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] xiii
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
xv
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] xvi
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] xviii
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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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
Natural composites
5
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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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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15
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) 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 (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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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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(a)
(b)
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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(b)
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(d)
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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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
1 0.8 0.6 0.4 0.2 0
0
0.5
1.2
1
1.5 2 Time (s)
2.5
5 × 106 cycles
1 0.8 0.6 0.4 0.2 0
3
0
0.5
1.2
7 × 106 cycles
1
Displacement (mm)
Displacement (mm)
1.2
106 cycles Displacement (mm)
Displacement (mm)
1.2
0.8 0.6 0.4 0.2 0
289
1
2 1.5 Time (s)
2.5
3
2.5
3
107 cycles
1 0.8 0.6 0.4 0.2 0
0
0.5
1
1.5 Time (s)
2
2.5
3
0
0.5
1
1.5
2
Time (s)
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
References
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de santis r, mollica f, ronca d, ambrosio l, nicolais l (2003), ‘Dynamic mechanical behavior of PMMA based bone cements in wet environment’, J Mater Sci: Mater Med, 14, 583–594, doi: 10.1023/A:1024014822277. 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. de santis r, ambrosio l, mollica f, netti p, nicolais l (2007), ‘Mechanical properties of human mineralised connective tissues’, in F Mollica, L Preziosi, K R Rajagopal, Modeling of biological materials, Boston, Birkhauser, 211–261. deligianni dd, missirlis yf, tanner ke, bonfield w (1991), ‘Mechanical behaviour of trabecular bone of the human femoral head in females’, J Mater Sci: Mater Med, 2, 168–175. deteresa s j, nicolais l (1988), ‘The contribution of thermal stresses to the failure of Kevlar fabric composites’, Polym Compos, 9, 192–197. dunne nj and orr jf (2001), ‘Influence of mixing techniques on the physical properties of acrylic bone cement’, Biomaterials, 22, 1819–1826, doi: 10.1016/ S0142-9612(00)00363-X. ebied a and journeaux s (2002), ‘Metal-on-metal hip resurfacing’, Curr Orthop, 16, 420–425, doi: 10.2106/JBJS.F.01563. evans sl, gregson pj (1998), ‘Composite technology in load-bearing orthopaedic implants’, Biomaterials, 19, 1329–1342. fam h, bryant jt, kontopoulou m (2007), ‘Rheological properties of synovial fluids’, Biorheology, 2, 59–74, ISSN: 0006-355X. fantner ge, hassenkam t, kindt jh, weaver jc, birkedal h, pechenik l, cutroni ja, cidade gag, stucky gd, morse de, hansma pk (2005), ‘Sacrificial bonds and hidden length dissipate energy as mineralized fibrils separate during bone fracture’, Nature, 4, 612–616, doi: 10.1038/nmat1428. fung yc (1993), ‘Viscoelastic properties of articular cartilage’, in Biomechanics: mechanical properties of living tissues, Second Edition, Springer–Verlag, New York, 520–524. galante jo (1998), ‘Overview of total hip arthroplasty’, in Callaghan JJ, Rosenberg AG, Rubash HE, The adult hip, Vol 2, Philadelphia, PA: Lippincott-Raven, 829–837. gilbert jl, ney ds and lautenschlager ep (1995), ‘Self-reinforced composite poly(methyl methacrylate): static and fatigue properties’, Biomaterials, 16, 1043– 1055, doi: 10.1016/0142-9612(95)98900-Y. hamilton jv, caldarise s, treacy dj (1994), ‘Composite prostheses’, patent US 5,314,492. hof al (2001), ‘The force resulting from the action of mono and biarticular muscles in a limb’, J Biomech, 34, 1085–1089, doi: 10.1016/S0021-9290(01)00056-2. howling gi, ingham e, sakoda h, stewart td, fisher j, antonarulrajah a, appleyard s, rand b (2004), ‘Carbon–carbon composite bearing materials in hip arthroplasty: analysis of wear and biological response to wear debris’, J Mater Sci: Mater Med, 15, 91–98, doi: 10.1023/B:JMSM.0000010102.26218.d1. huiskes r, janssen jd, sloff tj (1981), ‘A detailed comparison of experimental and theoretical stress analysis of human femur’. In Cowin, SC (Ed.), Mechanical properties of bone. ASME, New York, pp. 221–234. inman vt (1947), ‘Functional aspects of the abductor muscles of the hip’, J Bone Joint Surg Am, 29, 607–619.
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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,
1
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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.
13.10 References NIH Roadmap for Medical Research, October 29, 2008, Homepage of OPASI, NIH, US Dept. of Health and Human Services, Available at: http://nihroadmap.nih.gov/. ‘The Human Genome’, 2001, Nature, 409, (6822).
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Castle Island’s Worldwide Guide to Rapid Processing. A Brief Tutorial, What is Rapid Prototyping? July 2, 2008, Available at: http://home.att.net/~castleisland/rp_int. htm. agarwal, b. d. and broutman, l. j. 1980, Analysis and performance of fiber composites, John Wiley & Sons, New York. anderson, e. j. and knothe tate, m. l. 2007, ‘Design of tissue engineering scaffolds as delivery devices for mechanical and mechanically modulated signals’, Tissue Eng., 13, 10, 2525–2538, doi:10.1089/ten.2006.0443. anderson, e. j., kaliyamoorthy, s., iwan, j., alexander, d. and knothe tate, m. l. 2005A, ‘Nano-microscale models of periosteocytic flow show differences in stresses imparted to cell body and processes’, Ann. Biomed. Eng, 33, 52–62, doi:10.1007/s10439-005-8962-y. anderson, e. j., savrin, j., cooke, m., dean, d. and knothe tate, m. l. 2005B, ‘Evaluation and optimization of tissue engineering scaffolds using computational fluid dynamics’, Annual Meeting of the Biomedical Engineering Society, Baltimore, Maryland. ascenzi, m. g. and kabo, j. m. April 1, 2008, Method and system for modeling bone structure, US Patent Number 7,353,153 B2. bar-cohen, y. 2005, Biomimetics – Biologically Inspired Technologies, CRC Press. barocas, v. h. and tranquillo, r. t. 1997, ‘An isotropic biophasic theory of tissueequivalent mechanics: The interplay among cell traction, fibrillar network deformation, fibril alignment, and cell contact guidance’, J. Biomech. Eng., 119, 137. barthelat, f. 2007, ‘Biomimetics for next generation materials’, Phil. Trans. R. Soc. A, 365, 2907–2919, doi:10.1098/rsta.2007.0006. benjamin, m., kaiser, m. e. and milz, s. 2008, ‘Structure–function relationships in tendons: a review’, J. Anat., 212, 211–228. bibb, r. 2006, Medical modelling: the application of advanced design and development techniques in medicine, Woodhead Publishing Ltd, Cambridge, England. buckwalter, j. a., glimcher, m. j., cooper, r. r. and recker, r. 1995, ‘Bone Biology’, J Bone Joint Surg, 77-A(8) 1256. cooke, m. n., fisher, j. p., dean, d., rimnac, c. and mikos, a. g. 2002, ‘Use of stereolithography to manufacture critical-sized 3D biodegradable scaffolds for bone ingrowth’, J. Biomed. Mater. Res. Part B: Appl. Biomater., 64B(2), 65–69. crocker, j. c. 2008, ‘Golden handshake’, Nature, 451(7178), 528–529, doi:10.1038/ 451528a. curry, j. d. 2005, ‘Materials science-hierarchies in biomineral structures’, Science, 309, 253–254, doi:10.1126/science.1113954. dalton, d., klinkhammer, k., salber, j., klee, d. and moeller, m. 2006, ‘Direct invitro electrospinning with polymer melts’, Biomacromolecules, 7(3), 686–690, doi:10.1021/bmo5077q. de santis, r., ambrosio, l. and nicolais, l. 2000, ‘Polymer-based composite hip prostheses’, J. Inorg. Biochem., 79, 97–102. desantis, r., sarracino, f., mollica, f., netti, p. a., ambrosio, l. and nicolais, l. 2004, ‘Continuous fibre reinforced polymers as connective tissue replacement’, Compos. Sci. Technol., 64, 861–871, doi:10.1016/j.compscitech.2003.09.008. dg infso and dg jrc (ed) November, 2005, Towards virtual physiological human: multilevel modelling and simulation of the human anatomy and physiology – white paper, Available at: http://ec.europa.eu/information_society/activities/health/ docs/events/barcelona2005/ec-vph-white-paper2005nov.pdf.
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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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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)
• 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
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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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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.
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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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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
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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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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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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
475
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
477
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
479
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
481
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
800
1000
1200
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
483
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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Biomedical composites 11°
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
485
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
487
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)
488
Biomedical composites 10
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
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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
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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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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
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cell attachment. Journal of Materials Science – Materials in Medicine, 13, 465–468. kazanci, m., cohn, d., marom, g., migliaresi, c. & pegoretti, a. (2002b) Fatigue characterization of polyethylene fiber reinforced polyolefin biomedical composites. Composites Part A – Applied Science and Manufacturing, 33, 453–458. kdolsky, r. k., gibbons, d. f., kwasny, o., schabus, r. & plenk, h. (1997) Braided polypropylene augmentation device in reconstructive surgery of the anterior cruciate ligament: Long-term clinical performance of 594 patients and short-term arthroscopic results, failure analysis by scanning electron microscopy, and synovial histomorphology. Journal of Orthopaedic Research, 15, 1–10. kenny, s. m. & buggy, m. (2003) Bone cements and fillers: A review. Journal of Materials Science – Materials in Medicine, 14, 923–938. kikuchi, m., ikoma, t., itoh, s., matsumoto, h. n., koyama, y., takakuda, k., shinomiya, k. & tanaka, j. (2004) Biomimetic synthesis of bone-like nanocomposites using the self-organization mechanism of hydroxyapatite and collagen. Composites Science and Technology, 64, 819–825. kim, h. y. & yasuda, h. k. (1999) Improvement of fatigue properties of poly(methyl methacrylate) bone cement by means of plasma surface treatment of fillers. Journal of Biomedical Materials Research, 48, 135–142. kolarik, j., migliaresi, c., stol, m. & nicolais, l. (1980) Mechanical properties of model synthetic tendons. Journal of Biomedical Materials Research, 15, 147–157. kurtz, s. m. & devine, j. n. (2007) PEEK biomaterials in trauma, orthopedic, and spinal implants. Biomaterials, 28, 4845–4869. latour, r. a. & black, j. (1993) Development of FRP composite structural biomaterials – Fatigue-strength of the fiber–matrix interfacial bond in simulated in-vivo environments. Journal of Biomedical Materials Research, 27, 1281–1291. lewis, g. (2003) Fatigue testing and performance of acrylic bone-cement materials: state-of-the-art review. Journal of Biomedical Materials Research Part B – Applied Biomaterials, 66B, 457–486. lewis, g. (2008) Alternative acrylic bone cement formulations for cemented arthroplasties: Present status, key issues, and future prospects. Journal of Biomedical Materials Research Part B – Applied Biomaterials, 84B, 301–319. liljensten, e., gisselfalt, k., edberg, b., bertilsson, h., flodin, p., nilsson, a., lindahl, a. & peterson, l. (2002) Studies of polyurethane urea bands for ACL reconstruction. Journal of Materials Science – Materials in Medicine, 13, 351–359. lohbauer, u., frankenberger, r., kramer, n. & petschelt, a. (2006) Strength and fatigue performance versus filler fraction of different types of direct dental restoratives. Journal of Biomedical Materials Research Part B – Applied Biomaterials, 76B, 114–120. mandell, j. f. (1990a) Fatigue behavior of short fiber composite materials. In REIFSNIDER, K. L. (Ed.) Fatigue of composite materials. Amsterdam, Elsevier. mandell, j. f. (1990b) Fatigue behaviour of short fiber composite materials. In REIFSNIDER, K. L. (Ed.) Fatigue of composite materials. Amsterdam, Elsevier. mandell, j. f., samborsky, d. d. & cairns, d. s. (2002) Fatigue of composite materials and substructures for wind turbine blades. Albuquerque, New Mexico, USA, Sandia National Laboratories.
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marois, y., roy, r., vidovszky, t., king, m. w., belanger, a. y., chaput, c. & guidoin, r. (1993) Histopathological and Immunological Investigations of Synthetic-Fibers and Structures Used in 3 Prosthetic Anterior Cruciate Ligaments – Invivo Study in the Rat. Biomaterials, 14, 255–262. marrs, b., andrews, r. & pienkowski, d. (2007) Multiwall carbon nanotubes enhance the fatigue performance of physiologically maintained methyl methacrylate– styrene copolymer. Carbon, 45, 2098–2104. marrs, b., andrews, r., rantell, t. & pienkowski, d. (2006) Augmentation of acrylic bone cement with multiwall carbon nanotubes. Journal of Biomedical Materials Research Part A, 77A, 269–276. martin, r. (2003) Delamination fatigue. In HARRIS, B. (Ed.) Fatigue in composites. Cambridge, UK, Woodhead Publishing Ltd. matsuda, s., kishi, h. & murakami, a. (2004) New processing of apataite fiber and some application for medical use. Composites Science and Technology, 64, 909–914. mccabe, j. f., wang, f. & braem, m. j. a. (2000) Surface contact fatigue and flexural fatigue of dental restorative materials. Journal of Biomedical Materials Research, 50, 375–380. mccool, j. i., boberick, k. g. & baran, g. r. (2001) Lifetime predictions for resinbased composites using cyclic and dynamic fatigue. Journal of Biomedical Materials Research, 58, 247–253. mcgregor, w. j., tanner, k. e., bonfield, w., bonner, m. j., saunders, l. s. & ward, i. m. (2000) Fatigue properties of isotropic and hydrostatically extruded HAPEX (TM). Journal of Materials Science Letters, 19, 1787–1788. mendes, d. g., iusim, m., angel, d., rotem, a., mordehovich, d., roffman, m., lieberson, s. & boss, j. (1986) Ligament and tendon substitution with composite carbon fiber strands. Journal of Biomedical Materials Research, 20, 699–708. miettinen, v. & vallittu, p. (1997) Release of residual methyl methacrylate into water from glass fibre-poly(methyl methacrylate) composite used in dentures Biomaterials, 18, 181–185. migliaresi, c., nicoli, f., rossi, s. & pegoretti, a. (2004) Novel uses of carbon composites for the fabrication of external fixators. Composites Science and Technology, 64, 873–883. migliaresi, c. & pegoretti, a. (2002) Fundamentals of polymeric fiber-reinforced composite materials. In BARBUCCI R (Ed.) Integrated biomaterials science. New York, Kluwer Academic/Plenum Publishers. mohandesi, j. a., rafiee, m. a., barzegaran, v. & shafiei, f. (2007) Compressive fatigue behavior of dental restorative composites. Dental Materials Journal, 26, 827–837. moore, d. r. & turner, s. (2001) Mechanical evaluation strategies for plastics, Cambridge, UK, Woodhead Publishing Ltd. o’brien, f. j., taylor, d. & lee, t. c. (2007) Bone as a composite material: The role of osteons as barriers to crack growth in compact bone. International Journal of Fatigue, 29, 1051–1056. o’brien, t. k. (1990) Delamination of composite materials. In REIFSNIDER, K. L. (Ed.) Fatigue of composite materials. Amsterdam, Elsevier. oral, e., malhi, a. s. & muratoglu, o. k. (2006) Mechanisms of decrease in fatigue crack propagation resistance in irradiated and melted UHMWPE. Biomaterials, 27, 917–925.
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20 Nanostructured biocomposites for tissue engineering scaffolds D. M E N G, and A. R. B O C C AC C I N I, Imperial College London, UK
Abstract: Many biomaterials, from natural sources to synthetic products, are being used to fabricate scaffolds for tissue engineering and regenerative medicine. Materials including polymers, ceramics, glasses, metals and their composites are considered. Although the bulk materials chemistry (composition) of the scaffold is an important factor determining the scaffold performance, e.g. in terms of mechanical properties, structural integrity, degradation behaviour and overall biocompatibility, recent studies have shown that the roughness and topography of the scaffold surface play an essential role too, especially the presence of nano-scale features, which directly affects cell–material interaction, cell attachment, migration and proliferation. This chapter reviews current processing techniques proposed to introduce or engineer nanoscale topography on the surfaces of 3D scaffolds with the focus on biocomposites, and including bottom-up and top-down approaches. Various case studies are discussed, briefly highlighting the positive effect of nanoscale topographies on cell–material interactions and the need for further quantitative investigations. Key words: scaffolds, tissue engineering, nanocomposites, surface topography, porosity.
20.1
Introduction
The aim of tissue engineering and regenerative medicine (TERM) is to restore or regenerate a damaged tissue by combining cells, derived from a patient biopsy or from other sources, e.g. stem cells, with a 3D porous support, called a scaffold, that acts as a temporary extracellular matrix (ECM) during the process of new tissue growth.1 For its successful application in TERM strategies, a scaffold needs to exhibit biocompatibility, tailored biodegradability and/or bioresorbability, it should also have suitable mechanical properties, adequate physicochemical behaviour to direct cell– material or cell–cell interactions. Moreover, scaffolds should replicate the original morphology and structure of the damaged tissue promoting the integration with the surrounding biological environment.2,3 The goal is for the cells to attach to the scaffold, multiply, differentiate and organise themselves into healthy new tissue as the scaffold degrades. 509
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A wide range of biomaterials is being investigated for TERM scaffolds.2–4 For the case of bone and cartilage tissue engineering, in particular, typical biomaterials used include hydroxyapatite, calcium carbonate, bioactive glass and biodegradable polymers, such as poly(lactic acid) (PLA), poly(glycolic acid) (PGA), polycaprolactone (PCL) and poly(3-hydroxybutyrate (P3HB). However, none of these materials used on their own can satisfy all the goals required for creating optimal tissue scaffolding, such as suitable fracture strength, stiffness, toughness, osteoinductivity (the ability of a substance to initiate bone formation in a non-bony site), osteoconductivity (support bone growth and encourage the in-growth of surrounding bone) as well as controlled rate of degradation in vivo and in vitro. Therefore, the tailored combination of biomaterials to form biocomposites is being increasingly considered for the development of optimal scaffolds.5,6 An important aspect in the design of scaffolds for TERM applications is their required 3D porous architecture. Scaffold porosity is important for the development of new tissue as pores provide the space in which cells reside.7 It is now recognised that highly porous and interconnected structures are required to allow cell seeding and migration throughout the entire scaffold. The combination of different pore sizes in hierarchical 3D structures allows not only the growth and colonisation of the relevant cells but also capillary in-growth for vascularisation of the new tissue as well as providing the path for diffusion of nutrients and metabolic waste.8 Material composition (chemistry), microstructure, 3D porosity and structural properties are not the only factors determining the success of a particular scaffold type. It is well known that cell response in contact with solid surfaces is affected by the surface physicochemical parameters, such as surface energy, surface charges and chemical composition.9 Figure 20.110 shows a diagram demonstrating the complex relationship between cells and scaffolds, which is determined by the microenvironment created around cells adherent to the TERM scaffold. Nutrient transport brings growth factors, ligands, and other signals that can bind to cell receptors. The degrading scaffold can likewise release chemical signals that bind to membrane receptors and influence intracellular communication and cellular processes such as gene transcription. Cells also attach to the scaffold via integrin receptors. Integrin receptors are closely connected to the cell’s cytoskeleton and relay further information to the cell thereby affecting cell function.10 Surface topography is one of the most decisive physical cues for cells.11 The primary cell–material interaction, which affects how cells attach to the scaffold surface and their subsequent behaviour, is strongly dependent on surface roughness, with the presence of nanoscale features playing a key role.11–15 This is because although cells have micrometre dimensions, they evolve in vivo in close contact with the ECM, a substratum with topographi-
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Nutrients, ligands and growth factors are brought to cells within the scaffold via fluid flow where they are transported into the cell or bind to cell membrane receptors
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CO2 and other metabolic waste products are moved away from the cell
Integrin receptors interact with ECM molecules which are connected to the cell’s cytoskeleton, which can in turn influence cell behaviour
Growth factors and other molecular signals are released from the scaffold and bind to cell receptors Growth factors and other ligands bind to cell triggering intracellular communication membrane receptors initiating intracellar cascades which influence gene expression
20.1 Schematic diagram demonstrating the complex microenvironment created around a cell adherent to a tissueengineering scaffold. (Reproduced from reference 10, with permission of Elsevier.)
