip joint. This approach is particularly indicated to investigate the effects of kinematics and motion between the femoral head and the acetabular cup (Calonius and Saikko, 2002). Figure 12.5 shows the fatigue behaviour of the glass and carbon fibre reinforced PEI hip prosthesis loaded between 50 and 2000 N.
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The vertical displacement of the femoral head was measured after 1, 5, 7 and 10 million cycles; this composite prosthesis was stable up to 107 cycles. The stiffness of the composite prosthesis was 2050 N mm−1.
12.5
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13 Harnessing the properties of fiber-reinforced composites in the design of tissueengineered scaffolds A. T. D i B E N E D E T T O and L. P I N AT T I, University of Connecticut, USA
Abstract: Incorporating the effects of morphological structure on the cellular activity required for tissue regeneration optimizes the design of tissue-engineered scaffolds. A growing data base in cell biochemistry can be used to predict these effects using “in-silico” computational systems and advanced technologies of rapid prototyping (RP). The range of tools available to the designer of biomedical devices is described and examples are presented of how the principles of mechanical and computational analysis are used to design scaffolds for repair of loadbearing tissues in bone, tendons and articular cartilage. Key words: tissue engineering, in-silico computational analysis, rapid prototyping, fiber-reinforced scaffolds.
13.1
Introduction
The design of tissue-engineered scaffolds for the repair of load-bearing tissues has been influenced by more than fifty years of significant advances in the use of fiber-reinforced composites in the design of aircraft, bridges and other load-bearing commercial materials. The principles of classical mechanics and the use of computer-aided design and manufacturing techniques (CAD/CAM) have been directed toward harnessing the directional properties of anisotropic materials to optimize strength, stiffness and inservice fatigue life. These same principles have been applied in the design of functional tissue-engineered scaffolding for load-bearing tissues. Until recently, much of the work has, by necessity, been limited by the inability to incorporate directly into the design process specifications for optimization of tissue regeneration by appropriate cellular components. This limits the efficacy of the design process since cellular activity within an engineered medical device is highly dependent upon the morphology of the device and the external and internal mechanical and electrical forces experienced during use (Leong 2008; Ma 2006; Sun 2003, 2001). In typical load-bearing tissues, these interactions occur at multiple hierarchical levels of morphological structure and are not easily quantified. Furthermore, the traditional CAD/CAM processes employed in industry for fabrication of structures 296
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have not always been readily adaptable to the reproduction of anisotropic, non-homogeneous natural load-bearing tissues. There are rapidly growing, readily accessible databases in cell biochemistry that can now be used to predict the effects of biomechanical and bioelectrochemical forces on cellular activity. With the development of ‘insilico’ computational systems and advanced technologies of rapid prototyping (RP) that more closely mimic living tissues, the goal of optimizing both the biomechanical and biological behavior of a specific medical device is now well-within reach. It has been pointed out that: The landmark completion of the first draft of the map of the human genome . . . called attention to the wealth of empirical data at present being assembled to define the intricate schedules of gene expression and proteome activity in cells and developing organisms. . . . Increasingly these data pools are assembled into computer-based data ‘warehouses’, which in turn are readily accessible by internet software . . . The value of these remarkable sites continues to be enhanced by the spread of computer software designed for the analysis, or ‘mining,’ of the data and by software that allows the data to be imported into ‘in-silico’ models of cell biochemistry. [Semple 2005, p. 347]
The ability to include in the design process an expectation of how biological functions will be affected by the details of structural design is key to major improvements in design of fiber-reinforced composite scaffolds. In-silico computational models will provide materials and process designers with a more complete range of tools for optimizing the morphology and physical properties of a biomedical device. Efforts in this direction are being stimulated by world-wide, long-range strategic plans in both the USA and EU. After an introduction to the range of tools available to the designer of biomedical devices, examples of how the principles of mechanics and computational analysis are being used in the design of scaffolds for the repair of load-bearing tissues of bone, tendons and articular cartilage will be presented.
13.2
Harnessing directional properties of biomaterials
Ideally, the goal of the designer of an implanted device for the repair of a load-bearing tissue is to construct a ‘biosystem’ that precisely matches the morphology, mechanical properties and biological functions of the original load-bearing tissue. Nature has provided us with a dynamic skeletal system that is constantly changing in response to the demands placed upon it. Over the course of millennia, all living creatures have had the ability to continually redesign their hard and soft tissues in order to best sustain the survival of the species. By necessity, however, engineering design strategies are fundamentally different than those used by nature:
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While the engineer selects a material to fabricate a part to an exact design, nature goes the opposite way and grows both the material and the whole organism using the principles of self assembly. This provides control of the material at all levels of hierarchy and is certainly a key to the successful use of polymers and composites as structural materials. [Fratzel 2007]
In nature, the functional adaptation of a living species to the stresses and strains imposed upon it depends on the ability of its cellular make-up to induce repair and remodeling of its tissues. The modeling process may vary at different hierarchical levels of organization (Curry 2005; Fratzl 2007; Jeronimidis 2000, 1995; Lakes 1993; Weiner 1998), since the spatial organization of the tissues may be different at the nano (molecular), micro and macro levels of morphology (Barthelat 2007; Gao 2006; Gupta 2006; Sanchez 2005). Modern attempts to analyze the ability of tissues to adapt to their environment date as far back as the late 19th century (Wolff 1870). Mathematical theories of bone adaptation were proposed (Firoozbakhsh 1980) and tested experimentally by a number of researchers over the years (Hernandez 1997; Mattheck 1990; Taylor 2003). Adaptation theories relate the rates of bone resorption and remodeling to imposed stress and/or strain relative to a ‘normal’ range under which no changes in bone thickness and porosity (density) occur. The approach is similar to those used to describe crack growth and fatigue damage in structural materials and has been applied to the development of a computer-aided design (CAD) technique that includes modification of component shape to reduce points of stress concentration that are the primary source of crack growth and fatigue failure (Mattheck 1990). Experimental and theoretical aspects of tissue adaptation and optimum structural design have been discussed by Hernandez (1997) and Taylor (2003). Functional tissue engineering has the goal of mimicking the spatial character of the tissue to provide a favorable environment for cell activity. Although available fabrication techniques are not able to reproduce exactly the complex, non-homogeneous structures of most load-bearing tissues, man-made fiber-reinforced composites have been fabricated with controlled fiber orientations that share a degree of similarity with those of natural tissues (Bar-Cohen 2005). Synthetic biomaterials by themselves, however, cannot impart a biomimetic physiological effect. Biological effectiveness of a scaffold depends on the mobility of accessible cells within the porous volume of the scaffold and the interactions of the cells with the surfaces of its fibrous super structure. The multi-directional diffusion and flow of cells within the porous matrix of the composite depend upon the size and shape of the cells and the geometric details of the porous regions within the scaffold. With the aid of medical modeling techniques, finite element analysis (FEA) and CAD/CAM techniques, a design process can
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be implemented to build prototype scaffold models. If appropriate microbiological data are available, in-silico computational modeling can be employed to aid in the identification of appropriate prototype morphologies that imitate, as closely as possible, the physiological response of the tissue. Evaluation of the efficacy of the prototype scaffolds can then be made through in vitro and in vivo laboratory testing. The hierarchical structure of bone, tendons and articular cartilage are described in section 13.3, followed by a description in section 13.4 of the processes employed in manufacture of functional tissue-engineered scaffolds.
13.3
Morphology of load-bearing tissues
13.3.1 Bone All bone tissues have a common building block of mineralized collagen fiber (Weiner 1998). When new bone is regenerated, collagen fibers are produced, then oriented and, subsequently, carbonated apatite is crystallized in the interstices between the oriented fibers. Thus, nature enlists collagen fibers to harness the directional properties of new tissue in order to reproduce the properties of the original hard tissue. While in some respects bone tissue is similar to a manufactured fiber-reinforced composite, it has a more complex hierarchical structure and, in addition, the ability to remodel itself in response to the demands placed upon it. There are two distinct types of bone tissue, namely cortical bone with a porosity of 5–25%, and cancellous (trabecular) bone with a porosity of 40–90% (Isenberg 2006). A typical bone has a tubular structure with a cortex, i.e., tube walls, of cortical tissue consisting of longitudinally oriented osteons that are typically of the order of 200 μm in diameter and 10–20 mm in length (Isenberg 2006). The structure of the osteons is that of concentric lamellae 3–7 μm in thickness. At the center are blood vessels that supply the tissue with nutrients. Cortical bone tissue can adopt various levels of hierarchical organization and scale depending on species, age or the conditions under which it was formed (Martin 1998). At the ends of a long bone, the cortex becomes thinner and consists of a higher porosity cancellous tissue in the form of trabeculae with a changing orientation of the mineralized collagen fibers. The high tensile strength and stiffness of cortical bone is attributable to the longitudinal orientation of the osteons, whereas the need for a transfer of compressive load between the bone and its connection to other bones through a joint is satisfied by the changing orientation of the collagen fibers in the trabeculae. The space inside the cortical bone tube is filled with a hematopoietic marrow that is supported and surrounded by bone tissue and periosteum (Buckwalter 1995). The marrow serves as a source of bone cells, and the blood vessels in the marrow form a critical part of the circulatory system within the bone. In situations where the production
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of new bone is required rapidly, as is the case after fracture of a healthy bone, a morphologically different structure, known as woven bone, is produced. It is a porous material with the collagen fibers having no preferred orientation (Taylor 2003). In this case, the gap between the two surfaces of the fractured bone is filled with fibrous bone tissue that is subsequently remodeled.
13.3.2 Tendons and ligaments The primary function of a tendon is to transmit a tensile force exerted by a muscle to the bones of the skeletal system. In some instances, it is also subjected to compression and shear in limbs where adjacent bony surfaces are used as pulleys. Ligaments connect bone to bone and have a hierarchical structure and function similar to tendons, but with a slightly different composition of collagen fibrils and proteoglycan matrix. The tendon structure is a tough cord of dense white fibrous connective tissue that consists of a unidirectional array of collagen fibers surrounded by a matrix of proteoglycans (Kasteleic 1978). The fiber reinforcement is primarily type 1 collagen, but other types are also present (Benjamin 2008). The proteoglycan matrix contains fibroblast cells organized in longitudinal rows in close proximity to the oriented collagen fibrils. These cells secrete the extracellular matrix and orient the collagen assembly during remodeling. Oriented collagen molecules aggregate to form sub-fibrils that are organized into discrete units of collagen fibril bundles. Groups of primary fiber bundles are then organized at the next hierarchical level as secondary fiber bundles (fascicles), which, in turn, form a third hierarchy of tertiary fiber bundles of the tendon. Each of these structural units is surrounded by thin layers of connective tissue through which pass blood vessels, lymphatic vessels and nerves. A crimping of the oriented collagen fibrils optimizes the loadbearing capacity and fatigue resistance of the tendon by allowing an initial rubbery-elastic stretching of the tendon as the fibrils ‘un-crimp’. Variations of the fiber diameters, fiber bundle sizes and a helical organization of tendon components result in regional variations of strength and compliance that give the tendon its unique character. The ability of the different hierarchical units to slide independently against each other also allows tendons to change shape as a muscle contracts (Fallon 2002). Both the transmission of load and the sliding of fibrils and fascicles occur within an encapsulating proteoglycan-rich matrix, which also imparts a viscoelastic response to the mechanical behavior of the tendon.
13.3.3 Articular cartilage Articular cartilage is non-homogeneous and highly anisotropic with, typically, an overall composition of approximately 70% water, 18% collagen, 8% proteoglycans and 4% chondrocyte cells (Taylor 2003). The components
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form a multi-layered extracellular matrix (ECM) of a network of collagen fibers entangled in a proteoglycan hydrogel. The aqueous phase contains inorganic salts, primarily NaCl, whose ions interact with ionic charges along the proteoglycan chains. The composition and morphology of articular cartilage varies continuously with depth from the articular surface to contact with a subchondral bone surface (Sun 2003). The structural variation can be classified into four layers as the superficial tangential zone, middle zone, deep zone and the calcified cartilage zone. The calcified cartilage zone provides for a gradual transition of mechanical properties from cartilage to bone, thus minimizing stress concentrations at the cartilage-bone interface (Taylor 2003). The gradation of mechanical properties continues from a highly oriented array of collagen fibers perpendicular to the joint surface in the deep zone to the parallel array of collagen fibers in the thin superficial tangential zone. The complex morphology provides a low friction character to the tissue and a hard surface skin that is resistant to the predominantly compressive forces and induced lateral tensile strains developed in the tissue under stress (Taylor 2003, p. 51; Moutos 2007, p. 162). The articular cartilage is also an electrochemically charged tissue by virtue of the fixed net negative charges of carboxyl (COO−) and sulfate (SO−3) ions on the proteoglycan and other macromolecules in the ECM. The mechanical inhomogeneity and anisotropy of structure as well as the fixed charge density of the tissue have a profound effect on the mechanico-electrochemical environment that controls cellular response and metabolic rates (Sun 2003, 2001, 1999).
13.4
The designer’s tools
Design of a scaffold that mimics the mechanical properties and biological functions of a load-bearing tissue requires a thorough analysis of the tissue’s morphology. The first step in a modeling process is to obtain a scanned three-dimensional (3D) image of the living tissue in question using techniques such as magnetic resonance imaging (MRI) and computed tomography (CT) (Bibb 2006). These data are exported from the scanner to a FEA package for creation of a CAD model of the tissue. With access to an appropriate database, a simulation of cellular migration within the tissue can be created that serves as the baseline for evaluating proposed scaffold designs. Biomaterial components are chosen and all available mechanical and biomechanical data are used to create an appropriate CAD/CAM program for fabrication and evaluation of prototype scaffolds.
13.4.1 Medical modeling MRI and CT are used to create 2D and 3D visual representations of tissue structure. Both MRI and CT scanners generate multiple 2D slices, each of
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which displays an image of structural density. The 2D images are stacked to produce a 3D visualization of the tissue. MRI uses radiofrequency signals to map its images and is favored for soft tissue analysis because it has a higher capacity for tissue differentiation than that of CT. On the other hand, CT is a better tool for examining bone tissue and interfaces between bone and soft tissue because it has a higher resolution than that of MRI. As it uses x rays to obtain its images, however, it differentiates tissue solely on the basis of a difference in the gray-scaled value of tissue density. A combination of CT with positron emission tomography (CT/PET) provides a more refined image by enabling co-registration of data into a single, more precise image. An up-to-date overview of medical modeling and its application to tissue and organ design can be found in Bibb (2006). Data derived from medical scanners can then be exported to an FEA program capable of mapping the structure of anisotropic materials.
13.4.2 CAD/CAM processes FEA is used for building CAD/CAM systems capable of guiding free form fabrication of anisotropic scaffolds. To create a structural representation using a FEA, a mesh is generated that contains the physical properties of the components and the variables that affect the structure’s thermomechanical responses. The mesh represents the surface of a two-dimensional object by breaking it down into a large number of connected geometric elements, such as triangles, i.e., an STL file, each of which is defined by its normal and three points representing its vertices. Other meshes can be used to break down the whole object into discrete 3D elements. These data are then used to construct a CAD model for the generation and analysis of the resulting images and, ultimately, for creating a CAM process to guide the fabrication of the structure. The principles of FEA and strategies for mesh design are described in Nicholson (2003). A practical guide for design engineers on the integration of FEA with CAD/CAM techniques is presented in Kurowski (2004). CAD/CAM techniques have been the primary tools for mechanical and civil engineers in the design of load-bearing structures such as bridges, automobiles and aircraft. Of particular relevance to the design of functional tissue-engineered scaffolds for load-bearing tissues is the development of high-performance composites such as glass, Kevlar and carbon fiberreinforced polymeric materials. The same principles of mechanics that have been applied to their design and fabrication have been adapted to build CAD/CAM programs for the design of orthopedic devices and biocomposite scaffolds (Harris 2008; Hollister 2002; Li 2007; Liu, 2007; Matsudo 2005; Ramakrishna 2001; Sun 2005). An analytical model that includes biomechanical responses of both the structural and cellular components, preferably at different hierarchical levels of tissue morphology, can be included
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with the classical theories of mechanics to describe a scaffold that mimics the behavior of an injured tissue. A recently issued patent on a method and system for modeling bone structure, for example, ‘discloses a structural and mechanical model and modeling methods for human bone based on bone’s hierarchical structure and its hierarchical mechanical behavior’ (Ascenzi 2008). The invention provides a FEA model that reflects the geometry, distribution and orientation of components, and the variation of porosity characteristics within the structure.
13.4.3 Manufacturing of tissue-engineered scaffolds Rapid prototyping Rapid prototyping (RP) is a generic classification of processes that fabricate a physical object by building a sequence of thin-layers, one at a time, to form a complete 3D structure directly from CAD/CAM sources. Two excellent reviews of design and fabrication of tissue-engineered scaffolds are Weigel (2006) with a list of 284 references and Moroni (2008) with a list of 256 references. One of the more widely used RP technologies is selective laser sintering (SLS). In this case, a laser beam, guided by a computercontrolled scanner, is selectively traced over a thin layer of densely packed thermoplastic polymer powder. The laser beam heats the powder to a temperature far enough above its glass temperature to sinter it, thereby creating a thin layer of precisely oriented thermoplastic fibers. A new layer of powder is then rolled over the first and a second layer of oriented fibers is formed. The process is continued until a multilayered scaffold structure is created. After the remaining loose powder is removed, a non-homogeneous scaffold of predetermined fiber volume fraction and porosity results (Weigel 2006). Three dimensional printing is similar to SLS, except that the laser is replaced by an inkjet head (Sachs 1989). Initially, a layer of powder is spread over a building platform. A multi-channel ink-jet printhead then deposits a liquid adhesive binder on the powder layer to bind powder particles into a precisely determined pattern, resulting in a thin slice of the structure. A powder delivery piston is then raised incrementally to spread a second layer over the first one, a roller compresses the powder and the deposition of liquid adhesive binder is repeated to establish the fiber orientations in the second layer. The process is repeated until the complete 3D structure of the scaffold is established. After drying of the liquid binder, loose powder is removed by an air jet and the finished scaffold is removed (Weigel 2006). Under certain conditions, cells and biochemical and mechanical signals that promote tissue in-growth can be printed and incorporated into prefabricated layers before final assembly (Anderson 2007; Hutmacher 2004; Mironov 2003; Weigel 2006; Zhang 2005).
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Fused deposition modeling is an RP technique for producing 3D scaffolds using thermoplastic fibers (Zein 2002). Plastic filaments or pellets are fed to an extruder with a CAM-controlled extrusion nozzle mounted on a mechanical stage capable of moving in horizontal and vertical directions. A 3D scaffold is constructed by depositing molten filaments with predetermined fiber orientations layer-by-layer according to the pattern determined by a CAM design strategy. Each layer can be bonded to the layer below and the size and shape of pores can be manipulated by modifying the angle of deposition between successive layers (Moroni 2008). A brief tutorial (Castle Island 2008) describes these and a number of other RP processes, such as stereolithography (SLA), laminated object manufacturing and laser powder forming. Electrospinning Electrospinning (ELSP) may be somewhat more adaptable than RP processes for mimicking the non-homogeneous, complex anisotropic character of some load-bearing tissues. The spinning of fibers from a biomaterial solution is accomplished by using a high-voltage electrostatic field operated between a delivery nozzle and a collection platform (Dalton 2006; Ma 2005; Mitchell 2006; Yoshimoto 2003). Suspended droplets of polymer melts or solution are electrically charged and injected onto a collection platform. The melts are cooled (or the solvents evaporated) to form a layer of fibrous material. The orientation of the fibers in the resulting 2D non-woven structure is determined by controlling the path of the delivery nozzle or the direction of the electrical field (Pedrozo 2008). The 2D profiles can be extended to 3D fibrous meshes by multiple level deposition and prolongation of processing time (Nair 2004; Weigel 2006). The porosity volume, pore size distribution, fiber volume fraction, fiber dimensions and fiber orientations can be controlled to favor cell transport, cell attachment and the ingrowth of new tissue (Ma 2005; Moroni 2008; Weigel 2006). Fibers with diameters ranging from micrometers to nanometers can be produced, similar to the dimensions of those in ECM of load-bearing tissues, thus making it possible for cells to express the proper proteins and signals (Ma 2005; Moroni 2008; Sell 2007). A recent patent describes an ELSP method for organizing assemblies of collagen and other fibrillar proteins capable of self-assembly using electric fields or electric currents to produce natural protein-based scaffolding (Pedrozo 2008).
13.5
In-silico computational analysis
While general continuum theories of mechanics can accommodate to the non-linear elastic, anelastic and viscoelastic behavior of load-bearing tissues (Lakes 1993, 1991; Mindlin 1965; Provenzano 2001), a robust bio-
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mechanical model for functional tissue-engineered scaffolds is still difficult to obtain because of the complexity of the non-homogeneous, anisotropic nature of natural tissues and a paucity of information on cell biology within a synthetic, porous scaffold. The on-going development of computational analysis will enable continued improvement of FEA-based models for use in CAD/CAM processes. The incorporation of stress-dependent morphogenetic effects in tissue remodeling appeared in the mid-1990s (Rodriguez 1994; Semple 2005; Taber 1995; Tracqui 1995). An analytical model relating cell movement and contact navigation to interstitial fluid flow and deformation of an anisotropic ECM (Barocas 1997) and a study of morphogenetic processes in the growth of vascular networks (Kiani 1991) are examples of research efforts leading to the use of in-silico computational analysis of tissue-engineered scaffolds. While not a substitute for in vitro and in vivo laboratory testing, in-silico programs provide an economical path for finding an optimal prototype scaffold design for in vivo testing. The mapping of the human genome in 2003 resulted in an enormous data pool of proteome activity in cells and developing organisms that is being assembled in the form of readily accessible computer-based software. Recent trends indicate a significant advancement in the use of these data to support the design of functional tissue-engineered scaffolds. ‘In-silico computational models provide a platform that is equivalent to an in-vivo, in-vitro, and in-situ or ex-vivo model platform’ (Knothe Tate 2007). They are computer representations of a system in which all parameters and system variables are expressed in a mathematical format, using analytic theories and empirical representations. The key elements in building a tissue model are: the functions to be replaced, e.g., load-bearing and biological functions; a control volume, e.g., an abstract representation of the tissue to be replaced; the governing equations used to predict effects of system variables on model behavior; and boundary and initial conditions (Knothe Tate 2007). Mechanical representations for the load-bearing structure include matrix representations of linearly elastic orthotropic lamina and laminates of composites (Agarwal 1980, p. 152), and analytical and empirical formulas for non-linear elastic, anelastic and viscoelastic behaviors (Lakes 1999, 1993, 1991). Mathematical representation of the effects of fluid transport and imposed forces on the rates of change of cell density and ECM density are also required (Semple 2005). The structure and function of a tissue are interdependent, and the interdependence may be different at different length and time scales. If one chooses a continuum model, the properties and variables are averaged over the total representative volume and the variations in scale cannot be dealt with individually. For example, in articular cartilage the cell distribution varies from the articular end to the region in contact with bone. The
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organization of collagen fibers and the shapes of the chondrocyte cells that secrete ECM of cartilage vary. The variation in morphology is directly related to changes in the biomechanical functions at different locations (Kim 2003). Stochastic network modeling techniques, widely used in engineering design, were employed for discrete modeling of bone tissue at length-scales of tissue, cell and molecular dimensions (Anderson 2005A; Knothe Tate 2007; Sidler 2006). The effects of structural and compositional changes on the flow of interstitial fluid through a pericellular network and through the matrix porosity were studied to determine the influence of osteocyte density and osteocyte connectivity on cortical bone permeability. The stochastic models were used to quantify the decrease in bone permeability owing to decreasing osteocyte density, and to examine transport through the pericellular network in order to predict the depth of penetration of specific molecules. All of the above-mentioned tools have been brought together in a processing strategy called ‘computer-aided tissue engineering’ (CATE).
13.6
Computer-aided tissue engineering
Computer-aided tissue engineering (CATE) ‘integrates advances in biology, biomedical engineering, information technology, and modern design and manufacturing to tissue engineering application’ (Sun 2005). The three components of CATE are characterized as: computer-aided tissue and bio-modeling; scaffold informatics and biomimetic design; and bio-manufacturing for tissue and organ regeneration (Sun 2004A, 2004B, 2002). The first stage in the construction of a CAD model for a composite scaffold is to obtain a CT (or other) 3D image of the tissue to be repaired, and to characterize its morphology and physical properties. The exact reproduction of the structure of natural tissues is difficult because the conversion of CT voxel-based 3D images of tissues to the vector-based modeling environment of a CAD is a complex process (Sun 2005). The generation of a CADbased model using a FEA with 3D unit cells represented by common structural shapes found in tissue morphology, however, is a practical method of modeling a prototype scaffold. A library of CAD-based unit cells is illustrated in Fig. 13.1 (Sun 2005). Included are unit volumes that characterize various woven textile fiber patterns and other shapes that define pore size, shape and distribution within the scaffold. A common strategy for guiding biomimetic regeneration of damaged tissue has been to generate a porous, 3D fibrous structure containing cells that secrete and maintain the ECM of the target tissue (Leong 2008; Ma 2006). The mechanical requirement is met by providing structural support, internal surfaces for cell attachment and physical compatibility at the
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13.1 Library of designed scaffold unit cells based on various feature primitives. (Reprinted from Sun et al., 2005, with permission from Elsevier.)
damaged tissue site. The biological requirement is met by creating channels of interconnected pores through which cells, signals and nutrients can be transported. To provide a favorable environment for cell proliferation and new tissue in-growth, the CAD-based program should define the placement of living cells, nutrients and growth factors at appropriate places in the interstices of the prototype scaffold. The designer has a wide range of natural and synthetic biomaterials from which to choose (Ramakrisha 2004; Salgado 2004). Ideally, the biodegradation rate of the scaffold structure is selected as that which permits the growth of new tissue to pick up loadbearing capacity at a rate that maintains an adequate structural strength at the damage site. Selection of an appropriate CAD/CAM technique for the fabrication of prototypes (Sun 2004A, 2004B, 2002) completes the design process.
13.7
Case studies
13.7.1 Case study 1: a tissue-engineered scaffold for bone repair Injectable particulate composites of cross-linkable poly(propylene fumarate) (PPF) have been proposed as scaffold materials for orthopedic
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applications (Temenoff 2000). Osteoblasts were encapsulated in crosslinked gelatin microparticles and seeded into crosslinking and fully crosslinked PPF composites to promote cell proliferation and phenotypic expression (Payne 2002A, 2000B). Although in vivo tests had shown that the seeded, photo-crosslinked PPF elicited a bone healing response, it was recognized that the internal pore space, having been obtained by leaching of salt crystals (NaCl) from the scaffolds, could not be controlled (Fisher 2001). Optimization of the porous structure was accomplished by utilizing a stereolithographic, 3D printing technique to fabricate scaffold prototypes (Cooke 2002). A biodegradable resin mixture of diethylfumarate (DEF), PPF and a photoinitiator, bisacylphosphine, was chosen for the scaffold structure. A stereolithography machine, (SLA 250/40), (3D systems, Valencia, CA), with a servomechanism that controlled an ultraviolet laser was employed to photopolymerize the resin one horizontal layer at a time on a build platform controlled to move vertically. A 3D model of the prototype was designed using a CAD software package, Pro/ ENGINEER 2000i (PTC, Needham, MA), and exported as an STL file into a preprocessing software package, Maestro (3D systems, Valencia, CA). An in-silico computational model of the fluid transport in the ECM (porous volume) of the prototype scaffolds was employed to optimize cell viability and tissue in-growth. The resulting CAD/CAM processing system was used to construct the DEF/PPF prototype scaffolds (Knothe Tate 2007). The scaffold prototypes (Fig. 13.2) were 3D layered cylinders with nine circular and four semi-circular channels in the longitudinal direction (Anderson 2005B; Knothe Tate 2007). All channels were connected through seven transverse rectangular channels. A software package (CFD-ACE, Huntsville, AL) was employed to predict flow rates through the target design scaffolds. Flow through the scaffold volume was induced by a pressure gradient across a fluid medium ‘idealized’ as water. The mass flow rate through the mesh created by the software package was calculated, the permeability in the longitudinal direction was determined using Darcy’s law, and validated experimentally in the actual manufactured scaffolds, by using the same mass flow rate (Fig. 13.2) (Anderson 2005B; Knothe Tate 2007). To predict the mechanobiological behavior of cells seeded on the scaffolds, CDF software was used to simulate Navier–Stokes steady-state flow through the scaffolds. When flow was simulated in the longitudinal direction (i.e., from top to bottom in the cylindrical scaffold), it was observed that the velocity profiles in the through direction were typical of Poiseuille flow with a velocity higher than that in the larger transverse channels. The velocity profiles in the transverse channel were typical of jet-flow expansion. ‘The high-flow environment of the through channels produces high shearing stress along the longitudinal wall. In contrast, the low-velocity transverse
(a)
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Calculated permeability Permeability prediction – curve fit y = 3.6367x4 – 4.5575x3 + 2.2324x2 – 0.3116x
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13.2 Comparison of scaffold parameters: (a) computer-aided design (CAD) drawing of target scaffold geometry; (b) microcomputer tomographic (mCT) image of actual manufactured prototype geometry; (c) predicted permeability (k) of scaffold. (Reprinted from Knothe Tate, 2007, with kind permission of Springer Science and Business Media and of M. Knothe Tate.)
layers provide low-level shear stresses that promote cell adhesion, as well as mechanical stimuli conducive to osteoblastic differentiation.’ (Knothe Tate 2007). The computational analysis was able to distinguish between impermeable and permeable proposed prototypes, thereby confirming the usefulness of in-silico modeling. This case study, carried out by several laboratories working in concert, clearly demonstrate the power of in-silico computational modeling in the manufacture of functional tissue-engineered scaffolding.
13.7.2 Case study 2: design of composite femoral prostheses Over the past twenty years there have been many attempts to use fiberreinforced biomaterials in the design of proximal femoral prostheses (De Santis 2000, 2004; Huiskes 1993; Simoes 2000). In recent years, conventional mechanics-based studies have been augmented by using medical imaging data to construct CAD/CAM-based models. A critical step in the adoption of these programs is development of an effective conversion of CT data into a format required by CAD-based modeling. Three process paths for generating a suitable CAD model from medical imaging data were evaluated for femur model generation (Sun 2005). Sun and colleagues compared a commercial MEDCAD interface with a reverse-engineering interface approach and a STL-triangulated model approach. The MedCAD interface,
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developed by Materialise, Belgium, exports data from the imaging system to a CAD platform by fitting primitive unit cells, such as those shown in Fig. 13.1, for characterization of 2D segmentation slices of tissue. The reverse engineering interface process path employs a 3D voxel model created from segmentation and converts it to a point data format. The points are loaded into a reverse engineering software package and triangulated to form a faceted model. The STL-triangulated model process path converts the 3D voxel model to an STL file, which is then imported into the reverse engineering software as an STL-triangulated surface model. The three process paths were compared for the generation of a CAD model from CT images of a proximal femur bone from a small child. It was found that the MEDCAD interface lacked appropriate primitive features to accurately characterize the femur, which required them to employ freeform shapes to better describe its anatomy, and a primitive sphere to represent the top of the femur. The STL triangulated model had difficulty in representing the geometry of small and complex features without modifications of the modeling process. The reverse-engineering approach appeared to be the preferred modeling approach because of ‘the accuracy, structure fidelity, and the versatility in data transfer to STEP or IGES’ (Sun 2005). The Bio-CAD modeling process and its application to tissue scaffold design were illustrated by using CT histograms to characterize the spatial heterogeneity and porosity of the proximal femur. A ‘homogenization technique’ is described whereby the femur structure is divided into seven 10 mm layers and an average quantitative computed tomography number (QCT#) is determined for each layer from the CT histograms. The QTC# was then correlated with tissue density (ρ) and, in turn, the density was correlated with average elastic modulus (E) within each 10 mm layer. The structural non-homogeneity and porosity of the bone were thus modeled and appropriate unit cells were integrated to generate a bone tissue model scaffold. Hydroxyapatite and two different polymeric biomaterials, L-PGA and L-PLA, were chosen as the components for the scaffold. The component compositions and the sizes of the unit cells were varied using FEA software (ABAQUS) to model a range of porosities, and other morphological and mechanical properties. From QCT measurements of the original proximal femur, they determined that the bone Young’s modulus varies from 0.6–2.0 GPA, and from ‘a biological point of view’, the scaffold structure should have porosity in the range of 55–70%. From a range of FEA predictions, they found that ‘L-PLA based unit cells with around 40–60% porosity do give a better option as a scaffold material for the proximal femur’. The CATE process can then be completed by building an appropriate ‘BioCAD’ model for the development of a CAM program to guide an RP process for fabrication of proximal femur prototypes. The authors point out that Bio-CAD based CATE is still in an early stage of development, but
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eventually will provide ‘precisely controlled architecture and multi-material printing for different types of biological factors, cells and scaffold materials’ (Sun 2005, p. 1113).
13.7.3 Case study 3: design of functional articulate cartilage scaffolds Being a non-homogeneous soft tissue with a limited regenerative capacity, articulate cartilage offers difficult challenges in the design of a scaffold that supports the cell attachment and proliferation necessary for tissue repair (Mankin 2000; Woodfield 2004). A ‘state of the art’ microscale weaving technique was used to design a tissue-engineered scaffold for cartilage repair (Moutos 2007). A continuous multi-filament polyglycolic acid yarn was chosen to generate porous 3D orthotropic fiber-supported structures with fibers oriented in three orthogonal directions. The scaffold consisted of five woven layers oriented in the warp direction (lengthwise in the loom) and six layers oriented in the weft direction (90 ° relative to warp). Specification of fiber diameter, spacing and volume fraction throughout the volume of the scaffold allowed for a controlled anisotropy, a variable distribution of interconnected pores and a mechanical response mimicking that of the target tissue. A vacuum-assisted infusion of a biocompatible cell-loaded hydrogel filled the pore structure to form an ECM with a spatially uniform distribution of cells capable of promoting chondrocytic phenotype. The mechanical properties of two types of woven structure containing three types of hydrogels (in the absence of cells) were compared with those of native articular cartilage. With the exception of some tensile properties, the mechanical behavior of the 3D woven composite scaffolds ‘mimicked the anisotropy, viscoelasticity and tension–compression non-linearity of native articular cartilage’. In the absence of direct experiments measuring cell activity in the scaffolds, the similarity of the apparent hydraulic permeability of the 3D woven composites to that of native cartilage was used by the authors to attribute a biomimetic character to the scaffolds. In a similar study, the biochemical properties of 3D poly(ethylene glycol/ (polybutylene terephthalate) (PEGT/PBT) block copolymer scaffolds were evaluated (Woodfield 2004). The mechanical properties of various scaffolds were compared with normal articular cartilage, and their biochemical properties were assessed using in vitro and in vivo tests. A custom-built 3D deposition device was used to dispense molten polymer through a syringe and nozzle device using needles of varying diameter to allow changes in mechanical properties and porosities from layer to layer in the scaffold. Fiber layers were deposited in a 0–90 ° orientation pattern in the construction of 4-mm-thick scaffolds (approximately the thickness of human articular knee cartilage). Two scaffold types were generated for comparison using
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a hydrophilic copolymer 1000/70/30 PEGT/PBT with a swelling capacity of 18%, and a hydrophobic copolymer 300/55/45 PEGT/PBT with a swelling capacity of 2%. The measured volume percent porosities were approximately 55–72 vol% and 71–87 vol%, respectively. Mechanical properties were varied by controlling the scaffold architecture. The static equilibrium compressive moduli of the scaffolds varied from 0.05 to 2.5 MPa for fiber spacing of 2 to 0.5 mm. Reduced fiber spacing resulted in lower porosity and higher equilibrium compressive modulus. The dynamic stiffness moduli were 0.16 to 4.33 MPa at 0.1 Hz. The moduli were a function of porosity, the properties of the fiber components and the geometric arrangement from one layer to the next. The properties of the prototype scaffolds were compared with those of a human articular knee cartilage with an equilibrium compressive modulus of approximately 0.6 MPa and dynamic stiffness modulus of 4.5 MPa at 0.1 Hz. Upon examination of the available data, it appeared that the 300/55/45 hydrophobic PEGT/PBT with fiber separation at 1.0 mm most closely resembled the native cartilage from a dynamic stiffness point of view, but the hydrophilic copolymer 1000/70/30 PEGT/PBT with fiber separation between 0.5 to 1.0 mm was a better match with native cartilage with respect to static compressive modulus. In vitro cell culture experiments were conducted by seeding and culture of bovine chondrocytes on 300/55/45 scaffolds with a 1 mm fiber spacing. To study chondrogenesis in vivo, 4 mm × 4 mm thick cylinders cored from the 3D scaffolds were seeded, cultured and implanted in subcutaneous pockets of 6-week-old nude mice. The animals were sacrificed at 21 days after implant for histological study. Cell attachment onto and throughout the scaffolds seeded in vitro with bovine chondrocytes occurred rapidly with an early onset of chondrogenesis. ‘There was near-complete formation of cartilage-like tissue within the interconnecting pores after 21 days of culture.’ A thin capsule of fibroblast-like cells was also observed surrounding the periphery of the scaffold. The ECM of the subcutaneous implants exhibited a cartilage-like morphology throughout and fibroblast-like cells on the periphery with glycosaminoglycans beyond the border of the scaffold fibers. Scaffolds of 300/55/45 copolymer with a staggered spacing between successive layers of 1 mm were seeded with human articular chondrocytes isolated from patients undergoing hip replacement surgery. In vitro seeding resulted in rapid cell attachment and proliferation, a bridging of interconnecting pores between staggered fibers and a homogeneous distribution of living cells after 5 days’ dynamic seeding. The data suggest that a 3Ddeposited PEGT/PBT scaffold that matches the static and dynamic mechanical properties of human articular knee cartilage and is capable of supporting the cell attachment and proliferation necessary for tissue repair can be attained.
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13.7.4 Case study 4: tendon and ligament tissueengineered scaffolds The use of 3D braided structures of poly(l-lactic acid) (PLLA) fibers as a functional tissue-engineered scaffold for the repair of the anterior cruciate ligament (ACL) was reported by Laurencin and colleagues (Laurencin 2005). The ACL is a cable-like tissue composed of collagen fibers, elastin, proteoglycans, water and cells. The collagen fibrils display a periodic crimping pattern that repeats every 45–60 μm and the ligament is surrounded by a sheath of vascularized epiligament. The collagenous network is also twisted by approximately 180 ° from the femoral attachment to the tibial attachment site. In designing the prototype scaffold, the authors took into account the ligament’s triphasic behavior when exposed to strain. When an initial force is applied to the tissue, ‘un-crimping’ of the collagen fibers and a release of water occurs, resulting in a low stress per unit strain. Once the fibers are stretched, there is a stiffening of the structure and a more linear increase in stress until a point is reached where interfibrillar slippage occurs. In the third phase of mechanical behavior the collagen fibers fail by defibrillation causing a loss of stiffness and ligament failure. A 3D braiding technique was used to create a scaffold with ‘controlled pore size, integrated pores, resistance to wear and rupture, and mechanical properties comparable to natural ACL’. The PLLA fibers of the 3D braided structures were similar in diameter to collagen fibers and wound in bundles throughout the thickness of the scaffold to create a hierarchical structure similar to that of the ligament (Fig. 13.3). Variations in fiber orientation cause changes in the pore size within the three regions. The pore size distribution is meant to ‘encourage tissue (ligament and bone) in-growth and capillary supply’. The prototype scaffolds, designed for in vivo study in rabbits, had mechanical properties similar to those of rabbit ACL. After one week of in vitro exposure to ACL fibroblasts, the scaffolds exhibited extensive cellular proliferation and tissue growth along the longitudinal axes of the fibers with ECM bridging between fibers. Absorption of growth factor fibronectin on the surfaces of the PLLA fibers increased cell proliferation. The authors state that ‘their design allows these scaffolds to degrade while promoting tissue growth, enabling the body to fully regenerate lost or damaged tissue without risk of scaffold or neoligament rupture or stress shielding of the new tissue.’ The study illustrates the importance of hierarchical structure at different levels of organization in the design process. In another research study, the importance of fiber organization on the attachment, alignment and gene expression of human rotator cuff fibroblasts on poly(lactide-co-glycolide) (PLGA) scaffolds was examined (Moffat 2008). Nanofiber scaffolds were produced from a polymer solution using an electrospinning process. Fibroblast activity was compared on
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Bony attachment end
Intra-articular zone
Bony attachment end
PLLA 3-D braid
13.3 Design of 3-D braided ligament scaffolds showing the macrostructure of the bony attachment and intra-articular zone. (Reprinted from Laurencin and Freeman, 2005, with permission from Elsevier.)
unaligned (nearly randomly oriented) and longitudinally aligned nanofiber scaffolds. Both types had similar scaffold thicknesses, average fiber diameters, average pore diameters, percent porosities and permeabilities. Thus, differences in mechanical and biological behaviors were attributed to nanofiber organization. As expected, the longitudinal elastic modulus and tensile strength of the aligned scaffolds were three times and four times greater, respectively, than those of the unaligned scaffolds. The rotator cuff fibroblasts cultured in vitro on the aligned scaffolds were oriented in the direction of the longitudinal axis and adopted a phenotypic elongated morphology. In contrast, fibroblasts cultured on the unaligned fibers exhibited a polygonal morphology without preferential orientation. The alignments and orientations were quantitatively similar to the mean angular distributions of the underlying nanofiber substrates. There was no significant difference between the two scaffold types in human rotator cuff fibroblast proliferation rate. The gene expression of integrins and type I and type III collagen were also compared. The α2 expression detected on the aligned scaffolds was significantly higher over time than that on the unaligned scaffolds, while αV and β1 expression levels on the unaligned fibers were significantly higher over time than that on the aligned scaffolds. The authors point out that elevated expression and production of αV and β1 have been associated with healing of tendons and ligaments, ‘which may result in disorganized scar tissue formation’. In contrast, the aligned nanofiber scaffolds circum-
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vent such a healing response and directly promote the deposition of a biomimetic collagen matrix instead’.
13.8
Future trends
The last decade has marked a major advance in design and fabrication of functional tissue-engineered scaffolds. The primary reason has been the dramatic increase in interdisciplinary research and engineering in imaging, modeling and simulation of human anatomy and physiology, much of which has arisen from the first draft of the Human Genome Project on the DNA reference sequence for humans (Nature 2001). Long-range strategic plans for basic research in the biological sciences and medicine, and for development of computer and engineering technologies related to public health, were initiated in the USA by the National Institute of Health (NIH 2008), in the EU by the efforts of the Computational and Information Technologies Society (DG INFOSO&DG JRC, 2005) and in the UK by the Institution of Mechanical Engineers (Hoare 2004). Recent advances in medical imaging and computational analysis have provided the opportunity to meld biological function, structural design and rapid prototype manufacturing in the creation of a functional tissue-engineered scaffold. Medical imaging techniques continue to be refined. Electron beam CT scanners produce accurate, detailed images in fractions of a second, even permitting the scanning of a beating heart, for example, and requires less than 20 percent of the radiation dose of conventional CT scanners (Tallisetti 2004). Computed tomography (CT) can be combined with positron emission tomography (PET) to promote co-registration of data into a single, more precise image. The PET component also has the potential to measure tissue functions such as blood flow. Diffraction enhanced imaging is being developed that delivers a more detailed image of soft tissue than is possible with standard CT or MRI techniques (Zhong 2004). Continued exploration of new process paths for generating CAD/CAM-compatible data presentations from medical imaging data will add precision to the rapid prototyping of patientspecific scaffolds. As the empirical data pools on genetic and proteome activity continue to expand, the use of in-silico computational analysis will grow. The same tools and analysis strategies used in the human genome project are now being applied to analyze the behavior of proteins and cells in tissues and organs. Thus, computational biology will be a primary tool in regenerative medicine and tissue engineering (Semple 2005). Another aspect of the design of load-bearing scaffolds is the choice of biodegradable material used for the support structure. The primary objective has been to choose a biocompatible material that optimizes structural integrity during the formation of new tissue in-growth. New biodegradable biomaterials are being designed that elicit cellular responses at the
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molecular level (Hench 2002; Lutolf 2003; Salgado 2004). These are molecular-modified biopolymers that stimulate specific interactions with cell integrins that result in cell proliferation, differentiation and ECM production and organization. The long-term challenge of creating a patient-specific scaffold that precisely mimics the mechanical and biological nature of the original tissue is to overcome the primary difference between an engineering approach to functional tissue design and ‘nature’s way’. The engineer can select from a wide array of natural and synthetic biomaterials, but must fabricate a scaffold by building it from an idealized representation of the final product. While computational analysis and rapid prototyping are powerful tools, they cannot currently recreate accurately the smallest scale hierarchical features of natural load-bearing tissues or the complex, non-homogeneous features of a natural ECM. Nature starts at the molecular level and uses the principles of self-assembly to manufacture tissue from the molecules and cells available in the whole organism. This provides control at all levels of hierarchy and results in a self-reparable tissue. There is a considerable amount of research activity in the area of self-assembly at the nanoparticle and molecular level (Jakab 2004; Whitesides 2002A, 2002B). Progress in achieving directed self-assembly of nanoparticles carrying complementary strands of DNA has recently been reported (Crocker 2008; Nykypanchuk 2008; Park 2008). The particles were found to selectively adhere to one another in a precise geometry when grafted complementary DNA strands ‘hybridized’ to form a double helix conformation. The final architecture was determined by the lengths and nucleotide sequences of the DNA strands. It appears to be an assembly scheme that might be applicable to directional bonding in complex structures. ‘The ultimate dream is the creation of a DNA tool-kit that will make possible the self-assembly of nearly any material reliably at the nanoscale’ (Crocker 2008). The application of selfassembly principles to construction of a functional tissue-engineered scaffold is still in its infancy, but is one of the more likely routes to manufacturing truly biomimetic scaffolds.
