Biomedical Composites: Materials, Manufacturing and Engineering 9783110267488, 9783110266696

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J. P. Davim (Ed.) Biomedical Composites

Also of Interest Nanocomposites Materials, Manufacturing and Engineering Davim, Charitidis (Eds.), 2013 ISBN 978-3-11-026644-3, e-ISBN 978-3-11-026742-6

Membrane Systems For Bioartificial Organs and Regenerative Medicine De Bartolo, Curcio, Drioli, 2014 ISBN 978-3-11-026798-3, e-ISBN 978-3-11-026801-0

Biomimetics A Molecular Perspective Jelinek, 2013 ISBN 978-3-11-028117-0, e-ISBN 978-3-11-028119-4

Nanotechnology Reviews Kumar, Challa (Editor-in-Chief) ISSN 2191-9089, e-ISSN 2191-9097

Science and Engineering of Composite Materials Hoa, Suong V. (Editor-in-Chief) ISSN 0792-1233, e-ISSN 2191-0359

Biomedical Composites Materials, Manufacturing and Engineering

Edited by J. Paulo Davim

Editor Professor J. Paulo Davim University of Aveiro Department of Mechanical Engineering Campus Santiago 3810-193 Aveiro Portugal [email protected]

ISBN 978-3-11-026669-6 e-ISBN 978-3-11-026748-8 ISSN 2192-8983 Library of Congress Cataloging-in-Publication Data A CIP catalog record for this book has been applied for at the Library of Congress. Bibliographic information published by the Deutsche Nationalbibliothek The Deutsche Nationalbibliothek lists this publication in the Deutsche Nationalbibliografie; detailed bibliographic data are available in the Internet at http://dnb.d-nb.de. © 2014 Walter de Gruyter GmbH & Co. KG, Berlin/Boston The publisher, together with the authors and editors, has taken great pains to ensure that all information presented in this work (programs, applications, amounts, dosages, etc.) reflects the standard of knowledge at the time of publication. Despite careful manuscript preparation and proof correction, errors can nevertheless occur. Authors, editors and publisher disclaim all responsibility and for any errors or omissions or liability for the results obtained from use of the information, or parts thereof, contained in this work. The citation of registered names, trade names, trademarks, etc. in this work does not imply, even in the absence of a specific statement, that such names are exempt from laws and regulations protecting trademarks etc. and therefore free for general use. Typesetting: PTP-Berlin Protago-TEX-Production GmbH, Berlin Printing and binding: Hubert & Co. GmbH und Co. KG, Göttingen Cover image: gettyimages/thinkstockphotos, Abalone Shell ♾ Printed on acid-free paper Printed in Germany www.degruyter.com

Preface Recently, biomedical composites have contributed to great developments in applications in the modern medical industry. With the current application of implants in the human body, it is clear that the implants should be compatible with the tissues. Therefore, biomedical composites are particularly attractive because their properties are comparable to those of the tissues. Innovations in biomedical composites, design and fabrication processes are increasing the performance of the tissues as well as other medical applications. Biomedical composites are widely used to repair, for example, tooth, bone, cartilage skin. The present volume aims to provide recent information on biomedical composites (materials, manufacturing and engineering) in eight chapters. Chapter 1 of the book provides information on ceramic polymer composites for hard tissue applications. Chapter 2 is dedicated to HAp-metal based biocomposite coatings and characteristics of plasma deposited HAp-Ti/Ti6Al4V coatings. Chapter 3 describes hydrogels based on poly(vinylalcohol) for cartilage replacement. Chapter 4 contains information on polymer composites for cemented total hip replacements. Chapter 5 describes bioresorbable composites for bone repair. Chapter 6 describes bioactive glasses and glassceramics. Chapter 7 contains information on metal oxide-based one-dimensional titania nanostructures via electrospinning (characterization and antimicrobial applications). Finally, Chapter 8 is dedicated to hydrogels for biomedical applications. The present volume can be used as a research book for final-year undergraduate engineering courses or as a topic on biomedical composites at the postgraduate level. Also, this book can serve as a useful reference for academics, researchers, materials, mechanical and biomedical engineers, professionals in medical technology and related industries. The scientific interest in this book is evident for many important centers of research, laboratories and universities as well as the biomedical industry. Therefore, I hope that this book will inspire and enthuse others to undertake research in this field of biomedical composites. The Editor acknowledges De Gruyter for this opportunity and for their enthusiastic and professional support. Finally, I would like to thank all the chapter authors for their availability for this work. September 2013

J. Paulo Davim Aveiro, Portugal

Contents Preface | V List of Contributing Authors | XI

1 1.1 1.2 1.3 1.4 1.5 1.6 1.7

2 2.1 2.2 2.2.1 2.2.2 2.3 2.4 2.4.1 2.4.2 2.4.3 2.5 2.5.1 2.5.2 2.5.3 2.6 2.6.1 2.6.2 2.6.3 2.7

Sunita Prem Victor and Chandra P. Sharma Ceramic polymer composites for hard tissue applications | 1 Introduction | 1 Polyethylene based composites | 3 Polymethymethacrylate based composites | 6 Polyester based composites | 7 Chitosan based composites | 10 Future Scope | 11 Conclusion | 12 References | 12 Xuan Zhou and Ramesh K. Guduru HAp-metal based biocomposite coatings and characteristics of plasmadeposited HAp-Ti/Ti6Al4V coatings | 17 Introduction | 17 HAp-Ti/Ti6Al4V based composites | 18 Hydroxyapatite (HAp) | 18 Titanium and its alloys | 19 Plasma Spray of HAp-Ti/Ti6Al4V based composites | 20 Property requirement of biocomposites | 21 Mechanical properties | 22 Biocompatibility | 22 Bioactivity | 23 Property evaluation | 23 Bond strength | 23 Corrosion behavior evaluation | 24 Immersion test in simulated body fluid | 24 Plasma sprayed HAp-(Ti/Ti6Al4V) based composite coatings | 25 Bond strength of plasma-sprayed HAp-(Ti/Ti6Al4V) based composite coatings | 25 Electrochemical corrosion behavior of plasma-sprayed HAp-(Ti/Ti6Al4V) based composite coatings | 27 Immersion behavior of plasma sprayed HAp-(Ti/Ti6Al4V) based composite coatings | 27 Conclusions | 29 References | 29

VIII   

3 3.1 3.2 3.3 3.4 3.5 3.6 3.7 3.8 3.9 3.10

4 4.1 4.1.1 4.1.2 4.2 4.3

   Contents

Julieta Volpe, Lucía M. Masi, Vera A. Alvarez and Jimena S. Gonzalez Hydrogels based on poly(vinylalcohol) for cartilage replacement | 33 Hydrogels: General Ideas | 33 Main properties of hydrogels | 34 Hydrogels as biomaterials | 37 Polyvinyl alcohol (PVA) hydrogels: General characteristics | 38 PVA hydrogels for biomedical applications | 40 Cartilage: A brief description | 41 Articular cartilage: Architecture and composition | 41 Articular cartilage: Mechanical properties | 43 Frequent medical issues relating to cartilage: Degeneration and osteoarthritis | 44 Materials used as articular replacement | 44 Conclusions | 46 Acknowledgments | 46 References | 46 S. Arun, P. S. Rama Sreekanth and S.Kanagaraj Polymer composites for cemented total hip replacements | 53 Introduction | 53 Understanding hip joint prosthesis and fixation techniques | 53 Economic and clinical factors surrounding revision surgeries | 56 UHMWPE composites | 57 PMMA composites | 60 Summary | 63 Future scope | 63 References | 64

Sandra Pina and José M.F. Ferreira 5 Bioresorbable composites for bone repair | 69 5.1 Introduction | 69 5.2 Bioresorbable materials | 73 5.2.1 Polymers | 73 5.2.1.1 Polyglycolic acid – PGA | 73 5.2.1.2 Polylactic acid – PLA | 74 5.2.1.3 PGA-PLA copolymers | 76 5.2.1.4 Poly ε-caprolactone – PCL | 76 5.2.2 Bioactive ceramics | 77 5.3 Composites manufacturing methods | 78 5.4 Clinical applications of bioresorbable composites for bone repair | 79 5.5 Conclusions | 80 References | 81

Contents   

6 6.1 6.2 6.2.1 6.2.2 6.3 6.3.1 6.3.2 6.3.3 6.4 6.5 6.6 6.6.1 6.6.2 6.7 6.8 6.9

7 7.1 7.2 7.3 7.4 7.5 7.6 7.7 7.8 7.8.1 7.8.2 7.8.3 7.8.4 7.8.5 7.9

   IX

G. El-Damrawi and H. Doweidar Bioactive glasses and glass-ceramics | 89 Biodental metals, ceramics and bioactive glass-ceramics; historical background | 89 Metallic implant materials | 89 Gold alloys | 90 Dental amalgam | 90 Glass-ceramics and bioactive glass-ceramics | 90 Commercial glass-ceramic products | 91 Protective glass-ceramic | 91 Bioceramics | 92 Preparation techniques | 92 Structure of glass-ceramics | 94 Crystallinity enhancement | 96 By adding activator agents | 96 By sintering process | 98 Dental glass-ceramics | 98 Bioactive glass-ceramics | 99 In vitro and in vivo test for bioactivity | 101 References | 104 M. Shamshi Hassan, Touseef Amna, Mohamed Bououdina and Myung-Seob Khil Metal oxide-based one-dimensional titania nanostructures via electrospinning: Characterization and antimicrobial applications | 107 Introduction | 107 General routes/procedures for the synthesis of nanofibers | 109 Electrospinning process | 109 General applications of electrospun nanofibers | 111 Antimicrobial applications of metal oxide-based nanotextured materials/nanofibers | 112 Concept of doping and composite nanofibers | 113 Development of pristine TiO2 nanofibers via electrospinning technique | 114 Doping of titania with metal oxide | 117 Doping of titania with zinc | 117 Doping of titania with copper | 121 Doping of titania with nickel | 124 Doping of titania with cobalt | 126 Doping of titania with cerium | 128 Plausible antibacterial mechanism of TiO2 / doped-TiO2 nanostructures | 131

X   

7.10

   Contents

Concluding remarks | 133 Acknowledgment | 134 References | 134

Luisa Russo, Sabrina Zaccaria, Maria Assunta Autiello and Assunta Borzacchiello 8 Hydrogels for biomedical applications | 141 8.1 Hydrogels: Classification and basic structure | 141 8.1.1 In situ forming hydrogels | 143 Physical crosslinking methods | 143 Covalent crosslinking strategies for forming hydrogels in situ | 146 8.2 Structure-properties relationship | 147 8.2.1 Hydrogel mechanical properties | 147 Hydrogels’ time dependent properties | 147 Stress strain behavior | 149 8.2.2 Hydrogel swelling | 150 8.3 Biomedical applications | 152 8.3.1 Tissue engineering | 152 8.3.2 Drug delivery | 155 8.3.2.1 Design criteria for hydrogels in drug delivery | 156 Incorporation of drugs | 157 8.3.2.2 Drugs release from hydrogels formulations | 158 Dynamic hydrogels | 159 Composite hydrogels | 160 Micro-nanoscale hydrogels | 160 In situ forming hydrogel | 161 References | 162 Index | 169

List of Contributing Authors Vera A. Alvarez Composite Materials Group (CoMP) Research Institute of Material Science and Technology (INTEMA) CONICET National University of Mar del Plata (UNMdP) Solís 7575 7600 Mar del Plata Argentina Touseef Amna Department of Animal Science Institute of Rare Earth for Biological Applications Chonbuk National University Jeonju 561-756 South Korea Srinivasan Arun Department of Mechanical Engineering Indian Institute of Technology Guwahati Guwahati-781 039 Assam India [email protected] Maria Assunta Autiello Department of Chemical, Materials and Production Engineering University of Naples Federico II P.le Tecchio 80 80125 Napoli Italy

Assunta Borzacchiello Institute of Composite and Biomedical Materials-(IMCB-C.N.R) Viale Kennedy, 54 - Mostra d‘Oltremare Pad. 20 80125 Napoli Italy [email protected] Mohamed Bououdina Nanotechnology Centre and Department of Physics, College of

Science University of Bahrain PO Box 32038 Kingdom of Bahrain [email protected] Gomaa El-Damrawi Physics Department Faculty of Science Mansoura University Mansoura 35516 Egypt Hamdy Doweidar Physics Department Faculty of Science Mansoura University Mansoura 35516 Egypt [email protected]

XII   

   List of Contributing Authors

José M.F. Ferreira Department of Materials Engineering and Ceramics CICECO University of Aveiro 3810-193 Aveiro Portugal [email protected] Jimena S. Gonzalez Composite Materials Group (CoMP) Research Institute of Material Science and Technology (INTEMA) CONICET National University of Mar del Plata (UNMdP) Solís 7575 7600 Mar del Plata Argentina [email protected]

Ramesh K. Guduru Department of Mechanical Engineering University of Michigan Dearborn, MI – 48128 USA M. Shamshi Hassan Department of Organic Materials and Fiber Engineering Chonbuk National University Jeonju 561-756 South Korea Subramani Kanagaraj Department of Mechanical Engineering Indian Institute of Technology Guwahati Guwahati-781 039 Assam India [email protected]

Myung-Seob Khil Department of Organic Materials and Fiber Engineering Chonbuk National University Jeonju 561-756 South Korea Lucía M. Masi Composite Materials Group (CoMP) Research Institute of Material Science and Technology (INTEMA) CONICET National University of Mar del Plata (UNMdP) Solís 7575 7600 Mar del Plata Argentina Sandra Pina Department of Materials Engineering and Ceramics CICECO University of Aveiro 3810-193 Aveiro Portugal [email protected] Luisa Russo University of Naples Federico II P.le Tecchio 80 80125 Napoli Italy and Institute of Composite and Biomedical Materials-(IMCB-C.N.R) Viale Kennedy, 54 - Mostra d‘Oltremare Pad. 20 80125 Napoli Italy

List of Contributing Authors   

Chandra P. Sharma Division of Biosurface technology Biomedical Technology Wing Sree Chitra Tirunal Institute for Medical Science and Technology Poojappura, Thiruvananthapuram-695012 Kerala India [email protected] Pattela S. Rama Sreekanth Department of Mechanical Engineering National Institute of Science and Technology Palur Hills Berhampur-761 008 Orissa, India [email protected] Sunita Prem Victor Division of Biosurface technology Biomedical Technology Wing Sree Chitra Tirunal Institute for Medical Science and Technology Poojappura, Thiruvananthapuram-695012 Kerala India

   XIII

Julieta Volpe Composite Materials Group (CoMP) Research Institute of Material Science and Technology (INTEMA) CONICET National University of Mar del Plata (UNMdP) Solís 7575 7600 Mar del Plata Argentina Sabrina Zaccaria Institute of Composite and Biomedical Materials-(IMCB-C.N.R) Viale Kennedy, 54 - Mostra d‘Oltremare Pad. 20 80125 Napoli Italy Xuan Zhou Department of Electrical and Computer Engineering Kettering University Flint, MI – 48504 USA [email protected]

Sunita Prem Victor and Chandra P. Sharma

1 Ceramic polymer composites for hard tissue applications 1.1 Introduction Bone repair or regeneration is a common and complicated clinical problem in orthopedic and dental surgeries all over the world. Every year, millions of people suffer from bone defects generated by tumor resection, trauma, bone diseases and congenital abnormalities. A recent report on the world orthopedic implants and products industry noted that the total orthopedic drug, implant and device market is estimated to be to be approximately 18 billion dollars per year. Many patients who receive an orthopedic implant may have to undergo revision surgery in their lifetime. To decrease patient discomfort and costs, prostheses with improved clinical efficacy and longer effective life times are being studied. Synthetic bone replacements and prostheses are mainly comprised of polymers, metals and ceramics. It is well known that bone, in response to the surrounding mechanical stimuli, adapts its anatomical structure through natural growth and resorption processes [1]. Therefore, by virtue of their higher stiffness than bone tissue, metal-based implants results in stress shielding making it susceptible to fracture [2]. To avoid mechanically-induced bone resorption, scientists have striven to develop biomaterials partially mimicking the biphasic composition of bone and its structure. In general, bone is a natural composite material, which by weight contains about 45–60 % mineral, 20–30 % matrix, and 10–20 % water. Nanoscopically, the constitutional building blocks of bone are mineralized collagen fibrils of 80–100 nm thickness and a length of few microns [3]. The matrix is the organic component, which is primarily composed of the protein Type I collagen which is a triple helix that is highly aligned, yielding a very anisotropic structure. The major mineral constituent of bone is a substituted calcium phosphate similar in composition and structure to hydroxyapatite (HA, Ca10(PO4)6(OH)2). The organic matrix provides the flexibility and the inorganic mineral is predominantly responsible for the mechanical properties of bone [4]. Calcium phosphate based ceramics present a unique class of materials indispensable in biomaterial applications due to their chemical similarity to the mineral component of mammalian bones and teeth. Their bioactive property has received most attention in the field of hard tissue replacement. They provide fixation by biological ingrowth of the local tissue through the formation of a biologically active hydroxycarbonate apatite layer on their surfaces in vivo. They also have the additional benefits of biocompatibility; corrosion and high compression resistance respectively [5]. Calcium phosphates with their varying Ca/P ratios are given in Tab. 1.1. However, among the various calcium orthophosphates only certain compounds can be utilized

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for biomedical applications, because those having a Ca/P ionic ratio less than 1 are not suitable for implantation into the body due to their high solubility and acidity. The stable phases of calcium phosphates depend considerably upon temperature and the presence of water, either during processing or in the environment of their use. At body temperature, only two calcium phosphates are stable in contact with aqueous media, such as body fluids: at pH  4.2, the stable phase is Ca10(PO4)6(OH)2(Hydroxyapatite, HA). Since body fluids are at pH 7.4 and 37 °C, HA is the most stable calcium phosphate phase at that condition [6]. Tab. 1.1: Ca/P ratio of various calcium phosphates

Ca/P Ratio

Name

Formula

Acronym

2.0 1.67 1.50 1.33 1.0 1.0 1.0 1.0 0.7 0.67 0.5 0.5

Tetracalcium phosphate Hydroxyapatite Tricalcium phosphate (ά,β,γ) Octacalcium phosphate Dicalcium phosphate dihydrate Dicalcium phosphate Calcium pyrophosphate (ά,β,γ) Calcium pyrophosphate dihydrate Heptacalcium phosphate Tetracalcium dihydrogen phosphate Monocalcium phosphate monohydrate Calcium metaphosphate (ά,β,γ)

Ca4O(PO4)2 Ca10(PO4)6(OH)2 Ca3(PO4)2 Ca8H2(PO4)65H2O CaHPO4.2H2O CaHPO4 Ca2P2O7 Ca2P2O72H2O Ca7(P5O16)2 Ca4H2P6O20 Ca(H2PO4)2H2O Ca(PO3)2

TTCP HA TCP OCP DCPD DCPA CPP CPPD HCP TDHP MCPM CMP

Bone defects have been clinically treated by the implantation of prefabricated calcium phosphate blocks and granules. The commonly used ceramic bone substitute materials are hydroxyapatite (HA, Ca10(PO4)6(OH)2) and tricalcium phosphate (TCP, Ca3(PO4)2) which have different characteristics in vivo, although both forms have Ca/P ratios within the range known to promote bone ingrowth (1.50–1.67). The HA is known to bond with bone directly and can thus be used as a bone replacing material while the TCP (both α and β phases) is known to be a bone substituting material because it dissolves gradually and new bone will be formed where it is resorbed [7]. The calciumto-phosphate ratio (Ca/P) of HA varies approximately between 1.5 and 1.67. The pH range is 4.6 to 12.4, over which this variable composition is stable at 25 °C. However, ceramics have certain drawbacks like difficulty to fabricate, low mechanical reliability, lack of resilience, and high density. Synthetic or natural polymers have viscoelastic characteristics and may be used alone in sites where mild mechanical stresses exist, such as in soft tissues like cartilage or tendons. The requirements for a polymer material to be used for biomedical applications include fatigue resistance, resistance to ageing in saline aqueous media, biocompatibility, dimensional stability, absence of migrating harmful addi-

1.2 Polyethylene based composites   

   3

tives and being sterilizable by standard methods without loss of properties. Polymers by themselves are generally flexible and exhibit a lack of mechanical strength and stiffness, whereas inorganic materials such as ceramics and glasses are known to be stiff and brittle. The combination of polymers and inorganic phases leads to composite materials with improved mechanical properties due to the inherent higher stiffness and strength of the inorganic material. So in applications where stiffness is required together with damping abilities such as bone, filling polymers with ceramic particulates may be an interesting solution. Scientists have replicated in vitro the process of collagen mineralization and obtained materials chemically and hierarchically mimicking bone tissue [8], while others have synthesized fibrous silk macroporous blocks containing calcium phosphate [9]. Micro- and nano-HA has been added to synthetic polymers generating materials not only with mechanical characteristics comparable to those of bone [10] but also capable to act as framework for bone formation. Further, the addition of bioactive phases to bioresorbable polymers can alter the polymer degradation behavior of the composite scaffolds. Literature reports suggest that highly porous ceramic polymer composite scaffolding appears to be a promising substrate for bone tissue engineering due to its excellent mechanical properties and osteoconductivity [11]. These composite scaffolds supported uniform cell seeding, cell ingrowth, and tissue formation. So these composite scaffolds comprising ceramics like HA and TCP with polyethylene (PE), Polymethylmethacrylate (PMMA), Polypropylene fumarate (PPF) and chitosan are being widely used for bone augmentation and repair. The development of composite materials for biomaterial applications is thus attractive since their properties can be engineered to suit the mechanical and physiological demands of the host tissue by controlling the volume fraction, morphology, and arrangement of the reinforcing phase. This chapter discusses the chemistry of selected calcium phosphate polymer composite systems, their preparation techniques and their physical and biological properties. It also includes their current biomedical applications like bone fillers, coatings, cements, scaffolds and drug carriers.