cal and structural features of nanometer size.9 Moreover cells have many structures in the nanoscale, such as filipodia and cytoskeletal and membrane proteins that interact with the environment surrounding them.16 There is increasing interest in developing methods to modify biomaterial surfaces in order to mimic the nanoscale topographical features presented to cells. These nanoscale surface features should promote functions such as cell adhesion, cell mobility and cell differentiation.17–21 Studies have confirmed that nanotopographies, e.g. nanofibrous structures, can enhance protein adsorption, including fibronectin and vitronectin, hence increasing cell attachment.18 In addition, topographical changes on a bioactive surface can affect local ion exchange properties,11 e.g. the simultaneous effect of topography and topography-mediated ion exchange on cellular behaviour must be known in order to discern the different roles of topography and surface chemistry. In some applications, it is essential to use topography to guide the cell orientation in order to achieve a functional tissue, such as tendons, nerves and intervertebral-disc regeneration. It is, thus, now well established that the contact guidance of cells by combination of hierarchical
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micro- and nano-topographical patterns on scaffold surfaces is a promising perspective for TERM.22,23 This strategy requires, however, the development of material processing and engineering methods to induce combined micro- and nanostructured topographies on 3D scaffold surfaces, including the development of novel nanoscale biocomposites. A number of techniques can be used to create ordered and unordered nanoscale topographies on biomaterials. Electron beam lithography24 and photolithography25 are two standard techniques for creating ordered features on a variety of material surfaces, primarily used in miniaturisation strategies in microelectronics, however, only in 2D, e.g. on flat surfaces. Phase separation, colloidal lithography and laser assisted patterning26 are other techniques developed to create patterns in 2D. Indeed, most previous research investigating the effect of surface nanotopography on cell function and tissue formation has been carried out on flat surfaces.27 There are only a few demonstrated methods, such as polymer demixing and chemical etching, that can be used to create (unordered) surface patterns in both 2D and 3D structures.16 In several cases, a 3D structure can be made by combining previously patterned 2D surfaces, e.g. by rolling a flexible material (usually a polymer or composite film) to form tubular scaffolds or by assembling the 3D architecture in a ‘LEGO® fashion’ from patterned smaller units. In the last few years, however, a number of novel approaches have been developed with the specific aim of incorporating nanoscale features on the surfaces of 3D TERM scaffolds. These approaches involve top-down as well as bottom-up methods. Top-down methods manipulate matter at a very small scale to engineer a desired structure, as is the case with miniaturisation strategies in microelectronics, whereas bottom-up techniques are based on exploiting intermolecular forces to self-order or self-assemble a structure mimicking natural materials such as proteins and DNA.28 In this chapter, the current state-of-the art in the field of processing 3D scaffolds with engineered nano-scale surface topography is discussed. In several cases, this includes the development of composites, the extension or adaptation of techniques available for micrometre-scale topographies or the assembly of 3D architectures from previously patterned 2D surfaces. Some typical TERM applications of 3D nanostructured scaffolds and specific biological responses are highlighted when describing the different approaches taken from selected reports. These studies refer generally to osteoblast cells and bone regeneration, but the 3D nanostructured scaffolds described here are relevant for applications in several other areas of TERM. Two-dimensional nanostructured biomaterials are not covered in this chapter, except when they are used to fabricate 3D scaffolds, relevant reviews are reported elsewhere.29–31 It should also be mentioned that bionanopatterning of surfaces, i.e. the patterning of biomolecules on surfaces with nanometre resolution, is beyond the scope of the topics covered in this
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chapter. A recent review on bio-nanopatterning technology with particular focus on the recent advances on the use of nanolithographic techniques as tools for biomolecule immobilisation at the nanoscale is available.32
20.2
Processing of 2D topographies for assembly of 3D (biocomposite) structures
20.2.1 Photolithography Performing photolithography in the near-UV range is a standard fabrication method for micro-scale topographies, as precise geometries and patterns can be created on a variety of materials. This process usually involves exposing a silicon wafer coated with a photoresist to a near-UV light source (typically 365 or 405 nm in wavelength) through a mask that has the desired pattern on it. This mask selectively allows light through to the wafer thereby recreating the pattern in the photoresist. Owing to the limitation of this method, that is the wavelength of light used, nanometer-size structures can not usually be created.16 However, with recent developments, shorter wavelength lasers can allow the exposure of even smaller features using photolithography. Wavelengths in the deep UV (157 nm) and ‘soft’ x-ray regions (2–50 nm) may permit an extension of the current lower limits of photolithography to the natural limit of complementary devices based on metal oxide semiconductors.25 However, this technique has not yet been applied to create nanoscale topography in 3D tissue engineering (biocomposite) scaffolds to our knowledge. Photolithography is still an important method, as it can be used in combination with other techniques to create the desired nanoscale patterns. For example, Wood et al.33 created patterned colloidal nanopillared topographies on the base substrate of quartz or silicon N-(βaminoethyl)-Ψ-aminopropyltrimethoxysilane (AAPS) by combining and utilising lithography and conventional photolithography techniques, which made it possible to devise a method where planar and nanopillared surfaces can lie collaterally on the same device. In a TERM relevant study, Kenar et al.34 used photolithography technique (with UV light) to generate microgrooves and micropits on singlecrystal silicon wafers that were coated with positive photoresists to create a variety of templates for studying the influence of pattern types on cell adhesion. For example, poly(3-hydroxybutyrate-co-3-hydroxyvalerate) (PHBV) films were prepared using solvent casting on the master templates and, hence, the negative replication of the master patterns was generated. Their study showed that microtopography can improve osseointegration when combined with chemical cues (fibrinogen being the most effective one), and selective osteoblast adhesion and alignment can be guided by the microgrooves and cell adhesive protein lines on the PHBV films. Lee et al.35
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demonstrated the combination of photolithography (UV light) with protein microarraying to precisely controlling ‘on-the-spot’ cell-cell interactions. A micropatterned photoresist layer that was spin-coated on a silane-modified glass substrate served as a temporary stencil during the microarray printing. The micropatterns defined the dimensions and the geometry of the cell adhered inside the printed protein spots. It was observed that hepatic (HepG2) cells could attach on 300 or 500 μm diameter protein spots. The extent of contacts between cells within each spot varied in agreement with dimensions of the photoresist stencil, from single cells attaching on 30 μm diameter features to multi-cell clusters residing on 100 or 200 μm diameter regions. It has also been shown that the combination of photolithography and photochemical techniques on specific biomaterials, for example polyethylene glycol hydrogels,36 can lead to the patterning of non-planar surfaces and complex 3D environments.37 Another extension of photolithography for biomedical applications has been presented by Papenburg et al.,38 who showed the development of a new method called phase separation micromolding, a technique based on immersion precipitation on a micropatterned mould fabricated from a silicon wafer using photolithography.
20.2.2 Electron beam lithography Electron beam lithography (EBL), used in nanofabrication of high-density storage media and electronics, involves the use of high-energy electrons to expose a positive or negative electro-sensitive resist such as PMMA.24 The resist can be developed to produce nanotopographies. The mask used in photolithography is not required to pattern the surface with the beam of electrons in this case, because the pattern is programmed into the unit in order to precisely control how the beam travels over the resist-coated surface. EBL has the ability to create single surface features down to about 3–5 nm.39 However, single features are not of practical use in studying cellular growth and behaviour, in contrast to that, high density arrays of a single feature such as posts or channels are used, which allows for enough surface area to be covered with the pattern to observe a cell population’s response to the surface features. Vieu et al.39 extended the limit of EBL to approximately 30–40 nm in order to create a relatively large surface area of features on bulk silicon samples with a 200 nm thermally grown oxide layer and 50 nm spin coated PMMA resist layers, as shown in Fig. 20.2a. EBL can be time consuming and costly in order to fabricate these features. When dealing with microstructural surfaces, in order to mass produce the pattern, a master pattern is normally created on silicon wafer by EBL, then it is replicated by using some type of curable polymer, such as polydimethylsiloxane.40 Nanometer patterns can be replicated using the same principle, making the mass reproduction of the patterns much faster
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(a) A
B
C
D
(b)
20.2 (a) SEM image of a 20-nm line and space array obtained by EBL on silicon, the exposure dose was 11 nC cm−1.39 (b) SEM micrographs of the different stages of the fabrication process of a 3D vascular scaffold. Nanopatterned masters with pits (A) or pillars (B) were made by electron beam lithography. C The micropatterned master with cell guidance and spacers was made by photolithography. D Cross-section of the complete construct. Scale bars are (A + B) 1 μm and (C and D) 100 μm.42 (Images reproduced with permission of Elsevier.)
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and less expensive. This has been demonstrated in polycaprolactone using hot embossing.41 A combination of EBL and photolithography has been successfully applied to fabricate a three-layered vascular scaffold that supports three different cell types individually on each of the layers in a coaxial arrangement (Fig. 20.2b).42 Nano-pits (Fig. 20.2b-A) or pillars (Fig. 20.2b-B) created using EBL on the surface of the inner layer promoted endothelia cells, while the middle layer (Fig. 20.2b-C) arranged smooth muscle cells uniaxially with micro-scaled linear patterns created using photolithography, and an outer layer of collagen deposited by fibroblasts completed the complex composite layered structure. Although the final structure of the scaffold is in 3D (Fig. 20.2b-D), the topography of each layer was created in a two-dimensional manner, that is on flat surfaces first and then combined and rolled to form the 3D vascular scaffold. Gomez et al.43 fabricated 1 and 2 μm wide microchannels using electronbeam (e-beam) lithography and electropolymerisation on polypyrrole (PPy), a well-known electrically conducting polymer which has been also considered to fabricate composite scaffolds in combination with hydroxyapatite.44 The depth and the roughness of the microstructure increased as the e-beam writing current increased. It was found that embryonic hippocampal neurons cultured on patterned PPy polarise (i.e. defined an axon) faster on this modified material. Axon orientations were affected by the topography features, but this effect was not pronounced on the overall axon length.43
20.3
Direct fabrication of surface nanotopographies in 3D structures
20.3.1 Polymer demixing This technique is used to obtain randomly organised patterns on inorganic material surfaces. It is based on spontaneous phase separation of a polymer blend under the spin-casting process onto a rigid substrate,45 such as silicon or glass. Various shapes, for example, pit, island and ribbon, can be obtained, and the size of the features can be modulated by the careful adjustment of the polymer ratio or the polymer concentration, respectively.9,16 This technique is a fast and inexpensive method for creating 2D cell culture substrates with nanoscale features. These features are somewhat uncontrollable in the horizontal direction, and also in the vertical direction, as the features appear randomly placed over the surface area.46 Owing to the lack of control over the pattern organisation, polymer demixing is not an ideal method for creating model symmetrical and well-defined surfaces to study the interaction between cells and their nanoscale topography sur-
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roundings. Despite this limitation, some previous studies have shown fibroblast filipodia interacting specifically with 10 nm high islands on surfaces obtained by polymer demixing.46 Although fibroblasts were seen to interact quickly and strongly with these nano-islands, when the height of these features was >27 nm, the cells started to remain motile on top of the islands instead of being strongly in contact with the material, as they are on a flat surface. Dalby et al.47 observed that with 27 nm high nano-islands produced by polymer demixing, there was a rapid adhesion and cytoskeletal formation of InfinityTM Telomerase Immortalised human fibroblasts on the islands at 4 days of culture. However, after 30 days of culture, the cells grew less well on the nano-islands compared with flat control surfaces. Unlike the techniques described in sections 20.2.1 and 20.2.2, which might only be used to create nanoscale patterns on flat 2D surfaces, polymer demixing allows nano-topographies to be obtained also on 3D structures. Berry et al.,48 for example, demonstrated that nano-scale islands can be created inside thin tubes (diameter 8) for almost the whole process. Under these conditions, the fibril formation takes place before the precipitation of calcium phosphate, which is crucial since the collagen fibrils act as templates for the mineralization. Upon decrease of pH below ∼8, amorphous HA forms onto the fibrils before their assembling into fibres. As the pH approaches a neutral value, two processes enter in competition, involving
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the same binding chemical groups on the surface of the fibres: the organization of collagen fibres into a three-dimensional network and the proceeding HA crystallization (Gelinsky et al., 2007b; Mann, 2001; Tampieri et al., 2003). Finally, a biocomposite is obtained made by fibrillated collagen (Coll) embedding nanocrystals of doped HA; the water in excess is further eliminated by a freeze-drying process. The final porosity of the spongy mineralized composites is strongly dependent on the freezing temperature, heating/cooling ramp and the content of water in the starting gel. For example, the adoption of a freezing temperature of −25 °C yields the formation of large and anisotropic pores, with the largest dimension in the range 250–450 μm (Fig. 21.1). The lowering of the freezing temperature up to −40 °C induces a decrease of the average pore diameters and an even more evident pore anisotropy (Gelinsky et al., 2007b). X-ray micro-CT with a local resolution of 10 μm can be successfully employed to evaluate the total porosity of HA/Coll composites (usually ranging between 80–85%), and the pore wall thickness (usually 15–20 μm). Studies performed by Bigi et al. (1991) on turkey leg flexor tendon compared with HA/Coll composites, using TG–DTA analysis, gave evidence of the interaction between inorganic phase and collagen: the inorganic phase hampers the combustion of organic phase and the profile of TG–DTA curves for non-mineralized and mineralized collagen are perfectly superimposed respectively with those of young and mineralized turkey tendon (Tampieri et al., 2003).
100 μm
21.1 Porosity in self-assembled collagen.
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Intensity (a.u.)
The synthesis of the doped HA/Coll composite at room temperature (∼25 °C) allows the formation of a poorly crystallized HA (Bouyer et al., 2000), as illustrated by the XRD pattern reported in Fig. 21.2. Here, the nanometric size of the HA crystallites is responsible for the large broadening of the XRD reflections: such nanocrystals grow inside collagen fibres with their c axis preferentially oriented parallel to the direction of orientation of the fibres (see insert in Fig. 21.2). The TEM micrograph in Fig. 21.3 shows nanometric nuclei of HA formed onto the collagen fibrils and merged inside the growing collagen fibres.
0.344 nm 0.29 nm
Collagen fibre axis 24
28
32 2-theta (deg)
36
21.2 XRD spectrum of HA/Coll. The insert shows the orientation of the c axis of the HA phase parallel to the long axis of collagen.
100 nm
21.3 Nuclei of HA formed inside a collagen fibre.
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Table 21.1 ICP-OES quantitative analyses of apatite/collagen materials
HA/Coll (70/30)wt MgHA/Coll (70/30)wt MgHA/Coll (40/60)wt MgSiHA/Coll (70/30)wt
Ca2+ (mol)
Mg2+ (mol)
PO43− (mol)
SiO44− (mol)
Ca/P (mol)
Mg/Ca (mol) %
Si/P (mol) %
1.870 1.315 1.419 1.071
– 0.035 0.027 0.073
1.144 0.857 0.983 0.696
– – – 0.021
1.635 1.534 1.444 1.539
– 2.662 1.903 6.816
– – – 3.017
A confirmation of the stoichiometric deviation in the HA nucleated on collagen is provided by elemental analysis via inductively coupled plasma optical emission spectroscopy (ICP OES) of the HA/Coll composites (Table 21.1). Although the mineral part of bone can probably be described as a calcium hydroxyapatite [Ca10(PO4)6(OH)2] (Suchanek and Yoshimura, 1998), it is now established that the minor ionic species present in the mineral bone, such as CO32−, Mg2+, SiO44− and HPO42− are crucial in the biochemistry of bone remodelling (Bigi et al., 1997; Driessens, 1980). For this reason, the preparation of synthetic HA phases doped with foreign ions incorporated in their structure has been studied (Sprio et al., 2008) and the improvement of its bioactivity and bioreabsorbability when tested in vitro and in vivo has been assessed (Landi et al., 2008). Among substituting cations, magnesium is particularly relevant and widely studied: it has been verified that in calcified tissues, the amount of magnesium associated with the apatite phase is higher (about 5% at.) at the first stages of the bone remodelling process and decreases with increasing calcification and with the ageing of the individual (Bigi et al., 1992). The presence of magnesium increases the nucleation kinetic of HA on collagen fibres and temporarily retards its crystallization; for this reason the concentration of magnesium is higher in the cartilage and in young bone and there is growing evidence that it may be an important factor in the qualitative changes of the bone matrix that determines bone fragility. Magnesium depletion adversely affects all stages of skeletal metabolism, causing cessation of bone growth, decrease of osteoblastic and osteoclastic activities, and osteopenia. In addition, silicon is an essential element for the healthy skeletal and cartilage growth and development in the higher biological organisms (Carlisle, 1970, 1979; Pietak et al., 2007; Schwarz and Milne, 1972; Seaborn and Nielson, 2002). Silicon also influences the cartilage synthesis and the integrity of the extracellular matrix, as well as the biomineralization process. It has dose-dependent effects on the differentiation, proliferation and collagen synthesis of osteoblasts, with a direct influence on the bone remodelling process and osteoclast development and resorption activities.
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50 nm
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40 nm
c FT
200 nm (a)
500 nm (b)
21.4 TEM micrograph of freeze dried: (a) HA/Coll (70/30 wt%) composite; and (b) MgHA/Coll (70/30 wt %) composite. Inset (c): highresolution HREM images and Fourier transform.
Additional support for silicon’s metabolic role in connective tissue is provided by the finding that silicon is a major ion of osteogenic cells, especially high in the metabolically active state of the cell. Silicon is also known to bind to glycosaminoglycan macromolecules and has been shown to play a role in the formation of cross-links between collagen and proteoglycans (Schwarz, 1973), thus resulting in the stabilization of bone matrix molecules and preventing their enzymatic degradation. For the composites Si–Mg–HA/Coll 70/30% wt, the deviation from HA stoichiometry increases the similarity to natural bone tissue so that not only microstructure but also the composite composition resembles that of bone which causes an increase in the bioactivity and biodegradability of the composite itself (Guillemin et al., 1995; Rendey et al., 1999; Landi et al., 2006; Tampieri et al., 2004). The presence of magnesium and silicon further inhibits the crystallization of HA; the mineral phase has been evaluated from a typical set of highresolution electron microscopy (HREM) images (Fig. 21.4c). The HA crystals nucleated on the organic fibre exhibit a rod-like structure (Fig. 21.4a). In the presence of the substituting ions, the nuclei of mineral phase have a globular shape and smaller dimensions (Fig. 21.4b). The inset in Fig. 21.4c shows, at higher magnification, the Si–Mg–HA nucleus on collagen fibre. The analysis of the high-resolution images and its Fourier transform (FT) (inset in Fig. 21.4c) reveals that this shape is accompanied by a lack of order and the particles appear completely amorphous, as indicated by the intensely diffuse ring and by the complete lack of spots in the FT of the image. It
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must be stressed that the degree of disorder observed does not correspond to a short range order (as the case of low crystalline inorganic HA) but to a ‘true amorphous’ structure (Bertinetti et al., 2004). These data indicate that, despite similar synthesis conditions being used, the structural features of the mineral phase in the composite are completely different from that nucleated in the absence of an organic template (Celotti et al., 2006). In fact, during mineralization a structural control was accomplished through the preferential nucleation of a specific crystal face/axis, by molecular recognition, at the surface of the organic template. Since Mg has different polarity, structure and stereochemistry than Ca, the activation energy controlling the nucleation rate changes, and this is reflected in a modification of site specificity, mineral structure and crystallographic alignment (from which the definition of the biologically inspired process arises) (Bertinetti et al., 2006; Mann, 2001). Si–Mg–HA nucleated on collagen is spontaneously carbonated (as also stated for the HA/Coll composite) and therefore the theoretical formula of the mineral phase becomes: ( Ca, Mg )10 -x 2 ( PO 4 )6 -x -y ( CO3 )x ( SiO 4 ) y ( OH )2 -y
[21.1]
Thus, as in the natural mineralization process, it was possible to recognize a chemical control mechanism that occurred during the nucleation and crystals growth and was mainly achieved by the regulation of ion movement and the kinetic competition between the fibrils assembling and HA nucleus formation (Mann, 2001). Other evidence of the chemical interaction between HA and collagen fibres comes from the study of the FTIR spectra (Fig. 21.5a) in which a shift of the band corresponding to −COO− stretching from 1340 to 1338 cm−1 can be detected. The band at 873 cm−1 strengthens for the composite HA/Coll (70/30), indicating that the nucleation of HA into collagen implies carbonation of inorganic phase. Moreover, the carbonation can be assigned only to the B position as confirmed by the absence of the band at 880 cm−1 and by EDS analysis which reveals that the increase of the C content (in the CO32− groups) corresponds to a decrease in phosphorous concentration. The interaction of HA with collagen seems to prevent carbonation in the A position probably blocking the access to OH groups. In Fig. 21.5, a comparison between of HA/Coll, Mg–HA/Coll and Si–Mg– HA/Coll is reported. Commonly, in the FTIR spectra of crystalline HA, the low site symmetry of the PO43− tetrahedron splits the ν4 PO43− contour into three components (Fig. 21.5b); this triplet tends to merge into two bands and one broad band in the HA/Coll composite and in the doped HA/ Coll composite, respectively. In addition, the broad band centred between 550 and 560 cm−1, assigned to acid phosphate (HPO42−) in the mineral lattice, was much more intense in the doped HA nucleated on collagen.