13.9
Acknowledgement
The authors wish to thank The Institute of Materials Science of the University of Connecticut for providing funding and access to its facilities during the preparation of this chapter.
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14 The challenge of biocompatibility evaluation of biocomposites J. M. A N D E R S O N and G. V O S K E R I C I A N, Case Western Reserve University, USA
Abstract: The evaluation of biological response, i.e., biocompatibility, of biocomposites is complex, complicated and challenging. Biocompatibility studies of biocomposites present two major challenges: 1) identification of the overall biocompatibility (safety) of the biocomposite and identification of the contributions made by each of the distinct phases of the biocomposite structure at the tissue/biocomposite interface; and 2) surface characterization of the biocomposite structure where the distinct characteristics of each of the phases are appropriately and adequately characterized. The recent use of bioactive materials, tissue engineering, and nanotechnology in the development and/or utilization of biocomposites further enhance these challenges. This chapter focuses on the challenges that must be met in the future development of biocomposites. Key words: biocomposites, biocompatibility, biological response evaluation, nanotechnology, tissue engineering.
14.1
Introduction
Composites generally are considered to be materials consisting of two or more chemically distinct constituents having a distinct interface that separates the constituents. In addition, composites generally are composed of one or more discontinuous phases embedded within a continuous phase. From an engineering plastics perspective, the discontinuous phase is usually harder and stronger than the continuous phase and is called the reinforcement or reinforcing material, while the continuous phase is termed the matrix. In medical science, bone and tendon have long been considered to be biological composites with a number of levels of hierarchy or, organization. Synthetic composite materials have been fabricated to provide desired mechanical properties such as strength, stiffness, toughness, and fatigue resistance for biomedical application as substitutes for hard tissues. These properties of composites are strongly influenced by the properties of their constituent materials, their distribution and content, and the interaction between the discontinuous reinforcing material and the continuous matrix. 325
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To properly describe a composite material, specifications regarding the constituent materials and their properties must be identified in addition to the geometry of the reinforcement or reinforcing material, their concentrations, distributions, and orientations. These properties and characteristics ultimately determine the biocompatibility and functional success or failure of medical devices and prostheses that utilize composites in their construction. In the past, composites were generally characterized on the basis of their macro-scale properties with an emphasis on their bulk mechanical properties and little or no focus on their surface properties. Recent advance in tissue engineering, controlled release science and technology, and nanomaterials, as well as increasing knowledge of the mechanisms of molecular and cellular interactions with materials, has led to new insights, perspectives, and concepts for the development and application of composites in biomedical science. Biocompatibility studies of composites present two major challenges. These are 1) biological response evaluation to identify the overall biocompatibility (safety) of the composite and the contributions made by each of the distinct phases of the composite structure at the tissue/composite interface; and 2) surface characterization of the composite structure so that the distinct characteristics of each of the phases of the composites are appropriately and adequately characterized. In the past, biocompatibility and/or biological response evaluation of so-called ‘inert’ composites was easily carried out utilizing ISO, ASTM, and USP protocols. The recent advent of tissue engineering, regenerative medicine, and nanotechnology, and the development of bioactive materials, all of which are now being used in the development of a new generation of composites, has added new challenges to the determination of biocompatibility. The biocompatibility of a material is generally defined as the ability of a material to perform with an appropriate host response in a specific application. The host response generally is considered to be determined, controlled, or modulated by the surface of the implant. For composites with a single continuous material phase at the tissue/device interface, biocompatibility evaluation is straightforward. However, for composites with two or more distinct material phases (chemistries) at the tissue/device interface and with or without biodegradable and/ or bioactive components, adequate and appropriate biological response evaluation, i.e., biocompatibility, is much more complex, complicated, and challenging. As an example of the complex nature of biocomposites, nanocomposites containing bioactive materials have been suggested for hard-tissue engineering and composites utilizing Bioglass® have facilitated the formation of new bone with mechanically strong bonding to the implant surface. The in vitro dissolution products of bioactive glasses have been shown to upregulate seven families of genes in primary human osteoblasts (Hench 2002).
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These activated genes express numerous proteins that influence all aspects of differentiation and proliferation of osteoblasts: transcription factors and cell-cycle regulators; signal transduction molecules; proteins involved in the DNA synthesis, repair, and recombination; growth factors and cytokines that influence the inflammatory response to the materials; surface-cell antigens and receptors; extracellular matrix (ECM) components; and apoptosis (programmed cell death) regulators (Jones et al. 2007; Jell et al. 2008). The upregulation of these proteins generated by the solubilization of bioactive glasses clearly can affect the biocompatibility and biological responses of nanocomposites using bioactive glass as a component. While biocompatibility and biological response studies of the parent materials have suggested the potential use of these materials in macro-, micro-, and nanocomposites, biocompatibility studies of the composite structures themselves and identification of the specific roles played by the continuous and discontinuous phases at the tissue/composite interface remain to be investigated. Differentiating biological responses generated by each of the phases in these composite structures remains a significant challenge to their clinical application. Appropriate biocompatibility evaluation of these types of composites also would require a correlative surface analysis study for the determination of the geometric and compositional variations in the two different phases at the tissue/composite interface. A systematic approach to the multiple challenges of determining the biocompatibility of biocomposites requires an understanding of the initial inflammatory reactions, foreign body reaction, and wound healing responses. Following an overview of biocompatibility and the biological environment, a discussion of material effects and biological responses is presented with an emphasis on the surface characteristics of respective biocomposites. This review ends with a discussion of the challenges of elucidating symbiotic relationships between composite surface structure/chemistry and biocompatibility.
14.2
Biocompatibility and the biological environment
Implantation of a biocomposite, as with any other medical device or prosthesis, activates host defense systems. The continuum of host reactions following implantation of a biocomposite is presented in Table 14.1. The surgical implantation procedure itself induces injury, immediately followed by blood–material interactions that include protein adsorption to the surfaces of the biocomposite. The surface chemistry of the implant has been identified as the major modulator of the material surface adsorbed protein layer. The interaction of adsorbed proteins with adhesion receptors present on the inflammatory cell populations constitutes the major cellular recognition system for implantable materials. Protein adsorption remains a poorly
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Biomedical composites Table 14.1 Host reactions following implantation of biocomposites Injury Blood-material interactions Provisional matrix formation Acute inflammation Chronic inflammation Granulation tissue Foreign body reaction Fibrosis/fibrous capsule development
Table 14.2 Principal cell types in the inflammatory, foreign body, and wound healing response Response
Cell type
Acute inflammation Chronic inflammation Granulation tissue Foreign body reaction Fibrous encapsulation
Polymorphonuclear leukocytes (neutrophils) Monocytes and lymphocytes Fibroblasts and endothelial cells (capillaries) Macrophages and foreign body giant cells Fibroblasts
studied area in biomaterials science and little is known regarding the affinity, concentration, denaturation, and transient adsorption–desorption phenomena that occur on biomaterial surfaces (Anderson 2001; Andrade and Hlady 1987; Brodbeck et al. 2003; Chinn et al. 1992; Horbett 2004; Jenney and Anderson 2000a, 2000b; Tang and Hu 2005). Thus, it can be readily appreciated that the complexity of biocomposite biocompatibility is achieved within seconds of implantation of a biocomposite. Provisional matrix formation that occurs on the protein-adsorbed biocomposite surfaces is simply the formation of a platelet/fibrin blood clot/thrombus that, in essence, serves as a scaffold for the subsequent inflammatory, foreign body reaction, and wound healing responses. The principal cell types involved in the inflammatory, foreign body, and wound healing responses are presented in Table 14.2 (Anderson 2004). Acute inflammation is initiated almost immediately with the migration of polymorphonuclear leukocytes (neutrophils) into the implant site. These cells predominate during the first few days following injury and then they are replaced by monocytes and lymphocytes as the predominant cell type. Neutrophils are short-lived, disintegrating and disappearing within 24 to 48 h following implantation. The presence of monocytes and lymphocytes characterizes the chronic inflammatory response. Lymphocytes present at
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this stage of response are considered to be a component of the innate immune response and not the acquired immune response. Monocytes immigrate from the vasculature and may adhere to the protein-coated biocomposite surface leading to the foreign body reaction, or, may be present in the provisional matrix that is undergoing remodeling with dissolution of the platelet/fibrin blood clot and replacement by migrating fibroblasts that produce collagen and capillary ingrowth into the implant site (Anderson 2001, 2004; Shen et al. 2004). Monocytes at the biocomposite surface differentiate into macrophages that are very long-lived (up to months). The differentiation of monocytes to macrophages on implant surfaces leads to macrophage fusion with the formation of foreign body giant cells (Anderson 2001, 2004). These cellular reactions at the biocomposite surface occur within the first week following implantation and continue for extended time periods. The granulation tissue response surrounding the implant continues with the deposition of collagen and the formation of the fibrous capsule (Anderson 2004; Sharkawy et al. 1997). These parallel pathways are presented in Fig. 14.1. In general, these parallel responses occur in a simultaneous manner and timeframe; however, there are significant differences. The fibroblasts and capillaries in the granulation tissue eventually atrophy, may undergo
Injury, implantation Inflammatory cell infiltration PMNS, monocytes, lymphocytes Exudate/tissue
Biomaterial
Acute inflammation Mast cells
IL-4, IL-13 Monocyte adhesion Macrophage differentiation
PMNs Chronic inflammation Monocytes Lymphocytes Granulation tissue
Macrophage mannose Receptor upregulation Th2: IL-4, IL-13
Macrophage fusion
Fibroblast proliferation and migration Capillary formation Fibrous capsule formation
Foreign body giant cell formation
14.1 Tissue or cellular host response to injury.
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apoptosis, and lead to an acellular, avascular fibrous capsule. At the surface, macrophages and foreign body giant cells are present for the lifetime of the implant. Retrieved medical devices from humans have indicated that these cell types at the tissue/material interface are present for decades. The size, shape, and chemical and physical properties of the biomaterial or device are responsible for variations in the intensity and duration of the inflammatory and wound healing processes. Thus, intensity and time duration of the inflammatory reaction may characterize the biocompatibility of the material. With biocomposites, the size of the discrete material phases present at the surface may determine variations in the activity of these cells as well as their contributions to the overall inflammatory response, foreign body reaction, and fibrous capsule formation. In vivo tests for biocompatibility are indicated in Table 14.3 (Anderson 2004). These tests may use the biocomposites themselves or may utilize extracts of the biocomposites to determine the toxicity and bioactivity of potential leachables, dissolution products, or other materials utilized in the synthesis and production of the biocomposite. From a regulatory perspective, these tests are performed on the final product or ‘as used’ medical device following sterilization. Thus, the material characteristics for biocompatibility under a regulatory program are much more specific and controlled than the use of these tests to determine biological response evaluation for research and development purposes. It is significant that both types of biocompatibility or biological response evaluation programs require a broad and in-depth understanding of the surface properties of the materials or, in this case, biocomposites under investigation. Biocompatibility evaluation and surface analysis and characterization both are required for an adequate and appropriate evaluation of
Table 14.3 In vivo tests for biocompatibility Sensitization Irritation Intracutaneous reactivity Systemic toxicity (acute toxicity) Subchronic toxicity (sub-acute toxicity) Genotoxicity Implantation Hemocompatibility Chronic toxicity Carcinogenicity Reproductive and developmental toxicity Biodegradation Immune responses
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the biocompatibility or safety of the biocomposite under consideration. Furthermore, the unique characteristics of a biocomposite may require not only unique approaches to the utilization of these types of tests, but also may require additional tests developed specifically to address the design criteria under which the biocomposite has been developed and the proposed function of the biocomposite.
14.3
Surface effects and characterization
In composites where both the continuous (matrix) phase and the discontinuous phase are present at the surface or interface, appropriate and adequate characterization is required of both phases at the interface. The given dimensional and geometric parameters as well as the spatial distribution of these phases at the interface present significant challenges in characterization of the composite. While the continuous (matrix) phase may be at the micro- or macro-level, the discontinuous phase is usually micro or nano in scale. The advent of nanotechnology and the utilization of nano-materials in a wide variety of shapes and sizes offer real challenges to their characterization at the interface/surface of biocomposites. Surface evaluation in biocompatibility assessment is a fundamental requirement in the development of medical devices. Understanding the surface structure of materials leads to selecting the appropriate investigative analyses that will provide an in-depth assessment of the biocompatibility induced by the presence of materials in the body. The surface region of a material represents the interface between the bulk and biological environment. Its exposure to the surrounding biological environment assigns direct reactivity to it, as the surface directly interacts with the proteins, cells, tissues, and organs. The surface of the material has been recognized to be different from the bulk for a number of reasons, including: 1) spatial relevance (material surface, 2D, versus material bulk, 3D), 2) processing and manufacturing methods that tend to affect the exposed material surfaces rather than the bulk, 3) additives to enhance the properties of the materials, and 4) contamination (oxidative processes or impurities owing to environmental exposure outside of the intended use). For composites, there is a tendency for one of the components to dominate the surface, thus setting the stage for the perceived reactivity of the material by the biological environment (biocompatibility). This process is further complicated by considering the continually changing surface properties (physical and chemical) of biodegradable composites, where the bulk, over time, becomes the surface. Recognizing the potential complexity of a dynamic composite material surface represents the first step in intelligently selecting the adequate methods of analysis necessary for the assessment of its biocompatibility.
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A fundamental aspect that arises when speaking of material surface is the depth of what is defined as the surface. This critical aspect determines the types of analyses that can be performed as techniques are characterized by various limitations in the depth of analysis. Ratner (2004) has defined the surface depth as the zone where the structure and composition, influenced by the interface, differs from the bulk composition and structure. A generic guideline associates the surface depth of atomic materials to five atomic layers (0.5–1 nm) whereas that of polymers can extend from 10 to 100 nm, depending on the polymer system, the chain’s molecular weight and environmentally dependent chain mobility. Thus, it has been established that the material surface, as a transduction mechanism or facilitator, directly influences the response of proteins and cells. Consequently, manipulation of surface chemistry and topographic features has been identified as conducive to modulating protein and cellular response, thus, biocompatibility.
14.3.1 Surface chemistry The surface chemistry of a biomaterial is generally considered to influence protein adsorption and cellular adhesion (Jenney and Anderson 2000a, 2000b; Senaratne et al. 2005). In general, surface chemistry of any material could be characterized as hydrophobic or hydrophilic. It has been recognized that hydrophobic surfaces adsorb a high density monolayer of proteins leading to cellular adhesion, while hydrophilic surfaces are more likely to resist protein adsorption and, implicitly, resist cellular adhesion (Hoffman 1999; Vert and Domurado 2000). Within the context of biocompatibility, a hydrophobic or hydrophilic surface can be desirable, depending on the intended clinical application. Certain circumstances call for native tissue integration within and around the implanted material/device where the biocompatibility of the material is judged by its capability to attract and support native cells (e.g., orthopedic implants, surgical meshes for soft tissue reinforcement). In other cases, a non-fouling or stealthy surface may be required to prevent the process of cellular migration, adhesion, and proliferation from being established (e.g., urinary catheters, endotracheal tubes). Protein adsorption represents the first step in the process of material surface–biological environment interaction. This critical step sets the stage for the manner in which the material ultimately will be perceived by the local environment. This phenomenon has been the subject of many investigations. While its recount is not the subject of this chapter, understanding the process of protein adsorption is critical in the design and performance of biomaterials. Water molecules have been found to bind on hydrophobic (structured water) and hydrophilic (primarily via hydrogen bonds) surfaces
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(Antonsen 1992; Hoffman 1986). The capability of hydrophilic surfaces to allow the formation of water-initiated hydrogen bonds to neutral polar groups, as well as to ionic groups belonging to the surface chemistry of the composite, leads to the formation of a low reactivity water monolayer, thus reducing the protein adsorption and cellular adhesion (Luk 2000).
14.3.2 Surface morphology The surface hydrophobicity and hydrophilicity of a biocomposite may be determined by the composition and surface morphology of the composite. In particular, in the case of polymer based composites, the mobility and other characteristics of the polymer chain can interfere with the protein adsorption process. Such phenomena of entropic and osmotic repulsion by the surface polymer chains can certainly affect the biocomposite behavior towards cells (Ratner 2004). The surface polymer chain mobility of polymer-based composites also will affect the entropic and osmotic properties of the material. Recent investigations have strongly suggested that the chain segments near the free surface of a material, or in thin films, are not affected in the same way as the bulk material. Fackhraai and Forrest (2008) reported that surface relaxation was observed over a wide range of temperatures, providing evidence for enhanced surface mobility relative to the bulk. Furthermore, in a series of articles, Andersson and co-workers suggested a correlation between surface chain mobility and soft tissue reactions (Berglin et al. 2004; Andersson et al. 2007, 2008). Using the glass transition temperature (Tg) property of a polymer, they evaluated the effect of surface mobility on the inflammatory and wound healing response in vitro and in vivo and proposed that increased surface chain mobility led to significantly less cellular presence and prolonged cytokine production, which was inferred to ultimately extend the resolution time of the inflammatory and wound healing responses. For a composite consisting of a polymeric phase, it is possible to expect localized confinement of chain mobility in specific areas at the surface, which could lead to reduced overall mobility compared with the pure polymer. Based on the Andersson premise, such outcome would impact positively the inflammatory and wound healing response as it relates strictly to the polymeric component. The surface dynamic of chains in a polymer composite are strongly affected by the local environment. Local density has been recognized to play an important role in polymer chain mobility. Specifically, the lower the density of the chain surroundings, the higher is the chain mobility. Especially in a biodegradable polymer composite where fluctuations of the local density are expected over time, distribution of chain mobility and its associated effect on the surrounding biological environment is difficult to assess.
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The inherent phase separation characteristic to biocomposites leads to formation of discrete aggregates not just in the bulk, but also at the surface. The random variation in size and distribution of these aggregates at the surface renders it heterogeneic. Such phenomenon leads to nonuniform surface biocompatibility as the size and distribution of the aggregates lead to subsequent heterogeneity in protein adsorption and cellular adhesion.
14.3.3 Surface topography The concept of spatially controlled surfaces to modulate the adhesion, migration, and proliferation of cells onto material surfaces has been employed for many decades. Significant strides have been achieved in associating certain topographical features to specific cellular behavior. In vitro studies have been performed to investigate the responses of a wide number of cells to various topographical cues (Bettinger et al. 2006; Curtis et al. 2005; Recknor et al. 2006; Sarkar et al. 2006; Vernon et al. 2005). Those reports have identified topographical features on the order of 1–100 μm being instrumental in modulating the adhesion, proliferation, and function of the investigated cell lines (Kenar et al. 2006; Manwaring et al. 2004; Miller et al. 2001; Sutherland et al. 2005; Zhu et al. 2005). Recent advances in nanofabrication have been adopted by the biomaterial development sector to further enhance the modulation capabilities associated with topographical cues. Integration of nanoscale topographic features onto the material surface has addressed the previously unmet biological need of the cellular nanoscale features (filopodia and cytoskeleton) as well as extracellular matrix features at the nanoscale level (pores, fibers, and ridges). Through those efforts, it was recognized that cells are capable of detecting 50–100 nm topographical cues, and features as small as 10 nm were reported to have an effect on the cellular behavior (Liliensiek et al. 2006; Yim and Leong 2005). Thus, micro- and nanotopography could be used to modulate cellular behavior and, through it, biocompatibility.
14.4
Evaluation of biocompatibility: the relevance of employed analyses
Technological advances that have led to improved implantation performance have been supported by established and adequate strategies for testing the biocompatibility of these medical devices. Evaluation of the biological response to a biocomposite should follow established procedures allowing for adequate determination of its biocompatibility.
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14.4.1 Biological assessment The assessment of biocompatibility of a particular biocomposite entails the consecutive execution of well accepted scientific evaluations. Not discretely, but considered together, the results of these evaluations help provide an objective picture of the associated composite biocompatibility (Anderson 2001). In general, the sequence of evaluations increases in complexity and scope over time, from basic in vitro single cell culture cytotoxicity evaluation to in vivo large animal anatomically relevant evaluations of biocompatibility and biofunctionality. The task of evaluating the biocompatibility of biocomposites is more complex owing to the presence of more than one discrete phase, nonhomogeneous distribution of phases in the bulk compared with the surface of the biocomposite, and non-uniform surface chemistry, morphology, and topography. It is accepted that biocompatibility is defined by the exposed surface of the biocomposite to the local biological environment. In the case of biocomposites, there is potential for phase separation, mechanical fracture/cracking, and biodegradation followed by the release of leachables. The in vitro evaluation provides the initial biocompatibility screening. There are a number of advantages for which the initial use of in vitro systems in screening biocompatibility of biocomposites is desirable, such as comparative cost-effectiveness and evaluation time, as well as the potential for evaluation of a large number of prototypes (different materials, alternative architectures, chemistries). However, the critical aspect satisfied by the in vitro evaluation of biocompatibility is the potential for standardization, and thus, reproducibility of testing conditions (Anderson 2001; Cavaillon 1994). The in vitro evaluations are represented by target cell assays, making use of the most sensitive cells pertinent to the application for which the biocomposite is developed. Another important aspect of the in vitro screening evaluation of biocompatibility of a biocomposite is that these assays minimize the variables of metabolism, distribution, and adsorption, and maximize the cell line exposure to any potential toxicity. There are associated disadvantages to the in vitro biocompatibility evaluation, in particular, the extrapolation of these results to the clinical population. Consequently, any in vitro biocompatibility screening must be followed by an in vivo evaluation of biocompatibility and biofunctionality. The staggering complexity of the in vivo systems and, in some cases, the limited choice of cell lines should not detract from the initial in vitro screening evaluation, but be an incentive to such an approach that could provide critical cytotoxicity information in controlled environments. The use of cell lines in the ideal situation of in vitro evaluation of biocompatibility would rely on human non-transformed cells, specifically, primary isolated cells in early passage (Balkwill 2002; Beutler 1999). In most
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of the initial stages of biocompatibility screening, immortalized cell lines are used, with origins within transformed primary cultures. While these cells have a limited life, a continuous immortal cell line may be derived from different species and tissues (Hirano and Kishimoto 1989). It has been acknowledged that, in comparison with the primary cultures of the same type of cell, a cell line presents morphological alterations such as decrease in cell size, reduced adherence, and enlarged nucleus (Cohen and Cohen 1996; Van Snick 1990). Phenotype and genotype variations compared with the primary cell represent a real possibility associated with immortal lines (Cohen and Cohen 1996). Nevertheless, in vitro evaluation of biocompatibility using cell lines provides the necessary first-stage assessment of biocompatibility onto which more relevant and complex assessments will rest. These principles of biocompatibility screening have also been employed in the case of composite materials. It is critical to mention that an overwhelming majority of the composites developed are indicated for musculoskeletal application. Schneider et al. (2008) used cloned human mesenchymal stem cells (MSC) to investigate the toxicity of amorphous tri-calcium phosphate nanoparticle loading into the poly(lactide-co-glycolide) scaffold (Table 14.4). Differentiation of MSC onto the surface of the composite, along with the production of osteocalcin, an indicator of osteogenic differentiation, qualified the scaffolds as non-toxic and able to support pertinent cell attachment and differentiation (Table 14.4) (Schneider et al. 2008). Others have found the use of animal-derived cells for preliminary biocompatibility screening of composites more beneficial. In general, this approach hinges on the premise that an in vivo animal model had already been identified or established. Thus, the use of same animal cells/cell line in this preliminary step allowed for more reliable extrapolation of the data from in vitro screening to the in vivo animal model. Among others, Hong et al. (2007) used a primary rabbit cell line to evaluate the performance of a poly(lactideco-glycolide) modified with carbonated hydroxyapatite nanoparticles composite intended for osteogenic remodeling and regeneration (Table 14.4). The rabbit model is an established animal model in evaluation of material performance for musculoskeletal applications (Losken et al. 2008; Sumitomo et al. 2008; Wellisz et al. 2008). Thus, in this context, the use of the rabbit derived cell line is pertinent. Gupta et al. (2007) followed the same approach by using primary neonatal rat calvarial cells for the evaluation of another composite intended for bone remodeling/regeneration (Table 14.4). It is evident that the majority of in vitro composite evaluations were performed using one cell line, predominantly, differentiated or nondifferentiated human osteoblasts (Table 14.4) (Huang et al. 2007, 2008; Kim et al. 2005; Negroiu et al. 2008; Schneider et al. 2008). The specific assays included morphological, biochemical, and genetic evaluation. It is known that morphologic changes of cells interacting with various surfaces are an
• PEU: medical polymer of known biocompatibility • Au or Ag: enhance biostability of bulk through free radical scavenging • PLGA 80:20: biomedical and biodegradable copolymer • CHAP: enhances inherent biocompatibility of PLGA in orthopedic applications • PEU: medical polymer of known biocompatibility • Au: enhances biostability of matrix through free radical scavenging • PLGA 85:15: biomedical polymer used clinically • ATCP: enhances inherent biocompatibility of PLGA in orthopedic applications
• PET: medical polymer • HA: biological enhancement and mechanical reinforcement
Poly(ether)urethane and gold or silver nanoparticles (PEU–Au or PEU–Ag)
Polyethylene terephthalate and hydroxyapatite (PET–HA)
Poly(lactide-co-glycolide) 85:15 and amorphous tricalcium phosphate (PLGA–ATCP)
Poly(ether)urethane and gold nanoparticles (PEU–Au)
Poly(lactide-co-glycolide) and carbonate hydroxyapatite (PLGA–CHAP)
Role of each component
Composite1
Table 14.4 Biocomposites: biocompatibility
• In vitro • Cytocompatibility assay using human dermal fibroblasts • Monocyte activation assay • Bacterial adhesion • In vitro, human bone marrow cells • Examine cellular morphology by confocal laser scanning microscopy • Cellular differentiation and bioactivity evaluation by alkaline phosphatase activity, osteocalcin, and DNA content • In vitro • Composite extract: cytotoxicity by MTT assay using murine fibroblast culture and TNF-α release (ELISA) using monocyte culture • Direct contact: cell morphology and proliferation by SEM, Alamar Blue, and TNF-α release (ELISA)
• In vitro, rabbit osteoblast line • Contrast microscopy using Giemsa stain to assess attachment and morphology
• In vivo, swine model, subcutaneous • Composite analyses (see Table 14.1) • Histology (H&E)
Biocompatibility evaluation
(Dimitrievska et al. 2008)
(Schneider et al. 2008)
(Hsu et al. 2006)
(Hong et al. 2007)
(Chou et al. 2008)
Reference
• S: ceramic used in orthopedic applications • CPC: mechanical reinforcement material • CDG: high degree of biological functional groups and is being used in tissue scaffolds • HA: biological enhancement and mechanical reinforcement
• CS: biological function • PGA: mechanical reinforcement • PA/(bFGF: biological function
• PHEMA: biocompatible material • nanoCHA: biological enhancement and mechanical reinforcement
Silica-calcium phosphate nanocomposite (S–CPC)
Collagen and polyglycolic acid and peptide amphiphile with fibroblast growth factor (CS–PGA–PA/(bFGF)
Poly(2-hydroxy ethyl methacrylate) and carbonate substituted hydroxyapatite (PHEMA–nanoCHA)
Collagen gelatin and hydroxyapatite (CDG–HA)
• BCP: effective in bone repair and regeneration • HA: biological enhancement and mechanical reinforcement
Biphasic calcium phosphate nanocomposite (BCP–HA)
Table 14.4 Continued
• In vitro, human osteoblasts • Cell proliferation using an MTS assay • Cell differentiation through measuring the alkaline phosphatase activity and osteocalcin expressed by the cells • Protein adsorbtion using an assay kit (fetal bovine serum) • In vitro and in vivo • In vitro: (bFGF release assay using fluorescence spectrophotometry • In vivo ectopic bone formation: (bFGF release assay, x-ray absorptiometry, and bone mineral density • In vitro, acellular and cellular • Acellular: change in surface morphology by ESEM through sample immersion in simulated body fluid • Cellular using primary human monocytes: lactate dehydrogenase release, TNF-α • Cellular using human osteoblasts: growth and proliferation (Alamar Blue, immunofluorescence, and SEM)
• In vivo osteogenic differentiation in nude mice, subcutaneously; and in an ectopic defect in rats • Histology (H&E) for mice samples • Immunohistochemistry for rat samples • Expression of osteogenic specific genes using RT–PCR • In vitro, neonatal rat calvarial osteoblasts • Expression of osteogenic specific genes using RT–PCR
(Huang et al. 2008)
(Hosseinkhani et al. 2007)
(Kim et al. 2005)
(Gupta et al. 2007)
(Lin et al. 2007)
• PHEMA/PCL: biocompatible material • NanoHA: biological enhancement and mechanical reinforcement • PPF: biocompatible material • UStube: mechanical reinforcement • MP: controls the crystallization process of HA • HA: biological enhancement and mechanical reinforcement
Poly(2-hydroxy ethyl methacrylate)/ polycaprolactone and hydroxyapatite (PHEMA/ PCL–nanoHA) Poly (progylene fumarate) and ultra-short carbon nanotubes (PPF–UStube) Sodium maleate–vinyl acetate and hydroxyapatite (MP–HA)
• In vitro acellular using exposure of the composite to simulated body fluid and examining by SEM and IR • In vivo rabbit muscle implant and examination by SEM, IR, XRD, histology • In vitro acellular and cellular • Acellular degradation in a PBS solution evaluated by weight loss of composite • Cellular using stromal stem cells: cellular morphology (SEM) and cell proliferation (MTT assay) • In vitro acellular and cellular • Acellular bioactibity in water using ESEM • Biological response to human osteoblasts-like cells through actin filament analysis (confocal microscopy) • In vitro, rat mesenchimal stem cells • Scaffold cellularity assessed with PicoGreen assay kit (confocal microscopy) • In vitro, human embryonic kidney cells • Cellular viability was assessed using an MTS assay • Cellular proliferation was assessed by contrast microscopy using Trypan Blue • Morphology was determined using actin filament staining (confocal microscopy)
Note: the first name describes the matrix phase, while the second describes the reinforcing material.
• Gel: antigenic properties and bioabsorbability • MMT/CS: biological enhancement and mechanical reinforcement
Gelatin and montmorillonite– chitosan (Gel–MMT/CS)
1
• PA: biocompatible material • NA: biological enhancement and mechanical reinforcement
Polyamide and nanoapatite (PA–NA)
(Negroiu et al. 2008)
(Shi et al. 2007)
(Huang et al. 2007)
(Zhuang et al. 2007)
(Jie et al. 2007)
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indication of surface biocompatibility. Information is usually obtained by subjective visual inspection. Alterations such as shrinking of the cell nucleus, fragmentation of the cytoplasm, rounding off, and cell detachment have been used as indicators of surface-induced cell morphology alteration (Akira et al. 1993; Borish and Steinke 2003). To aid in such assessment, staining assays have been used to distinguish between alive and dead cells leading to expanding the visual analysis to cellular viability. An overwhelming number of reports on the biocompatibility of composites have used such evaluation techniques (Table 14.4) (Dimitrievska et al. 2008; Hong et al. 2007; Hsu et al. 2006; Huang et al. 2007, 2008; Jie et al. 2007; Negroiu et al. 2008; Schneider et al. 2008; Shi et al. 2007). Another approach to in vivo biocompatibility evaluation involves the use of biochemical assays that investigate cell membrane integrity and appear to be considered the most useful methods of quantitative assessment of cytotoxicity. In many cases, a combination of cellular morphology and biochemical evaluation provides a clearer picture of the overall cytotoxicity of a composite surface (Table 14.4) (Huang et al. 2008; Kim et al. 2005; Negroiu et al. 2008; Shi et al. 2007; Zhuang et al. 2007). Finally, genotoxicity has been used recently in the evaluation of biocomposites. The genotoxicity assays evaluate whether a chemical has the potential to cause somatic or germ cell effects in animals (Moser and Loetscher 2001; Yoshie et al. 2001). This approach has been expanded to evaluation of genetic expression within study cells (Table 14.4) (Dimitrievska et al. 2008; Gupta et al. 2007; Lin et al. 2007; Schneider et al. 2008). The challenge in evaluating the biocompatibility of a composite is secondarily related to the choice of an appropriate cell line, yet, this is primarily affected by the distribution of two or more composite phases within the overall structure and, in particular, their distribution at the exposed surface to the biological environment. The issue of size and distribution of different phases at the surface raises additional questions regarding biocompatibility evaluation. Is the reporting of homogeneous cellular distribution and morphology an indirect indication of the homogeneity of the exposed surface? Should parallel controls representing each distinct phase, as an entity, be included in the evaluation of the composite to enhance the understanding of its surface interaction with the biological environment? Is the overall cytotoxicity effect sufficient in the evaluation of biocompatibility of a biocomposite? Following successful cytotoxicity screening, it is time for the biocomposites to be evaluated in vivo. In vivo biocomposite studies are extremely sparse. Few researchers have reached the evaluation point in in vivo animal models. The process of in vivo evaluation requires attention to a number of critical details that ensure the success of the study. Implantation procedures have been developed to obtain the most adequate systems for appropriate
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evaluation of biocompatibility. It is critical to recognize that the inflammatory and wound healing response in vivo is influenced not only by the presence of the biocomposite, but also by the surgical technique, as issues related to non-aseptic techniques or accidental tissue damage could affect the overall outcome. Such a system does not allow for the highly quantitative analyses employed by the biochemical and genetic in vitro assays, however, it provides information on the composite biocompatibility in the complex environment of the body, which is generally assessed through immunohistochemical and histological analyses (Table 14.4) (Chou et al. 2008; Hosseinkhani et al. 2007; Lin et al. 2007). A concern related to the evaluation of biocomposites in vivo is similar to that outlined in the case of their in vitro evaluation: specifically, whether the biocomposite surface (chemistry, morphology, topography) is assumed to be homogeneous and, furthermore, whether the structural evaluation of the bulk can be assumed to be representative of the surface.
14.5
Surface characterization
The traditional methods employed in analyzing the bulk structure of biocomposites, in most cases, are not suitable for surface evaluation. Also, it is important to recognize that methods particularly developed to analyze the surface have the potential of altering the surface (Ratner 2004). Therefore, it is advisable to employ more than one method of surface characterization to ensure the reliability of the information extracted. Artifacts of surface analysis tend to affect hard materials such as metals, ceramics, glasses, and carbons less than organic and polymeric materials. Chain mobility at the surface of the biocomposite consisting of a polymeric phase could also introduce an artifact of surface characterization, as the properties of the polymeric phase will be affected by the chain mobility.
14.5.1 Contact angle The surface energy of any material has been demonstrated to be directly related to wettability, which, in turn, has been found to be instrumental in modulating biological responses. Specifically, through contact angle measurements, a surface is rendered either hydrophilic or hydrophobic. In general, hydrophilic surfaces lead to less cellular adhesion of inflammatory cells than hydrophobic surfaces (Brodbeck et al. 2003). The contact angle measurement is highly operator dependent and affected by surface contamination, surface roughness heterogeneity, and surface rearrangements (chain mobility). These limitations especially come into focus in the case of biocomposites, where, owing to surface non-homogeneity, a contact angle measurement could differ from region to region.
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Hsu et al. (2006) reported on the difference in contact angle between poly(ether)urethane and the poly(ether)urethane–gold nanocomposite (Tables 14.5 and 14.6). No consistent correlation between increasing the gold content within the nanocomposite and contact angle values emerged. Atomic force microscopy (AFM) analysis of the nanocomposite surface clearly showed the presence of gold micelles of 1–4 nm in diameter and quasi-homogenously distributed (Tables 14.5 and 14.6) (Hsu et al. 2006). The question that then arises is exactly what is the purpose of contact angle measurements in this study.
14.5.2 Fourier transform infrared spectroscopy – attenuated total reflectance Fourier transform infrared spectroscopy – attenuated total reflectance (FTIR–ATR) provides information related to the presence or absence of specific functional groups, as well as the chemical structure of polymer materials. Shifts in the frequency of absorption bands and changes in relative band intensities indicate changes in the chemical structure or changes in the environment around the sample. Thus, FTIR–ATR spectroscopy can be used to determine the resultant surface chemistry especially following induced chemical or physical modifications. Gupta et al. (2007) used a combination of SEM and FTIR to evaluate the morphology and surface chemistry of silica–calcium phosphate nanocomposites (Table 14.5 and Table 14.6). While one SEM photomicrograph was included, FTIR results were not reported. Considering the nature of the composite, relevant information on the surface chemistry could have been collected to provide some understanding of the component composition at the surface and possibly in the bulk, as these scaffolds were highly porous. Kim et al. (2005), however, used and reported on FTIR of gelatin– hydroxyapatite nanocomposites (Table 14.5 and Table 14.6). The nanocomposite was compared with the two pure components, gelatin and hydroxyapatite, and clear differences in FTIR spectra were identified, intermediate between those of the pure components. The FTIR helped determine the chemical nature of the nanocomposite at its surface. A similar approach was employed by Huang et al. (2007) in their evaluation of poly-2-hydroxyethylmethacrylate / polycaprolactone in combination with hydroxyapatite nanoparticles (Table 14.5 and Table 14.6). While both these studies reported the presence of characteristic groups of interest for their respective biocomposites, such information is qualitative and does not provide a measure of the percentage representation at the surface of the biocomposite of each component and/or surface distribution of each phase.
• Cylindrical, 30–40 nm in diameter and 100–200 nm long • Variable loading, 2–15 wt%
Poly(lactide-co-glycolide) 80:20 and carbonate hydroxyapatite (PLGA – CHAP) Poly(ether)urethane and gold nanoparticles (PEU–Au)
• <100 nm particle size and specific surface area of ~68 m2 g−1 • Loading of matrix: 30 wt% HA
Collagen gelatin and hydroxyapatite (CDG–HA)
Silica-calcium phosphate nanocomposite (S–CPC)
• Shape and size information not provided • Loading of matrix: 0–10 wt% HA • Nanofibers, no other data provided • Loading of matrix: 70/30 HA/BCP by weight • 50 nm crystals • Loading of the matrix: unclear
Polyethylene terephthalate and hydroxyapatite (PET–HA) Biphasic calcium phosphate nanocomposite (BCP–HA)
Poly(lactide-co-glycolide) 85:15 and amorphous tricalcium phosphate (PLGA–ATCP)
• Spherical • 5-nm particles with 43.5 ppm Au or 30.2 ppm Ag loading
Poly(ether)urethane and gold or silver nanoparticles (PEU–Au Or PEU–Ag)
• Spherical • 5-nm particles with 17.4–174 ppm Au loading • Spherical • 25–50 nm with a 40% weight loading (PLGA–ATCP 60:40)
Size of reinforcing material
Composite1
Table 14.5 Biocomposites: the materials
• Matrix: optical microscopy • Reinforcing material: TEM • Composite: optical microscopy, SEM, FTIR-ATR • Matrix: SEM • Reinforcing material: SEM, x-ray (EDS), FTIR–ATR • Composite: SEM • Matrix: AFM, S-CA • Reinforcing material: TEM • Composite: AFM, S-CA • Matrix: SEM • Reinforcing material: nitrogen adsorption, FTIR-ATR, TEM, SEM • Composite: SEM, XRD • Matrix: XRD, optical microscopy • Reinforcing material: none • Composite: SEM, optical microscopy • Matrix: none, deferred to a previous publication • Reinforcing material: none • Composite: none • Matrix: none • Reinforcing material: none • Composite: TEM, SEM, FTIR-ATR, x-ray CT • Matrix: none • Reinforcing material: none, based in information provided by the manufacturer • Composite: XRD, FTIR-ATR, TEM
Analyses
(Kim et al. 2005)
(Gupta et al. 2007)
(Lin et al. 2007)
(Dimitrievska et al. 2008)
(Schneider et al. 2008)
(Hsu et al. 2006)
(Hong et al. 2007)
(Chou et al. 2008)
Reference
1
• Size and shape: not provided • Loading of the matrix: 0.5 or 0.83 wt% • Size and shape were not provided • Loading of the matrix: unclear
• PGA: 20 μm in length and 0.5 μm in thickness fibers • Loading of matrix: 1.5, 3. 6 or 12 mg per 0.75 mL CS (3 mg mL−1) • PA: not disclosed • Loading of matrix: unclear • 20–30 nm in width and 50–80 nm in length • Loading of the matrix: 10, 20 or 40 wt% • Size and shape of the NA particles not provided • Loading of matrix: 65 wt% • Intercalated layer of 4.82 nm in thickness • Loading of the matrix: unclear, however, it is mentioned that MMT/CS is added to 6 wt% gelatin • 20–30 nm in width and 50–80 nm in length fibers • Loading of the matrix: 20 or 40 wt% • • • • • •
Matrix: SEM Reinforcing material: SEM Composite: SEM, microCT, Matrix: none Reinforcing material: none Composite: none
• Matrix: XRD, SEM • Reinforcing material: TEM, XRD • Composite: FTIR-ATR, SEM
(Negroiu et al. 2008)
(Shi et al. 2007)
(Huang et al. 2007)
(Zhuang et al. 2007)
(Jie et al. 2007)
• • • • • •
Matrix: none Reinforcing material: none Composite: none Matrix: XRD Reinforcing material: XRD Composite: XRD, SEM
(Huang et al. 2008)
(Hosseinkhani et al. 2007)
• Matrix: none • Reinforcing material: TEM, XRD, FTIR-ATR • Composite: FTIR-ATR
• Matrix: SEM • Reinforcing material: SEM, SIMS (PA only) • Composite: SEM
Note: the first name describes the matrix phase, while the second describes the reinforcing material.
Poly(2-hydroxy ethyl methacrylate)/ polycaprolactone and hydroxyapatite (PHEMA/ PCL–noHA) Poly(propylene fumarate) and ultra-short carbon nanotubes (PF–UStube) Sodium maleate–vinyl acetate and hydroxyapatite (MP–HA)
Gelatin and montmorillonite– chitosan (Gel–MMT/CS)
Poly(2-hydroxy ethyl methacrylate) and carbonate substituted hydroxyapatite (PHEMA–nanoCHA) Polyamide and nanoapatite (PA–NA)
Collagen and polyglycolic acid and peptide amphiphile with fibroblast growth factor (CS–PGA–PA/bFGF)
Table 14.5 Continued
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Table 14.6 Methods of surface analysis Depth analyzed
Spatial resolution
Analytical sensitivity
Contact angle
320 Å
1 mm
FTIR-ATR
1–5 μm
10 μm
Low or high depending on chemistry 1 mol%
SEM
5Å
40 Å
High but not quantitative
AFM
Å
SIMS
10 Å–1 μm
0.1 nm or better 100 Å
High but not quantitative Very high
ESCA (XPS) Auger ES STM SIMS
10–250 Å 50–100 Å 5Å 10 Å–1 μm
10–150 μm 100 Å 1Å 100 Å
0.1 at% 0.1 atom% Single atoms Very high
Method
Reference (Hsu et al. 2006)
(Kim et al. 2005; Gupta et al. 2007; Huang et al. 2007) (Gupta et al. 2007; Hong et al. 2007; Hosseinkhani et al. 2007; Huang, Lin et al. 2007; Shi et al. 2007; Zhuang et al. 2007; Chou et al. 2008; Dimitrievska et al. 2008; Schneider et al. 2008) (Hsu et al. 2006) (Hosseinkhani et al. 2007) – – – –
14.5.3 Scanning electron microscopy Scanning electron microscopy (SEM) uses a beam of electrons to scan the surface of a sample and create detailed 2D and 3D images of a specimen at nanometer resolution. It is usually used as a corroborative method to investigate surface properties. Most of the studies reporting on biocomposites have used SEM to investigate surface topography (Table 14.5 and Table 14.6) (Chou et al. 2008; Dimitrievska et al. 2008; Gupta et al. 2007; Hong et al. 2007; Hosseinkhani et al. 2007; Huang et al. 2007; Schneider et al. 2008; Shi et al. 2007; Zhuang et al. 2007). SEM has been found useful by some investigators in evaluating the surface degradation of some biocomposites, in particular, the presence of pitting or cracking (Table 14.5 and Table 14.6) (Chou et al. 2008; Zhuang et al. 2007). Others have employed this imaging technique in investigating basic surface topography (Table 14.5 and Table 14.6) (Gupta et al. 2007; Hong et al. 2007; Huang et al. 2007; Schneider et al. 2008; Shi et al. 2007). This technique, while providing topographical feedback, does not have the capability of conferring quantitative information on the chemistry or morphology at the surface of the biocomposite.
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Furthermore, unless one knows how each phase looks by imaging appropriate controls side by side with the composite sample, SEM does not provide the information required to make a clear distinction between surface represented phases. Yet others have employed SEM to investigate cellular adhesion onto various biocomposites (Table 14.5 and Table 14.6) (Dimitrievska et al. 2008; Hosseinkhani et al. 2007). While some qualitative information can be extracted, quantitative analysis is outside the capabilities of this technique. Again, it must be employed as a corroborative method in supporting quantitative techniques such as cell attachment/proliferation analyses.