1.2 Polyethylene based composites Hydroxyapatite (HA) reinforced high density polyethylene (HDPE) composite (HAPEXTM) has been developed since the early 1980s as an analogue for bone replacement and a load bearing material in joint endoprostheses [12]. Bonfield et al. [13–15] introduced the bone analogue concept and proposed the use of composites of HDPE with hydroxyapatite (HA). The efforts to obtain mechanical properties similar to bone were based on the reinforcement of a ductile polymeric matrix with a bonelike ceramic. The ceramic ensures the mechanical reinforcement of the polymer along with its bioactive character and biocompatibility. The presence of the ceramic further led to improvements in the creep resistance of the material. The advantages offered

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   1 Ceramic polymer composites for hard tissue applications

by UHMWPE include [16] very good sliding properties, good impact strength, good fatigue resistance and good biocompatibility. It has been shown that an optimum combination of mechanical and biological properties has been achieved with HAPEXTM containing 40 % by volume of HA [17]. This composite has a modulus value similar to that of cortical bone, has better toughness and considerable bioactivity. The obtained modulus value shows promise in solving the problem of bone resorption occurring to implants of conventional materials which have much higher modulus values than bone. Implants composed of HAPEXTM composite encouraged bone apposition which is favorable from a compatibility point of view rather than the fibrous encapsulation which was encountered with other implant materials [18]. Mechanical coupling of the reinforcement and the matrix, resulting from shrinkage of the HDPE matrix around the HA particles during composite processing, has produced modulus and strength which are adequate for minor load-bearing applications and have resulted in its successful clinical use. The production methods play a very important role in controlling the mechanical properties of composites. HDPE composites can be processed by compounding and extrusion when compared to UHMWPE which can only be compression molded from powder. The hydrostatic extrusion of HAPEXTM has indicated that composite with high modulus and strength similar to those of cortical bone can be manufactured for major load-bearing skeletal implants [19]. Another study [20] reports the use of Twin screw compounding extrusion generating high shear forces to add HA into HDPE. Around 50 vol% HA could be compounded into the HDPE. Tensile testing showed that the increase in HA content led to increases in the Young modulus from around 1.3 GPa for the unfilled PE to 7.7 GPa for 50 vol%. However, at HA content > 40 % there was a reduction in tensile strength which resulted in brittle fracture. This is in contrast to human bone which has a mineral content > 45 % and mineral particles in the nanometer range. Investigators went on to study the effect of altering HA particle size, morphology and homogenous distribution on the tensile strength of the composite. Two grades of spray-dried HA with mean particle sizes of 4 and 8 micron respectively having slight variation in specific surface area were studied to monitor the differences in tensile strength [21]. Minor differences in the tensile strength were observed for the different filler particles and the filler contents investigated. However, the tensile and torsional stiffness were higher with the smaller sized HA particles resulting in stiffer composites. Furthermore, the stiffness of HDPE/HA composites is proportional to the volume fraction of HA. Nevertheless, though the HA particles increase the material stiffness and enhance the creep behavior, the higher the HA content, the higher the number of interfaces between the polymer and the ceramic. The number of interfaces has to be taken into account since failure can preferentially occur at the interface when the implant is under mechanical loading [22]. Another study by Joseph et al. [23] investigated the effect of spray-dried and sintered HA powder with similar particle sizes but differences in specific surface area. The composite was compound extruded and its rheological properties characterized. The composite manufactured

1.2 Polyethylene based composites   

   5

with spray-dried HA had substantially higher viscosity than that using sintered HA. The difference was attributed to the effect of specific surface area which correlates with the amount of free PE available to act as a matrix between the coated particles [24]. Studies were then carried out to improve the impact properties of the composite. Composites with 20, 30 and 40 vol% fillers were manufactured and subjected to impact testing. For all volume fractions, both the initiation and propagation energies were higher with the sintered HA with a lower surface area. Several studies by Bonfield and co-workers [25, 26] have indicated that the low efficiency of the HA particles as reinforcement agents is due to their inherent low aspect ratio and low degree of chemical interaction with the HDPE phase. There have been attempts to enhance the mechanical performance by coupling with silane agents and grafting with acrylic acid to increase the adhesion of the HA particles towards the polymeric matrix. Silane coupling treatments depend on factors like particle surface area, particle size distribution and chemical reactivity of HA particles. Coupling reactions have also been carried out using titanate and zirconate coupling agents which led to the conclusion that the positive effect of these agents on the stiffness and strength of the composite results from their dominant effect as HA dispersion promoters. Solid state processing techniques like hydrostatic extrusion to develop bone matching mechanical performance (modulus up to 13 GPa) has helped in the attainment of significant improvements in composite stiffness [27]. A complimentary method carried out by reinforcing HDPE/HE composite with high modulus HDPE fibers resulted in modulus values up to 17 GPa. Shear controlled orientation injection molding was successful in inducing anisotropic character to high density PE and in the respective composites reinforced with HA. These studies demonstrated the relative importance of the molecular weight characteristics of the HDPE on the attainment of high anisotropic bone-analogue composites. Recently porous scaffolds of HA/PE composites have been prepared by selective laser sintering [28]. Twin screw extrusion was used to process the composites which were ground down to desired particle size and then were selectively laser sintered. Scaffolds of 20 vol% HA/PE and 30 vol% HA/PA could be sintered successfully by optimizing particle size and laser power. Scanning electron microscopy revealed that the material was all open celled which is essential for tissue ingrowth. The first PE/HA composite to be assessed for bioactivity was the HAPEX composite comprising PE with 40 vol% HA. The initial biological assessment of HAPEX performed in vivo [29] showed that the material was compatible. In vitro tests carried out revealed that the cells proliferated faster on 40 vol% HA/PE than on 20 vol% HA/PE. Confocal laser microscopy showed actin stress fibers in the cells that could be seen with vinculin staining adhesion plaques, although the number of adhesion plaques and actin fibers were higher and more organized on the 40 vol% HA/PE than on the 20 vol% HA/PE. Transmission electron microscopy (TEM) of the cell culture showed that more collagen was formed between the cells on the higher HA content material. HAPEX was implanted into patients and used by Downes et al. [30] in orbital floor

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   1 Ceramic polymer composites for hard tissue applications

applications and results demonstrated that they encouraged the supporting bone to bond to the implant material. Another study by Meijer et al. [31] further reported the clinical application of these implants. Scanning electron microscopy was used to review the response to 11 HAPEX implants which had been removed owing to recurrence of the original clinical problem, rather than for device failure. In most of the cases, the implants were covered with fibrous tissue, and in half of these there was a thin epithelial outer layer. The response to implantation was considered to be good even after 30 months of implantation.

1.3 Polymethymethacrylate based composites Polymethylmethacrylate (PMMA) has been employed as the best-known polymeric material for bone reconstructions and as acetabular cups for hip prostheses, for over 40 years. Polymethylmethacrylate (PMMA) was first used in paste form for the fixation of prosthesis as the material can be quickly adjusted to the contours of the defects due to its plasticity and its rapid liquid to solid phase transition [32]. Commercially available bone cements are made of (a) a powder component consisting of PMMA polymer or related block copolymers and (b) a liquid component made of methylmethacrylate (MMA) or related monomer liquids. Further, the commercially available PMMA cements generally contain a benzoyl peroxide, barium sulfate or zirconium dioxide, in addition to oligomers of PMMA [33]. Before being utilized for clinical procedure, bone cement is fabricated by combining the powder and the liquid component together. In fact the powder component dissolves in the monomer phase and, thereafter, it autopolymerizes within 10–15 min, resulting first in moldable viscous dough within the first few minutes, and then very dense cured PMMA-based bone cement. Structurally, PMMA-based bone cements are formed of typical linear polymers. However, semi interpenetrating network structures are also possible by modification with crosslinkers like ethylene glycol dimethacrylate. However, the exothermic setting reaction of more than 70° C, the toxic effects of the monomer residues and its complete inability to bond to bone directly has led to severe clinical complications [34]. Further, in joint replacement surgery, the main function of PMMA-based cement is to transfer body weight in order to increase the load-bearing and fixation capacity of the reconstruction area. On the other hand, PMMA-based cement is morphologically very dense, thus not permitting bone ingrowth. Additionally, as PMMA degrades slowly over the years, small particles of the implant leads to pronounced macrophage activation. This leads to inflammation, activation of osteoclasts and aseptic loss of the surrounding bone [35]. However, in most cases, PMMA-based cements have been successfully utilized in orthopedic surgery for many decades. Therefore, interest has arisen in developing more biocompatible PMMA-based cements. One of the options studied was combining ceramic with polymer to prepare composite methacrylate cement with improved bioactivity [36]. Basically, it was proposed

1.4 Polyester based composites    

   7

to make the PMMA-based bone cements more tissue-friendly and bioactive in nature. Bioactive bone cements are normally successfully obtained when the PMMA matrix contains an appropriate concentration of bioactive ceramics. It has been reported that bioactive PMMA cements containing hydroxyapatite have significantly higher bonebonding strength than plain PMMA cement [37]. However, the presence of fairy large concentrations of ceramics in PMMA can lead to detrimental mechanical properties, especially if there is poor adhesion between the PMMA matrix and the ceramic filler. Harper [38] reinforced PMMA bone cement with 17 wt% HA in an attempt to confer bioactivity without reducing the mechanical properties. The cellular response of this composite was further evaluated and it was noted that the addition of HA increased the cellular response to the PMMA with higher actin organization observed in the cells grown on HA containing PMMA. So, incorporation of many types of bioactive fillers such as hydroxyapatite (HA), β-tricalcium phosphate (β-TCP) [39], titania [40], bioactive glass [41] and bioactive glass-ceramics [42] has been carried out with PMMA as matrix. These bioactive composites are proven to be advantageous in producing an intermediate region between the bone and prosthesis through the formation of an apatite layer on their surface [43]. A study by Ishhara et al. [44] demonstrated that 4-methacryloyloxyethyl trimellitic anhydride (4-META) containing MMA bone cement could adhere to HA with improved tensile bond strength. The adhesion was explained by the formation of 4-methacryloyloxyethyl trimellitic acid (4-MET) which gets firmly absorbed onto the surface of HA followed by copolymerization with MMA monomer. Misra et al. [45] found that when the HA surface had been surface treated with zirconyl methacrylate, the diametral tensile strength of dental composites was increased by 50 %. The stiffness of PMMA-based composites was further enhanced by covalent bonding with hydroxyapatite. Introduction of covalently bonded HEMA to nonstoichiometric hydroxyapatite was carried out by co-precipitation of apatitic octacalcium phosphate (AOCP) in the presence of hydroxyethylmethacrylate phosphate [46]. The phosphoHEMA apatite obtained was then used to copolymerize with either HEMA or methyl methacrylate (MMA) to form chemical bonds between the mineral filler and the polymer matrix. This chemical interaction resulted in better mechanical properties. A recent micro-CT analysis [47] study on the quantitative bone healing parameters at bone implant interface using PMMA containing mixtures of calcium phosphate revealed complimentary morphological information and histology and clear formation of newly formed bone.

1.4 Polyester based composites Among various materials investigated for bone applications and treating skeletal defects, polyester based composites have been extensively studied. Poly(propylene fumarate) (PPF) is an unsaturated linear polyester, whose degradation occurs by means of hydrolysis with the generation of propylene glycol and fumaric acid as end

8   

   1 Ceramic polymer composites for hard tissue applications

products. This biodegradable polymer has many unsaturated double bonds capable of crosslinking and has been shown to be biocompatible, noncytotoxic and osteoconductive. Jayabalan et al. [48] studied (Poly(propylene fumarate-co-ethylene glycol) copolymer/hydroxyapatite as an injectable bone cement for the tissue engineering of bones. The group further went on to study the use of polypropylene fumarate/ phloroglucinol triglycidyl methacrylate blend/hydroxyapatite as partially degradable polymeric cement for orthopedic applications. They reported that the blend having 50:50 ratios of these oligomers and hydroxyapatite undergoes fast setting and binds bones with appreciable tensile strength. The in vitro biodegradation studies reveal slow degradation. Previous studies have also shown that solid PPF has mechanical characteristics similar to those of bone and ideal compressive strength for orthopedic applications. When TCP is added to the PPF the resulting composite cement has mechanical properties similar to those of trabecular bone. Further, the heat generated in the crosslinking reaction in PPF is lower than that obtained in commercialized bone cement, PMMA. He et al. [49] studied the effect of double-bond ratio and TCP content on crosslinking and mechanical characteristics of PPF-based injectable composites. The increase in double-bond ratio and the addition of TCP resulted in an increase of mechanical properties. In vitro studies in formulations including a porogen agent showed that an increase in PPF molecular weight increased the mechanical properties and that the formulations maintained the minimum requirement for replacement of human trabecular bone during 7 weeks [50]. Gerhart et al. [51] compared the antibiotic release and mechanical properties of a biodegradable bone cement made of PPF and MMA with a commercially available MMA cement. Higher and sustained antibiotic levels for a longer duration were obtained for the bone cement made of PPF and MMA. However, although the initial mechanical properties of the cements were good, they decreased by a factor of 12 just 4 weeks after implantation. To overcome the problems associated with mechanical properties they developed a composite with TCP and CaCO3 as fillers, which hardens in 24–36 h, had properties much higher than human trabecular bone and was workable enough to be packed into defects of complex shape. They found that the compressive strength and resistance to degradation increased with an increase in MMA concentration. The dissolution of calcium from the composite helped to maintain pH and promote bone remodeling. In another study, PPF and calcium phosphate cement (CPC) were combined to provide appropriate mechanical strength for core-decompressed femoral heads and offer the properties of osteoconductivity. The effects of different ratios of CPC to PPF on mechanical strength and cytotoxicity were investigated. The results demonstrate that the bone cement is less cytotoxic and the increment of the CPC proportion increases the mechanical strength, reduces the crosslinking temperature and diminishes excessive swelling of the cement. An angiogenic effect was also observed with addition of ginsenoside Rg1 and the angiogenecity of released Rg1 was confirmed by the assay of tube formation in human umbilical vein endothelial cells. This newly developed angiogenic bone cement composite possesses remarkable

1.4 Polyester based composites    

   9

development potential for application to treating osteonecrosis of the femoral head. Zhongyu et al. [52] prepared and evaluated the biodegradable behavior of PPF/CaSO4/ β-TCP composite designed for bone tissue engineering. The aim of this study was to investigate the effects of PPF molecular weight and CaSO4/β-TCP molar ratio on the in vivo degradation of PPF/(CaSO4/β-TCP) composite and the bone tissue response to PPF/(CaSO4/β-TCP). A total of 36 composite samples were implanted into segmental defects, harvested at 2, 4 and 8 weeks after the operation, and analyzed using radiographic and histological analysis to assess the in vivo degradation of the composites as well as tissue response to the implants. The in vivo degradation results show that all the samples maintained their original shape. There was tissues penetration in the pores formed by the degradation of CaSO4/β-TCP spheres near the surface of the composites. The rate of in vivo degradation and pore forming increased with a decrease in PPF molecular weight and an increase in CaSO4/β-TCP molar ratio. The composites were found to be capable of in situ pore formation with pore forming rate varying with the composition of the composite. No inflammatory reaction was observed after implantation which indicated that PPF/ (CaSO4/β-TCP) create a promising osteogenic scaffold for its controllable degradation rate and excellent biocompatibility. The effect of HA on the performance of nanocomposites of an unsaturated polyester, i.e., hydroxy-terminated high molecular weight poly(proplyene fumarate), was investigated [53]. Biodegradable hydroxyl terminated (HT-PPF) nanocomposites were prepared with different HA nanoparticles, i.e., (i) calcined HA nanoparticles ( DCPD > DCPA > OCP > α-TCP > β-TCP >> HAp. The solubility of amorphous calcium phosphate (ACP) cannot be measured precisely, but usually it is reported to be between OCP and HAP [17–18]. The mechanical properties of bulk HAp were also listed in Tab. 2.2 for comparison. Tab. 2.2: Common calcium phosphate compounds [11, 14, 15]

Symbol

Name

Formula

Ca/P

MCPM or MCPH MCPA or MCP DCPD DCPA or DCP OCP α-TCP β-TCP ACP HA or HAp OHA OA TTCP

Monocalcium phosphate monohydrate Monocalcium phosphate anhydrous Dicalcium phosphate dehydrate (Brushite) Dicalcium phosphate anhydrous (Monetite) Octacalcium phosphate α-Tricalcium phosphate β-Tricalcium phosphate (Whitlockite) Amorphous calcium phosphate Hydroxyapatite Oxyhydroxyapatite Oxyapatite Tetracalcium phosphate

Ca(H2PO4)2⋅H2O Ca(H2PO4)2 CaHPO4⋅H2O CaHPO4 Ca8(HPO4)2(PO4)4⋅5H2O α-Ca3(PO4)2 β-Ca3(PO4)2 Cax(PO4)y⋅nH2O Ca10(PO4)6(OH)2 Ca10(PO4)6(OH)2–2xOx Ca10(PO4)6O Ca4(PO4)2O

0.5 0.5 1.0 1.0 1.33 1.5 1.5 1.2–2.2 1.67 1.67 1.67 2.0

2.2.2 Titanium and its alloys Titanium and its alloys have been employed in a wide range of applications such as the aerospace industry and biomaterials. In recent decades, titanium and its alloys have attracted much more attention as biomaterials owing to their good biocompatibility, high corrosion resistance and balanced mechanical properties (low specific gravity) [20]. Titanium is an allotropic material which exhibits a hexagonal close packed crystal structure (hcp, α-Ti) at low temperature and a body-centered cubic structure (bcc, β-phase) at high temperatures (>880 °C) [21, 22]. This α-β transformation temperature can be altered by addition of alloying elements such as aluminum, zirconium, vanadium, molybdenum, tantalum and niobium etc. [23]. Overall, the mechanical properties of titanium alloys can be controlled by optimizing the microstructures through addition of alloying elements or heat treatment. The most important titanium alloys that have been developed and studied in recent years are listed in Tab. 2.3 along with their mechanical properties [20, 24–26]. Among all these materials, commercially

20   

   2 HAp-metal based biocomposite coatings

pure (cp) titanium and Ti-6Al-4V alloy are still the most widely used biomaterials in biomedical applications [21]. Tab. 2.3: Mechanical properties of titanium and its selected alloys [20, 24–26]

Material

Standard

Hardness (HV) Modulus (GPa) Tensile Strength (MPa) Type

Ti CP1-4 Ti6Al4V Ti6Al6V2Sn Ti6Al7Nb Ti12Mo6Zr2Fe Ti15Mo Ti29Nb13Ta4.6Zr

ASTM 1341 ASTM F136 – ASTM F1295 ASTM F1813 ASTM F2066 –

120–260 300–400 300–400 – – – –

102–104 101–110 110–117 110 74–85 78 80

240–550 860–965 1000–1100 900–1050 1060–1100 874 911

a a+b a+b a+b b b b

2.3 Plasma Spray of HAp-Ti/Ti6Al4V based composites Plasma is a complicated phenomenon, which is basically an electrically neutrally mixed medium of charged particles (positive ions and negative electrons or ions). In the plasma spray process, a torch (or gun) is used to generate the plasma either by means of direct current (DC) arc or radio frequency discharge (RF). The DC arc is the most commonly used heat source and more than 95 % of industrial installations use a DC plasma torch. Usually, the plasma gun consists of a circular anode made of copper, and a cathode made of thoriated tungsten as shown in Fig. 2.1.

Powder Carrier Gas Coating Plasma Gas (Ar, He) Anode Cathode

Plasma Substrate Fig. 2.1: Schematic view of atmosphere plasma spraying system.

2.4 Property requirement of biocomposites   

   21

An electric arc is created between the anode and the cathode within the plasma gun. The plasma jet is produced and expanded at atmosphere by passing a plasma gas through the arc which is formed between the tip of the cathode and the wall of the anode. According to the different atmospheres, the plasma spray can be divided into three categories: atmospheric plasma spray (APS), vacuum plasma spray (VPS) and controlled atmosphere plasma spray (CAPS) [27–29]. Typically Ar+H2, Ar+N2 or Ar+He gaseous mixtures are used as working gases to form the plasma jet. Each gas has its own effect. Ar stabilizes the arc formed inside the gun. H2, He or N2 enhance the heat transfer to the particles due to their high heat conductivity. The typical temperature in which a plasma gun works is 10,000 to 16,000 K and the velocity of the plasma jet at the exit nozzle can reach 800 m/s [30]. Atmospheric plasma spray process is used to spray different types of materials, including liquids and solids (e.g. metals, ceramics and polymers etc.) and these materials are usually fed into the plasma plume as shown in Fig. 2.1 along with a carrier gas. The advantages of the plasma spray process are high processing temperatures and easy densification of the coatings. With increasing demand for implants and artificial bones in recent years, plasma-sprayed HAp coatings have attracted more attention because of good adhesion characteristics. However, due to a large difference in the thermal expansion coefficients between the ceramic coatings and the metal implant substrates, residual stresses arise at the ceramic/metal interfaces. These residual stresses often cause cracks and reduce the adhesion strength of the ceramic coatings. Therefore, gradient coatings and porous films have been demanded to prevent these cracks in the coatings. In addition, porous HAP has good bioactivity, so that the new bone is formed inside the pores easily. RF and DC plasma spray techniques have been employed to spray HAp-based coatings for implant applications, and usually the adhesion strength of RF-sprayed coatings tends to increase with increasing the applied RF jet power. But, decomposition of HAp with RF jet power is also very high compared to the DC jet. Therefore, using a DC plasma jet with relatively low power to reduce the HAp decomposition and obtain enhanced mechanical properties is more preferred because, in general, the velocity of injected HAp powders is higher in the DC compared to the RF plasma jet. In addition, the control of the RF plasma jet is not so easy compared to the DC plasma jet [31].

2.4 Property requirement of biocomposites The property requirements of biocomposites vary for different applications. However, their selection must meet all the requirements for a particular application. Generally, the following attributes are necessary for an implant coating in load-bearing applications.

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   2 HAp-metal based biocomposite coatings

2.4.1 Mechanical properties Upon implantation, biomaterials are commonly subjected to cyclic loading in the corrosive body fluid environment. Mechanical properties of biomaterials ensure the successful long-term stability of the implant without fracture/failure. Among various properties, the important mechanical properties required for a load-bearing implant are: hardness, Young’s modulus, ultimate tensile strength, elongation and wear resistance. Hardness describes the resistance of a material to penetration and is usually obtained by an indentation test. Young’s modulus is a measurement of the stiffness of a material and is defined as the ratio of linear stress to linear strain in the elastic deformation range. Ultimate tensile strength is the maximum tensile stress that a material can withstand before it fails. Elongation describes the capability of a material to extend under tensile strength. It is usually expressed as a percentage. Wear resistance describes how capable a material is to resist the gradual wearing away caused by abrasion and friction. Hardness can be used as one of the criteria for judging the wear resistance of a material. Usually, hard materials show better wear resistance [24, 32]. Besides the above mentioned properties, the interfacial bond strength is also very important for load-bearing biocomposite coatings on the implants. A metallic implant is commonly coated with biocompatible and/or bioactive agents (e.g. HAp) for favorable tissue integration. Therefore, the mechanical stability of the interface between the coating and the metallic substrate becomes crucial, and the bond strength is an important measurement of this particular property.