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(a) 0.24 MgHA/Col (70/30)wt 0.22
Absorbance
0.20
MgSiHA/Col (70/30)wt
0.18 COO– band 1338 cm–1
0.16
CO32–band 873 cm–1
0.14
HA/Col (70/30)wt
0.12 0.10 0.08
Crystalline HA
0.06 2000
1800
1600
1400 1200 1000 Wavenumbers (cm–1)
800
600
(b) 0.26
Absorbance
0.24 0.22
MgHA/Col (70/30)wt
0.20
MgSiHA/Col (70/30)wt
0.18 0.16 0.14
HA/Col (70/30)wt
0.12 0.10 0.08
Crystalline HA 700
650
600 550 Wavenumbers (cm–1)
500
450
21.5 FTIR spectra of doped HA/Coll: (a) showing shift of –COO− band and (b) the three components of the ν4 PO42− contour.
These peculiarities of Si–Mg–HA/Coll can be similarly observed in the FTIR spectrum of young bone; whereas they tend to disappear in mature bone as well as in highly crystalline synthetic apatite (Fowler et al., 1966; Miller et al., 2001; Rey et al., 1990). Correlation was also noted between the fractional intensity of ∼1050 cm−1 band related to PO43− group and the
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Table 21.2 Mechanical properties of HA/Coll composites with different porosities Sample porosity (%)
Flexural strength σf (MPa)
Young’s modulus E (GPa)
45.3 45.9 46.7 51.5 52.4 54.1 62.0 62.6 63.3
17.56 23.23 7.46 7.86 8.77 9.05 4.68 4.92 6.11
6.853 4.514 6.286 3.438 2.005 3.906 1.809 1.654 1.504
crystal size of the apatite: a reduction of the apatite crystal size induced an increase of the percentage area of this component (Paschalis et al., 1996; Pleshko et al., 1991). All these features indicate that the mineral phase containing Si and Mg ions and nucleated on the natural template (collagen) is non-stoichiometric and remarkably amorphous resembling very well the features of newly deposited bone mineral. Flexural strength (σf) and elastic modulus (Young’s modulus E) were determined on the HA/Coll 70/30 % wt dry composites revealing a pseudoplastic behaviour. The σf and E values were calculated on the basis of regression curve and are listed in Table 21.2 vs. the sample porosity. The elastic modulus determined on the mineralized layer well reproduces the value found for trabecular bone at correspondent values of porosity (Table 21.2).
21.4
Scaffolds with hierarchically organized structure: inspiration from nature
The morphology of human bone is characterized by a complex hierarchic structure exhibiting both gradients of density and anisotropic properties, which confer its unique properties of lightness, resistance, elasticity and capacity to self-regenerate following trauma of limited severity. The peculiar structure of bone also allows an efficient distribution of the mechanical loads up to the smallest trabeculae, where the strain-sensing mechanism exhibited by the osteocyte cells can take place and activate the permanent bone remodelling process. This mechanism allows the continuous rearrangement of the bone tissue in the presence of biophysical stimuli from the environment, so that the local damage is minimized and the functionality of the bone as a whole is optimized. Thus, only a scaffold exhibiting structures with a high degree of hierarchy is able to yield the optimal bio-
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physical response, after implantation in vivo, particularly in the healing of long bone disease, where the biomechanical loads are more complex and strong. At present, the development of hierarchically organized bone scaffolds would represent a breakthrough, surpassing the current solutions for repairing bone and osteochondral tissues. The vegetable kingdom includes a great variety of hierarchical tissue structures with impressive biomechanical properties; for example, some tree branches can support the weight of an entire house, and others can be bent and stretched into any shape. In this respect, native or semi-processed wood and plants may be successfully used as templates for generating ceramic materials through multistep transformation processes, involving pyrolysis and chemical reactions (with solid, liquid or gaseous phases). Using this material, exhibiting a hierarchic structure very close to the ones of specific regions of natural bone, lightweight scaffolds can be developed that have improved biomechanical properties. Furthermore, the integration of this technology with bio-hybrid composites obtained by biologically inspired mineralization, strongly increases the therapeutic potentiality. Particularly, devices for the healing of large osteochondral defects (more than half of a condyle), are going to be developed using external retentive structures made of resistant hierarchically organized HA ceramics (produced as described previously in this section) and internal soft core made of HA/Coll bio-composite as subchondral scaffold covered by a cartilage-like layer. Preliminary investigation on ceramization processes of native wood structures has led to the development of porous bodies in silicon carbide (Esposito et al., 2004; Greil et al., 1998a, 1998b, 2001; Sieber et al., 2000, 2002) by wood pyrolysis and subsequent infiltration with melted silicon; BioSiC® technology was also developed (Feria, 2005; Gonzalez et al., 2004; Kardashev et al., 2005; Parfen’eva et al., 2005) with similar technology, to manufacture porous bioinert compatible ceramic bodies, characterized by excellent mechanical performances and morphology very well resembling the graded morphology of cortical-spongy human bone. Similarly, starting with the carbon preform, through a sequence of hydrothermal treatments and vapour/liquid infiltration processes (Ruffini et al., 2008), bioactive ceramic materials showing hierarchically organized pore structures can be obtained. In more detail, starting from vegetal raw materials, selected on the basis of their structure, a pyrolysis treatment yields carbon templates, which are subsequently infiltrated by vapour-phase calcium to produce calcium carbide (carburization) (formula [21.2]); (Fig. 21.6). The conversion of carbon into CaC2 induces an important volume increase of the solid and the formation of micropores where the gas can get in. Two possible steps for vapour-phase process are: the infiltration of the
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30 μm
60 μm (a)
(b)
21.6 SEM micrographs of pyrolysed wood (a) before and (b) after the transformation in CaC2.
gaseous calcium into the carbon preform and the heterogeneous gas/solid reaction between gaseous calcium and carbon to form CaC2 on the structure surface, followed by the diffusion of gaseous calcium through the CaC2 layer just formed. Subsequently, the oxidation is performed and the gas–solid CO2–CaO reaction proceeds and it is controlled by the kinetics of the heterogeneous surface chemical reaction. Subsequently, a porous calcium carbonate structure is obtained from calcium oxide via the addition of compressed CO2 into a high-pressure closed autoclave. The chemical reaction involved in the processes is shown in equation [21.4]. Following this initial stage, as a compact layer of CaCO3 is formed on the surface of CaO, the rate of reaction decreases owing to the diffusion limitation of CO2 through the CaCO3 layer. Finally, the phosphation treatment transforms porous calcium carbonate into hydroxyapatite (formula [21.5]). 2C + Ca → CaC2
[21.2]
CaC2 + H 2 O → CaO + C2 H 2
[21.3]
CaO + CO2 → CaCO3
[21.4]
10CaCO3 + 6KH 2 PO 4 → Ca10 ( PO 4 )6 ( OH )2 + …
[21.5]
The x-ray diffraction (Fig. 21.7) of the materials resulting from the phosphation treatment shows the formation of pure hydroxyapatite as the final product, whereas Fig. 21.8 shows the hierarchically organized structure of HA as grown during the hydrothermal process.
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Counts CaCO3 + KH2PO4 200 °C 12 h piece 3600
1600
400
0 10
20
30
40 50 Position (2Theta, deg)
60
21.7 XRD analysis after hydrothermal treatment of calcium carbide.
100 μm
20 μm
21.8 Micrographs of wood transformed in porous hydroxyapatite.
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21.5
Development of the three-layered osteochondral scaffold
Along the direction of designing biomimetic osteochondral composite scaffolds resembling the composition of the extracellular matrices of cartilage and bone tissue, a biologically inspired approach has been considered the most effective. Therefore, chemically and morphologically graded hybrid materials have been generated, built by stacking a lower mineralized layer (see section 21.3), corresponding to the subchondral bone, an intermediate layer with reduced amount of mineral to mimic the tidemark and an upper layer formed by collagen and hyaluronic acid (HY), reproducing some cartilaginous environmental cues (Tampieri et al., 2008). The cartilage-like layer can be prepared by adding pure collagen with hyaluronic acid to create bridges between the collagen fibres, then the introduction of this saccharidic structure rich in polar hydrogen atoms, will improve the hydrophilicity of the system. The porosity of the cartilaginous layer is represented by pores more isotropic than the mineralized composite and the average pore diameter is in the range of 100–150 μm. The presence of the HY acid within the structure is also confirmed by solid state NMR analysis performed on the collagenous layer with and without hyaluronic acid. The addition of 1,4-butanediol diglycidyl ether (BDDGE) as a cross-linking agent able to stabilize collagen and retard its degradation kinetic, is necessary in all the three different layers constituting the scaffold. However, the effect of the cross-linking agent on the material morphology is more evident on collagen when the mineral phase is absent. Enzymatic tests carried out on the different layers using collagenase reveal that the mineral phase retards the degradation of collagen and BDDGE further stabilizes the HA/Coll composite: HA/Coll (70/30) % wt degrades completely in 38 h, whereas HA/Coll (70/30) % wt + BDDGE takes 78 h. Figure 21.9 displays the ESEM image of the graded composite: it is possible to distinguish a disordered lower layer that corresponds to the mineralized one, a second layer, with lower mineralization extent, which corresponds to the tidemark and a third layer in which the propagation of a planar ice front (Schoof et al., 2001) during the freeze–drying cycle causes the formation of a columnar-like structure converging towards the external surface where it forms horizontal flat ribbons (see insert), resembling the morphology of the lamina splendens. In the lower layer, the higher density and scarce flexibility of the mineralized fibres hamper the formation of a directionally ordered structure during freeze-drying.
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500 μm
500 μm
21.9 ESEM micrograph of the osteochondral scaffold morphology: three different layers are distinguishable for the different density of mineral phase. In detail the columnar morphology of cartilage-like layer.
21.5.1 In vitro and in vivo tests In the preliminary biological assessment, there was an investigation to find out whether the resulting composite materials would differentially support cartilage and bone tissue formation in the various layers, when loaded with articular chondrocytes or bone marrow stromal cells. Expanded human articular chondrocytes was resuspended and statically loaded onto the composite scaffolds, either from the cartilaginous or from the sub-chondral bone layer. Chondrocyte-seeded scaffolds were transferred to agarose-coated dishes and cultured for two weeks in complete medium. Immediately after cell seeding or following a 2-week-long culture, chondrocyte-scaffold constructs were fixed in 4% formalin, embedded in paraffin, cross-sectioned and stained with Safranin-O. After 2 weeks in chondrogenic medium, cells in the cartilaginous layer display a chondrocytic morphology and generate a tissue intensely stained with Safranin O (Fig. 21.10a), whereas cells in the subchondral bone layer remain fibroblastic and do not accumulate histologically detectable amounts of glycosaminoglycans (Fig. 21.10b). The scaffold regions initially void of cells, especially those in the central core, remain essentially necrotic. It is possible to conclude that the cartilaginous layer of the composite scaffold
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100 μm
21.10 Representative safranin O-stained cross-sections of constructs after two weeks of culture in chondrogenic medium of scaffolds seeded from (a) the cartilaginous layer or from (b) the subchondral bone layer. Scale bar = 100 μm.
(a)
(b)
21.11 Histological evaluation performed on decalcified samples (a and b) taken from the implantation site after 3 months: fibro-cartilaginous tissue accompanied by connective tissue is visible.
is permissive to human articular chondrocyte differentiation and cartilaginous matrix deposition, whereas only a fibrous tissue could develop in the subchondral bone layer. Conversely, bone tissue formation is supported within the subchondral layer of the composite, but not in the cartilaginous region. In vivo tests were performed on large animals (horses, n = 2); in the knee articulation, a defect type 4 was created and repaired with the osteochondral substitute. Figure 21.11a shows the defect and the natural formation of tissue after 3 months. Histological evaluation was performed on decalcified samples taken from the implantation site after 3 months: fibro-cartilaginous tissue accompanied by connective tissue is shown in Fig. 21.11b filling completely the gap, whereas bone tissue is evident underneath.
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Future trends
21.6.1 New magnetic scaffold In the development of bio-composites for bone and cartilage repair one of the major concerns is to increase their integration and remodelling rate. One of the driving ideas is to create a conceptually new type of scaffold able to be manipulated in situ by means of magnetic forces. This approach is expected to generate scaffolds with such characteristics as multiple use and possibly multipurpose delivery in order to repair large bone defects, subchondral and osteochondral lesions in the articular surface of the skeletal system. The magnetic moment of the scaffolds gives them the fascinating potential of being continuously controlled and reloaded from an external supervising centre with all needed scaffold materials and various active factors. Such a magnetic scaffold can be imagined as a fixed ‘station’ that offers a long-living assistance to tissue engineering, providing thus a unique opportunity to adjust the scaffold activity to the personal needs of the patient. In addition to the reloading option, magnetic scaffolds will give rise to a few very important options that cannot be adequately addressed by other methods. One such aspects is the possibility of achieving efficient scaffold fixation via magnetic forces. This provides a very elegant and simple solution to the fixation problem. Indeed, scaffold fixation represents a problem owing to difficulties in obtaining a stable interface between the host bone and the scaffold because of their different physical characteristics. As a basic scaffold, a bio-hybrid composite hydroxyapatite/collagen (HA/Coll) can be selected and prepared as described in section 21.4 adding magnetite or other magnetic compatible compounds such as Ti1-x(Fe2+, Fe3+)xO2 and Zr1-x(Fe2+, Fe3+)xO2. The chemico-physical, compositional and magnetic characteristics of the final composites depend on the different procedures tried for performing the magnetic additivation (Wen et al., 2008). TEM analysis of a HA/Coll 62/30% + 8% magnetite composite shows the intimate interaction among the three phases (Fig. 21.12): magnetic nanoparticles (dark) are embedded in the collagen fibres, which are, in turn, mineralized with nanoparticles of apatite. The possibility of allowing preferential interactions among the three phases is under investigation, opening perspectives of graded magnetization design.
21.6.2 Towards the third generation of tissue-engineering solutions Bone and cartilage healing is a complex process involving a number of different factors such as bone morphogenetic proteins (BMPs), signalling/
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50 nm
21.12 TEM analysis of HA/Coll/magnetite composites.
transcription factors, nuclear transcription factors as well as extracellular matrix components. In this respect, a third generation of tissue engineering solutions are represented by the addition of such factors to biomaterials to induce and accelerate tissue in-growth. Although the delivery of BMPs as recombinant proteins can induce local bone formation and healing of bone defects, it was previously demonstrated that local gene transfer of BMP-2 at the site of critical size femoral and cranial defects in the rabbit and rat resulted in more rapid and efficient bone healing. More recently, it was also shown that gene transfer of the LIM mineralization protein (LMP), a novel intracellular positive regulator of the osteoblast differentiation program, can induce efficient bone formation (Lattanzi, 2008). In humans, three different LMP splice variants have been identified, termed LMP-1, LMP-2, and LMP-3. Gene transfer of human LMP-1 and LMP-3 induces expression of genes involved in bone formation including certain bone morphogenetic proteins (BMPs). Direct injection of viral vectors, such as adenovirus, adeno-associated virus or implantation of plasmid DNA in a matrix that express certain BMPs or related osteogenic factors, has resulted in efficient new bone formation in animals models. However, concerns about the
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immune response to the viral vectors and possible dissemination of the vectors have limited the clinical use of this approach. As an alternative to direct in vivo gene transfer, cell-based approaches, using cells genetically modified in culture to express an osteoinductive gene before implantation, have produced efficient new bone formation in animal models. Similarly, gene transfer of intracellular proteins such as Osterix (OSX) and LIM mineralization protein (LMP) also have been shown to induce osteogenesis. It has previously been demonstrated that adenoviral gene transfer of human LMP-3 (Ad.hLMP-3) facilitates ectopic bone formation following direct injection as efficiently as BMP-2 (Pola et al., 2004). However, concerns about the dissemination of the Ad.LMP-3 vector have prevented the clinical application of Ad.LMP-3 for facilitating healing of critical size bone defects. Thus, a clinically relevant approach has been developed to induce bone formation using an ex vivo approach to deliver LMP-3. For this task, genetically modified primary, autologous dermal fibroblasts seeded on HA/Collagen composite before implantation, were used. It was demonstrated that implantation of autologous dermal skin fibroblasts transduced ex vivo with Ad-LMP-3 and adsorbed on HA/Coll nanocomposite (hydroxyapatite/collagen 70/30%) can facilitate ectopic bone formation in mandibular bone defect healing in rats. The efficient induction of osteogenesis by Ad.LMP-expressing fibroblasts clearly shows that LMP-3 is a potent osteoinductive agent. Moreover, given that dermal fibroblasts are easy to isolate and expand, the ex vivo approach using genetically modified dermal fibroblasts could be applied clinically. A mandibular ovalar defect, with a diameter of about 0.5 cm, was created posteriorly to the root of the incisor and filled with the biomaterial alone or with biomaterial plus LMP-3. Radiographic findings 12 weeks after treatments show no bone formation in rats with mandibular defect and treated with biomaterial alone; while the gap is partially filled with local bone formation in the experimental site in the rats treated with LMP-3. Histological analysis showed that the surgical hole was totally healed 4 months after implantation in the HA/Coll scaffold + autologous dermal fibroblasts LMP-3-modified, while the defects were still well evident in all the other treatments.
21.7
Acknowledgements
We thank very much Dr E. Pola and G. Logroscino (Department of Orthopaedics, Università Cattolica del Sacro Cuore, Roma, Italy) for transgenic therapy, collaboration and valuable discussion, European Commission (Projects NMP3-CT-2003-505711, NMP4-CT-2006-033277), the Ministry of Industry, University and Research (Project FIRB RB1P068JL9) and Finceramica SpA (Faenza, Italy) for funding parts of this research work, Dr I. Martin (Departments of Surgery and of Biomedicine, Univer-
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sity Hospital of Basel, Switzerland) and Dr R. Quarto (Advanced Biotechnology Center, Stem Cell Laboratory, University of Genova, Italy) for biological assessment and valuable discussion and Professor M. Marcacci (Rizzoli Orthopaedic Institute, Bologna) for pre-clinical experiment.
21.8
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22 Developing targeted biocomposites in tissue engineering and regenerative medicine J. A. P L A N E L L and M. N AVA R R O, Institute for Bioengineering of Catalonia (IBEC), Spain
Abstract: Regenerative medicine is a relatively new field with new requirements for smart materials, where composites will have a strong role to play. The new paradigm of regenerative medicine and tissue engineering requires biomaterials with high specificity, where physical and chemical properties are duly tailored and combined with appropriate mechanical and degradation features in order to trigger specific cell events and functions involved in the regenerative process. In this chapter, the chemical, physical and biological elements that have to be targeted by biocomposites in regenerative medicine are described. Key words: biocomposite, regenerative medicine, tissue engineering, scaffolds, cell/material interactions.