14.5.4 Atomic force microscopy Atomic force microscopy (AFM) provides 3D information on the surface features of a sample. Its capabilities span from examining any rigid surface to examining the surface of specimens immersed in liquid. It is more sensitive than scanning electron microscopy (SEM) in being able to resolve minor differences between relatively smooth surfaces, at the level of a single atom. A profile of ‘hills’ and ‘valleys’ is created upon scanning of the surface. Hsu et al. (2006) used AFM in combination with contact angle to characterize the surface of poly(ether)urethane–gold nanocomposites (Table 14.5 and Table 14.6). Formation of gold micelles within the poly(ether)urethane phase, at the surface of the composite, was evident. AFM is significantly more useful as an investigative method when the biocomposite phases are clearly distinct visually, such as in the composite discussed above. However, if topographical distinction is lacking, it is difficult to use AFM in evaluating the topography of each phase, and with it the overall presence of each component at the surface of the biocomposite.
14.5.5 X-ray techniques Energy-dispersive x-ray (EDX) represents a SEM feature that allows for elemental chemical identification at the surface of a specimen. Hong et al. (2007) implemented such a method to identify the component surface elements of the biocomposite under investigation (Table 14.5). However, the method must be used as a corroborative approach to evaluating the surface chemistry, for example, in combination with FTIR. X-ray diffraction (XRD) patterns include information about the crystal structure as well as crystal size, orientation disorders, and lattice disorders. Its use is limited to composites or phases of composites that can form crystalline structures. Schneider et al. (2008) used this method to ensure the amorphous structure of the crystalline component, tricalcium phosphate, before
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preparation of the biocomposite (Table 14.5). Kim et al. (2005) appeared to use XRD for a similar purpose as Schneider et al., but no reporting on the XRD analysis was included as part of the study (Table 14.5). Zhuang et al. (2007) implemented the same technique to measure the change in crystalline segment distance before and after intercalation of the montmorillonite crystals with the non-crystalline phase of chitosan (Table 14.5). In x-ray computed tomography (x-ray CT), digital geometry processing is used to generate a three-dimensional image of the inside of an object from a large series of two-dimensional x-ray images taken around a single axis of rotation. Gupta et al. (2007) employed this method to evaluate the porosity of the created silica–calcium phosphate nanocomposites (Table 14.5). This method can provide information on the topography of the surface compared with that of the bulk, but does not provide any information on the phase distribution of the two phases forming the biocomposite.
14.5.6 Transmission electron microscopy Transmission electron microscopy (TEM) technique is applicable only to thin films resulting in a 3D image of the sample. Gupta et al. (2007) and Kim et al. (2005) used this method to determine the crystal size within their respective biocomposites (Table 14.5). However, as this method is not considered to be a surface characterization technique, it must be queried whether the authors have assumed that the measured crystal size was similar on the surface and in the bulk.
14.5.7 Secondary-ion mass spectroscopy Secondary-ion mass spectroscopy (SIMS) produces a mass spectrum of the outermost 10 Å of the surface through ion bombardment. Hosseinkhani et al. (2007) employed this method for the surface characterization of a peptide–amphiphile (PA) as a component of a collagen-based bone regenerating scaffold (Table 14.5 and Table 14.6). Unfortunately, the method was only used to evaluate the PA and none of the composite samples fabricated.
14.5.8 Other surface characterization methods Surface characterization methods such as electron spectroscopy for chemical analysis (ESCA), Auger electron spectroscopy (Auger ES), and scanning tunneling microscopy (STM) represent established methods of characterizing surface chemistry, morphology, and topography. However, none of these methods were actually used in any of the studies of biocomposites reviewed for this work (Table 14.5 and Table 14.6).
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14.6
Future trends
As the field of biocomposites continues to expand by integration with new technologies, the challenges for appropriate and adequate biocompatibility testing of biocomposites also will be increased. Significant among the many challenges will be the correlative elucidation of symbiotic relationships between composite surface structure/chemistry and biological response evaluation. In particular, design parameters of biocomposites that require activity and interaction between the surface structure/chemistry and protein, cell, tissue, and organ interactions must be quantitatively evaluated in the context of the proposed design criteria that form the basis of the composition of the biocomposite. A truly interdisciplinary and integrated approach to biological response evaluation will be required to verify stated claims incorporated in the design parameters. At present, these correlations are lacking in current biocompatibility evaluation of biocomposites. The following examples illustrate the complexity that will be faced in determining biocompatibility of multifunctional biocomposites. Composite systems consisting of multiple polymer phases with distinct release kinetics have been developed for application in tissue engineering and cancer chemotherapy. Hydrophilic and hydrophobic porous biodegradable polymers containing the therapeutic agent are used to provide the continuous phase of the composite whereas microcapsules containing a different therapeutic agent have been embedded within the continuous phase. These systems can provide physicochemical, biological, and spatial control of the therapeutic agents depending on what materials are used for the continuous (matrix) and discontinuous (micro/nanosphere) phases (Sands and Mooney 2007). In general, the scaffold matrix phase can provide for early and rapid release of the therapeutic agents whereas the embedded micro/nanospheres provide for a delayed interaction with the biological environment. Modification of the surfaces of the porous matrix phase has been developed where integrin adhesion sequences, e.g., RGD, have been utilized to facilitate cellular migration and adhesion within the porous structure of the tissueengineered scaffold (Ennett et al. 2006; Fischbach and Mooney 2007). Drugloaded composites also have been suggested for anticancer therapies where low doses of multiple antitumor drugs could significantly enhance the therapeutic efficacy when compared with the delivery of single medications. Mooney and colleagues have exploited these concepts in developing various systems for the development and control of angiogenesis through modulation of spatial and temporal gradients. Their efforts clearly have demonstrated that these systems can initiate angiogenesis, induce maturation of the vascularity, and maintain the integrity of the vascularity (Sands and Mooney 2007). In the area of orthopaedics, Mikos and colleagues have used gelatin microspheres containing a therapeutic factor that was then
Biocompatibility evaluation of biocomposites
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incorporated into oligo(poly(ethylene glycol)fumarate) hydrogels containing a second factor for cartilage tissue engineering. Biocomposites also have used plasmid DNA to provide systems where the bioactive agent must be provided in a more delayed and sustained fashion and provide for longterm availability of therapeutic proteins via cellular synthesis (Christenson et al. 2007; Holland et al. 2005). Whereas in vitro and in vivo studies have clearly demonstrated the feasibility of these systems to function appropriately, their complex architecture as well as the time-dependent variation provided by biodegradation, presents real and complex challenges in the evaluation of the biocompatibility of these systems. Not only must the dimensional and geometric factors/ variables be considered in biocompatibility evaluation, but the resultant effect of the released therapeutic agents also must be addressed in a biocompatibility evaluation program. With such systems, the unique features of each respective drug-releasing biodegradable biocomposite ultimately will dictate approaches and methods to adequately and appropriately evaluate biocompatibility. With these systems, it is obvious that not only in vivo biological response evaluation tests will play a significant role in the determination of the biocompatibility, but also that the unique nature of the drug-delivery biocomposite under study will require new and novel approaches to biological response evaluation, i.e., biocompatibility. Clinical application of biocomposites requires both broad and in-depth evaluation of the biocompatibility of each biocomposite system. This, in turn, forms the basis for regulatory approval of the intended biocomposite system in its designated application. Demonstration of adequate and appropriate function, i.e., efficacy, will be necessary before these systems can be used to improve the health and welfare of patients.
14.7
References
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15 Cellular response to biocomposites P. JAYA K U M A R and L. D I S I LV I O, King’s College London Dental Institute, UK
Abstract: The modern medical challenge of tissue repair, replacement and regeneration in degenerative and pathological processes may be met by the use of biocomposite materials. Biocomposites are intrinsically multifunctional and may contend with the demands of reconstructing natural tissues beyond conventional biomaterial strategies. An understanding of the cellular responses to the physicochemical and mechanical structure of these biomaterials is essential in achieving appropriate functional characteristics in tissue environments. We discuss definitions of biocomposites, cellular interactions, their role in skeletal regeneration and reconstruction, and the state of the art in in vitro experimentation in this field. Key words: cell interactions, biocomposites, nanocomposites, biomimetics, bone, tissue engineering, osteoinductive agents.
15.1
Introduction
The repair, replacement and regeneration of tissues involved in injury, degeneration and pathology using biomaterials remains a modern medical challenge. There continues to be a high demand for degradable and nondegradable materials with appropriate structural and functional properties enabling restoration of damaged tissues. In particular, skeletal damage leading to bone loss is a major healthcare concern with massive socioeconomic implications. The advances in surgical practice combined with a rise in active and ageing populations has led to an increase in trauma and elective orthopaedic surgery. In 2006, approximately 542 000 total knee replacements and 231 000 total hip replacements were performed in the US.1 In 2008, the National Joint Registry recorded over 144 000 primary and revision hip and knee procedures in England and Wales.2 Combined with other musculoskeletal elective and trauma treatments, the costs are immense. Fracture repair and implant technology continues to improve, but, there are a significant proportion of patients with complications such as fracture non-union and bony defects not amenable to healing by direct fixation alone. Consequently, there is a greater demand for bone repair and regeneration systems beyond current conventional strategies. Autografts, allografts, xenografts and implants composed of metal, plastic and ceramic materials remain the traditional implant options. They are all associated 354
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with various limitations. Autografts may cause donor site morbidity and are limited in terms of donor stock. Allografts and xenografts may carry the risk of disease transmission and adverse immunochemical reactions. Metals, polymers and ceramics have been extensively used in implantology and continue to contend with the complications of wear processes, loosening, stress shielding and mechanical failure to varying extents. Limitations in materials lead to limitations in longevity of implants which progress from structural failure to clinical failure and revision surgery, increasing cost and morbidity. Synthetic bone graft materials in the context of skeletal tissue engineering are the potential solution. Composite materials have been developed over several decades in the aerospace, construction and automotive industry. There has been resurgence in their design and application in health and non-healthcare areas. The drive has been influenced by increased environmental and healthcare concerns, alongside a trend for lower energy consumption and more sustainable manufacturing methods producing lighter weight structures.3 Biocomposites are specifically composed of two or more constituents, combined to produce a unique material with complex physical–chemical, biological, and mechanical properties that are applicable to biomedical applications. There has been widespread use in the field of trauma and orthopaedic reconstructive surgery, as joint replacement and fracture fixation devices. It is the development of biocomposites as synthetic bone graft substitutes in bone tissue engineering and drug delivery systems that presents an exciting future for skeletal reconstruction. The multifunctionality of biocomposites is a general characteristic of most living systems and tissues such as skin, bone, cartilage, dentine and collagen. This feature may supersede the often single-phased homogeneous and isotropic nature of materials such as polymers, metals and ceramics in isolation. This chapter provides an outline of the basic cellular requirements of skeletal reconstruction and regeneration, with bone structure and function being the biological blue print and target for replication. The traditional options and the state of the art in biocomposite research, development and manufacture in bone tissue engineering are discussed. We emphasise polymer/ceramic composite design and fabrication and examine the variety of tailor-made biocomposites and nanocomposites used as bone tissue engineering scaffolds.
15.2
Skeletal regeneration and reconstruction
15.2.1 Definitions Skeletal regeneration is essentially achieved by the interaction of three elements: cells, growth factors (GFs) and a scaffold material. Skeletal
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reconstruction is successfully achieved by an implant possessing the appropriate biological and mechanical characteristics for the target environment. A biomaterial is a material used in medical devices or implants that interacts with biological systems.4 If a biomaterial is being used to fulfil the role of the scaffold material in order to achieve these objectives, it must by definition be biocompatible and bioactive. Biocompatibility is a fundamental requirement of biomaterials and is defined as the ability of a material to perform with an appropriate host response in a specific application.5,6 Bioactivity is an associated feature, which implies the facilitation of an appropriate cellular response. Wintermantel and Mayer in 1995 modified the definition of biocompatibility into surface and structural compatibility of implants.7 Surface biocompatibility involves appropriate chemical, biological and physical characteristics, including surface morphology e.g. for cell anchorage and guidance. Structural biocompatibility involves appropriate mechanical properties including elastic modulus, load transmission and deformation characteristics. Optimal biomaterial–host tissue interaction is achieved when both surface and structural elements enable appropriate adaptation of the cellular and mechanical behaviour of the host tissue. In practice, this also requires reproducible and quality-controlled manufacturing and sterilisation processes, optimal surgical technique and patient health. In the skeletal environment, the appropriate cellular responses include the phenomena of osteoconduction, osteoinduction and osseointegration to variable degrees dependent on the target application. Osteoinduction is defined as the process by which osteogenesis (i.e. new bone formation from osteocompetent cells in connective tissue or cartilage) is induced.8 In effect, this phenomenon features in most bone-healing processes. Osteoinductive materials provide a biological stimulus for induction, recruitment, stimulation and differentiation of primitive, undifferentiated and pluripotent stromal cells into osteoblast or preosteoblasts, the initial cellular phase of a bone-forming lineage. The original definition of osteoinductive materials by Marshal Urist states that these biomaterials are capable of inducing bone to form when placed into an extraskeletal site.9 Osteoinductive materials include autografts, demineralised bone matrix (DBM) and specific bone morphogenetic proteins (BMPs) which naturally form bone within the skeleton as well as extraskeletally.10,11 Osteogenic materials are defined as those that contain living cells and are capable of differentiation into bone. Osteoconduction is defined as the process of bony ingrowth from local osseous tissue onto surfaces. The original definition was not strictly restricted to biomaterials,12 however, the contemporary concept of an osteoconductive material is one where bone formation is promoted to appose and conform to its surface, when the material is placed into bone, by virtue of its composition, shape or
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surface texture.8 In effect, these materials act as receptive scaffolds that facilitate enhanced bone formation. Purely osteoconductive biomaterials e.g. hydroxyapatites (HA) are not usually associated with bone formation outside bone.11 Osseointegration is the process of achieving stable direct anchorage and contact between bone and implant. The phenomenon of osseointegration has was first described by Brånemark in 197713 and first defined by Albrektsson et al. in 1981.14 The definition stated direct contact, at the light microscope level, between living bone and implant. It has also been defined at the histological level as the direct anchorage of an implant by formation of bony tissue around the implant without the growth of fibrous tissue at the bone-implant interface.15 A biomechanical definition has also evolved stating that osseointegration is ‘A process whereby clinically asymptomatic rigid fixation of alloplastic materials is achieved, and maintained, in bone during functional loading’.16 This mode of anchorage between bone and implant has been highly successful and well demonstrated in the field of craniofacial implantology. Osteoinduction, osteoconduction and osseointegration are powerful inter-related phenomena in bone regeneration and intrinsically related to bone morphogenetic proteins, bone growth factors and direct bone anchorage factors respectively. It should be noted that these are relative, not absolute, terms. A materials position on the scale of osteoinductivity and osteoconductivity is intrinsically dependent on processing, donor site, surface chemistry, geometry, texture and extrinsically dependent on the biological factors, and the chemical and mechanical tissue environment.
15.2.2 Cellular components, factors and bone healing Synthetic bone graft substitutes including biocomposites require the appropriate cellular responses and interaction with factors to produce successful bone healing. Osteoinduction is the basic biological mechanism of recruiting and stimulating immature cells to develop into preosteoblasts. This phenomenon forms part of the initial phase of the bone morphogenesis cascade and commences immediately after skeletal injury or biomaterial implantation, being particularly active during the first week post-trauma. Actions of the newly developed preosteoblasts are more clearly observed several weeks later in the callus stage of fracture healing.17,18 Bone morphogenesis cascade Bone morphogenesis is a tri-phase sequential cascade: chemotaxis and mitosis of mesenchymal cells, differentiation of the mesenchymal cells initially into cartilage, and replacement of the cartilage by bone. The natural
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repair mechanism of bone utilises bone forming cells, a cartilaginous scaffold upon which the new woven bone is formed, and bioactive molecules to direct the repair process. Cells Bone healing is primarily dependent on the osteoinductive potential of the tissue environment. The natural skeletal tissue environment consists of differentiated bone cells, including osteoblasts, osteoclasts and osteocytes, as well as undifferentiated mesenchymal cells. In the area of bone tissue engineering, four cell types have been used, unfractionated fresh bone marrow; purified, culture expanded mesenchymal stem cells (MSCs), embryonic stem cells (ESCs) and differentiated osteoblasts. The popular theory of response to musculoskeletal injury by Frost states bone, marrow, and soft tissue injury trigger repair and healing by early sensitisation of surviving cells simultaneously with the release of local, biochemical and biophysical messengers directing an appropriate cellular response in the first week post-injury. Some messengers guide cellular differentiation whilst others encourage mitogenesis. This primitive healing response involves the stimulation of a variety of cell types during weeks two to three, including fibroblasts, capillaries and mesenchyme.19,20 However, the proper repair and healing of bone or anchorage of an implant is achieved by the injury-induced recruitment and stimulation of the undifferentiated mesenchymal cells to form osteoprogenitor cells and further development to differentiated bone cells.21 Preexisting, pre-injury differentiated cells, such as osteoblasts, express only a minor contribution toward new bone formation in fracture-healing situations.19,20 MSCs are immature and undifferentiated cells that can be isolated in bone marrow and periosteum. They have the capacity for extensive replication without differentiation and possess multilineage developmental potential.22 Understanding the stimulation, induction and control of differentiation of these cells is the key to harnessing the regenerative power of musculoskeletal tissue. MSCs are either pluripotent or multipotent depending on their lineage. They are capable of differentiation into fibroblastic, adipogenic, reticular and most osteogenic cells.23,24 Stem cells can also be classified into adult and embryonic types. Embryonic stem cells (ESCs) are usually isolated from the inner wall of the pre-implantation blastocysts. These embryonic cells in early developmental stage are more proliferative and pluripotent owing to their indefinite amplification, without the risk of de-differentiation.25 ESCs are non-autologous as they contain a haploid set of chromosomes from a different non-self genetic parent, hence they can be immunoreactive.
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However, recent advances in gene transfer, MHC manipulation and nuclear cloning allows for the formation of autologous-like ESCs.24 This technology elicits a variety of ethical and social implications.26 Future research may see stem cell generation of entire skeletal tissues without biomaterial scaffolds.27 Adult MSCs can also transform through multiple passages without loss of characteristics and tend not to de-differentiate unless exposed to specific biochemical and mechanical cues where they can then be directionally differentiated.28 Jaiswal et al. (1997) reported passing MSCs through 30 population doublings in vitro without loss of osteogenic potential.29 They have been used extensively in experimental bone tissue engineering. A number of animal studies have demonstrated that tissue engineering techniques using autologous MSCs, combined with a suitable scaffold, can create new bone in the defect site. A study by Korda et al.30 examined whether autologous MSCs seeded in an allograft were able to survive a range of normal clinical impaction forces. The study showed that MSCs survived normal impaction forces; this could have clinical significance in the treatement of revision hip surgery. However, the long-term biological effects of the stem cells at implant sites and interactions with biomaterials as issues concerning the phenomenon of cell plasticity, remain largely unknown.31 Stem cell research and technology is still in its infancy. Stem cells require osteoinductive and developmental cues which can be physical and chemical. Osteoinductive stimuli, such as mechanical stress and electrical signals directly or indirectly influence bone induction.32,33 Ultimately, osteoinductive agents are the instigators in transforming undifferentiated mesenchymal cells into preosteoblasts and subsequently differentiated bone cells. This classical cellular sequence is naturally observed in orthotopic environments but is also demonstrated experimentally in heterotopic bone formation, in sites such as muscle, to assess the osteoinductive potential of materials or agents.34 Osteoinductive agents Growth factors (GFs) are proteins secreted by cells that act as signalling and regulation molecules triggering specific target cells to execute specific functions. This cellular communications network influences critical functions such as cell proliferation, matrix production and differentiation of tissues. A GF may affect multiple cell types and may induce several different cellular responses in a variety of tissues. Once a GF binds to a target cell receptor, it induces a ligand–receptor interaction. This induces an intracellular signal transduction system to the nucleus and results in a biological response. These interactions can be very specific, where a specific GF binds
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to a single cellular receptor, or complex, with one or more GFs binding to one or more receptors in order to produce an effect. A number of GFs have been shown to be expressed and play significant roles in bone and cartilage formation, fracture healing, and the repair of other musculo-skeletal tissues. These include somatomedins, insulin-like growth factor I, II (IGF-1, IGF-II), fibroblast growth factor (FGF), transforming growth factor-β (TGF-β), platelet derived growth factor (PDGF), insulin-like growth factor (IGF) and bone morphogenetic proteins (BMPs).35 Osteoinduction is therefore mediated by numerous GFs and, on this basis, it is thought that these GFs may have potential use as therapeutic agents in bone healing and tissue engineering. Hauschka et al.36 (1988) showed that osteogenesis is in part the result of the combination of actions of several growth factors acting at specific stages in different cells. Bone is rich in many growth-stimulating factors but the most important group is the transforming growth factor-beta (TGF-β) superfamily; out of which the bone morphogenetic proteins (BMPs) are the most significant. Bone morphogenetic proteins In 1965, Urist9 was able to elicit new bone formation from an intramuscular injection of demineralised bone matrix. In the late 1970s further research from this original work led to the isolation of the soluble glycoprotein bone morphogenetic protein (BMP) originating from the transforming growth factor-β (TGF-β) superfamily.37,38 This superfamily of growth factors plays an important role during embryogenesis and postnatal tissue repair with BMPs specifically having the greatest osteogenic and osteoinductive potential.39,40 The process of osteoinduction was originally thought to be the result of a single BMP, but 15 BMPs (BMP 2-16) have been discovered to date, mostly distinguished by their potential to induce bone formation de novo via stem cell recruitment, proliferation and differentiation.39–41 Evidence for proposed use in tissue-engineered bone was provided by Wang et al. (1993), who showed that BMPs caused commitment and differentiation of multipotential stem cells in to osteoprogenitor cells.42 Bone healing Understanding and controlling these cellular responses to bone graft materials and biocomposites is necessary to achieve the desired positive biological interaction between graft material and host site. The phenomena of osteoinduction, osteoconduction and osseointegration, resulting in new bone formation, union and incorporation between graft and host leading to
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a skeletal tissue environment that is biologically and mechanically matched to the original. The healing response triggered by placement of bone grafts in skeletal tissue include the host inflammatory reaction to trauma associated with surgical preparation and implantation of the graft, the inflammatory and immune reaction to the graft material itself, and the cellular processes such as proliferation, migration and differentiation. The speed and degree of this response and, thus, graft incorporation into host skeletal tissue is dependent on local and systemic factors, namely the type of graft material and surrounding local tissue integrity, and the systemic physiological state of the host, respectively.11 The cellular events occurring at the graft–tissue interface and graft material itself during its incorporation involve five phases, namely haematoma formation, inflammation, vascularisation, resorption, and regeneration and remodelling. Haematoma formation occurs on trauma related to host-site preparation and graft implantation which also triggers the release of cytokines and growth factors. The inflammatory process involves proliferation, differentiation and migration of mesenchymal cells. Neutrophils, lymphocytes and monocytes are also attracted and recruited to the zone of trauma, migrating to the site and into organizing haematoma. This process also involves platelet adherence, degranulation with release of several growth factors (e.g. FGF-2, PDGF, TGF-β) into a fibrin meshwork. Fibroblasts also produce collagen in response to various stimulatory factors (e.g. TGF-β) combined with the response to the extent of collagen degradation by factors such as metalloproteinases also released by fibroblasts, as well as macrophages.43 The vascularisation process involves the development of fibrovascular tissue within the graft and in the immediate tissue periphery. Moreover, there is vascular invasion into the graft which may also originate from preexisting Haversian and Volkmann canal systems. This phase is essential for provision of the graft environment with nutrients and cell populations. The resorptive process involves focal osteoclasts acting on graft surfaces before bone regeneration is achieved via intramembranous or endochondral ossification. Once mechanically stable, the remodelling process occurs to optimise functionality. The rate of skeletal remodelling, dependent on age, ethnicity, genetics, metabolic and pathological conditions, will influence long-term graft turnover and integrity. The most important factor, however, is mechanical load. Mechanical loads, specifically strain, influences bone remodelling and causes bone to physically adapt in relation to the mechanical loads placed across it.44 Experimental studies have suggested that no strain leads to bone resorption and increasing strain positively alters the anabolic activity of osteocytes directly or indirectly.45 Direct effects have been stated to involve strain-induced changes in the shape of processes
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within the osseous cannaliculi. Indirect effects are considered to involve strain-induced alterations in hydrostatic pressures which may subsequently exert anabolic or catabolic effects on cells, specifically osteoblasts.46 In general, most of these phases are orchestrated by a variety of inflammatory mediators such as kinins, prostaglandins, vasoactive amines, nitric oxide, interleukins, complement, angiogenic and growth factors. In association with the appropriate cellular response, a state of mechanical stability needs to be achieved in order to prevent granulation and fibrosis at the graft–host interface allowing graft incorporation. The important integral factors within local tissues in the zone of implantation include the vascularity of the graft bed alongside the number and competence of the endothelial and osteo-progenitor cells. Limitations in these cellular ingredients will result in a poor response to cues from osteoinductive and osteoconductive materials. Environments deficient in progenitor stem cells limiting graft incorporation include scar or ischaemic tissue, large bony defects, acute or chronic infection, immunosuppresion and radiotherapy.
15.2.3 Traditional options and state of the art Autografts and allografts currently play an important role in the treatment of non-union, bridging diaphyseal defects and filling metaphyseal defects. Vascularised and cancellous autografts are considered the ‘gold standard’ bone graft owing to their excellent capacity for skeletal incorporation. However, they are limited in terms of supply when extensive grafting is required, such as in spinal arthrodesis and repair of large bony defects, and they are associated with post-operative pain and host donor site morbidity.47,48 Allografts such as demineralised bone matrix demonstrate osteoinductive potential and prove useful alternatives to autografts, but the majority are fraught with problems of low osteogenicity, increased immunogenicity, high resorption rates compared with autogenous bone and risk of disease transmission.49 Osteoinductive growth factors, such as bone morphogenetic protein, continue to act as powerful elements in bone regeneration science. Synthetic bone graft substitutes have been developed over several decades to contend with the limitations associated with autograft and allografts. The challenge to develop these biomaterials with high biocompatibility, osteogenicity and osteoinductivity on demand remains. A wide variety of these SBGS materials have been introduced which vary with respect to composition, biological properties and mechanical strength. There is also great variation in level of osteoconductivity and osteoinductivity. Current synthetic biomaterials have usually undergone extensive preclinical testing to negate a significant inflammatory or immune response. The inflammatory reactions to SBGSs are rather dependent on material
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composition, size and surface chemistry. However, the majority do not provoke a major inflammatory reaction at implantation. It is important to note that, in the field of total joint arthroplasty, there is a concern related to osteoconductive synthetic biomaterials causing a wear debris induced inflammatory reaction and secondary bone resorptive process. Particulate debris, of the order of 2 μm, in contact with or phagocytosed by macrophages are thought to trigger the release of cytokines, such as TNF-α and IL-1β, which induce osteolysis. This bone resorptive process occurs by direct or indirect cytokine-induced osteoclastic chemotaxis, maturation, attachment, and activation. The theoretical risk of SBGSs producing particulate matter, secondary to motion or degradation in vivo, and invoking this erosive response exists. Practically, however, the structural composition of SBGSs as large granules, blocks or cements does not appear to create a significant particulate load and the acid solubility of most of these materials implies macrophages are able to contend with these particles without creating a sustained inflammatory response. At present, SBGSs do not appear to cause significant particle-induced bone resorption in clinical practice. In modern surgical practice, composite preparations involving bone marrow or cancellous autografts, demineralised allografts and synthetic grafts (e.g. osteoconductive calcium-based biomaterials, recombinant osteoinductive proteins) are more readily used. To date, no SBGS exists that encompasses all the qualities of autograft. There is a continuing challenge to produce SBGSs in combination with various factors to achieve these qualities, such as by incorporating extracted or synthesised protein growth factors. The challenge to utilise elements and basic materials to replace hard and soft tissues remains. Metals commonly used in implantology include stainless steel and alloys of cobalt chromium and titanium. These generally possess higher strength, ductility and wear resistance, and metals such as titanium and its alloys possess excellent biocompatibility. However, the higher elastic modulus and tensile strength compared with host bone tissues leads to greater stiffness and a host-implant mechanical and load-sharing mismatch. This ‘stress-shielding’ phenomenon has implications on the skeletal environment, leading to necrosis, reduced bone density and loosening of implants or fracture, complications presenting a continual challenge in orthopaedics today. It is also established that metals do not possess surface characteristics conducive for expeditious new bone growth. Furthermore, wear processes may lead to metallosis and production of particulate matter and metal ions which can trigger an adverse immune response. Polymers are generally versatile substances capable of producing a variety of forms with variable properties and compositions e.g. solids, gels, films, fabrics and fibres. They tend to have a comparatively lower modulus and tensile strength than metals, ceramics and most hard tissues, and are too flexible and weak
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for any significant mechanical demand or direct load-bearing applications. Further limitations include fluid absorption, which may lead to leaching of constituents such as monomers and structural alteration during sterilisation processes. Ceramics generally possess good biocompatibility, corrosion and high compression resistance. However, alone, they have a relatively low strength, high brittleness and poorer resilience.
15.3
Biocomposites
15.3.1 Definitions Composites are defined as structures consisting of two (or more) distinct constituents or phases, which in combination form a complex material with different properties from those of the individual components. Biocomposites are composite materials composed of one or more phase(s) derived from a biological origin.50 The resultant physicochemical, mechanical, and biological response enables medical applications that cannot be achieved by the individual components alone.50 The basic constituent phases include a reinforcement phase, such as a particle or fibre, and a continuous matrix phase. Reinforcement phase materials include plant fibres e.g. cotton, flax, hemp, or fibres from recycled papers, woods and food crop by-products along with natural nano-fibrils e.g. cellulose and chitin. Biocomposite continuous matrix phase materials have largely involved natural or synthetic polymers. Natural polymers include those derived from renewable resources, such as vegetable oils, starch,51 protein from grain,52 polylactic acid,53 polyhydroxy alkanoates54 and natural rubber.55 However, commercially available natural-fibrereinforced composites are limited. Synthetic polymers include thermoplastics e.g. polyethylene, polypropylene, polystyrene, polyvinyl chloride as well as thermosets e.g. unsaturated polyesters, phenol formaldehyde, maleates, isocyanates and epoxies. These dominate the commercial biocomposite production market. The overall biocomposite material is anisotropic and multi-functional in nature and allows modulation not only of the physical or mechanical properties, but also of the cellular response. The latter can also be optimised by alterations in specific morphology as well as chemical and physical signals, such as hydrophobic and/or hydrophilic molecules, growth factors, protein and drugs.
15.3.2 Skeletal tissues: the natural composites Tissues in general are ‘natural’ biocomposite structures at the molecular and microscopic level, consisting of organic–inorganic phases which are
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co-ordinated and geared toward their function. Bones and teeth are hard tissues with a complex, hierarchical structure over multiple levels.56 They are essentially composed of an extracellular matrix which has been represented as a two-phase composite composed of a organic type I collagen fibre frame-work matrix (16%) reinforced with inorganic (40–50 vol.%; 50–60 wt.%) hydroxyapatite [(Ca10(PO4)6(OH)2)] crystals alongside ground substance or natural bioactive cement (proteins, polysaccharides, mucopolysaccharides) (2%) and water (22%) containing the bulk of skeletal cells.57,58 Apatite crystals are nanoscale, elongated, plate-like and orientated in relation to the directions of primary stress.56 Macroscopically, the relative fractions of these constituent components vary in relation to bone type, such as cortical (compact, dense) bone compared with cancellous (trabecular, spongy, porous) bone, and patient condition and pathology. Ultimately, skeletal tissue, especially cortical bone, is anisotropic and biologically and mechanically inhomogeneous in nature. Cortical bone has loading fibres orientated along the loading axis with highly organised structures such as the Haversian system. Exceptions, per se, include immature cancellous bone which has been viewed as structurally isotropic, although the chemical content is diverse. Mechanically, hard tissues have a greater elastic modulus and strength compared with soft tissues. Cortical bone has greater apparent modulus of elasticity, tensile strength, and fracture toughness than trabecular bone. However, owing to the highly porous structure (trabecular bone 75–90%; cortical bone 5–10%), relative properties of the extracellular matrix are similar.56 Studies have thus used the mechanical properties of cortical bone as a benchmark before factoring porosity into the picture. Porosity is important for vascularisation and bony in-growth.
15.3.3 Musculoskeletal applications Biomaterials have been synonymous with orthopaedic surgery and found application in joint replacement and internal and external fracture fixation.59 Musculoskeletal tissue engineering and regeneration requires cells, growth factors and a scaffold material with appropriate overall biological and mechanical properties. The challenge involves determining the most successful combination of fabricated three-dimensional scaffold to modulate appropriate cell expression, interaction and function to allow optimal bone regeneration. A variety of biomaterials have been produced which fulfil the requirements of biocompatibility, bioactivity and controlled biodegradability alongside appropriate physicochemical and biological requirements. These include the capability to elicit appropriate cellular responses including adhesion, migration, proliferation, differentiation, formation of extracellular matrix, allowing tissue and vascular ingrowth into interconnecting
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pores, and withstanding physiological loads. In practice, this also requires an inexpensive, reproducible, high-resolution fabrication process, producing materials with a level of structural and functional complexity as close as possible to natural tissue.60 Biocomposites may provide a more intelligent replacement and regeneration material with greater structural biocompatibility than homogeneous, isotropic, monolithic materials alone. Tailoring biocomposites to fulfil the role of a scaffold material may more accurately match the anisotropic nature of tissues and mimic the tissue environment per se. Scaffold-guided tissue engineering provides a transient three-dimensional structure essential for guidance and support of cellular activity. The phase components of biocomposites, especially with the development of nanofibres, may produce a nanocomposite material which not only provides a scaffold matrix with optimal stability, but one with a constitutional make-up at nanoscale levels similar to extracellular matrix. Mimicking this architecture is a major challenge in tissue engineering. A variety of biocomposites have been widely developed and used in musculoskeletal applications and bone regeneration over the last three decades in response to the limitations of conventional materials. Biocomposites have been developed in musculoskeletal applications as plates, screws, bone replacement materials, bone cements and in joint arthroplasty. Classification can be based on primary reinforcement phase forms, including short fibres, continuous fibres and particulate powders. In musculoskeletal application, a useful classification involves the degree of biodegradability, namely fully, partially or non-resorbable compositions. The majority of contemporary biocomposite research and development in this area involves polymer and ceramic-based materials and combinations. Non-resorbable biocomposites make up the majority of materials in orthopaedic surgery and have the inherent advantage of long-lasting mechanical strength and support, particularly useful in large bony injuries and defects. There are a great variety of materials that have been used in total hip and knee joint replacement, internal and external fixation and dental implants and wires. They have also commonly been used as coatings to promote bone growth and osteoconduction on implant surfaces.59 Resorbable biocomposites may be fully or partially degradable and present an attractive concept given the potential to reduce costs, risks and morbidity of reoperation. They may also provide a role as drug delivery systems for factors such as BMPs for skeletal regeneration. Fully resorbable biocomposites in practice are often limited in terms of mechanical strength and ability to withstand greater forces, an important feature in the musculoskeletal system. This is offset against the advantage of complete elimination of the implant from the body. Partially resorbable biocomposites include a reinforcement phase in a matrix phase where at least one component is fully biodegradable in physiological environments. This usually consists of
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a non-absorbable reinforcement material and an absorbable matrix material. These composites are considered more suitable for bridging small bone fractures based on their mechanical limitations. Polymer/ceramic biocomposites Polymer/ceramic composites have been successfully developed for orthopaedic implant applications based on their similarities to the chemical composition of bone, alongside their suitable mechanical and degradation characteristics (see Table 15.1).57,59 Polymer matrix composite materials were one of the first types of biocomposites developed for medical application. The biodegradability of most resorbable synthetic biocomposites and scaffolds is conferred by the polymer matrix, commonly aliphatic polyesters, such as poly(glycolic) acid (PGA), poly(lactic) acid (PLA) and co-polymers poly(lactic-co-glycolic) acids (PLGA).61,62 They have been widely applied owing to their biocompatibility, safety, non-toxicity, controllable biodegradation and metabolisation.61 They produce natural degradation products by non-enzymatic hydrolysis of ester bonds and are eliminated by natural metabolic pathways. PGA is relatively hydrophilic in nature and degrades rapidly in aqueous solutions compared with the more hydrophobic PLA scaffold. This is because an extra methyl group provides increased hold and mechanical integrity.63 Degradation rates can be controlled by varying molecular weight, alteration in enantiomers and PLA–PGA copolymerisation. Poly-l-lactic
Table 15.1 Polymer/ceramic biocomposite components: degradation, mechanical characteristics and uses Biocomposite components
Degradation characteristics
Mechanical/other properties
Aliphatic polyesters; PGA, PLA, co-polymers PLGA, PCL, PPF Collagen, gelatin Ceramic composites; calcium phosphates (HA, TCP, CTP)
Natural degradation products, non-enzymatic hydrolysis of ester bonds
Mechanical properties highly variable and dependent on crosslinking, additives, chain length and branching and enhanced using reinforcement phases Mechanical properties can be enhanced using reinforcement phases e.g. combinations with metals (e.g. fibres), polymers, ceramics
Variation in Ca/P ratio regulates the degradation profiles
Use/application Bone tissue engineering Drug delivery systems
Bone tissue engineering Drug delivery systems
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acid in combination with β-tricalcium phosphate has been used successfully in vivo as fully resorbable biodegradable interference screws in anterior cruciate ligament reconstruction with good functional results.64 A less common aliphatic polyester used in biocomposite technology for bone tissue engineering is poly(ε-caprolactone) (PCL) which degrades at a significantly slower rate than PLA, PGA and PLGA,65 making it more appropriate for developing longer term fixation systems. PCL and poly(propylene fumarate) (PPF) are synthetic biodegradable polymers being increasingly developed to improve polymer/ceramic biocomposite properties. PPF is similar to PLA and PGA polymers in terms of degradation characteristics and can be variably synthesised to produce a variety of mechanical properties and orthopaedic applications.66 Collagen, a naturally derived polymer and major fibrous protein in extracellular tissue matrices, has also been popularly used in biocomposite manufacture owing to its biodegradability, enhanced cell adhesion and high biocompatibility. In isolation, it has poor mechanical properties, unreliable sampling and uncontrolled biodegradability. Denatured collagen, gelatin, has also been used in bone tissue engineering, particularly in drug delivery and porous scaffold production. Ceramic composite materials constitute the inorganic component of resorbable polymer/ceramic biocomposites. They are commonly used owing to their similarities with natural bone. The most popular ceramics used in bone tissue engineering and controlled drug delivery systems include calcium phosphates [e.g. calcium tetraphosphate (Ca4P2O9) (CTP), tricalcium phosphate [Ca3(PO4)2,TCP], hydroxyapatite [Ca10(PO4)6(OH)2,HA], and their derivatives and combinations. A similar Ca/P ratio to bone and the ability to form bonds with soft and hard tissues classifies them as true bioactive materials. They are also osteoconductive, with surface characteristics that support osteoblast adhesion, proliferation and differentiation, and osteoinductive with the ability to bind BMPs and osteopromotive factors (Figs. 15.1 and 15.2).67,68,69 The variation in Ca/P ratios between the different calcium phosphates regulates the degradation profiles. They vary in degrees of crystallinity from low, such as in TCP, to high, such as in HA, which are inversely related to dissolution rates. HA has an excellent osteogenic potential and TCP has a relatively faster degradation rate.70 HA/ceramic biocomposites involve use of ceramic reinforcement phases e.g. TCP, Si3N4, HA whiskers, Al2O3 platelets and ZrO2 particles.71 HA/metal biocomposites include HA dispersed with silver, using conventional sintering methods, producing composites of increased fracture toughness with the benefit of the inert, ductile, and anti-bacterial nature of silver.72 HA/polymer biocomposites appear to offer the most promising and robust system in terms of synthetic bone engineering and substitution.58 Bonfield et al. were the first to develop HA-reinforced polymer biocomposites as bone analogue biomaterials tailored to replicate skeletal tissues.73 The original design
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5 KV
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15.1 Human mesenchymal stem cells growing on the surface of hydroxyapatite – β-tricalcium phosphate. Cell filopodia can be seen anchoring onto the material surface.
rationale was to toughen and reinforce biocompatible polymer matrix with a bioactive HA filler. Several studies have investigated polymer/HA composition, molecular orientation, morphology, and interface through modified or novel processing.58 Bone cement biocomposites have shown increasing popularity and use as osteoconductive bone graft substitutes over the past 10 years. Their preparation involves a non-polymerisation, near non-exothermic, acrylic cement-like production using powders e.g. monocalcium phosphate, tricalcium phosphate and calcium carbonate, mixed in a solution e.g. sodium phosphate. Products demonstrate high strength (10–100 MPa in compression; 1–10 MPa in tension) and weakness under shear.71 Bone cement composites e.g. calcium phosphate cements, are used in fracture management, and have been applied to improve compression strength of vertebral bodies in fractures and injuries related to osteoporosis.74,75 (Nano)-composites and biomimetic / hybrid biocomposites: the next generation Nanotechnology is the science of creating and utilising materials with dimensions under 100 nm in at least one direction.50 Biomimetic and nanocomposite materials are super-specialised, dense hybrid structures that
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(a)
(b)
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15.2 Human osteoblast cells on hydroxyapatite – tricalcium phosphate at 14 days (a) Cells growing on the surface and within the macropores of the biocomposite material; (b) Dense cell layers were observed ‘capping’ the pores.
more extensively express the intrinsic advantage of biocomposites, namely the ability to match the nano-architecture of natural, living tissues. Bone is a tough hybrid tissue with strong interactions between its constituent organic and inorganic phases. It is considered to be a nano-structured composite material consisting of a polymer matrix, namely the organic phase
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of nanometre-scale collagen fibres, reinforced with nanometre-scale ceramic particles, namely the inorganic compound of carbonated HA.50 Synthetic polymer/ceramic biocomposites are ideal choices for mimicking skeletal tissue. Experiments have shown they may also recreate the heterogeneous nanosurface roughness and surface nano-scale particles required for appropriate cellular interaction including osteoconductivity.76–80 Replication at this level using polymer composites with ceramics (e.g. HA, zinc oxide, alumina) and metals (e.g. CoCrMo, Ti, Ti6Al4V) is better than using micrometre-sized particles.77,78,80,81,82 Polymer-nanoceramic biocomposites have demonstrated increased osteogenesis, osteoblast function and protein adsorption, quantitatively assessed by fibronectin and vitronectin levels, on the nano-biocomposites compared with micrometre-sized ceramics.77,78,81 Ultimately, nano-composites consist of increased atomic levels, electron distribution, defects and overall properties at their surfaces. This more closely replicates physiological bone at the nano-scale level biologically and mechanically. It is widely believed to enhance surface reactivity of the materials compared with traditional, conventional options with surface particles that are micrometre-size in scale.77–79,81 High-resolution nano-scale analyses, by e.g. atomic force microscopy, of bone revealed the complex organic–inorganic biomineralized composite architecture.83 Structure is intrinsic to function and mechanical fracture models show that the organic phases often act plastically as energy dissipation networks that form bridges along crack propagation at a nanoscopic level. Understanding and modelling practical behaviour of biocomposites as substitutes for skeletal tissues is essential for future clinical applications.