2.4.2 Biocompatibility When biomaterials are implanted into the body, the body could respond in different ways. Some initial inflammatory bursts are usually expected. In order to describe the compatibility between biomaterials and the body, the word “biocompatibility” was first used in 1970 by RJ Hegyeli and CA Homsy [33]. Over the years, the evolution of concepts of biocompatibility experienced several steps and many definitions were offered accordingly [34]. From 1940 to 1980, people believed that the materials that are least reactive in implant situations give the best performance. Thus materials that are non-toxic, non-immunogenic, non-thrombogenic etc. are selected for biomaterials, and the biocompatibility is defined as the least negative response. This definition mainly considers the stability of biomaterials and ignores some situations in which the biomaterials are designed to react with the body. Thus, this definition was reevaluated by Williams in 1987, which is also the most cited one, as: “Biocompatibility refers to the ability of a material to perform with an appropriate host response in a specific situation” [35]. In order to further articulate the concept regarding the mechanism of biocompatibility, Williams redefined biocompatibility in 2008 as: “Biocompatibility refers to the ability of a biomaterial to perform its desired function

2.5 Property evaluation    

   23

with respect to a medical therapy, without eliciting any undesirable local or systemic effects in the recipient or beneficiary of that therapy, but generating the most appropriate beneficial cellular or tissue response in that specific situation, and optimizing the clinically relevant performance of that therapy” [34]. Recently, a definition which places more emphasis on biological reactions has been offered by Ratner as: “the ability of a material to locally trigger and guide non-fibrotic wound healing, reconstruction and tissue integration” [33].

2.4.3 Bioactivity The concept of bioactive materials was first defined in 1969 as “the materials that elicit a specific biological response at the interface of the materials which results in the formation of a bond between the tissues and the materials” [36]. Obviously, there is a slight difference between the meaning of biocompatibility and bioactivity as the latter emphasizes the active and positive reactions between the tissues and the implant materials. In 1980, Hench et al. [37] confirmed the formation of a calcium phosphate film on bioglass in the body environment. In 1988, Hench introduced an index IB, although not a standard parameter, to evaluate the level of bioactivity by the time elapsed until 50 % of the interface is boned [36]. Kokubo et al. [38, 39] proposed the essential requirement for a specific material to bond to a living bone is the “apatite formation on its surface inside the body”. From then on, examination of formation of apatite film in simulated body fluid (SBF) became a popular method to evaluate the bioactivity of implant materials.

2.5 Property evaluation 2.5.1 Bond strength The bond strength describes the degree of adhesion of a coating to a substrate when a tensile strength is acting normal to the surface of the coating. An ASTM standard test procedure (ASTM C633-01) is commonly used for evaluating the bond strength of thermal spayed coatings. In order to test the bond strength, the required coating is made on a cylindrical pin substrate and then glued to another similar pin using an epoxy. The whole assembly, consisting of the two pins, is then placed in a furnace for 1 hr at 120 °C, which allows the epoxy to cure. After curing the epoxy the whole assembly is loaded into a tensile testing machine using appropriate fixtures and a tensile load is applied on the assembly by moving the crosshead of the machine at a speed of 0.05 in/min. The value of the load at which the whole assembly fails is noted as the bond strength, which is calculated as the ratio between maximum load and the cross section area of the pin. Thus calculated, bond strength values give adhesive

24   

   2 HAp-metal based biocomposite coatings

strength of the coatings if the failure occurs entirely at the coating-substrate interface, and otherwise it would be cohesive strength, if the failure is within the coating. On the other hand, if the failure occurs within the epoxy then it implies that the coating is stronger than the epoxy, and, therefore, glue with a higher tensile strength should be chosen for the test [40].

2.5.2 Corrosion behavior evaluation In terms of corrosion, the human body presents a more aggressive environment compared to water. Degradation of biomaterials in the human body limits the functionality and fatigue life of the implants and thereby leads to their mechanical failure. Besides, the released corrosion products may elicit adverse biological reactions and cause discoloration, immunologic disease, foreign body response etc. [21]. Thus, corrosion resistance has become an important parameter in determining the suitability of a material for biomedical applications. There are certain standard electrochemical testing techniques in use to determine the corrosion behavior of biomaterials, such as open circuit potential (OCP), potentiodynamic polarization and electrochemical impedance spectroscopy (EIS). For OCP measurements, the biomaterial is exposed to a simulated physiological solution and the open circuit potential is commonly recorded for 1 hour. For potentiodynamic polarization studies, the material is polarized from negative potential to positive potential at very low scan rates and the variation of current with respect to potential is recorded [41]. EIS measurement has also been extensively used as an efficient technique for corrosion behavior evaluation in recent years. It is one of the linear response methods. The impedance spectrum is measured as a function of frequency of a small-amplitude sinusoidal potential perturbation superimposed on a direct potential bias. Small amplitude perturbations bring two advantages. First, the effects of perturbations on the studied system will be small, which is beneficial to describe the system more accurately. Second, the response of the system will be linear, namely it contains only the first-order terms of the Taylorexpanded nonlinear current-voltage relationship, which is beneficial to the analysis of the data [49, 50].

2.5.3 Immersion test in simulated body fluid The immersion test in simulated body fluid (SBF) has been accepted by a large part of the scientific community as a simple and effective technique for bioactivity evaluation. SBF is an artificial solution with ion concentrations nearly equal to those of human blood plasma. The specimens are immersed in the SBF and the experiments are commonly performed at 37 °C. After certain periods of time, the specimens are removed from the SBF solution and then washed with distilled water and dried at

2.6 Plasma sprayed HAp-(Ti/Ti6Al4V) based composite coatings    

   25

room temperature. The morphological analysis of the newly-formed bonelike apatite layer on the surface of the specimens can be studied to understand the effect of the SBF as well as the bioactivity of the implant material. The time period required for apatite formation and the thickness of the apatite layer indicate bioactivity of the material.

2.6 Plasma sprayed HAp-(Ti/Ti6Al4V) based composite coatings Plasma spray is currently the only technique which has been commercially used for deposition of HAp coatings on dental and orthopedic implants due to easy manufacturability and cost effectiveness [44]. The advantages of plasma-sprayed HAp coatings include early fixation between the bone tissues and the implant materials with great bone ingrowth and reduced metal ion release from the metallic implants. However, the composition and morphology of HAp in the coatings could be significantly different compared to the feedstock due to the high temperatures involved in the plasma spray process [18]. As previously mentioned, some important properties that are generally expected of biomaterials are good biocompatibility, bioactivity and mechanical properties. Therefore, in the following discussion, these three properties are particularly articulated for plasma-sprayed HAp coatings. Although the concept of biocompatibility has evolved to a broader definition, the term biocompatibility here only refers to the stability of biomaterials in the bio-environment. Thus, the corrosion behavior of the biomaterial in vitro can be used to indicate its biocompatibility. The active reactions between biomaterials and the body are described by bioactivity which can be examined by an immersion test in SBF. Although the plasma-sprayed HAp coatings show excellent biocompatibility and bioactivity, the adhesion of a HAp coating to its substrate and the long-term stability of the substrate/coating interface are always concerned. Various mechanically strong secondary phases (e.g. Ti/Ti6Al4V) [5, 10, 13], even a third phase (e.g. Yittria Stabilized Zirconia) [45] has been added to enhance the interface properties.

2.6.1 Bond strength of plasma-sprayed HAp-(Ti/Ti6Al4V) based composite coatings Although the mechanical properties (e.g. hardness, tensile strength, Young’s modulus, fracture toughness etc.) are very important for load-bearing implant applications, the bond strength between the biocomposite coating and the metallic implant also becomes an important factor. Thermal-sprayed HAp coatings are commonly observed to fail at the coating and implant interface rather than the interface between the coating and the bone or within the coating [18, 46], therefore evaluation of their bond strength is very critical.

26   

   2 HAp-metal based biocomposite coatings

The factors that influence adhesive bond strength of the plasma-sprayed HAp coatings include porosity, surface roughness, residual stress etc. The effects of particle size and spray distance on the bond strength were studied by Kweh et al. [47]. Usually small particles undergo complete melting in the plasma plume before deposition on the substrate. Thus, fully melted particles form a dense lamellar structure and result in enhanced bond strength of the coatings. Another way to increase adhesion of the HAp coatings is to increase the surface roughness of the substrate. Although sand or grit blasting has been extensively used, the roughness can be increased only by a factor of 2 to 3. Tsui et al. [48] precoated a 100 μm Ti layer using large particles to increase the surface roughness by a factor of more than 20. By doing so, the adhesion of HAp coating was improved significantly. They also investigated the effect of residual stress on the adhesion, and found that the residual stress in the coating was always tensile, which may be sufficient to promote cracking in the coating and/or at the interfaces. Therefore, reducing the residual stress by minimizing the substrate temperature during the spray process or employing heat treatment after spraying is suggested [46]. Zheng et al. [49] studied the bond strength of plasma-sprayed HAp/Ti composite coatings and found an improvement in the mechanical properties of the coatings with the addition of Ti. The bond strength was observed to increase from 12.9 to 14.5 MPa when 20 wt% of Ti was added, and with the increase of Ti content up to 60 wt%, the bond strength was further enhanced to 17.3 MPa. Poor adhesion of pure HAp coatings is mainly due to mismatch of the coefficients of thermal expansion between the Ti substrate and the HAp coatings. This is where reinforcement of HAp coatings with Ti in HAp/Ti composite coatings was observed to help enhance the bond strength and reduce the residual stresses within the coatings. The influence of the addition of Ti6Al4V on the bond strength was investigated by Khor et al. [13]. Their results demonstrated a significant increase in the bond strength of the composite coatings with the addition of 50 wt% of Ti6Al4V. The mechanical stability of the composite coatings subjected to physiological medium was also much better than that of pure HAp coatings. Even after soaking in SBF for 8 weeks, the bond strength of the composite coatings was observed to be reduced only by ~32 % of the original value compared to the pure HAp coatings that showed a 75 % reduction. Yttria stabilized zirconia (YSZ) was also added to HAp/ Ti6Al4V composite coatings as the third phase to further improve the bond strength by Khor et al. [45]. The bond strength improved further from ~27.5 MPa to ~32.5 MPa by addition of 15 wt% of YSZ, and after 4 weeks of immersion in the SBF, the decrease in the bond strength was around 27.7 %.

2.6 Plasma sprayed HAp-(Ti/Ti6Al4V) based composite coatings    

   27

2.6.2 Electrochemical corrosion behavior of plasma-sprayed HAp-(Ti/Ti6Al4V) based composite coatings The corrosion of HAp coated biomaterials has been extensively studied and compared with non-coated ones in recent years. Usually, the electrochemical performance of the coating systems is mainly determined by the dissolution and passivation characteristics of the underlying metallic substrate [50]. Souto et al. [50] reported that the overall corrosion resistance of the implants was not much affected by the presence of pure HAp coatings on the surface, although there was a slight improvement in the corrosion resistance with HAp coatings. On the other hand, Valereto et al. [51] reported a reduced corrosion resistance of HAp coated implants by two times or more than that of uncoated ones. Therefore, the overall performance of the implant-coating systems is strongly dependent on the quality of the HAp coatings deposited, and their porosity, crystallinity and thickness, which are also important factors to influence the electrochemical characteristics. The microstructural parameters of the coatings are greatly dependent on the deposition systems employed, spray parameters and stock powder size etc. used. Therefore, a close tailoring of the coating and the process control as well as coating optimization are necessary to obtain enhanced overall corrosion characteristics of the implant-coating systems in biomedical applications.

2.6.3 Immersion behavior of plasma sprayed HAp-(Ti/Ti6Al4V) based composite coatings Bone-like apatite layer is very important to establish the bone-bonding at the interface between the bioactive materials and living tissues. Apatite deposition on the surface of the HAp-Ti/Ti6Al4V composites after immersion in SBF has been widely used to evaluate the bioactivity of the implant materials. Plasma-sprayed pure HAP or reinforced HAP composite coating showed good bioactivity in the in vitro studies. Based on the previous researchers’ observation, formation of an apatite layer commonly happens through three stages. In the first stage, the top surface of HAp coating appears to dissolve during immersion and some spherical particles precipitate. During the second stage, more spherical particles precipitate, and the size of the initially precipitated spherical particles increases as immersion duration increases. At this stage, the surface of the HAp-Ti/Ti6Al4V composite is covered by a dune like apatite layer with the spherical particles embedded inside. Finally, in the third stage, many granular particles nucleate and precipitate in the initially formed apatite layer, which indicates a continuous growth of the apatite layer [52, 53]. For HAp composite coatings, with the increase in immersion time, the fourth stage of apatite layer formation is also observed, in which the apatite patterns initially formed on the top surface of the HAp coating spread out to cover the Ti matrix too [53]. Fig. 2.2 shows the cross section and surface morphology of a plasma-sprayed HAp-Ti composite coating and

28   

   2 HAp-metal based biocomposite coatings

its immersion behavior in SBF for different time periods [53]. The composite coating exhibits a lamellar structure (see Fig. 2.2a) and the dark gray and light white areas are HAp and Ti, respectively. After soaking in SBF for various time periods, the as-sprayed coating shows formation of a dune like apatite layer (Fig. 2.2b,c) on the top surface of the HAp part coating in the second stage and a thick, tortoise-shell-like apatite layer on the surface of the entire coating in the fourth stage (Fig. 2.2d) [53].

a

b

c

d

Fig. 2.2: Cross section and surface morphology of plasma-sprayed Hap-Ti composite coating and its immersion behavior in SBF solution: (a) cross section, (b), (c) surface morphology variation in 3 weeks, second stage, (d) surface morphology variation in 8 months, fourth stage [53].

Formation of amorphous phase and decomposition of HAp are unavoidable during the preparation of HAp composite coatings by plasma spray technique. Thus the plasma-sprayed HAp coatings are composed of crystalline and amorphous phases of HAp along with some secondary phases, such as CaO, TTCP and TCP. A positive result from the dissolution of amorphous phase was observed by Weng et al. [54], in which it promoted the deposition of apatite layer even in the early stages. However, a complete dissolution of secondary phase was also observed after a two week immersion in the SBF [52], which on the other hand deteriorated the long-term stability of the HAp

References   

   29

coating. Therefore, proper design of coatings and especially control of spray parameters is very important to achieve optimal coating properties and is still one of the critical research topics in the field. In addition, the formation of apatite layer on the surface of the HAp coating in the SBF was not distinctly affected by the addition of Ti up to 60 wt%, which indicates a very good bioactivity of HAp composite coatings [49].

2.7 Conclusions In this chapter we have reviewed the important role of metallic second phase in the HAp-Ti/Ti6Al4V based composite coatings for implant applications as well as their property requirements from the biomedical applications point of view. The advantages of the plasma spray process for depositing biocomposite coatings has been elicited along with the processing artifacts. The literature indicates a requirement for close control of the parameters in the atmospheric plasma spray process and the addition of metallic phase to improve the adhesion, mechanical and corrosion characteristics while maintaining the bioactivity of the composites. Therefore, in order to obtain the optimized overall property, deep and extensive research is necessary on plasma spray deposition of HAp-based biocomposite coatings.

References [1]

[2]

[3] [4] [5] [6] [7] [8] [9]

M.F. Morks, N.F. Fahim, A. Kobayashi, Structure, mechanical performance and electrochemical characterization of plasma sprayed SiO2/Ti-reinforced hydroxyapatite biomedical coatings, Appl. Surf. Sci. 255 (2008) 3426. Yanfeng Dai, Min Xu, Junchao Wei, Haobin Zhang, Yiwang Chen, Surface modification of hydroxyapatite nanoparticles by poly(l-phenylalanine) via ROP of l-phenylalanine N-carboxyanhydride (Pha-NCA), Appl. Surf. Sci. 258 (2012) 2850. C.Q. Ning, Y. Zhou, In vitro bioactivity of a biocomposite fabricated from HA and Ti powders by powder metallurgy method, Biomaterials 23 (2002) 2909. R.S. Lima, K.A. Khor, H. Li, P. Cheang, B.R. Marple, HVOF spraying of nanostructured hydroxyapatite for biomedical applications, Mater. Sci. Eng. A 396 (2005) 181. Won-Gi Kim, Han-Cheol Choe, Surface characteristics of hydroxyapatite/titanium composite layer on the Ti-35Ta-xZr surface by RF and DC sputtering, Thin Solid Films 519 (2011) 7045. Diana Garcia-Alonso Garcia, Plasma spray deposition of hydroxyapatite based composites as a step towards bone scaffolds, thesis, Dublin City University, Ireland, 2009. M. Metikos-Hukovic, E. Tkalcec, A. Kwokal, J. Piljac, An in vitro study of Ti and Ti-alloys coated with sol–gel derived hydroxyapatite coatings, Surf. Coat. Technol. 165 (2003) 40. Congqin Ning, Yu Zhou, Correlations between the in vitro and in vivo bioactivity of the Ti/HA composites fabricated by a powder metallurgy method, Acta Biomaterialia 4 (2008) 1944. Shinn-Jyh Ding, Properties and immersion behavior of magnetron-sputtered multi-layered hydroxyapatite/titanium composite coatings, Biomaterials 24 (2003) 4233.

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[10] Chun-Cheng Chen, Tsui-Hsien Huang, Chia-Tze Kao, Shinn-Jyh Ding, Characterization of functionally graded hydroxyapatite/titanium composite coatings plasma-sprayed on Ti alloys, J. Biomed. Mater. Res. Part B: Appl Biomater 78B (2006) 146. [11] Tanya. J. Levingstone, Optimization of Plasma Sprayed Hydroxyapatite Coatings, thesis, Dublin City University, Ireland, 2008. [12] R. Narayanan, S.K. Seshadri, T.Y. Kwon, K.H. Kim, Calcium phosphate-based coatings on Titanium and its alloys, J. Biomed. Mater. Res. Part B: Appl. Biomater. 85B (2008) 279. [13] Y.W. Gu, K.A. Khor, P. Cheang, In vitro studies of plasma-sprayed hydroxyapatite/Ti-6Al-4V composite coatings in simulated body fluid (SBF), Biomaterials 24 (2003) 1603. [14] Mehdi Sadat-Shojai, Mohammad-Taghi Khorasani, Ehsan Dinpanah-Khoshdargi, Ahmad Jamshidi, Synthesis methods for nanosized hydroxyapatite in diverse structures, Acta Biomaterialia (2013), http://dx.doi.org/10.1016/j.actbio.2013.04.012. [15] Jeffery O. Hollinger, An Introduction to Biomaterials, Second Edition, Taylor & Francis Group, 2011. [16] Leng, Y. and Qu, S., TEM examination of single crystal hydroxyapatite diffraction, J. Mater. Sci. Lett. 21 (2002) 829. [17] Sergey V. Dorozhkin, Nanosized and nanocrystalline calcium orthophosphates, Acta Biomaterialia 6 (2010) 715. [18] Limin Sun, Christopher C. Berndt, Karlis A. Gross and Ahmet Kucuk, Material Fundamentals and Clinical Performance of Plasma-Sprayed Hydroxyapatite Coatings: A Review, J. Biomed. Mater. Res. 58 (2001) 570. [19] Robert B. Heimann, Thermal spraying of biomaterials, J. Surf. Coat. Technol. 201 (2006) 2012. [20] Mitsuo Niinomi, Tomokazu Hattori, Toshihiro Kasuga, Hisao Fukui, Titanium and its alloys, Encyclopedia of Biomaterials and Biomedical Engineering, Taylor & Francis Group, 2005. [21] Seeram Ramakrishna, Murugan Ramalingam, T.S.Sampath Kumar, Winston O. Soboyejo, BIOMATERIALS: A Nano Approach, Taylor & Francis Group, 2010. [22] N. Poondla, T.S. Srivatsan, A. Patnaik, M. Petraroli, A study of the microstructure and hardness of two titanium alloys: Commercially pure and Ti-6Al-4V, J. Alloy Compd. 486 (2009) 162. [23] Mechanical Engineers’ Handbook, Second edition, John Wiley & Sons, Inc, 1998. [24] M. Geetha, A.K. Singh, R. Asokamani, A.K. Gogia, Ti based biomaterials, the ultimate choice for orthopaedic implants – A review, Prog. Mater. Sci. 54 (2009) 397. [25] Mitsuo Niinomi, Mechanical properties of biomedical titanium alloys, Mater. Sci. Eng. A 243 (1998) 231. [26] C. Veiga, J.P. Davim, A.J.R. Loureiro, Properties and applications of Titanium alloys: A brief review, Rev. Adv. Mater. Sci. 32 (2012) 133. [27] Davis, J. R. (ed.), Handbook of Thermal Spray Technology, ASM Thermal Spray Society, 2005. [28] P. Fauchais, Understanding plasma spraying, J. Phys. D: Appl. Phys. 37 (2004) 86. [29] J. Webster (ed.), Thermal Spray Coating, John Wiley & Sons, Inc., 2007. [30] Lech Pawlowski, The Science and Engineering of Thermal Spray Coatings, John Wiley & Sons, Inc., 2008. [31] Masahiro Kou, Takafumi Toda, Osamu Fukumasa, Production of fine hydroxyapatite films using the well-controlled thermal plasma, Surf. Coat. Technol. 202 (2008) 5753. [32] Teoh Swee Hin, Engineering Materials for Biomedical Applications, World Scientific Publishing Co. Pte. Ltd., 2004. [33] Buddy D. Ratner, The biocompatibility manifesto: biocompatibility for the twenty-first century, J. of Cardiovasc. Trans. Res. 4 (2011) 523. [34] David F. Williams, On the mechanisms of biocompatibility, Biomaterials 29 (2008) 2941. [35] Williams D.F. Definitions in Biomaterials, Amsterdam: Elsevier, 1987. [36] Wanpeng Cao, Larry L. Hench, Bioactive materials, Ceram. Int. 22 (1996) 493.