22.1
Introduction
Composite materials have been widely used in various medical applications and devices. They are here to stay in the medical devices area and specifically in the implants field. However, regenerative medicine is a relatively new field with new requirements for smart materials, where composites will have a strong role to play. Tissue engineering and regenerative medicine explore the repair and regeneration of organs and tissues using the natural signaling pathways and components such as stem cells, growth factors and peptide sequences among others, in combination with synthetic scaffolds (Hardouin et al. 2000). In addition, after the implantation of an appropriate combination of the basic tissue engineering triad (cells, signaling and scaffold), there are some processes such as angiogenesis and nutrients delivery that are crucial to stimulate tissue regeneration and that must take place immediately after implantation. Probably the key issue when trying to engineer the regeneration process is first to understand the dialogue between the cells (preferably undifferentiated) that will eventually biosynthesize the regenerated extracellular matrix, with its microenvironment. The rationale is that environmental biochemical cues as well as other biophysical signals are read by cells as messages that they translate into intracellular commands that modify their behavior in 573
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response to such signals and which make them react with appropriate responses. Cells write their own messages into the extracellular space either by biosynthesizing new extracellular matrix, and/or remodeling the existing one. This bidirectional communication of the cell with its extracellular matrix and its microenvironment is in fact the language that should be mastered in order to control, guide and engineer the regenerative process. When analyzing the different elements that participate in the regenerative process, it is necessary at least to consider the following ones: a) interstitial transport that will be responsible for bringing oxygen, electrolytes, nutrients and molecules to the cells playing the regenerative role, b) cell mobility and transport that has to explain how cells will move to the regeneration site, c) the biomaterial or the scaffold that has to be designed to be biodegradable and to have the appropriate stiffness to favor the regenerative process, d) the metabolic environment that will be responsible for the degradation of the implanted products in the regeneration site, e) the cell– material interactions, where the surface properties of the material have a direct and strong effect on cell behavior, f) the biochemical and biophysical stimuli from the microenvironment, such as chemical cues that signal cells for a regenerative response, or mechanical stimuli that can also regulate cell behavior, g) angiogenesis control that is responsible for the in vivo regeneration of the tissue, as in certain tissues like bone, nervous or cardiac muscle, vessel formation is necessary, while for others such as cartilage, ligament or cornea, vessel formation should be avoided, h) the immunological and inflammatory response that is responsible for the healing and eventual regeneration of the tissue, i) the cell behavior that includes attachment, proliferation and differentiation and that will control the regenerative process. All these factors play a decisive role of their own and by their mutual interaction in the regeneration process, and, ideally, in order to engineer the regenerative process, it would be convenient to control and to guide them. Because such a complex combination is required, the type of biomaterials to be considered should be tailored appropriately. Within this complex context, we will now concentrate on the materials issue. Historically, biomaterials have evolved from the concept of bioinert materials, where biomaterials were expected not to interact, or in other words to be as neutral as possible with the biological environment, to the present concept of smart biomaterials, where they are expected to elicit favorable and desired biological responses. Usually, the concepts of first, second and third generation are used to describe their evolution up to the present day. The biomaterials philosophy has evolved from the one existing in the 1960s and 1970s, in the first generation biomaterials, where nothing biologically really positive could be expected from the materials side, to the present one in the third generation of biomaterials, where it is expected that biomaterials can be tailored and made smart enough in order that they
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can elicit the positive and desired response into and from the biological environment. At this point, it is convenient to introduce an important new idea; historically, in materials science, the concept has been that a material does not constitute a device and that a device may be made from different materials. This is the case for implants, devices and biomaterials. Consequently, the paradigm has been that biomaterials could be selected and used in completely different applications, i.e. in implants meant to interface with different tissues, like stainless steel that is being used in orthopedic implants as well as in cardiovascular stents. Our point of view is that the new paradigm of regenerative medicine will not allow this crude approach and biomaterials will have to be duly tailored for the desired clinical application and for the tissue to be regenerated. Nowadays, we already see that specific surface modifications are performed depending on the use of the biomaterial. This specificity and need to tailor the material by combining physical and chemical properties, with appropriate mechanical and degradation features, make composites probably the best type of material as the required properties can be engineered by combining the appropriate components. This is the reason why it is believed that composite materials are intended in the near future to be the materials specifically designed for scaffolds to be used in very specific regenerative applications. To date, a wide range of materials has been developed for fixation, repair and regeneration of tissues; however, as previously stated, depending on the applications and the tissue requirements, there are some cases where it is difficult to achieve all targeted properties from a single material. Consequently, great opportunities exist in the development of hybrid materials, which include more than one type of material, often with the incorporation of a synthetic polymer component. The necessity for unique material property combinations with respect to both biological and physiochemical functionality often demands material solutions that rely on the generation of multicomponent hybrid, composite, and otherwise complex biomaterials. For example, most natural ‘biomaterials’ exhibit a composite microstructure, in the sense that they contain two or more chemically and structurally defined components, which play specific roles and combine to generate the physical and biological properties characteristic of the particular matrix or tissue. Thus, the creative use and combination of natural and synthetic biopolymers, ceramics, and inorganic materials offers a convenient bridge between chemical and biosynthetic approaches. For biocomposites, an important requirement is the synergy between adequate mechanical properties and the biological compatibility of both components. In general, the ideal materials for tissue engineering application must fulfill the following requirements (Boccaccini et al., 2002; Spaans et al., 2000):
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• The material has to be biocompatible and its degradation by-products non-cytotoxic. • The scaffold should be biodegradable and should resorb at the same rate as the tissue is repaired. • The scaffold should have a highly interconnected porous network, formed by a combination of macro- and micropores that enable proper cellular and tissue in-growth, vascularization and nutrient delivery. • The mechanical properties of the scaffold should be appropriate to fill the cavity and sustain the loads applied in the implantation site in order to eventually regenerate the required tissue. Moreover, the material needs to keep its structural integrity during the initial stages of the new tissue formation. In addition to these standard requirements, as mentioned before, it would be desirable that the developed materials could also control and trigger specific events and cell functions involved in the regenerative process. In order to achieve these objectives, physical and chemical material properties will play the leading role and, consequently, the development of smart composite materials aimed to display both chemical and physical signals able to activate specific cellular events is envisioned.
22.2
Cell/material interactions
Once a material is implanted, a bidirectional complex combination of events takes place between the material and the biological environment. The success of a biomaterial strongly depends on its interactions with the physiological fluids and cells. The first nanoseconds of contact between the material and the biological environment are crucial and will affect further stages of the material/cell interaction process. The initial contact takes place between the material surface and water molecules which form a water coating layer. This first stage is highly dependent on the surface properties of the material and will determine the type of biomolecules that will interact with the surface in subsequent stages. The next stage occurs from some seconds up to hours after implantation. This stage consists of the interaction of all the macromolecules existing in the physiological medium, namely sugars, lipids and proteins, with the material surface. The ‘Vroman effect’ takes place during this stage and the material surface is covered by an adsorbed layer of proteins (McFarland et al., 1999). Finally, a third stage follows after time periods ranging from minutes up to days after implantation. During this stage, cells get in contact with the surface and interact with it. This third stage is characterized by multiple complex interactions between the extracellular matrix proteins, cell membrane proteins and cytoskeleton proteins,
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surface chemistry and topography, the micro- and macrostructure of the material (porosity, pore size and geometry, interconnectivity) and the released degradation by-products of the material. Overall, biocompatibility and material interactions with biological entities are dependent on the proteins adsorbed to the surface during these first stages of contact. Thus, the protein adsorption process is of paramount importance. Moreover, this complex process is highly influenced by the materials’ surface properties; indeed, it is known that surface features such as its chemical composition, and surface energy determine the nature of the proteins adsorbed to the surface and their orientation and conformation. Both protein conformation and orientation are two important parameters in cell attachment and adhesion. In fact, the availability of protein peptide sequences for cell attachment on the surface varies depending on the orientation and conformation of the adsorbed proteins. Only those peptide sequences exposed to the cell-material interface are accessible to cell membrane receptors like integrins, while those located in the interior of the protein are not. Material/cell interactions take place through cell adhesion proteins known as integrins that interact with particular peptide motifs from the proteins adsorbed on the material surface. Integrins are cell transmembrane proteins that possess two glycoproteic units (α and β) and three domains (cytoplasmic, transmembrane and the extracellular one). The extracellular domains of the α and β units possess receptors for the specific recognition of cell adhesive peptide motifs that are contained in some adhesive proteins present in the extracellular matrix (ECM) (Siebers et al., 2005). Thus, cell attachment and proliferation are integrins-mediated processes. The cell adhesion process is very relevant given that it triggers some mechanical and chemical signals that affect further cell events such as proliferation and differentiation that in turn determine cell functionality (Anselme et al., 2000). Although integrins are the first contact point between cells and the material surface; cells interaction with the external medium and specifically, with the material surface is carried out through their cytoplasm, specifically, through cell structures known as lamellipodia or pseudopodia depending on the cell type, which are cell extensions formed by actin filaments (Anselme 2000). Fillopodia are the actuators of the adhesion, morphology, spreading and motility processes. Integrins located within these long and thin cytoplasm extensions interact with the substrate surface creating focal contacts that are points where several integrin receptors meet to form stronger adhesion points. Depending on the surface conditions, the fillopodia will receive more or fewer signals allowing the cell to attach or separate from the surface, to move in one direction or other, to get a very spread or rounded morphology, etc. (Beningo et al., 2001; Margel et al., 1993).
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Thus, in general, living cells have great abilities to sense, respond to, and even remodel their surrounding environment. All these processes take place through cell membrane receptor proteins, which are the most important channels for transduction of external information into the cell. Biocomposites aimed at different regenerative medicine applications should be able to activate specific processes by the generation of signals that trigger different cascades through these surface receptors. In this sense, the composite material should be able to deliver biochemical and/or biophysical signals which could be in fact constituents of the composite. In addition to surface properties there are other physical and chemical features related to the intrinsic properties of the materials such as the material mechanical properties, degradation rate and by-products release among others, and other properties resulting from the elaboration techniques such as the materials morphology and architecture that also affect to a high degree the tissue regeneration process.
22.3
Physical aspects of materials in tissue engineering
22.3.1 Mechanical properties One criterion that a material must fulfill to act as temporary scaffold for tissue engineering purposes is the one related to its mechanical properties. The mechanical properties of the scaffold must be appropriate to regenerate the tissue. Moreover, the material must keep its structural integrity during the first stages of the new tissue formation and undertake a progressive and gradual degradation. Mechanical properties are of paramount importance in scaffolds designed for the regeneration of load-bearing tissues such as bone, cartilage, or blood vessel, and also in the case of non-load bearing applications such as skin grafts and wound dressings. Depending on the application, the requirements for each type of material vary. The idea is to fulfill the mechanical needs of every tissue and to create a material that provides support and that matches the mechanical properties of the tissue to be regenerated. As stated in previous chapters, matching the mechanical properties of the material with those of biological tissues is of great importance for most tissues. Such is the case of bone, where materials with appropriate elastic modulus are crucial in order to avoid the stress shielding effect. In this sense, composite materials exhibit clear advantages over single phase materials; the combination of several materials allows the mechanical properties to be tuned according to the specific requirements of every case. For instance, biological materials such as bones and teeth are characterized by a layered structure consisting of strong inorganic platelets embedded
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in a soft, ductile organic matrix (Lowenstam and Weiner, 1989). In spite of the inherently weak inorganic constituents (e.g., silica, calcium carbonates, and phosphates), the high strength of the inorganic building blocks is ensured by limiting at least one of their dimensions to the nanoscale (Gao et al., 2003). These tiny building blocks are usually organized into a hierarchical structure spanning over various length scales. Changes in the fraction of inorganic phase (i.e., degree of mineralization) lead to biocomposites ranging from soft tissues such as calcified tendons to strong, hard structures such as bone. Other approaches involve the elaboration of biocomposites by means of interpenetrating networks of protein elastomers with synthetic hydrophilic polymers or with natural polymers in order to design mechanically improved replacements for the ECM (Bonderer et al., 2008).
22.3.2 Surface topography Numerous works have shown that cell morphology correlates closely with function. Nowadays, it is possible to control cell shape and therefore the expression of differentiated cell phenotypes by creating specific patterns of surface chemistry and/or topography. Several studies have demonstrated that biomaterial surface roughness and patterning can be used to control cellular activity (Chesnel et al., 1995; Healy et al., 1996). Furthermore, it has been accepted that cell attachment, adhesion and proliferation on a surface can be guided by microtopography. In fact, numerous studies have confirmed that topography influences cell adhesion (Anselme et al., 2000; Huang et al., 2004), morphology (Chen et al., 1997), migration (Tan and Saltzman 2002), orientation, focal adhesion (Diener et al., 2005) and differentiation (Zinger et al., 2005). In order to influence the different structures involved in cellular events, several length scales have to be used. The use of micro- or nanotopographic features depends on the size of the biological structures to be activated or affected by the topographic changes. For instance, proteins require features in the nanoscale that are closer to their dimension scale and better adjusted to their size; whereas, cells, which are larger entities, can be stimulated using microscale topographic structures. Indeed, some studies have reported clear cell stimulation when using features ranging between 10 and 100 μm. Biocomposites, where one of the phases consists of inorganic particles, present a certain type of surface texture owing to the presence of these particles that may be fully imbedded in the matrix or partially exposed in the material surface generating specific surface topographies depending on the geometry, dimensions and percentage of the inorganic phase. Although both micro- and nanotopography appear to have an important effect of protein adsorption and subsequent cell events, the complete picture and understanding of the relationship between micro- and
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nanotopography and cell response is not completely clear. Moreover, it must be highlighted that not all cell phenotypes react in the same manner to topographical changes; it seems that topographical stimulation is highly cell-dependent. Most of these principles have been demonstrated in 2D systems. Their extrapolation to 3D patterned or textured structures represents a major challenge and will probably require the extension of fabrication efforts derived from the fields of computer-aided design/computer-aided manufacturing and rapid prototyping fabrication processes. Even more challenging is the patterning of biocomposites where materials with different properties are combined.
22.3.3 Architecture The architecture of biocomposites constructs for tissue regeneration plays a very important role in cell in-growth, migration, maintenance and further structure of the new tissue. Pore distribution, interconnectivity and size distribution are key issues. Scaffolds present a highly interconnected porous network, formed by a combination of macro, micro and even nanopores that enable proper tissue vascularization, nutrient diffusion and waste release. The specific pore size depends on the type of tissue to be regenerated. In the case of bone, a pore size between 100 and 350 mm has shown to be optimum one (Klawitter and Hulbert, 1971). The micro- and macrostructure of the scaffolds depend significantly on the processing technique. Numerous elaboration techniques and approaches to the development of three-dimensional scaffolds that combine biodegradability and bioactivity are currently under study (Coombes and Heckman 1992; Mikos et al., 1993; Mooney et al., 1996; Wake et al., 1996). Each manufacturing technique provides particular and different structural characteristics to the scaffold. Hence, the choice of the technique depends on the requirements of the final application. In the case of composite materials, an extra difficulty is added to the elaboration techniques owing to the incorporation or combination of two or more different materials. Thus, the different phases have to be properly combined and processes have to be adapted to suit each of them. Some of the most used techniques for the processing of such scaffolds are gel casting, solvent casting and particulate leaching, laminated object manufacturing, phase separation, gas saturation, fibre bonding, electrospinning, rapid prototyping and membrane lamination among others (Yang et al., 2001). Among them electrospinning and rapid prototyping appear to be very promising (Agarwal et al., 2008; Pfister et al., 2004). Moreover, a polymer scaffold should be easily and reproducibly processed into a desired shape and structure, which can be maintained after
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implantation, thus defining the ultimate shape of the regenerated tissue. Therefore, developing new chemical and physical techniques for the attainment of well-defined biomimetic and biocompatible 2D and 3D structures within synthetic and natural materials remains as a significant challenge in the field.
22.4
Chemical aspects of materials in tissue engineering
22.4.1 Material degradation Many strategies in regenerative medicine have focused on the use of biodegradable polymers as temporary scaffolds for cell transplantation or tissue formation. In fact, one of the most important criterions a 3D scaffold must fulfill is a degradation rate that matches the tissue regeneration process in order to maintain the mechanical strength and to avoid collapse or stress shielding. Besides, the material has to be degradable over an appropriate timescale into products that can be metabolized or excreted. Degradation rate is highly dependent on the intrinsic properties of the material. Polymers biodegradability is mainly originated by hydrolysis of the polymer chain backbone and to a lesser extent by enzymatic activity (Li and McCarthy 1999; Vert and Li 1992). Degradation times depend on multiple parameters such as polymer crystallinity, molecular weight, thermal history, porosity, monomer concentration, geometry and the location of the implant. In an aqueous environment, water penetrates the bulk of the polymer sample and preferentially attacks the chemical bonds of the amorphous phase, shortening the polymer chains (Gopferich 1996). Crystalline regions remain and temporarily support the physical properties of the device until water attacks the crystalline regions. In a second stage, enzymatic attack of the bulk takes place (Li and Vert 1994). Degradation in the internal part of the device is accelerated owing to the presence of the acidic degradation by-products that autocatalyze the degradation process of the material. Thus, during degradation the amount of reactive hydrolitic and acidic end groups increase while the material crystallinity and size decrease; these factors accelerate the degradation process (Grizzi et al. 1995; Middleton and Tipton 2000). This type of degradation is known as ‘bulk degradation’ (Li and Vert 1994; Li 1999). Several studies have shown that the incorporation of an inorganic phase into a bioabsorbable polymer matrix not only modifies the mechanical behavior of the porous structure (Navarro et al. 2004) and enhances the bioactivity of the tissue engineering scaffolds, but also modifies the degradation pattern of the polymer (Spaans et al. 2000; Navarro et al. 2005). On
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the one hand, the addition of an inorganic component may increase the wettability of the material and also generate interstices at the interface of both materials allowing in this way, the penetration of fluid to the materials’ bulk, which, in turn, may accelerate the degradation process. On the other hand, the incorporation of inorganic particles may have a buffering effect and control the pH and type of compounds released to the medium. Thus, the right combination of organic and inorganic phases in biocomposites may be also used to tune the biodegradability which depends on each specific application, and the by-products from the degradation process.
22.4.2 Surface chemistry Surface chemistry together with surface topography is of paramount importance in cell/material interactions. Functional groups exposed at the material surface are responsible for surface properties such as wettability, surface electrical charges and free energy that, in turn, influence protein adsorption and cell behaviour. Thus, by tailoring the functional groups available at the material surface, it is possible to modify and enhance protein/surface interactions. Nevertheless, at present, there is no methodology that allows a full control of the protein conformation and orientation after adsorption (Roach 2007). Some natural polymers such as hyaluronic acid, collagen and chitosan have proved to have some intrinsic bioactive effect in some tissues such as cartilage tissue. In the case of polymers, bioactivity depends on the functional groups and binding sites available at the material surface. Surface modification methods to improve the interactions between the material surface and cells have evolved during recent decades. Surface bioactivation can be achieved by functionalizing surfaces with different biomolecules by means of a variety of methods where both chemical bonding and physical adsorption take place. Dip-coating techniques, the formation of self-assembled monolayers (SAMs) and binding polymer chains or brushes to the surface are some of the methods used to enhance cell adhesion, to influence proliferation and differentiation rates, and to achieve faster and more stable integration between the material and the tissue as in the case of dental implants and some orthopaedic prostheses (Blawas et al. 1998; Scotchford et al., 1998). More sophisticated ‘bottom-up’ and ‘top-down’ techniques have been developed to engineer surfaces with high levels of specificity as well as the synthesis and tailoring of new biomolecules for specific applications. The development of more complex biopolymers and biomolecules such as elastin-like biopolymers including peptide sequences that induce mineralization and cell adhesion, or self-assembled amphiphilic peptides that include cell signalling cues are new approaches to mimic the natural process
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by which collagen induces the assembling of calcium phosphate, and hydroxyapatite crystallites within bone in order to generate its mineral rigid phase (Rodríguez-Cabello et al., 2007; Sargeant et al., 2008). The combination of these new biomolecules and biopolymers with a CaP inorganic bioactive phase lead to the attainment of new biocomposites with enhanced bioactivity and mineralization potential. As with surface topography, the majority of the studies proving the effect of surface chemistry, optimal distribution and concentration of the chemical cues on the surface, and other parameters, have been performed in 2D surfaces rather than in 3D scaffolds. Thus, again, the extrapolation of functionalization methods to such porous structures for regenerative medicine implies numerous studies and the development of new techniques or adjustment of previous ones to functionalize 3D porous structures and analyse their effect on cells activity. Biocomposites present an additional difficulty attributed to the fact that each of the phases forming the composite material may have different surface chemistry and functional groups, thus, in these cases the presence of different surface types must be carefully considered when selecting functionalization techniques. Thus, chemical modification of 3D biocomposites is a complex task, which involves the hierarchical assembly of molecular signaling ligands into materials with appropriate physical properties.
22.4.3 Release of biomolecules In addition to insoluble signals mimicking ECM molecules or domains, cells can also respond to soluble bioactive molecules such as cytokines, growth factors, and angiogenic factors. Although these molecules alone can be used for tissue induction, it is possible that combining cell therapies with drug and/or growth factor delivery may accelerate the tissue regeneration process. Materials science provides the possibility of linking drug delivery to reparative medicine. Antigen–antibody hydrogels, temperature- and pH-sensitive gels for drug release, polymer-protein conjugates with temperature-switchable ligand binding or mechanical properties, and celladhesive delivery systems are among many examples (Sakiyama et al., 1999). Moreover, tissue–inductive factors can be incorporated into biodegradable polymers during scaffold processing (Babensee and Mikos, 2000; Bancroft et al., 2001). A different approach involves the use of biodegradable microparticles or nanoparticles loaded with biomolecules that can be embedded into the substrates to create biocomposite materials able to release inductive chemical signals. The release of bioactive molecules in vivo is then modulated by both diffusion and scaffold degradation rate.
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Maintenance of bioactivity has been a major concern in delivering large molecules such as proteins, and improvements in this area are clearly needed. Current challenges include the design of tissue-targeted polymers with control over subcellular transport in order to optimize drug effects.