15.4
Cellular responses and experimental testing
Brief examples of in vivo experimentation include various ceramic–polymer biocomposite configurations. A biphasic calcium phosphate–poly-dl-lactide composite with and without biostimulative agents has been synthesised and successfully used as a replacement and regeneration material for osteoporosis damaged bone in rats.84 A nano-HA and polyamide composite implanted in rabbits demonstrated bioactivity and direct bonding with cortical bone without fibrous tissue and development of a bone-like apatite layer.85 Partially resorbable HA biocomposites have also been developed to deliver growth factors in vivo, providing a novel solution for tackling the major orthopaedic problem of fracture non-union.86 A HA-based biocomposite combining fibrin incorporating and delivering biologically active fibroblast growth factor (FGF-1) has been developed.86 The aim was to promote site-specific induction and coordination of neovascularisation and
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Table 15.2 Biocomposites: examples tested in vitro HA–polymer HA–high-density polyethylene (HAPEXTM) HA–hyaluronic acid (HYAFF-11HA ®) HA–Ceramic HA–ZrO2 HA–Al2O3 HA–metal HA–silver Other Bioglass ® HA-bone cements, e.g. HA-PMMA, HA-PEMA, HA-PHEMA, HA-PHEMA/PCL. Nanocomposites +/− HA and combinations, e.g. nHA–PVA, Nha–collagen, nHA– polyamide, nHA–silk fibroin, nHA–PLLA, nHA–PCL
osteogenesis, the processes considered to be sub-optimal and implicated in cases of skeletal non-union. A sustained delivery profile was observed alongside positive cellular responses. An extensive body of work has done on the characterisation, development and analysis of polymer / ceramic biocomposites in vitro. Several studies have been conducted and we provide some examples, including HAPEXTM, HA-bone cements, HYAFF® 11-HA, HA–nanocomposites, BIOGLASS® and metal composites (Table 15.2). HAPEXTM An HA-reinforced high-density polyethylene (HDPE) composite (HAPEXTM) has been developed as a second-generation orthopaedic biomaterial tailored to augment and substitute cortical bone.73 The optimisation, mechanical and biological characterisation of HAPEXTM is well established.58,73 Studies have shown maintenance of tensile strength and Young’s modulus in physiological conditions alongside positive effects on stimulation, proliferation and differentiation of human osteoblast cell lines.73,87–89 HA particles within HAPEXTM were originally observed to act as micro-anchors, providing favourable cell attachment sites and enhancing the potential of the material for osseointegration.87 Variations in HA volume concentrations on HDPE, specifically 40% HA content, with controlled surface topography, showed greater bioactivity with enhanced primary human osteoblast proliferation, differentiation, expression and cytoskeletal organisations rates alongside increased focal contact points (Fig. 15.3).88,90,91 Surface topography was found to have a significant influence on bioactivity and cellular interactions such as attachment, and pos-
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15.3 Layers of adhered and actively dividing human osteoblast cells on HAPEXTM.
tulated to determine the success of skeletal tissue regeneration, and longevity of implants.92 Topographical optimisation by polishing followed by roughening of the surface of non-machined HAPEXTM demonstrated enhanced osteoblast recruitment, proliferation and phenotype expression.73 Another study observed better cellular response to rough-surfaced HA filled in HA/HDPE composites than with smooth-surfaced implants.89 Overall, HAPEXTM has shown optimal interaction with osteoblasts with studies showing better bioactivity than with poly(methyl methacylate) bone cement, a well established orthopaedic composite since the 1960s.93 Clinical applications of HAPEXTM include use as middle ear prostheses and orbital floor implants. Positive interactions with osteoblastic cells originating from crania has stimulated research and development as a skull reconstruction implant.89 HA–bone cements Poly(methyl methacrylate) (PMMA) bone cement was introduced to orthopaedic surgery by Sir John Charnley in the 1960s and continues to remain the standard for cemented arthroplasty procedures in orthopaedics today. In cemented joint replacements, the implant–cement interface is known to
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form a strong bond, however, the bone–cement interface is considered weaker with a greater stimulatory effect on fibroblastic cells than direct bone contact. This interface contributes to implant loosening and implant failure, and is implicated in periprosthetic fracture, often requiring revision surgery. In isolation, cement is shown to have lower bioactivity and a poorer effect on the osteoblast populations than with HA composites such as HAPEXTM.93 Investigators have successfully introduced HA as a bioactive phase to cement, and developed (8.8 and 17.5 wt.%) HA–PMMA biocomposites in an attempt to produce stronger implants with greater direct bone apposition and minimal fibrous tissue encapsulation.94–96 The studies showed greater osteoblastic proliferation, differentiation and phenotype expression alongside preferential anchorage, focal contact formation, cytoskeletal organisation and overall interaction of cells to HA particles exposed at the cement surface than PMMA itself.94 HA–poly(ethyl methacrylate) (PEMA) bone cement biocomposites have also been developed and they show preferential anchorage to surfaceexposed HA but do not enhance cellular adhesion, proliferation or differentiation.97 The authors attributed this to increased residual monomer with HA incorporation. Ethylmethacrylate has also been developed in a cross-linked, polyhydroxylated form (PHEMA) incorporating linear poly-(ε-caprolactone) (PCL) and reinforced with HA to produce a semi-interpenetrating polymer network.98 Studies show that PHEMA/PCL/HA [PHEMA/PCL 70/30 (w/w) + 50% (w/w) HA] in a hydrated state had a positive mechanical and biological performance with appropriate swelling behaviour, undetectable weight loss, no release of toxic leachables and an environmental pH acceptable for cell growth (Figs. 15.4 and 15.5).98 The incorporation of HA provided mechanical enhancement providing an implant elastic modulus similar to that of trabecular bone suggesting its use as a filler or substitute in skeletal zones predominated by spongy bone. HYAFF®11–HA HYAFF®11–HA is a novel biocomposite composed of a benzyl ester of hyaluronic acid polymer matrix reinforced with HA.99 The material was developed as a biocompatible, biodegradable composite with a similar configuration to the mineral phase of bone, favouring osteoconduction. Studies also showed a grainy topography with indications of HA on the surface, a feature considered optimal for cellular interaction with implants in skeletal environments.94,99 A lack of cytotoxicity or inhibition of cell proliferation with acceptable pH for cell viability and improved mechanical compressive properties suggested use of this biocomposite in trabecular bone reconstruction such as in vertebral defects.99
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15.4 Mesenchymal stem cells growing on the surface of hydroxyapatite reinforced PHEMA/PCL biocomposite.
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15.5 Human osteoblasts at 28 days in culture growing on nanocarbonate substituted hydroxyapatite reinforced PHEMA/PCL composite.
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HA nanocomposites Nanocomposites developed to date have shown promising bioactivity and biocompatibility for application in skeletal tissue regeneration. Biocomposite nano-particles of HA (n-HA) combined with binders including poly(vinyl alchol) (PVA) and collagen,100 polyamide,101 as well as silk fibroin102 scaffolds amongst others have been developed. The n-HA was synthesised at ultimate size ranges of 10–50 nm and incorporated or in situ synthesised with PVA/collagen binders demonstrating relatively high elasticity and linear viscoelastic function.100 Collagen incorporation provided internal porosity to the biocomposite, a feature essential for vascular and bony ingrowth. This relationship has been further observed in studies involving HA–collagen composites of various macropore sizes demonstrating greater cell migration, infiltration and differentiation of human osteoblasts from collagen into HA in scaffolds with larger pores.102 The larger pore composites were recommended as more appropriate models for testing biofunctionality and cell interactions at trauma sites in implantation in the skeletal system.103 An apatite–collagen nanocomposite, produced by a co-precipitation method, was successfully fabricated and aimed at more accurately mimicking developing bone by incorporation of silicone as an osteopromotive agent.104 Variations in the apatite–collagen ratios demonstrated markedly different morphologies ranging from ceramic-like particles to rope-like macrofibrils macroscopically, but nanostructurally indistinct features of apatite nanocrystals randomly distributed in an amorphous collagen matrix. HA–PVA biocomposites have also been successfully developed and characterised as tissue engineering scaffolds using a selective laser sintering rapid prototyping process.105,106 Studies show this technique provides greater control, increased spatial and microstructure precision and reproducibility of biocomposites, features often limited in conventional manual scaffold fabrication techniques.106 This has been considered vital in producing complex objects with pre-defined macro- and micro-strucutres, in order to replicate bony and craniofacial defects.105 Collagen and polymers, such as poly(l-lactic acid) (PLLA), have also been fabricated with apatite as in an accelerated biomimetic process.107 The investigators observed a complex composite coating on polymer-containing submicrometre bone-like apatite particles and collagen fibrils, similar in composition to natural human bone. This chemical–physical surface coating was considered as an optimal feature to facilitate cell interaction and osteoconductivity. n-HA–polyamide nanocomposites were developed and studied in vitro and in vivo with implantation in rabbit femora.101 Investigators observed the development of a bone-like apatite layer containing carbonate on implant surfaces and highlighted the importance of this feature in establish-
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ing a direct, non-fibrous, bone-implant bond which enhances bioactivity in the skeletal environment. Another novel porous tissue engineering scaffold of n-HA and silk fibroin was developed by a freeze-drying co-precipitation technique.102 The biocomposite was observed to have good homogeneity, interconnecting porosity and excellent mechanical properties, including compressive modulus and strength, attributed to the beta-sheet structure of silk and n-HA volumes up to 70% (w/w). Biocomposite nano-fibres have also been developed using polycaprolactone / n-HA and type I collagen (PCL/nHA/ Col) fabricated using a simple electrospinning technique.108 The variably sized nanofibrous composites were shown to provide mechanical support and developed into a scaffold using a deposition process. The interconnected fibre network provided porosities greater than 80% and an openpore structure for direct growth and colonisation of human fetal osteoblasts, and vascular in-growth and transportation channels for nutrients and metabolic waste. PCL/nHA/Col also promoted expeditious mineralisation and surface functionalisation of hFOB activities. Other works on calcium-based nanocomposites include a nano-sized calcium silicate (n-CS) combination with poly(ε-caprolactone) (PCL) fabricated using a n-CS slurry in a solventcasting method.109 In vitro testing of bioactivity and biocompatibility demonstrated excellent surface apatite formation and positive interaction, respectively, with mouse fibroblast. Mechanical properties were also enhanced with the incorporation of n-CS, which was stated to improve and optimise hydrophilicity, compressive strength and elastic modulus. Biological properties were reportedly improved with n-CS with improved cell attachment and proliferation than with PCL alone. Bioglass® Bioglass® reinforced with HDPE was developed and showed good in vitro bioactivity with formation of a biologically active hydroxylcarbonate apatite layer on the biocomposite surface similar to bioactive glassceramics.110,111 Biocompatibility with primary human osteoblast-like cells was also demonstrated supporting evidence for its use as a skeletal tissue engineering biocomposite. Metal composites Titanium dioxide (TiO2) and TiO2 glass composites with PCL were fabricated using a sol–gel process and powder compaction.112 Greater mechanical properties were a function of manufacturing conditions, specifically the compression moulding conditions alongside the PCL weight added to TiO2. The study demonstrated production of a biocompatible composite with
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strength and stiffness properties spanning those of spongy and cortical bone that could be tailored by varying the polymeric phase, which was essentially capable of toughening TiO2-based materials.
15.5
Conclusions
Tissues are intrinsically composite structures. Biocomposites are an ideal, state of the art option in tissue engineering targeted at matching the natural microstructures of tissues. Nanocomposites, the next generation of biocomposites, are aimed at mimicking tissues at the nano-scale level. Individual constituents of biocomposites can be suitably combined to tailor-make a biomaterial for repair, regeneration and replacement of tissues by eliciting the appropriate cellular response. In terms of musculoskeletal applications, biocomposites can be composed of strong synthetic molecules which are similar to skeletal tissues in structural properties, such as porosity and surface topography, and function. In order to reach this goal, biodegradable polymers and bioactive ceramics have been extensively studied and optimised to produce novel hybrid biocomposite systems. The challenges facing development, production and application of biocomposites in orthopaedic surgery involve optimising fabrication and processing technologies, biocompatibility and bioactivity. Investigators must understand and improve the mechanical and biological properties such as strength, elastic modulus, nano-scale and nano-surface architecture, composite–tissue interface, porosity and overall osseous cellular response. Tissue engineering using polymer–ceramic and other biocomposites also demands increased definition and control of biodegradation profiles alongside drug and factor release kinetics. The resolution and three-dimensional precision of biodegradable polymer – bioactive ceramic biocomposites requires further expansion. In this chapter, the place of biocomposites has been highlighted in relation to current options for bone substitution, cellular requirements and contemporary work on cellular responses in relation to these implants.
15.6
References
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84. ignjatovic n, ajdukovic z, uskokovic d. New biocomposite [biphasic calcium phosphate/ poly-dl-lactide-co-glycolide/biostimulative agent] filler for reconstruction of bone tissue changed by osteoporosis. J Mater Sci Mater Med, 2005, 16(7), 621–626. 85. yang k, wei j, chao yuan, li bao. A study on in vitro and in vivo bioactivity of nano hydroxyapatite/polymer biocomposite, Chin Sci Bull, 2007, 52(2), 267–271. 86. kelpke ss, zinn kr, rue lw, thompson ja. Site-specific delivery of acidic fibroblast growth factor stimulates angiogenic and osteogenic responses in vivo. J Biomed Mater Res A, 2004, 71(2), 316–325. 87. huang j, di silvio l, wang m, tanner ke, bonfield w. In vitro mechanical assessment of hydroxyapatite-reinforced polyethylene composite. J Mater Sci Mater Med, 1997, 8, 809–813. 88. di silvio l, dalby m, bonfield w. In vitro response of osteoblasts to hydroxyapatite-reinforced polyethylene composites. J Mater Sci Mater Med, 1998, 9, 845–848. 89. zhang y, tanner ke, gurav n, di silvio l. Osteoblastic response to 30 vol% hydroxapatite–polyethylene composite. J Biomed Mater Res Part A, 2007, 18-A(2), 409–417. 90. dalby mj, di silvio l, davies g, bonfield w. Surface topography and Ha filler volume effect on primary human osteoblasts in vitro. J Mater Sci Mater Med, 2000, 12, 805–812. 91. di silvio l, dalby mj, bonfield w. Osteoblast behaviour on HA/PE composite surfaces with different HA volumes. Biomaterials, 2002, 23, 101–107. 92. dalby mj, di silvio l, gurav n, annaz b, kayser mv, bonfield w. Optimising Hapex TM topography influences osteoblast response. Tiss Eng, 2002, 8(3), 453–467. 93. dalby mj, bonfield w, di silvio l. Enhanced HAPEX topography: comparison of osteoblast response to established cement. J Mater Sci Mater Med, 2003, 14, 693–697. 94. dalby mj, di silvio l, harper ej, bonfield w. In vitro evaluation of a new cement polymethylmethacrylate cement reinforced with hydroxapatite. J Mater Sci Mater Med, 1999, 10, 793–796. 95. dalby mj, di silvio l, harper l, bonfield w. Increasing the hydroxyapatite incorporation into poly(methylmethacrylate) cement increases osteoblast adhesion and response. Biomaterials, 2002, 23, 569–576. 96. dalby mj, di silvio l, harper ej, bonfield w. Initial interaction of osteoblasts with the surface of a hydroxyapatite–poly(methylmethacrylate) cement. Biomaterials, 2001, 22, 1729–1747. 97. opara tn, dalby mj, harper ej, di silvio l, bonfield w. The effect of varying percentage hydroxyapatite in poly(ethylmethacrylate) bone cement on human osteoblast-like cells. J Mater Sci: Mater Med, 2003, 14(3), 277–282. 98. giordano c, causa f, di silvio l, ambrosio l. Chemical–physical and preliminary biological properties of poly(2-hydrotheylmethacrylate)/poly-(εcaprolactone)/hydroxyapatite composite. J Mater Sci Mater Med, 2007, 18(4): 653–660. 99. giordano c, sanginario v, ambrosio l, di silvio l, santin m. Chemical– physical characterization and in vitro preliminary biological assessment of
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Biomedical composites hyaluronic acid benzyl ester-hydroxyapatite composite. J Biomater Appl, 2006, 20, 237–252. degirmenbasi n, kalyon dm, birinci e. Biocomposites of nanohydroxyapatite with collagen and poly(vinyl alcohol). Colloids Surf B Biointerfaces, 2006, 48(1), 42–49. jie w, hua h, lan w, yi h, yubao l. Preliminary investigation of bioactivity of nano biocomposite. J Mater Sci Mater Med, 2007, 18(3), 529–533. liu l, liu j, wang m, min s, cai y, zhu l, yao j. Preparation and characterization of nano-hydroxyapatite/silk fibroin porous scaffolds. J Biomater Sci Polym Ed, 2008, 19(3), 325–338. tsiridis e, gurav n, bailey g, sambrook r, di silvio l. A novel ex-vivo culture system for studying bone repair. Injury, 2006, 37S, S10–S17. lynn ak, nakamura t, patel n, porter ae, renouf ac, laity pr, best sm, cameron re, shimizu y, bonfield w. Composition-controlled nanocomposites of apatite and collagen incorporating silicon as an osseopromotive agent. J Biomed Mater Res A, 2005, 74(3), 447–453. chua ck, leong kf, tan kh, wiria fe, cheah cm. Development of tissue scaffolds using selective laser sintering of polyvinyl alcohol/hydroxyapatite biocomposite for craniofacial and joint defects. J Mater Sci Mater Med, 2004, 15(10), 1113–1121. wiria fe, chua ck, leong kf, quah zy, chandrasekaran m, lee mw. Improved biocomposite development of poly(vinyl alcohol) and hydroxyapatite for tissue engineering scaffold fabrication using selective laser sintering. J Mater Sci Mater Med, 2008, 19(3), 989–996. chen y, mak af, wang m, li j. Composite coating of bonelike apatite particles and collagen fibers on poly l-lactic acid formed through an accelerated biomimetic coprecipitation process. J Biomed Mater Res B Appl Biomater, 2006, 77(2), 315–322. venugopal j, vadgama p, sampath kumar ts, ramakrishna s. Biocomposite nanofibres and osteoblasts for bone tissue engineering. Nanotechnology, 2007, 18(5), 511–518. wei j, heo sj, liu c, kim dh, kim se, hyun yt, shin jw, shin jw. Preparation and characterization of bioactive calcium silicate and poly(epsilon-caprolactone) nanocomposite for bone tissue regeneration. J Biomed Mater Res A, 2008, 18 (Epub ahead of print). huang j, di silvio l, wang m, tanner ke, bonfield w. In vitro assessment of hydroxyapatite and Bioglass® reinforced polyethylene composites. In: Bioceramics, Vol. 10 Eds L Sedel and C. Ray (Proceedings of the 10th International Symposium on Ceramics in Medicine, Paris, France). Elsevier Science Ltd. 1997, 519–522. huang j, di silvio l, wang m, rehman i, ohtsuki c, bonfield w. Evaluation of in vitro bioactivity and biocompatibility of Bioglass-reinforced polyethylene composite. J Mater Sci Mater Med, 1997, 8, 809–813. santis rde, catauro m, di silvio l, manto l, raucci mg, ambrosio l, nicholais l. Effects of polymer amount and processing conditions on the invitro behavior of hybrid titanium dioxide/p-polycaprolactone composites. Biomaterials, 2007, 28, 2801–2809.
16 Testing the in vivo biocompatibility of biocomposites R. G I A R D I N O, Rizzoli Orthopaedic Institute and Bologna University Medical School, Italy; M. F I N I, N. N I C O L I A L D I N I, A. PA R R I L L I1, Rizzoli Orthopaedic Institute, Italy
Abstract: The evaluation of the biocompatibility of a material before clinical use is a long trial that requires a well planned sequence of steps. These procedures are now standardized at an international level. In particular, in the European Community, the rules of EN ISO 10993 are the guideline for each investigation in this field. The in vivo tests of biocompatibility are complementary to those performed in vitro and require compliance to the ethical and legal rules on animal experimentation. Biocompatibility tests must be performed on the final product or material, taking into account the nature, duration and conditions of the exposure in the human body, the physical and chemical features of the product, the toxicological activity of the chemical elements or compounds, and the presence of leachable materials. The tests should be applied with interpretation and judgement by the appropriate professionals qualified by training and experience taking into consideration factors relevant to the device/material, its intended use, and the current knowledge of the device/material provided by scientific reports and previous clinical experience. Key words: biocompatibility, in vivo animal models, preclinical studies, biomaterials, biocomposites.
16.1
Introduction
The assessment of biocompatibility is a mandatory prerequisite for all materials and devices before their clinical application. The definition of the concept of biocompatibility is the absence of an adverse response after contact with the human body. According to the 2005 Consensus Meeting of the European Society of Biomaterials ‘Biocompatibility refers to the ability of a material to perform with an appropriate (host) response in a specific application’.1 This statement emphasizes the fact that the evaluation of the biocompatibility of a material must be closely related to its applications. According to Williams2 the response of a host to a material is extremely varied, and is controlled by factors involving the host, the material and the surgical procedure. 385
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The evaluation of the biocompatibility of a material assigned to clinical use is a long trial that requires a well planned sequence of steps. These procedures are now standardized at an international level. In particular in the European Community, the rules of EN ISO 10993 provide the guidelines for each investigation in this field. This chapter deals with in vivo biocompatibility tests; these tests must be considered as complementary of in vitro tests and they usually represent an extension of those texts. Indeed, the capability of a biomaterial to perform a specific function cannot be evaluated by single in vitro or single in vivo methodology. In vitro methods provide necessary and useful results3 that precede and complete the in vivo testing.4 They decrease the number of animals required for in vivo biocompatibility evaluation. By adopting both in vitro and in vivo methods, it is possible to reduce limits and increase benefits of each single approach. Moreover, in vivo testing allows long-term investigations, which, by reproducing the complex biological environment, are suitable for the evaluation of biofunctionality. But there are also many drawbacks that must be considered: there is the need for appropriate facilities and complex professional expertises, prolonged times and high costs. Performing in vivo studies requires, first of all, compliance to ethical and legal rules on animal experimentation. For this reason, before the presentation of in vivo tests for biocompatibility, it is necessary to dwell upon these prerequisites of animal research.
16.2
Preclinical in vivo experimentation: ethical and legal requirements
Biomedical research using animals is regulated worldwide by precise rules, whose cornerstones are the appropriate use of the animals and the humane treatment of them.5 In the United States the Animal Welfare Act dates from 1966 and this was followed in 1986 by the Public Health Service Policy on Human Care and Use of Laboratory Animals and by the NIH Guide for the Care and Use of Laboratory Animals. All these documents contain guidelines about the animal environment, housing and management as well as veterinary care. In the European Union, the Directive 86/609 – Protection of Experimental Animals give to the state members of the European Community indications to be applied in the regulation of animal research. According to this directive, the Italian Government with the Law by decree number 116 of January 27, 1992, on the protection of animals used for scientific purpose, settled the guidelines for this matter.6 Considering, in particular, the biological evaluation of medical devices, part 2 of ISO 10993 ‘Animal Welfare Requirements’7 contains the direc-
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tives to be followed when animal tests are to be performed. The tests are justified only when the resulting data are not otherwise available and essential for the material characterization, when no other scientifically validated method not involving animals is available, and when strategies to minimize pain, suffering, distress and lasting harm have been identified and implemented. Animal experiments must be carried out in authorized laboratories, with appropriate facilities for animal stabling and welfare. Cages should be appropriate for small, medium and large animals. Surgical procedures must be performed in general anaesthesia and in well equipped operatory rooms, under the control and responsibility of well trained scientists. Strategies shall be adopted to minimize the number of animals involved in the tests. In particular, animal tests shall not be performed before appropriate preliminary in vitro tests have been carried out with favourable results. ISO 10993-2 emphasizes also the humane requirements of animal welfare such as the conditions of stabling and the procedures of euthanasia. The study protocol shall refer to the specific ISO 10993 requirements, and must describe the scientific aims of the test; the notices about the material to be investigated; the rationale for the use of animals; a detailed and complete study documentation. The research protocol must also be approved by an Ethical Committee and by the public authority, according to the rules of each country.
16.3
ISO 10993 and biocompatibility tests
The International Organization for Standardization (ISO) is a worldwide federation of national standards bodies for the biological evaluation of medical devices.8 ISO 10993 refers to the fundamental principles governing the biological evaluation of medical devices, the definition of categories of devices based on the nature and duration of the contact with the body, and the selection of appropriate tests to be performed. Table 16.1 summarizes the rules contained in ISO 10993. In the ISO 10993 introduction, it clearly states that ‘the protection of humans is the primary goal’. The guidelines originate from the international and national standards about the biological evaluation of medical devices. ISO tests should be applied with interpretation and judgement by the appropriate professionals qualified by training and experience considering the factors relevant to the material, its expected applications, and the current knowledge provided by scientific literature and previous clinical experience. In particular, the biocompatibility tests must be performed on the final product or material, taking into account the type, duration and conditions of the exposure in the human body, the physical and chemical features of the product, the toxicological activity of the
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Table 16.1 ISO Rules 10993 1. Guidance to the selection of tests-evaluation and testing 2. Animal welfare requirements 3. Tests for genotoxicity (in vitro and in vivo), carcinogenicity (in vivo) and reproductive toxicity (in vivo) 4. Selection of tests for interactions with blood (in vitro) 5. Tests for in vitro cytotoxicity (in vitro) 6. Tests for local effect after implantation (in vivo) 7. Ethylene oxide sterilization residuals 9. Framework for the identification and quantification of potential degradation products 10. Test for irritation and sensitization (in vivo) 11. Tests for systemic toxicity (in vivo) 12. Sample preparation and reference materials 13, 14, 15. Identification and quantification of degradation products from polymers, ceramics, metals and alloys 16. Toxicokinetic study design for degradation products and leachables
chemical elements or compounds, the presence of leachable materials. Knowledge of current information from preclinical and clinical studies is also recommended. With respect to the type of contact with the body, ISO 10993-1 classifies the devices into three main groups: •
Surface devices may be in contact with intact skin, intact mucosal membranes and wounded or compromised surfaces (i.e. ulcers, burns, granulation tissue). • External communicating devices may have an indirect contact with the blood (i.e. set for fluid administration), a contact with soft and hard tissues (i.e. dental cements) and a contact with circulating blood (i.e. intravascular catethers, dyalizers). • Implant devices may be in contact with bone (i.e. screws, pins and plates, orthopaedic prostheses), with soft tissue (i.e. breast implants), or with blood (heart valves, vascular grafts). With respect to the duration of the contact with the body, there are three time periods: limited contact (<24 h), prolonged contact (24 h to 30 days) and permanent contact (>30 days). It is not required that a material or device will be submitted to all the tests, but only to those related to the specific contact with the body environment.9 According to this guideline, in the following paragraphs, we consider, in general, the in vivo tests; the test most appropriate for the purpose of the investigations will be selected.
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Extraction and sample preparation
ISO 10993-12 deals with the preparation of the samples for the biocompatibility tests.10 The toxicity of a material is mainly related to its soluble components, therefore some of the biocompatibility tests are performed using extracts of the material. For this reason, it is advisable to treat this subject before considering each single test. The extraction is performed to provide a test sample suitable for determining the biological reactivity and the hazard potential of any leachable and potentially dangerous substances that may be released from the device into the surrounding tissues. The quantity of the extract is related to the period of extraction, temperature, ratio between the surface area of the material, volume of extractant and the nature of extractant. The extraction medium is selected according to the foreseen use and the nature of the device. The samples of the device or material are incubated in the extraction medium. There are definite ratios between surface area and extraction medium volume that are related to the thickness of the device or material. A common extraction procedure is to incubate the test device or material at 37 °C for 24 h at a ratio of either 60 or 120 cm2, per 20 ml of extraction medium. The device should be a representative specimen of the final device and it should be processed (i.e. sterilized) in the same way of the final device. For this reason, composite materials are tested as finished materials. Elastomers, coated materials and composites must be tested intact whenever possible, because of potential differences in extraction between intact and cut surfaces. The extraction must be performed in clean, chemically inert and closed containers. Polar and nonpolar solvents can be used for the extraction. Polar solvents are water, physiological saline and culture media without serum. Non-polar solvents are vegetable oils. Additional extraction media are ethanol/water, ethanol/ saline, polyethylene glycol 400, dimethyl sulfoxide and culture media with serum. The extraction time is chosen to maximize the amount of material extracted: the use of standard conditions is recommended. Liquid extracts will be used immediately after preparation to prevent changes in composition.
16.5
Irritation and sensitization test
These tests are designed to determine the allergic response involving immunological systems and the localized inflammatory response to single, repeated or continuous application of the test substance, without involvement of an immunological mechanism. ISO 10993-10 recommends both these tests.11 At the present time, there are no validated in vitro tests to detect irritant or sensitizer materials.
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The irritation reaction is the primary and direct effect of a substance on the skin. On the contrary the sensitization reaction is an immuno mediated contact dermatitis.
16.5.1 Irritation test The potential of the material to be tested to produce dermal irritation can be evaluated using healthy young albino rabbits of either sex. One day before the test, the back of the animal is shaved at both sides of the spine. The test material is applied directly on the test site, covering the site with a non-occlusive dressing such as a gauze patch, and then wrapping the application site with a semi-occlusive bandage. At the end of the scheduled contact time the dressing is removed. For acute tests (single exposure), the appearance of the site of application must be checked at 1, 24, 48 and 72 h following the removal of the patches. For repeated exposures, the appearance of application site will be recorded 1 h after the removal of the patches and, after the last exposure, at 1, 24, 48 and 72 h. The score for the evaluation of the results is based on two clinical parameters, the erythema and eschar formation and the oedema formation. The reaction is graded as follows: a) Erythema and eschar formation: no erythema 0; slight erythema 1; well defined erythema 2; moderate erythema 3; and severe erythema – slight eschar formation 4; b) Oedema: no oedema 0; slight oedema 1; well defined oedema 2; moderate oedema 3; and severe oedema 4. Another test is for intracutaneous reactivity, that is the assessment of the potential of the material to produce irritation following intradermal injection of extracts. For this test healthy young albino rabbits are used. The back of the animal is shaved on both sites of the spine one day before the injections. Intracutaneous injections are divided into four groups, each of them of five injections: 1) extracts in polar solvent; 2) polar solvent alone; 3) extracts in non–polar solvent; 4) non-polar solvent alone. The appearance of each site must be noted immediately after the injection and at 24, 48, 72 h. Tissue reaction is graded for erythema and oedema according to the above described score.
16.5.2 Sensitization tests Sensitization tests are performed in guinea pigs. They include the Magnusson– Kligman maximization test and the Buehler closed-patch test. Only one of these methods is required; the preferred method is the maximization test. This test requires intradermal injections; therefore, if the material is not obtained in an injectable form, other methods can be used, such as the closed patch test. A positive test does not exclude the material from use; it merely indicates the need for further data to evaluate the risk of exposure
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to humans. The sensitizing capability of substances is defined as the minimum concentration required to induce an assigned level of sensitization. The maximization sensitization procedure includes preliminary tests and a main test that is divided into an intradermal induction phase, a topical-induction phase and a challenge phase. The preliminary tests are performed to determine the concentration of the test materials to be used in the main test. In the intradermal induction phase of the main test, a pair of 0.1 ml intradermal injections of each of the following: a) a mixture of complete Freund’s adjuvant (water-in-oil emulsion with killed mycobacteria) and the chosen solvent (50 : 50 volume ratio); b) the test material at the concentration used in the preliminary test; c) the test material emulsified in a 50/50 mixture of Freund complete adjuvant and the solvent. The topical induction phase is performed seven days after the end of the intradermal induction phase. The test material is applied in the intrascapular region of the animal, using filter paper patches, secured with occlusive dressing. The dressing and the patches are removed after 48 h. In the control animals, the solvent alone is used. After 14 days, the challenge phase is carried out. The test material is administered by topical application at the selected concentration on one flank of each animal using soaked patches. The patches are secured with an occlusive dressing, and are removed after 24 h. The challenge skin sites are observed in test and control animals at 24, 48 and 72 h after the removal of the dressing. The skin reaction is evaluated for erythema and edema according to the above presented grading scale. The closed patch test is performed on young adult albino guinea pigs. A minimum of 10 animals are treated with the test material, and a minimum of five animals are used as control. For topical application, a patch of appropriate dimensions is soaked with the test material or with an extract. The patch is applied to the skin under an occlusive dressing for 6 h. Preliminary tests are performed to determine the appropriate concentration of the test material to be used in the main test. For this purpose, four different concentrations of the test material or of a material extract are applied to the flanks of at least three animals using patches. After 6 h, the patches are removed and the application sites are assessed for erythema and edema according to the scoring system presented above. The main test includes an induction phase in which the test material or extract is applied on the back of the animals using appropriate patches soaked at the scheduled concentrations. Occlusive dressings and patches are removed after 6 h. This procedure is repeated for three weeks at weekly intervals. The challenge phase starts 14 days after the last application. The test material is applied using patches by a single topical application to an intact area of the skin. Dressings and patches are removed after 6 h and, 24 h after the primary challenge, the animals are depilated or shaved. The test sites are assessed a minimum of 2 h after the removal of fur and the topical reaction is graded
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according to the score for erythema and oedema; the evaluation is repeated 48 h after the removal of the challenge patch.
16.6
Systemic toxicity
Systemic toxicity is related to the possible adverse effects that can derive from the absorption, distribution or metabolism of leachates that originate from a material or device, involving parts of the body or organs with which they are not in direct contact.12 The tests are scheduled to determine systemic adverse effects that can be observed after administration of a single dose or multiple doses of a test sample per day during a period of 24 hours (acute toxicity), or by repeated exposures for a prolonged time (subacute, subchronic and chronic systemic toxicity). ISO 10993-11 requires the test for acute systemic toxicity for all the device categories exposed to blood contact. The test sample can be administered intravenously or intraperitoneally, but also dermal and implantation studies can be considered. ISO 10993-11 states that there are not absolute criteria for selecting a particular animal species. Mouse or rats are preferred, with the option of rabbits for dermal and implantation studies. The precision of studies on systemic toxicity is related to the number of animals per dose group. With respect to the controls to be performed, they comprise monitoring of body weight and food water consumption of the animal, clinical observation to detect toxic reactions, haematology and clinical chemistry analysis, and pathologic observations (gross and histological) on organs and tissues.
16.6.1 Acute systemic toxicity These investigations provide general information on the risk that can derive from an acute exposure to the material.13 They can be the first step in establishing the appropriate dosages for further subacute or subchronic studies. Usually a rodent species is used for these investigations. Animals in the control group must be managed in an identical way to that of the test group, with the exception of the exposure to the material to be tested. The period of observation for an acute systemic toxicity study can be of at least 3 days or longer if necessary. The animals are observed for evidence of changes in the skin and fur, in the eyes and mucous membranes, in the respiratory, circulatory and nervous system, and in the motion and behaviour.
16.6.2 Subacute, subchronic and chronic toxicity The most common exposure to many medical devices is characterized by continuous or repeated contacts with the body. The effects that can derive
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may therefore be related to the accumulation of chemicals in tissues or other mechanisms that it is important to evaluate by long-term studies. The determination of subchronic and chronic toxicity is carried out after the collection of information on acute toxicity. These tests give information about the target organs and the possibility of toxic accumulation. In subchronic or chronic toxicity studies, the material to be tested is administered daily for a period of 3 to 6 months. A single dose level is often adequate. The study of systemic toxicity with repeated exposures provides information on potential toxic effects, target organs, and the reversibility of the effects before clinical use of a device. Rodents such as mice or rats are used for these studies. The dose used for a systemic toxicity test must be defined after a preliminary risk assessment, depending on the clinical application of the device. For studies of longer duration, it is advisable to schedule at least three dose levels, with appropriate numbers of control animals. Besides the above listed controls on the animals, ISO 10993-11 indicates that full histopathology must be performed on organs and tissues. Lungs, liver and kidney are routinely examined. Further histopathological examination may be carried out in those organs where there is evidence of lesions.
16.6.3 Pirogenicity tests Pyrogens, substances that cause a febrile reaction, can be present in a device and, therefore, leachates of such materials can cause febrile responses (material-mediated pyrogenicity). Assessment of material-mediated pyrogenicity can be included in biocompatibility tests. ISO 10993-11 recommends the evaluation of the pyrogenicity potential of extractable substances derived from material leaching and includes the pyrogen test in the category of systemic toxicity testing. The baseline temperature of the rabbits before injection provides a basis for comparison. Material-mediated pyrogenicity may be determined by many factors, as endogenous pyrogens, (i.e. interleukines), prostaglandin, inducers such as polyadenilic and polyuridic acids, drugs active on thermoregulatory centers (such as cocaine and morphine), bacterial exotoxins, neurotransmitters and some metal salts, such as nickel. To evaluate material-mediated pyrogenicity, the rabbit pyrogen test is recommended. The test is performed by injecting extracts of the substance into rabbits and recording the temperature for a scheduled period of time.
16.7
Genotoxicity, carcinogenicity, reproductive and development toxicity
Genotoxicity (mutagenicity), carcinogenicity, and reproductive and development toxicity are considered in the same ISO 10993-3 standard.14
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16.7.1 Genotoxicity The possible mutagenic activity of materials can be considered and investigated, according to ISO 10993-3 for all the devices with prolonged (>30 days) body contact. The test can be performed both in vitro and in vivo. The simplest and sensitive in vitro assays are those involving bacteria and cultured mammalian cells. Only if scientifically indicated and if in vitro test results indicate potential genotoxicity are the in vivo tests performed. Two procedures are available: a search for chromosomal aberration in bone marrow or peripheral blood cells of rodents, and the analysis of chromosomes in metaphase cells.
16.7.2 Carcinogenicity These tests explore the carcinogenetic potential of materials to either a single or multiple exposures over a period of the total life-span of the test animal. These tests may require a prolonged period of observation, and can be designed to evaluate in a single experimental study both carcinogenicity and chronic toxicity. They should be considered for a device that will have permanent contact (longer than 30 days) with tissues. According to the ISO rules ‘Carcinogenicity tests should be conduced only if there are suggestive data from other source’ and ‘should be appropriate for the route and duration of exposure or contact’. The number of animals used for these studies is high. ISO 10993-3 also refers to the American Society for Testing and Materials document ASTM F 1439-92: ‘Performance of lifetime bioassay for tumorigenic potential of implanted materials’, in which a minimum of 60 male and 60 female rodents are scheduled for both the treatment and the control group. The final evaluation is based on the gross and histopatologic evaluation of the organs and tissues.
16.7.3 Reproductive and development toxicity These tests are used to evaluate the potential effects of the material/device on reproductive function, embryonic development and prenatal and early postnatal development. These tests should be performed only when the device has a potential impact on the reproductive ability of the subject, so they are mostly indicated for intrauterine devices and resorbable or leachable materials and devices.
16.8
Hemocompatibility
The effects of the interaction of materials with blood can be both at cellular and plasmatic levels. They include the mechanisms of coagulation, throm-
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bosis, embolism, inflammation, the activation of complement, and fibrinolytic systems. The cellular and humoral pathways are closely connected and reciprocally influenced. The hemocompatibility tests are grouped into five categories: thrombosis, coagulation, platelets, hematology, immunology. Standardization requires an evaluation of hemocompatibility for any medical device that has direct or indirect contact with circulating blood. Guidelines are presented in ISO 10993-4, ‘Selection of tests for Interactions with blood’.15 The tests required are related to the type of blood contact of the material or device (external communicating devices – blood path indirect, external communicating devices – circulating blood, and implant devices). For each contact category, primary and optional tests are recommended from a list of general test categories: thrombosis, coagulation, platelet function, hematology, and immunology. Hemocompatibility tests can be performed both in vitro16 and in vivo.17 In vitro studies allows a first screening, but relevant informations are collected in vivo after the implant of the device (i.e. vasculard grafts, heart valves) into the vascular system for an assigned period. After the removal of the device, it can be examined for the detection of signs of thrombus formation on its surface. In some instances it is possible to apply ex vivo methods, avoiding the sacrifice of animals. In this case, the blood is diverted from an artery into a shunt, in which the test device is placed, and returns to an animal vein (closed systems) or is collected into a chamber in which it contacts the test material (open systems). Coagulation tests are performed in this way.18
16.9
Tests for local effects after implantation
These tests are described in part 6 of ISO 1099319 and are designed to determine the local effects of an implant material on living tissue, at both macroscopic and microscopic level. The tests requires that the specimen is implanted into an appropriate site or tissue and the implant is not intended to be subjected to mechanical or functional loading; local effects are evaluated in comparing the tissue response caused by materials used for clinically accepted medical devices. The placement of a material or device in the tissues by injection, insertion or surgical implantation elicits a host response that is dynamic and changes with time.20 Surgical technique must be applied carefully to limit tissue damage, avoiding excessive pressure on tissues, bleeding and dead spaces. An inaccurate technique can lead to migration or mobility of the implant. When the implant is positioned in the back of the animal, to prevent the displacement of samples in the postoperative period it is recommended to avoid the handling of the dorsum skin.21 After the healing of the implanted site, tissues are explanted and examined for macroscopic and microscopic responses. Several macroscopic and
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microscopic parameters are considered, including fibrosis, degeneration, presence of phagocytic cells, necrosis, fatty infiltration, and foreign debris. The scheduled time for short-term implants is less than 12 weeks. For long-term implants, the scheduled time is more than 12 weeks. In the evaluation of the local effects of an implant, the surgical procedure required for placement of the sample must also be considered and also the reparative process that follows the surgical intervention. These conditions relate to the early period after implantation. Short-term tests (up to 12 weeks) can be performed in mice, rats, guineapigs and rabbits with implants in subcutaneous tissue and muscle. The scheduled implantation periods are of 1, 3, 4, 9, 12 weeks. For long-term studies in subcutaneous tissue, muscle and bone, rats, guinea-pigs, rabbits, sheep, goats and pigs are the considered animal species. The implantation periods are 12, 26, 52, 78 and 104 weeks. Subcutaneous implants are placed in pockets obtained in the dorsal subcutaneous tissue22,23 through a midline incision in at least three animals and in a sufficient site number to place 10 samples for each material and experimental time. For muscle implants, the paravertebral muscles of rabbits are the preferred implant sites. The use of a minimum of three rabbits per time period is recommended with sufficient test and control implants to provide at least eight implant specimens for each implantation period. Bone implantation is required when the material is proposed for the production of orthopaedic devices. The seats of implant can be the diaphisis (cortical bone) (Fig. 16.1) or the epiphysis (cancellous bone) of long bones.24
16.1 Screw implants positioned in the cortical bone; sheep femoral diaphysis, operatory view.
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The samples can be screw shaped or cylindrical. The implant size is based on the size of the animal species selected for the tests and of the bone in which the implant is placed. Cylindrical implants in rabbits are 2 mm in diameter and 6 mm long, while those in sheep and goats are 3 mm in diameter and 12 mm long. Screw type implants in rabbits, sheep, goats and pigs vary from 2 to 4 mm. Equivalent sites of implant will be used for test and control specimens using the controlateral bones. In rabbits, a maximum of 6 implants (three tests and three controls) are placed. In other species (sheep, goat, pig), there is a maximum of 12 implant sites (6 test and 6 controls). Femur and tibia are the bone segments used for the implants.25,26 The preparation of holes in which the implants are placed is performed using a low drilling speed with profuse saline irrigation. Cylindrical specimens are inserted by simple pressure. When screw-shaped specimens are inserted, the use of a device for the record of the insertion torque is recommended. There are, however, other possible designs for the samples, according to the nature of the material to be tested. For the evaluation of cement composites, discs of the material implanted subcutaneously and in cranial defects were used;27 disc shaped samples were also employed to evaluate the buccal vestibular tissue reaction to some materials for dentures.28 For histological assessment, the parameters given in Table 16.2 are considered. For bone implants, the interface between the tissue and the material is of particular interest. Osteointegration is defined as the direct structural and functional connection between living bone and the surface of the implant. The specimens containing implants must be processed for the histology following the procedure for undecalcified bone. After fixation in 4% buffered paraformaldehyde and dehydration in a graded series of alcohols, they are included in methyl methacrylate. To obtain sections of the required thickness, the most useful procedure is by a cutting–grinding system.29 This system comprises a cutting device based on a diamond saw and a rotatable clamp on which the sample is hold. The thick slices so obtained are glued
Table 16.2 Parameters for histological assessment of local effects after implantation 1. Extent of fibrosis, fibrous capsule and inflammation 2. Degeneration: changes in tissue morphology 3. Presence, number and distribution of the inflammatory cells in relation with the distance from the material/tissue interface 4. Presence of necrosis 5. Material debris, fatty infiltration, granuloma 6. For porous materials, quality and quantity of tissue in-growth 7. For bone implants, study of the bone implant interface
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on a Plexiglas slide, and then thinned to the required thickness by using a grinding system with a graded series of abrasive paper. To evaluate osteointegration, the histological observation of samples is performed to measure some morphological parameters. The histomorphometry measures should be obtained with an image analysis system based on the connection of the microscope with a personal computer equipped with appropriate image analysis software.30 The most commonly used parameters evaluated with histomorphometry are the Affinity Index (the length of bone directly opposed to the implant / the total length of bone implant interface ×100%) the bone in-grow (the area of the bone tissue grown into the screw thread /area between the screw thread),31,32 and the Bone Mirror Area that indicates the quality of bone in the region closely contiguous to the implant (Fig. 16.2).33,34 The extent of fibrous tissue between bone and the implant, if present, is also be evaluated. The parameters related to the quality of cancellous bone are evaluated according to Parfitt:35 among them the most significant are trabecular bone volume, BV/TV (%), defined as the whole bone area, expressed as a percentage of the total tissue area in the sampling site and converted to a volume; the trabecular thickness, Tb.Th (μm), the trabecular number, Tb.N (mm−1) and the trabecular separation, Tb.Sp (μm2) that were calculated from BV/TV and the perimeter of trabeculae using Parfitt formulae.36 The rate of bone deposition and mineralization is evaluated by dynamic histomorphometry. For this purpose, fluorochromes (i.e. tetracycline, aliza-
16.2 Histological appearance of vertebral pedicular implant of an hydroxyapatite coated stainless-steel screw in sheep lumbar spine. The penetration of bone (bone in-growth) into the threads of the screw is evident. (Fast Green, 2.5x)
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rin red, calcein blue, xylenol orange) are administered to the animal at scheduled times before the sacrifice. The fluorescence seen on histological sections under microscopic observation with appropriate filters marks the new deposition of bone (Fig. 16.3).37 It seems be useful to remember that, beside these investigations, more recently it is possible to perform computerized microtomography on the explanted parts; this technique is not destructive for the samples the time required to capture the images is short, and the same sample can therefore be used afterwards for the histology and histomorphometry. More detailed information about this technique is now given. Computerized microtomography is a non-destructive technique that is very useful for a three-dimensional study of bone and biomaterials because it is able to supply structural and densitometric information. Microtomography is based on the same principles of a normal x-ray tomography used in clinical medicine but avails itself of a microscopic level resolution (Fig. 16.4a, b). This type of equipment allows images to be obtained of the internal structure of a small object with high spatial resolution and high speed. The latest generation of this type of system is able to reach a spatial resolution of 5 μm corresponding to near 1 × 10−7 mm3 voxel size. In this way, we can visualize particulars of an object down to a size of 1 μm. The typical CT scanners used in medicine have a spatial resolution in the range of 1–2.5 mm, which corresponds to 1–10 mm3 voxel size. The x-ray tomography technique allows a 3D object to be visualized and its internal structure analysed without sample preparation or chemical fixation. The evolution of image analysis techniques with specialized hardware and software has been useful in studying the images obtained with
16.3 Dynamic histomorphometry is obtained labeling the bone with fluorochromes (i.e. tetracyclines). The fluorescence bands observed on histological sections marks the deposition of new bone.