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[37] Makata Ogino, Fumio Ohuchi, L.L. Hench, Compositional dependence of the formation of calcium phosphate films on bioglass, J. Biomed. Mater. Res.14 (1980) 55. [38] Tadashi Kokubo, Bioactive glass ceramics: properties and applications, Biomaterials 12 (1991) 155. [39] Tadashi Kokubo, Hiroaki Takadama, How useful is SBF in predicting in vivo bone bioactivity? Biomaterials 27 (2006) 2907. [40] ASTM INTERNATIONAL: Standard Test Method for Adhesion or Cohesion Strength of Thermal Spray Coatings, ASTM C633-01 2008. [41] ASTM INTERNATIONAL: Standard Test Method for Conducting Cyclic Potentiodynamic Polarization Measurements to Determine the Corrosion Susceptibility of Small Implant Devices, ASTM F2129. [42] Mark E. Orazem, Bernard Tribollet, Electrochemical Impedance Spectroscopy, John Wiley & Sons, Inc., 2008. [43] Robert S Rodgers, An introduction to electrochemical impedance spectroscopy (EIS), ACS Princeton Local Section, 2009. [44] R. Petit, the use of hydroxyapatite in orthopedic surgery: a ten-year review, Eur. J. of Orthop. Surg. & Traumatol. 9 (1999) 71. [45] Y.W. Gu, K.A. Khor, D. Pan, P. Cheang, Activity of plasma sprayed yttria stabilized zirconia reinforced hydroxyapatite/Ti-6Al-4V composite coatings in simulated body fluid, Biomaterials 25 (2004) 3177. [46] Y.C. Tsui, C. Doyle, T.W. Clyne, Plasma sprayed hydroxyapatite coatings on titanium substrates Part 1: Mechanical properties and residual stress levels, Biomaterials 19 (1998) 2015. [47] S.W.K. Kweh, K.A. Khor, P. Cheang, Plasma-sprayed hydroxyapatite (HA) coatings with flamespheroidized feedstock: microstructure and mechanical properties, Biomaterials 21 (2000) 1223. [48] Y.C. Tsui, C. Doyle, T.W. Clyne, Plasma sprayed hydroxyapatite coatings on titanium substrates. Part 2: optimisation of coating properties, Biomaterials 19 (1998) 2031. [49] Xuebin Zheng, Minhui Huang, Chuanxian Ding, Bond strength of plasma-sprayed hydroxyapatite/Ti composite coatings, Biomaterials 21 (2000) 841. [50] Ricardo M. Souto, Maria M. Laz, Rui L. Reis, Degradation characteristics of hydroxyapatite coatings on orthopaedic TiAlV in simulated physiological media investigated by electrochemical impedance spectroscopy, Biomaterials 24 (2003) 4213 . [51] I.C. Lavos-Valereto, I. Costa,S. Wolynec, the electrochemical behavior of Ti-6Al-7Nb Alloy with and without Plasma-Sprayed Hydroxyapatie Coating in Hank’s Solution, J. Biomed. Mater. Res. 63 (2002) 664. [52] S.W.K. Kweh, K.A. Khor, P. Cheang, An in vitro investigation of plasma sprayed hydroxyapatite (HA) coatings produced with flame-spheroidized feedstock, Biomaterials 23 (2002) 775. [53] Xuan Zhou, Raj Siman, Lin Lu, Pravansu Mohanty, Argon atmospheric plasma sprayed hydroxyapatite/Ti composite coating for biomedical applications, J. Surf. Coat. Technol. 207 (2012) 343. [54] J. Weng, Q. Liu, J.G.C. Wolke, D. Zhang, K. De Groot, The role of amorphous phase in nucleating bone-like apatite on plasna-sprayed hydeoxyapatite coatings in simulated body fluid, J. Mater. Sci. Lett. 16 (1997) 335.

Julieta Volpe, Lucía M. Masi, Vera A. Alvarez and Jimena S. Gonzalez

3 Hydrogels based on poly(vinylalcohol) for cartilage replacement 3.1 Hydrogels: General Ideas Over the last fifty years, hydrogels have become of great interest for numerous biomedical and pharmaceutical applications such as soft contact lenses, surface coatings, dressings, and drug delivery systems [1–4]. The first hydrogel polymer with potential biomedical uses was discovered by Lim and Wichterle in 1955 [1]. That was a copolymer of 2-hydroxyethyl methacrylate and ethylene dimethacrylate, with high stability under different pH and temperature conditions, which makes it suitable for contact lens production. Many characteristics and the significant progress made in designing and synthesizing these kinds of material have made them highly attractive for tissue engineering and bionanotechnology applications [2–4]. There is no precise definition of the term hydrogel [1]. They are considered water swollen, crosslinked polymers, which consist of over 90–99 % of water. Hydrogels are prepared by chemical polymerization or physical assembly of synthetic or natural resources [1–2, 5]. Collagens and crosslinked guar gum are examples of natural polymers that are modified to produce hydrogels, while poly(vinyl alcohol) (PVA), poly(acrylonitrile)/poly(pyrrole) (PAN/PPY), poly(hydroxyethyl methacrylate) (PHEMA), N-isopropylacrylamide (NIPA), poly(acrylamide) (PAM), poly(acrylic acid) (PAA), and poly(acrylonitrile) (PAN) are examples of synthetic ones [6]. Hydrogels are hydrophilic multiphase materials that display excellent biocompatibility and both solid-like and liquid-like properties [3, 6]. Their structure consists of a three-dimensional network of polymer chains, usually described as a mesh, with the interstitial space filled up with water; biological fluid and/or ion species (Fig. 3.1). The network holds the fluid in place providing solidity to the hydrogel, while the fluid phase makes it wet and soft [6]. Due to this particular structure, these polymers show viscoelastic and sometimes pure elastic behavior [7].

Crosslinks Mobile ion Fixed ion Undissociated ionizable group

Interstitial fluid phase

Polymer chain

Fig. 3.1: Schematic representation of the internal structure of a hydrogel.

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   3 Hydrogels based on poly(vinylalcohol) for cartilage replacement

Hydrogels can absorb a large amount of water without dissolving, which allows easy exchange of nutrients and wastes with the surrounding environment [2]. Proteins and polysaccharides in the extracellular matrix can be considered as hydrogels. This similarity with the extracellular matrix gives hydrogels the ability to mimic human tissue [2]. It is possible to classify hydrogels in several categories, based on: – Kind of polymer: homopolymer, copolymer, multipolymer, or interpenetrating polymeric hydrogels. – Ionic charges: neutral, anionic, cationic, or ampholytic hydrogels. – Physical structural features: hydrogen-bonded and complexation structures, amorphous, or semicrystalline hydrogels [8]. – Nature of their crosslinks: chemical or physical hydrogels. Chemical hydrogels are composed of three-dimensional networks produced by chemical crosslinking, resulting in insoluble structures with covalent bonds between the chains. The most common method used to obtain this kind of hydrogels is radical crosslinking, initiated by temperature or UV irradiation [9]. The procedure is complex and not versatile and residual amounts of toxic radicals and monomers should be carefully removed after preparation. If the residual monomers remain in the hydrogel, the biocompatibility of this material will be affected [9]. In physical hydrogels, the network structure is built from small amphiphilic molecules to macromolecules self-assembled through non-covalent crosslinks, such as electrostatic forces, hydrogen bonds, hydrophobic interactions, or chain entanglement. This technique avoids the use of crosslinking agents, providing compounds that are free of toxic [5, 7].

3.2 Main properties of hydrogels Hydrogels are a unique class of macromolecular networks that may contain a large fraction of aqueous solvent within their structure [10]. Their capacity to swell without dissolving is called swelling. In order to achieve high swelling capability, it is important to generate a structure with good affinity with the aqueous solvent. This affinity depends on several factors: structural modification (porosity), hydrophilicity, and chain flexibility [11]. The presence of functional groups (amide, hydroxyl, and carboxyl) and monovalent ions (potassium, sodium, and ammonium) favors the dissolution process thus improving the swelling properties of the polymer. Hydrogels can absorb different amounts of water, even several thousand times their weight, depending on the number and types of ion and functional groups as well as the network density, solvent nature and polymer solvent interaction [11–12]. The ratio between the weight of the sample in its

3.2 Main properties of hydrogels   

   35

swollen state (Ms) and the weight of the sample in its dry state (Md), expresses the degree of swelling, Mt%, (Eq. (3.1)). 𝑀𝑡 % =

𝑀𝑠 − 𝑀𝑑 × 100 𝑀𝑑

(3.1)

This parameter can also be calculated as the ratio of the sample volume in the swollen state to the volume in the dry state [2]. The swelling mechanism can be considered as a diffusion process followed by a relaxation process [11, 13]. At the beginning, the rate of swelling is determined by the rate of water diffusion which depends on the solution temperature, the molecular weight of the solvent and the extent of porosity within the hydrogel structure. In the following stage, the absorption rate slows down due to polymer chain relaxation. Fig. 3.2 shows

A typical Swelling curve Diffusion

Relaxation

Mt %

Equilibrium

Time

Swelling degree

Swelling Kinetics

Crosslink density

Time Fig. 3.2: Schematic swelling curves typical of hydrogels. Water absorption as a function of time; and effect of crosslinking density on the swelling kinetic.

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   3 Hydrogels based on poly(vinylalcohol) for cartilage replacement

that as the crosslinking density increases, the diffusion-controlled stage is reduced and equilibrium swelling is reached in a shorter time [11]. The final water content of hydrogels depends on both kinetics and thermodynamics parameters [13]. Swelling influences most of the properties of hydrogels’ materials such as: mechanical, optical, barrier and surface properties [2]. Usually, large swelling corresponds to poor mechanical properties. However, for most applications is important to reach a compromise between large swelling and good mechanical behavior. This can be achieved by different strategies such as varying the polymerization conditions or increasing the crosslinking density, which improves the elastic modulus and the structural properties but decreases the water absorption capacity. Polymerizing hydrophilic monomers with hydrophobic ones is also an alternative to enhance both the swelling and the mechanical properties [14]. Typically, hydrogels are mechanically characterized by different techniques such as uniaxial compression and tensile tests, creep and stress relaxations tests, and indentation tests. Usually, the mechanical properties of the hydrogels are evaluated in compression configuration. The test consists in applying a specific load to the hydrogel in its swollen state, at a constant rate. The polymer deforms itself until it reaches the breaking point, where the applied load becomes stronger than the resistance of the hydrogel, causing its failure [14]. Analyzing the stress-deformation curve (Fig. 3.3), it is possible to identify three different zones. The first one shows a parabolic relationship between the stress and the strain; the second region displays linear behavior, whose slope corresponds to the hydrogel elastic modulus; the sharper the

Parabolic zone Linear zone

Stress, σ

Breaking point

Strain, ε Fig. 3.3: Typical mechanical curve, stress as a function of strain, of hydrogel in compression test.

3.3 Hydrogels as biomaterials   

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slope, the higher the modulus. The last region includes the maximum strength point (breaking point) [11]. Due to hydrogels’ viscoelastic behavior, elastic properties are measured by compression and tensile testing while creep and stress relaxation experiments evaluate the material’s response over time [15]. Swollen hydrogels exhibit viscoelastic properties that can be measured by means of a rheometer or a dynamical-mechanical analysis (DMA). The storage (E’) and loss modulus (E’’) characterize the elastic and the viscous components of the hydrogel properties, respectively. Indentation tests are also commonly performed to analyze the mechanical properties. The method consists of quantifying the reaction force resulting from the indentation of the hydrogel at a single point with an established penetration depth [16]. However, the available testing instruments were originally designed for harder materials, so the study of hydrogels presents some difficulties [15]. In this sense, several researchers have conducted indentation and nanoindentation tests on hydrogel materials, adapting the experimental parameters [17–18], techniques [19] and analysis methods [20]. The understanding of the hydrogels’ tribological properties is critical for their use in tissue engineering applications. Different tests like reciprocating sliding tests [21–22], ball-on-plate tribometer [22–24] and reciprocating pin-on-plate methods [25] have been frequently used. Recently, investigations have focused on the frictional behavior of the hydrogels against many materials [22, 26–29] including natural cartilage [29]. Sardinha et al. [29] performed a study in which natural articular cartilage was in direct contact with a PVA hydrogel implant, in order to assess the tribological properties of PVA-H. With the purpose of mimicking in vivo conditions, the tests were executed on a pin-on-plate tribometer in the presence of distilled water with a buffered solution [30].

3.3 Hydrogels as biomaterials Biomaterials can be defined as “nonviable materials used in a medical device intended to interact with biological systems” [31]. They are natural or synthetic materials designed to treat, augment, or replace any tissue, organ, or function of the body [32]. The performance of a biomaterial can be evaluated by its biocompatibility, which refers to the ability to act with an appropriate host response in a particular application [33]. Nowadays, biomaterials are being used in a large number of devices and implants, such as bone plates, heart valves, biosensors, blood tubes, sutures, dental implants, artificial hearts, joint replacements, intraocular lenses, among others [33–34]. They are made of different materials that may be grouped into ceramics (zirconia, titania, and hydroxyapatite (HA)), metals (stainless steel, gold, and NiTi, Co-Cr and Ti alloys), polymers and composites. Polymers such as polymethylmethacrylate (PMMA), poly-

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ethylene (PE), silicone rubber (SR), polyvinyl alcohol (PVA), polytetrafluoroethylene (PTFE) are the objects of growing interest for various biomedical applications [33]. This is mainly due to the fact that it is possible to obtain them with a wide range of properties, compositions and shapes. They can be readily fabricated into complex structures and shapes. Furthermore, they are capable of absorbing large amounts of liquid and swelling, mimicking human tissue. In some cases polymers are reinforced with different fillers, such as carbon and ultra-high molecular weight polyethylene (UHMWPE) fibers [33], producing composites, in order to improve their mechanical properties. In this chapter, the focus is centered on the study of a composite hydrogel based on PVA, UHMWPE and HA as a potential biomaterial for cartilage replacement.

3.4 Polyvinyl alcohol (PVA) hydrogels: General characteristics Polyvinyl alcohol (PVA) is a hydrophilic, biocompatible and biodegradable synthetic polymer which has a relatively simple structure with a pendent hydroxyl group (Fig.  3.4). The corresponding monomer (vinyl alcohol) does not exist in a stable form. For this reason, PVA is prepared in two stages: the polymerization of vinyl acetate in an alcoholic solution via a free-radical mechanism followed by the partial hydrolysis of poly(vinyl acetate) [35–37]. Different grades of PVA can be obtained by varying the content of acetate groups in the polymer which is called the degree of hydrolysis and affects the polymer solubility, chemical properties, crystallinity and behavior of the final material [34]. Industrially, PVA is produced in the order of several hundred kton per year all over the world, and hydrolysis grades between 70 and 99 % are available for different applications [38].

Fig. 3.4: Chemical structure of poly(vinylalcohol) (PVA).

PVA is an insipid and odorless, white, translucent or cream colored granular powder, developed in 1924 by Hermann et al. [39]. It has a melting point of 180 to 190 °C, and it is soluble in hot water but insoluble in common organic solvents [40]. PVA is widely used in several applications such as adhesive agent, inorganic binder, fibers, coatings, lacquers, films and acetal resins, textiles, the food and the

3.4 Polyvinyl alcohol (PVA) hydrogels: General characteristics   

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cosmetic industries. Due to its excellent biocompatibility it has raised special interest for applications in medicine and pharmaceutical fields [39, 41]. In this sense, PVA hydrogels (PVA-H) have become attractive to the area of ‘tissue engineering’ for regenerating and repairing a large variety of organs and tissues [35, 42–44]. PVA hydrogels are stable and elastic gels with multiple attachment sites for biological molecules. They can also induce cell orientation and enhance the transmission of mechanical stimuli to the cells. Hydrogels can be produced by chemical, irradiative or physical crosslinking methods [3, 36]. Fig. 3.5 shows a freeze-thawed (F-T) hydrogel of PVA and its internal porous structure (scanning electron microscopy image).

a

b

Fig. 3.5: PVA hydrogel (a) macroscopic aspect and (b) SEM image.

For obtaining chemically crosslinked PVA-H, it is necessary to promote the formation of bridges between the pendent hydroxyl groups (OH) present in the PVA chains. This is achieved by using crosslinking agents, such as glutaraldehyde [45] or other monoaldehydes, in conjunction with methanol or acetic or sulfuric acid [46]. As explained previously, the main disadvantage of this method is the residual amounts of toxic radicals and monomers, which make this kind of hydrogels unacceptable for pharmaceutical or biomedical applications. Although it is possible to remove these noxious residues, their extraction involves tedious and expensive procedures [3]. The use of electron-beam or gamma irradiation appears to be an alternative to the chemical crosslinking mechanism. Even though this is a costly method it does not leave toxic residues and improves cell adhesion in tissue-engineering applications [3]. The third technique, physical crosslinking, entails crystallite formation by strong interchain hydrogen bonding. For this procedure, concentrated PVA solutions experience repeated freezing and thawing cycles leading to harmless hydrogels, commonly known as cryogels. They exhibit tissue-like viscoelasticity, a high degree of swelling in water and biological fluids, biodegradability, and stability at room temperature [34, 37]. The studies have shown that the morphology, structure, and stability of the resulting physical PVA-H are affected by diverse parameters: conditions of cryogenic treatment (number of cycles, duration and temperature of freezing/thawing

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   3 Hydrogels based on poly(vinylalcohol) for cartilage replacement

rates), PVA grade (degree of hydrolysis and molecular weight), and concentration of aqueous solution [47]. By manipulating some of these variables it is possible to control the final hydrogel properties, such as water content, diffusion properties, adhesive characteristics and mechanical strength and stiffness [47–48]. Cryogels contain crystallites along a three-dimensional structure are capable of distributing mechanical loads. This characteristic confers improved mechanical stiffness and strength to the physical gels compared to hydrogels obtained by chemical and irradiative methods [34, 37].

3.5 PVA hydrogels for biomedical applications PVA hydrogels are non-carcinogenic and non-toxic biomaterials that have been used for a great number of medical and pharmaceutical applications: lining for artificial hearts [49], drug delivery systems [50], artificial pancreas [51], contact lenses, catheters, hemodialysis membranes, and artificial skin [37]. These hydrogels present a simple chemical structure, viscoelastic behavior and high degree of swelling in water [37]. These characteristics give them the possibility to mimic natural human tissue, making them excellent candidates for tissue engineering. One of the first investigations of PVA hydrogels was centered on the reconstruction of vocal cords [34]. In this study, blood compatibility, elastic behavior, surface and physical properties were examined. The results indicate that the reconstruction was possible and that this type of material was also adequate for kidney biomembranes [37]. In the following years, numerous studies about PVA biocompatibility were carried out [52–56]. Some of them focus on the production of physical crosslinked PVA-H by the freezing-thawing technique [7, 48, 57–60]. The resulting material exhibited characteristics similar to natural tissue like high water content, rubber-like elasticity and high mechanical strength. In order to evaluate the response of the body to this kind of hydrogel, it was implanted into rabbits. Moreover, these implants showed no adhesion to surrounding tissue and good bioinertness, demonstrating their feasibility for medical application [37]. Continuing with the analysis of biocompatibility, some researchers [34, 45] center their investigations on the interaction between blood components and PVA hydrogels. Gels obtained by annealing in the presence of glycerol display a modified surface and an increase in crystallinity. These changes lead to a reduction in platelet adhesion and protein adsorption, achieving that the blood gets protected from direct contact with the material’s surface [37, 61]. Over the years, diffusive properties of PVA hydrogels also have gained importance in the medical field. In this area, Peppas N. [62] was a pioneer who further analyzed the potential drug-delivery applications of this material. Theophylline and oxprenolol release studies were conducted using physical crosslinked PVA-H leading to the

3.7 Articular cartilage: Architecture and composition   

   41

conclusion that by controlling the freezing/thawing cycles, the drug liberation and mucoadhesive characteristics would be improved [37, 50]. Recent studies tend to modify PVA hydrogels structure blending these gels with different substances such as, chitosan, gelatin, poly(acrylic acid), polyvinylpyrrolidone (PVP) and clays, among others [34, 63–66]. The aim of these investigations had been to synthesize biomaterials as wound dressings, drug carriers, and as artificial muscles and tissues [35]. Currently, in the area of tissue engineering, the development of hydrogels for articular cartilage replacement is a topical issue. To understand and to mimic articular cartilage, it is first necessary to learn and study the structure, organization and properties of the natural cartilage.

3.6 Cartilage: A brief description Cartilage is a highly organized connective tissue adapted to resist tensile, compressive and shearing forces. It lacks the capability of self-recovery due to the absence of nerves, blood and lymphatic vessels [67]. This tissue is rigid enough to provide support for structures where elastic deformability is required: the ear, respiratory system and part of the rib cage, among others. Three varieties of cartilage can be distinguished in the human body [67–68]: – Elastic cartilage: it is found in the external ear, the epiglottis, and the auditory tube of the middle ear. Its structure is composed of proteoglycans, collagen, and elastic fibers, which make it the most flexible type of cartilage. – Fibrocartilage: is a very tough tissue due to the high proportion of collagen. It is present at the insertions of ligaments and tendons and forms the intervertebral disks. – Hyaline cartilage: is possible to find in the nose, trachea, larynx, bronchi and at the end of the ribs. It also allows the movement of one bone against another and is therefore known as articular cartilage [67–68]. In the next section this type of cartilage will be described with more detail.

3.7 Articular cartilage: Architecture and composition Articular cartilage is a specialized soft tissue that lines the ends of articulating bones, forming the smooth, low-friction surface of the diarthrodial joint (site of the junction). This tissue can be described as a tough, resilient, porous, and viscoelastic material, capable of undergo years of cyclic loading without fail [67–70]. The only resident cells of the cartilage are the chondrocytes, which synthesize the extracellular matrix (ECM). The ECM is composed of a dense network of collagen fibers, non-collagenous proteins, proteoglycans (PGs), and tissue fluid [69, 71].