22.5
Specific processes in regenerative medicine
22.5.1 Inflammatory and immunological responses Inflammation is an inevitable consequence of surgery and is a natural pathway to remodeling neo-tissue at the site of injury. It has been defined as a ‘vital response’ of tissue injury (Houck 1963). Taking into account this definition, inflammation could be considered as a protective and normal response to any kind of toxic stimulus. This stimulus may alter the normal physiological process of the host, varying from the acute transient and highly localized response to simple mechanical injury or to the complex persistent response involving the whole organism. This initial response may initiate further a series of biochemical, immunological and cellular events, which may range in time from recognition of the toxic stimulus through mobilization of natural defense mechanisms ending with physical repair and restoration of function of the injured tissue. Thus, inflammation is a complex combination of vascular, lymphatic and local tissue reactions elicited by the presence of viable and non-viable irritants (Naik and Sheth, 1976). Although the materials that are used clinically are non-immunogenic, non-toxic, and chemically inert, they trigger acute, potentially chronic, inflammatory responses. The inflammatory response is induced by the materials in an indirect way, through proteins that adsorb to them and is highly dependent on surface properties (Wilson et al., 2005). Indeed, the surface chemistry and hydrophobicity strongly influence the composition of this adsorbed layer (Tang and Eaton, 1993; Thull 2002). In addition, the surface charge of the implanted material influences PMN (polymorphonuclear neutrophils) and macrophage adhesion (Hunt et al., 1996; DeFife et al., 1999). Consequently, the foreign body reaction (FBR) varies depending on the type of surface chemistry that is used, and the intrinsic material characteristics affect the path of the FBR. Thus, the activation of the FBR against implanted materials partly depends on the physicochemical characteristics of the material. The adsorption of fibrin, proteins from the complement system, and antibodies further activates the inflammatory process that is initiated by tissue damage (Luttikhuizen et al., 2006). One of the research goals within the biomaterials field is to create new devices that can prevent foreign body reactions and promote normal healing.
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The foreign body reaction is a serious limitation especially in those cases where materials have to be in direct contact and integrated with the surrounding tissues; it can also lead to chronic pain and eventual rejection of the device and failure. The result of all the previous considerations is that the implant or medical device surface plays a crucial role in its interaction with the biological environment. Consequently, the study and characterization, modification and functionalization of biomaterials surfaces are probably the main strategies for success of implants and the tissue regeneration. Immunity refers to the ability of an organism to resist disease by identifying and destroying foreign substances or organisms (Fabre et al., 2001; Janeway 1996). Cells and molecules involved in such mechanisms constitute the immune system and the response resulting from the introduction of a foreign agent is known as the immune response. The implantation of any medical device can be considered an external invading element that might induce an immune response, mostly but not exclusively dependent on the properties of the device (Anderson 1993; Parker et al., 2002; Zhang et al. 1996). There is some controversy about which properties of the surface of the materials stimulate particular cell/tissue reactions. It has been hypothesized that not only wettability and surface charge of the surface of the materials, but also the presence of certain functional groups have importance for the adhesion and activation of immunological cells in vitro. Furthermore, the degradation rate and mechanisms of biodegradable devices can also modulate and might allow control of tissue responses in vivo. For degradable materials, the FBR, in general, becomes chronic, until final degradation. For non-degradable materials, on the other hand, the reaction continues until a capsule is formed around the implant, shielding it from the nonspecific immune system. Therefore, the development of new biocomposites with controlled degradation and surface chemistry is of paramount importance to evade the adsorption of proteins related to negative inflammatory and immune responses while enhancing the successful integration of the implant. In fact, the objective should not be to suppress the inflammatory response, but to control it. Since an inflammatory process always exists whenever there is wound that requires healing, because inflammation is the first step towards healing, smart materials should be able to control and guide the process for the best benefit in the regenerative process. Again here, composite materials offer the possibility to contain adequate stimuli for the guiding of the inflammatory process.
22.5.2 Mechanotransduction Mechanotransduction is the process by which mechanical energy is converted into electrical and/or biochemical signals (Burger and Klein-
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Nulend 1999). At the cellular scale, normal tissue cells experience different types of forces as they anchor and pull on their surroundings. The transmission of forces is a bidirectional process that takes place between cells and their surroundings through cell membrane adhesion molecules such as integrins and cadherins. A normal tissue cell not only applies forces but also responds through the cytoskeleton organization (and other cellular processes) to the resistance that the cell senses, regardless of whether the resistance comes from normal tissue matrix, synthetic substrate, or even an adjacent cell. Cells adjust their cytoskeleton, focal adhesive points and overall state depending on the resistance of the substrate. Epithelial cells and fibroblasts were the first cells reported that responded in different ways to soft versus stiff substrates. This study was performed on gels with different stiffness and coated with specific ligands (Pelham and Wang, 1997). Although molecular pathways are still only partially known, muscle cells, neurons, and many other tissue cells have since been shown to sense substrate stiffness (Deroanne et al., 2001; Engler 2004; Wang et al., 2000). In general, it has been shown that there is a trend to more organized cytoskeletons and larger, more stable adhesions with increasing elastic modulus of the material (Engler et al., 2004; Georges and Janmey 2005; Raeber et al., 2005). However, the responses appear to be specific to anchorage-dependent and/or relatively contractile cells. (Discher et al., 2005). One of the tissues that better illustrates the mechanotransduction process is bone. Bone is a biocomposite that possesses a hard extracellular matrix composed of collagen fibers and calcium phosphate mineral as structural elements (Carter and Hayes, 1977). In trabecular as well as in compact bone the three-dimensional organization of its structure depends on the direction of the principal mechanical stresses during daily loading and movement (Fiala and Hert 1993; Hert 1994; Petrtyl et al., 1996). The loading information is communicated to effector cells that can make new bone or destroy old bone. Thus, mechanical stimulation activates osteocytes to produce signals such as PGE2 and IGF that recruit new osteoblasts from periosteum and bone marrow stroma, or osteoclasts according to the specific case (Feyen et al., 1985). Thus, as previously discussed, mechanical properties play a remarkable role in tissue regeneration, not only because of the temporary support function they have but also because of their effect at the cell scale owing to the mechanotransduction phenomenon. The mechanical effect of biocomposites on cells may be tailored by combining their components in such a way that they are able to induce specific cell cytoskeletal changes that stimulate the production of certain chemokines and factors responsible of particular cell events.
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22.5.3 Angiogenesis Angiogenesis, the formation process of new blood vessels, is critical to tissue functionality for the delivery of nutrients and oxygen among other functions, and depends on the tightly coordinated interplay between a specific peptide sequence known as VEGF and other signaling molecules. VEGF induces endothelial cells which are the ones lining the lumen of blood vessels to proliferate, migrate and sprout neovessels, and provides a survival signal until these immature vessels become stabilized by mural cells as pericytes and smooth muscle cells, and ECM deposition (Carmeliet 2000; Blau and Banfi 2001). Simultaneous interactions between different molecules such as angiopoietin-2 and basic fibroblast growth factor (BFGF) contribute to the induction of this process, whereas sequential collaboration with platelet derived growth factor (PDGF) and angiopoietin-1 (Ang-1) mediate recruitment of mural cells and stabilization of the endothelium, respectively (Hirschi et al., 2002). Furthermore, collaboration with anti-angiogenic molecules (e.g., thrombospondin) may play an important role in normal angiogenesis, and contribute to the functionality of the newly developed vasculature (Sund et al., 2005). Delivery from polymeric systems that mimic the ECM binding characteristics of the biological VEGF delivery system may mimic microenvironmental aspects crucial to the efficacy of pro- and anti-angiogenic therapies (Ozawa et al., 2004). In this sense, composite systems consisting of multiple polymer phases with distinct release kinetics may be designed for applications in which the sequential delivery of factors is desired. The sequential collaboration of VEGF and PDGF inherent to angiogenesis may, for instance, be mimicked by dual growth factor release from biocomposite porous scaffolds (Richardson et al., 2001). Localization of VEGF, predominantly in a compartment close to the surface of the scaffold pores leads to rapid release, whereas localization of PDGF in a second compartment more deeply embedded within the matrix leads to a delayed release dependent on polymer degradation rate (Ennett et al., 2006). As contrasted to individual delivery of either factor, this approach leads to the formation of mature vascular networks, as characterized by larger vessels associated with mural cells (Richardson et al., 2001). Sequential release vehicles may also be designed by embedding particulate systems into hydrogel matrices. For example, incorporation of gelatin microspheres encapsulating a first factor into oligo[poly(ethylene glycol) fumarate] hydrogels containing a second factor may lead to sequential supply of the respective molecules (Holland et al., 2005). The release profile of these composite systems may be tuned by varying the location of each factor and the extent of crosslinking of each polymer.
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Some examples include the elaboration of composite materials where VEGF was incorporated into three-dimensional porous scaffolds made of poly(lactide-co-glycolide) copolymer for localized protein delivery, then coated the surface with a bioactive glass to enhance the osteoconductivity, and potentially the osteoinductivity of the construct (Leach et al., 2006), or the combination of VEGF with PDGF, where VEGF is incorporated in the material while PDGF is incorporated in microspheres (Peters et al., 2002). Other approaches involved the elaboration of silk scaffolds by embedding parallel silk fibers within a collagen-HA solution in order to enhance bone formation and new vessels formation (Seo et al., 2008). In general, the main goal of these systems is to provide therapeutically efficient concentrations of VEGF or its inhibitors in a biologically inspired manner that controls their temporal availability, and interactions with other signaling molecules.
22.6
Conclusions
When looking into composites for application areas, it is necessary to say that practically all the areas of substitution, repair and regeneration are suitable for their application. While composites have been used in certain specific applications as conventional implants, either biodegradable or nonbiodegradable, it seems that they have an even brighter role to play with the development of regenerative medicine. Regenerative medicine will require smart biomaterials able to send signals preferably to undifferentiated cells (stem cells) that, as a consequence, will modify their behavior and will respond by regenerating the new tissue. We do not have yet the appropriate materials for every given regenerative process. However, it is possible to foresee that composite materials may play a key and leading role, since composite materials can be tailored in terms of composition and properties. As an example, it is possible to think of the design of a composite that combines adequate mechanical properties, with the existence of specific cellular attachment sequences on its surface and also with a proper release rate of a certain growth factor. Only composite materials allow thinking in terms of such a flexibility of design and this is the reason why it should be expected that a wide variety of composites will be available in the future for the regeneration of tissues and organs.
22.7
References
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lowenstam ha, weiner s (1989) On biomineralization (Oxford University Press, New York, 1989), pp. 336. luttikhuizen dt, harmsen mc, van luyn mja (2006) ‘Cellular and molecular dynamics in the foreign body reaction’ Tisse Eng 12:1955–1970. margel s, vogler ea, firment l, watt t, haynie s, sogah dy (1993) ‘Peptide, protein and cellular interactions with self-assembled monolayer model surfaces’ J Biomed Mater Res 27:1463–1476. mcfarland cd, mayer s, scotchford c, dalton ba, steele jg, downes s (1999) ‘Attachment of cultured human bone cells to novel polymers’ J Biomed Mater Res 44(1):1–11. middleton jc, tipton aj (2000) ‘Synthetic biodegradable polymers as orthopedic devices’ Biomaterials 21:2335–2346. mikos ag, sarakinos g, leite sm, vacant jp, langer r (1993) ‘Laminated threedimensional biodegradable foams for use in tissue engineering’ Biomaterials 14:323–330. mooney dj, baldwin df, suh np, vacanti jp, langer, r (1996) ‘Novel approach to fabricate porous sponges of poly(d,l-lactic-co-glycolic acid) without the use of organic solvents’ Biomaterials 17:1417–1422. naik sr, sheth uk (1976) ‘Inflammatory process and screening methods for antiinflammatory agents – A review’ J Postgrad Med 22:5–21. navarro m, ginebra mp, planell ja, barrias cc, barbosa ma (2005) ‘In vitro degradation behavior of a novel bioresorbable composite material based on PLA and a soluble CaP glass’ Acta Biomater 1:411–419. navarro m, ginebra mp, planell ja, zeppetelli s, ambrosio l (2004) ‘Development and cell response of a new biodegradable composite scaffold for bone guided regeneration’ J Mater Sci Mater Med 15:419–422. ozawa cr, banfi a, glazer nl, thurston g, springer ml, kraft pe, et al. (2004) ‘Microenvironmental VEGF concentration, not total dose, determines a threshold between normal and aberrant angiogenesis’ J Clin Invest 113(4):516–527. parker ja, walboomers xf, von den hoff jw, maltha jc, jansen ja (2002) ‘Softtissue response to silicone and poly-l-lactic acid implants with a periodic or random surface micropattern’ J Biomed Mater Res 61:91–98. peters mc, polverini pj, mooney dj (2002) ‘Engineering vascular networks in porous polymer matrices’ J Biomed Mater Res 60(4):668–678. petrtyl m, hert j, fiala, p (1996) ‘Spatial organization of the Haversian bone in man’ J Biomech 29:161–169. pelham rj, wang y (1997) ‘Cell locomotion and focal adhesions are regulated by substrate flexibility’ Proc Natl Acad Sci USA 94:13661–13665. pfister a, landers r, laib a, hübner u, schmelzeisen r, mühaupt r (2004) ‘Biofunctional rapid prototyping for tissue engineering applications: 3D bioplotting versus 3D printing’ J Polym Sci: Part A: Polym Chem 42:624–638. raeber gp, lutolf mp, hubbell ja (2005) ‘Molecularly engineered PEG hydrogels: a novel model system for proteolytically mediated cell migration’ Biophys J 89:1374–1388. richardson tp, peters mc, ennett ab, mooney dj (2001) ‘Polymeric system for dual growth factor delivery’ Nat Biotechnol 19(11):1029–1034. roach p, eglin d, rohde k, perry cc (2007) ‘Modern biomaterials: a review – bulk properties and implications of surface modifications’ J Mater Sci: Mater Med 18:1263–1277.
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23 Ethical issues affecting the use of biocomposites L. T R O M M E L M A N S and K. D I E R I C K X, Centre for Biomedical Ethics and Law, Belgium
Abstract: The development and application of biomedical composites for medical uses necessitates reconsideration of how pertinent ethical issues should be addressed. These innovations create the prospect of better biocompatibility, making the devices that contain them more appropriate to their task. They are also more complex and their interaction with the human body is more intricate and longer lasting. These features need to be considered when deciding on the ethically responsible way of developing and applying these products. In this chapter we focus on the following elements: (1) the identification of risks and weighing of risks and benefits, (2) the protection of the recipients of these biocomposites over the long-term, (3) the possibility of using these products for enhancement purposes and (4) the impact of these products on the health-care system and the just allocation of these products. Key words: biocomposites, ethics, risk–benefit analysis, nanotechnology, safety, health care allocation.
23.1
Introduction
The notion of good science and good medicine has an ethical undertone: the pursuit of human flourishing through science, engineering and medicine is an unquestionable good and should be encouraged. Most people would also claim that this end does not justify any means. Science and medicine conducted at the expense of human dignity or causing needless suffering is considered to be bad science and bad medicine and is therefore widely and rightly condemned. The scientific and medical community is increasingly aware that fundamental ethical principles should govern research and medicine. Nearly every person engages in ethical discourse, but the ethical analysis of complex biomedical innovations and the development of appropriate ethical guidance are usually considered to be the province of (bio)ethicists. While ethicists are addressing these issues, it is unrealistic to expect that they will provide simple answers to complex questions, or that they will present the community with clear and prohibitive rules for shaping human conduct. 593
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Moreover, ethicists do not operate in a vacuum, but work within particular paradigms and narratives that influence their analyses and conclusions. Therefore ethicists will seldom be able to ‘solve’ ethical problems, such as the use of human embryonic stem cells or nanotechnology, the same way that scientists, engineers or physicians are able to solve the challenges they face. What ethicists do is analyse the ethical relevance of a medical/scientific innovation, of the circumstances under which it is applied and of its ends and ramifications. They argue their position and advise those who must eventually make their own decisions. Sometimes the advice may become a consensus and is formalised later in guidelines, codes of conduct, conventions or legislation at national or supranational level (Council of Europe, 2005; World Medical Association, 1997; The European Parliament and the Council, 2001; 2004). In this chapter, we will focus on some of the ethical issues generated by new developments in the field of biomedical composites. The rapid development of the field, the broad range of possible composites, and the wide variety of applications of these products make it impossible to provide an exhaustive overview of all products and the various ethical challenges they present. We will therefore highlight some general ethical issues that may arise as a result of the development of biomedical composites and the products that are derived from them. We will firstly investigate which features of biomedical composites are ethically relevant and secondly go deeper into some ethical questions and the impact of these issues on their future development and application. Biocomposites have been developed to widen the range of mechanical and biological properties of biomaterials and, consequently, to optimise the structure and performance of the medical devices that are based upon them, and to improve the interaction of these devices with the surrounding tissue. Through their structural characteristics they aim to attain a better structural compatibility with human tissues than earlier single-phase homogeneous and isotropic biomaterials (Ramakrishnan, 2004). However, as research into biocomposites has progressed, their complexity and concomitantly their functional possibilities have grown. Not only has their development given rise to new views on biocompatibility, they are also increasingly being applied in more complex medical devices such as biodegradable stents (Ramcharitar, 2008) and as elements of even more complex products such as human tissue engineered products. The introduction of nanomaterials adds another layer to the complexity of these products (Resnik, 2007). It is these developments that prompt us to investigate what the ethical issues of the development of novel biocomposites are. This analysis covers a broad range: the development of the materials themselves, the application of the materials in medical devices and the testing of these materials in
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clinical trials, the final application of these products and the impact that these (expensive) products may have on the allocation and distribution of health care. Every new development in science and medicine, including this one, should prompt two questions: does it raise new ethical issues? And if not, how should we address well known ethical concerns in the light of these new technologies? When technology evolves, new elements may play a crucial role in judging how ethical principles such as the reduction of risks and the protection of vulnerable persons or the just allocation of scarce resources can be sustained. These concerns require that we investigate how exactly and to what degree the specific features of this innovation may influence the ethically responsible conduct of research and therapy (European Group on Ethics, 2004; Gordijn, 2004). The moral baseline of any form of research, no matter how promising, advanced or new, should be the respect for and well being of those who are the most vulnerable in this process: those who participate in clinical trials and those who receive devices containing biomedical composites and who may, especially in the fledgling state of a development, run serious risks, notwithstanding the assumed benefits. The development of these products may also have an impact on our fundamental anthropological views and the design of health care systems.
23.2
Developments in biomedical composites
We focus on three characteristics of biomedical composites that carry ethical relevance: (1) the changing views on biocompatibility, (2) the combination of composites with other innovative technologies such as nanotechnology and the development of human tissue engineered products, and (3) the use of biomedical composites for enhancement purposes. The third part of this chapter will reflect on the implications of these features for the ethically responsible conduct of research and medicine, not only at the level of the individual patient, but also at societal level.
23.2.1 Changing views on biocompatibility Perhaps the most important evolution in the research into biomaterials is our changing understanding of the concept of biocompatibility (Williams, 2008). Where initially the biocompatibility of medical devices mostly referred to events that had to be avoided, such as toxicity, carcinogenicity, irritation and thrombogenicity (Williams, 1987), it now refers to the ability of a material to perform adequately in the body with an appropriate host response in a specific situation. The emphasis has shifted from producing products that are selected predominantly for
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their mechanical and structural properties and that are as inert as possible to products that actively initiate a complex process in which interaction is not something to be avoided, but something that is to be deliberately encouraged, albeit in a well controlled manner. The material has not simply to exist in the body without causing harm; it has to perform in a way that is appropriate for that particular application and for that specific situation (Williams, 2008). This of course has consequences for the ethical evaluation of these products: the fundamental principle ‘do not harm’ has to be reassessed in the light of the new aims of biomaterials and of biocompatibility.
23.2.2 Combining biomedical composites with other new technologies The development of novel biomedical composites goes alongside other innovations, such as the development of tissue engineering and of nanotechnology. Nanoparticles may become key elements in the development of new composites. However, the field of nanotechnology itself is not devoid of ethical challenges; applying nanotechnology in biocomposites therefore will introduce these issues into the field of biomaterials. The ethical questions that have been raised in the nanotechnology debate are wide ranging, and not all of these may be relevant in the context of biomedical composites, but some elements have to be taken up such as the assessment of safety and toxicity, the use of nanomaterials for enhancement purposes, the just allocation of new therapies, and the protection of the autonomy and privacy of the patient (Health Council of the Netherlands, 2006). Already biomedical composites form the backbone of some human tissue engineered products, allowing the combination and integration of various cell types with these biomaterials. Again, tissue engineering gives rise to a complex network of ethical issues, ranging from the use of specific cell types such as embryonic stem cells or animal derived cells; to the question of whether the ultimate quality of the final human tissue engineered product is sufficiently superior that it merits the production costs compared with the next-best therapy (Trommelmans, 2007a).