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(a)
(b)
16.4 (a) Section of a rat lumbar vertebra acquired with micro-CT (image pixel of 4 μm.) (b) Coronal and sagittal sections of the same lumbar vertebra and their three-dimensional positions.
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16.5 Porosity of a polymeric biocomposite. Image acquired with micro-CT and analyzed with a CTAn software for image analysis to evaluate the pore size.
micro-CT (Fig. 16.5) because they have a large applicability, an high calculation intensity and a rigorous statistic approach. For these reasons micro-CT is an important and valid non-destructive method especially for the preclinical evaluation of implant materials. It is also extremely helpful when it comes to characterizing devices in the preimplant phase and evaluating possible deformations and/or degradations in the explant phase. It is therefore useful for the analysis of both tissue engineering scaffolds and bone tissue regeneration after the preclinical application of innovative biocomposites (Fig. 16.6 and 16.7). Micro-CT could be used to measure bone characteristics in some pathologies such as osteoporosis and primitive and metastatic tumors (lung, prostate, multiple myelome), to study medicines that operate on bone remodeling, to research and produce new porous biomaterials and to study tissue engineering scaffolds. Compared with classical histology micro-CT has some advantages such as maintenance of sample integrity, reduced time for image acquisition, a large improvement of analysable sections and the possibility of three-dimensional analysis (Fig. 16.8).
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16.6 Bone growth inside a biocomposite scaffold analyzed with a micro-CT system.
16.10 Biocompatibility evaluation in pathological conditions The use of healthy animal models for the in vivo testing of biocompatibility is a currently used procedure, but is not sufficient for a complete evaluation of materials or devices. The clinical situation in which the final product is placed is often characterized by a morbid condition that may affect the outcome. This is particularly evident when the material is intended for use in the bone apparatus. For this reason, experimental models were developed to obtain pathological conditions of clinical relevance. One of the most frequent of such disease is osteoporosis, that is related to age, post-menopausal status, or is secondary to systemic diseases or pharmacological therapies. Osteoporotic animal models are today widely used in the preclinical studies of biomaterials. Bilateral ovariectomy is the better surgical procedure to reproduce the clinical situation of post-menopausal osteoporosis. The female rat and sheep are the species most suitable for this purpose. Bone implants of the testing materials were placed in bone 16 weeks after ovariectomy in rats.38 In sheep, the development of an osteoporotic state was still present 12 months after ovariectomy, and was more evident after 24 months (Fig. 16.9).39 Besides the long bones such as femur and tibia, the sheep model is suitable also for studies on the spine, where the decrease of trabecular and
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16.7 Maximal intensity projection of cortical bone with three biocomposite implants in rabbit tibia.
cortical bone mass was observed in the vertebral bodies.40,41 It is necessary to bear in mind that the studies on small animals such as the rat are acceptable only in the first phase of a research, because of the differences between rat bone quality compared with humans. Large animal models, such as sheep, should be adopted in future developments, for a better approach to the clinical situation. Figure 16.10 shows the histological appearance of implants in cortical bone of titanium samples in femur diaphysis of sheep, in healthy and osteoporotic conditions. The bone-to-implant contact is clearly higher in healthy bone. In Fig. 16.11 bioglass implants in trabecular bone of rat femoral condyle, in healthy and osteoporotic conditions can be observed. Healthy bone shows a good bone-to-implant contact, whereas in
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16.8 3D model of a scaffold implant in a condyle of a rabbit femur obtained from micro-CT images.
osteoporotic bone a connective fibrous layer is interposed between the material and bone trabeculae. The use of pathological bone-derived cell culture for in vitro tests,42 in association with pathological animal models allows the evaluation of the different response of bone to candidate orthopaedics materials by comparing healthy and pathological conditions.
16.11 Biofunctionality The ability of a device to perform its functions after the implantation in the living body and to maintain them in the course of time is called biofunctionality. The biological environment is aggressive and the in vivo performance of a final product can be different from that expected when only laboratory tests are perforned. Laboratory simulations with appropriate machines can mimic the mechanical stress to which the devices such as joint prosthesis and heart valves are submitted, but no simulation test is able to repeat the behaviour in the body over a long time span. For this reason,
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(a)
(b)
16.9 Bone biopsies of the iliac crest in (a) young healthy sheep and (b) osteoporotic ovariectomized sheep. The osteopenic condition is still present 12 months after ovariectomy. (Fast Green, 4x)
after the laboratory tests, the in vivo implantation of the device is required to collect more complete information. Biofunctionality must be evaluated using models that are as much as possible similar to the final destination of the device.43 Accepting that no ideal animal analogue of man is available for preclinical testing, large species as pig and sheep or goats are considered as suitable for this purpose.44
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16.10 Histological appearance of implants in cortical bone (titanium), sheep femur diaphysis, in (a) healthy and (b) osteoporotic conditions. The bone-to-implant contact is clearly higher in healthy bone. (Fast Green, 4x)
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16.11 Histological appearance of implants in trabecular bone (bioglass), rat femoral condyle, in (a) healthy and (b) osteoporotic conditions. Healthy bone shows a good bone- to-implant contact, whereas in osteoporotic bone a connective fibrous layer is interposed between the material and bone trabeculae. (Fast Green, 4x)
16.12 References 1. european society of biomaterials conference symposium, Sorrento, Italy, 2005. 2. williams dh. General concepts of biocompatibility. In Black J and Hastings G. (eds) Handbook of biomaterial properties, Chapman & Hall, London, 1998.
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3. zhang y, hao l, savalani mm, harris ra, di silvio l, tanner ke. In vitro biocompatibility of hydroxyapatite-reinforced polymeric composites manufactured by selective laser sintering. J Biomed Mater Res A, 2008, Dec 23. Epub ahead of print. 4. mattioli-belmonte m, giavaresi g, biagini g, virgili l, giacomini m, fini m, giantomassi f, natali d, torricelli p, giardino r. Tailoring biomaterial compatibility: in vivo tissue response versus in vitro cell behavior. Int J Artif Organs, 2003, 26(12), 1077–1085. 5. smith ac, fasse rt, swindle mm. Ethics and regulations for the care and use of laboratory animals. In: An YH, Friedman RJ. (eds) Animal models in orthopaedic research, CRC Press, 1999. 6. repubblica italiana gazzetta ufficiale, January 27, 1992. 7. iso 10993 Biological evaluation of medical devices Part 2 Animal welfare requirements, 2006. 8. iso 10993 Biological evaluation of medical devices Part 1 Evaluation and testing, 2004. 9. gatti am, knowles jc. Biocompatibility and biological tests in R Barbucci (ed) Integrated biomaterial science, Kluwer Academic/Plenum Publishers, New York, 2002. 10. iso 10993 Biological evaluation of medical devices Part 12 Sample preparation and reference materials, 2007. 11. iso 10993 Biological evaluation of medical devices Part 10 Tests for irritation and delayed type hypersensitivity, 2002. 12. iso 10993 Biological evaluation of medical devices Part 11 Tests for systemic toxicity, 2006. 13. ren j, zhao p, ren t, gu s, pan k. Poly (d,l-lactide)/nano-hydroxyapatite composite scaffolds for bone tissue engineering and biocompatibility evaluation. J Mater Sci Mater Med, 2008, 19(3), 1075–1082. 14. iso 10993 Biological evaluation of medical devices Part 3 Tests for genotoxicity, carcinogenicity and reproductive toxicity, 2003. 15. iso 10993 Biological evaluation of medical devices Part 4 Selection of tests for interactions with blood, 2002. 16. seyfert ut, biehl v, schenk j. In vitro hemocompatibility testing of biomaterials according to the ISO 10993–4. Biomol Eng, 2002, 19(2–6), 91–96. 17. arabi h, mirzadeh h, ahmadi sh, amanpour s, rabbani s, abdi a. In vitro and in vivo hemocompatibility evaluation of graphite coated polyester vascular grafts. Int J Artif Organs, 2004, 27(8), 691–698. 18. ambrosio l, peluso g, davis p. Biomaterials and their biocompatibilities in Wise DL Ed. Human biomaterials application, Humana Press, Totowa, NY, 1996. 19. iso 10993 Biological evaluation of medical devices Part 6 Tests for local effects after implantation 2007. 20. anderson gm. Soft tissue response. In: Black J and Hastings G (eds) Handbook of biomaterial properties, Chapman & Hall, London, 1998. 21. jansen ja. Animal Models for studying soft tissue biocompatibility of biomaterials. In: An YH, Friedman RJ. (eds) Animal models in orthopaedic research, CRC Press, 1999. 22. pesáková v, smetana k jr, balík k, hruska j, petrtýl m, hulejová h, adam m. Biological and biochemical properties of the carbon composite and polyethylene implant materials. J Mater Sci Mater Med, 2003, 14(6), 531–537.
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23. yao ch, liu bs, hsu sh, chen ys, tsai cc. Biocompatibility and biodegradation of a bone composite containing tricalcium phosphate and genipin crosslinked gelatin. J Biomed Mater Res A, 2004, 69(4), 709–717. 24. giavaresi g, ambrosio l, battiston ga, casellato u, gerbasi r, fini m, nicoli aldini n, martini l, rimondini l, giardino r. Histomorphometric, ultrastructural and microhardness evaluation of the osseointegration of a nanostructured titanium oxide coating by metal-organic chemical vapour deposition: an in vivo study. Biomaterials, 2004, 25(25), 5583–5591. 25. oudadesse h, derrien ac, mami m, martin s, cathelineau g, yahia l. Aluminosilicates and biphasic HA–TCP composites: studies of properties for bony filling. Biomed Mater, 2007, 2(1), S59–S64. 26. giavaresi g, fini m, chiesa r, giordano c, sandrini e, bianchi ae, ceribelli p, giardino r. A novel multiphase anodic spark deposition coating for the improvement of orthopedic implant osseointegration: an experimental study in cortical bone of sheep. J Biomed Mater Res A, 2008, 85(4), 1022–1031. 27. ruhé pq, hedberg el, padron nt, spauwen ph, jansen ja, mikos ag. Biocompatibility and degradation of poly(dl-lactic-co-glycolic acid)/calcium phosphate cement composites. J Biomed Mater Res A, 2005, 74(4), 533–544. 28. ebadian b, razavi m, soleimanpour s, mosharraf r. Evaluation of tissue reaction to some denture-base materials: an animal study. J Contemp Dent Pract, 2008, 9(4), 67–74. 29. donath k, brenner g. A method for the study of undecalcified bone and teeth with attached soft tissue. J Oral Pathol, 1982, 11, 318. 30. nicoli aldini n, fini m, giavaresi g, torricelli p, martini l, giardino r, ravaglioli a, krajewski a, mazzocchi m, dubini b, ponzi-bossi mg, rustichelli f, stanic v. Improvement in zirconia osseointegration by means of a biological glass coating: an in vitro and in vivo investigation. J Biomed Mater Res, 2002, 61(2), 282–289. 31. ramires pa, wennerberg a, johansson cb, cosentino f, tundo s, milella e. Biological behavior of sol–gel coated dental implants. J Mater Sci Mater Med, 2003, 14(6), 539–545. 32. giavaresi g, chiesa r, fini m, sandrini e. Effect of a multiphasic anodic spark deposition coating on the improvement of implant osseointegration in the osteopenic trabecular bone of sheep. Int J Oral Maxillofac Implants, 2008, 23(4), 659–668. 33. slotte c, lundgren d, sennerby l, lundgren ak. Influence of preimplant surgical intervention and implant placement on bone wound healing. Clin Oral Implants Res, 2003, 14(5), 528–534. 34. tavares mg, de oliveira pt, nanci a, hawthorne ac, rosa al, xavier sp. Treatment of a commercial, machined surface titanium implant with H2SO4/ H2O2 enhances contact osteogenesis. Clin Oral Implants Res, 2007, 18(4), 452–458. 35. parfitt am. Bone histomorphometry: proposed system for standardization of nomenclature, symbols, and units. Calcif Tissue Int, 1988, 42(5), 284–286. 36. parfitt am, drezner mk, glorieux fh, kanis ja, malluche h, meunier pj, ott sm, recker rr. Bone histomorphometry: standardization of nomenclature, symbols and units. J Bone Miner Res, 1987, 2, 595–610. 37. fini m, giavaresi g, aldini nn, torricelli p, botter r, beruto d, giardino r. A bone substitute composed of polymethylmethacrylate and α-tricalcium
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phosphate: results in terms of osteoblast function and bone tissue formation. Biomaterials, 2002, 23, 4523–4531. fini m, giavaresi g, torricelli p, krajewski a, ravaglioli a, belmonte mm, biagini g, giardino r. Biocompatibility and osseointegration in osteoporotic bone. J Bone Joint Surg Br, 2001, 83(1), 139–143. fini m, giavaresi g, torricelli p, borsari v, giardino r, nicolini a, carpi a. Osteoporosis and biomaterial osteointegration. Biomed Pharmacother, 2004, 58(9), 487–493. fini m, giavaresi g, greggi t, martini l, nicoli aldini n, parisini p, giardino r. Biological assessment of the bone-screw interface after insertion of uncoated and hydroxyapatite-coated pedicular screws in the osteopenic sheep. J Biomed Mater Res A, 2003, 66(1), 176–183. nicoli aldini n, fini m, giavaresi g, giardino r, greggi t, parisini p. Pedicular fixation in the osteoporotic spine: a pilot in vivo study on long-term ovariectomized sheep. J Orthop Res, 2002, 20(6), 1217–1224. torricelli p, fini m, giavaresi g, giardino r. In vitro models to test orthopedic biomaterials in view of their clinical application in osteoporotic bone. Int J Artif Organs, 2004, 27(8), 658–663. ignatius aa, betz o, augat p, claes le. In vivo investigations on composites made of resorbable ceramics and poly(lactide) used as bone graft substitutes. J Biomed Mater Res, 2001, 58(6), 701–709. donati d, di bella c, lucarelli e, dozza b, frisoni t, nicoli aldini n, giardino r. OP-1 application in bone allograft integration: preliminary results in sheep experimental surgery. Injury, 2008, 39 Suppl 2, S65–S72.
17 The mechanics of biocomposites L. N I C O L A I S, University of Naples ‘Federico II’, Italy; and A. G L O R I A and L. A M B R O S I O, National Research Council, Italy
Abstract: An overview is presented of the theoretical approach to the design of polymer composite materials. Composite materials with polymeric matrices are strong candidates for structural applications where high strength and stiffness to weight ratios are required. The fundamental analysis of the mechanical response of composite materials involves investigation on two levels, micro- and macro-scale. This approach encompasses micromechanics and macromechanics, extends to lamination theory and makes clear the physical significance of the basic concepts of composite material design. First, the mechanical behaviour is described of a unidirectional fibre-reinforced lamina as the basic building block of a laminate, showing stress–strain relationships in various coordinate systems. The design of laminates having suitable properties for a particular application is explained and short fibrereinforced composites and particulate composites are examined. Finally, polymer nanocomposites are discussed as an interesting strategy to improve the mechanical properties of polymers. Key words: micromechanics; macromechanics; lamination theory; fibre-reinforced composites; particulate composites; polymer nanocomposites.
17.1
Introduction
The term ‘composite materials’ refers to the combination, on a macroscopic scale, of two or more materials, differing in composition or morphology, to obtain specific chemical, physical and mechanical properties. The constituents almost always differ chemically, and they are essentially insoluble in each other. The advantage enjoyed by the resulting composite materials is that they may exhibit a combination of the best properties of their constituents, and often some beneficial properties which are not shown by the constituents. (Jones, 1999; Mallick, 1997; Schwartz, 1992). Over the past years, composite materials with polymeric matrices have emerged as suitable candidates for load-bearing structural applications in several fields. However, unlike the metals which the composites are replacing in structural applications, their processing characteristics and their final properties are strongly affected by the chemical composition, the load conditions and the various environments in which they perform. 411
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More recently, fibre-reinforced polymer (thermoplastic or thermoset) matrices have attracted much attention, since they show high strength and stiffness-to-weight ratios. However, their application in various industrial sectors (such as biomedical and construction) has extended interest in their use also to particulate reinforcing polymer composites and, lately, to nanocomposites. The polymer-based matrix, and its interaction with a reinforcing phase in the form of continuous or discontinuous high stiffness fibres and particles, provide one of the major controlling factors in a composite’s properties. A wide variety of polymeric materials, both thermoplastic and thermosetting, can be selected to meet specific performance characteristics. Thus, the identification of property and processing requirements of the material to be used as matrix in composites is important at the design stage. Many studies and uses have highlighted that the polymeric matrix plays a crucial role in composite performance. The polymeric matrix and its interaction with reinforcing phase, in the form of continuous, discontinuous fibres and particles, is one of the major controlling factors in the processing and properties characteristics of composites. Recent, more demanding applications of composites have required the polymeric matrix to make an increasingly important contribution to composite performance (Jones, 1999; Seferis and Nicolais, 1983). In these composites, the rôle of the matrix is variously to transfer stresses to the reinforcement, to provide a barrier for containing or excluding, to protect the surface of the fibres from mechanical abrasion, to be biocompatible, or to confer any other attribute demanded in a specific application. Composite materials are commonly divided into several classifications, including the following: ‘fibrous composite materials’ consisting of fibres embedded in a matrix; ‘laminated composite materials’, made up of layers of composite materials; ‘particulate composite materials’ which are composed of particles in a matrix; and combinations of some of the preceding examples (Fig. 17.1). Various polymer composites have also been investigated for various biomedical applications (Table 17.1) Specific advantages have been gained through using polymer composite biomaterials (Ramakrishna et al., 2001), for which a recently coined alternative term is ‘biocomposites’. Polymer composite materials represent an alternative choice to overcome drawbacks related to homogeneous materials. For example, it is well known that one of the major problems in orthopaedic surgery is the mismatch of stiffness between the bone and metallic or ceramic devices, which may negatively affect the bone remodelling and healing process. Thus, implants should be biocompatible as well as characterized by suitable mechanical properties. In this respect, polymer composite biomaterials have
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17.1 Schematic representation of fibre-reinforced and particulate composites: (a) continuous and aligned fibres; (b) discontinuous and aligned fibres; (c) discontinuous and randomly oriented fibres; (d) particles.
attracted much attention because of their modifiable properties, which may be tailored to match those of the host tissues. The composites concept has been applied also to the design of scaffolds for tissue engineering, with the aim of improving their functionality (Guarino et al., 2007). In this context, this chapter will highlight the basic principles in designing composite materials. The mechanical properties and behaviour of a lamina will be analyzed with reference to its constituents (matrix and fibre). Since an individual lamina can be considered the basic building block of a laminate, the knowledge of its stress–strain relations is crucial to the understanding of laminated fibre-reinforced structures.To this end, micromechanics, macromechanics and lamination theory will be considered. The chapter will also overview some approaches related to short-fibre composites and particulate composites. As advanced technologies expand, so the need for novel functional materials significantly increases. These materials typically possess materials with a specific combination of properties designed to meet the demands of a particular application. Then we will discuss polymer nanocomposites as a strategy to improve the mechanical properties of polymer systems, and show also a three-phase model which could be used to describe the tensile modulus of these materials, taking into consideration the crucial role of the interface.
17.2
Basic concepts in designing composite materials: lamina and laminate properties
Studies of composite materials emphasize the relationship of structural performance to the properties of a ply. A ‘ply’ (also defined as ‘lamina’ or ‘layer’) is a thin sheet of material consisting of an oriented array of fibres
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Table 17.1 Various applications of different polymer composite biomaterials (Ramakrishna et al., 2001) Applications
Polymer composite biomaterials
External fixation Bone plates and screws
CF/epoxy CF/PEEK, CF/epoxy, CF/PMMA, CF/PP, CF/PS, CF/PLLA, CF/PLA, KF/PC, HA/PE, PLLA/PLDLA, PGA/PGA CF/UHMWPE, UHMWPE/UHMWPE PET/PU, PTFE/PU, CF/PTFE, CF/C Bone particles/PMMA, titanium/PMMA, UHMWPE/PMMA, GF/PMMA, CF/PMMA, KF/PMMA, PMMA/PMMA, Bioglass/bis-GMA PET/PHEMA, KF/PMA, KF/PE, CF/PTFE, CF/PLLA, GF/PU CF/LCP, CF/PEEK, GF/PEEK CF/epoxy, CF/C, CF/PS, CF/PEEK, CF/PTFE, CF/UHMWPE, CF/PE, UHMWPE/UHMWPE PET/SR, CF/UHMWPE PET/PU, PET/collagen
Total knee replacement Cartilage replacement Bone cement
Tendon/ligament Intramedullary nails Total hip replacement Finger joints Abdominal wall prosthesis Vascular graft Spine cage, plate, rods, screws and disc Bone replacement material Dental implants Dental post Arch wire and brackets Dental bridges Dental restorative material
Cells/PTFE, Cells/PET, PET/collagen, PET/gelatin, PU/ PU-PELA CF/PEEK, CF/epoxy, CF/PS, Bioglass/PU, Bioglass/PS, PET/ SR, PET/hydrogel HA/PHB, HA/PEG-PHB, CF/PTFE, PET/PU, HA/HDPE, HA/ PE, Bioglass/PE, Bioglass/PHB, HA/PLA, Bioglass/PS CF/C, SiC/C CF/C, CF/epoxy, GF/polyester GF/PC, GF/PP, GF/nylon, GF/PMMA UHMWPE/PMMA, CF/PMMA, GF/PMMA, KF/PMMA Silica/bis-GMA, HA/2,2′(4-methacryloxydiethoxy phenyl) propane
BIS-GMA bis-phenol A glycidyl methacrylate, C carbon, CF carbon fibres, GF glass fibres, HA hydroxyapatite, KF Kevlar fibres, LCP liquid crystalline polymer, PC polycarbonate, PEA polyethylacrylate, PEEK polyetheretherketone, PEG polyethylene glycol, PELA block co-polymer of lactic acid and polyethylene glycol, PET polyethyleneterephthalate, PGA polyglycolic acid, PHB poly(hydroxyl butyrate), PHEMA poly(2-hydroxyethyl methacrylate), PLDLA poly(l-dl-lactide), PLLA poly(l-lactic acid), PMA polymethylacrylate, PMMA polymethylmethacrylate, PP polypropylene, PS polysulfone, PTFE polytetrafluoroethylene, PU polyurethane, SR silicon rubber, UHMWPE ultra-high-molecular-weight polyethylene.
embedded in a continuous matrix. These plies are stacked one upon another in a defined sequence and orientation, and bonded together yielding a laminate with tailored properties. The properties of the laminate are strongly related to those of the ply, namely specification of the ply thickness, stacking sequence, and the orientation of each ply. The properties of the ply are specified by the properties of the fibres and matrix, their volumetric concentration, and geometric packing.
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Generally, the ply material is obtained in a continuous tape or sheet form which is in a chemically semi-cured condition. The preparation of structural items involves using this ‘prepreg’ material and either winding it on to a mandrel or cutting and stacking it on to a mould after which heat and pressure or tension is applied to complete the chemical process (Nicolais, 1975). The basic building block of a laminate is a lamina which is a flat (or curved in a shell) arrangement of unidirectional fibres or woven fibres in a matrix. Consequently, the knowledge of the mechanical behaviour of a lamina is crucial to the understanding of laminated fibre-reinforced composites. The basis for engineering design of such a laminate is the properties of the cured ply or lamina. A unidirectional fibre-reinforced lamina may be treated as a thin two-dimensional item and it is mechanically characterized by its stress–strain response to loadings: 1. in the direction of the fibres (indicated as ‘1’) which exhibits a nearly linear response up to a fracture stress; 2. in the direction transverse to the fibres (labelled as ‘2’) which exhibits a significantly decreased moduli and strength values; 3. in plane shear load. Unlike metals, a unidirectional fibre-reinforced lamina, in the form of a thin sheet, is not isotropic. Thus, to specify its elastic properties in its natural orientation it requires four constants in the case of plane stress state (Jones, 1999; Mallick, 1997; Nicolais, 1975). In particular, for a unidirectional fibre-reinforced lamina in the 1–2 plane, a plane stress state may be defined by setting σ3 = τ13 = τ23 = 0. However, for orthotropic materials, imposing a plane stress state results in implied out-of-plane strain of ε3 ≠ 0 and γ13 = γ23 = 0. Accordingly, for orthotropic materials the strain-stress relation may be reduced to: ⎡ ε1 ⎤ ⎡ S11 ⎢ ε 2 ⎥ = ⎢ S12 ⎢⎣γ ⎥⎦ ⎢⎣ 0 12
S12 S22 0
0 ⎤ ⎡ σ1 ⎤ 0 ⎥ ⎢σ 2 ⎥ S66 ⎥⎦ ⎢⎣τ 12 ⎥⎦
S12 = −
ν12 ν = − 21 E1 E2
(17.1)
where S11 =
1 E1
S22 =
1 E2
S66 =
1 G12
(17.2)
E1, E2, G12 and v12 (= −ε2/ε1) are Young’s moduli, the shear modulus and the Poisson’s ratio, respectively, of the composite lamina, whilst 1 and 2 represent the principal material axes (Fig. 17.2). The strain–stress relations in equation (17.1) can be inverted in order to obtain the following stress–strain relations:
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1
θ
X
17.2 Fibre-reinforced lamina: principal material axes (1–2) and generic x–y axes.
⎡ σ 1 ⎤ ⎡Q11 ⎢σ 2 ⎥ = ⎢Q12 ⎢⎣τ ⎥⎦ ⎢⎣ 0 12
Q12 Q22 0
0 ⎤ ⎡ ε1 ⎤ 0 ⎥ ⎢ ε2 ⎥ Q66 ⎥⎦ ⎢⎣γ 12 ⎥⎦
(17.3)
where the Qij represent the elements of the reduced stiffness matrix [Q] for a plane stress state in the 1–2 plane, which can be obtained as the components of inverted compliance matrix [S] reported in equation (17.1). Accordingly, for the orthotropic lamina, these Qij elements are expressed as follows: Q11 =
E1 1 − ν12 ν 21
Q12 =
ν12 E2 ν 21 E1 = 1 − ν12 ν 21 1 − ν12 ν 21
Q22 =
E2 1 − ν12 ν 21
Q66 = G12 (17.4)
It may be noticed that there are four independent elastic constants (E1, E2, n12 and G12), if the equations (17.1) and (17.3) are considered taking into account the reciprocal relation:
ν12 ν 21 = E1 E2
(17.5)
From this, the above mentioned stress–strain relations are defined in terms of the principal materials coordinates for an orthotropic material (Jones, 1999). The ‘mechanics of materials approach’ provides a convenient means to determine the composite elastic properties. It is assumed that the composite is void free, the fibre–matrix bond is perfect, the fibres are of uniform size and shape and are spaced regularly, and the material behaviour is linear and elastic.
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The elastic constants of a lamina (E1, E2, n12 and G12) may be obtained by micromechanics considerations in terms of the moduli, Poisson’s ratios and volume fractions of the fibre (Vf) and the matrix (Vm), using the following equations: E1 = Ef Vf + EmVm Ef Em E2 = Ef Vm + EmVf ν 12 = ν f Vf + ν mVm Gf Gm G12 = Gf Vm + GmVf
(17.6)
In this context, another interesting approach is represented by Halpin– Tsai equations (Halpin and Tsai, 1969; Jones, 1999; Nicolais, 1975): E1 = Ef Vf + EmVm ν 12 = ν f Vf + ν mVm p 1 + ξηVf = pm 1 − ηVf
(17.7)
where ( pf pm ) − 1 ( pf pm ) + ξ p = E2 , G12 , ν 23 pf = Ef , Gf , ν f pm = Em , Gm , ν m
η =
(17.8)
and ξ is a measure of fibre reinforcement of the composite material that depends on the fibre geometry, packing geometry and loading conditions. It is worth noting that the expressions for E1 and n12 are the same generally accepted rule of mixtures results which are reported in equations (17.6). In particular, the Halpin–Tsai equations can be equally applied to fibre, ribbon or particulate composites, and the only difficulty in using them is related to the determination of an appropriate value of ξ also through the use of curve-fitting techniques. Thus, for example, as already discussed, ξ values for transverse modulus (ξE2) and for shear modulus (ξG12) are generally different functions of the cross-section aspect ratio (Jones, 1999; Nicolais, 1975). The above reported relations enable prediction of lamina properties which are defined as ‘apparent properties’, since they should be compared with those measured through mechanical tests. However, the prediction of these properties represents an important step in designing composites with specific macroscopic properties.
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If the lamina is rotated with respect to the applied stress or strain, a method of transforming stress–strain relations from one coordinate system to another has to be involved. For this reason, if a rotation of principal material axes from generic x–y axes is considered it is possible to recall transformation equations for expressing stresses (or strains) in a x–y coordinate in terms of stresses (or strains) in a 1–2 coordinate system, knowing the angle (θ) from the x-axis to the 1-axis (Fig. 17.2). 2 −2 mn ⎤ ⎡ σ 1 ⎤ n2 ⎡σ x ⎤ ⎡ m 2 ⎢σ y ⎥ = ⎢ n 2 2 mn ⎥ ⎢σ 2 ⎥ m ⎥⎢ ⎥ ⎢ ⎥ ⎢ ⎣τ xy ⎦ ⎣ mn −mn m 2 − n2 ⎦ ⎣τ 12 ⎦
(17.9a)
⎡ ⎤ ⎡ ⎤ n2 −2 mn ⎤ ⎢ ε1 ⎥ ⎢ ε x ⎥ ⎡ m2 ⎢ ε ⎥ = ⎢ n2 m2 2 mn ⎥ ⎢ ε 2 ⎥ ⎢ y ⎥ ⎢ ⎥ 2 2⎥⎢ ⎢ γ xy ⎥ ⎣ mn −mn m − n ⎦ ⎢ γ 12 ⎥ ⎢⎣ 2 ⎥⎦ ⎣ 2 ⎦
(17.9b)
where m and n are cos θ and sin θ, respectively. Moreover, introducing the following Reuter’s matrix: ⎡1 0 0 ⎤ [ R ] = ⎢0 1 0 ⎥ ⎢⎣0 0 2 ⎥⎦
(17.10)
the engineering strain vectors ⎤ ⎡ ⎡ ε1 ⎤ ⎡ 1 0 0 ⎤ ⎢ ε 1 ⎥ ⎢ ε 2 ⎥ = ⎢0 1 0 ⎥ ⎢ ε 2 ⎥ ⎢⎣γ ⎥⎦ ⎢⎣0 0 2 ⎥⎦ ⎢ γ ⎥ 12 ⎢ 12 ⎥ ⎣ 2 ⎦
(17.11a)
⎡ ⎤ ⎡ ε x ⎤ ⎡1 0 0 ⎤ ⎢ ε x ⎥ ⎢ ε y ⎥ = ⎢0 1 0 ⎥ ⎢ ε y ⎥ ⎥ ⎢ ⎥ ⎢0 0 2 ⎥ ⎢ ⎦ ⎢ γ xy ⎥ ⎣γ xy ⎦ ⎣ ⎢⎣ 2 ⎥⎦
(17.11b)
may be considered instead of the tensor strain vectors in the strain transformations as well as in stress–strain law transformations. Thus, the stress–strain relations in the x–y coordinate system becomes: ⎡σ x ⎤ ⎡Q11 ⎢σ y ⎥ = ⎢Q12 ⎢ ⎥ ⎢ ⎣τ xy ⎦ ⎢⎣Q16
Q12 Q22 Q26
Q16 Q26 Q66
⎤ ⎡ εx ⎤ ⎥⎢ ⎥ ⎥ ⎢ εy ⎥ ⎥⎦ ⎣γ xy ⎦
(17.12)
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— ¯ ] matrix represent the transformed reduced where the Qij elements of the [Q stiffnesses instead of the so-called reduced stiffnesses (Qij). In contrast to ¯ ] matrix shows non-zero terms in all nine positions, the [Q] matrix, the [Q but there are still four independent elastic constants as the lamina is orthotropic. Considering equation (17.12), in the x–y coordinate system, an orthotropic lamina behaves as anisotropic as there is a shear-extension coupling effect, which means that there is a coupling between shear strain and normal stresses and between shear stress and normal strains. Dealing with equation (17.12) and taking into account the orthotropic characteristics of such a lamina in the principal material coordinates, it may be defined as ‘generally orthotropic lamina’, which is a lamina with its principal material axes not aligned with the x–y coordinate system. Consequently, equation (17.12) represents the stress– strain relations for a lamina of an arbitrary orientation, and the transformed — reduced stiffnesses Qij result functions of the four independent elastic constants (E1, E2, n12 and G12) as well as of the θ angle. In order to show readily the effects of rotating a lamina in a laminate, Tsai and Pagano (1968) decided to write the transformed reduced stiffnesses in the following form: Q11 = U1 + U 2 cos 2θ + U3 cos 4θ Q12 = U 4 − U3 cos 4θ Q22 = U1 − U 2 cos 2θ + U3 cos 4θ 1 Q16 = U 2 sin 2θ + U3 sin 4θ 2 1 Q26 = U 2 sin 2θ − U3 sin 4θ 2 Q66 = U 5 − U3 cos 4θ
(17.13)
where U1 = U2 = U3 = U4 = U5 =
3Q11 + 3Q22 + 2Q12 + 4Q66 8 Q11 − Q22 2 Q11 + Q22 − 2Q12 − 4Q66 8 Q11 + Q22 + 6Q12 − 4Q66 8 Q11 + Q22 − 2Q12 + 4Q66 8
(17.14)
However, it must be underlined that Tsai and Pagano’s rotation angle (θ) is oppositely defined to that reported in Fig. 17.2, hence the sine terms reported in equation (17.13) are of opposite sign (Jones, 1999).
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The classical lamination theory, which is based on several simplifying assumptions, and the stress–strain relations of an individual lamina, which is seen as the basic building block, allow the constitutive equations of a thin N-layered laminate to be obtained (Fig. 17.3). By using the classical lamination theory and considering z as the direction of the normal to the middle surface, the stresses in the kth lamina of a thin N-layered laminate (plate) can be expressed in terms of the laminate middle surface strains and curvature as follows: ⎡σ x ⎤ ⎡Q11 ⎢σ y ⎥ = ⎢Q12 ⎢ ⎥ ⎢ ⎣τ xy ⎦ k ⎢⎣Q16
Q12 Q22 Q26
Q16 Q26 Q66
⎤ ⎥ ⎥ ⎥⎦ k
0 ⎡⎡ ε x ⎤ ⎡ kx ⎤ ⎤ ⎢ ⎢ ε y0 ⎥ + z ⎢ ky ⎥ ⎥ ⎢⎢ 0 ⎥ ⎢ ⎥⎥ ⎣ kxy ⎦ ⎦ ⎣ ⎣γ xy ⎦
(17.15)
where the middle surface strains are ⎡ εx ⎤ ⎢ ε y0 ⎥ ⎢ 0 ⎥ ⎣γ xy ⎦ 0
(17.16)
and the middle surface curvatures are ⎡ kx ⎤ ⎢ ky ⎥ ⎢ ⎥ ⎣ kxy ⎦
(17.17)
The resultant forces (Nx, Ny, Nxy) and moments (Mx, My, Mxy) ‘per unit width’ acting on a laminate are obtained by integration of the stresses in
t 2 Z0 Z 1
Z2
Middle surface Z
Zk-1
t Zk
K
Zn-1 Zn
N Layer number (a)
(b)
17.3 (a) Unbounded view of a generic three-layered laminate; (b) basic geometry of an N-layered laminate (Jones, 1999).
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each lamina through the laminate thickness, thus providing the following equations: ⎡ N x ⎤ ⎡ A11 ⎢ N y ⎥ = ⎢ A12 ⎢ ⎥ ⎢ ⎣ N xy ⎦ ⎣ A16
A12 A22 A26
A16 ⎤ ⎡ ε0x ⎤ ⎡ B11 A26 ⎥ ⎢ ε0y ⎥ + ⎢ B12 ⎢ ⎥ A66 ⎥⎦ ⎣ γ 0xy ⎦ ⎢⎣ B16
B12 B22 B26
⎡ M x ⎤ ⎡ B11 ⎢ M y ⎥ = ⎢ B12 ⎢ ⎥ ⎢ ⎣ M xy ⎦ ⎣ B16
B12 B22 B26
B16 ⎤ ⎡ ε x0 ⎤ ⎡ D11 B26 ⎥ ⎢ ε y0 ⎥ + ⎢ D12 ⎢ ⎥ B66 ⎥⎦ ⎣γ x0y ⎦ ⎢⎣ D16
D12 D22 D26
B16 ⎤ ⎡ kx ⎤ B26 ⎥ ⎢ ky ⎥ B66 ⎥⎦ ⎢⎣ kxy ⎥⎦ D16 ⎤ ⎡ kx ⎤ D26 ⎥ ⎢ ky ⎥ D66 ⎥⎦ ⎢⎣ kxy ⎥⎦
(17.18a)
(17.18b)
where the Aij elements represent extensional stiffnesses, the Bij elements are bending-extension coupling stiffnesses, and the Dij elements are bending stiffnesses. These generic elements of the [A], [B] and [D] matrices are: N
Aij = ∑ ( Qij )k ( zk − zk −1 ) k =1
1 N ∑ (Qij )k ( zk2 − zk2−1 ) 2 k =1 1 N Dij = ∑ ( Qij )k ( zk3 − zk3 −1 ) 3 k =1 Bij =
(17.19)
where zk is the directed distance to the bottom of the kth lamina and zk−1 represents that to the top of the kth lamina. However, integrating through the laminate thickness to obtain equations ¯ ]k for each kth ply (17.18), it has been considered that the stiffness matrix [Q is constant within the lamina, unless the lamina presents temperature- or moisture-dependent properties and a temperature gradient or a moisture gradient exists across the lamina. Thus, if the elevated temperature or mois¯ ij)k are ture is constant through the thickness of the layer, the values of (Q constant in the layer but probably degraded because of the presence of temperature and/or moisture. Accordingly, the stiffness matrix may go outside the integration over each layer, but is within the summation of forces and moments resultant for each layer. On the other hand, if there is a temperature and/or moisture gradient in the kth ply and the lamina has material properties which are temperature¯ ij)k values are a function of z and dependent and/or moisture-dependent, (Q must be left inside the integral. This situation leads to a non-homogeneous laminate within each ply, and complicated numerical solutions are needed to solve the problem (Jones, 1999). Analogously to that described for an individual lamina, at a laminate level shear–extension coupling exists owing to the presence of A16 and A26 terms in equation (17.18a). Furthermore, it may also be observed that
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bending–extension coupling are related to Bij, whereas D16 and D26 are responsible for bend–twist coupling. Equations (17.18) suggest that a laminate has a maximum of 18 stiffnesses which are the Aij, Bij and Dij elements. Considering that the number of elastic constants for a laminate result in a maximum of four per lamina, if this laminate is made of the same material in each lamina, it is characterized by only four elastic constants (Jones, 1999; Mallick, 1997; Nicolais, 1975; Tsai et al. 1968). Accordingly, the classical lamination theory allows evaluation of the strains and curvatures of the middle surface once forces and moments are known (or vice versa), hence of the layers stresses in the global laminate coordinate system. These stresses may be transformed from laminate coordinates to lamina principal directions, then a failure criterion may be applied to each lamina in its own principal material directions. All of the above discussed basic principles can be taken into consideration for designing special cases of laminated fibre-reinforced composites, in order to avoid the coupling effects described previously. Although there are more complicated cases, in common practice many laminates are constructed from laminae which present the same material properties and thickness, but with different orientation of the principal material directions relative to one another and relative to the laminate axes. For this reason, in order to better understand how to design specific laminates, some examples related to laminae with the same material properties and thickness may be briefly reported (Mallick, 1997; Jones, 1999). For example, all elements in the extension-bending coupling matrix [B] are zero when the laminate is designed to be symmetric. To produce a symmetric laminate, each lamina of +θ orientation above the midplane must be matched with an identical ply (in thickness as well as material) of +θ orientation at the same distance below the midplane. A16 and A26 are zero (i.e. no shear-extension coupling) for a balanced laminate, which is obtained when for each ply of +θ orientation, there is an identical ply (in thickness as well as material) of −θ orientation in the laminate. Accordingly, for symmetric and balanced laminates there is no bending– extension coupling or shear–extension coupling. Thus, whenever possible, laminates should be designed with a balanced and symmetric layup. Stress and strain analyses of balanced and symmetric laminates are considerably easier than those of unbalanced and/or asymmetric laminates. Moreover, symmetric laminates are dimensionally more stable than the asymmetric ones. However, it should be noted that laminates that are designed to minimize the bending–twisting coupling cannot be symmetrical unless they contain only 0 ° and 90 ° plies (Mallick, 1997). As a result, there are also other engi-
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neering solutions which can help to obtain special cases of laminate stiffnesses through an appropriate sequence of stacking. Eventual thermal and hygroscopic (moisture) stress analyses must be also considered (Jones, 1999; Mallick, 1997; Nicolais, 1975). An important aspect in designing fibre-reinforced composite materials is clearly related to the stress transfer from matrix to fibre. Kelly and Tyson (1965) initially analysed and described this aspect. The main assumptions made by the authors were that the matrix yields plastically, the fibre/matrix adhesion is perfect and the shear stress along the fibre is constant and equal to the shear yield strength of the matrix, τy. With these assumptions, τy is given by the following equation:
τy =
σ fb D 2lc
(17.20)
where σfb represents the average tensile strength of the fibre, D the fibre diameter and lc the fibre critical length. In this context, single fibre techniques have been developed for analysing stress transferability and for evaluating the effectiveness of fibre surface treatments. The embedded single fibre composite technique requires the mounting of a single fibre in a matrix that has an elongation-to-break higher than that of the fibre and the axial stressing of the sample to an elongation that results in fragmentation of the embedded fibre. The distribution of fragment lengths may be observed by means of an optical microscope, and then an average interfacial shear stress transferability, 〈τ〉, may be calculated as reported in equation (17.20) (Bian et al., 1991). In conclusion, by tailoring geometry, material properties and the orientation of each lamina it is possible to design a laminate with appropriate properties for specific applications.
17.3
Short-fibre composites
In short-fibre composites, the applied load is transmitted through the matrix to the fibre by the shear-transfer mechanism. In particular, taking into account the fibre critical length, the average fibre stress may be evaluated and then substituted into an expression to obtain the ultimate strength for short-fibre composites (Kelly and Tyson, 1965; Iroh and Berry, 1996). However, if short fibres are considered, the properties may be also provided by Halpin Tsai equations (1968), where the moduli in the fibre direction is a sensitive function of the aspect ratio (l/d) at small aspect ratios, giving the same properties as a continuous fibre composite at large but finite aspect ratios. Random or nearly random distributions of fibres, finite in length and arranged in a matrix, constitute many naturally occurring and synthetic materials.