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   3 Hydrogels based on poly(vinylalcohol) for cartilage replacement

There are several types of collagenous fibers. However, 90–95 % of total collagen corresponds to the type II. This class of protein, in conjunction with type I collagen, provides tensile strength maintaining tissue structure integrity. The function of the other fibrillar and globular collagen types, such as V, VI, IX, and XI, is not totally known, but they are believed to play a role in intermolecular interactions [71–72]. Between 5–10 % of the cartilage wet weight is composed of large protein-polysaccharide molecules, called proteoglycans. The agglomeration of various proteoglycans into large macromolecules is important for water absorption. They are essential to the functionality of cartilage tissue, conferring resilience and elasticity [71–73]. About 80 % of the cartilage wet weight is composed of fluid and around 30 % of the fluid would have a strong association with the collagen fibers and is thus significant for the structural organization of the ECM. This interaction with the ECM gives the ability to resist and recover from compression. The rest of the fluid can move freely during loading [74]. Articular cartilage is structured in four different zones whose composition varies between the surface and the subchondral bone plate: the superficial, the middle, the deep and the calcified zone (Fig. 3.6). The superficial zone is formed by flat and small chondrocytes that are orientated with the collagen fibers. It contains a large amount of water and collagen fibers, which run parallel to each other and to the joint surface, conferring high tensile and

Articular Surface Colagen Superficial region

Middle or Transitional region

Chondrocyte

Deep or Radial region Tride Mark

Calcified region Subchondral Bone

Fig. 3.6: Organization of articular cartilage.

3.8 Articular cartilage: Mechanical properties   

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shear strength [71]. In the middle zone, chondrocytes appear spherical and larger, and the number of proteoglycans is increased. The collagen fibers are randomly orientated. The deep zone presents the lowest amount of chondrocytes that form columns lying perpendicular to the cartilage surface [69, 71]. The collagen fibers are larger than in the other zones and they are orientated as the chondrocytes. These fibers, in conjunction with the great amount of proteoglycans, provide the cartilage’s resistance to compressive loading [71]. Finally, the lowest zone is calcified and transitions into the subchondral bone, forming an intermediate layer between the elastic cartilage and the rigid bone. The chondrocytes in this zone exhibit low metabolic activity [69, 71].

3.8 Articular cartilage: Mechanical properties Articular cartilage can be described as an anisotropic, viscoelastic, and nonlinear composite material. Basically, it is composed of a fluid-saturated, permeable, and porous matrix, which is reinforced with fibers [72, 75]. In engineering terms, the cartilage provides a wear-resistant and low-friction surface. Moreover, it allows translations and rotation between bones, giving support and load transfer. This tissue is subjected to a combination of compressive, tensile and shear stresses [72, 76]. Articular cartilage mostly experiences cycles of compressive loads during normal joint movement. Under compression, tissue fluids transpose the interconnected pore structure, generating reversible volumetric changes. The combination of the elasticity of the matrix with the fluids’ imbibition allows the tissue to recover its initial dimensions when the load is removed [77]. The cartilage compressive modulus varies from 0.08 to 2 MPa depending on the location on the joint and the depth in the tissue [72, 78]. Tension stresses are also perceived by the articular cartilage, even when it is compressed; during the tensile loading process, collagen fibers stretch and align in the direction of the load application. The tensile modulus of articular cartilages fluctuates between 5 and 25 MPa, according to its location in the body [72]. As translational and rotational movements occur in the joint, proteoglycans interact with collagen fibers, causing shear stresses within the articular cartilage. The equilibrium shear modulus for articular cartilage has been found to vary from 0.05 to 0.25 MPa [72]. Articular cartilage exhibits a low friction coefficient (0.005 approximately) [72]. Elasto-hydrodynamic lubrication, fluid pressurization, boundary lubrication and squeeze film lubrication are the possible mechanisms that justify this low value [76].

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3.9 Frequent medical issues relating to cartilage: Degeneration and osteoarthritis Articular cartilage function is affected by several factors such as genetic characteristics, aging, health and lifestyle. In this respect, cartilage damage is an increasing and persistent problem, which affects millions of people all over the world. An acute trauma, congenital or acquired joint diseases and osteoarthritis can frequently injure this tissue [23, 79]. Osteoarthritis (OA) is a chronic, progressive affection of the synovial joint which produces the thinning or loss of the cartilage. Over time, this loss causes a reduction in the joint space that may lead to the direct contact between bone ends. OA induces anatomical changes to the joint structure. Typical signs of the illness include joint pain, deformity, crepitus with motion, joint effusions and restriction of motion. Cartilage degeneration does not progress uniformly with time, and depends on the kind of joint and the individual. The most affected parts of the body are knee, spine, hip, hand joints and the foot [79–80]. Diverse techniques have been proposed to regenerate tissue function due to the articular cartilage inability to self-repair [22, 46, 60]. Autologous chondrocyte implantation, microfracture, osteochondral autograft transfer, mosaicplasty are some of the alternatives. Although in the short term these options result in a pain reduction and functional improvement, tissue function is not fully recovered. For this reason, there has been increased interest in the development of synthetic articular cartilage materials that fulfill the native tissue requirements [23, 63, 81].

3.10 Materials used as articular replacement Over the last few years, rigid materials such as ceramic, metals, and some polymers, have been used for artificial joint replacement. These traditional materials are stiffer than articular cartilage. This condition in addition to the absence of lubrication could cause excessive wear that may lead to osseous dissolution [22]. Titanium fibers are commonly used to couple articular implants with bone; however, this technique results in a painful, troublesome recovery. Other alternatives consist in implants based on alumina, silicone, carbon-reinforced polyether etheroketone (PEEK), and UHMWPE [22, 82]. Despite the success of UHMWPE as a prosthesis material, several studies have revealed that its excessive wear leads to osseous dissolve, and thus implant loss [83]. For this reason, hydrogels have been proposed as an appropriate material for articular cartilage replacement. Their viscoelastic nature, high water content, and biocompatibility allow them to mimic this human tissue even though their mechanical properties are insufficient. Hydrogels are capable of reducing stresses at the implantbone, increase contact area, encourage fluid film lubrication and lowering wear [15,

3.10 Materials used as articular replacement   

   45

84–85]. Therefore, the possibility of prosthesis failure is minimized [24, 86]. Much research has focused on polyhydroxyethyl methacrylate (PHEMA) [27, 85, 87–89] and, especially, on PVA hydrogels obtained by different methods and combined with other materials [22–23, 30, 63, 81, 90–92], in order to reach the required performance. Some of these new composite materials correspond to PVA with a stainless steel fiber mesh and bone cement, blend hydrogels based on PVA and PVP, PVA hydrogels filled with synthetic or natural substances like gelatin, carboxymethyl cellulose (CMC), starch, sodium salt of CMC, CM-chitosan, and proteolytic enzyme papain, among others [63, 80, 93–94]. In recent years, PVA cryogels have increasingly attracted interest as a biomaterial for artificial articular cartilage. These gels show excellent properties such as high elasticity, microporous structure, good biocompatibility, hydrophilicity, high elastic modulus, good absorption and exudation of body fluids, and exceptional biotribological properties [24, 95]. Their availability and low processing cost are also some of the main advantages of this material. All these characteristics, together with its non-degradability within the body, make PVA-H the optimum alternative for tissue replacement [92, 96]. Poor mechanical properties and non-bioactivity are the major limitations of PVA cryogels for this application [39, 90]. Non-bioactivity avoids the adhesion between the gel and the bone, thus the fixation method is a very important problem that needs to be resolved. In this sense, hydroxyapatite (HA, Ca10(PO4)6(OH)2 ) has gained a great deal of attention due to its bioactivity, osteoconductivity [97–98], biocompatibility and osteoproductivity [99]. The HA is a calcium phosphate, which crystallizes in a hexagonal system. It is chemically similar to the mineral component of natural bone (containing 69 wt% of calcium phosphate) and other hard tissues in mammals. Hydroxyapatite has been used in several medical applications due to its ability to bond living bony tissue [100]. The incorporation of nano-HA to the PVA hydrogels enhances mechanical properties and promotes the implant’s osteointegration [24, 96, 99]. Several investigations have focused on the influence of different contents of HA on the properties of PVA/HA composite cryogels. The results show that the composites with 3 wt% of HA content display the optimum properties for articular cartilage replacement [90, 95]. PVA hydrogels with an HA gradient are also being evaluated. The final goal is synthesizing a material where one side is high HA content, and the opposite side contains a great amount of free water, improving its lubrication function [101]. Researchers began to look for new methods in order to achieve adequate mechanical properties for PVA/HA composite implants. Therefore, UHMWPE and polypropylene (PP) fibers are being tested as PVA hydrogel reinforcement. These novel composites successfully increase the tensile modulus of the biomaterial, simulating the collagen network structure within the cartilage [102].

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Conclusions The results included in the present chapter demonstrated that hydrogels based on PVA are promising materials for the development of cartilage replacement. Although several contributions were made in this area; there is still a lot of work to do since limitations continue to appear, especially when it comes to clinical applications of these materials.

Acknowledgments We appreciate funding by CONICET (National Scientific and Technical Research Council), UNMdP (National University of Mar del Plata) and ANPCyT (National Agency for Science and Technology Promotion).

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S. Arun, P. S. Rama Sreekanth and S.Kanagaraj

4 Polymer composites for cemented total hip replacements 4.1 Introduction The Total Hip Replacement (THR) plays an important role for people having hiprelated problems associated with age, accident, osteoarthritis, wear of acetabular etc. Statistics have shown that more than 230000 total hip arthroplasty (THA) and 490000 total knee arthroplasty (TKA) were performed annually in USA alone [41], and the number of total joint replacement (TJRs) is expected to increase substantially over the next two decades [42]. Though it was originally intended for older patients, THA is now being used in younger and much more active patients, who require prostheses that can last longer and endure a much more active lifestyle without the need for a revision. Thus, the design goals have advanced from simply alleviating pain and discomfort for older patients to returning to physically demanding lifestyles including sports activities for young active patients [43]. Though different materials are used in THR, the artificial hip prosthesis implanted in the patient is not a once-in-a-lifetime procedure; the chances of revised surgeries may vary depending on the lifestyle, activities and other post-operative care of the patient. Although there are many possible reasons for revision surgeries, the aseptic loosening (primarily due to micromotion between implant and bone cement interface), implant breakage and wear of acetabular caused 58 % of total revisions [4]. The chances of revised surgeries can be minimized if the material-related problems associated with acetabular cup and bone cement can be addressed. This chapter henceforth discusses the various problems and developments related to polymer materials used as an acetabular cup (UHMWPE) and bone cement (PMMA).

4.1.1 Understanding hip joint prosthesis and fixation techniques The hip joint is one of the load-bearing joints besides knee and ankle joints. Most of the other joints such as elbow, shoulder, wrist etc. pertain to the non-load-bearing class. The joints may vary in their shape, size and loading pattern but their typical lubrication mechanism remains the same. The joints may lose their natural lubricating function due to accidents or some degenerative joint diseases like arthritis, or due to excessive wear leading to severe pain in the joint location. Now, the patient needs to be operated for joint replacement with artificial implant materials paving the way for TJR surgeries. Although TJRs are common in cases of knee, shoulder, elbow, wrist, ankle and finger joints, these are the recent developments compared to

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that of hip joints. The details of the different parts used in total hip replacement are shown in Fig. 4.1. A total hip prosthesis consists of four major components namely femoral stem, femoral head, acetabular cup and metallic back-up. The acetabular cup is fitted into the metallic back-up either by mechanical fixation or by using PMMA bone cement, [44]. The femur bone is sectioned at the greater trochanter at a suitable angle and the femoral stem is fixed either by press fit or cemented fixation, [45]. The femoral head is fixed on the stem and then anatomically positioned into the acetabular cup.

Acetabular Cup Outer Shell METAL Liner POLYETHYLENE Femoral Head METAL Femoral Stem METAL

Fig. 4.1: Assembly of artificial hip prosthesis [1].

It should be noted that bone cement was used for fixation of either femoral stem in the femur or to fix the acetabulur cup into the metallic back-up or both in the THR. The history of bone cement began in the early 20th century with Ottorohm, who synthesized PMMA and used it for dental applications. Later, Charnley used bone cement for orthopedic surgeries and for filling the medullar cavity of the bone and the fixation of total joint prosthesis during the 1950s. The fixation was done by two methods: mechanical fixation (cementless fixation), and cemented fixation, which are shown in Fig. 4.2a and 4.2b. It was observed that there was an interface between stem and bone made of bone cement in the case of cemented fixation. In cementless fixation, a porous coating was used to increase the stability of the stem. Wahab et al. [46] studied the usage of uncemented Total Hip Replacement (UTHR) and found that it was suitable for younger patients. Refior et al. [47] found that the usage of cementless fixation was limited to patients having an irregular shaped femur and large medullar cavi-

4.1 Introduction   

   55

ties where it required less operative time, whereas the usage of cemented Total Hip Replacement (CTHR) for older patients was necessary to stabilize the THR. Unnanuntana et al. [48] compared the effect of cemented and cementless fixation and found that the life of cemented implant fixation was less compared to that of cementless fixation and there was a higher risk of revision in male than the female gender. The overall comparison between CTHR and UTHR suggested that the number of revised surgeries was more in CTHR than the UTHR. After surgery, the weight-bearing can begin on the next day with crutches for CTHR and walking can be done within two days after surgery, but the weight-bearing of UTHR was based on its stability.

a

b

Fig. 4.2: (a) Uncemented total hip replacement. (b) Cemented total hip replacement [2].

The comparison of cemented and cementless fixation is shown in Tab. 4.1. It compares different aspects such as: life of the implant, number of revised surgeries, age of the patient and cost of the surgery etc. Tab. 4.1: Comparison of cemented and cementless THR [3]

Parameters

Cemented fixation

Cementless fixation

Life of the implant Number of revised surgeries Initial walk Normal walk

Life of THR is less More

Life of THR is more Less

Pain relief Cost of the surgery Physical factors Age limit

Within 2 days Within 5 days

Within 6–12 weeks Varies from 1 month to a year based on the stability Within 5 days Based on the stability Less Comparatively high Suitable for patients with irregular Not suitable for patients with irregular medullar cavities and voids medullar cavities and voids Can be used above 65 years Encouraged below 65 years of the patient

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4.1.2 Economic and clinical factors surrounding revision surgeries The failure of acetabular cup and the bone cement fixation of the implant led to the revision of THR. UHMWPE has been in use as one of the materials for acetabular cup due to its chemical resistance, lubricity and extraordinary resistance to wear, which are all crucial parameters for the longevity of the implant. However, the number of revised surgeries being performed due to wear, osteolysis, oxidation of material, infection etc. is a major concern among the materials science and medical community. Wear has been accepted as a major cause of osteolysis in THA. Submicron size wear particles migrated into the effective joint space and stimulated a foreign-body response resulting in bone loss [49]. In many patients, the bone around the implant became worn away leading to reduced joint function, which was caused by wear debris induced osteolysis in the implant [50]. The wear debris generated from the acetabular cup against femoral head causes serious damage to the living tissues. Though UHMWPE is biologically an inert material, the presence of large amounts of biologically inactive particulate materials can trigger an exceptionally strong macrophage response leading to osteolysis around the implant in general and bone–implant interfaces in particular [51]. The various factors that enforce the revision surgeries of total hip replacement are shown in Fig. 4.3, where it is observed that the aseptic loosening was found to be the leading factor along with dislocation, osteolysis, fracture, infection etc. Though PMMA is being used as a bone cement, it has inherent limitations such as lack of bioactivity, radiolucent properties, dimensional change, and poor thermal and mechanical properties. Arens et al. [52] increased the bioactivity of bone cement by

10,5 %

7,3 % 6,8 %

Aseptic loosening

3,5 % 2% 1,7 % 0,8 % 1,5 %

13,5 %

Dislocations Osteolysis Fracture Infection Implant breakage acetabular wear acetabular Pain Implant breakage stem others

52,4 % Fig. 4.3: Various factors leading to revision of hip surgeries [4].

4.2 UHMWPE composites   

   57

combining PMMA with freshly harvested bone marrow from sheep, which paved the way to use PMMA-cement in cancellous bone augmentation for osteoporotic patients with increased bioactivity. Urabe et al. [53] evaluated the bactericidal activity, and the profile of eluates by adding Vancomycin (VCM) with calcium phosphate (CP) and VCM with PMMA. Eluates obtained on 1, 3, 7, and 14 days and 4, 8, and 12 weeks were evaluated by high-performance liquid chromatography (HPLC) and by microbiological assay (MBA). It was found that CP composite released more VCM over a longer period than PMMA/VCM composite. Makita et al. [54] analyzed the changes in radiopacity and strength of PMMA bone cement with the addition of various concentration of BaSO4 such as 10, 20, 30 and 40 wt%. Though radiopacity of the composites was increased with BaSO4 concentration, the strength of the same was observed to be reduced. Nuno et al. [55] measured the transient and residual stresses and peak temperature during the polymerization of PMMA. The peak temperature of 70 °C was observed during the curing process and the maximum radial residual stresses of 0.6 MPa in compression confirmed the shrinkage of bone cement towards the stem. From the literature, it is observed that various factors including bioactivity, wear debris generation and the dimensional stability led to revising the primary surgeries of TJRs. An increase in the number of revised surgeries enforces increased hospital infrastructure, government burden, professional burden and total expenditure for the joint replacement surgeries. According to Kim [56], the total annual hospital cost for primary and revision hip and knee arthroplasty procedures in the USA was estimated to be $9.1 billion in 2004 and it is expected to exceed $80 billion by 2015. The range of price difference between CTHR and UTHR varied from $179 to $1311 and it is expensive for the cementless stem [48]. Lavernia et al. [57] showed that 59 % of total expenditure on THA surgery went to hospital facilities and infrastructure. An increase in TJRs is not only a burden to the patient but also to the government and medical care units in order to meet the economic factors surrounding the THR and provide the necessary facilities to improve patient comfort levels. Based on the above discussion, it is observed that although CTHR has more chances of revision, nonetheless it has specific advantages over UTHR. A combination of good material for acetabular cup with increased longevity and improved properties of bone cement can lead to greater success of THR. Several researchers have attempted to improve the properties of UHMWPE and PMMA by reinforcing them with different fillers. The associated problems and the enhancement of the properties of UHMWPE composites and PMMA bone cement composites are discussed in detail in the subsequent sections.

4.2 UHMWPE composites Several research groups have attempted to increase the mechanical properties such as tensile strength, strain at fracture, yield strength and fatigue strength of UHMWPE by reinforcing different filler materials. The fillers are also aimed to improve the antimi-

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crobial and antioxidant properties and wear resistance of UHMWPE. Although other techniques such as radiation crosslinking and vitamin E blending are available, both of them have limitations such as oxidation of material, reduction of material properties and inhibition of crosslinking. Thus UHMWPE composites with suitable fillers as reinforcement are being explored. The selection of a reinforcing material used to prepare the composites is based on the requirement of the final product in biomedical applications. This section is devoted to presenting and discussing the potential of UHMWPE composites to improve mechanical properties, wear resistance, infection retardation and biocompatibility using various fillers as reinforcing materials. A variety of fillers and blends used with UHMWPE is given in Tab. 4.2. The mechanical properties and tribological characteristics of several composites are discussed below.

Tab. 4.2: List of different fillers/blends used with UHMWPE

Filler/Blend

Reference

UHMWPE fibers Quartz High density polyethylene (HDPE) Carbon nanotube (CNT)

Cohen et al. [5], Deng et al. [6] Liu et al. [7], Xie et al. [8] Jacobs et al. [9], Lucas et al. [10]

Carbon fiber Hydroxyapatite (HA) Wollastonite Coconut shell Kaolin Epoxy Silver Zirconium Natural coral

Ruan et al. [11], Xue et al. [12], Bakshi et al. [13, 14], Morlanes et al. [15], Samad et al. [16, 17], Maksimkin et al. [18], Fonseca et al. [19, 20] Sreekanth et al. [21, 22, 23] Farling [24], Wood et al. [25], Dangsheng [26], Sui et al. [27] Fang et al. [28, 29], Long et al. [30], Wang et al. [31], Xiong et al. [32] Tong et al. [33] Pradhan et al. [34] Guofang et al. [35] Chand et al. [36] Morley et al. [37] Schwartz et al. [38], Plumlee et al. [39] Ge et al. [40]

UHMWPE was reinforced with a variety of fillers to achieve the desired properties. However, none of them were a clinical success. Although carbon fiber reinforced UHMWPE was clinically used due to its superior wear behavior [24], it was later withdrawn due to its reduced ductility, decreased crack resistance, and the fiber-matrix interface of the composite. Tong et al. [33] reinforced UHMWPE with wollastonite fibers having the aspect ratio of 10:1, 15:1 and 20:1. It was observed that the addition of wollastonite fibers up to 20 wt% improved the abrasive wear resistance of UHMWPE significantly. Xie et al. [8] investigated the physical, physiological and mechanical properties and the wear resistance of UHMWPE/quartz composites, which were found to increase with concentration of the reinforcement. They also reported the mild toxicity of wear debris generated from the composites. Guofang et al. [35] studied the tribo-