23.2.3 Use of technologies for therapeutic or enhancement purposes A third aspect of composite-based products is their eventual use. While they are designed and developed for therapeutic purposes, it is not inconceivable that, given some slight alterations, their features might bring the final product beyond the therapeutic, leading not to a cure for a specific condition, but to the enhancement of a specific function. The search for the
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best possible solution and the best possible product is commendable, as it would be unacceptable to treat patients with inferior products if better products could be designed. However, the barrier between therapy and enhancement is rather flimsy, and the acceptability of enhancement on moral grounds is generally rather limited. The development of novel biocomposites that carry a realistic prospect for use for enhancement purposes must encourage us to think about the desirability of certain developments (Parens, 1998a; 1998b). Given these three elements of the development and (potential) use of biomedical composites, it is necessary to go deeper into their ethical implications. The aforementioned issues should not be answered exclusively by the scientists who develop these composites, but their profound knowledge of the field, of its possibilities and impossibilities and likely developments gives them an important voice in the debate. Their collaboration with ethicists and philosophers may lead to a better understanding of these issues and therefore to a better approach to the problems that may be generated by biomedical composites.
23.3
Ethical challenges in the development of biomedical composites: risks, benefits and safety
The development of these materials is an enabling technology, allowing the creation of more efficient medical devices, of human tissue engineered products or even other products, not an end in itself (Health Council of the Netherlands, 2006). They are developed to address the shortcomings of the older monophasic materials, and in that sense they cannot be labeled as entirely new products that form a category of their own, with some kind of unique set of ‘ethical problems’. Issues such as risk evaluation, determining the safety of the product, the impact that the development has on health care budgets or the long-term effects of these products on our concept of health, normality and what it is that makes us human are neither new nor unique for these materials and their application. These are well established ethical subjects that have been debated in various medical contexts. Yet the development of these materials urges us to reassess these challenges because a profound knowledge of the ethical questions can ensure that choices made in the early stages of a promising development lead to morally relevant and acceptable research trajectories while avoiding ethically questionable advances.
23.3.1 Identifying and weighing risks If Primum non nocere (first do no harm) is a fundamental principle of the ethical conduct of medical research and medicine, than the main concern
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of every physician and researcher should be the avoidance and minimisation of risks. However, as risks can never be completely excluded, and as small risks can go alongside great benefits, it is important to identify these risks and to weigh them against the prospective benefits. Only when this exercise is performed can one inform patients and clinical trial participants adequately and responsibly about the medical devices or other products that are going to be used. Clinical research can be justified only if three conditions are fulfilled: the potential risks to the trial participant are minimised, the potential benefits to individual subjects are enhanced, and the potential benefits to individual subjects and society are proportionate to or outweigh the risks (Emanuel, 2000). Allowing products subsequently for therapy is only acceptable when the benefits clearly override the risks, and when the new treatment is superior to existing treatments. As we argued earlier, the changed views on biocompatibility and/or the tendency to incorporate nanotechnology or living cells in products containing biomedical composites increases the complexity of the products and creates risks that were not present in earlier monophasic biomaterials such as metal alloys, ceramics or polymers (Ramakrishnan, 2004). To illustrate this higher level of complexity, we will elaborate on three examples: (1) the new views on biocompatibility, (2) the development of human tissue engineered products (HTEPs) and (3) the implications of nanotechnology. Biocompatibility In the past, the safety of novel materials could be sufficiently assessed by judging the inertness of the material in the human environment and over a certain period of time. We now have to assess safety of a product that actively interacts with the body and where this interaction is moreover changing over time, especially in the case when the biomedical composite is biodegradable and its activity may depend on the exact site where the integration is envisaged (Williams, 2008). The interaction should not elicit undesirable local or systemic effects. Secondly, the actual behaviour of the biocomposite in one location or form may not be a fully reliable predictor of the behaviour of the same biocomposite in another location or in another device, making extrapolation of trial results more difficult. Thirdly, the actual behaviour of the implant may vary from patient to patient, as no two patients will interact in exactly the same way with a given material. So while the fundamental requirement of safety remains the same, the complexity of the problem and consequently also the difficulty of guaranteeing the safety of patients has changed dramatically (Williams, 2008). If the biocomposite is biodegradable, one has not only to assess the safety of the product itself but also the effects of the metabolites of the degrading products. Degradation mechanisms in the body are not always well under-
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stood, and our knowledge of the behaviour in situ and over a longer period of time of some products and their metabolites is still rather restricted. Consequently, as the focus of biomaterials research is shifting from merely addressing physical and mechanical deficiencies to interacting with the body and boosting the regenerative capacity of the body, the importance of understanding the biochemical and biological interactions between implant and body cannot but grow considerably. Human tissue engineered products The desirability of a beneficial interaction between body and implant is perhaps most clear in the development of human tissue engineered products (HTEPs). HTEPs are perhaps the most complex man-made constructs that are available today for medical purposes. HTEPs are inherently variable products. Even when the design and culture conditions are aimed at the construction of products as uniform as possible, individual HTEPs for a specific application will always show a certain amount of variability because they contain metabolically active cells in a dynamic extracellular environment. Secondly, in tissue engineering, the regeneration of the body is attempted through a physiologically active structure of cells and extracellular material that is implanted in the body, with the intention of coaxing the body into a deliberate and ongoing interaction with the HTEP through cues produced by the inserted product (Williams, 2006a). The ensuing dynamic is threefold: (a) within the HTEP itself, (b) in the body of the recipient which may influence the incorporation of the HTEP (Rothwell, 2005) and (c) the recipient’s body and the HTEP interact with each other (Kuijer et al., 2007). A mismatch between the dynamic generated in the HTEP and the dynamic going on in the patient can render the product at best inefficacious, at worst unsafe. For instance, the ideal implantation time for the HTEP may be dependent on the developmental stage of the HTEP. If it is too ‘green’ or by contrast too far developed, it may not integrate with the surrounding tissue (Ahsan et al., 2005; Caplan, 2007). Understanding the interactions between product and recipient is therefore essential for the successful development of HTEPs (Stocum, 1998). Once the process of integration and regeneration is initiated, it is impossible to reverse it completely, because of this interactivity. If the process goes awry, at best the HTEP can be removed. If, however, cells of the HTEP migrate to other parts of the body and generate unwanted secondary effects elsewhere, or if biomolecules affect other cells then those of the HTEP, mere reversion of the procedure may be pointless in order to solve the newly generated problem. A number of potential hazards can be identified (Halme, 2006; EMEA, 2007). Cell choice, the potential alteration in the genetic
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make-up during cell culture and expansion, and the presence of unwanted cell types or contaminants in the administered HTEP need to be taken into consideration. Standard treatments to enhance safety such as sterilisation or irradiation may not be applicable because HTEPs are dynamic, living structures. Establishing Good Tissue Practices and Good Manufacturing Practices will consequently be crucial for the successful development of HTEPs (European Commission, 2006). The evolution of the HTEP in the body must be monitored and controlled in order to detect and prevent undesired cell migration, propagation or (de)differentiation, and not to evoke rejection or inflammatory responses by the host’s body. All these new risk factors come on top of more ‘common’ risks associated with cell culturing and surgery. The introduction of specific cell types, including animal-derived cells or human embryonic stem cells, and the conditions under which cells are obtained will equally bring ethical issues to the production of these HTEPs (De Vries, 2008). Products containing nanoparticles A third example of the increased complexity of the new biomedical composites is the incorporation of nanoparticles. Nanotechnology has been welcomed by some as one of the most promising technologies, while others have focused on the risks involved (Resnik, 2007). Nanoparticles show a high reactivity and interactivity, which make them interesting subjects of research for use in biomaterials. One issue in the nanodebate concerns the safety of these particles. Materials that have been tested and found safe at the microscale have shown other characteristics at the nanoscale. The possibility of nanoparticles leaching from devices and migrating to other parts of the body, and outside of the body is the subject of investigation as they may considerably harm the patient and the environment, even when the nanoparticles in the composite are very effective. The potential toxicity of a material when it is used at the nano level requires that all materials that have already been found safe at micro level should be reassessed when their incorporation as nanoparticles is considered. Especially their capacity to migrate (and maybe pass the cell membrane and interfere with cellular metabolical processes or pass the blood–brain barrier) should be investigated carefully before incorporating these nanoparticles into the human body. Automatically concluding that nanoparticles made from material A are safe because material A in its conventional form is safe, is not an option. The risks of leaching and migration also implies that – again – deciding on the safety of the biocomposite will have to be done with a systemic view of the action of the product in the entire body and over a long period of time.
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Weighing risks and benefits The three examples we have given here show that the development of new biocomposites present us with new challenges for weighing risks and benefits. Before starting a trial (with the exception of an exploratory trial) investigators should be in clinical equipoise i.e. it should not be clear beforehand which of two interventions – an existing one or the innovation – is best (Freedman, 1987). Being in equipoise, however, requires that one is able to assess to a certain degree the efficacy, the benefits and risks of both interventions in order to be able to compare them. A new kind of trade-off has to be made: against an enormous benefit stand potential large risks, probably even emergent risks owing to the process of regeneration that have never before been encountered. What’s more, some adverse events may only become apparent after the trial. In determining whether there is equipoise in the use of these new materials, both risks and benefits are clearly taken to a higher level. Before the medical devices that contain these new materials are made available, their therapeutic superiority has to be proven – not only in the short term, but also in the long term. However, determining long-term superiority is even more tricky than shortterm superiority. Relying on the results of clinical trials that usually follow patients for a rather short period, compared with the often life-long presence of the device in the human body is an essential, but not sufficient factor in determining the efficacy and safety of the product. Products that are destined to interact dynamically with the body therefore need to be monitored beyond the cut-off point of the trial. Successful application of these materials in devices or HTEPs will be greatly beneficial for the patient, returning him or her to greater functionality and postponing the need for further interventions or even making them redundant. If, however, the attempted intervention fails, especially when efficacious therapies are available, even when they are theoretically ‘second best’ to the new products, the intervention may turn into a hazard. The probable risks and benefits of a biomaterial consequently need to be compared with available therapies whose efficacy has been demonstrated. Unfortunately, consensus on which therapy is best for some conditions has not necessarily been reached, which makes the comparison all the trickier. One of the challenges of these new approaches is therefore the identification, weighing and minimisation of risks. These risks are partly unknown owing to our lack of understanding of the long-term evolution of these products in the body and of the induced regeneration of human tissue in vivo in the case of tissue engineering (Wood, 2006). Both the kind of risks and the magnitude of the risks challenge us. While many of the risks cited here are also present in other medical technologies, it is the complexity of them that leads to an inflation of uncertainties,
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making the risk assessment for novel biomedical composites difficult to perform.
23.3.2 Short- and long-term safety issues A key feature of the new biomedical composites is, as pointed out earlier, the new view on biocompatibility that it has generated. Demanding biocompatibility ‘new style’ of a biomaterial does, however, have some important consequences for the ethically responsible treatment of the recipients of these devices, especially in the long term. When biodegradable materials are used, or when the possibility of particle leaching or cell migration cannot be excluded, the issue of long-term safety and the protection of the recipient come to the fore. The unforeseeability of the various reactions that may occur, and not necessarily at the site of the implant, raises the risks and, more importantly, makes their identification and quantification more difficult. The interaction between device and body will hopefully result in the gradual acceptance by the body of the ‘foreign’ device through which it gradually loses its ‘foreignness’. This is a long-term process that may continue well beyond the duration of the trial. These events, especially adverse events, will still have to be identified and analysed (European Parliament, 2007; van Zuijlen, 2001; Wood, 2006). Trials with their early cut-off points alone will therefore not suffice to make comprehensive projections about the future safety, efficacy and quality of the therapy; neither will they enable researchers to fully appreciate the long-term change in quality of life. Extrapolations of data gathered during trials will allow some prognoses, but they will always contain a degree of uncertainty. The actual long term effects can only be judged during the follow-up process. As has been pointed out by the National Institute for Clinical Excellence (NICE) for one of the few well researched clinical applications of HTEPs, autologous cartilage implants, the use of these products has to be considered as part of an ongoing clinical trial, because of uncertainties about the long-term effectiveness and the potential adverse effects of the procedure (National Institutes for Clinical Excellence, 2005). A similar argument could be applied to the more complex biomedical composites. When developing new technologies it would therefore probably be better to consider trials as an integrated part of the chain of production, testing and application of the devices and their constituents. This approach may allow us to better integrate various findings. Information gathered from clinical trials and from the post-clinical trial phase can lead to an adjustment of the production of the biomaterial and its possible applications in order to provide a better adapted product for a specific application. It might therefore be useful to think of these devices as ‘therapies under control’ or of ‘technology under investigation’ (Wasiak, 2006), in which the producers
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of the basic biomaterials have an important role to play. This also implies that good governance of post-trial events needs to be established (Trommelmans, 2007b).
23.3.3 Informing participants and patients about short- and long-term effects A fundamental requirement for research and medicine is respect for the autonomy of the trial participant and patient. For persons to be able to decide autonomously on their participation in a trial or to choose the treatment that is the most suited for their lifestyle and that does not go against their world views or convictions, it is essential that they should be adequately informed about the various aspects of the product, such as its composition and of the risks and benefits that can be anticipated. Interventions that involve a large amount of uncertainty, or that can produce severe unwanted or unpredictable side effects are usually restricted to life-threatening diseases for which no alternative is available. In those cases, the potential benefit (saving life) weighs up against the risks. However, for many medical devices containing innovative biomaterials there are perfectly valid and tried alternatives, even if they are not the perfect solution to the patient’s problem. The complexity and the newness of the materials, especially when combined with other techniques, the problematic assessment of risks and benefits and the availability of reliable alternatives present a complex puzzle. Trial designers, investigators and physicians will therefore have to demonstrate convincingly that the benefits of the new devices or HTEPs will be at least as good as the existing treatment, while the risks involved in applying them will not exceed those of the conventional treatment. If the risks exceed those of conventional treatment, it is reasonable to expect that the proposed benefits will indeed be realised and that they would render the increased risks acceptable. Given the limited knowledge of both risks and benefits and given that there often is no general consensus on what actually is the best treatment among existing therapies, this calculus will be extremely difficult (Trommelmans, 2008).
23.4
Therapy or enhancement?
Developing increasingly better biocomposites may eventually lead to their use, not for therapeutic purposes, but for enhancement purposes. The ethical debate about the moral (in)acceptability of enhancement is rekindled whenever a new medical development comes under attention (Parens, 1998b). The distinction between therapy and enhancement is rather difficult to make: both approaches aim at the same goal; to improve a particular
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condition and/or the quality of life of a specific person. Especially in the case of medical devices, the difference seems rather slim: implanting a device will, if it is successful, certainly improve the patient’s condition. It may also enhance his or her condition in the sense that he or she will not only perform better than before, but also outperform comparable persons. The example of the South African athlete Oscar Pistorius who had to fight the International Association of Athletics Federations (IAAF) to be allowed to compete against able-bodied athletes in the Olympics is a case in point: the IAAF moved to block him from the Olympics, with a new ruling banning ‘technical aids’ (his carbon fibre prosthetics that replace his amputated lower legs). Pistorius has been dogged by claims that the blades give him an extra long stride; that his performance was somehow enhanced through the use of his prosthetics, something he denies. Likewise one could imagine the development of biomedical composites that make the body that little bit more resistant to wear and tear, that provide that little bit more strength. So where lies the ethically relevant distinction between therapy and enhancement, because there seems to be a consensus that treatment of ill or disabled persons is morally imperative but enhancing them is not, maybe even morally illicit? One of the most used arguments is the ‘normal functioning model’ of Norman Daniels (Daniels, 1992). In this model, the fundamental criterion for the ethical acceptability of an intervention is the restoration or preservation of the species specific functions, possibly with corrections for age, gender etc. Everything beyond this criterion is to be considered as an enhancement and is not allowed. However, this criterion is not so easily sustained: the boundary between enhancement and therapy is rather vague if we use the normal functioning model. Also, this does not answer the question why enhancement should or should not be allowed. Applying the same product in different persons can moreover lead to one person being cured and another one being enhanced by the same intervention, based on their respective condition before the implant. Should we then provide ‘inferior’ products to those who have a less good condition, so as not to enhance them, and better devices to those who enjoy a good condition in order to cure them? It is likely that enhancement is in some sense nearly unavoidable when developing and applying new biomedical composites. But should we consciously aim to produce biomedical composites to enhance persons? One of the most heard arguments against enhancing persons is that such an intervention undermines some of our fundamental intuitions. Enhancement is then perceived as a form of cheating: one allows people to obtain results not through a virtuous life pattern or sustained exercise, but through dubious methods (Parens, 1998) in which there is no merit on the part of the recipient, except perhaps of having a well-furnished wallet. That type of enhancement is usually viewed with suspicion, because it seems to make the effort ‘less authentic’ (Juengst, 1998).
Ethical issues affecting the use of biocomposites
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Another argument against enhancement is the slippery slope argument (Launis, 2002), in which one argues that applying a new technology which is not morally problematic in one specific context may eventually lead to other applications that are morally problematic in another context. So we may not find it problematic to provide a wounded soldier with prostheses to replace his missing limbs, but may find it problematic to apply the same techniques to provide an athlete with these prostheses and giving him an unfair advantage in track and field competitions. The slippery slope argument, although popular, does not carry much ethical weight. However, the argument is certainly an expression of a deeper underlying feeling of uneasiness with new technologies and their social and psychological impact. The development of new biomedical composites and their capacity for exploitation in a wide variety of applications may certainly influence our future view on the limits of tissue degeneration and on tissue repair or even regeneration through the use of biocomposites. Especially in the use of medical devices a grey area between therapy and enhancement exists. While some composites are clearly solely designed for and aimed at therapeutic purposes, and will never be used to enhance specific features; some composites may be designed to answer the desire to enhance the body, e.g. some implants used in cosmetic surgery. The inherent possibility of the materials not only being used for therapeutic but also for enhancement purposes requires at least that we are attentive to the underlying concerns of persons about the meaning of health and illness and the capacities of the human body. Whether we allow these applications or not will not primarily depend on our views on the essence of what makes us human and how we can and should promote individual and societal flourishing, but the fact that these products can be designed and produced, makes the researchers and physicians who are involved in them important voices in the debate, as they are the persons who can inform ethicists and the general public about the (im)possibilities of specific methods and technologies.
23.5
Implications of the application of biocomposites in therapy
23.5.1 Best standard of care If these novel materials prove to be superior to other materials, logic compels us to adopt them as the standard of care. One should after all not knowingly deny superior treatment of net advantage to the recipient (Van der Graaf and Van Delden, 2009). However, while these products may from a purely therapeutical viewpoint turn out to be superior to any other intervention, their price may be prohibitive on a larger scale. The financial
606
Biomedical composites
sustainability of these products will then have to be taken into consideration (Trommelmans, 2009). The demand for all types of medical devices and of tissue-engineered products is very likely to increase considerably over the next years. Failing and degenerating tissues and organs are an increasing challenge to health care, not only in the developed world, but also in the developing world (Greenwood, 2006). A growing life expectancy, changing life styles and medical advances lead to an increasing incidence of degenerative diseases whose impact on patients and society will grow considerably over the coming decades. Interventions such as the use of prostheses, replacement therapies and the development of human tissue engineered products will become more common. The demand for long lasting, ‘intelligent’ materials to be incorporated in these devices will as a consequence grow alongside. It is likely that these materials will not only be focused on restoring functionality, but, wherever possible, also on regeneration of the body.
23.5.2 Access to therapy The increased introduction of ever more complex products may have a beneficial effect on the life quality and life expectancy of individual recipients, provided these products and materials meet strict safety and efficacy standards, otherwise there is no point in applying them. The other side of the coin is that some of these products and materials will be markedly more expensive and will lay a serious burden on our already overstretched health care budgets. The ethical issue of the just distribution of health care will therefore become once again more urgent. The debate on the just allocation of health care is not recent, but as technology becomes more expensive, its relevance grows considerably. Providing every person with the best possible device, based on the best material available may either lead to a financially untenable situation for public health care, or to a system in which only the wealthiest citizens can afford the best care, while other, less affluent citizens have to be satisfied with something less qualitative. The debate on how to allocate services and products is far from concluded. Defining strategies for setting priorities in who gets which material is rapidly becoming a pressing ethical and health policy issue and, although some may argue that the development and application of these novel materials will in time lead to lower prices for the devices and will have a beneficial impact on the health care budget, experiences with other medical innovations have taught us not to overestimate this effect. The debate on the allocation of health care is not only restricted to who gets which product, it should also address the difficult question on how good material X or Y should be compared with material A or B in order to merit
Ethical issues affecting the use of biocomposites
607
their introduction and/or reimbursement by public health agencies. Even when X or Y has proven to be better than A or B, the qualitative difference between both types may be so small that it does not merit the eventual introduction. If these materials eventually become available for therapy, their actual application will therefore not solely depend on the clinical evidence but also on other factors, including the evaluation of the economic viability of these products (Bock, 2003; McAteer, 2007) and their reception by end users, both surgeons and patients. These products will unquestionably be very expensive, which implies that the decision to apply them will also have repercussions for health care organisation. Besides the economic aspects, the ethical (and political) issue of the just distribution of health care is at stake. If we want to uphold the principle that those most in need of a specific treatment should be able to get it, whatever its cost and regardless of the patient’s financial resources, we will have to apply some reimbursement strategy. Which products will be ultimately available, and which conditions will be treated will partly depend on the initial choices that are made in the construction of the biomedical composites. These choices should therefore not be made solely from a technical point of view, but also with the desired ethical outcome in mind. Such an approach has, however, not yet been studied in depth.