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Usually, the spatial orientation of the discontinuous fibres results in a configuration intermediate between a truly random array in three dimensions and a two-dimensional random array of fibres. Generally, these materials possess internal orientation which is not dependent on the thickness, and often the distribution of fibres may vary through the thickness. Halpin and Pagano (1969) modelled these materials as laminated systems. This laminate model consists of layers of unidirectional short-fibre composites with fibre volume fractions in a layer oriented at an angle θ, determined by the percentage of fibres at the same angle in the actual material. If the material to be modelled exists in a sheet form, wherein the thickness is considerably less than the average lengths of fibres, then the reinforcement may be considered as a two-dimensional random array of fibres. This is described in laminated plate theory as ‘quasi-isotropic’. The success of the proposed lamination approximation is strongly related to the assumption of physical volume averaging in real material systems combined with the ability to estimate the stiffness for an oriented short fibre-sheet. Comparing theoretical predictions with experimental data obtained from mechanical tests on random, quasi-isotropic and oriented short nylon fibre composites, it has been possible to predict the strong dependence of longitudinal modulus (E1) on aspect ratio (l/d) through Halpin–Tsai equations (Halpin, 1969; Nicolais,1975). In particular, as the aspect ratio became large, E1 values became identical to those of oriented continuous fibres. The other moduli (E2, G12) and n12 seemed not to be sensitive functions of the aspect ratio, and might be approximated by the continuous fibre results. Moreover, to test the laminate model, ‘quasi-isotropic’ laminates of two different aspect ratios were realized and then tested. The stiffness of these two laminates was similar to the random composites, and theoretical predictions were in agreement with experimental data. In particular, experimental data and predictions for random short boron fibres in an epoxy matrix showed that stiffness reached a plateau as the aspect ratio increased. The critical value of l/d at which the plateau occurs is strongly dependent upon the ratio Ef/Em (Nicolais, 1975). In this analysis, it was assumed that there was no bias in fibre orientation and that all the fibres have the same aspect ratio. However, in the preparation of short fibre composites only partial orientation is usually achieved, and most fabrication procedures result in considerable fibre breakage (Nicolais, 1975; Schierding and des Deex, 1969), thus producing a distribution of fibre lengths. The corresponding theory employing the lamination analogy requires that the fibre orientation can be handled as a complex laminate of weighed angle plies ±θ, assuming a symmetric angular distribution f(θ) to match the required fibre distribution. Employing the laminate analogy, we may assume that the material behaves mathematically as a
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biased laminate obtained from layers of oriented short-fibre composite material of volume fraction Vf. The percentage a(θk)/h of the material orientedπat the ±θ is obtained from the experimental angular distribution function ∫ f (θ )dθ = 1.0 . In this way, 0 different orientations can contribute to the overall response in proportion to their fractional thickness in the laminate or: a (θ k ) Aij′ (θ k ) h k =1 N
Aij = ∑
(17.21)
where Aij is the laminate stiffness and Aij′ (θ k ) represents the stiffness of the plies oriented at θ angles, and N is the number of θ orientations. The stiffness moduli are specified in terms of E1, E2, n12 and G12 for a ply and the angle of orientation θ. In turn, the lamina properties are related to fibres and matrix properties and to the aspect ratio l/d as specified in Halpin–Tsai equations. Then, ‘the engineering stiffness properties’ of the laminate can be derived from the Aij elements. If the laminate analogy is extended to the variable aspect ratio problem, it would be required to divide each layer for each specific θ direction into another series of laminates having different aspect ratios in each layer. In this case, a simple solution can be to use an average aspect ratio. The ‘first moment’ of the distribution is defined as the ‘number average aspect ratio’, 〈l/d〉 which is inserted into the micromechanics calculations, the Halpin–Tsai equations, as if it were a monochromatic aspect ratio (Nicolais, 1975). The maximum strain theory may be modified to predict randomly oriented short-fibre composites (Jerina et al., 1973; Nicolais, 1975). The Halpin– Tsai equations have established relations for the stiffness of an oriented short-fibre ply from the matrix and fibre properties. These equations show that the longitudinal stiffness of an oriented short-fibre composite is a sensitive function of the aspect ratio. The short-fibre stiffness asymptotically approaches to the continuous filament at large aspect ratios. As with stiffness, strength is a function of aspect ratio and approaches an asymptotic limit as the aspect ratio becomes large. However, the strength limit for short fibres does not approach the continuous filament strength. Consequently, the oriented short-fibre material fails at an ultimate longitudinal strain which is less than the ultimate longitudinal strain of the continuous fibre material. Studies have highlighted that the asymptotic short-fibre strength limit is less than the continuous filament composite strength (Chen, 1971; Nicolais, 1975). Finite element analysis of the discontinuous fibres has shown a stress concentration factor owing to the fibre ends, which becomes constant at sufficiently large aspect ratios (Barker and MacLaughlin, 1971). A plateau
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value for the strength has also been reported when the fibres become sufficiently long (Chen, 1971; Nicolais, 1975). The strength of high-aspect-ratio short-fibre randomly oriented composites may be predicted with maximum strain by reducing the longitudinal strength allowable to reflect the reduced strength in the fibre direction owing to the discontinuous reinforcements. For example, it has been reported that at sufficiently large aspect ratios, the strength of short fibre glass/epoxy composites was 60% of the strength of the continuous filament composite (Chen, 1971; Nicolais, 1975). This fact may be incorporated into maximum strain theory by reducing the continuous filament longitudinal strain allowable by 60%. The effect of fibre volume fraction of the quasiisotropic strength can be included by using the Halpin–Tsai equations for evaluating the orthotropic short-fibre ply moduli.
17.4
Particulate composites
For polymer manufacturing processes such as injection moulding and compression moulding, several theoretical models have been developed to predict the tensile modulus and tensile strength of the particulatereinforced polymers. However, the theories for predicting the strength of particulate reinforced systems are less developed than those for predicting the modulus (Fu et al., 2008; Nielsen and Landel, 1994; Sahu and Broutman, 1972; Yan et al., 2006). The behaviour of materials obtained with spherical particles included in polymeric matrix may be considered isotropic, and the Young’s modulus (EC) of the composite can be easily predicted using the Halpin–Tsai equations for E2 and G12 described earlier, or the following Kerner equation (Nicolais, 1975; Schapery, 1968): Ec = Em
1 + ACVf 1 − CVf
(17.22)
where 7 − 5ν m 8 − 10ν m Ef −1 Em C= Ef +A Em A=
(17.23)
and nm represents the Poisson’s ratio of the matrix, Vf is the volume fraction of the filler, Em and Ef are the Young’s moduli of matrix and dispersed material, respectively.
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In addition to the elastic moduli, the other tensile properties of this class of materials are sensitive to the properties of the matrix and the adhesion between filler and matrix. If a matrix is able to craze, it contributes to an inhomogeneous deformational mechanism that leads to an increase in elongation and work-to-break for the resulting composites owing to the formation and propagation of crazes through the polymer. The crazes normally observed in thermoplastic are not cracks, but rather localized occurrences of highly oriented polymer (Nicolais, 1975). For example, results from styrene acrylonitrile (SAN)/glass bead composites have shown a high elongation, which may be explained assuming that the growth of crazes can be terminated by the glass beads or vacuoles around them (Lavengood et al., 1973; Nicolais, 1975). If the propagating craze should encounter a glass sphere to which the matrix is not strongly adherent, interfacial debonding can effectively blunt the tip of the craze and prevent, or at least slow down, further craze propagation. Therefore, the elongation at break may be enhanced by the fact that in a glass-bead composite the crazes, which in unfilled materials may be nucleated only at singular defects, nucleate throughout the entire specimen. Thus, filler particles act as stress risers allowing multiple volume elements to reach the critical stress for craze formation. Moreover, the strength of these composites decreases as the volume content of beads increases. This phenomenon is also observed in limiting the transverse tensile strain and in-plane shear strain capabilities of the unidirectional prepreg materials; variations in reinforcement volumetric packing is not a material design variable. This strength decrease is simply a reflection of the decreased cross-sectional area of the polymer bearing the load and can be expressed as a function of concentration by the following Nicolais–Narkis equation (Nicolais and Narkis, 1971): σc = σm (1 − 1.21 Vf2 / 3 )
(17.24)
where σc and σm represent the strengths of the composite and the polymer, respectively. The Nicolais–Narkis equation describes structures where the adhesion is poor, because the weighting factor (1.21) is believed to be dependent on the adhesion quality between the matrix and the inclusion. In particular, a value of 1.21 for the weighting factor is stated to be valid for the extreme case of poor adhesion and spherical inclusions (Kunori and Geil, 1980). A semiempirical single parameter equation describing the moduli of particulate systems was developed by Narkis (Narkis et al., 1978): Ec 1 = Em K (1 − Vf1 / 3 )
(17.25)
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where K is an empirical parameter apparently related to a stress concentration factor, with usual values in the range of 1.4–1.7, whereas Vf still represents the bead volume fraction. As for the modulus of particulate composites, another interesting approach is represented by the empirical and useful Eilers equation (Eilers, 1941). In the rubbery state, it is well known that the rubbery elastic property of the matrix is dominant. In this context, the Eilers equation was based on the relationship between the relative viscosity and the volume fraction of the disperse phase (Eilers, 1941; Yang et al., 2006) and is often used to describe the effect of filler on the rubbery modulus. The Eilers expression for the modulus of the composite may be written as: EC ⎛ 1.25Vf ⎞ = ⎜1 + ⎟ Em ⎝ 1 − Vf φm ⎠
2
(17.26)
in which φm represents the maximum packing factor of particulate reinforcement. Other theoretical models for tensile modulus and strength of particulate composites are summarized in Table 17.2 (Fu et al., 2008; Guth, 1945; Nielsen, 1966; Nielsen and Landel, 1994; Verbeek, 2003; Verbeek and Focke, 2002; Yan et al., 2006). It may be noticed that the application of these models for predicting the composite strength and modulus requires the values of some important parameters or constants. For example, an important parameter is the maximum packing factor of particulate reinforcement, φm, which can be obtained either theoretically or experimentally once the particle distribution condition has been determined. The Einstein coefficient, KE, represents another important factor which can be calculated by knowing the Poisson’s ratio of the matrix material and the relative Einstein coefficient ratio, KE/2.5, where 2.5 is usually the KE value of a material with a Poisson’s ratio of 0.5 (Nielsen and Landel, 1994). However, the strength of the composite depends on the weakest fracture path present through the material. Hard particles affect the strength in two ways. One of these is a weakening effect, owing to the stress concentration they cause. Another is a reinforcing effect, since the hard particles may serve as barriers to crack growth. In some cases, the weakening effect is predominant and hence the composite strength is lower than the matrix. In other cases, the reinforcing effect is more significant and then the composites will have strengths higher than the matrix (Fu et al., 2008). The prediction of the strength of composites is difficult, and the difficulty comes from the consideration that the strength of composites is determined by the fracture behaviours which are associated with the extreme values of such parameters as interface adhesion, stress concentration and defect size/ spatial distributions.
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Table 17.2 Other theoretical models for tensile modulus and strength of particulate composites (Yan et al., 2006) Name
Models
Nomenclature
Guth model
For spherical particles:
EC = tensile modulus of the reinforced polymer Em = tensile modulus of the matrix KE = Einstein coefficient Vf = reinforcement volume fraction α = reinforcement aspect ratio A = a constant which takes into account such factors as geometry of the reinforcing phase and Poisson’s ratio of the matrix B = a constant which takes into account the relative modulus of the reinforcing and matrix phases Ef = tensile modulus of the reinforcement vm = Poisson’s ratio of matrix Vm = volume fraction of polymeric matrix (associated with zero void) ϕ = a factor which depends on the maximum packing fraction φf φf = maximum packing fraction of reinforcement φ = void content of composite χ = modified void content (void relative to polymer phase) MRF = modulus reduction factor Gm = shear modulus of the polymeric matrix σC = tensile strength of the composite τm = shear strength of the polymeric matrix K3 = correction factor MPF = strength reduction factor σm = tensile strength of the polymeric matrix K = stress concentration factor
EC = Em (1 + K EVf + 14.1Vf2 ) For non-spherical particles: EC = Em (1 + 0.67αVf + 1.62α 2Vf2 ) Nielsen model for tensile modulus (based on Halpin–Tsai and Kerner models)
E C = Em
1 + ABVf 1 − BϕVf
Ef −1 Em B= Ef +A Em
ϕ ≅ 1+
1 − φf Vf φf
with spherical reinforcements: A=
7 − 5ν m 8 − 10ν m
with nearly spherical shaped reinforcements: A = KE − 1 Verbeek model for tensile modulus
(1 − Vm ) φm 2
φ= χ=
1 − Vmφm
φ Vm (1 − φ ) + φ (1 − χ ) Gm 3
ϕ =α
Ef
MRF = 1 −
Vf 1 − Vf
tan h (ϕ ) ϕ
E c = VfE fMRF + VmEm Verbeek model for tensile strength
12
⎛ GmVf ⎞ u = α⎜ ⎟ ⎝ E f (1 − Vf ) ⎠ MPF = Vf
Nielsen model for tensile strength
( αu ) ⎛⎜⎝ tan 1h u
( )
σ C = Vmσ m + K 3τ mMPF σC = σm (1 − Vf2 3 ) K
−
1⎞ ⎟ u⎠
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The load-bearing capacity of a particulate composite therefore depends on the strength of the weakest path throughout the microstructure, rather than the statistically averaged values of the microstructure parameters. Hence, although numerous theories of composite strength have been published before, there is no universally accepted theory to date. Some theoretical models for composite strength have already been reported above, whilst other empirical and semi-empirical equations are reported in what follows. These models are easy to use in practice, and can give correct predictions in appropriate cases (Fu et al., 2008). The Nielsen model (Table 17.2) is often used for the prediction of strength of particulate composites where there are poorly bonded particles. In this model, the parameter K represents weakness in the structure caused by discontinuities in the stress transfer and generation of stress concentration at the particle/polymer interface. When there is no stress concentration, the value of K will be its maximum, equal to unity. Under the hypothesis that there is no adhesion between filler and polymer, which means that the load is sustained only by the polymer, the above described Nicolais–Narkis equation gives the lower bound strength of the composite. An upper bound is immediately obtained by considering that, for perfect adhesion, the strength of the composite is simply equal to the strength of the polymer matrix. Thus, the strength is intermediate between these two bounds and cannot be higher than that of the matrix. When some adhesion exists between particle and matrix, the interface can transfer a part of the stress to the particles, making a contribution to the composite strength. Accordingly, in order to better predict the composite strength, the following equation may be considered (Bigg, 1987; Fu et al., 2008):
σ C = σ m (1 − aVfb + cVfd )
(17.27)
where c and d are constants. The Nicolais–Narkis equation can be also modified to include the situation of some adhesion:
σ C = σ m (1 − 1.07Vf2 3 )
(17.28)
Arguing that the assumption of a uniform filler distribution as used in most models was unlikely in practice, Piggott and Leidner (1974) proposed an empirical equation:
σ C = gσ m − α ′Vf
(17.29)
where α′ is a coefficient of the particle/matrix adhesion, and g is a constant. It may be observed that in the case of nil or some adhesion, the addition of particles generally leads to a reduction in strength. However, for strong
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filler–matrix adhesion, where the stress is transferred through shear from matrix to particles, the strength of the composite may be written as:
σ C = 0.83 pfVf + kσ m (1 − Vf ) + σ a H (1 − Vf )
(17.30)
where f represents the coefficient of particle/matrix friction, p the residual pressure, k the relative change in the strength of the matrix owing to the presence of the filler, σa the particle/matrix adhesion strength, and H a constant (Fu et al., 2008; Leidner and Woodhams, 1974). In general, particulate composites can be used in many applications owing to the fact that the addition of rigid particulate fillers to a polymeric matrix leads to a material with particular properties such as high modulus, lower creep, more resistance to abrasion.
17.5
Polymer nanocomposites
Polymer nanocomposites are commonly defined as the combination of a polymer matrix and additives that present at least one dimension in the nanometer range (Liao et al., 2006; Nicolais and Carotenuto, 2005). These additives can be one-dimensional (nanotubes and fibres), two-dimensional (layered minerals like clay), or three-dimensional (spherical particles). Polymer nanocomposites have attracted the attention of the scientific community, owing to their outstanding mechanical properties, like elastic stiffness and strength, with only a small amount of the nanoadditives. This is because of the large surface-area-to-volume ratio of nanoadditives when compared with the micro- and macro-additives. Other advantageous properties of polymer nanocomposites include barrier resistance, flame retardancy and scratch/wear resistance, as well as optical, magnetic and electrical properties. In this context, carbon nanotubes (CNTs) have played an important role. They may be seen as seamlessly rolled sheets of a hexagonal array of carbon atoms with diameters ranging from a few angstroms to several tens of nanometers across. CNTs exist in two forms, a single-walled carbon nanotube (SWNT) in which the tube is formed from only a single layer of graphitic carbon atoms, and multi-walled carbon nanotubes (MWNT), in which the tube is made of several layers of coaxial carbon tubes. The great interest in CNTs arises from their unique structural and physical properties, such as their small size in the nanometer scale and their unique electronic behaviour, since they can be either metallic or semiconducting depending on their geometrical structure. Further potentially interesting characteristics are their exceptional properties of ballistic transport, their high thermal conductivity and high optical polarizability. In particular, great attention has also been focused on their mechanical properties such as high elastic modulus and tensile strength. By stretching SWNT
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Biomedical composites
bundles through an atomic force microscope (AFM), Walters et al. (1999) found that the tensile strength of SWNT was 45 ± 7 GPa, whilst by pulling 2-mm-long MWNT ropes, Pan et al. (1999) obtained Young’s modulus of 0.45 ± 0.23 TPa and tensile strength of 1.72 ± 0.64 GPa. Yu et al. (2000) carried out tensile tests on individual MWNT and SWNT ropes through AFM, reporting that the tensile strengths of SWNT rope and MWNTs varied from 11 to 63 GPa and 13 to 52 GPa, respectively. For SWNT bundles values of tensile strength ranging from 2.3 ± 0.2 GPa to 14.2 ± 1.4 GPa was evaluated by Li et al. (2000), whilst Zhu et al. (2002) characterized the Young’s modulus and tensile strength of SWNT strands, highlighting that the tensile strength ranged from 49 to 77 GPa. However, in their studies the lowest value obtained for elastic modulus of SWNT was only about 100 GPa, owning to inter-nanotube defects. However, the high elastic modulus and tensile strength of CNTs, when incorporated in polymer matrix, does not necessarily provide a composite with enhanced mechanical properties. In this context, studies on using CNTs in polymer composites have highlighted both encouraging and discouraging results. Some researchers have shown that there is no enhancement in mechanical properties when CNTs are added to a polymer matrix. In contrast, other studies have shown a moderate increase in elastic modulus with slight or no enhancement in mechanical strength for CNT/polymer composites. Xu et al. (2002) studied the mechanical properties and interfacial characteristics of MWNT/epoxy composite thin films from a spin-coating process. Although compared with neat resin thin films, an increase of 24% in elastic modulus was found when 0.1 wt% MWNTs was added, the fracture load of the composite film is somewhat lower than that of the neat epoxy film. Studies on an aligned SWNT/poly(methyl methacrylate) (PMMA) composite have shown moderate increase in elastic modulus and yield strength with an increase in nanotube loading and draw ratio (Haggenmueller et al., 2000). As is well known, PMMA is the basic ingredient of acrylic bone cement and dental prostheses. Bone cement is a grouting agent, which plays an important interfacial role between metallic total joint prostheses and host bone, and may also be used as an injectable material in the field of vertebroplasty. However, this polymer continues to present a low resistance to mechanical fatigue and impact. For this reason, a great variety of materials have been considered in order to enhance its mechanical strength, but none of these materials has provided good results (Marrs et al., 2006). To improve the mechanical properties of bone cement, many researchers have proposed to incorporate small amounts of different fibres (glass, carbon, stainless steel, Kevlar-related) into this polymer, the aim being to bridge incipient fatigue cracks and arrest their propagation.
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Drawbacks related to large fibre size, fibre–bone cement matrix debonding, viscosity increase, ductile fibre deformation and fracture, and nonuniform additive material distribution have led to situations that are far from the ideal solution. Consequently, the introduction of nanosized materials, particularly carbon nanotubes seems to offer new promise for enhancing the properties of polymer, including PMMA-based bone cement. Thus, carbon nanotubes, in both single and multiwall varieties, have emerged as one of the most exciting solution. For example, Marrs et al. (2006) have clearly shown that small percentages of multiwall carbon nanotubes improve the mechanical properties of bone cement. In particular, results from quasi-static tests have shown that a 2 wt% MWNT concentration was nearly optimal for enhancing the mechanical properties of bone cement in 3-point bending. This amount of MWNT enhanced flexural strength by 12.8%. Moreover, the 2 wt% loading of multiwall carbon nanotubes was associated with substantial enhancement of the fatigue performance of bone cement, showing a 3.1-fold increase in the mean actual fatigue life. Larger concentrations of MWNTs did not seem to produce the enhancements observed using a 2 wt% concentration, hence suggesting that the optimal concentration was near 2 wt%. Generally, it may be concluded that specific loadings of nanosized fillers may favourably alter the static and fatigue mechanical properties of polymer, even if many technical aspects have to be considered in designing polymer nanocomposites. In fact, the mechanical properties of polymer nanocomposites obtained from experimental results reported in literature suggest further research in terms of processing methods, understanding and control of nanoparticles–polymer interfacial adhesion, eventual nanosized particles treatments, suitable dispersion techniques in the matrix, in order to realize polymer nanocomposites with specific and predictable performance. Moreover, great efforts have been performed to find suitable models for describing the modulus of polymer nanocomposites. Generally, the modulus of polymer composites has been extensively studied experimentally and treated with a two-phase model by various researchers. In particular, the two-phase model suggested by Takayanagi (1964) has been widely used to explain the modulus of polymers, polymer blends and composite materials. However, recent studies on several polymer/inorganic nanocomposites have highlighted that a two-phase model would not be suitable to calculate their moduli (Ji et al., 2002). In this context, studies performed on nylon 6/montmorillonite nanocomposites have underlined the possibility of considering a three-phase model as a powerful tool to describe the tensile modulus of polymer nanocomposites (Ji et al., 2002). In fact, based on Takayanagi’s two-phase model, a three-phase model, which includes the matrix, interfacial region and fillers, has been proposed. In this model, fillers (sphere-, cylinder- or plate-shape)
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are randomly distributed in a matrix. If the particulate size is in the range of nanometers, the interfacial region clearly plays an important role in the evaluation of moduli of polymer nanocomposites. Important system parameters include the dispersed particle size, shape, thickness of the interfacial region (δ ), particulate-to-matrix modulus ratio (ED/EM), and a parameter (k) describing a linear gradient change in modulus between the matrix and the surface of particle on the modulus of nanocomposites. The role of the interface in the tensile modulus of polymer nanocomposites may be traduced by the fact that there is a thin layer between the rigid dispersed phase and the matrix that should possess different mechanical behaviour both from the matrix and from the dispersed phase. Thus, this three-phase model considers an interface region in order to provide an expression of the tensile modulus of polymer nanocomposites as a function of volume fraction of dispersed phase, interface and matrix. The moduli of matrix, dispersed phase, and interface may be given as EM, ED and Ei, respectively, and their volume fractions are VM, VD and Vi, respectively, with VM + VD + Vi = 1. This system could be schematically represented as a three-phase model in which the three phases are connected to each other in series and in parallel (Fig. 17.4) (Ji et al., 2002). The response of these three phases to a stress may be reduced to that of three regions A, B, and C connected in series. The elongations of these three regions under the stress loaded on specimen T can be indicated as εA, εB, and εC, respectively. In region A, only the matrix with volume fraction VAM exists. In region B, the interface with volume fraction VBi and matrix with
T
l
M
i1
D
A
i2
B
j b M a
M
C
17.4 Three-phase model representation for polymer nanocomposites: the equivalent model (Ji et al., 2002).
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435
volume fraction VBM coexist in a parallel arrangement. In region C, the matrix with volume fraction VCM, the interface with volume fraction VCi, and the dispersed phase with volume fraction VD coexist in a parallel arrangement. It is well known that VAM + VBM + VCM = VM, VBi + VCi = Vi. The elongation in region A is:
εA =
T T + EA EM
(17.31)
where EA is the modulus of the region A, i.e. EM. The elongation in region B is:
εB =
T (1 − α ) EM + α EBi
(17.32)
whilst the elongation in region C is:
εC =
T (1 − α ) EM + (α − λ ) ECi + λ ED
(17.33)
The overall elongation of these three regions is:
ε = (1 − β ) ε A + ( β − ϕ ) ε B + ϕε C β −ϕ ϕ ⎛ 1− β ⎞ =T⎜ + + ⎟ (1 − α ) EM + α EBi (1 − α ) EM + (α − λ ) ECi + λ ED ⎠ ⎝ EM (17.34) λϕ = Vd αβ − λϕ = Vi
(17.35)
Since ε/T = 1/Ec, where Ec is the modulus of the composite: 1 1− β β −ϕ ϕ = + + (1 − α ) EM + α EBi (1 − α ) EM + (α − λ ) ECi + λ ED EC EM (17.36) It is worth noting that: a) when α − λ = β − ϕ = 0, or α = λ and β = ϕ, the above reported equation becomes: 1 1−ϕ ϕ = + EC EM (1 − λ ) EM + (α − λ ) ED
(17.37)
it represents just the Takayanagi equation, describing the modulus of the two-phase composite with one homogeneous rigid phase and one homogeneous matrix phase, as ordinarily used in previous studies.
436
Biomedical composites
b) when λϕ = VD = 0, which means that there is no dispersed phase and, hence, no interfacial region, only a matrix exists: EC = E M
(17.38)
c) when 1 − β = 0, 1 − α = 0, i.e. α = β = 1 and VM = 0, i.e., the matrix phase does not exist. Thus, there is no interfacial region and a dispersed phase giving: EC = E D
(17.39)
As for the moduli of the interface, in regions 1 and 2 they are in quite different forms as shown in Fig. 17.4 Ei1 = Ei1 ( l )
(17.40)
where Ei1(l) is the modulus at region 1 with the distance l from the surface of the dispersed phase. If a linear gradient distribution of the modulus along the normal direction is assumed: dEi1 =
Ei1 ( 0) − E M dl δ
(17.41)
then, after integration: Ei1 (l ) = Ei1 ( 0) − ⎡ ⎣⎢
Ei1 ( 0) − E M ⎤ l ⎦⎥ δ
(17.42)
where Ei1(0) is the modulus of the interface at the surface of the dispersed rigid phase and τ represents the thickness of interface region. Consequently, when l = 0, Ei1(0) stands for the modulus of the interface neighbouring to the surface of the fillers and it is a constant.; when l = δ, Ei1(τ) is the modulus at the edge of interface next to matrix, i.e. EM. Accordingly, the modulus of the interface region 1 with thickness of δ is: Ei1 =
Ei1 ( 0 ) + EM 2
(17.43)
As shown in Fig. 17.4, region 2 connects with the dispersed phase at the top, and connects with the matrix at the bottom. From top to bottom, the modulus of this phase varies from Ei2(0) to EM. This ‘block i2’ may be treated as a series arrangement of a number of volume units varying as a function of l: δ
1 1 1 1 E ( 0) ⎡ ⎤ = dl = ln ⎡ i 2 ⎤ ⎢ Ei 2 δ ∫0 Ei 2 (l ) ⎣⎢ E M ⎦⎥ ⎣ Ei 2 ( 0) − E M ⎥⎦
(17.44)
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437
where Ei2(0) is also the modulus at the interface of the dispersed phase. It should be noted that EBi = Ei2, ECi = Ei1. Of course, Ei1 or Ei2 could be another type of function depending on the interaction of the macromolecules with the surface of the dispersed phase. If the modulus of the interfacial region takes the form of the function with a linear gradient decrease along the normal direction of the surface of dispersed phase, the tensile modulus of polymer nanocomposites (EC) may be given by: 1 1− β = + EC EM
β −ϕ ⎡ ⎛ EM ⎞ [ ( ) ⎤ (1 − α ) EM + α ⎢ ln ⎜ ⎟ Ei2 0 − EM ]⎥ ⎣ ⎝ E12 ( 0 ) ⎠ ⎦ ϕ + (1 − α ) EM + (α − λ ) [ Ei1 ( 0 ) + EM ] 2 + λ ED
(17.45)
If a random orientation of a plate-like dispersed phase with a thickness t and both length and width are ξ >> t, the following conditions may be assumed: α = β, ϕ = λ. Each particle presents two interface regions, Vd = λ2 and α2 − Vd = (2δ/t)Vd, which can be arranged to α = [(2δ/t + 1)Vd]1/2. The modulus of the interfacial region may be assumed to be: Ei1 ( 0 ) = Ei2 ( 0 ) = kEM
(17.46)
where k represents the modulus ratio of the interface neighbouring on the surface of a particle and describes a linear gradient change in modulus between the matrix and the surface of particle. Analogously, if the dispersed phase has a spherical form, α may be written in a form which considers the radius r of the dispersed sphere. Using these considerations, equation (17.45) can be simplified. Employing an iteration method and experimental data, it is possible to obtain δ for the investigated nanocomposite.
17.6
Conclusions
Polymer composite technology is a technology whose time has come in many industrial sectors and, specifically, in the biomedical sector. Design methodology and control of related parameters are important to define different processing techniques and applications. The extension of this, from an initial micro- and macro-scale approach to the actual nano-scale level, promises to lead to understanding and definition of what were previously phenomenological aspects related to naturally made composites, leading to the design of a new series of multi-functional materials.
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17.7
References
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mallick pk (1997), Composites engineering handbook, New York, Marcel Dekker, Inc. marrs b, andrews r, rantell t, pienkowski d (2006), ‘Augmentation of acrylic bone cement with multiwall carbon nanotubes’, J Biomed Mater Res A, 77, 269–276, DOI: 10.1002/jbm.a.30651. narkis m, nicolais l, joseph e (1978), ‘The elastic modulus of particulate-filled polymers’, J Appl Polym Sci, 22, 2391–2394, DOI: 10.1002/app.1978.070220829. nicolais l, narkis d (1971), ‘Stress–strain behaviour of SAN/glass bead composites in the glassy region’, Polym Eng Sci, 11, 194–199, DOI: 10.1002/app.1971.070150220. nicolais l (1975), ‘Mechanics of composites’, Polym Eng Sci, 15, 137–149. nicolais l, carotenuto g (2005), ‘Metal polymers nanocomposites’, New Jersey, Wiley. nielsen le (1966), ‘Simple theory of stress–strain properties of filled polymers’, J Appl Polym Sci, 10, 97–103. nielsen le, landel rf (1994), ‘Mechanical properties of polymers and composites’, New York, Marcel Dekker Inc. pan zw, xie ss, lu l, chang bh, sun lf, zhou wy, wang g, zhang dl (1999), ‘Tensile tests of ropes of very long aligned multiwall carbon nanotubes’, Appl Phys Lett, 74, 3152–3154, DOI: 10.1063/1.124094. piggott mr, leidner j (1974), ‘Misconceptions about filled polymers’, J Appl Polym Sci, 18, 1619–1623. ramakrishna s, mayer j, wintermantel e, leong kw (2001), ‘Biomedical applications of polymer–composite materials: a review’, Compos Sci Technol, 61, 1189–1224. sahu s, broutman lj (1972), ‘Mechanical properties of particulate composites’, Polym Eng Sci, 12, 91–100. schapery ra (1968), ‘Thermal expansion coefficients of composite materials based on energy principles’, J Compos Mater, 2, 380–404, DOI: 10.1177/ 002199836800200308. schierding rg, des deex o (1969), ‘Factors influencing the properties of whisker-metal composites’, J Compos Mater, 3, 618–629, DOI: 10.1177/ 002199836900300403. schwartz mm (1992), Composite materials handbook, 2nd ed, New York, McGraw-Hill. seferis jc, nicolais l (1983), ‘Introductory remarks’ In Seferis JC, Nicolais L, Eds, The role of the polymeric matrix in the processing and structural properties of composite materials, Plenum Press, New York and London. takayanagi m, uemura s, minami s (1964), ‘Application of equivalent model method to dynamic rheo-optical properties of crystalline polymer’, J Polym Sci Part C, 5, 113–122. tsai sw, pagano nj (1968), ‘Invariant properties of composite materials’, Composite Materials Workshop Papers, St. Louis, 233–253. verbeek cjr, focke ww (2002), ‘Modelling the Young’s modulus of platelet reinforced thermoplastic sheet composites’, Compos A: Appl Sci Man, 33, 1697–1704. verbeek cjr (2003), ‘The influence of interfacial adhesion, particle size and size distribution on the predicted mechanical properties of particulate thermoplastic composites’, Mater Lett, 57, 1919–1924, DOI:10.1016/S0167-577X(02)01105-9. walters da, ericson lm, casavant mj, liu j, colbert dt, smith ka, smalley re (1999), ‘Elastic strain of freely suspended single-wall carbon nanotube ropes’, Appl Phys Lett, 74, 3803–3805, DOI: 10.1063/1.124185.
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xu xj, thwe mm, shearwood c, liao k (2002), ‘Mechanical properties and interfacial characteristics of carbon nanotube reinforced epoxy thin film’, Appl Phys Lett, 81, 2833–2835, DOI: 10.1063/1.1511532. yan w, lin rjt, bhattacharyya d (2006), ‘Particulate reinforced rotationally moulded polyethylene composites – mixing methods and mechanical properties’, Compos Sci Technol, 66, 2080–2088. yang dg, jansen kmb, ernst lj, zhang gq, bressers hjl, janssen jhj (2006), ‘Effect of filler concentration of rubbery shear and bulk modulus of molding compounds’, Microelectron Reliability, 47, 233–239. yu mf, files bs, arepalli s, ruoff rs (2000), ‘Tensile loading of ropes of single wall carbon nanotubes and their mechanical properties’, Phys Rev Lett, 84, 5552–5555, DOI: 10.1103/PhysRevLett.84.5552. yu mf, lourie o, dyer mj, moloni k, kelly tf, ruoff rs (2000), ‘Strength and breaking mechanism of multiwalled carbon nanotubes under tensile load’, Science, 287, 637–640, DOI: 10.1126/science.287.5453.637. zhu hw, xu cl, wu dh, wei bq, vajtai r, ajayan pm (2002), ‘Direct synthesis of long single-walled carbon nanotube strands’, Science, 296, 884–886, DOI: 10.1126/ science.1066996.
18 Tribology of biocomposites S. K A N AG A R A J, Indian Institute of Technology, Guwahati, India; M. S. A. O L I V E I R A and J. A. D E O L I V E I R A S I M Õ E S, University of Aveiro, Portugal
Abstract: The development of orthopaedic implants in total joint replacements (TJR) includes the use and study of a variety of materials. Though a variety of natural and metallic biomaterials were employed, ultra-high-molecular weight polyethylene (UHMWPE) has been amongst the most widely used material for the last 30 years for THR. Despite the recognized success and worldwide acceptance, it has the limitation of producing very fine particles, which causes an undesirable biological response. It is estimated that an approximately 10–20% of the hip and knee replacement surgeries are being revised every year for a variety of reasons. Thus, the methods to improve the characteristics of implants are the major focus of research. As the tribological characterization of the material plays an important role in the life of the component, an attempt has been made to overview the tribological properties of various materials that are currently being investigated as potential combinations for TJR implants. Key words: tribology, biomaterials, total joint replacement, nanocomposites, UHMWPE.
18.1
Introduction
Any material that is used in the repair and reconstruction of lost, damaged or deceased tissues can be regarded as a biomaterial that is used as implants intended to interact with biological systems owing to its biocompatibility in addition to other required properties. A large number of polymer composites, biocomposites, are widely used in various medical applications because of their great versatility, attractive properties and variety of compositions. One of the most prominent applications of biocomposites is for orthopedics. Wear of any biocomposites and resulting wear debris which is leading to osteolysis is a major cause of failure in both hip and knee prostheses, Ingham and Fisher (2000). It is estimated that an approximately 10–20% of the hip and knee replacement surgeries are being revised every year for a variety of reasons. While continued development of biomaterials has increased the success rate of total joint replacement (TJR) surgeries, longer human life expectancy and implantation in younger patients has driven bioengineers from original implant concerns to consideration of long-term 441
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limitations of the materials. The ideal material for TJR prostheses should exhibit a biocompatible chemical composition to avoid adverse tissue reactions, an excellent resistance to degradation or corrosion in the human body environment, acceptable strength to sustain the cyclic loading endured by the joint, a low modulus to minimize bone resorption, and a high wear resistance to minimize the generation of debris, Hoeppner and Chandrasekaran (1994). Ramakrishna et al. (2001) summarized several important factors to be considered in selecting a material for a biomedical application. They listed various biocomposites used in a human body. It is observed that carbonfibre-based polymer composites are mainly used in total joint, finger and disc replacement surgeries. Liang and Shi (2004) and Park et al. (2000) suggested various combination of materials used in the human body tribological system, which are listed in Table 18.1. It gives comprehensive details of Table 18.1 Summary of tribo-biomaterials (Liang and Shi, 2004 and Park et al., 2000) Major properties, description
Materials
Application
Metal Brass, stainless steel, Ni-plated steel, Zn-plated steel
Inserts
Wear and corrosion resistance
Alloys Titanium alloys, Ti–Al–V alloy, Co–Cr–Mo alloy, Co–Cr alloy, tantulum
Total joint replacement, stem, ball head, cup, porous coating and metal backing
Wear and corrosion resistance, heavy, hard, low stiffness
Inorganic Diamond-like carbon
Biocompatible coating
Reduced friction and increased wear resistance
Ceramics Al2O3, ZrO2, Si3N4, SiC, B4C, quartz, bioglass, sintered hydroxyapatite
Bone joint coating, ball, cup
Wear and corrosion resistance
Polymers UHMWPE, PTFE, Polyglycolic acid, Polyurethan, PMMA
Joint socket, cup, interpositional implant, temporomandibular joint, joint bone, leaflet heart valve, bone cement fixation
Wear, abrasion and corrosion resistance Low coefficient of friction Elastic with less wear Highly biocompatible, high strength, and dynamic range of breathability
Composites Specialized silicone polymers, carbon fibre, glass fibre and Kevlar fibre-reinforced composites
Bone joint
Wear, corrosion and fatigue resistance
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different metals, metal alloys, inorganic compounds, ceramic compounds, polymer and polymer composites used in various joints of the human body where tribological characterization of the material plays an important role in the life of the component. Since the 1960s, the tribological characterization of implant materials that are generally used for hip, knee and shoulder joints, mated with ceramic and metallic materials have been studied. It is well known that the long-term usage of prosthesis materials in total joint replacements is limited by the wear debris which is associated with an incidence of non-specific pain and prosthesis loosening. Thus, attempts are made to understand the problem of total joint replacement by studying the different influencing variables on tribological properties of biocomposites used in a human body. Synthetic biomaterials such as stainless steel, titanium alloys, polymers, ceramic and composites undergo degradation through fatigue and corrosive wear because of load bearing and salty environment of the human body. Despite the success of surgical implants such as artifical hip and knee joints, the materials used still do not meet specific requirements. The potential causes of failure for total hip arthroplasty are deficiency in design of size and shape of the device for a particular patient, surgical problems, host abnormalities or diseases, infection and materials fracture, wear and corrosion, Wang et al. (1997). Amongst these causes, this chapter describes the tribological characterization of polymer composites that are used in biological systems mainly in synovial joints.
18.2
Experimental consideration on tribological characterization of composites
Stachowiak et al. (2004) summarized the main factors such as a list of operating variables, material properties, testing parameters and lubrication characteristics to be considered in order to achieve a close agreement between test and real tribological data. The parameters shown in Table 18.2 need to be considered to characterize the tribological properties of biocomposites under real conditions. Though many tribometers are currently available, it is essential to accurately select the most appropriate type of tribometer to achieve a particular type of wear mechanism. Stachowiak et al. (2004) listed various types of tribometer to derive the required mechanism of wear and they are listed in Table 18.3. Selection of a tribometer mainly depends on the type of wearing contact and, therefore, operating conditions need to be studied. They also summarized various techniques available to measure the wear volume of the test sample and discussed the advantages and limitation of each techniques which are shown in Table 18.4. In order to develop a new material for a human tribological system, the advantages and limitations of the existing bearing materials used in the system have to be known and these are
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Table 18.2 Commonly used parameters in the characterization of tribological contacts (Stachowiak et al., 2004) Variable and related parameter Operating parameter Material parameter
Materials parameter for abrasive and erosive wear tests Environmental parameter
Lubrication parameter
Load, contact stress, sliding speed, sliding distance Hardness, shear strength, micro-structure, toughness, limiting strain and average grain size, melting point, glass transition temperature and softening point, thermal diffusivity, thermal shock resistance, thermal conductivity, specific heat, electrochemical potential and current density Grit hardness, mineral type, size and shape characteristics of grits, fractal dimension, spike parameter, acidity of carrier fluid for slurries Relative humidity, local air pressure, partial pressure of oxygen and other active gases, radiation level in nuclear environments Viscosity, flow rate, pressure, velocity, thermal conductivity, thermal diffusivity, specific heat, chemical reactivity, acidity, boiling point, solidification point, heat of oxidation and water/ oxygen solubility
shown in Table 18.5, McKellop (2001). In order to reduce the functional limitations and enhance the merits of the bearing couples, new biocomposites are being developed and aspects of their tribological and biological systems are being studied.
18.3
Tribology of polymer composites
Improvements in tribological characteristics of polymer can be achieved by the deliberately engineered addition of strengthening and lubricating agents. There are a wide variety of composite materials available, but the main types of polymer composites are: a) polymer containing a lubricating polymer, b) polymer containing metal/inorganic/polymeric powder/carbon nanotubes, and c) fibre-reinforced polymer composites.
18.3.1 Polymer containing a lubricating polymer When polyamide-6 (PA-6)–UHMWPE composite was sliding against a counterface, the tiny UHMWPE particles in the PA-6 matrix may play a role of a lubricating agent and the bond between the UHMWPE and the PA-6 may prevent UHMWPE particles being transferred into the counterface. The relative role of contact pressure is reduced by the wear of the composites owing to changes in the physical and mechanical properties of
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Table 18.3 Prevailing wear conditions and type of tribometer (Stachowiak et al., 2004) Mechanism of wear
Tribometer
Abrasive wear Erosive wear
Apparatus involving abrasive paper or bed of sand Specimens mounted in steam of air or liquid jet mixed with abrasive particles Specimens mounted in steam of fluid or else mounted on vibrating platform immersed in fluid Specimens slid against a moving counterface (which may also be a specimen) either by rotation or reciprocating movement. Lubricant supply system fitted for lubricated tests Specimens and moving counterface enclosed in chamber fitted with heating, refrigeration and/or vacuum pumping system to maintain a specialized environment Specimens slid against the counterface at very small amplitude. Apparatus can be fitted with chamber for specialized environments Specimens in form of rollers and spheres are constrained to move at specified speeds. Apparatus can be fitted with chamber for specialized environments Apparatus containing hammer to impact wearing specimen. Can be fitted with enclosing chamber for non-ambient tests Cutting tool as a specimen. Cutting performed at high surface speed. Specimen immersed in molten test material and rotated to accelerate wear Idealized conformal and non-conformal contacts devised with purpose of displaying mechanism involved in the formation of lubricating films Apparatus involving sliding wear at low velocities allowing the friction coefficient during the tests to be monitored
Cavitational wear Dry or lubricated sliding wear in ambient conditions Dry or lubricated wear under non-ambient conditions
Fretting / fatigue
Wear under combined rolling and sliding or pure rolling Impact wear
Diffusion or solubilization wear
Hydrodynamic and elastohydrodynamic lubrication Boundary and solid lubrication
the composites, Liu et al. (2001). It was concluded that contact pressure is the main controlling parameter followed by sliding distance on the wear of the PA-6/UHMWPE composite. The results from Liu et al. (2004) indicate that when polypropylene (PP) concentration is lower than 25 wt.%, the friction coefficient and wear scar width of the UHMWPE/PP blend was reduced with increasing PP content. A further increase in PP content results in the friction coefficient increasing abruptly, and the wear scar width of the blends increases linearly. During sliding, when the temperature of the UHMWPE contact exceeds a critical value, wear proceeds in a series of discrete steps caused by the sudden loss of a molten or softened layer of polymer, Barrett et al. (1992). Owing to
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Table 18.4 Relative merits of wear measurement techniques (Stachowiak et al., 2004) Techniques
Advantages
Disadvantages
Weighing
Simple and accurate
In situ measurement of change in length of worn surfaces Stylus profilometry
Accurate and allows continuous record of wear rates Very accurate; gives distribution of wear between specimens
Data corrupted by displaced or transferred materials No discrimination between wear of either specimens
Laser scanning profilometry
Very accurate and fast; gives distribution of wear between specimens Simple and rapid
Optical profilometry
Surface activation
Ultrasonic interference
In situ measurements of wear in closed machinery; possibility of simultaneous measurements of wear rates of various parts Sensitive to small changes in dimension
Slow and mostly suitable for the end of the test; expensive equipment is required Expensive equipment is required Method impossible when specimen has complex shape or its shape is distorted by wear or creep under load Inaccurate and difficult to ensure safety of personnel
Specialized technique that requires expertise
the low value of surface friction coefficient of UHMWPE, the surface temperature of the composite is less, which indicates better wear resistance of the composites. Partly oxidized transfer film was formed on the counterface when sliding with UHMWPE and fine powders was found on the counterface sliding with UHMWPE/PP. Debris of UHMWPE/PP composites existing between the two friction surfaces acts as a rolling lubricant, thereby reducing the friction coefficient and wear to a much lower level. Hashmi et al. (2001) studied the effect of applied pressure on the wear rate of PP, UHMWPE and their blends. It is observed that the wear volume increases with applied pressure in all cases and the maximum wear rate was observed in PP samples. UHMWPE shows the minimum wear loss at various pressures. It is observed that the wear rate of PP is more sensitive to pressure than UHMWPE. Addition of 15 wt.% of UHMWPE in PP demonstrated almost similar wear behaviour of pure UHMWPE. It is interesting to note that a small weight fraction of UHMWPE in PP improves the wear resistance of PP to a significant extent. This improvement may be
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Table 18.5 Advantages and disadvantages of current bearing choices (McKellop, 2001) Bearing combination
Advantages
Disadvantages
Alumina on alumina
Usually very low wear, high biocompatability
Cobalt–chrome on cobalt–chrome
Usually low wear, can self-polish moderate surface scratches Some additional protection against third-body abrasion Lower wear of polyethylene than with conventional metal–polyethylene Some additional protection against third-body abrasion No short- or long-term oxidation
Sometimes high wear, component fracture Higher cost Technique-sensitive surgery Question of long-term local and systemic reactions to metal debris and/or ions Hardened layer can wear off Higher cost
Hardened cobalt– chrome on polyethylene Ceramics on polyethylene
Polyethylene sterilized with ethylene oxide or gas plasma Polyethylene sterilized with gamma in low oxygen Cross-linked, thermally stabilized polyethylene
Some cross-linking, some wear reduction
Minimal polyethylene wear rate, No short- or long-term oxidative degradation
Component fracture Difficulty of revision (e.g., if Morse taper is damaged) Higher cost
No cross-linking so polyethylene wear not minimized Polyethylene wear not minimized Residual free radicals (long-term oxidation) Newest of low-wear bearing combinations, only early clinical results available, questions remain as to optimum cross-linking level and optimum method for thermal stabilization
attributed to the morphology of the blend. When sliding starts, the increased temperature at the contact surface owing to frictional heat softened PP, which deforms and loses its structural integrity more easily than UHMWPE, which does not flow owing to highly entangled molecular chains. As the softened portion of the sample, PP is removed and transferred to the counterface, more UHMWPE particles are exposed and the new surface of PP having a low surface temperature comes into contact with the counterface. Owing to the low coefficient of friction, UHMWPE reduces the frictional heat and temperature, hence providing wear resistance to the blend. The microstructure of a composite, which can be modified by the processing
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conditions, is a major variable in the outcome of its tribological behavior as reported by Sung and Suh (1979) and Peers et al. (2006). Liu et al. (2006) studied the effect of contact pressure on the wear rate of PA, UHMWPE, and PA/UHMWPE under dry and lubricated sliding conditions. It is observed that both contact pressure and lubricating condition have strong effects on the wear loss of the specimens. The specimens show higher wear loss under higher contact pressure at dry-sliding conditions. The dry-sliding wear rate of the specimens is about five to six times higher than that of lubricated wear. The wear rate under 1 MPa pressure is about 75–85% higher than the sliding wear under 2.5 MPa and dry-sliding conditions. However, it was reduced to about 48–60% under lubricated conditions. It is observed that UHMWPE demonstrated the highest drysliding wear rate, under both 1 and 2.5 MPa of contact pressure, followed by PA/UHMWPE and PA specimens; whereas, PA shows the highest wear rate in lubricated sliding condition, followed by UHMWP and PA/ UHMWPE. It is also found that the specific wear rate under 2.5 MPa is less than that of 1 MPa contact pressure, although the wear volume loss under 2.5 MPa is greater than that of 1 MPa. This suggested that increasing contact pressure by 2.5 times did not result in an increase of wear loss by a factor of 2.5. The tribological characterization of HARUHMWPE (Hydroxyapatite+ UHMWPE) composites with 5, 10 and 30% hydroxyapatite reinforcements and also with an unfilled UHMWPE under various loading conditions was studied by Kalin et al. (2002). It was shown that the wear resistance of the hydroxyapatite-reinforced UHMWPE composites are noticeably higher than that of the parent polymer. Furthermore, the wear resistance of the composites increases to an extent with the increase in hydroxyapatite content.