4.2 UHMWPE composites   

   59

logical characteristics of kaolin filled UHMWPE composites obtained by different processing technique under dry condition. It was concluded that the wear rate was much lower for the polymerization-based composites than the melt-mixing process, because of the smaller crystal size and strong interface, but a reverse trend was observed in the case of strength of the composites. Fang et al. [28] reported that the chain mobility of UHMWPE and UHMWPE/HA composites was increased by swelling them in paraffin oil and the resulting composite showed a 90 and 50 % increase in Young’s modulus and yield strength, respectively, compared to that of unfilled UHMWPE. Fang et al. [29] observed that the tensile strength of UHMWPE/HA composite was increased by 5 %, while the fracture strain was reduced by 22 % compared to that of virgin UHMWPE at 22.8 vol.% of reinforcement. It was also reported that the stiffness of the composite was closer to the lower bound value of cortical bone. Xiong et al. [32] reinforced UHMWPE with nanocrystalline hydroxyapatite up to 7 wt%. The hardness of the composites was found to be maximum at 1 wt% of filler content and 50 kGy irradiation dose. The wear volume of 1 wt% nanocomposites was reduced by 27 and 81 % at 0 and 150 kGy irradiation doses, respectively. Ge et al. [40] used 5–30 wt% of natural coral particles of < 50 μm size as a reinforcement with UHMWPE and studied the wear resistance of the composites in a hip joint simulator lubricated with bovine serum against CoCrMo femoral head. It was reported that the hardness and wear resistance of 30 % natural coral composite were increased by 95 and 70 %, respectively as compared to virgin UHMWPE. They concluded that the reinforcement of natural coral in UHMWPE changed the severity of adhesive wear of UHMWPE acetabular cups and it generated a lesser number of microcracks on the worn surface. Plumlee et al. [39] prepared UHMWPE/Zr composites with 10 and 20 wt% Zr. A dual axis wear simulator was used to study the wear behavior of the composites for 250000 cycles. They observed that the wear volume of the composites was reduced with an increase of Zr concentration. It was concluded that the optimum inclusion of 10 wt% Zirconium particles into UHMWPE matrix effectively reduced the wear rate of the component without the reduction of impact toughness. Morley et al. [37] studied tribological characteristics of UHMWPE/Ag nanocomposites using a horizontal tribometer under synovial fluid condition with different sliding velocities. It was reported that no detectable variation of either temperature or friction coefficient was observed over the range of sliding velocity (10–80 mm/s). Wood et al. [25] observed improved wear strength and mechanical properties of UHMWPE–carbon nanofibers (CNF) composites. Though the stiffness of the polymer was not influenced by the reinforcement, the yield load, fracture load and toughness of the composites were increased by 13, 10.7 and 10.8 %, respectively at 3 wt% of CNF in UHMWPE. The static coefficient of friction was also reported to be reduced by 78 %. Ruan et al. [11] reported that the Young’s modulus and yield stress of UHMWPE were increased by 38 and 50 %, respectively by reinforcing 1 wt% multi-walled carbon nanotubes (MWCNTs). A significant increase of ductility and ultimate tensile strength of the composite was also observed. Xue et al. [12] blended UHMWPE (80 %) with HDPE

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(20 %) and the blend was reinforced by MWCNTs (0.2–2.0 wt%) in order to improve the creep and wear resistance of the test sample. It was concluded that the presence of MWCNTs did not show any significant influence on the creep behavior but the wear resistance of the composites was found to be improved. The yield strength and Young’s modulus of the blend were enhanced by 20 and 26 %, respectively at 2 wt% MWCNTs in the polymer blend. Bakshi et al. [13] tested UHMWPE/MWCNTs (5 wt%) nanocomposite films prepared by electrostatic spraying technique. The Young’s modulus of the composite was found to be increased by 82 %, but failure stress and strain were decreased by 13 and 64 %, respectively. Bakshi et al. [14] also evaluated the UHMWPE/MWCNTs nanocomposites using a Nanoindenter, where it was reported that the hardness and Young’s modulus of the nanocomposites were increased by 11 and 10 %, respectively at 5 wt% of MWCNTs in UHMWPE. It was also observed that the plasticity index of the polymer was not influenced by the presence of MWCNTs. Morlanes et al. [15] irradiated 3 wt% UHMWPE/MWCNTs composites at 90 kGy in air and it was observed that the Young’s modulus, yield stress and fracture stress of composites were increased by 34, 10 and 4 %, respectively, whereas the fracture strain was reduced by 44 % compared to that of unirradiated composites. Fonseca et al. [20] reported that the Young’s modulus, strain at fracture and toughness of 0.2 wt% of MWCNTs in UHMWPE were increased by 80, 26 and 35 %, respectively. Maksimkin et al. [18] reinforced UHMWPE with 1 wt% MWCNTs and reported that the ultimate strength and yield strength of the composites were increased by 286 and 35 %, respectively compared to that of virgin UHMWPE. Tribological tests showed the decrease of dry friction coefficient from 0.24 to 0.14 for 1 wt% UHMWPE/MWCNTs composites. Sreekanth et al. [21] reinforced UHMWPE with MWCNTs up to 5 wt% and reported that the optimum concentration of MWCNTs in UHMWPE was found to be 2 wt%, where the work to failure, fracture stress, strain at fracture, and yield stress of the medical grade UHMWPE were enhanced by 176, 93, 70 and 44 %, respectively. The hardness of virgin UHMWPE was found to be enhanced by 75 % and the plasticity index was reduced by 17 % for 2 wt% of MWCNTs, while the Young’s modulus was enhanced by 170 % for the UHMWPE/MWCNTs nanocomposites. Sreekanth et al. [22] also reported that the presence of MWCNTs reduced the irradiation induced oxidation of UHMWPE. The biocompatibility of UHMWPE/MWCNTs nanocomposites was also confirmed by Reis et al. [58], Ormbsy et al. [59] and Sreekanth et al [23].

4.3 PMMA composites Nguyen et al. [60] characterized the fracture toughness and fatigue crack growth behavior of PMMA bone cement in air and Ringer’s solution. It was observed that the cement was found to be more resistant to fatigue-crack propagation in Ringer’s solution than air. Walker et al. [61] designed an automated mechanical mixing device to prepare bone cement and it was observed that the bending modulus of the system-

4.3 PMMA composites   

   61

produced cement was found to be increased by 16.6 % compared to the conventional mixer-produced sample. Chaplin et al. [62] recovered the PMMA bone cement sample used for the fixation of hip prostheses and examined its density, porosity and microhardness. It was reported that the Young’s modulus of the recovered sample was found to be less than that of the freshly prepared PMMA cement. It was also reported that the modulus of elasticity and micro-hardness of the specimens were increased with PMMA density. Giddings et al. [63] used the small punch test to characterize the elastic modulus and fracture behavior of PMMA at room temperature and body temperature. It was observed that the specimens exhibited ductile crack initiation at body temperature, whereas brittle crack initiation was observed on the sample at room temperature. Pelletier et al. [64] investigated the mechanical properties and pore distribution of 10 commercially available medical grade cements cured in the molds at 20, 37, 40 and 50 °C. An increase of mechanical properties was found in the bone cement cured at 50 °C. It was found that the pores were observed to gather near the surface of cooler molds and near the center of warmer molds for all cement brands. Small pores were more often present in cements cured at lower temperatures, whereas higher temperature molds produced more large sized pores. Serbetci et al. [65] studied the mechanical, thermal and biological properties of PMMA by mixing various concentration of hydroxyapatite (HA). The optimum value of HA was found to be 14.3 wt%, where the fatigue life of the composite was found to be increased significantly. Zamarron et al. [66] increased the bioactivity of PMMA bone cement by varying the wollastonite content by soaking the samples in a simulated body fluid (SBF) with an ionic concentration nearly equal to that of human blood plasma. It was reported that the cement with 39 wt% of wollastonite showed both bioactivity and improved compressive strength. Fukudu et al. [67] examined the in vivo bone bonding strength of PMMA bone cement by adding the Titania particles with 10, 20, and 30 wt%. It was reported that the mechanical and thermal properties of Titania–PMMA composite cements were substantially higher than those of commercially available PMMA cement. Silva et al. [68] studied the gamma irradiation effect on the PMMA/HA and PMMA/seaweed residues (SW) composites and it was reported that the presence of HA, and SW in PMMA controlled the irradiation induced degradation rate successfully. Kurtz et al. [69] studied the mechanical behavior of PMMA bone cement with elevated barium sulfate content, where the properties were found to be decreased beyond 36% of barium sulphate/PMMA. Ginebra et al. [70] used three different types of radiopacifying agents, two of them were inorganic (BaSO4, ZrO2) and one was organic (an iodine containing monomer, IHQM). The Zirconia added bone cement showed an improvement of tensile strength along with the enhancement of fracture toughness and fatigue crack propagation resistance. However, the bone cement with iodine containing monomer recorded an increase of tensile strength and fracture toughness but the fatigue crack propagation resistance was found to be decreased. Kotha et al. [71] evaluated the mechanical properties of PMMA reinforced with short Titanium fiber. The fracture toughness of composites was found to be increased by 30 % by the

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addition of fibers. Kane et al. [72] compared the fatigue life of PMMA reinforced with straight and variable diameter Zirconia fibers (VDF). The composites were prepared with 5, 10, 15, and 20 vol. % of the filler material using vacuum-mixing technique. The optimum concentration of reinforcement for maximum fatigue life was found to be 10 vol. % of filler. Moreover, the VDF reinforced composites had a fatigue life of a greater magnitude when compared to that of straight fiber reinforcement. Lopes et al. [73] evaluated the mechanical properties of bioactive glass (3CaO–P2O5–MgO–SiO2) reinforced with PMMA. The composites were prepared by melt-mixing technique with 0, 30, 40 and 50 wt% of bioactive glass. The maximum flexural strength and the elastic modulus of composite were found at 30 wt% bioactive glass in PMMA. Chan et al. [74] analyzed the mechanical properties of nanoclay-reinforced PMMA. The composites were prepared with 1, 2, 4, 5 and 7 wt% of nanoclay and the mechanical properties of nanocomposites were found to be optimum at 5 wt% of nanoclay, where the Young’s modulus and tensile strength of nanocomposites were found to be increased by 34 and 25 %, respectively, compared to that of pure PMMA. Khaled et al. [75] reinforced Titania nanotubes with PMMA to study the effect of mechanical properties. The functionalized nanotubes were dispersed by ultrasonic agitation method for making the nanocomposites of 1, 2, 4 and 6 wt% of Titania nanotubes. The fracture toughness and flexural strength of the composites were found to be increased by 20 and 40 %, respectively at an optimum concentration of 2 wt% Titania nanotubes. Jia et al. [76] reinforced the PMMA by MWCNTs and studied the mechanical properties of the nanocomposites. The mechanical properties of the untreated MWCNTs with PMMA were found to be decreased whereas the nanocomposites prepared with treated MWCNTs showed enhanced mechanical properties. Coleman et al. [77] reviewed the PMMA/MWCNT composites and found that the method of processing of nanocomposites such as solution-processing method, melt-processing technique and rotary-mixing technique was an important factor in making the nanocomposites. The best method was found to be the solution-dispersing technique for pellet raw materials and the rotary-mixing technique for powder raw materials. Ormsby et al. [78] reinforced PMMA with 0.1 wt% of MWCNTs and studied the effect of different dispersion methods such as magnetically stirring, dry blending using a small-scale turbo blender, and ultrasonic disintegrator. It was concluded that the dry blender showed proper dispersion for optimum mechanical properties. Lee et al. [79] proposed an injection molding and film casting processes to make PMMA/MWCNTs nanocomposites. The nanocomposites made by the injection molding process showed reduced mechanical properties compared to those of virgin PMMA. However, the tensile strength of the sample prepared by the film casting process was found to be increased by 17 % for 0.01 wt% of MWCNTs. Ormsby et al. [80] reinforced the PMMA with various concentration of MWCNTs such as 0.1, 0.25, 0.5, 0.75 and 1.0 %, and investigated the effect of unfunctionalized, carboxyl and amine functionalized MWCNTs reinforced PMMA. It was observed that the mechanical properties of composites were found to be increased in carboxyl and amine functionalized nanocomposites, where the carboxyl

Future scope   

   63

functionalized MWCNTs provided significant improvement in mechanical properties. Nien et al. [81] reinforced PMMA with MWCNTs having 0.1, 0.2, 0.27, 0.43, 0.59 and 0.75 wt% of MWCNTs. The tensile and compressive strength of the composites were found to be increased by 18 and 23 %, respectively at an optimum concentration of 0.27 wt% of MWCNTs. The studies related to the temperature raise and volumetric shrinkage of PMMA were found to be very limited. Moreover, the usage of MWCNTs and single-walled carbon nanotubes (SWCNTs) in the enhancement of mechanical properties as well as the exothermic parameters has not been documented to date. These results are still to be explored in detail.

Summary It is clearly understood from the above studies that UHMWPE and PMMA composites created a consistent place in the THA cemented fixation technique. The usage of UHMWPE composites has considerably increased the wear resistance and toughness leading to increased life of the implants. The issues related to stress shielding, fretting wear between the stem-cement interface, and temperature raise and volume shrinkage of PMMA during its polymerization were found to be decreased for the PMMA composites. Recent studies have shown that both UHMWPE and PMMA reinforced by MWCNTs have exhibited superior properties and also possess reasonable biocompatibility.

Future scope MWCNTs composites should be explored more diligently with specific focus on in vivo biocompatibility and hip joint simulator studies. Although several authors reported on the enhancement of wear resistance of the composites, very few of them have explored the wear testing in hip joint simulators. An attempt in that direction will help to estimate the longevity of developed composites. Fatigue studies and estimation of other mechanical properties of the composites at human body temperature and physiological conditions have been focused on least by the researchers. Such a study will enable to estimate the strength of the material and address the root problems in a systematic way. Rheological characteristics and dimensional stability of PMMA/MWCNTs composites have also not been thoroughly reported. The exothermic parameters such as temperature raise and volumetric shrinkage of PMMA composites bone cement should be studied in detail in order to explore them for its applications.

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[21] Sreekanth PSR, Kanagaraj S. Assessment of surface and bulk properties of UHMWPE/MWCNT nanocomposites using nanoindentation and microtensile testing. Journal of Mechanical Behavior of Biomedical Materials 2013, 18, 140–151. [22] Sreekanth PSR, Kanagaraj S. Restricting the ageing degradation of the mechanical properties of gamma irradiated UHMWPE/MWCNT nanocomposites. Journal of Mechanical Behavior of Biomedical Materials 2013, 21, 57–66. [23] Sreekanth PSR, Kanagaraj S. Biocompatibility studies on MWCNT reinforced ultrahigh molecular weight polyethylene nanocomposites. Trends in Biomaterials and Artificial Organs 2013, 27 (1), 1–9. [24] Farling G. Human body implant of graphitic carbon fiber reinforced ultra–high molecular weight polyethylene. United States Patent No. 4,055,862, 1977. [25] Wood WJ, Maguire RG, Zhong WH. Improved wear and mechanical properties of UHMWPEcarbon nanofibers composites through an optimized paraffin-assisted melt-mixing process. Composites, Part B 2011, 42, 584–591. [26] Dangsheng X. Friction and wear properties of UHMWPE composites reinforced with carbon fiber. Materials Letters 2005, 59, 175–179. [27] Sui G, Zhong WH, Ren X, Wang XQ, Yang XP. Structure, mechanical properties and friction behavior of UHMWPE/HDPE/carbon nanofibers. Materials Chemistry and Physics 2009, 115, 404–412. [28] Fang L, Leng Y, Gao P. Processing of hydroxyapatite reinforced ultrahigh molecular weight polyethylene for biomedical applications. Biomaterials 2005, 26, 3471–3478. [29] Fang L, Leng Y, Gao P. Processing and mechanical properties of HA/UHMWPE nanocomposites. Biomaterials 2006, 27, 3701–3707. [30] Long LJ, Yuan ZY, Qingliang W, Shirong GE. Biotribological behavior of ultrahigh molecular weight polyethylene composites containing bovine bone hydroxyapatite. Journal China University of Mining & Technology 2008, 18, 0606–0612. [31] Wang Q, Liu J, Ge S. Study on biotribological behavior of the combined joint of CoCrMo and UHMWPE/BHA composite in a hip joint simulator. Journal of Bionic Engineering 2009a, 6, 378–386. [32] Xiong L, Xiong D, Jin J. Study on tribological properties of irradiated crosslinking UHMWPE nano-composite. Journal of Bionic Engineering 2009, 6, 7–13. [33] Tong J, Ma Y, Jiang M. Effects of the wollastonite fiber modification on the sliding wear behavior of the UHMWPE composites. Wear 2003, 255, 734–741. [34] Pradhan SK, Dwarakadasa ES, Reucroft PJ. Processing and characterization of coconut shell powder filled UHMWPE. Materials Science and Engineering A 2004, 36, 757–762. [35] Guofang G, Huayong Y, Xin F. Tribological properties of kaolin filled UHMWPE composites in unlubricated sliding. Wear 2004, 25, 688–694. [36] Chand N, Dwivedi UK, Sharma MK. Development and tribological behaviour of UHMWPE filled epoxy gradient composites. Wear 2007, 262,184–190. [37] Morley KS, Webb PB, Tokareva NV, Krasnov AP, Popov VK, Zhang J, Roberts CJ, Howdle SM. Synthesis and characterisation of advanced UHMWPE/silver nanocomposites for biomedical applications. European Polymer Journal 2007, 43, 307–314. [38] Schwartz CJ, Bahadur S, Mallapragada SK. Effect of crosslinking and Pt–Zr quasicrystal fillers on the mechanical properties and wear resistance of UHMWPE for use in artificial joints. Wear 2007, 263, 1072–1080. [39] Plumlee K, Schwartz CJ. Improved wear resistance of orthopaedic UHMWPE by reinforcement with zirconium particles. Wear 2009, 267, 710–717. [40] Ge S, S Wang, X Huang. Increasing the wear resistance of UHMWPE acetabular cups by adding natural biocompatible particles. Wear 2009, 267, 770–776.

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[41] Kurtz SM, Ong K, Lau E, Halpern M. Projections of primary and revision hip and knee arthroplasty in the United States from 2005 to 2030. Journal of Bone & Joint Surgeries 2007, 89, 780–785. [42] Kurtz SM, Mowat F, Ong K. Prevalence of primary and revision total hip and knee arthroplasty in the United States from 1990 through 2002. Journal of Bone &Joint Surgeries 2005, 87, 1487–1497. [43] Lewis G. Properties of crosslinked ultra-high-molecular-weight polyethylene. Biomaterials 2001, 22, 371–401. [44] Ritter MA. All polyethylene versus metal backing. Clinical Orthopedics and Related Research 1995, 311, 69–75. [45] Ducheyne P. Prosthesis fixation for orthopedics. In: Encyclopedia of medical devices and instrumentation. Ed. J. G. Webster, Wiley–Interscience, New York, 1988. [46] Wahab A, Quinlan JF, Sherif S, Kelly IP. Are cementless asetabular components contra-indicated in the elderly?. Eur J Orthop Surg Traumatol 2007, 17, 263–266. [47] Refior HJ, Parhofer R, Ungethiim M, Mer WB. Special problems of cementless fixation of total hip-joint endoprostheses with reference to the PM type. Arch Orthop Trauma Surg 1988, 107, 158–171. [48] Unnanuntana A, Dimitroulias A, Bolognesi MP, Hwang KL, Goodman SB, Marcus RE. Cementless femoral prostheses cost more to implant than cemented femoral prostheses. Clin Orthop Relat Res 2009, 467, 1546–1551. [49] Zhu YH, Chiu KY, Tang WM. Review Article, Polyethylene wear and osteolysis in total hip arthroplasty. Journal of Orthopaedic Surgeries 2001, 9, 91–99. [50] Amstutz H, Campbell P, Kossovsky N, Clarke I. Mechanism and clinical significance of wear debris-induced osteolysis. Clinical Orthopaedics and Related Research 1992, 276, 7–18. [51] Peter DF, Campbell PA, Amstutz HC. Metal versus polyethylene wear particles in total hip replacements, A review. Clinical Orthopaedics & Related Research 1996, 329, S206–S216. [52] Arens D, Rothstock S, Windolf M, Boger A. Bone marrow modified acrylic bone cement for augmentation of osteoporotic cancellous bone. Journal of the Mechanical Behavior of Biomedical Materials 2011,4 2081–2089. [53] Urabe K, Naruse K, Hattori H, Hirano M, Uchida K, Onuma K, Park HJ, Itoman M. In vitro comparison of elution characteristics of vancomycin from calcium phosphate cement and polymethylmethacrylate. J Orthop Sci 2009, 14, 784–793. [54] Makita M, Yamakado K, Nakatsuka A, Takaki H, Inaba T, Oshima F, Takeda HKK. Effects of barium concentration on the radiopacity and biomechanics of bone cement, experimental study. Radiat Med 2008, 26, 533–538. [55] Nuno N, Madrala A, Plamondon D. Measurement of transient and residual stresses during polymerization of bone cement for cemented hip implants. Journal of Biomechanics 2008, 41, 2605–2611. [56] Kim S. Changes in surgical loads and economic burden of hip and knee replacements in the US, 1997–2004. Arthritis and Rheumatism 2008, 59(4), 481–488. [57] Lavernia CJ, Hernandez VH, Rossi MD. Payment analysis of total hip replacement. Current Opinion in Orthopedics 2007, 18, 23–27. [58] Reis J, Kanagaraj S, Fonseca A, Mathew MT, Silva FC, Potes J, Oliveira MSA, Simoes JAO. In vitro studies on multiwalled carbon nanotubes/ultrahigh molecular weight polyethylene nano composites. Brazilian Journal of Medical and Biological Research 2010, 43, 476–482. [59] Ormsby R, McNally T, O’Hare P, Burke G, Mitchell C, Dunne N. Fatigue and biocompatibility properties of a poly (methyl methacrylate) bone cement with multi–walled carbon nanotubes. Acta Biomaterialia 2012, 8, 1201–1212.