23.6
Conclusion
The development of better biocomposite materials may improve the design of medical devices, and increase the quality of life for an increasing number of persons. Research into these materials is therefore ethically commendable. The growing complexity of these products, the changing views on biocompatibility and the combination of the technology with other new technologies such as tissue engineering and nanotechnology lead us to reinvestigate some of the conditions that need to be fulfilled in order to make this development ethically acceptable. The benefits of the products may be very large, and further research has to be encouraged. However, more research is needed to determine the various safety issues we have highlighted in this chapter. One has further to investigate the type and size of the possible risks, and to critically investigate the benefit/risk ratio of these products, not only for the recipients themselves, but also for society at large. One has equally to be aware that the development of these products may lead to undesirable applications and that their application may have a considerable impact on the organisation of health care services and the allocation of health care. Given the increasing health care costs, and the growing need for medical devices by an increasing proportion of the population, defining the criteria at societal level of who gets which device, based on a specific material, will become as urgent in the future as defining
608
Biomedical composites
the risks and benefits for an individual person in the setting of the doctor– patient relationship. Further research will also have to focus on the development of appropriate informed consent procedures and forms, and on the influence that the development of these technologies may have on our ideas of what makes us human (Health Council of the Netherlands, 2006; Khushf, 2007; Resnik, 2007).
23.7
References
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Index
ABAQUS, 310 abrasive wear, 453–4 access to therapy, 606–7 Achilles tendon, 16–17 Acroflex-100, 185 acrylic bone cements, 210–29 see also poly(methyl methacrylate) composition, 214–18 for joint replacement, 210–29 hip and knee joint, 210–12 joint replacement, 212–13 properties, 218–22 mechanical properties, 219–22 porosity, 218–19 acrylonitrile–butadiene–styrene, 488 adhesions, 85 affinity index, 398 Al–Cu–Fe, 452 3-aminopropyltriethoxysilane, 519 angiogenesis, 65, 587–8 angiopoietin-1, 587 Animal Welfare Act, 386 anterior cruciate ligament, 15–17, 235–6 antibiotic laden bone cements, 223–5 apatite, 365 articular cartilage, 4 case study on tissue-engineered scaffold, 311–12 mechanical properties, 11–14 swelling behaviour, 13 tensile–compressive properties, 14 morphology, 300–1
ASTM F1612–95, 288 ASTM F 2077, 190 ASTM F 1439–92, 394 atomic force microscopy, 7, 346 Auger electron spectroscopy, 347 autoclaving, 29–30 autologous cartilage implants, 602 autologous chondrocyte implantation, 88 autologous nerve grafting, 89 axial skeleton external applications, 54 internal applications, 50–4 Cambridge Cup, 53 lateral malleolus fracture, 52 response to carbon fibre reinforced epoxy resin fracture fixation plates, 51 BAK interbody fusion system, 181–2 Bard composix mesh, 84 barium sulfate, 222 basic fibroblast growth factor, 587 bending-extension coupling, 421 bending tests, 8 beta cells, 64 bioactive coating, 132–3 chemical bonding, 132–3 mechanical bonding, 132 bioactive glass ceramics, 104 bioactivity, 356 Biobrane, 81 bioceramics, 128
611
612
Index
biocompatibility, 356, 595–6, 598–9 evaluation for biocomposites, 325–49 and the biological environment, 327–31 future trends, 348–9 relevance of employed analyses, 334–6, 340–1 surface characterisation, 341–2, 345–7 surface effects and characterisation, 331–4 in vivo testing of biocomposites, 385–407 biofunctionality, 404–5, 407 bone growth inside a biocomposite scaffold, 405 3D model of scaffold implant, 407 ethical and legal requirements in preclinical experimentation, 386–7 extraction and sample preparation, 389 genotoxicity, carcinogenicity, and reproductive and development toxicity, 393–4 haemocompatibility, 394–5 irritation and sensitisation test, 389–92 ISO10993, 387–8 maximal intensity projection, 406 pathological conditions, 402–4 systemic toxicity, 392–3 test for local effects, 395–9, 401 biocomposite fabrication inorganic materials, 105 polymer matrix and reinforcing crystals, 106 polymers, 103 biocomposites, xxi–xxiii biocompatibility evaluation, 325–49 biological assessment, 335–6, 340–1 biological environment, 327–31 host reactions after implantation, 328
principal cell types in the inflammatory, foreign body, and wound healing response, 328 tissue/cellular host response to injury, 329 in vivo tests, 330 for bone repair, 101–21 component selection and general design considerations, 102, 104–6 composite scaffolds, 116–20 fabrication of particulate composites, 106–9 key challenges, 120–1 nanocomposites fabrication, 109–14 template-mediated formation of nanocomposites, 114–16 cellular response, 354–78 definitions, 364 experimental testing, 371–4, 376–8 musculoskeletal applications, 365–71 skeletal regeneration and reconstruction, 355–64 skeletal tissues, 364–5 in vitro tested examples, 372 composite interbody fusion devices, 188–90 poly(ether imide) fibre-reinforced cage, 189 composite intervertebral disc prostheses, 190–5 composite biomimetic total IVD prostheses, 192 PHEMA/PCL compressive properties, 194 stress–strain curve of swollen PHEMA-based composite hydrogels, 193 design and fabrication, 25–41 conventional composite processing techniques, 27–30 influence of processing parameters, 38–40 medical applications, 40
Index production techniques, 27 solid free-form fabrication technologies, 33–4, 36–8 solution-based techniques, 30–3 technology and materials applications, 41 types of constructs, 26 development for regenerative medicine, 547–67 bio-hybrid composites for bonelike scaffold, 550–8 future trends, 565–7 hierarchical structure, 558–60 new approach, 549–50 three-layered osteochondral scaffold, 562–4 ethical issues in usage, 593–608 developments, 595–7 implications of application in therapy, 605–7 risks, benefits and safety, 597–603 therapy or enhancement, 603–5 fatigue behaviour, 465–95 behaviour in hard tissue applications, 473–81, 483–91 behaviour in soft tissue applications, 491–5 fundamentals of failure, 466–9, 471–3 future trends, 348–9 hard tissue applications, 44–56 advantages, 55 axial skeleton applications, 50–4 disadvantages, 55 future trends, 55–6 head and neck applications, 45–50 mechanics, 411–37 applications of biomaterials, 414 fibre-reinforced and particulate composites, 413 lamina and laminate in basic design concepts, 413–23 particulate composites, 426–8, 430–1 polymer nanocomposites, 431–7 short-fibre composites, 423–6
613
phases continuous matrix phase, 364 reinforcement phase, 364 relevance of employed analyses, 334–6, 340–1 soft tissue applications, 59–92 composite biomaterials for soft tissue repair, 70–9 composite nature of extracellular matrices, 61–70 use in clinical intervention, 80–90 for spinal implants, 178–97 basic concepts, 185–7 surface characterisation, 341–2, 345–7 atomic force microscopy, 346 contact angle, 341–2 Fourier transform infrared spectroscopy – attenuated total reflectance, 342 other methods, 347 scanning electron microscopy, 345–6 secondary-ion mass spectroscopy, 347 transmission electron microscopy, 347 X-ray techniques, 346–7 surface effects and characterisation, 331–4 surface chemistry, 332–3 surface morphology, 333–4 surface topography, 334 in tissue engineering and regenerative medicine, 573–88 cell/material interactions, 576–8 chemical aspects of materials, 581–4 physical aspects of materials, 578–81 specific processes in regenerative medicine, 584–8 tribology, 441–60 experiment on composite characterisation, 443–4 polymer composites, 444–60 tribo-biomaterials, 442
614
Index
in vivo biocompatibility testing, 385– 407 biofunctionality, 404–5, 407 bone growth inside a biocomposite scaffold, 405 3D model of scaffold implant, 407 ethical and legal requirements in preclinical experimentation, 386–7 extraction and sample preparation, 389 genotoxicity, carcinogenicity, and reproductive and development toxicity, 393–4 haemocompatibility, 394–5 irritation and sensitisation test, 389–92 ISO 10993, 387–8 maximal intensity projection, 406 pathological conditions, 402–4 systemic toxicity, 392–3 tests for local effects, 395–9, 401 biodegradability, 581–2 biofunctionality, 404–5, 407 Bioglass, 107, 113, 128, 326, 377, 518 biomaterials, 356 composite coatings design, 131 biomimetic approach, 535–6, 551 biomimetic deposition for implants, 156–9 apatite/TiO2 composite coating cross-section, 160 surface analysis of an H2O2oxidized NiTi SMA sample, 157 XRD patterns of NiTi SMA samples, 159 for tissue engineering scaffolds, 163–7 apatite/collagen composite coating, 164 MTT assay, 166 Saos-2 cells cultured on apatite/ collagen coated PLLA scaffolds, 165 BioScaffolder, 36
bis-GMA–TEGDMA see bis-phenolA-diglycidyl dimethacrylate triethylene glycol dimethacrylate bis-phenol-A-diglycidyl dimethacrylate triethylene glycol dimethacrylate, 476–8 bone, 4–11 see also specific bones hierarchical composite, 5–7 human bone structure, 6 mechanotransduction, 586 mineral component, 5 morphology, 299–300 organic matrix, 5 structure, composition and properties, 4–5 structure–property relationships, 7–11 bone tissue elastic constants, 10 bone analogue biomaterials, 129 bone healing, 360–2 bone implants test for local effects, 395–9, 401 bone biopsies, 400 dynamic histomorphometry, 399 histological appearance of cortical bone implants, 401 histological appearance of trabecular bone implants, 402 histological appearance of vertebral pedicular implant, 398 histological assessment parameters, 397 lumbar vertebra section, 403 polymeric biocomposite porosity, 404 screw implant position, 396 bone mirror area, 398 bone morphogenesis cascade, 357–8 bone morphogenetic proteins, 360 bone remodelling, 361 bone repair biocomposites for, 101–21 component selection and general design consideration, 102, 104–6 composite scaffolds, 116–20 key challenges, 120–1
Index nanocomposites fabrication, 109–14 particulate composites fabrication, 106–9 template-mediated formation of nanocomposites, 114–16 case study on tissue-engineered scaffold, 307–9 scaffold parameters comparison, 309 fatigue behaviour, 473–76 S–N curves for CF/PEEK and stainless steel bone plates, 475 injectable composites, 255–70 biological behaviour, 267–8 bone tissue engineering, 268–9 classification, 257–63 stability, rheology and injectibility, 264–6 bone resorption, 361 bottom-up approaches, 536 braiding process, 285–6 Buehler closed-patch test, 390–2 burn treatment, 80–2 calcar bone resorption, 282 calcium carbide, 559 XRD analysis after hydrothermal treatment, 561 calcium phosphate, 256, 550 calcium phosphate cement, 256 calcium phosphate minerals, 104 calcium phosphate suspensions, 256 calcium salts, 228 Cambridge cup, 52–4 cancellous bone, 9, 299 carbon, 459 carbon fibre reinforced nylon-12, 483–4 S–N compressive fatigue curves, 485 carbon nanotubes, 78, 431–2, 448, 449–52, 530–4 Bioglass scaffold coated with CNT, 532 electrophoretic deposition cell, 531 polyurethan foam coated with CNT, 533 wear coefficient of CNT/HDPE composites, 452
615
carbon–carbon double bonds, 203 carburisation, 559 carcinogenicity, 394 cartilage, 87–8 see also specific types of cartilage cell adhesion, 577 cemented fixation, 212 cements, 263 ceramics, 26 ceramisation, 559 CharitéIII, 184–5, 190 chemical bonding, 132–3 chemical etching, 517–19 dissolution process of a bioactive glass-ceramic scaffold, 519 chemical stability test see potentiodynamic polarisation tests chitosan, 74, 114 chondrogenesis, 312 chondroitin sulfate, 71 chondroitin sulfate glycoforms, 64 circumferential tensile test, 495 cobalt–chrome metal sponge, 182 cobalt–chromium–molybdenum alloy plates, 183 coefficient of thermal expansion, 137, 207 collagen, 66–9, 235, 300, 550–8 FTIR spectra of doped HA/Coll, 557 ICP-OES quantitative analyses, 554 mechanical properties of HA/Coll, 558 nuclei of HA formed inside a collagen fibre, 553 porosity in self-assembled collagen, 552 TEM micrograph of freeze milling Ha/Coll, MgHA/Coll, HREM images and Fourier transform, 555 type I collagen, 66–7 type II collagen, 67–8 type IV collagen, 68 type V and type VI collagen, 69 XRD spectrum of HA/Coll, 553 composite acetabular cups, 281–2
616
Index
composite cements, 225–8, 280–1 composite coatings for implants and tissue-engineering scaffolds, 127–69 coatings for implants, 148–61 coatings for tissue-engineering scaffolds, 161–7 design, 130–3 technologies for biomaterials surface modification, 133–47 composite femoral protheses case study on tissue-engineered scaffold, 309–11 composite films, 77 composite hip, 284–90 modelling, 286–8 femoral bone elastic properties, 287 stem technologies, 284–6 basic prototyping technologies, 284 continuous-fibre-reinforced hip prostheses, 286 in vitro testing, 288–90 fatigue test set-up, 288 glass and carbon fibre reinforced PEI fatigue behaviour, 289 composite hip stem, 282–3 composite laminate theory, 244 composite materials tissue replacement and tissue engineering scaffolds, 242–50 filament wound structure, 244 loss modulus vs static load, 247 storage modulus vs static load, 246 tendon and ACL fibre-reinforced prosthesis vs Achilles’ tendon & ACL, 245 composite powder, 149 composite scaffolds, 116–20 HAP/PCL scaffolds, 118 composite technology, xxi–xxii compression, 13 compression moulding, 28–9, 284–5 compression–compression fatigue tests, 484
computed tomography, 301–2 computer-aided system, 302–3 tissue engineering, 306–7 library of CAD-based unit cell, 307 conformal evaporated- film-byrotation, 535 contact angle, 341–2 conventional composite processing techniques, 27–30 see also specific techniques autoclaving, 29–30 compression, 28–9 extrusion and injection for thermoplastic materials, 27–8 filament winding, 28 infusion, 29 cortical bone, 9, 299, 365 screw implant position, 396 crimp pattern, 235–6 CultiSphers, 90 cytotoxicity, 340 3D fibre deposition technique, 72–3 Dacron, 86, 495 Darcy’s law, 308 delamination, 471 dental composites, 201–9 aging time and fatigue life, 482 classification, 205–7 filler particle size, mechanical and physical properties, 206 hybrid, 207 microfilled, 206–7 small particle, 206 traditional, 206 fatigue behaviour, 476–81, 483 fibrous composites overview, 208–9 flexural strength of Vectris–Pontic dental composites, 482 general property requirements, 207–8 matrix monomers, 202–4 methacrylate-based monomers, 203 property requirements, 204 reinforcing agents, 204–5
Index restorative dentistry, 201–2 shear S–N curves for dental restorative materials, 479 silane coupling agents, 205 trade names, manufacturer and composition of restorative composites, 478 dental post, 480 dip-coating techniques, 582 direct growth methods, 534–5 direct polymer melt deposition, 527 dual energy x-ray absorptiometry, 54 dynamic torsional loading, 477–8 E-glass, 459 Eilers equation, 428 Einstein coefficient, 428 elastic constants, 417 elastic modulus, 8, 186, 287 elastin, 70, 235 elastin-like polypeptides, 79 elastokines, 70 electrochemical deposition, 142–3, 160 electrochemical impedance spectroscopy, 142 electron beam lithography, 514, 516 three-layered vascular scaffold supporting 3 different cell types, 515 electron spectroscopy, 347 electrophoretic deposition, 531 electrospinning, 31–3, 73, 165, 304 hybrid process, 528 nanofibres and nanofibrous scaffolds, 525–8 process diagram, 32 embryonic stem cells, 358–9 emulsion freezing / freeze-drying, 165 endodontic treatment, 208 endotendon, 15 endotenon, 235 energy dispersive x-ray, 156, 346 engineering strain vectors, 418 environmental temperature, 217
617
epitenon, 236 etching, 132 ethics biocomposite developments, 595–7 combination with other new technologies, 596 technologies for therapeutic or enhancement purposes, 596–7 views on biocompatibility, 595–6 implications of biocomposite application in therapy, 605–7 access to therapy, 606–7 best standard of care, 605–6 risks, benefits and safety of biocomposite development, 597–603 identifying and weighing risks, 597–602 informing participants and patients of treatment effect, 603 short- and long-term safety issues, 602–3 therapy or enhancement, 603–5 use of biocomposites, 593–608 Exact family, 494 extracellular matrix composite nature, 61–70 collagen, 66–9 elastin, 70 fibrin, 63 fibronectin, 69 glycosaminoglycans and proteoglycans, 64–5 laminin, 70 extraction, 389 extrusion, 27–8 Factor XIII, 63 fascicles, 15 fatigue, 220–1 behaviour in hard tissue applications, 473–81, 483–91 bone repair and replacement, 473–76 dental applications, 476–81, 483 joint replacement, 483–8 spine surgery, 489–1
618
Index
behaviour in soft tissue applications, 491–5 tendons and ligaments, 491–4 vascular grafts, 494–5 of biocomposites, 465–95 fundamentals of failure, 466–9, 471–3 damage during fatigue life, 470 Goodman diagram, 469 S–N diagrams, 468 stress-controlled fatigue loading, 467 typical Paris plot, 470 unidirectional carbon/epoxy laminates S–N diagrams, 472 stress cycle to fracture, mixing technique and antibiotic inclusion, 221 fatigue delamination testing, 471 fatigue resistance, 479–80 fatigue tests, 183–4 femoral bone, 277 FGC see functionally graded coating FiberKor posts, 480 fibre-reinforced composites, 237–42, 453–60 case studies, 307–15 articulate cartilage scaffolds design, 311–12 composite femoral prostheses design, 309–11 tendon and ligament tissueengineered scaffolds, 313–15 tissue-engineered scaffold for bone repair, 307–9 chopped-fibre composites, 453–4 computer-aided tissue engineering, 306–7 design tools, 301–4 CAD/CAM processes, 302–3 medical modelling, 301–2 tissue-engineered scaffolds manufacturing, 303–4 harnessing directional properties of biomaterials, 297–9
morphology of load-bearing tissues, 299–301 articular cartilage, 300–1 bone, 299–300 tendon and ligament, 300 orthotropic lamina, 238 in silico computational analysis, 304–6 soft composite design principles, 239–42 structure and angle variation, 240 triangular unit element, 240 unidirectional and woven fabric reinforcements, 455–60 composite cups wear rates, 458 effect of load on friction and wear, 456 volume wear rates, 458 wear resistance ratios, 459 wear volume plotted against sliding distance, 457 fibrin, 63 fibroblasts, 60 fibronectin, 69 filament winding, 28, 187, 239–42, 285 soft fibre-reinforced structure and angle variation during winding, 240 triangular unit element, 240 finite element model, 267 flucloxacillin, 224 fluoride ion, 208 foetal bovine serum, 521 foreign body reaction, 585 Fourier Transform Infrared Spectroscopy, 154 Fourier transform infrared spectroscopy – attenuated total reflectance, 342 surface analysis methods, 345 surface chemistry, 343–4 freeze-drying process, 71 functionally graded coating, 130 fused deposition modelling, 304 γ-methacryloxypropyltrimethoxysilane, 205