18.3.2 Polymer containing a metal, inorganic, and polymeric powder and carbon nanotubes It is generally observed that wear resistance of the polymer is improved with an addition of a reinforcing element such as fillers and plasticizers. The carbon nanotubes (CNTs) are found to produce a valuable increase in wear resistance and a reduction in the friction coefficient when blended with polymers. It is found that the abrasive wear of the filled polymer composite depends on the particle size. For nanoscale level fillers, the nanocomposites showed an average weight loss about half of that of pure PMMA by adding only 2 wt.% of CaCO3 Avella et al. (2001). Blending alumina nanoparticles of 38 nm with PTFE resulted in a very large reduction in the wear rate of about 600 times and only a minor increase in the friction coefficient, Sawyer
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et al. (2003). When blended with silica, silica carbide, silicon nitride and zirconium oxide nanoparticles, PEEK composite exhibited a much lower frictional coefficient and wear rate than pure PEEK, Wang et al. (1996 and 1997). A thin, uniform and tenacious transfer film was observed on the counterface, which improved the tribological behaviour of the nanoparticles-filled composites. Wang et al. (1997) studied friction coefficient and wear rate of nanosized SiO2 filled-PEEK composites as a function of SiO2. It is observed that the friction coefficient decreased sharply when the SiO2 content was below 10 wt.%. Then it decreased gradually with increasing SiO2. The wear rate decreased sharply when SiO2 content was less than 2.5 wt.%. There was a linear decreasing trend when the SiO2 content increased from 2.5 wt.% to 7.5 wt.%. When the filler content was above 7.5 wt.%, the wear rate gradually began to rise with increasing SiO2, although it was still considerably lower than that of the unfilled PEEK. From friction coefficient and wear rate point of view, it might be rational to suggest that the optimal content of SiO2 in the composite should be 7.5 wt.%. Liu et al. (1999) studied abrasive wear behaviour of various composites of UHMWPE reinforced with quartz powder under dry conditions. The main objective was to study the influence of filler particle size, sliding speed and abrading particle size on the abrasive wear performance of UHMWPE composites. The presence of filler in UHMWPE tends to reduce the formation of deep furrows. It was observed that the wear resistance of the composites was greatly improved after reinforcement with quartz powder, an increase of about two to four times compared with unfilled UHMWPE was observed. When quartz powder was mixed with UHMWPE, it leads to increased surface hardness of the specimen and an enhancement of ploughing and cutting resistance. At the same time, there was sufficient resistance to shear of the filler–matrix interface to prevent easy removal of the filler particles. The smaller fillers may readily be transferred into the transfer film formed on the abrasive section, therefore the effect of wear resistance improvement in UHMWPE reinforced by smaller fillers is not as obvious as that for the larger filler particles, Papanicolaou and Bakos (1992) and Liu et al. (1997). The variation of weight loss of the UHMWPE/20 vol.% quartz composites was studied when organosiloxane content varied from 0 to 1.0 phr (part of reagent per hundred parts of UHMWPE) by Xie et al. (2003). When both distilled water and fetal bovine serum were used as lubricants, the results showed that the weight loss of the UHMWPE/quartz composites decreased with the content of organosiloxane up to about 0.5 phr. The organosiloxane-induced coupling between UHMWPE and quartz and cross-linking of UHMWPE leads to an improvement in the composite’s wear resistance. However, with a further increase in the content of organosi-
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Wear rate (× 10–8 mm3 N–1 m–1)
loxane, the high density of cross-linking leads to a deterioration in the wear resistance. Thus, the weight loss of the UHMWPE/quartz composites increases significantly. Therefore, the organosiloxane content at 0.5 phr can produce maximum wear resistance for the UHMWPE/quartz composites. Similar findings about other particulate composites, such as PTFE/alumina and PTFE/CoCr have also been reported by Phipatanakul et al. (1998). Xue et al. (2006) observed that the wear rate of the UHMWPE/HDPE composite can be significantly reduced by adding carbon nanotubes (CNTs). The wear rates of all CNT composites, both pre-treated CNTs and untreated CNTs, decreased with increasing CNTs content. The addition of only 0.5 wt.% CNTs to the 80% UHMWPE/20% HDPE COP blend caused about a 50% reduction in the wear rate. In addition, the composites reinforced with untreated CNTs had a better wear performance than the composites with pre-treated CNTs. The UHMWPE/HDPE composite without CNTs had a slightly higher wear rate (44.3 × 10−8 mm3 N m−1) than the pure UHMWPE (40.8 × 10−8 mm3 N m−1), as shown in Fig. 18.1. HDPE, owing to its higher crystallinity and lower wear resistance, helps to increase creep performance and reduce wear strength of UHMWPE. However, the CNTsreinforced composites gave better wear results than pure UHMWPE and the UHMWPE/HDPE blend without CNTs. The above results were obtained when the test sample was sliding against the austenitic stainless steel X5CrNi18-10. To check that these findings hold for the usual martensitic bearing steel 100Cr6, pure UHMWPE, cop, cop2%p, and cop2%u were tested against 100Cr6 balls. Fig. 18.1 compares these results with the respective data for X5CrNi18-10. The wear rates of specimens slid against 100Cr6 are higher than those against X5CrNi18-10, except for pure UHMWPE.
90 80
X5CrNi18-10 100Cr6
70 60 50 40 30 20 10 0 UHMWPE
COP
COP + 2.0% pretreated CNT
COP + 2.0% untreated CNT
18.1 Specific wear rates of four specimens against X5CrNi18-10 and 100Cr8 (Xue et al., 2006).
Tribology of biocomposites
Wear coefficient
1.4 × 10–4
1.0 × 10–1
1.2 × 10–4
8.0 × 10–2
1.0 × 10–4 8.0 × 10–5
6.0 × 10–2
6.0 × 10–5 4.0 ×
10–2
4.0 × 10–5
2.0 × 10–2
2.0 × 10–5
0.0
C
N
T23
2. 2 C m N T96 5. 6 m
.4
m
m .2 N C
N C
T86
-T43
2. -2 3
W PE
U
H
M
M H
2
m .4
m -8 6
.2 -4 3
W PE
PE W
m
0.0
U
M H U
Wear coefficient (mm3 N–1 m–1)
Wear volume (mm3)
1.6 × 10–4
Wear volume
1.2 × 10–1
451
Sample and sliding distance
18.2 Variation of wear volume and wear coefficient of UHMWPE and composites against sliding distance.
The variation of wear volume of UHMWPE and UHMWPE/0.2 wt.% CNT composites with sliding distance is shown in Fig. 18.2, Kanagaraj et al. (2001). It is observed that the wear volume decreases by adding CNTs. The decrease of wear volume is observed to be 28% at 43.2 m, 28% at 86.4 m and 35% at 232.2 m, owing to the enhanced interfacial strength between the polymer and CNTs and other mechanical properties, resulting in a good load transfer effect to the nanotube network from polymer. The variation of the coefficient of wear resistance of the test samples is also shown in Fig. 18.2. It is observed that the wear coefficient decreases with reinforcement of CNTs in UHMWPE and it follows the same trend of the wear volume. It also decreases with the increase of sliding distance approximately following a linear model. It follows that the suggested relationship between wear volume or wear coefficient and sliding distance, Liu et al. (2006). When CNTs were introduced into the UHMWPE matrix, the surface hardness was enhanced, as was the ploughing and cutting resistance of the specimen and they prevented easy removal of the wear particles. The interesting phenomenon is the restriction of plastic displacement of the composite by the counter face which was the result of the plasticity index and plastic contact between the composite and counterface.
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Wear coefficient (10–4)(mm3 N–1 m–1)
8 7 6 150 N
5
120 N
4
90 N 3 60 N
2
30 N 1 0 0.0
0.2
0.4 0.6 0.8 CNT concentration (wt%)
1.0
18.3 Variation of wear coefficient of CNT/HDPE composites against CNT concentration and load.
The wear coefficient of HDPE/CNT composites as a function of CNTs concentration and load is shown in Fig. 18.3, Kanagaraj et al. (2008). It is observed that the wear volume decreases with an increase of CNT concentration up to 0.5 wt.% of CNTs, which is owing to the good interface between the polymer and CNTs; this results in a good load transfer effect to the nanotube network from polymer. It is observed that there is no significant change of wear volume with CNTs loading at 30 N. However, the wear volume increases with the increasing applied load. It is observed that the trend of the wear coefficient curve varies with load. With an increase of load, the wear coefficient increases with CNT concentration up to 0.5 wt.% for 120 and 150 N. However, the opposite trend was observed for the same loading condition with further increases of CNT concentration. A sine wave trend was observed for the normal loading of 60 and 90 N with an increase of CNT loading. It is also observed that 0.5 wt.% of CNT with HDPE composite produces critical characteristics that affect tribological behaviour. From the results of Anderson et al. (2002) shown in Fig. 18.4, it is observed that Al–Cu–Fe filled UHMWPE disks show enhanced wear resistance to volume loss, as compared with unfilled UHMWPE and alumina-filled UHMWPE. Al–Cu–Fe/UHMWPE samples showed approximately a 35% decrease in volume loss compared with the UHMWPE samples. The volume loss of the unfilled UHMWPE was the result of both deformation and removal of the polymer during wear. The improved wear resistance of
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0.8 0.7
Volume loss (mm3)
0.6 0.5 0.4 0.3 0.2 0.1 0 Alumina/UHMWPE
UHMWPE
Al–Cu–Fe/UHMWPE
18.4 Volume loss from UHMWPE, Al–Cu–Fe/UHMWPE, and alumina/ UHMWPE samples.
the Al–Cu–Fe/UHMWPE composites has been attributed to the high hardness, high Young’s modulus, and low coefficient of friction of the Al–Cu–Fe quasi-crystalline filler and the increased strength of the polymer composite compared with unfilled UHMWPE.
18.3.3 Fibre-reinforced polymer composites Chopped-fibre composites The limits of sliding speed and contact stress can be raised and the wear rate be reduced by the incorporation of chopped fibres, Friedrich (1986). He concluded that chopped fibre reinforcements are effective in reducing wear provided that there is a strong adhesion between fibre and matrix. The major problem with fibre reinforcement, in particular with chopped fibre, is that the wear resistance depends more on fracture occurring between the fibre and the matrix than on the bulk properties of materials, Bahadur (1991). This is because wear occurs on a microscale where there is no mechanism of material strengthening similar to bulk fracture which involves cracks passing through many fibre–matrix interfaces. The optimum concentration of chopped fibre for maximum wear resistance is found to be close to 10 vol.% which is less than the concentration required for optimum toughness of the composite, Bahadur (1991). The inherent brittleness of most of the reinforcement fibres results in rapid fibre damage under abrasive
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wear, Bijwe et al. (1990). It was found that abrasive wear is more severe when the short chopped fibre is used rather than small sphere and that the improvement in adhesion between the filler particles and the polymer matrix results in the increase of abrasive wear resistance, Yang et al. (1991). The steady-state wear rates of UHMWPE/HDPE composites were compared by Jacobs et al. (2000) and it is observed that UHMWPE gives the best wear results but that the composites with more than 30 vol.% UHMWPE have the highest wear resistance. Nevertheless, the scatter of the fibre-reinforced composites is smaller than for the pure HDPE and is of the same order as that for UHMWPE. UHMWPE-fibre-reinforced HDPE composite promises to be a good substitute for UHMWPE as a bearing material. It combines a wear resistance similar to UHMWPE with very good mechanical properties. The wear resistance of the PE/PE composites was found to increase continuously with UHMWPE fibre content. Lu and Friedrich (1995) observed the strong influence of carbon fibre on the wear rate of CF/PEEK composites. The specific wear rate was reduced by more than one order of magnitude when at least 10 vol.% of short carbon fibres was added. Above this fibre content, further improvement was only slight. The correlation between the wear behaviour and microstructure of graphite/PTFE composites was studied over a wide range of graphite content from 0 to 50 vol.%, Yan et al. (1996). It was found that the wear behaviour of graphite/PTFE composites corresponds with the physical properties of imperfections and this is reflected by the interfacial properties between crystalline and amorphous regions in PTFE. Bahadur et al. (1994) observed that wear was found to reduce considerably when nylon was reinforced with carbon fibre, and the wear reduction was even more significant when carbon fibre was used for reinforcement along with CuS as filler. The effects of various short-fibre reinforcements on the sliding wear characteristics of the various polymer composites were studied by Friedrich (1986) and Friedrich et al. (1995). It is observed that the carbon fibres often improve wear resistance of a material by a factor of 5 or more than the same volume fraction of glass fibres. From recent studies carried out with PAN and Pitch based carbon fibres by Friedrich et al. (1992), it is concluded that the composite’s wear resistance improves with decreasing sensitivity of the fibres reaction against fibre–fibre friction and early breakage. In addition, a strong bonding to the matrix helps to keep pieces of broken fibre on the composite surface, thus preventing early formation of third body abrasives and enhanced wear. Friedrich et al. (1991 and 1992) observed that the effect of counterface roughness on the wear rate of shortfibre-reinforced PEEK is not as pronounced as for the virgin matrix. A summary about how short fibres and internal lubricants can affect the specific wear rate and the coefficient of friction at room temperature is discussed in detail by Friedrich et al. (1995).
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Unidirectional and woven fabric reinforcements Fibre orientation is critical to the tribology of polymer composites. Wear of the matrix and fibre proceeds at the same rate until the depth of about half of the fibre diameter is worn away and the fibres start to detach in short segments from the matrix. It is observed that the wear debris originating from the fibres ranges from fine powder to complete segments of fibre as the wear proceeds. In contrast, wear debris from the matrix tends to be uniformly fine. It is possible that a fine transfer film of the matrix polymer may cover the exposed fibre and reduce the overall coefficient of friction, Gong et al. (1989). The tribological behaviour of a unidirectional and oriented E-glass-fibre-reinforced epoxy composite against a cylindrical counterface, friction and wear were found to be improved when sliding took place against either a clean or a wet counterface, Tayeb and Gadebrab (1996). Polymer composites with the normal fibre orientation give a low wear rate since partially worn fibres remain firmly attached in the matrix but at the risk of sudden seizure because the exposed normal fibres tend to gouge into the counter face and initiate severe wear mechanism of seizure, Tsukizoe and Ohmae (1986). During the process of wear, the fibres are subjected to repeated bending which causes them to gradually debond from the matrix, Jacobs (1991). A simultaneous process of cracking and fragmentation at the fibre ends allows material to be eventually released as wear debris. Ovaert and Cheng (1991) studied tribological characterization of PEEK-CF composites with the fibres oriented parallel to the sliding direction and it is observed that the wear of the composites depends on the counterface surface. Both wear and the friction coefficient decreased as a result of the addition of carbon fibre. The results from Zhang et al. (1991) show that the friction coefficient and the specific wear rate of the PEEK composite against a steel ring were decreased significantly under sliding, which comes from the enhancement of the interface strength of carbonfibre-reinforced PEEK composites. Dangsheng (2005) studied the effects of carbon fibre content on tribological properties of UHMWPE at different concentration. Figure 18.5 shows the effect of load on friction and wear of the 20% CF-UHMWPE composite under dry sliding. The decrease in the friction coefficient that is observed in CF-reinforced UHMWPE under increased loading is attributed to surface softening arising from frictional heating. The wear volume loss increased with load, because of the surface softening from frictional heating under high load. Dangsheng (2005) also observed that the wear volume loss of pure UHMWPE was the highest under both dry and distilled water lubrication conditions. The volume loss under dry sliding was higher than that under distilled water lubrication for each sample. The volume loss decreased with the content of CF, and the
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2.5
0.30 2.0 0.28 1.5 0.26 1.0 0.24
Wear volume (mm3)
Friction coefficient
0.32
0.5 0.22 98
196 294 Applied load (N)
392
0.0
18.5 Effect of load on friction and wear of 20%CF-UHMWPE under dry sliding (Dangsheng 2005).
rate of decrease of wear of samples at CF content from 0 to 20% was faster than that from 20 to 30% under both the dry and the distilled water lubrication conditions. It can be seen in Fig. 18.6 that the wear volume increases with sliding distance and at a larger sliding distance wear rate falls, Roychowdhury et al. (2004). After the initial film formation, the wear rate at higher sliding distances would fall depending on the nature and orientation of the transfer film, which depends on the counterface roughness. It is expected that roughness normal to the sliding direction is an important factor in deciding the wear rate, Cooper et al. (1993). Wear rates of HDPE and KFRP10 in Fig. 18.6 remain nearly constant. The HDPE composites with the exception of KFRP10 show relatively higher levels of wear compared with the unfilled HDPE and UHMWPE. The wear debris for composites is partly composed of torn out fibres and they act as abrasive elements. The generation of wear particles from the composite materials is probably the result of breakage of the bond between the matrix and the fibre. This may be caused by the rise in temperature and roughness of the counterface. Similar observations have been made by Ramasubramanian et al. (1993) and Stachowiak (1993) during wear characterization of the composites. The primary mechanism of polymer composites wear against metallic counterface is the result of adhesion, transfer film mechanism, abrasion caused by pulled out fibres from the matrix and the effect of counter surface roughness. The wear of UHMWPE was found, unexpectedly, to be higher than HDPE and KFRP10. From these results, KFRP10 seems to be a potential material for acetabular
Tribology of biocomposites Sliding velocity = 0.24 m s–1
HDPE UHMWPE KFRP10 KFRP15 KFRP20 CFRP20
12
10 Wear volume (mm3)
457
8
6
4
2
0
0
2
4
6 8 Sliding distance (km)
10
12
18.6 Wear volume of polymer and composites plotted against sliding distance (Roychowdhury et al., 2004).
cups since it shows favourable tribological characteristics along with the advantage of a composite structure. Figure 18.7 shows the effect of carbon fibre content on the wear resistance of the PAN-based carbon fibre/PEEK composite cups against alumina heads, Wang et al. (1998). Both the 20 wt% and the 30 wt% fibre-reinforced composite cups showed more than an order of magnitude decrease in the wear rate compared with UHMWPE cups. However, the unreinforced PEEK cups exhibited an extremely high wear rate which was eight times higher than UHMWPE cups. The wear rate was decreased drastically with an increase of the carbon fibre content. Increasing the fibre content from 0 to 20% resulted in a 100-fold reduction in the wear rate. From 20 wt% to 30 wt%, the wear rate was decreased by a factor of 2 again. Figure 18.8 shows the wear behaviour of the 30 wt% PAN-based carbon fibre/PEEK composite with that of a 30 wt% pitch based carbon fibre/PEEK composite against various counterface materials, Wang et al. (1998). Both the PANbased and the pitch-based composites showed significantly higher wear rates against CoCr heads than against zirconia or alumina heads. For instance, for the same feloral head, the PAN-based composite showed a higher wear rate than the pitch-based composite. These results indicate that CoCr head is not suitable for articulation against either one of the two composite materials while alumina heads work with both composite
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Wear rate (mm3/106 cycles)
1000 N=2
100
10
1 UHMWPE
20 wt% PAN– CF/PEEK Acetabular cup material PEEK
30 wt% PAN– CF/PEEK
18.7 Volumetric wear rates of the PAN-based carbon fibre/PEEK composite cups against 32 mm alumina heads as a function of the fibre content (Wang et al., 1998).
Cup wear rate (mm3/106 cycles)
50
UHMWPE 30 wt% PAN–CF/PEEK 30 wt% Pitch–CF/PEEK
40
N=2
30
20
10
0 CoCr
Alumina
Zirconia
Femoral head material
18.8 Wear rates of the composite cups as a function of carbon fibre type (PAN vs pitch) and femoral head material (Wang et al., 1998).
materials. Zirconia heads work much better with the pitch-based carbonfibre-reinforced PEEK composite cups. Since zirconia heads possess a far better fracture resistance than the alumina heads and the pitch-based carbon fibre composite/zirconia combination yielded the lowest wear rate. The experimental results indicate that carbon-fibre-reinforced PEEK composite materials can be used as a viable bearing surface for acetabular cups
Tribology of biocomposites
459
in a total hip replacement. It is possible to produce a composite acetabular insert that offers far superior wear performance over conventional UHMWPE. Potential wear rate reductions of almost two orders of magnitude are achievable with a pitch based carbon fibre PEEK composite cup in articulation against a zirconia head. Tribological properties of plain-weave fabric-reinforced poly(vinyl butyral)-modified phenolic resin–matrix composites with three different reinforcing fabric materials as components (E-glass, high-strength carbon and Kevlar-49) were investigated, against a cast iron disc under dry sliding conditions, Vishwanath et al. (1993). Of the three reinforcing fabric materials, the composite made from glass fabric reinforcement turns out to have poor wear resistance, whereas the best wear resistance was discovered in the Kevlar-49 fabric-reinforced composite. To develop a high-performance composite with optimum wear resistance, various unidirectional composites were tested as a function of matrix, type of fibre reinforcement, and orientation of the fibres relative to the sliding direction, Cirino et al. (1987 and 1988). Various matrices, e.g. epoxy resin, PA, and PEEK and three typical fibre materials, e.g. glass, carbon, and aramid, were chosen for the investigation and the results are shown in Table 18.6. It is observed that PEEK composites have a higher wear resistance than the epoxy resin and carbon fibre composites give better results than glass fibres, with aramid fibres somewhere in between in any sliding direction. The wear resistances of these materials can be utilized to systematically design a composite material with a generally good wear performance. An optimum wear resistance can be expected from a composite consisting of a PEEK matrix and carbon fibres parallel to the sliding direction. If multidirectional sliding is required, additional carbon fibres in the antiparallel orientation would be helpful. Table 18.6 Wear resistance ratios of various unidirectional composites, as normalized to the epoxy resin (w ˙ s−1 = 1.79 × 104 m mm−3) Fibre
Orientation1
EP1
EP2
PA66
amPA
PEEK
Carbon
N P AP N P AP N P AP
65.4 99.8 44.1 0.5 2.1 1.1 122.4 44.1 46.3 1
102.4 197.5 90.7 – – – 104.3 45.7 75.8 1
41.3 122.9 80.2 1.8 16.3 8.9 145.5 60.8 73.7 27.7
117.7 128.6 57.0 – – – 145.5 65.1 102.4 0.07
65.1 158.1 94.8 5.1 5.5 2.0 61.4 46.1 44.1 11.8
Glass
Kevlar
Unreinforced 1
N, Normal; P, Parallel and AP, Antiparallel to the fibre direction.
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Biomedical composites
PEEK with carbon fibre woven fabric reinforcement resulted in even better wear factors than the two basic orientations of the unidirectional material. This synergistic effect is a result of fibre interlocking in the contact area owing to the woven structure of the reinforcement, Mody et al. (1988). The form of reinforcement and its orientation with respect to the sliding direction played a vital role in determining the degree of enhancement and the woven fibres are proving to be far better than unidirectional fibres. The antiparallel orientation for the unidirectional composite was the least wear resistant compared with the other two orientations. Although the parallel and normal orientations were close in wear rates, damage was less in the former case which proved to be the most wear resistant orientation. Wear rates for the woven composite were an orderof-magnitude lower than for the unidirectional composite. However, the parallel orientation which comprised 80% in-plane fibres parallel and 20% in-plane fibres transverse to the direction of sliding constituted the most wear-resistant composite. Generally, the wear behaviour of composite specimens possessing normal oriented fibres was worst.
18.4
Conclusion
An overview of various polymer composites has been presented showing which materials provide better wear resistance. A summary is provided of the results of tribological characterization of composites such as a polymer containing a lubricating polymer, polymers containing metal, inorganic, polymeric powder and carbon nanotubes, and chopped fibre, unidirectional and woven fabric composites used in biological system.
18.5
Acknowledgements
The authors express their gratitude to scientists whose contribution are referred to in this article.
18.6
References
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19 Fatigue behaviour of biocomposites A. P E G O R E T T I, University of Trento, Italy
Abstract: Fatigue fracture has been identified as one of the major causes of implant failure of medical devices, and it represents a serious limitation for the intended usage of polymer composites for hard and soft tissue applications. This chapter presents an overview on the fatigue behaviour of polymeric composite biomaterials for bone repair and replacement, dental applications, joint replacement, spine surgery, tendons and ligaments augmentation devices and vascular grafts. Some recent advances in testing techniques are outlined and the future development of fatigue-resistant biocomposites is discussed. Key words: polymer composites, biomaterials, fatigue, fracture crack propagation.
19.1
Introduction
Materials of natural or man-made origin used to direct, supplement, or replace the functions of living tissues of the human body are generally called ‘biomaterials’ (Black, 1992). Over the last 30 years, polymer composites for biomedical applications have gained an increasing importance as structural biomaterials (Migliaresi and Pegoretti, 2002; Ramakrishna et al., 2001). One of the main reasons for the introduction of polymer composites in biomedical applications is that their properties can be varied and tailored to suit the mechanical and physiological conditions of the host tissue. In fact, several different polymeric matrices and reinforcing elements such as fibres or particles can be combined in various volume fractions and arrangements thus generating materials with an extended range of mechanical properties. Fibre-reinforced polymer composite materials have been proposed to replace metallic alloys in certain medical devices because of their lightness, tunable mechanical properties, and radiolucency (Baidya et al., 2001). Moreover, composite materials offer a unique chance to tune the material properties in different regions of a given component thus opening challenging opportunities for integrated biomimetic approaches (Kikuchi et al., 2004, Zhou, 1998). The current utilization of polymer composites as biomaterials encompasses both hard and soft tissue applications. Among hard tissue applications polymer composites have been successfully used for bone fracture repair, bone plates, intramedullary nails, spine instrumentation, joint 465
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replacements, total hip and knee replacements, bone cements, bone replacements (synthetic bone grafts), and dental applications. Applications of polymer composites as biomaterials for soft tissue applications include bulk space fillers, encapsulants and carriers, such as wound dressings, ureter prostheses and catheters, functional load-carrying and supporting implants such as tendons and ligaments, vascular grafts and hernia patches. Other biomedical applications of polymer composite materials include prosthetic limbs and medical instrumentation. During daily activities such as standing, sitting, walking, jogging, stretching and climbing, hard and soft tissues are normally subjected to repetitive stresses that fluctuate in time (Black, 1992). Fatigue fracture has been identified as one of the major causes of implant failure of medical devices (Teoh, 2000), and it surely represents a serious limitation also for the intended usage of polymer composites for hard and soft tissue applications. In fact, in several potential clinical applications, materials are subjected to cyclic loads. For example, in the acetabulofemoral joint (hip joint) half to two million cycles per annum can be estimated, depending on the patient’s age and activity (Wallbridge and Dowson, 1982), while roughly one million cycles per annum can be estimated for the finger joint (Unsworth, 1991). This chapter presents an overview on the fatigue fracture problems in polymeric composite biomaterials, with a description of some recent testing methodologies, and an outline of the future development of fatigueresistant biocomposites.
19.2
Fundamentals of fatigue failure in polymer composites
Fatigue is a general phenomenon in most engineering materials, and refers to microstructural damage and failure caused by applied cyclic loads or strains (Ritchie, 1999). Depending on the service conditions, a given component can be exposed to cyclic loads periodically oscillating between a lower and an upper value (load-controlled fatigue) or it can be cyclically deformed within certain limits of displacement or strain (strain-controlled fatigue). From an experimental point of view, the cross-section variations of unnotched specimens are generally ignored during fatigue experiments and, at a first approximation, tests performed under load-control are generally considered under stress-control. The key parameters describing a typical stress cyclic loading are represented in Figure 19.1. The following variables define a given fatigue condition: maximum stress σmax minimum stress σmin
Fatigue behaviour of biocomposites
sB
467
Tensile strength
Stress
smax sA
smean
Δs
0 smin Time
19.1 Key parameters for stress-controlled fatigue loading.
σ min σ max 1 mean stress σ mean = (σ max + σ min ) 2 1 stress amplitude σ A = (σ max − σ min ) 2 stress range Δσ = σmax − σmin stress ratio R =
Analogous quantities can be defined for strain- (or displacement-) controlled fatigue experiments. The load history is completely defined when at least two of the above parameters are fixed and when the waveform (e.g. sinusoidal, triangular or square) and the frequency (in Hz) of the cyclic load are specified. Depending on the testing configuration, loads (or displacements) can be generated either by rotational bending, reciprocal bending, reciprocal torsion, or by pulsating strokes of an axial actuator. Fatigue experiments can be broadly divided into two main categories (Hertzberg and Manson, 1980; Moore and Turner, 2001): i) those aimed at obtaining a S–N (or Wöhler) curve reporting the applied stress as a function of the number of cycles to failure for un-notched specimens, or ii) more rigorous fracture mechanics approaches, in which a crack growth (or delamination) rate is evaluated as a function of the applied stress intensity factor (or strain energy release rate) on specimens containing an initial sharp notch (or a delamination starter, in case of composite laminates). One common way to represent the results of an uniaxial fatigue test on un-notched samples is by S–N diagram such as those represented in Fig.
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Biomedical composites
400 Carbon–epoxy laminates
sA (MPa)
300
200 Glass–epoxy laminates 100
0 –1
0
1
2
3
4
5
6
7
log Nf
19.2 S–N diagrams for (䉱) carbon/epoxy and (䊉) glass-epoxy [0/± 45/0] laminates tested at a frequency of 10 Hz and a stress ratio R = 0.1.
19.2, in which the alternating stress amplitude is reported as a function of the number of cycles to failure at a constant stress ratio R (Stefani, 2008). In some materials, the S–N curve flattens out eventually, so that below a certain endurance limit (σe) failure does not occur no matter how long the loads are cycled. For other materials, notably non-ferrous alloys, no endurance limit exists and, for design purposes, a fatigue strength (σN) is considered, which represents the stress level at which failure is likely to occur for a given number of cycles (Nf). The maximum test frequency is generally limited to approximately 10 Hz since inertia in components of the testing machine and heating of the specimen often become problematic at higher frequencies. At that speed it takes about 12 days to reach 10 million loading cycles. Consequently, obtaining a full S–N curve is a rather tedious and expensive procedure, and it will surely be impractical to determine whole families of curves for every combination of mean and alternating stress. A possibility to overcome this problem for design purposes is based on the use of Goodman plots such as that schematically depicted in Fig. 19.3. The plot is constructed with mean stress as the abscissa and stress amplitude as the ordinate, and a straight ‘lifeline’ is drawn from a value corresponding to the endurance limit σe on the ordinate axis to the tensile strength σB on the abscissa. For any given mean stress, the endurance limit can be read directly as the ordinate of the lifeline at that value of the mean stress. Alternatively, if the design application requires a given ratio of σA to σM (considering that
Fatigue behaviour of biocomposites
469
Stress amplitude sA
sE (or sN) Failure Integrity sB
0 Compression
Mean stress sM
Tension
19.3 The Goodman diagram.
σA 1− R = ), a line is drawn from the origin with a slope equal to that ratio. σM 1+ R Its intersection with the lifeline then gives the effective endurance limit for that fatigue loading condition. A great deal of experimental evidence indicate that the crack growth rate can be correlated with the cyclic variation in the stress intensity factor (Paris law equation) (Paris, 1962): da = A ΔK m dN
(19.1)
da is the fatigue crack growth rate per cycle, K = Kmax − Kmin is the dN stress intensity factor range during the cycle, and A and m are parameters depending on material, frequency, stress ratio, and environmental conditions (temperature, humidity) (Hertzberg and Manson, 1980). A typical Paris plot is reported in Fig. 19.4 for injection moulded polypropylene (PP) and PP-short glass fibre (SGF) composites with three different fibre contents (Pegoretti and Ricco, 2000). When subjected to fatigue conditions, composite materials generally suffer ‘fatigue damage’ consisting of a cycle-dependent degradation of internal integrity (Reifsnider, 1990). For composite laminates reinforced with continuous fibres, the mechanisms and entities of fatigue damage markedly depend upon the stacking sequence, the ply thickness, the materials types and the loading conditions (Stinchcomb and Bakis, 1990). For composite laminates containing both 0° and off-axis (θ) plies, the progressive development of damage during fatigue life can be schematically depicted as in Fig. 19.5, which traces the damage process at various percentages of life under uniaxial tensile loading.
where
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Biomedical composites
log (da/dN) (mm cycle–1)
–2
–3
–4 PP –5
PP + 10% SGF PP + 20% SGF
–6 PP + 30% SGF –7 0
0.2
0.4
0.6
log ΔK (MPa
0.8
1
m0.5)
19.4 Typical Paris plot representing the fatigue crack propagation in neat injection moulded PP, and relative composites reinforced with 10, 20 and 30 wt% of short glass fibres.
1 - Matrix cracking fibre breaking
Damage
0°
3 - Delamination fibre breaking
0°
0° 0°
0°
2 - Crack coupling, interfacial debonding, fibre breaking 0
5 - Fracture
0°
0° 0°
0°
4 - Delamination growth, fibre breaking (localized) 100
Life time (%)
19.5 Damage modes during fatigue lifetime of composite laminates (Stinchcomb and Bakis, 1991). Reprinted from Reifsnider K.L., 1991 with permission from Elsevier.
0°
Fatigue behaviour of biocomposites
471
The damage process can be considered as composed of the following three main stages. During the initial 10–15% of life, damage develops very rapidly (stage I) involving prevalently matrix cracking but also some fibre breaking. During the next 70–80% of the fatigue life (stage II) the damage evolves at a lower rate than during stage I, with turning and growth of matrix cracks along the interfaces between plies. Crack coupling produces interfacial debonding that can eventually grow in the plane of the laminates to form delaminations. In the last 10–15% of fatigue life, the damage increases at a higher rate with an extended delamination growth and final fracture. Fibre–matrix debonding may also occur at stresses as low as 15% of the ultimate bond strength of the interface (Latour and Black, 1993). Fatigue failure under multiaxial state of stress can be approached by criteria similar to those adopted for static failure, except that the material properties are not constants but depend on the number of cycles, stress state and stress ratios (Shokrieh and Lessard, 2003). Perhaps the most commonly observed failure mode in composites is delamination, which implies a separation between individual plies (O’Brien, 1990). Fatigue delamination testing of composite laminates can be conducted under mode I by using double-cantilever-beam, or under mode II on end-notched flexure (ENF), four-point end-notched flexure (4ENF), or end-loaded split (ELS) specimens (Martin, 2003). The above-mentioned specimens are used to characterize fatigue delamination by monitoring the da delamination growth per fatigue cycle ( ) as a function of the applied dN cyclic strain energy release rate (ΔG). A power law equation has been proposed to relate the delamination crack growth rate with the applied ΔG (Hojo et al., 1987): da = DΔG n dN
(19.2)
For composite materials, the values of the exponent n are typically in the range from 3 to values in excess of 15 (Martin, 2003). Information on the fatigue behaviour of short-fibre-reinforced composite materials are mostly limited to random chopped glass fibres strand reinforced thermosets and injection-moulded thermoplastics reinforced with short glass or carbon fibres (Mandell, 1990a). The fatigue damage in chopped-strand-reinforced thermosetting matrices is analogous to that of continuous fibre laminates (Caprino, 2003). Damage generally develops in the form of matrix and interface cracking, which usually intensifies with increasing cycles finally resulting in a reduction of residual mechanical properties and eventually to complete failure. Damage prevalently arises from matrix cracking and/or fibre–matrix debonding. In short-fibrereinforced injection-moulded thermoplastics, several fatigue damage
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Biomedical composites
mechanisms can be observed depending on the characteristics of the matrices and the composite microstructure. For ductile matrices: i) matrix yielding or ii) fibre–matrix debonding may occur. For brittle matrices microcracks or crazing can develop, most probably around transversely oriented fibres. Available results on the fatigue response of short-fibre composites consist of both traditional S–N curves and of crack propagation curves (Hertzberg and Manson, 1980). Analyzing a great deal of data, Mandell (1990b) verified that the S–N curves of composites with long or short fibres, with various matrices and orientation can be reasonably well fitted by the following equation:
σ max = σ B − B log N f
(19.3)
where the maximum stress in a fatigue cycle, σmax, is linearly related to the logarithm of the cycles to failure Nf, being σB the tensile strength and B the slope of the S–N curve. Because of the great differences in the manufacturing processes, a direct comparison of the fatigue response of otherwise identical composites reinforced with continuous and short fibres of the same type are relatively rare (Harris, 2003). However, Harris et al. (1990) were able to produce unidirectional XAS/914 carbon/epoxy composites with the same volume fraction (0.35) of long or short fibres. In particular, composites reinforced with continuous fibres were obtained from conventional continuous prepregs, while composites containing chopped 3-mm-long fibres were produced from a glycerol-alignment method. As reported in Fig. 19.6, although the strength of the short-fibre composite is much lower than that of long-fibre laminate, its fatigue
2.5 R = 0.1 Long fibres
smax (GPa)
2.0 1.5 1.0
Short fibres
0.5 0 –1
0
1
2
3 4 log Nf
5
6
7
8
19.6 S–N diagrams for unidirectional carbon/epoxy laminates reinforced with continuous and discontinuous fibres (Harris, 2003).
Fatigue behaviour of biocomposites
473
response is better, since the slope of the S–N curves, i.e. B in equation (19.3) is lower. A more convenient way to represent equation (19.3) is the following:
σ max = 1 − b log N f σB
(19.4)
The advantage of this approach is that materials with markedly different static strength values can be compared on the same normalized S–N curves whose slope is given by the normalized stress fatigue parameter b.
19.3
Fatigue behaviour of biocomposites for hard tissue applications
One of the main problems of metallic or ceramic implants is the mismatch of stiffness with respect to hard tissues (Park et al., 2007). Polymer composites may surely offer possibilities for a better match of their elastic properties with those of natural hard tissues. Moreover, if properly designed, polymer composites may display an anisotropic behaviour, typical of some hard tissues. For example, the mechanical properties of cortical bone markedly depends on the loading direction. The studies on the fatigue response of polymer biocomposites are mostly focused on the application of such materials in bone repair and replacement, dentistry, joint replacement and spine surgery.
19.3.1 Bone repair and replacement Fatigue loading of bone generally causes damage consisting of the formation of microcracks (O’Brien et al., 2007). This damage acts as a stimulus for bone remodelling (Schaffler and Jepsen, 2000), even if the rate of damage accumulation can in some circumstances exceed the self-repair ability of the bone (Burr et al., 1985) and fractures result. This may occur during highly intensive activities such as marching or running, or under ‘normal’ circumstances because the bone mechanisms is deficient as in osteoporotic bones. Currently available polymer composites do not have this inherent ability to repair damages, but intensive research efforts in the field of self-healing polymers and composites are in progress (Pegoretti, 2009; Yuan et al., 2008). Therefore, the ability to survive the fatigue loading experienced under service conditions is a crucial aspect in polymer composites for applications in bone repair and replacement. Bone fractures can be treated (anatomic reduction) by internal or external fixation devices. Polymer composites have been successfully investigated for both types of applications. In internal fixation, bone fragments are held in alignment by various implants such as screws, plates, wires, pins
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and intramedullary nails. Several hydroxyapatite (HA) reinforced polymers has been proposed for bone internal fixation or replacement both for non-resorbable and for biodegradable devices (Roeder et al., 2008). HA reinforced polymers were pioneered by Bonfield et al. (Bonfield et al., 1981), with the study of HA-reinforced high-density polyethylene (HDPE). HDPE reinforced with 40 vol% of HA particles was commercialized under the trade name HAPEXTM. The fatigue behaviour of HAPEXTM at 37 °C in saline was determined under uniaxial (fully reversed axial tension– compression and fully reversed torsion) and biaxial (various combinations of axial and torsional loads) loading conditions (Ton That et al., 2000a and 2000b). S–N curves under uniaxial conditions indicated fatigue limits at between 37 and 25% of the ultimate strength of the material. Superposition of torque on axial loading considerably reduces the fatigue life of HAPEXTM. Owing to molecular alignment, hydrostatically extruded HAPEXTM exhibited anisotropic mechanical properties and a significantly improved flexural and fatigue life compared with conventional, isotropic HAPEXTM (McGregor et al., 2000). According to a recent investigation, HDPE reinforced with HA whiskers exhibited a four- to five-fold increase in fatigue life compared with an equiaxed powder at either 20 and 40 vol% filler content (Kane et al., 2008). These results confirm the observations of Joseph and Tanner (2005) who found a marked dependency of the fatigue behaviour on the surface area and morphological features of filler particles on the fatigue behaviour of HA-filled HDPE composites. Hydroxyapatite has been also proposed as a possible filler for other non-resorbable matrices such as poly(ether ether ketone) (PEEK) (Abu Bakar et al., 2003a), polysulfone (PSU) (Wang and Chua, 2003) and ultra-high-molecular-weight polyethylene (UHMWPE) (Fang et al., 2006). HA particle reinforced PEEK composites with various filler contents up to 40% (by volume) were subjected to tension–tension fatigue under load-controlled mode (Abu Bakar et al., 2003b, Tang et al., 2004). The fact that all of the specimens survived cyclic load at 50% ultimate tensile strength suggests that HA/ PEEK composite is a promising fatigue-resistant material for biomedical applications. Wang investigated the fatigue performances of HA/PSU composite under different combinations of axial and torsional loads (Wang and Chua, 2003). It was found that the fatigue life of unfilled PSU and HA/PSU composite was reduced as the shear component of biaxial stress field increased. The addition of particulate HA in the polymer led to a shorter fatigue life of composite in low shear stress conditions. However, in high shear stress conditions, the effect of shear stress became dominant and the fatigue life of unfilled PSU and HA/PSU composite was similar. HA particles have been also used to reinforce resorbable matrices, such as poly(llactide acid) (PLLA) (Shikinami and Okuno, 1999, Shikinami and Okuno, 2001) or poly(lactide-co-glycolide acid) (PLGA) (Higashi et al., 1986) for
Fatigue behaviour of biocomposites
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bone repair and replacement applications. Very limited information is available on the fatigue response of these materials. The fatigue behaviour of miniscrews and miniplates made of forged composites composed of up to 40 wt% of raw hydroxyapatite particles in a PLLA matrix were investigated by Shikinami and Okuno and compared with those of similar titanium devices (Shikinami and Okuno, 2001). In spite of their approximately 2 or 3 times lower absolute strengths, a remarkable distinction that makes the composite miniplates better than the titanium ones was their fatigue resistance to alternate bendings. In fact, they retained 70% of their initial strength even after 60 cycles without revealing any damage, whereas the metallic devices fully broke off after only eight cycles. Carbon fibres have been also investigated as possible reinforcement to produce polymer composites for fixation devices (Fujihara et al., 2004), and a limited amount of information is available on the fatigue behaviour of these materials. Fujihara et al. (2007) investigated the fatigue bending properties of braided carbon/epoxy laminated composites with a target of surgically used compression bone plate to repair diapheseal fractures. Interestingly enough, they reached the conclusion that for an optimal braiding angle of 20° the fracture resistance under fatigue loading was maximized. Braided carbon fibres were also used to reinforce PEEK bone plates (Schambron et al., 2008). As documented in Fig. 19.7, preliminary fatigue experiments indicate that CF/PEEK bone plates are likely to perform well under physiological loads and when subject to a saline environment. In fact,
1.1
Normalised load [-]
1.0 0.9 0.8 0.7 CF/PEEK, non-aged CF/PEEK, aged in saline A316L, tested in air A316L, tested in saline
0.6 0.5 100
101
102
103
104
105
106
107
Cycles to failure
19.7 Normalized S–N curves for CF/PEEK and stainless steel bone plates. Reprinted with permission from Schambron et al., 2008, with permission from Elsevier.