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[60] Nguyen NC, Maloneyà WJ, Dauskardt RH. Reliability of PMMA bone cement fixation, fracture and fatigue crack–growth behavior. Journal of Materials Science, Materials in Medicine 1997, 8, 473–483. [61] Walker GM, Daly C, Dunne NJ, Orr JF. Liquid monomer–powder particle interaction in acrylic bone cement. Chemical Engineering Journal 2008, 139, 489–494. [62] Chaplin RPS, Lee AJC, Hooper RM, Clarke M. The mechanical properties of recovered PMMA bone cement, A preliminary study. J Mater Sci Mater Med 2006, 17, 1433–1448. [63] Giddings VL, Kurtz SM, Jewett CW, Foulds JR, Edidin AA. A small punch test technique for characterizing the elastic modulus and fracture behavior of PMMA bone cement used in total joint replacement. Biomaterials 2001, 22, 1875–1881. [64] Pelletier MH, Lau ACB, Smitham PJ, Nielsen G, Walsh WR. Pore distribution and material properties of bone cement cured at different temperatures. Acta Biomaterialia 2010, 6, 886–891. [65] Serbetci K, Korkusuz F, Hasirci N. Thermal and mechanical properties of Hydroxyapatite impregnated acrylic bone cements. Polymer Testing 2004, 23, 45–155. [66] Zamarron DR, Hernandez DAC, Aragon LB, Lara WO. Mechanical properties and apatite-forming ability of PMMA bone cements. Materials and Design 2009, 30, 3318–3324. [67] Fukuda C, Goto K, Imamura M, Neo M, Nakamura T. Bone bonding ability and handling properties of a titania–polymethylmethacrylate (PMMA) composite bioactive bone cement modified with a unique PMMA powder. Acta Biomaterialia 2011, 7, 3595–3600. [68] Silva P, Albano C, Perera R, Dominguez N. Study of the gamma irradiation effects on the PMMA/HA and PMMA/SW. Radiation Physics and Chemistry 2010, 79, 358–361. [69] Kurtz SM, Villarraga ML, Zhao K, Edidin AA. Static and fatigue mechanical behavior of bone cement with elevated barium sulfate content for treatment of vertebral compression fractures. Biomaterials 2005, 26, 3699–3712. [70] Ginebra MP, Albuixech L, Barragan EF, Aparicio C, Gil FJ, San RJ, Vazquez B, Planell JA. Mechanical performance of acrylic bone cements containing different radiopacifying agents. Biomaterials 2002, 23, 1873–1882. [71] Kotha SP, Li C, McGinn P, Schmid SR, Mason JJ. Improved mechanical properties of acrylic bone cement with short titanium fiber reinforcement. J Mater Sci Mater Med 2006, 17, 1403–1409. [72] Kane RJ, Yue W, Mason JJ, Roeder RK. Improved fatigue life of acrylic bone cements reinforced with Zirconia fibers. Journal of the Mechanical Behavior of Biomedical Materials 2010, 3, 504–511. [73] Lopes P, Corbellini M, Ferreir BL, Almeida N, Fredel M, Fernandes MH, New RC. PMMA–co–EHA glass-filled composites for biomedical applications, Mechanical properties and bioactivity. Acta Biomaterialia 2009, 5, 356–362. [74] Chan M, Lau K, Wong T, Ho M, Hui, David. Mechanism of reinforcement in a nanoclay/polymer composite. Composites, Part B 2011, 42, 1708–1712. [75] Khaled SMZ, Charpentier PA, Rizkalla AS. Synthesis and characterization of poly (methyl methacrylate)-based experimental bone cements reinforced with TiO2–SrO nanotubes. Acta Biomaterialia 2010, 6, 3178–3186. [76] Jia Z, Wang Z, Xu C, Liang J, Wei B, Wu D, Shaowen Z. Study on poly (methyl methacrylate),carbon nanotube composites. Materials Science and Engineering. 1999, A271, 395–400. [77] Coleman JN, Khan U, Blau WJ, Gunko YK. Small but strong, A review of the mechanical properties of carbon nanotube-polymer composites. Carbon 2006, 44, 1624–1652. [78] Ormsby R, McNally T, Mitchell C, Dunne N. Incorporation of multiwalled carbon nanotubes to acrylic based bone cements, Effects on mechanical and thermal properties. Journal of the Mechanical Behavior of Biomedical Materials 2010, 3, 136–145.

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[79] Lee WJ, Lee SE, Kim CG. The mechanical properties of MWNT/PMMA nanocomposites fabricated by modified injection molding. Composite Structures 2006, 76, 406–410. [80] Ormsby R, McNally T, Mitchell C, Dunne N. Influence of multiwall carbon nanotube functionality and loading on mechanical properties of PMMA/MWCNT bone cements. J Mater Sci Mater Med 2010, 21, 2287–2292. [81] Nien YH, Huang C. The mechanical study of acrylic bone cement reinforced with carbon nanotubes. Materials Science and Engineering B 2010, 169, 134–137.

Sandra Pina and José M.F. Ferreira

5 Bioresorbable composites for bone repair 5.1 Introduction During the last century, new materials and surgical techniques have drastically changed the lives of millions of patients. Biomedical implants have made an important contribution to modern health care and will expand further increasing the aging population. Each material considered for potential clinical applications has specific chemical, physical and mechanical properties, which might originate variations in host/material response. However, the biological characteristics and the anatomic site of implantation are also important for the behavior and outcome of the implant. Autografts and allografts have been used by orthopedic surgeons to repair fractures and other bone defects. Nonetheless, limitations including risk of disease transfer, donor-site morbidity, potential immunogenicity, and insufficient supply have led investigators to search for alternative bone repair materials [1]. Bioresorbable composite materials involving a biodegradable polymeric matrix and bioactive/resorbable fillers have been seen as a promising approach for internal fixation in orthopedic surgeries for hard tissue repair and reconstruction, to attach soft tissues or tissue grafts to bone. These resorbable composites have the potential to replace the commonly used metal implants, with advantages including no necessity of a second surgical intervention for implant removal, radiolucency, no corrosion, no accumulation of metal in tissues, less pain and reduced stress-shielding [2–8]. The combination of polymers with calcium phosphates has been widely studied since this concept offers improved mechanical properties, while maintaining the favorable osteoconductivity and biocompatibility of calcium phosphates [9]. Moreover, with the addition of calcium phosphates, the degradation is delayed and the pH of the surrounding solution remains stable for longer periods [10]. Bioresorbable composite materials are most frequently manufactured from biodegradable polymers, as poly α-hydroxy acids, poly alkanoates, poly urethanes, poly orthoesters, polycarbonates, or copolymers of these [11–14]. The most widely researched of these in the biomedical field are poly α-hydroxy acids, such as polylactic acid PLA), polyglycolic acid (PGA), their copolymer polylactic-co-glycolic acid (PLGA), and polycaprolactone (PCL) [15–17]. Their degradation process in the body comprises hydrolysis followed by metabolization [18]. The process of hydrolysis involves the breaking of the polymer chains within the implant material to produce by-products that include lactic acid and glycolic acid single molecules. These degradation products are metabolized in the liver and produce carbon dioxide as a byproduct, which the body can eliminate [19]. The ideal biodegradable material provides appropriate strength whilst degrading in a predictable fashion throughout the healing process without causing adverse reactions [20].

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Bioresorbable composite materials are most common in screws and suture anchors for soft-tissue fixation and bone-tendon-bone grafts during anterior cruciate ligament (ACL) reconstruction. Commercially available bioresorbable composites are presented in Tab. 5.1. Several studies have failed to observe adverse clinical reactions with bioresorbable composites [21–23] or have shown only a mild reaction [24, 25] when compared with documented inflammatory response seen with biodegradable polymers. Histologically, less inflammatory response was seen in PLLA/hydroxyapatite (HA) composites than with PLLA alone [25]. The release of basic salts by the degradation of the bioceramics may buffer the acidic breakdown products of the polymers. β-tricalcium phosphate (TCP) has been shown to buffer the pH near PLGA implants undergoing degradation [26], and pH buffering causes less toxicity [27]. HA has also been shown to buffer the acidic breakdown products of PLLA [28]. This chapter presents an overview on the available bioresorbable composite materials for bone repair with a particular focus on new developments and trends in the field. Tab. 5.1: Commercially available bioresorbable composite materials

Manu- Product facturer Arthrex

Material

Indication

Properties

IntrafixTM

PLA/β-TCP

Tibial fixation

n.d.*

Corkscrew FT Suture anchor

PLA/β-TCP (85/15)

Rotator cuff repair Provides excellent contact between tendon and bone and stability in rotation and protects a broad healing zone from synovial fluid infiltration.

SutureTak Suture anchor

PLA/β-TCP (85/15)

Shoulder repair

PushLock

PLA/β-TCP (85/15)

Rotator cuff repair Allows the surgeon to adjust the amount of tension on the tissue intraoperatively, allowing for precise suture and tissue reduction.

Biocomposite Interference screw

PLDLA/ BCP (70/30)

ACL and PCL reconstruction

Features a unique molded-in suture eyelet that maintains its strength throughout most of the degradation cycle and eliminates suture abrasion during knot tying. The flexible eyelet eliminates the need to orientate the eyelet during insertion to optimize suture sliding.

Early bone formation can be connected to the favorable osteoconductive and bioresorbable properties within biphasic calcium phosphates.

5.1 Introduction   

Manu- Product facturer ArthroCare

Biomet

Material

Indication

DoubleplayTM Suture anchor

PLLA/β-TCP (70/30)

Tendon to tendon The unique “eyeless” design provides increased strength to the implant, and eliminates any risk of eyelet breakage.

Bilok® Parallel Interference screw

PLLA/β-TCP (70/30)

Bilok® Tapered Interference screw

PLLA/β-TCP (70/30)

Fixation of softtissue and bonetendon-bone grafts during ACL reconstruction

ComposiTCPTM30 PLDLA/β-TCP Bone-tendonInterference (70/30) bone and soft screw tissue fixation ComposiTCPTM60 PLDLA/β-TCP Soft-tissue fixaInterference (40/60) tion screw

ConMed Matryx® Linvatec Interference Screw

Mitek

   71

SR PLDLA (96%L/ 4%D)/β-TCP (75/25)

Properties

Osteoconductive healing response that essentially converts the screw itself into bone which creates a stronger fixation of the tendon, permitting patients to more quickly return to normal activity and restoring the operative site to a more natural condition. Unique star-shaped drive mechanism that limits stress and distributes torque evenly through the screw during insertion.

ACL and PCL graft Absorption over time and faster new fixation bone formation around the repaired ligament.

Biocryl® Interfer- PLGA/β-TCP ence Screw (70/30)

Fixation of soft- Osteoconductive and radiopaque. tissue and bone- Absorbs and allows for ossification of tendon-bone the implant site in as little as 3 years. grafts during ACL reconstruction

HealixTM BR Suture anchor

PLGA/β-TCP (70/30)

Bone fixation

n.d.*

LupineTM BR Suture anchor

PLGA/β-TCP (70/30)

Soft Tissue Repair/Arthroscopy

Complete, faster resorption and demonstrated bone in-growth with optimal pullout strength.

BioKnotlessTM BR Suture anchor

PLGA/β-TCP (70/30)

Gryphon TM BR Suture anchor

PLGA/β-TCP (70/30)

Rapid resorption profile, rapid bone in-growth, and demonstrated pullout strength while maintaining its proven design characteristics. Rotator Cuff, Instability

Provides audible feedback and confidence of fixation.

Bio-Intrafix Inter- PLA/β-TCP ference screw

ACL reconstruction

Excellent fixation, with >1000N pullout strength.

Milagro® BR Interference screw

Soft tissue repair n.d.* in ACL reconstruction

PLGA/β-TCP (70/30)

72   

   5 Bioresorbable composites for bone repair

Manu- Product facturer

Material

Indication

Properties

Smith & OsteoRaptor Nephew Suture anchor

PLLA/HA (75/25)

Hip labral repair

2.3 mm anchors allow for precise positioning, with a reliable track record of three years on the market.

BioRCI-HA Interference screw

PLLA/HA (75/25)

Fixation of soft- Includes reverse threaded screws to tissue and bone- keep the soft tissue grafts in optimal tendon-bone position. grafts during ACL and PCL reconstruction

BioSure HA Interference screw

PLLA/HA (75/25)

ACL reconstruction

Designed to reduce screw breakages by optimizing stress distribution and force transfer. Both the screw well and driver have been lenthened, so the driver inserts all the way to the tip. The screw also features a tapered body for ease of insertion, and a consistent wall thickness throughout for added durability.

Biosteon Wedge Interference screw

PLLA/HA (75/25)

ACL reconstruction, soft tissues and allografts procedures

Increase bone formation and decrease inflammatory cells. Improve implant / bone integration and thereby reduced tunnel widening and risk of graft slippage post-operatively.

Inion OTPS™ Biodegradable fixation system

PLLA/PLDA/ Fixation of trimethylene cancellous carbonate bone fractures, osteotomies or arthrodeses of the upper extremity, ankle and foot in the presence of appropriate plaster cast immobilization

Implants gradually loose their strength during 18-36 weeks in vivo and bioresorption takes place within two to four years. Implants are contraindicated for: (i) active or potential infection, (ii) high-load bearing applications and (iii) patient conditions including limited blood supply, insufficient quantity or quality of bone; and where patient cooperation cannot be guaranteed (e.g., alcoholism, drug abuse).

Duosorb®

PLDLA/β-TCP Bone (40/60) reconstruction

After 6 months, implants can partially become fragmented as TCP particles are progressively released from the polymer matrix.

Stryker

SBM

5.2 Bioresorbable materials   

Manu- Product facturer Takiron

Fixsorb®

   73

Material

Indication

Properties

PLLA/u-HA (60/40)

Fracture bone fixation

Its osteoconductive property and bioactive property of fusing directly with bone have contributed to improvement of its early fixation ability. It is degraded and resorbed in a uniform manner and rarely induces strong tissue reactions that result from rapid degradation. It is radiopaque and the device can be observed after surgery.

PLA: poly lactic acid; TCP: tricalcium phosphate; PLGA: poly lactide glycolide; PLLA: poly-L-lactide; BCP: biphasic calcium phosphate; SR: self-reinforced; HA: hydroxyapatite; u-HA: unsintered HA; ACL: anterior cruciate ligament; PCL: posterior cruciate ligament; n.d.*: not-documented.

5.2 Bioresorbable materials 5.2.1 Polymers In orthopedic applications, for bone repair and regeneration, polymers with very high strength and stiffness are required. Polymers most widely researched for orthopedic applications are poly α-hydroxy acids [29]: polyglycolic acid (or polyglycolide – PGA), polylactic acid (or polylactide  – PLA), poly lactide-co-glycolide (PLGA) and poly ε-caprolactone (PCL), and their copolymers, in part because they can be self-reinforced to gain better strength properties [30]. Another common bioresorbable polymer for orthopedic applications is poly glyconate, a copolymer consisting of two parts glycolic acid and one part trimethylene carbonate. Many other bioresorbable polymers have been studied but have yet achieved little or no clinical application, including poly hydroxyl butyrate, poly orthoesters and pseudo polypeptides based on tyrosine [31–33].

5.2.1.1 Polyglycolic acid – PGA Polyglycolic acid is a brownish, hard crystalline polymer with an average molecular weight of 20,000 to 145,000, a melting point between 224–230 °C, and a glass transition temperature of 36 °C [34] (Fig. 5.1). It is degraded in hydrolysis, and is broken down by nonspecific esterases and carboxyl peptidases. Its mechanical strength is lost in 6 weeks, and it is totally resorbed in a few months depending on the molecular weight, purity, crystallinity and the size and shape of the implant [35, 36].

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   5 Bioresorbable composites for bone repair

O O

Catalyst + Heat

O O

O

CH2 C

O O

CH2

C

n

O Glycolide

Poly glycolide

Fig. 5.1: Structure of PLG polymer and their respective monomer.

However, adverse tissue responses to fixation implants made of PGA have been reported [37–39], with the incidence rate varying from 2.0 to 46.7% [40]. The highest incidence has been observed in fractures of the distal radius and the scaphoid bone [41–43]. Another work reported adverse tissue reaction in 5.3% (107 reactions) of operations using self-reinforced-PGA implants [37]. Nevertheless, the frequency of foreign-body reactions significantly decreased when the dye was omitted from the PGA implant material [37, 44]. The risk of adverse tissue reactions has deterred the use of PGA implants in favor of PLA, for example, which have lower rate of degradation. PGA was the first bioresorbable polymer used in bone surgery as screws, plates and reinforcing pins, suggested by Schmitt and Polistina in 1969 [45]. PGA has been used mostly in sutures, rods and screws in fracture fixation of cancellous bone due to the rapid loss of mechanical strength of the implants [46, 47].

5.2.1.2 Polylactic acid – PLA PLA is a pale-colored semicrystalline polymer with molecular weights of 180,000 to 530,000, a melting point of about 174 °C, and glass transition temperature of 57 °C [48] (Fig. 5.2).

O H3C

Catalyst + Heat

O O O Lactide

CH3 O O CH C

CH3

CH3 O O CH C n

Poly lactide

Fig. 5.2: Structure of PLA polymer and their respective monomer.

PLA diverges from PGA that is a stronger acid and behaves more hydrophylically than PLA which is more hydrophobic because of its methyl groups. The decrease of pH values in the tissues adjacent to degrading biodegradable polymers may contribute

5.2 Bioresorbable materials   

   75

to adverse effects, an issue that could be addressed by the incorporation of basic salts within the polymer [49]. PLA can exist in several distinct forms, such as PLLA and poly-D-lactide (PDLA) [50], depending on the L and D configuration, and it is also degraded via hydrolysis. P(L/D)LA: PLLA is hydrophobic and crystalline and thus resistant to hydrolysis and degradation. By adding D-isomers into an L-isomer based polymerization system, polymer chains widen and cannot be packed as tightly as PLLA polymer chains. This results in a less crystalline and more rapidly degraded material [51]. PLLA implants have been successfully used to fixate and heal tissue and bone, for injuries such as ligament damage and skeletal fractures. PLLA materials have excellent biocompatibility and they are relatively strong because of its slow rate of complete resorption into the body, although it does not have sufficiently high strength characteristics for use in the fixation of larger fractures such as those in the humerus and femur. Much of the referenced PLLA research has focused on veterinary applications using rabbits [52, 53]. PDLLA (poly-DL-lactide) also shows characteristics that could be employed in high-strength situations, but PLLA is the preferred material for use in fracture fixation implants due to its higher strength compared with PDLLA [54]. Good biocompatibility has been observed with PLA implants [55, 56]. Also, PDLLA and PLLA were well-tolerated and the tissue response inside muscle was similar to that of stainless steel [57]. PDLLA pins were compared with PDLLA (70:30) with β-TCP (10%) and no different reaction in synovial membrane, lymph nodes, or bone formation was observed with either polymer [58]. Complete degradation of both materials occurred within 36 months. However, some problems related to foreign-body reactions were reported although they should not be generalized to all PLLA materials. Eitenmüller et al. [59] observed that 52% of the patients demonstrated an aseptic soft tissue problem caused by delayed clearance of the degrading PLA particles using PLLA plates for fixation of ankle fractures. In a second protocol, smaller plates and screws did not cause any soft tissue reactions. Bergsma et al. [60] reported a late tissue response to PLLA bone plates and screws used in the fixation of ten zygomatic fractures in humans. Intraosseally implanted self-reinforced-PLLA screws and pins have been shown to cause similar, mild foreign-body reactions as corresponding metallic devices, without signs of inflammatory reactions during follow-up of 48 weeks [56]. The total resorption time of PLA is considerably longer than PGA [57]. Cifuentes et al. [61] developed a composite consisting of PLLA loaded with 30 wt% of Mg particles for osteosynthesis implants with improved mechanical properties (hardness up to 340 MPa and yield strength up to 100 MPa). PGA screws have been shown to completely disappear within 6 months while PLLA has a very long degradation time and has been shown to persist in tissues for as long as five years post implantation [40]. Therefore, many resorbable orthopedic implants are currently manufactured from PLLA.

76   

   5 Bioresorbable composites for bone repair

5.2.1.3 PGA-PLA copolymers The problem of degradation has led to the development of the copolymers. PGA and PLA can be combined to form a full range of PLGA copolymers. Their properties can be controlled by varying the ratio of glycolide to lactide at different compositions [62], which offers the potential to develop site-specific bone fixation and soft tissueanchoring devices [63–65]. The degradation time of the copolymers is related to the ratio of monomers as the higher content of glycolide, the faster the rate of degradation. For example, the degradation time is 5 months for a 85:15 PDLA:PGA copolymer [66]. However, an exception to this rule is the 50:50 ratio of PGA:PLA, which exhibits the fastest degradation [67]. There are some distresses about the potential aseptic inflammatory wear debris generated during implant resorption. For example, Andrews and Veznedaroglu evaluated the incidence of infection in a series of 296 patients in which 146 received craniotomy fixation with titanium implants and 150 received craniotomy fixation with a PDLLA copolymer [68]. 43 patients in the titanium group and 37 patients in the polymer group also received postoperative irradiation. The incidence of infection was 4.6% for the titanium group and 4.0% for the resorbable polymer group. Caminear et al. [69] used 82:18 PLLA:PGA copolymer implants to fix distal chevron osteotomies in 15 patients and only one patient developed postoperatively a giant cell granuloma needing debridement. Resorbable membranes, made of PLA:PGA copolymers have been used for guided bone regeneration (GBR) [70, 71]. These membranes generally start to resorb between 4 and 6 weeks but, their stiffness and duration have been questioned. Sandberg et al. [72] noted that some resorbable membranes used in their study showed a lack of stiffness, resulting in the collapse of the membrane into the defect area, causing the newly formed bone to take on an hourglass shape.

5.2.1.4 Poly ε-caprolactone – PCL PCL (Fig. 5.3) has been widely used in the medical field for bone repair, for the past 30 years [73–75] with FDA (United States Food and Drug Administration) approval. The good solubility of PCL, its low melting point and exceptional blend-compatibility has stimulated extensive research into its potential application in the biomedical field. It is a hydrophobic and semi-crystalline polymer, with molecular weights of 3,000 to 90,000, melting point between 59 and 64 °C and glass transition temperature of –60 °C [76]. PCL has a good mechanical property and can be easily processed when compared with PLA and PGA [77]. In particular, it is especially interesting for the preparation of long-term implantable devices, owing to its degradation (up to 2 years in vivo) which is even slower than PLA, due to its hydrophobic character and high crystallinity [78]. Also, cell adhesion and growth on PCL implants is limited due to its hydrophobicity, though PCL is known to cause only a minor inflammatory response [79, 80]. In order

5.2 Bioresorbable materials   

O

O

Catalyst + Heat

   77

O O n

Caprolactone

Poly caprolactone

Fig. 5.3: Structure of PCL polymer and their respective monomer.

to improve the hydrophobicity of PCL, a copolymer of PCL with hydrophilic material such as poly(ethylene glycol) (PEG) has been performed [81, 82]. This approach has led to enhanced cell attachment, and accelerated hydrolytic degradation. PCL/PEG/ PCL triblocks constitute an excellent support for human endothelial cell adhesion and growth [83]. PCL-PEG copolymers have also been shown to support both human, and rat marrow stromal cell proliferation [84]. PCL has been used as a drug delivery system and as a potential material for bone repair, such as an injectable bone substitute [85], scaffold for bone-tissue engineering [86, 87], and an in-situ-polymerized for maxillofacial applications [88]. Marra et al. [87] reported that PCL is a comparable substrate for supporting cell growth resulting from two-dimensional bone marrow stromal cell culture. Also, PCL/PLA blend disc incorporated with hydroxyapatite is feasible as scaffolds for bone tissue engineering. In another study it was demonstrated that PCL scaffolds in combination with periosteum may serve as suitable biocomposites for neo-cartilage formation and subchondral bone regeneration for the reconstruction of large osteochondral lesions [89].