Index genipin, 74 genotoxicity, 340, 394 gentamicin, 224 glass bead composites, 427 glass transition temperature, 333 glial-cell-line-derived neurotrophic factor, 64 glycosaminoglycan, 64–5, 555 Gore-Tex dual mesh, 84–5 Gore-Tex prosthesis, 237 grit-blasting, 132 growth factors, 359–60 HA see hydroxyapatite HA/Coll composites, 552–8 HA nanocomposites, 376–7 HA–bone cements, 373–4 human osteoblasts on PHEMA/PCL biocomposite, 375 mesenchymal stem cells on PHEMA/ PCL biocomposite, 375 haematoma, 361 haemocompatibility, 394–5 haemolysis, 283 Halpin–Tsai equations, 417, 424 HAPEX, 107, 372–3, 474 human osteoblast cells, 373 middle ear prosthesis, 49 orbital floor implant, 46 hard tissues axial skeleton applications external applications, 54 internal applications, 50–4 biocomposites applications, 44–56 advantages, 55 disadvantages, 55 future trends, 55–6 head and neck applications, 45–50 aural, 48–9 dental applications, 50 maxillo-facial, 45–8 Haversian system, 6, 277, 299 HDPE see high density polyethylene heart valves, 86–7 heparan sulfate, 64 heparin, 110 hierarchial composite, 5
619
high density polyethylene, 106–8, 474 hindered settling, 265 hip arthroplasty, 279–83 composite acetabular cups, 281–2 composite bone cements, 280–1 composite hip stem, 282–3 healthy hip joint, 280 hip joint, 277–9 anisotropy in the yield and failure properties, 278 elastic constants, 278 resultant force, 279 hip joint prostheses composite materials, 276–90 composite hip, 284–90 hip arthroplasty, 279–83 properties, 277–9 HPPA. see hydroxyl poly calcium sodium phosphate human tissue engineered products, 599–600 HYAFF11, 250 HYAFF11–HA, 374 hyaline cartilage see articular cartilage hyaluronan, 65 hydrofluoric acid, 517 HydroThane, 248 hydroxyapatite, 5, 47, 72, 74, 106–8, 128, 448, 550 nanoparticles, 259 wood transformation during hydrothermal process, 561 hydroxyl poly calcium sodium phosphate, 154 Ilizarov external fixator, 54 image analysis, 19 immunity, 585 immunological analysis, 74 implants composite coatings, 127–69 biomimetic deposition, 156–9 ion beam assisted deposition, 153–6 other techniques, 160–1 plasma sprayed coatings, 148–51 spraying and sintering, 151–3
620
Index
in silico computational analysis, 304–6 inductively coupled plasma optical emission spectroscopy, 554 Infinity, 517 inflammation, 65, 584 injectability, 266 injectable bone substitute background, 256 backscattered scanning electron micrographs of new bone formation, 261 biological behaviour, 267–8 animal models, 267–8 cellular behaviour, 267 for bone repair, 255–70 classification, 257–63 particulate scaffold in a matrix, 259–60, 262 porous hardening matrix, 263 microtomographic images, 260 objectives bone in-growth, 257 wound healing, 257 schematics of physical presentation, 258 usage, 256–7 injection moulding, 27–8 inorganic materials, 105 Integra, 81–2 integrins, 577 interbody spacers, 181–2 interdigitation, 222 InterFix, 190 interstitial exclusion, 61 intervertebral disc prostheses, 182–5 intervertebral discs, 180–1 ion beam assisted deposition for implants, 153–6 HPPA–Ti FGC depth profiles, 155 HPPA–Ti FGC surface morphology, 155 irritation test, 390 ISO 10993, 387–8 Animal Welfare Requirements, 386–7 ISO Rules, 388 ISO 7206-1, 288
ISO ISO ISO ISO ISO ISO
7206-3, 288 10993-3, 394 10993-4, 395 10993-6, 395 10993-11, 392 10993-12, 389
joint replacement acrylic bone cements for, 210–29 antibiotic laden, 223–5 composite cements, 225–8 hip and knee joint, 210–12 properties, 218–22 radiopacifiers, 222–3 fatigue behaviour, 483–8 effect of MWCNT on flexural quasi-static strength and failure cycles, 488 fatigue crack propagation curves, 487 MWCNF effect on Weibull mean parameter, 489 PE–butene matrix and winding angle of UHMWPE fibres effect fatigue limit and degradation, 489 S–N compressive fatigue curves of CF/PA12 cylindrical tubes, 485 static and fatigue tests setup, 484 Kennedy Ligament Augmentation Device, 237 Kennedy model, 492 keratinocytes, 60 Kerner equation, 426 Kevlar, 281 Kevlar-29, 54 Kevlar-49, 459 KFRP10, 456–7 knee joint, 211 lamina composite materials design, 413–23 fibre-reinforced, 416 generic three-layered laminate, 420 laminate, 415 laminate model, 424
Index lamination theory, 242, 420 three-layered laminate, 420 laminin, 70 LEGO, 512 ligaments, 14–18 biocomposites fatigue behaviour, 491–4 single ACL matrix fibroin cords, 493 hierarchial spatial arrangement, 15 mechanical properties, 16–18 uniaxial tension stress–strain for anterior cruciate ligament and Achilles tendon, 17 ligaments and tendons composite materials, 234–51 devices used in replacement, 236–7 fibre-reinforced polymeric composites, 237–42 future trends, 250–1 tissue biology and anatomy, 235–6 tissue replacement materials and tissue engineering scaffolds, 242–50 ligamentum nuchae, 18 LIM mineralization protein, 567 line-of-sight, 134 LINK SB Charité prosthesis, 184–5 liquid-to-powder ratio, 265 lower urinary tract, 87 lymphocytes, 328–9 lyophilisation, 72 macrophages, 329–30 macroporosity, 263 magnesium, 554 magnetic resonance imaging, 301–2 magnetic scaffold, 565 Magnusson–Kligman maximization test, 390–1 marginal leakage, 208 martensitic bearing steel 100Cr6, 450 matrix metalloproteinases, 62 maxillo-facial region, 45–8 maximum strain theory, 425 maximum test frequency, 468 mechanical bonding, 132
621
mechanics of materials approach, 416 mechanotransduction, 585–6 MedCAD, 309–10 medical modelling, 301–2 mesenchymal stem cells, 358–9 metals, 26 methyl methacrylate, 202 microfabrication techniques, 74 microglia, 69 microindentation technique, 150 microtopography, 579 mineralisation, 551 molecular mechanotransduction, 66–7 monocytes, 328–9 montmorillonite nanocomposites, 433 mould heating, 284 MTS Bionix 858 Test System, 193 multi-walled carbon nanotubes, 78, 227, 431–3 nanocomposites, xxii fabrication, 109–14 HAP crystal surface modification, 110 HAP/PHBV solvent cast films produced using PAA stabilised HAP nano crystals, 111 template-mediated formation, 114– 16 HAP/chitosan composite, 115 nanoindentation technique, 11, 150 nanoparticles, 520–4, 600 fabrication process of porous chitosan scaffolds, 524 SEM of bioactive glass nanoparticles, 523 for tissue engineering applications, 520 total protein adsorption study, 521 nanostructured biocomposites bottom-up approaches, 536 direct fabrication of surface nanotopographies in 3D structures, 516–19 chemical etching, 517–19 polymer demixing, 516–17 future trends, 537–8
622
Index
nanoparticles, nanotubes and nanofibres, 520–34 carbon nanotubes, 530–4 nanocomposite approach, 520 nanofibres and nanofibrous scaffolds by electrospinning, 525–8 nanoparticles, 520–4 phase separation and particle leaching, 528–30 processing 2D topographies for 3D structures assembly, 513–14, 516 electron beam lithography, 514, 516 photolithography, 513–14 sol-gel, direct growth and biomimetic approaches, 534–6 biomimetic processes, 535–6 direct growth methods, 534–5 sol-gel methods, 534 for tissue engineering scaffolds, 509–538 nanotechnology, 369–71, 596, xxiii nanotopographies, 511 National Institute for Clinical Excellence, 602 natural composites structure–property relationships in bone, cartilage, ligament and tendons, 3–19 implications for tissue regeneration and tissue repair, 18–19 Nicolais–Narkis equation, 427, 430 Nielsen model, 430 NiTi SMA, 128–9 nitrogen ion beam, 138 non-Hookean behaviour, 187, 237–8 non-resorbable biocomposites, 366 normal functioning model, 604 nylon 6, 433 Omniflow II, 86 Opsite, 80 OPTIMA LT1CA30, 490 organosiloxane, 449–50
orthotropic lamina, 238 osseointegration, 357 osteoblasts, 267 osteocalcin, 336 osteochondral scaffold, 562–4 ESEM micrograph of scaffold morphology, 563 in vitro and in vivo tests, 563–4 histological evaluation, 564 tissue stained with safranin O, 564 osteoconductivity, 133, 144, 226, 259, 356–7, 588 osteogenesis, 548 osteogenic materials, 356 osteoinduction, 356, 588 osteoinductive agents. see growth factors osteon see Haversian system Osterix, 567 oxygen, 217 Palacos, 224 paratenon, 236 Parientene composite mesh, 84 Parietex composite mesh, 84 Paris law equation, 469 particle leaching, 528–30 particulate composites, 106–9, 426–8, 430–1 theoretical models for tensile modulus and strength, 429 particulate scaffold in a matrix, 259–60, 262–3, 268–9 calcium phosphate ceramics with a hardening matrix ceramics in resorbable mineral matrix, 262 ceramics in resorbable organic matrix, 260, 262 non-hardening IBS: suspensions, 259 PEEK see poly(ether ether ketone) peptide amphiphile, 536 peripheral nerves, 88–90 phase separation, 30–1, 528–30 process diagram, 31
Index photolithography, 513–14 pirogenicity tests, 392–3 plasma immersion ion implantation, 160 and deposition, 160–1 plasma polymerisation, 112 plasma sprayed coatings for implants, 148–51 cross-section of microhardness variation, 150 HA-TCP composite coating structure, 149 plasma spraying, 134–7 plaster of Paris, 54 Plastipore, 49 platelet derived growth factor, 587–8 PLGA see poly(lactic acid-co-glycolic acid) ply see lamina poly(3-hydroxybutyrate-co-3hydroxyvalerate) films, 513 poly(2-hydroxyethyl methacrylate), 191–5, 243 poly(acrylic acid), 110 polyamide-6, 444–5 poly(aryl ether ketone ether ketone ketone), 188 poly(aryletherketones), 188 polycaprolactone, 33 polydimethylsiloxane, 514 poly(ε-caprolactone), 104 poly(ether ether ketone), 188–9, 281, 449, 454, 474–5 poly(ether)urethan–gold nanocomposites, 346 poly(ethylene terephthalate), 243 poly(glycolic acid), 104 polyhydroxybutyrate, 144 poly(hydroxybutyrate-cohydroxyvalerate), 144 poly(lactic acid), 104, 129, 164–6, 313, 376, 529 poly(lactic acid-co-glycolic acid), 104, 144 poly(lactide-co-glycolide), 248–9, 313, 336
623
polymer, 103 biocomposites for spinal implants, 187–95 composite interbody fusion devices, 188–90 composite intervertebral disc prostheses, 190–5 polymer/ceramic biocomposites, 367–9 ceramic composite materials, 368–9 degradation, mechanical characteristics and uses, 367 human mesenchymal stem cells in hydroxyapatite – β-tricalcium phosphate, 369 human osteoblast cells in hydroxyapatite – tricalcium phosphate, 370 polymer matrix composite materials, 367–8 polymer demixing, 516–17 polymer nanocomposites, 431–7 three-phase model representation, 434 polymerisation, 203–4, 215–16 MMA reaction sequence, 216 poly(methyl methacrylate), 50, 53, 182, 214–18, 280–1, 432, 485–6 bone cement setting, 214–15 MMA polymerisation reaction, 216 curing parameters, 215–17 typical parameters of some commercial bone cements, 217 residual monomer and shrinkage, 217–18 stress-strain characteristics, 219–22 typical components, 215 polypropylene, 445–7 poly(propylene fumarate), 162, 307 polypyrrole, 516 poly(vinyl alcohol), 376 porogen leaching, 165 porosity, 218–19, 510 porous hardening matrix, 263, 269 post-surgical adherence, 82–5
624
Index
potentiodynamic polarisation tests, 157–8 Primum non nocere, 597–8 Pro/ENGINEER 2000i, 308 PROCEED surgical mesh, 84 ProDisc, 185, 190 Proplast, 49 Protection of Experimental Animals, 386 protein adsorption, 332 proteoglycan, 64–5, 235, 300 proteoglycan network, 12 pyrolysis, 559 SEM micrographs of pyrolysed wood, 560 quartz, 204–5, 449–50 quasi-isotropic, 424 radiopacifiers, 222–3 rapid prototyping, 33–4, 303–4 technologies, 163 regenerative medicine developing biocomposites as scaffolds, 547–67 bio-hybrid composites for bonelike scaffold, 550–8 future trends, 565–7 hierarchical structure, 558–60 new approach, 549–50 three-layered osteochondral scaffold, 562–4 developing targeted biocomposites, 573–84 cell/material interactions, 576–8 specific processes, 584–8 angiogenesis, 587–8 inflammatory and immunological responses, 584–5 mechanotransduction, 585–6 reinforcing agents, 204–5 reproductive and development toxicity, 394 resorbable biocomposites, 366–7 Reuter’s matrix, 418 rheology, 265
rotation bending cantilever fatigue tests, 477 salt/particulate leaching, 117 sample preparation, 389 Saos-2 cells, 163, 165 scanning electron microscopy, 345–6 scanning tunneling microscopy, 347 secondary-ion mass spectroscopy, 347 selective laser melting, 36 selective laser sintering, 34, 165, 303 self-assembled monolayers, 582 self-reinforced PLLA plates, 47, 51–2 sensitisation test, 390–2 Seprafilm, 83–4 Sepramesh, 83 shear-extension coupling, 419, 422 shear thinning, 38, 266 short-fibre composites, 423–6 silane coupling agents, 205 silica, 205 silica–calcium phosphate nanocomposites, 342, 347 silicon, 554 Simplex P, 224 simulated body fluid, 107 single-walled carbon nanotube, 431–2 skeletal tissue, 364–5 SLA 250/40, 308 slippery slope, 605 S–N diagrams, 467–8 sodium hydroxide, 517 soft tissue repair composite biomaterials, 70–9 biocompetent nanocomposites, 78–9 engineering methods and biological interactions, 71–6 hybrid and biomimetic materialbased composites, 76–8 poly-(2-hydroxyethylmethacrylate) hydrogels, 75 use in clinical intervention, 80–90 burn treatment, 80–2 cardiovascular applications, 86–7 lower urinary tract, 87
Index other clinical applications, 87–90 post-surgical adherence, 82–5 wound dressing, 80 sol-gel methods, 534 solid free-form fabrication technologies, 33–4, 36–8 auger screw dispensing head, 37 BioScaffolder for 3D plotting, 36 general file principle for rapid prototyping apparatus, 34 PLA scaffold, 38 SFF technology in computerintegrated manufacturing principle, 35 sintering principle, 35 solution-based techniques, 30–3 electrospinning, 31–3 phase separation, 30–1 solvent casting, 30 solvent casting, 30 Spectra 1000 UHMWPE fibres, 494 spinal fusion, 180–1 spinal implants biocomposites, 178–97 basic concepts, 185–7 future trends, 195–7 materials and design, 181–5 polymer-based materials, 187–95 spinal structure and function, 179–81 spine surgery fatigue behaviour, 489–91 S–N curves of long-term implantable CF/PEEK (PEEK OPTIMA LT1CA30), 490 spinneret, 32 spraying and sintering for implants, 151–3 HA-glass FGC coating structure, 152 nanoindentation loaddisplacement curve, 153 stability, 264–5 staining assays, 340 stainless steel X5CrNi18–10, 450 Stamey’s procedure, 87
625
standard of care, 605–6 steric stabilization, 264 stochastic network modeling techniques, 306 Stokes’s law, 264–5 strain–stress relations, 415–16 strength, 8 stress shielding, 186, 363 strontium oxide, 223 structural biocompatibility, 356 Stryker-Dacron ligament prosthesis, 237 styrene acrylonitrile, 427 surface bioactivation, 582 surface biocompatibility, 356 surface fatigue tests, 477 surface hydroxyl groups chemical reaction, 112 surface modification, 109, 133–47 biomimetic deposition, 140–2 schematic diagram, 141 Ti plate morphology, 143 ion beam assisted deposition, 137–40 process diagram, 139 other techniques for implants, 142–3 apatite coating surface morphology, 144 plasma spraying, 134–7 forming coatings, 135 layered structure of HA coating, 136 spraying and sintering, 137 glass functionally graded coating, 138 tissue engineering scaffolds, 143–7 biomimetic deposition technique, 148 bone-like apatite coating, 147 surfactant adsorption, 109–10 synthetic bone graft substitutes, 362–3 systemic toxicity, 392–3 acute systemic toxicity, 392 pirogenicity tests, 392–3 subacute, subchronic and chronic systemic toxicity, 392–3
626
Index
T sling, 87 Takayanagi equation, 435 Takayanagi’s two-phase model, 433–4 target cell assays, 335 TCP see tricalcium phosphate Teflon, 522, 529 Tegaderm, 80 Tenascin-R, 69 tendons, 14–18 biocomposites fatigue behaviour, 491–4 hierarchial spatial arrangement, 15 and ligament case study on tissue-engineered scaffold, 313–15 morphology, 300 mechanical properties, 16–18 uniaxial tension stress–strain for anterior cruciate ligament and Achilles tendon, 17 TGF-β isoform 1, 64 TGF-β isoform 2, 64 thermoplastic materials, 27 thermoplastic polymer materials, 26 thermosets, 26 thixotropy, 265 three-dimensional braiding, 249 thrombin, 63 Ti–6Al–4V springs, 183 tissue engineering chemical aspects of materials, 581–4 material degradation, 581–2 release of biomolecules, 583–4 surface chemistry, 582–3 developing targeted biocomposites, 573–84 cell/material interactions, 576–8 injectible bone substitutes, 268–9 physical aspects of materials, 578–81 architecture, 580–1 mechanical properties, 578–9 surface topography, 579–80 and regenerative medicine, 509–13
tissue engineering scaffolds biomimetically deposited coatings, 163–7 apatite/collagen composite coating, 164 MTT assay, 166 Saos-2 cells cultured on apatite/ collage coated PLLA scaffolds, 165 composite coatings, 127–69 composite materials, 24250 fibre-reinforced composites properties, 296–316 case studies, 307–15 computer-aided tissue engineering, 306–7 design tools, 301–4 future trends, 315–16 harnessing directional properties of biomaterials, 297–9 load-bearing tissue morphology, 299–301 in silico computational analysis, 304–6 materials and manufacture, 161–3 nanostructured biocomposites, 509–538 bottom-up approaches, 536 complex microenvironment, 511 direct fabrication of surface nanotopographies in 3D structures, 516–19 future trends, 537–8 nanoparticles, nanotubes and nanofibres, 520–34 processing 2D topographies for 3D structures assembly, 513–14, 516 sol-gel, direct growth and biomimetic approaches, 534–6 other techniques, 167 tissue engineering solution, 565 titanium dioxide, 377–8 total hip arthroplasty, 279 total joint replacements, 212–13 trabecular bones, 277–8 Transcyte, 81–2
Index transmission electron microscopy, 347 tribology biocomposites, 441–60 experiment on composite characterisation, 443–4 advantages and disadvantages of current bearing choices, 447 common parameters, 444 wear conditions and type, 445 wear measurement techniques, 446 polymer composites, 444–60 fibre-reinforced polymer composites, 453–60 lubricating polymer, 444–8 metal, inorganic and polymeric powder and carbon nanotubes, 448–53 tribo-biomaterials, 442 tricalcium phosphate, 148–51 triethylene glycol dimethacrylates, 202–3 tropocollagen, 5 tubes-by-fibre templates, 526 UHMWPE see ultrahigh molecular weight polyethylene ultrahigh molecular weight polyethylene, 48–9, 53, 182, 281, 444–8, 474, 484–5 specific wear rates, 450 volume loss, 453 wear volume and wear coefficient, 451 ULTRAPRO, 83 uniaxial tension, 13, 16 urethane dimethacrylate, 202
627
vacuum infusion, 29 vacuum mixing, 225 vapour–liquid–solid based technique, 535 vascular endothelial growth factor, 587–8 vascular grafts, 86 biocomposites fatigue behaviour, 494–5 vascularisation, 361 Vectris-Pontic, 481 vertebroplasty, 223 Viasorb, 80 Vicryl, 263 viscoelasticity, 17–18, 246–7 viscoplasticity, 265 vitreous carbon, 50 Vroman effect, 576 VYPRO II, 83 Wallerian degeneration, 88 water absorption, 207 Weibull model, 487–8 Wöhler curve see S–N diagrams wollastonite ceramics, 226 wound dressing, 80 woven bone, 300 X-ray computed tomography, 347 X-ray diffraction, 154, 346–7 XAS/914 carbon/epoxy composites, 472 Young’s modulus, 108, 113, 242 Zeta potential analysis, 264 zirconia, 458 zirconium dioxide, 222