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at 75% of the static strength, failure was not reached even after more than 300 000 load cycles, and the devices performed better than those made of the commonly used A316L stainless steel. Fatigue properties of screws made of biostable (PSU) or bioresorbable (PGLA) polymers reinforced with short carbon fibres have been investigated by Chlopek and Grzegorz (2005) by means of displacement controlled fatigue tests with R = 0. CF/PSU composites had better fatigue properties than CF/PLGA composites. Nevertheless, none of the investigated materials manifested an endurance limit. Polymer composites reinforced with carbon fibres have been also proposed for the fabrication of external fixators for bone fractures (Migliaresi et al., 2004), but no information is available on the fatigue behaviour of such devices. The use of polymer composites has also been proposed for mechanical analogue bone models for the evaluation of implants and surgical procedures (Chong et al., 2007a). A major problem with these models is that the current epoxy— short-fibreglass-based composite used as the cortical bone substitute is prone to crack formation and failure in fatigue or repeated quasistatic loading of the model. Fatigue crack propagation rate in short fibre reinforced epoxy composites for analogue cortical bone was recently investigated by Chong et al. from a fracture mechanics point of view (Chong et al., 2007b).
19.3.2 Dental applications Polymer composites are widely used in dentistry, mainly as restorative materials, dental posts, orthodontic archwires and brackets or bridges (Ramakrishna et al., 2001, Fujihara et al., 2004). Dental restorative materials are used to fill the tooth cavities, to mask discoloration or to correct contour and alignment deficiencies. Development of dental composite restorative materials started in the late 1950s when Bowen began to explore the usage of epoxy resins filled with inorganic particles (Drummond, 2008). Modern dental composites consist of polymerizable resin matrix, reinforcing glass or silica fillers, and silane coupling agents (Ferracane, 1995). Self-cure or light-activated polymerizable matrices typically contain one or more monomers, such as bis-phenolA-diglycidyl dimethacrylate (bis-GMA), triethylene glycol dimethacrylate (TEGDMA), and urethane dimethacrylate (UDMA). For most dental restorative materials, the main cause of failure is the breakdown of the resin matrix and/or the filler-matrix interface (Drummond, 2008). Fatigue failure manifests itself in dental prostheses and restorations as wear, fractured margins, delaminated coatings, and bulk fracture (Baran et al., 2001). Various testing strategies have been developed for the assessment of the fatigue behaviour of dental restorative materials. In most cases a flexural
Fatigue behaviour of biocomposites
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test is adopted on rectangular bars under three- or four-point bending configuration (Yap and Teoh, 2003). Flexural S–N curves and flexural fatigue limits have been experimentally measured for several commercially available dental restorative materials (Abe et al., 2005; Boberick et al., 2002; Braem et al., 1994 and 1995; McCool et al., 2001; Yoshida et al., 2004). Under environmentally controlled conditions, the fatigue behaviour is characterized by a well-defined endurance limit, that decreases after water sorption (Braem et al., 1994). A similar effect was also observed by McCool et al. (2001), who reported that aging in a water–ethanol solution reduced the survival stresses of dental restorative composites by a factor of four to five. In order to simulate the complex stress state commonly encountered under service conditions, various testing configurations have been developed to assess the fatigue response of dental restorative materials. Fujii et al. investigated the surface contact fatigue of two dental filling materials by a rolling-ball device (Fujii et al., 2004). The surface fatigue tests were carried out using loads from 1 to 5 N on a microfilled composite and a glass ionomer. The surface contact fatigue of the glass filled Bis–GMA– TEGDMA composite was at least 100 times greater than that of the glass ionomer. For both materials, the fatigue life was reduced significantly by increasing the test load. Contact versus flexure fatigue behaviours of dental composites have been also compared (Al-Turki et al., 2007; McCabe et al., 2000). Within the tested composite products, a hybrid composite material had a significantly greater flexural fatigue limit than a microfilled one, but the latter material had a significantly greater surface contact fatigue life, indicating that wear behaviour cannot be predicted from bulk fracture characteristics and vice versa. A rotation bending cantilever fatigue (RBCF) tests was adopted by Scherrer et al., in which specimen rotated around its longitudinal axis while the protruding end was loaded via ball-bearing (Scherrer et al., 2003). Interestingly enough, the authors claimed that correlations between monotonic flexure strength and resistance to fatigue loading were weak. Because fatigue tests are considered more pertinent than monotonic tests as to their predictive value, it is concluded that flexure strength data alone may not provide relevant information for long-term clinical performance. A laser-based micro-torsion apparatus was recently developed to measure the fatigue properties of a series of dental restorative composites (Papadogiannis et al., 2007). In particular, fatigue tests were performed on four different restorative composite resin, whose trade name and composition are reported in Table 19.1. Dynamic torsional loading was conducted at a resonance frequency of 30–50 Hz. All the investigated materials showed a loss of strength following repeated stress, owing to material fatigue. The material with the highest shear modulus had the lowest damping and the
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Table 19.1 Trade names, manufacturer and composition of restorative composites whose fatigue performances are summarized in Fig 19.8 (Papadogiannis et al., 2007) Trade name
Manufacturer
Composition
Alert
Jeneric/Pentron Inc., Wallingford, CT, USA
Admira
Voco Gmbh, Cuxhaven, Germany
Synergy
Coltene Whaledent Gmbh, Konstanz, Germany
Filtek P60
3 M Dental Products Inc., St Paul, MN, USA
Matrix: functional dimethacrylates of ethoxylated bisphenol-A polycarbonate resins (EBRADMA–PCDMA), crushed glass fibres (CS2). Fillers: barium borosilicate glasses, combined with silanated colloid silica (83.5 wt.%, 67 vol.%). Particle size: 0.8 μm. Photoinitiator, amine accelerator, UV absorber, inorganic pigments. Matrix: anorganic – organic copolymers (Ormocers), bis-GMA, diurethane dimethacrylates, BHT, TEGDMA. Filler: anorganic microfillers 56 vol.% (78 wt.%). Particle size: 0.7 μm. Matrix: BisGMA, BisEMA, TEGDMA resins. Filler: strontium glass, silanized barium glass, silanized amorphous silica, hydrophobed, 59 vol.% (74 wt.%). Particle size: 0.04–2.5 μm (range), Average 0.6 μm. Matrix: bis-GMA, UDMA, Bis-EMA resins. Filler: zirconia/silica (83 wt.%, 61 vol.%). Particle size range of 0.01–3.5 μm. Initiators, inorganic pigments.
highest fatigue strength. S–N curves reporting the maximum applied peak shear stress as a function of the number of cycles to failure are reported in Fig. 19.8. Limited information exists on the compressive fatigue behaviour of dental restorative materials (Mohandesi et al., 2007). Some authors investigated the effect of filler content on the fatigue resistance of dental restorative composites (Htang et al., 1995, Lohbauer et al., 2006). Htang et al. prepared a series of experimental composites incorporating a silanized quartz filler (3–5 micron in size) into a light cured Bis-GMA–TEGDMA matrix, with a filler content varying from 40 to 85 wt%. The specimens were stressed with an impact load (16 mJ) repetitively applied from 50 000 to 150 000 times. The cracks induced by cyclic loading were observed on the sectioned surfaces of the tested specimens. The fatigue impact resistance, calculated from the median fatigue limit, revealed an optimal filler content at 75 wt%, with a decrease of the fatigue impact resistance beyond this limit (Htang et al., 1995).
Fatigue behaviour of biocomposites
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18 16
tmax (MPa)
14 12 10 8 6 4 3
4
5
6
7
8
log Nf
19.8 Shear S–N curves for various commercial dental restorative materials whose composition is reported in Table 19.1: (䉱) Alert, (䊉) Admira, (䉬) Synergy, (䊏) Filtec P60 (Papadogiannis et al., 2007).
A fracture mechanics approach was also adopted to investigate the effect of cyclic loading and environmental ageing on the fracture toughness of dental resin composites with different filler composition: a fibre filler, a hybrid filler, and a microfiller (Ravindranath et al., 2007). Both unaged and aged specimens were subjected to cyclic loading at a frequency of 5 Hz with sinusoidal loads cycling for 1, 1000, and 100 000 cycles at a load level equal to 60% of the fracture load of non-cycled specimens. For all the investigated dental resin composites, results show that ageing for 4 months in a 50/50 alcohol–water mixture lowered the fracture toughness, which was further reduced by cyclic loading. Truong et al. (1990) investigated the fatigue crack propagation in notched specimens under cyclic loading for a series of experimental composite resins and six commercial composites. The fatigue resistance was measured as a function of filler content, water absorption, and post-curing temperature. Constants A and m of the Paris law, equation (19.1), have been obtained by linear regression analysis. The plot of fatigue resistance against volumetric filler content, showed a sigmoidal relationship with a rapid increase in the range of 40–50 vol% then levelling off at higher values. Fatigue crack propagation curves in aqueous environments were also evaluated by Takeshige et al. (2007) for three commercial resin composites and one unfilled resin. Contrary to common perceptions, a retardation of crack propagation in composites under aqueous environments was determined by a significant increase in the threshold
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value of the applied stress intensity factor range (ΔKth), representing the asymptotic value of ΔK at which da/dN approaches zero. Nano-structured dental resin composites are currently attracting an increasing amount of interest as, theoretically, they have better wear and fatigue resistance compared with microfilled composites and may favour the achievement of restoratives with better long-term performance. Nevertheless, in a recent study Truoss et al. (2006) found that nano-structured composites did perform either similarly or comparatively worse than a microfilled composite, in terms of wear and fatigue resistance. The main role of a dental post is to provide retention to the core of an endodontically treated tooth. The materials traditionally used for dental posts, such as stainless steel, titanium and ceramics, are about 10–17 times stiffer than dentin (12 GPa) . The modulus mismatch between dentin and post is one of the prime reasons to cause stress concentration at the root of the teeth (Fujihara et al., 2004). In the last twenty years, several composite posts have been proposed, mostly based on undirectional carbon/epoxy (Reynaud et al., 1994) or glass/epoxy (Pegoretti et al., 2002) composite materials. A dental post is essentially a composite rod with diameter ranging from 0.8 to 2.0 mm and a length of approximately 20 mm. The composition of composite posts can be tailored in order to reach a stiffness similar to dentine, and even functionally graded in order to possess a high stiffness in the coronal region gradually reducing to the value of dentine at the apical end (Fujihara et al., 2004). Another possible advantage of composite materials relies on their excellent fatigue properties as compared with metal or ceramic posts. While the fatigue properties of metal (Cohen et al., 1997; Huysmans and Vandervarst, 1993; Huysmans et al., 1993) and ceramic (Heydecke et al., 2002) posts are relatively well assessed, somewhat limited experimental data exist on the fatigue behaviour of composite dental posts (Drummond and Bapna, 2003, Grandini et al., 2005, Wiskott et al., 2007). Eight different commercial endodontic posts were evaluated by Drummond and Bapna under static and cyclic flexural loading (Drummond and Bapna, 2003). They reported that carbon/graphite fibre-reinforced resin posts and the glass FiberKor posts were significantly stronger than ceramic (zirconia) and the other glass-reinforced resin materials. Further, for glass fibre-reinforced test bars, no significative differences between testing in air and water was observed, but a marked lowering of strength (30–38%) was revealed when they were cyclic loaded. Grandini et al. also investigated the fatigue resistance and structural characteristics of eight different commercially available fibre posts (Grandini et al., 2005). Fatigue tests have been performed under three-point bending at 3 Hz. Despite the large variation in the response of different kinds of fibre-reinforced posts, the authors were not able to find a correlation between the results of fatigue testing and the material microstructure, and attributed the observed differences to random
Fatigue behaviour of biocomposites
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variations in the post-manufacturing process. The behaviour under rotational fatigue of several commercially available titanium-, ceramic-, glassfibre/epoxy- and glass-fibre/acrylic posts was recently compared in a rotating beam configuration (Wiskott et al., 2007). Interestingly enough, the fatigue resistance of the two fibrous posts with the highest fatigue resistance was twice that of any of the ceramic or metal posts. Dentistry applications of composite materials also encompass their usage in orthodontic archwires, brackets and bridges (Fujihara et al., 2004). The appearance of metal archwires and brackets gives an unnatural appearance to patients. Although the aesthetic appearance of clear polymer and ceramic brackets is satisfactory, they are easily fractured. These factors made researchers pay attention to polymer composite materials for the development of archwires and brackets. Composite archwires have been made of continuous poltruded glass-fibre-reinforced composites with matrices such as epoxies (Watari et al., 1998), polyethyleneterephthalate glycol (PETG) (Jancar et al., 1994), poly(methyl methacrylate) (Watari et al., 1998). Composite materials may find applications also as bridges, i.e. partial denture (false teeth) used to replace one or more teeth completely. The high cost and time-consuming preparation of current gold bridges has led to the development of relatively inexpensive composite bridges and dentures (Bjork et al., 1986, Miettinen and Vallittu, 1997). The transverse fatigue resistance of glass fibre reinforced acrylic resin for dentures has been evaluated by Vallittu et al. in a three-point bending configuration under displacement control (Vallittu et al., 1994). The results revealed that, compared with unreinforced specimens, continuous glass fibres at a concentration of 58 wt% enhanced the fatigue resistance during 5 × 105 loading cycles. The fatigue behaviour of a bis-GMA resin diluted with TEGDMA and reinforced with 65 wt% long E-glass fibres (Vectris-Pontic by Ivoclar), specifically developed for prosthetic dental bridges, have been tested both on dry and aged samples (Pegoretti and Migliaresi, 2002). The residual flexural strength of specimens subjected to 104 fatigue cycles at 1 Hz, under stress levels oscillating between a minimum value of 70 MPa and a maximum value varying from 895 to 1070 MPa is reported in Fig. 19.9. Ageing in water at 37 °C and under accelerated conditions (70 °C) induced a marked reduction of the materials fatigue life, as documented in Fig. 19.10. This latter aspect is actually one of the weak points affecting the usage of composite materials in clinical applications. In fact, it is well known that the presence of moisture, or direct immersion in water reduces the fatigue life of epoxy-based composite materials (Zaffaroni et al., 1998). In recent years, the usage of braided preforms have been proposed to produce aesthetic composite brackets (Fujihara et al., 2004). The advantage of braided preforms is that fibres are arranged in three directions, one of
Biomedical composites
Residual strength (MPa)
1400
1350
1300
1250
1200 0
200
600
400
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smax (MPa)
19.9 Residual flexural strength of unaged Vectris–Pontic dental composites after 104 fatigue cycles under various maximum stress values at 1 Hz frequency (Pegoretti and Migliaresi, 2002).
7 6 5 log Nf
482
4 3 2 1 0 0
5
10
15
20
25
30
35
Ageing time (weeks)
19.10 Fatigue life at 1 Hz frequency vs. the ageing time. Open symbols refer to samples still unbroken after 106 cycles, circles refer to samples aged at 37 °C and tested with αmax = 1001 MPa, squares refer to samples aged at 70 °C and tested with αmax = 572 MPa (Pegoretti and Migliaresi, 2002).
Fatigue behaviour of biocomposites
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which is the axial direction along the length of the reinforcing structure, and the other two, termed braiding yarns, are at predetermined sets of angles (such as ±30° and ±45°), with the yarns intertwined. The structure is such that no two yarns are twisted around each other and the axials are not crimped, enabling full transfer of their longitudinal reinforcing efficiency (Karbhari and Wang, 2007). In a recent study, the intrinsic nature of the triaxial braid resulted in very little decrease in flexural strength as a result of fatigue cycling at 75% of static strength (Karbhari and Wang, 2007).
19.3.3 Joint replacement Over the years, a number of artificial joints have been designed to replace or augment many joints in the body. Unlike those used to treat bone fractures, artificial joints are placed permanently in the body. The most commonly used artificial joints are total hip replacement (THR) and total knee replacement (TKR) (Ramakrishna et al., 2001). In England, if current trends continue, there will be almost 47 000 primary hip and 54 000 primary knee operations annually by 2010 (Dixon et al., 2004). Primary total hip arthroplasty (THA) in the USA is estimated to grow from 208 600 in 2005 to 572 100 in 2030 (Campbell et al., 2008) A typical THR consists of a cup type acetabular component, and a femoral stem whose head is designed to fit into the acetabular cup, thus allowing joint articulation. Conventional THRs use stainless steel, Co–Cr and titanium alloys for the femoral shaft and neck and Co–Cr alloys or ceramics such as alumina and zirconia for the head. In order to reduce the mismatch of stiffness of the femur bone and the prosthesis, the usage of composite materials has been considered, and carbon-reinforced composite stems have been introduced (Chang et al., 1990, Akay and Aslan, 1996). Limited information exists on the fatigue behaviour of such devices. In a preliminary study on composite hip prostheses made by resin transfer moulding, Reinhardt et al. (1999) performed some fatigue tests. The preforms of the prototype consisted of braided high-strength-carbon-fibre socks wrapped around a balsa wood insert, and a vinyl ester matrix was adopted. The set-up adopted for static and fatigue testing of the composite hip implant is schematically shown in Fig. 19.11. For a load of 4 kN, which is four times the body weigh of a 100-kg person, the prothesis deflected for 1 mm, and the ultimate load occurred at 7.5 kN with a deflection of 3.6 mm in the stem direction. Fatigue testing showed that the prosthesis survived two million load cycles under a maximum applied load of 5 kN. In a recent study, Campbell et al. (2008) analysed the mechanical behaviour of composite femoral stems made of carbon fibre reinforced nylon-12 (CF/PA12). The novel hip stem consisted of a hollow structure made of a 3-mm thick CF/PA12 reinforced composite in the form
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Load cell
Joint ball
Hip stem
Housing Cured resin Joint
Load tube
P
19.11 Setup for static and fatigue tests on composite prostheses made by resin tranfer moulding. Reprinted with permission from Reinhardt et al., 1999, copyright Sage Publications.
of braided sleeves of commingled fibre of PA12 and CF. The prototypes were consolidated at 250 °C under a pressure of 480 kPa for 5 min using an inflatable bladder compression moulding method. Compression– compression fatigue tests were performed on cylindrical samples machined from the prototypes at a frequency of 6 Hz under load control. The resulting S–N curve is reported in Fig. 19.12. Results indicated that the fatigue performance of CF/PA12 composite surpassed by far the required fatigue performance for total hip prosthesis (THP) stems. For acetabular components, ultra-high-molecular-weight-polyethylene (UHMWPE) substituted poly(tetrafluoroethylene) (PTFE), was introduced in the early 1960s, but it is now no longer used owing to its insufficient wear
Fatigue behaviour of biocomposites
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200
smax (MPa)
150
100
50 0
1
2
3
4
5
6
7
log Nf
19.12 S–N compressive fatigue curves of CF/PA12 cylindrical tubes for femoral stems applications (Campbell et al., 2008).
resistance. The fatigue resistance of UHMWPE has been widely investigated, mainly from a fracture mechanics point of view aimed at assessing the fatigue crack propagation (Baker et al., 2003, Gencur et al., 2006, Oral et al., 2006). At the early stages of material development, some attempts have been also made to improve the fatigue crack propagation resistance of UHMWPE by reinforcing with short carbon fibres (Connelly et al., 1984). The fatigue crack propagation resistance of the reinforced UHMWPE was found to be significantly worse than that of plain UHMWPE. Recently, some attempts have been made to improve the resistance to fatigue crack propagation of UHMWPE by radiation-induced crosslinking (Gencur et al., 2006; Oral et al., 2006). It turned out that crosslinking actually reduced the ability of UHMWPE to resist crack initiation and propagation under cyclic loading, probably owing to the concurrent decrease in the crystallinity content. Proper fixation to bones is as important as the design of joint replacement and the usage of proper materials for stem and acetabular cup. The most common approach is to press-fit the joint prosthesis into the bone using a grouting material called bone cement (Ramakrishna et al., 2001). The most widely used bone cement is based on poly(methyl methacrylate) (PMMA), also called acrylic bone cement (Kenny and Buggy, 2003), although alternative bone cement formulations for cemented arthoplasties have been proposed (Lewis, 2008). Over the past three decades or so, there have been a large number of reports on the impact of an assortment of variables on the fatigue lifetimes of a large number of acrylic bone-cement formulations, and the main results has been reviewed by Lewis (2003). From the
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Biomedical composites
examination of the reports on the subject, unanimity of support emerged for the notion that an increase in the molecular weight of the powder constituents of the fully cured cement leads to an increase in the cement’s fatigue life. On the other hand, disagreement emerged as to whether vacuum mixing the cement constituents leads to an increase in the fatigue life of the fully cured cement (relative to the hand-mixed counterpart). Some efforts have been also conducted for improving the fatigue resistance of PMMA-based bone cements by adding fibres, particles or nanofillers. The effects of centrifugation and titanium fibre reinforcement on fatigue failure mechanisms in PMMA bone cement were investigated by Topoleski et al. (1995). Ti fibre addition increased the notched fatigue life at high initial stress intensities, and the reinforcing effect increased with decreased stress intensity. In both un-notched and notched fatigue experiments, the combination of porosity reduction and Ti fibre reinforcement (C-R-PMMA) led to substantial improvements in fatigue life at each stress and stress intensity factor values, suggesting that the effects of Ti fibre reinforcement and porosity reduction are additive. Fibre reinforcement affects the fatigue crack propagation phase of failure in bone cement, enhancing the fatigue crack propagation resistance. A self-reinforced PMMA (SRC–PMMA) has been developed by Gilbert et al., consisting of high-strength, high-ductility PMMA fibres embedded in a PMMA matrix (Gilbert et al., 1995). These authors performed fatigue testing under a three-point bending configuration at 5 Hz and R = 0.1 on SRC–PMMA samples and on pure PMMA and on a commercial acrylic bone cement. The fatigue strength at 106 cycles of SRC–PMMA samples was significantly greater (80 MPa) compared with bone cement and PMMA, both of which had fatigue strengths of about 18 MPa. Among the other types of fibres that have been successfully adopted to improve the fatigue resistance of PMMA bone cements, one can mention surface-treated UHMWPE fibres (Andreopoulos et al., 1991) and milled carbon fibres (Kim and Yasuda, 1999). Hydroxyapatite can also be added to PMMA bone cement matrix in order to modify its resistance to cyclic loading conditions (Harper et al., 1995; Matsuda et al., 2004). Harper et al. (1995) observed no beneficial effects on the fatigue strength of an acrylic bone cement owing to the presence of HA particles, whereas Matsuda et al. (2004) reported a marked reduction of the fatigue crack propagation rate owing to the introduction of HA fibres. As shown in Fig. 19.13, the fatigue crack propagation curves shifted to higher applied ΔK values as the HA content increased. Marrs et al. (2006, 2007) recently explored the possibility of improving the fatigue resistance of a methyl methacrylate–styrene copolymer (MMAco-Sty) acrylic bone cement through the addition of multiwalled carbon nanotubes (MWCNTs). Various contents of as-produced MWCNTs were dispersed throughout the molten matrix of pre-polymerized commercial
Crack growth rate, da/dn (mm cycle–1)
Fatigue behaviour of biocomposites
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10–2 10–3 10–4 10–5 10–6
1
2 3 4 5 Stress intensity factor range, ΔK (kgf mm1.5)
6
19.13 Fatigue crack propagation curves for commercial acrylic cement unfilled (䊐) or reinforced with (䉱) 6 wt%, (䊊) 10 wt%, (䊏) 15 wt%, (䊉) 20 wt% and (䉭) 25 wt% of HA fibres. Reprinted with permission from Matsuda et al., 2004, copyright Elsevier.
bone cement powder in the heated (220 °C) chamber of a high-shear mixer. Bar specimens of the composite materials were fatigue tested to failure in 4-point bending at a test frequency of 5 Hz. The effect of MWCNT content on the fatigue cycles to failure is summarized in Fig. 19.14 that also contains the quasi-static flexural strength data. Unlike the small but significant improvement of quasi-static flexural strength, a 2 wt% loading of MWCNTs was associated with a substantial enhancement of the bone cement’s fatigue performance. In fact, specimens with a 2 wt% content of MWCNTs showed a 3.1-fold increase in the mean actual fatigue life. Although the nanotube–matrix interactions are not yet fully understood, the authors speculated that the observed reduction in material property enhancement with increasing MWNT loading could be related to the beneficial effects of MWCNT material augmentation competing against the adverse effects of sporadic inadequately dispersed (still agglomerated) ‘clumps’ of MWCNTs. These MWCNT clumps, whose number and extension is naturally increasing with the MWCNTs content, may act as fracture initiation sites. The same workers provided experimental data also on the fully reversed (R = −1) tension–compression fatigue resistance of these materials under three different peak stress amplitudes (20, 30 and 35 MPa) in a 37 °C saline environment (Marrs et al., 2007). By using a three-parameters Weibull model, the cycles to failure (Nf) for each sample were converted into a probability of failure P(Nf), as follows: β
⎡ N − N0 ⎞ ⎤ P ( N f ) = 1 − exp ⎢ − ⎛⎜ f ⎟ ⎥ ⎣ ⎝ α − N0 ⎠ ⎦
(19.7)
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8
80
6
4
Cycles to failure (×104)
Flexural stength (MPa)
90
70 2 0
2
4
6
8
10
MWCNT content (wt%)
19.14 Effect of MWCNT content on the flexural (䊊) quasi-static strength and (䊉) cycles to failure of an acrylic bone cement (Marrs et al., 2006).
where N0 is the minimum fatigue life (the baseline number of cycles for the sample at a given load amplitude), α is the location parameter (the number of cycles to failure below which 63.2% of the specimens fail), and β is the Weibull shape parameter (an indicator of the variance within the sample). In order to compare the fatigue resistances of the investigated materials, the authors proposed to adopt the Weibull mean (NWM) parameter calculated as follows: 1 N WM = N 0 + (α − N 0 ) Γ ⎛⎜ 1 + ⎞⎟ β⎠ ⎝
(19.8)
The higher NWM, the better the material fatigue resistance. Table 19.2 summarizes how NWM values depend on the MWCNTs content, for the investigated nanocomposites. The data reported in Table 19.2 clearly indicate that the fatigue resistance of MMA-co-Sty-MWCNF nanocomposites stressed under physiologically relevant conditions depend on the concentration of MWCNTs, the dispersion of the MWCNTs, and the peak stress of the dynamic loading cycle. In particular, testing at peak stress amplitudes of 20 and 30 MPa confirmed the existence of an optimal MWCNTs content (of about 2 wt%). The effect of adding an elastomeric second phase, acrylonitrile– butadiene–styrene, on the fatigue crack propagation behaviour of PMMA bone cement was investigated by Vila et al. (1999), who observed a decrease in the crack propagation rate between 1 and 2 orders of magnitude when comparing the plain with the modified cement.
Fatigue behaviour of biocomposites
489
Table 19.2 Effect of MWCNF content on the Weibull mean parameter (NWM) representing the fatigue resistance of MMA-co-Sty–MWCNF nanocomposites tested at various amplitudes of the peak stress (Marrs et al., 2007) MWCNT content (wt%)
NWM at σmax = 20 MPa
NWM at σmax = 30 MPa
NWM at σmax = 35 MPa
0.0 0.5 1.0 2.0 5.0 10.0
124 754 507 953 714 810 830 028 863 216 563 449
20 097 29 943 49 782 69 784 54 417 26 668
16 055 15 209 10 959 13 467 3 508 –
19.3.4 Spine surgery The human spine is a linked structure consisting of 33 vertebrae separated by fibrocartilaginous intervertebral discs and united by articular capsules and ligaments (Huang and Ramakrishna, 2004). Spine disorders may have several different origins and causes, such as births deformities, ageing, tumorous lesions, and mechanical loads induced by work and sport activities. When spine defects are limited to some vertebrae, treatments such as spinal fusion and disc replacement are used. The surgeons who helped develop the posterior lumbar interbody fusion cage (Arthur Steffee, MD, and John Brantigan, MD) initially envisioned a titanium device, owing to the mechanical loading requirements for these permanent implants. There were two perceived drawbacks with the initial proposed design, the first being the stiffness of the titanium device, which might promote stress shielding and inhibit bone growth, and the second being the radiopacity of the device, which would hinder diagnostic assessment of the bone growth (Kurtz and Devine, 2007). Several polymer composites based on highperformance thermoplastics have been considered for the Brantigan cage, the most successful ones being based on poly(aryl ether ketones) (PAEKs) reinforced with both chopped and continuous carbon fibres (Brantigan et al., 1991). Carbon-reinforced PEEK and poly(ether ketone ether ketone ketone) (PEKEKK) were successfully evaluated to aid the interbody lumbar fusion on 26 patients in a two-year pilot clinical study launched in 1989 (Brantigan and Steeffe, 1993). PEEK-based composites have almost two decades of successful clinical history in load-sharing fusion applications in the spine (Kurtz and Devine, 2007). In 1998, Victrex launched PEEKOPTIMA, which was specifically designed for long-term implantable applications. Both neat and carbon-fibre-reinforced PEEK-OPTIMA compounds have undergone extensive biocompatibility testing to meet the criteria of
490
Biomedical composites 300 250
smax (MPa)
200 150 100 50 0 –1
0
1
2
3
4
5
6
7
log Nf
19.15 S–N curves of long-term implantable CF/PEEK (PEEK OPTIMA LT1CA30). Open points refer to specimens that survived 106 cycles (run-out).
Food and Drug Administration (FDA) Master Files. In 2001, Victrex established Invibio Biomaterials Solutions that are currently providing several implantable PEEK grades and related composites for European Community (CE) and FDA approved devices. The outstanding fatigue resistance of CF/PEEK composites surely contributed to their widespread adoption for spinal applications. In Fig. 19.15, S–N curves of PEEK reinforced with 30 wt% of chopped carbon fibre (OPTIMA LT1CA30) are reported. Tests have been performed on injection-moulded specimens under tension– tension fatigue conditions at 5 Hz and R = 0. A linear fitting of the S–N curve with equation (19.4) furnished a stress fatigue parameter b equal to 0.065, which indicated that the fatigue resistance of this material is similar or even better that those of structural carbon/epoxy (b = 0.04) or glass/ epoxy (b = 0.10–0.14) composites (Mandell et al., 2002). Moreover, it is worthwhile to report that fatigue-proved (run-out) specimens (Fig. 19.15) preserved 97 and 88% of the initial tensile strength after 106 fatigue cycles under a maximum stress of 48 and 61% of the tensile strength, respectively. Osteoconductive and totally bioresorbable spinal/cervical interbody fusion cages were also fabricated from a composite of raw particulate hydroxyapatite/poly (l-lactide) (HA/PLLA) with an HA content of 40 wt% (Shikinami and Okuno, 2003). The mechanical strengths of three types of cages, designed for open-box, screw and cylinder constructs, were compared with those of existing metal and carbon-fibre/ polymer cages. Com-
Fatigue behaviour of biocomposites
491
pressive strengths of these composite cages surpassed those of existing metal and carbon-fibre cages. Fatigue resistance to alternate and static compressive loading persisted for longer than the minimum period (6 months) necessary for spinal devices in simulated body fluid at 37 °C.
19.4
Fatigue behaviour of biocomposites for soft tissue applications
Several types of implants are currently used in surgery to correct soft tissue deformities or defects, either congenital or acquired as a consequence of trauma or tumors (Ramakrishna et al., 2001). Depending on the intended application, the soft tissue implants may be required to solve various functions, such as filling the space from some defects; enclose, store or isolate something in the body; support mechanical loads; serve as scaffolds for tissue engineering. For functional load-carrying and supporting implants, the fatigue performances of composite materials may constitute an important parameter determining the success or failure in intended biomedical devices. This is typically the case of applications such as tendons, ligaments and vascular grafts.
19.4.1 Tendons and ligaments Tendons and ligaments are probably the best examples of load-bearing soft tissue implants. Tendons are strongly fibrous bands of tissue extending from a muscle to the periosteum of a bone, while ligaments are connective tissue bands linking bones in the proximity of every synovial joint. They are substantially composite materials consisting of undulated bundles of collagen fibres immersed in a complex matrix made of elastine and mucopolysaccharide hydrogel. A unique mechanical feature of tendons and ligaments is their non-linear S-shaped convex stress–strain curve. Seeldhom reported that estimated loads on the anterior cruciate ligament (ACL) of the knee joint are 196 N for level walking, 72 N for ascending stairs, 93 N for descending stairs, 67 N for ascending a ramp and up to 445 N for descending a ramp (Seedhom, 1993). Despite the fact that the mechanical testing of soft tissue is fraught with practical difficulties involving gripping the specimen, the measurement of cross-sectional area and the degree of hydration during both specimen preparation and testing, some attempts have been made to evaluate the in vitro fatigue resistance of tendons and ligaments (Schechtman and Bader, 1997, Schechtman and Bader, 2002, Devkota and Weinhold, 2003, Adeeb et al., 2004, Saweeres et al., 2005). The remarkable resistance to alternating loads of synthetic fibres make their use for tendon/ligament repair/replacement particularly attractive. Synthetic biomaterials used so far include fibres, in multifilaments or braided
492
Biomedical composites
forms, such as carbon (Jenkins, 1977, Forster et al., 1980, Mendes et al., 1986, Blazewicz et al., 1996), poly(ethylene terephthalate) (PET) (Ambrosio et al., 1998; Iannace et al., 1995; Karamuk et al., 2004; Kolarik et al., 1980), polyurethan (PU) (Liljensten et al., 2002; Peterson et al., 1985), UHMWPE (Kazanci et al., 2001, 2002b; Marois et al., 1993; Shalom et al., 1997; Ratner et al., 2005), aramids (Durselen et al., 1996), PP (Kdolsky et al., 1997). Moreover resorbable fibres such as polydiaxonone (PDS), PLLA and PGLA (Durselen et al., 2006) have been also considered. A silk-fibre matrix was also studied as a suitable material for tissue engineering ACL (Altman et al., 2002). The long-term clinical performances of a braided PP augmentation device (Kennedy model) in reconstructive ACL surgery was investigated by Kdolsky et al. (1997). One hundred and one partial or total breakages of the ligament augmentation device (LAD) were found. Scanning electron microscopic analysis of 24 retrieved failed LAD devices showed fatigue to be the principal failure mode, together with local abrasion at the fracture. In a more recent study, Guidoin et al. (2000) performed a retrospective analysis of 117 surgically excised anterior cruciate ligament in order to elucidate the mechanisms of failure of synthetic ligamentous prostheses. They were harvested from young and active patients (26 ± 7 yrs) at various orthopaedic centres in France between 1983 and 1993. The average duration of implantation of augmentation and replacement prostheses were 21.5 ± 12.6 and 33.2 ± 25.3 months, respectively. Fourteen types of ACL prostheses were analysed, each fabricated using a different combination of polymers, fibres and textile constructions. Three main mechanisms were involved in the failure of ACL prostheses: (1) inadequate fibre abrasion resistance against osseous surfaces; (2) flexural and rotational fatigue of the fibres, and (3) loss of integrity of the textile structure owing to unpredictable tissue infiltration during healing. Despite its crucial importance, the fatigue behaviour of fibre-based prostheses for ligament and tendons replacement or augmentation is a relatively little explored subject. Blasewicz et al. (1996) investigated the tensile fatigue response of carbon braids in air and immersed in isotonic solution at 37 °C. The effect of immersing carbon braids in an isotonic solution simulating body conditions was an increase in number of cycles to failure as compared with air conditions. A probable reason for this observation may be a decrease of abrasion between the filaments owing to the presence of isotonic solution. Moreover, The fatigue resistance of carbon braid increases with a decrease in the applied stress ratio, and a considerable improvement in lifetime is achieved below the maximum force of 500 N when energy dissipation per one cycle tends to zero. In order to match the complex and demanding mechanical requirements of a native human ACL, including adequate fatigue performance, a silk-fibre ACL prosthetic device was successfully designed by Kaplan and coworkers (Altman et al., 2002). As schematically depicted in Fig. 19.16a,
Fatigue behaviour of biocomposites 1 ACL matrix
6 Cords
6 Bundles
3 Strands
2 twists cm–1 clockwise
2 twists cm–1 counter-clockwise
493
30 Fibres
Parallel 30 (0) (a)
×
6 (2)
×
3 (2)
×
Parallel 6 (0)
×
1
3000
Load (N)
2500 2000 1500 1000 500 0 1 (b)
10
100
1000
10000 100000 1000000
Cycles
19.16 (a) Schematic of an 6-cord silk-based ACL prosthetic device; and (b) cycles to failure of the device. Reprinted with permission from Altman et al., 2002, copyright Elsevier.
single ACL matrix fibroin cords of geometry 30 × 6(3) × 3(3) × 6 were tied together with 2-0 Ethicon polyester sutures. Cycles to failure at maximum applied loads of 1680 and 1200 N (five specimens for each load) were determined in neutral phosphate buffered saline (PBS) at 23 °C using a sine wave function at 1 Hz. The obtained results are summarized in Fig. 19.16b, where the fatigue data trendline is extrapolated to a physiologically relevant cyclic load of 400 N (Chen and Black, 1980), indicating a fatigue life of 3.3 million cycles. Flat filament-wound polyethylene composites have been also evaluated as possible candidates for tendons and ligaments prostheses (Kazanci et al., 2001, 2002a, 2002b; Ratner et al., 2005). Filament wound flat strip composites of polyethylene fibre reinforced ethylene–butene copolymers were produced and their fatigue behaviour under cyclic loading was studied by Kazanci et al. (2002b). Composite materials were produced from Spectra 1000 UHMWPE fibres (Allied Signal) embedded in an ethylene–butene copolymer matrix of the Exact family (ExxonMobil). Fatigue tests were performed at room temperature under tension–tension sinusoidal load at
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Biomedical composites
Table 19.3 Effect of branching of PE–butene matrix and winding angle of UHMWPE fibres on the fatigue limit and fatigue degradation rate of flat strips (Kazanci et al., 2002b) Branching of PE–butene copolymer matrices (per 1000 main chain C atoms)
Winding angle (°)
Fatigue limit (MPa)
Fatigue degradation rate (MPa/logN)
50 42 50 66
28 42 42 42
46.4 28.3 28.4 24.6
5.7 4.1 2.9 2.8
R = 0.1 and a frequency of 1 Hz. The main results obtained from fatigue experiments are summarized in Table 19.3. In particular, three different copolymer compositions and two different winding angles were employed in order to study the effects of branching density in the polymeric matrix and of reinforcement angle on the fatigue response of the composite. The results were in agreement with published fatigue models, showing that the short-term fatigue behaviour, at relatively high stress levels, was controlled by the static properties of the materials, exhibiting better fatigue resistance for lower branching density of the copolymer and for a smaller reinforcement angle. However, the long-term fatigue behaviour, at moderate stress levels, was governed by the fatigue rate of degradation, which decreased with the branching density and winding angle. The fatigue-induced creep resulted in fibre reorientation in the loading direction, which, in turn, generated high residual properties. It was concluded that various polymer/angle combinations could result in fatigueproof composites of significant residual properties at 106 fatigue cycles.
19.4.2 Vascular grafts Arterial blood vessels are multilayered structures comprising collagen and elastin fibres, smooth muscle, ground substance and endothelium (Ramakrishna et al., 2001). As with almost all biological soft tissues, blood vessels are known to have a highly non-linear stress–strain relationship, strongly anisotropic mechanical properties, and significant viscoelastic features (Zhang et al., 2007). Segments of the natural cardiovascular system can be replaced by vascular grafts, with an increased success probability for blood vessels with lumen diameter of over 5 mm. Man-made prosthetic tubes (stents) can also be inserted into a natural passage/conduit in the body to prevent, or counteract, a disease-induced, localized flow constriction (Thurnher and Grabenwoger, 2002). Most widely used vascular grafts are woven or knitted fabric tubes of PET fibres or extruded porous wall tubes
Fatigue behaviour of biocomposites
495
of poly(tetrafluoroethylene) (PTFE) or PU (Ramakrishna et al., 2001). Dacron© (PET) prostheses are widely used as a durable and valuable vascular conduit since their introduction in 1957 (Szilagyi et al., 1958). In a review on the intrinsic structural failure of polyester (Dacron) vascular grafts, Van Damme et al. indicated that the most probable causative factor is material fatigue, leading to gradual breakdown and fragmentation of individual fibres, and subsequent biodegradation of the basic material (Van Damme et al., 2005). Moreover, in vitro and in vivo calcification of polymers such as PTFE, PU and silicone has been found to seriously reduce their flexibility, thereby causing their mechanical failure and degradation (Park et al., 2001). Early testing was conducted in the absence of standards and guidance specific to endovascular grafts, and references available for vascular grafts and stents did not adequately account for the complexity of endovascular graft systems. Failure of early-generation devices suggested that the testing being conducted was inadequate and that there was a lack of understanding of the in vivo environment (Abel et al., 2006). These concerns led to several efforts to improve preclinical testing. The Food and Drug Administration (FDA) sponsored a workshop to discuss the limitations inherent in testing of endovascular grafts, and an ISO standard for endovascular grafts was developed (ISO, 2003). Publication of the standard in 2003 succeeded in standardizing testing and reporting across device manufacturers; however, several clinical failure modes, such as migration and fractures, continued to be unpredicted by current preclinical testing. Despite fatigue failure being recognized as one of the main causes of mechanical failure of vascular prostheses, the investigation of the in vitro fatigue response of vascular grafts has been scarcely investigated (Botzko et al., 1979, Biedermann and Flora, 1982, Snyder et al., 1986). Thin walled tubes can be cyclically pressurized, while porous tubes can be lined with thin-walled latex balloons, and it has been demonstrated for PET knits that such tests yield results similar to animal implants in terms of dilatation and loss in strength (Botzko et al., 1979). Alternatively, cyclic mechanical testing can be performed by a circumferential tensile test, in which two metal halfcylinders are introduced into the prosthetic tube and moved apart (Dieval et al., 2008). Obviously, the environment plays an important role in determining the fatigue life of a vascular prosthesis. In fact, hydrolysis may lower the molecular weight and remove some plasticizers, thus decreasing the fatigue life.
19.5
References
abe, y., braem, m. j. a., lambrechts, p., inoue, s., takeuchi, m. & van meerbeek, b. (2005) Fatigue behavior of packable composites. Biomaterials, 26, 3405–3409. abel, d. b., dehdashtian, m. m., rodger, s. t., smith, a. c., smith, l. j. & waninger, m. s. (2006) Evolution and future of preclinical testing for endovascular grafts. Journal of Endovascular Therapy, 13, 649–659.
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abu bakar, m. s., cheang, p. & khor, k. a. (2003a) Tensile properties and microstructural analysis of spheroidized hydroxyapatite-poly(ether ether ketone) biocomposites. Materials Science and Engineering A-Structural Materials Properties Microstructure and Processing, 345, 55–63. abu bakar, m. s., cheng, m. h. w., tang, s. m., yu, s. c., liao, k., tan, c. t., khor, k. a. & cheang, p. (2003b) Tensile properties, tension–tension fatigue and biological response of polyetheretherketone–hydroxyapatite composites for load-bearing orthopedic implants. Biomaterials, 24, 2245–2250. adeeb, s. m., zec, m. l., thornton, g. m., frank, c. b. & shrive, n. g. (2004) A novel application of the principles of linear elastic fracture mechanics (LEFM) to the fatigue behavior of tendon tissue. Journal of Biomechanical Engineering– Transactions of the ASME, 126, 641–650. akay, m. & aslan, n. (1996) Numerical and experimental stress analysis of a polymeric composite hip joint prosthesis. Journal of Biomedical Materials Research, 31, 167–182. al-turki, l. i., drummond, j. l., agojci, m., gosz, m., tyrus, j. m. & lin, l. h. (2007) Contact versus flexure fatigue of a fiber-filled composite. Dental Materials, 23, 648–653. altman, g. h., horan, r. l., lu, h. h., moreau, j., martin, i., richmond, j. c. & kaplan, d. l. (2002) Silk matrix for tissue engineered anterior cruciate ligaments. Biomaterials, 23, 4131–4141. ambrosio, l., de santis, r., iannace, s., netti, p. a. & nicolais, l. (1998) Viscoelastic behavior of composite ligament prostheses. Journal of Biomedical Materials Research, 42, 6–12. andreopoulos, a. g., papaspyrides, c. d. & tsilibounidis, s. (1991) Surface treated polyethylene fibers as reinforcement for acrylic resins. Biomaterials, 12, 83–87. baidya, k. p., ramakrishna, s., rahman, m. & ritchie, a. (2001) Quantitative radiographic analysis of fiber reinforced polymer composites. Journal of Biomaterials Applications, 15, 279–289. baker, d. a., bellare, a. & pruitt, l. (2003) The effects of degree of crosslinking on the fatigue crack initiation and propagation resistance of orthopedic-grade polyethylene. Journal of Biom