5.2.2 Bioactive ceramics Bioactive ceramic materials used for bone repair are β-tricalcium phosphate (β-TCP), hydroxyapatite (HA), calcium silicate and phosphate-based glasses, due to their biocompatibility and osteoconductivity [90, 91]. Phosphate-based glasses and fibers show great potential as reinforcing materials as they degrade completely in aqueous media [92, 93] and their degradation products consist of ions already present within the body (mainly calcium and phosphates) [94]. The most commonly used are β-TCP and HA. β-TCP is biodegradable and able to promote osteogenesis and new bone formation. HA is highly crystalline and is the most stable and least soluble CaP in an aqueous solution down to a pH of 4.2 [91]. The resorption of a ceramic HA is believed to be slow (1 to 2 % per year), and once implanted into the body, HA may remain integrated into the regenerated bone tissue, while β-TCP is completely reabsorbed [95, 96]. The main drawback of these materials is their limitation to load-bearing applications, which arises from their poor mechanical properties. Therefore, recent research efforts are focused on combining these materials with polymers generally bioinert,

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to provide the composite bioactivity, in order to form a composite with improved mechanical properties. The composite implants are able to form a chemical bond with the host tissue, and the fixation of implants is accelerated [97–99]. For example, CaP/ poly-DL-lactide-co-glycolide composites exhibit good adhesion onto human cells, indicating a high level of biocompatibility [98, 99]. Additionally, previous studies with composite materials consisting of PLLA/TCP or PLLA/HA showed a rapid resorption and replacement by newly-formed bone tissue [100, 101]. Also, PLA/HA composite is a better substrate for osteoblast cell cultures than PLA alone [102, 103]. Such composites, with xenogeneic osteoblasts grown in them, become the site of intense osteogenesis when implanted into athymic mice [104]. Studies on composites made of PLGA copolymers and calcium silicate demonstrated that these materials are able to promote bone regeneration due to their excellent bioactivity and biocompatibility [105–107]. For example, Su et al. [107] produced nanoporous calcium silicate with PLGA and showed that the composites could improve the cell attachment rate on its surface and promote cells growth as compared with PLGA. Composites made of PCL reinforced with phosphate-based glass have also been reported [108–111]. Ahmed et al. [108] investigated PCL reinforced with binary calcium phosphate glass fibers, and showed that the flexural strength and modulus values obtained were 25–30 MPa and 2.5 GPa respectively for an approximately 18% fiber volume fraction sample. It was stated that these values were in line with human trabecular bone properties and not suitable for cortical bone. Prabhaker et al. [109] investigated PCL composites reinforced with phosphate-based glass (PBG) particulates and showed that it was possible to bestow control over the composites’ properties via incorporation of particulates with differing degradation profiles.

5.3 Composites manufacturing methods The methods of manufacturing biocomposites are melt molding and solid state drawing [20]. The melt-molding process can be subdivided into three methods to process implants from the melted raw material: compression molding, injection molding and extrusion. In the melt-molding process, a pre-impregnated (or prepeg) matrix imbedded with uniaxially-oriented reinforcing fibers is arranged in a twopiece mold which is heated under pressure to produce laminated parts. This process is not very expensive and not suitable for small medical implants [112]. In injection molding, the material mixture is injected into a mold at elevated temperature and pressure, and the composite is removed after cooling. This process allows rapid, high volume and economical production but requires expensive tooling. Sterilization of bioresorbable composite implants may require special processes due to their sensitivity to hydrolysis and to decrease of molecular weight at high temperatures. Conventional steam, ethylene oxide, and gamma irradiation at selected

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doses in combination with sufficient cooling during irradiation, allow effective sterilization of the products without adverse effects [113]. Another method to produce extremely strong composite materials is the selfreinforcing (SR) technique, invented by Törmälä and Rokkanen [114]. In this technique a composite structure is produced by a partially crystalline or amorphous polymeric material made of orientated fibers, fibrils or chain crystals and binding matrix [115]. The SR technique led to better mechanical properties (higher reinforcement degree) and eliminated the problem of toxic adhesion promoters. It results in initial strengths 5 to 10 times higher than those implants manufactured with melt-molding technique [116]. To manufacture screws, the SR polymer can be compression molded or machine cut, which has improved significantly the torque and bending strengths of the screws [117]. Finally, SR materials can be sterilized by gamma irradiation, thus eliminating toxic residues that remain after other methods of sterilization.

5.4 Clinical applications of bioresorbable composites for bone repair Bioresorbable composite implants are mainly used for stabilizing fractures, osteotomies, bone grafts and fusions mostly in trabecular bones, as well as to reattach ligaments, tendons, meniscal tears and other soft tissue structures [118, 119]. The midfacial skeleton would seem to be an acceptable location for the use of bioresorbable implants, given the relatively easy access to fractures of this region and the low biomechanical stresses to which they would be exposed. Orthognatic surgery: Bioresorbable fixation implants have been used for the fixation of facial bones in orthognathic surgery [120]. It was reported that adverse effects were observed in two plate exposures between the third and ninth months, and occurred mainly in the posterior maxillary region [120]. Shoulder: Implants were successfully applied in the repair and reconstruction of many intra-articular and extra-articular abnormalities in the shoulder, such as shoulder instability, rotator cuff tears, and biceps lesions or biceps tendon tenodesis [121]. Knee: Bioresorbable implants have been used in treating knee [24, 122], wrist [123] and hand [124] injuries. Soft tissue reconstruction in complex knee injuries was performed by using meniscal tacks and biodegradable suture anchors [122]. Foot and ankle: Composite implants have been employed for management of foot and ankle fractures [38, 59, 125, 126]. For example, Eitenmuller et al. [59] investigated the suitability of PLLA screws and plates for the treatment of ankle fractures.

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Fractures healed within 6 weeks, but 52% of the patients experienced an aseptic soft tissue problem caused by delayed clearance of the degrading PLA particles. Prompted by these problems, the authors treated 7 patients with volume-reduced plates and screws with flat heads, and none of the patients experienced any soft tissue reactions. The authors concluded that the use of PLLA implants is acceptable for the fixation of ankle fractures, and soft tissue inflammatory reactions can be avoided by using implants with reduced volume of biodegradable material. Spine: Applications of bioresorbable implants in spinal reconstructive surgery have been reported [127–129]. For example, PLA screws were used for anterior cervical decompression and fusion procedures [130, 131]. Deguchi et al. [128] evaluated the biomechanical stability of PLLA pins in the posterior lumbar spine and showed that the PLLA pin construct provided improved stability to the spine, although it was not as stiff as the screw construct. Pediatric orthopedics: Bioresorbable materials are used in pediatric orthopedics [132–136]. Svensson et al. [137] reported the use of biodegradable osteosynthetic materials in 50 children with transphyseal or osteochondral fractures. Two patients had non-union of articular radial head fractures, possibly related to a foreign-body reaction. In a study by Eppley et al. [133], where resorbable PLLA-PGA (LactoSorb) plate and screw fixation for craniofacial surgery was applied in 1883 infants and young children, it was observed that device-related complications requiring reoperation occurred in less than 0.5% of the implanted patients, which is less frequent than that reported for metallic bone fixation. Significant infectious complications occurred in 0.2%, device instability primarily resulting from postoperative trauma occurred in 0.3%, and self-limiting local foreign-body reactions occurred in 0.7% of the treated patients. The overall reoperation rate attributable to identifiable device-related problems was 0.3%. Bioresorbable membranes are used in GBR for the treatment of periodontal intraosseous defects [138, 139] and some oral surgical procedures, such as sinus lifts [140, 141].

5.5 Conclusions Bioresorbable composite materials have a bright future as implants for bone repair. They are the favored material to work with for a growing number of surgeons because of the postoperative advantages and potential for customization. These materials retain their strength long enough to support healing of bone, and then gradually and harmlessly disintegrate in the patient’s body. Also, they can be engineered to alter their material properties and degradation characteristics. Conversely, bioresorbable

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composites are still in stages of research and development. Future developments of these materials for bone repair and regeneration should be focused on the reduction of the foreign-body reaction and enhancement of the mechanical strength.

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[97] Elgendy H, Norman M, Keaton A, Laurencin C. Osteoblast-like cell (MC3T3-E1) proliferation on bioerodible polymers: An approach towards the development of a bone-bioerodible polymer composite biomaterials. Biomater 1993, 14, 263–9. [98] Ignjatovic N, Ninkov P, Kojic V, Bokurov M, Srdic V, Krnojelac D. Cytotoxicity and fibroblast properties during in vitro test of biphasic calcium phosphate/poly-dl-lactide-co-glycolide biocomposites and different phosphate materials. Microsc ResTech 2006, 69, 976–82. [99] Durucan C, Brown P. Calcium-deficient hydroxyapatite-PLGA composites: Mechanical and microstructural investigation. J Biomed Mater Res 2000, 51, 726–34. [100] Higashi S, Yamamuro T, Nakamura T, Ikada Y, Hyon S, Jamshidi K. Polymer-hydroxyapatite composites for biodegradable bone fillers. Biomater 1986, 7, 183–7. [101] Prokop A, Helling H, Fischbach R, Wollsiefer M, Dietershagen M, Reif D, et al. Neue biodegradierbare Tricalciumphosphat-Polylactidstifte zur Refixation osteochondraler Fragmente. Erste radiologische Ergebnisse einer tierexperimentellen. 3rd European Trauma Congress, Amsterdam, 1998. [102] Ma PX, Zhang R, Xiao G, Franceschi R. Engineering new bone tissue in vitro on highly porous poly(alpha-hydroxyacids)/hydroxyapatite composite scaffolds. J Biomed Mater Res 2001, 54, 284–93. [103] Tsigkou O, Hench LL, Boccaccini AR, Polak JM, Stevens MM. Enhanced differentiation and mineralization of human fetal osteoblasts on PDLLA containing Bioglass composite films in the absence of osteogenic supplements. J Biomed Mater Res A 2007, 80, 837–51. [104] Yang XB, Webb D, Blaker J, Boccaccini AR, Maquet V, Cooper C, et al. Evaluation of human bone marrow stromal cell growth on biodegradable polymer/bioglass composites. Biochem Biophys Res Commun 2006, 342, 1098–107. [105] Cheng W, Li H, Chang J. Fabrication and characterization of β-dicalcium silicate/poly(d,l- lactic acid) composite scaffolds Mater Lett 2005, 59, 2214–8. [106] Li H, Chang J. Preparation and characterization of bioactive and biodegradable Wollastonite/ poly(D,L-lactic acid) composite scaffolds. J Mater Sci-Mater Med 2004, 15, 1089–95. [107] Su J, Wang Z, Yan Y, Wu Y, Cao L, Ma Y, Baoqing, Y, Ming L. Nanoporous calcium silicate and PLGA biocomposite for bone repair. J Nanomater 2010, 2010 Article ID 181429. [108] Ahmed I, Parsons AJ, Palmer G, Knowles JC, Walker GS, Rudd CD. Weight loss, ion release and initial mechanical properties of a binary calcium phosphate glass fibre/PCL composite. Acta Biomater 2008, 4, 1307–14. [109] Prabhakar RL, Brocchini S, Knowles JC. Effect of glass composition on the degradation properties and ion release characteristics of phosphate glass-polycaprolactone composites. Biomater 2005, 26, 2209–18. [110] Ahmed I, Jones IA, Parsons AJ, Bernard J, Farmer J, Scotchford CA, et al. Composites for bone repair: Phosphate glass fibre reinforced PLA with varying fibre architecture. J Mater Sci- Mater Med 2011, 22, 1825–34 [111] Ahmed I, Cronin PS, Abou Neel EA, Parsons AJ, Knowles JC, Rudd CD. Retention of mechanical properties and cytocompatibility of a phosphate-based glass fibre/polylactic acid composite. J Biomed Mater Res B Appl Biomater 2008, 89, 18–27. [112] Iftekhar A. Biomedical composites. Standard handbook of biomedical engineering and design, New York, McGraw-Hill, 2004. [113] Bernkopf M. Sterilisation of bioresorbable polymer implants. Medical Device Technology 2007, 18, 26–29. [114] Tormala P. Biodegradable self-reinforced composite materials: Manufacturing structure and mechanical properties. Clin Mater 1992, 10, 29–34. [115] Rokkanen PU. Absorbable materials in orthopaedic surgery. Annuals Med 1991, 23, 109–15.

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[116] Tormala P, Vainoipaa S, Kilpikari J, Rokkanen P. The effects of fibre reinforcement and gold plating on the flexuraland tensile strengths of PGA/PLA copolymer materials in vitro. Biomater 1987, 8, 42–5. [117] Pohjonen T, Helevirta P, Törmälä P, Koskikare K, Pätiälä H, Rokkanen P. Strength retention of self-reinforced poly-L-lactide screws. A comparison of compression moulded and machine cut screws. J Mater Sci Mater Med 1997, 8, 311–20. [118] Bostman O. Current concepts review: Absorbable implants for the fixation of fractures. J Bone Joint Surg Amer 1991, 73, 148–53. [119] Tunc DC, Rohovsky MW, Zadwadsky JP, Spieker JE, Strauss ED. Evaluation of body absorbable screw in avulsion type fractures. Annual Meeting of the Society for Biomaterials, 1986. [120] Fedorowicz Z, Nasser M, Newton T, Oliver R. Resorbable versus titanium plates for orthognathic surgery. Cochrane Database of Systematic Reviews 2007, 2, Art. No.: CD006204. DOI: 10.1002/14651858.CD006204.pub2. [121] McFarland EG, Park HB, Keyurapan E, Gill HS, Selhi HS. Suture anchors and tacks for shoulder surgery. Part 1: Biology and biomechanics. Amer J Sports Med 2005, 33, 1918–23. [122] Burkhart SS. The evolution of clinical applications of biodegradable implants in arthroscopic surgery. Biomater 2000, 21, 2631–4. [123] van Manen CJ, Dekker ML, van Eerten PV, Rhemrev SJ, van Olden GD, van der Elst M. Bio-resorbable versus metal implants in wrist fractures: A randomised trial. Arch Orthop Trauma Surg 2008, 128, 1413–7. [124] Hughes TB. Bioabsorbable implants in the treatment of hand fractures: An update. Clin Orthop 2006, 445, 169–74. [125] Bostman, Hirvensalo E, Vainionpaa S, Max.Ela A, Vihtonen K, Tormala P, et al. Ankle fractures treated using biodegradable internal fixation. Clin Orthop 1989, 238, 195–203. [126] Hirvensalo E. Fracture fixation with biodegradable rods: Fourty-one cases of severe ankle fractures. Acta Orthop Scand 1989, 60, 601–6. [127] Coe JD, Vaccaro AR. Instrumented transforaminal lumbar interbody fusion with bioresorbable polymer implants and iliac crest autograft. Spine 2005, 1, 76–83. [128] Deguchi M, Cheng B, Sato K, Matsuyama Y. Biomechanical evaluation of translaminar facet joint fixation. Spine 1998, 23, 1307–18. [129] Vaccaro AR, Singh K, Haid R, Kitchel S, Wuisman P, Taylor W, et al. The use of bioabsorbable implants in the spine. Spine J 2003, 3, 227–37. [130] Brunon J, Duthel R, Fotso MJ, Tudor C. Anterior osteosynthesis of the cervical spine by phusiline bioresorbable screws and plates. Neuro-Chirurgie 1994, 40, 196–202. [131] Subach BR, Haid RW, Rodts GE, Branch CE, Alexander JT. Posterior lumbar interbody fusion (PLIF) using an impacted, bioabsorbable device. Amer Assoc Neur Surg, Congress of Neurological Surgeons, February 28 to March 2, Orlando, FL, 1999. [132] Bostman, Makela EA, Tormala P, Rokkanen P. Transphyseal fracture fixation using biodegradable pins. J Bone Joint Surg 1989, 71, 701–7. [133] Eppley BL, Morales L, Wood R, Pensler J, Goldstein J, Havlik R, et al. Resorbable PLLA-PGA plate and screw fixation in pediatric craniofacial surgery: Clinical experience in 1883 patients. Plastic Reconst Surg 2004, 114, 850–6. [134] Hope PG, Williamson DM, Coates CJ, Cole WG. Biodegradable pin fixation of elbow fractures in children: A randomized trial. J Bone Joint Surg 1991, 73, 965–8. [135] Partio ES, Merikanto J, Heikkila JT. Totally absorbable screws in fixation of subtalar extraarticular arthrodesis in children with spastic neuromuscular disease: Preliminary report of a randomized prospective study of 14 arthrodeses fixed with absorbable metallic screws. J Paediatr Orthop 1992, 12, 646–50.

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[136] Illi OE, Gitzelmann C, Gasser B, Misteli F, Ruedi M. Five years of experience with biodegradable implants in paediatric surgery. J Mater Sci Mater Med 1994, 5, 417–23. [137] Svensson P, Janarv P, Hirsch G. Internal fixation with biodegradable rods in pediatric fractures: One year follow-up of 50 patients. J Pediatr Orthop 1994, 14, 220–4. [138] Listl S, Tu Y-K, Faggion CJ. A cost-effectiveness evaluation of enamel matrix derivatives alone or in conjunction with regenerative devices in the treatment of periodontal intra-osseous defects. J Clin Periodontol 2010, 37, 920–7. [139] Zybutz M, Laurell L, Rapoport D, Persson G. Treatment of intrabony defects with resorbable materials, non-resorbable materials and flap debridement. J Clin Periodontol 2000, 27, 169–78. [140] Listl S, Faggion CMJ. An economic evaluation of different sinus lift techniques. J Clin Periodontol 2010, 37, 777–87. [141] Pjetursson BE, Tan WC, Zwahlen M, Lang NP. A systematic review of the success of sinus floor elevation and survival of implants inserted in combination with sinus floor elevation. Part I – Lateral technique. J Clin Periodontol 2008, 35, 216–40.

G. El-Damrawi and H. Doweidar

6 Bioactive glasses and glass-ceramics 6.1 Biodental metals, ceramics and bioactive glass-ceramics; historical background The use of biomaterials became practically known in surgery with J. Lister in the 1860s. Earlier trials were generally unsuccessful due to various problems accompanying the presence of foreign bodies. Among these is that the body’s immunological cells cannot access the implanted metal part. The earliest modern successful implants were used in the trials of fixation of fractured bone. However, most bone plates broke as a result of unsophisticated mechanical design. Material such as vanadium steel was considered for its good mechanical properties, but on the other side it suffers from fast corrosion in the body. Following the introduction of stainless steel and cobalt chromium alloys in the 1930s, greater success was achieved in fracture fixation and successful joint replacement surgeries could be done. Owing to further advances in materials and in surgical techniques, blood vessel replacements were tried in the 1950s. Heart valve replacement and cemented joint replacements were achieved in 1960s. More recent years have seen many further advances in different branches of applications [1–4].

6.2 Metallic implant materials Metals have been used as start materials for manufacturing some types of implants. “Sherman Vanadium steel” was the first metallic material, which was developed specifically for human use. This material is used to repair bone fracture, in the form of plates and screws. Some other metallic elements like iron are considered to be useful and essential in cell functions. The metallic materials cannot be tolerated in large amount in the body. This is because of their limited compatibility and their corrosion ability in the hostile body fluids. The corrosion products can diffuse into tissue resulting in unwanted effects. Corrosion is an undesired chemical reaction between metals and their environments. The human body represents an aggressive environment to implanted metals. The reaction between metals and body fluids results in degradation of metal to oxides or other compounds in the human body that contains water. Therefore, more attention must be paid toward corrosion durability of a metallic implant as an important aspect of its applicability as a biocompatible material. In the following sections we present some examples of biocompatible metallic alloys which can resist corrosion processes.

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6.2.1 Gold alloys It is well known that pure gold is relatively soft. The use of this metal as a restorative material is therefore limited to areas not subjected to too much loading or stress. On the other hand, gold alloys are useful in dentistry because of their corrosion resistance and stability. Corrosion resistance is retained in alloys containing 75 wt% gold and other noble metals. Moreover, gold alloys have mechanical properties that are superior to those of pure gold. As an example, introducing copper into gold matrix significantly increases the strength of the resulting alloy and as a consequence improves its mechanical properties.

6.2.2 Dental amalgam Amalgam is an alloy in which mercury is one of the component metals. The reason for using amalgam as a tooth filling material is that mercury is a liquid at room temperature. For this reason, it can react with other metals such as silver and tin and form a semi-plastic mass that can be easily packed into the cavity and can harden with time. Then, in order to fill the tooth cavity the dentist mixes solid alloy with mercury. The product is a readily formable material and can be used to fill the prepared cavity. The solid alloy that is used as a tooth filling material was reported to contain about 3 % mercury, 65 % silver, 2 % copper and not more than 26 % tin [2]. Until recently, metals, polymers and non-compatible glasses have been the major materials used as implants. The success of these non-bioactive implants has been very limited, especially for long-term applications. The disadvantage of such materials as implants is that they never become an integrated component of the biological system. The biological materials accumulate around the foreign implant material forming a capsule. Isolation by these capsules often prevents the implant from attaching to the required bone or tissue and this causes failure of the implant. These complications can be avoided if the implant materials are bioactive and can interact with the biological system [3, 4]. So, it is important to focus on the bioactive glasses and glass-ceramics, their properties and applications. Reviews on some specific types of biodental materials are also considered.

6.3 Glass-ceramics and bioactive glass-ceramics Natural glass-ceramic materials, like obsidian, have always existed in nature, but synthesized glass-ceramics (Pyroceram) were firstly prepared in 1953 and were found to fit their uses in different applications. The glass-ceramics were produced when lithium disilicate glass sample was annealed in a furnace at 600 °C and overheated to about

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900 °C. As a result, a white sample that had no change in its shape was observed. It was realized that this sample did not shatter when it accidentally dropped on the floor, contrary to what might normally have been expected from a piece of glass. This observation led to the conclusion that the thermally treated processes improve the properties and the structure of the material. This revealed that the treated glass is characterized by unusual toughness. The first synthetic research on glass-ceramic eventually led to the development of CorningWare in 1957. It also opened the field for the development of Vision, a transparent cookware. CorningWare entered the consumer marketplace in 1958 and became a well-known product [5].

6.3.1 Commercial glass-ceramic products The first known application of glass-ceramics was in the late 1950s. A certain type of glass-ceramics was used in the nose cones of rockets and aircrafts to protect their radar equipment. Such a glass-ceramic must have special electric, thermal and mechanical properties that make it fit for the application [6]. There is another very interesting class of glass-ceramics that is called patterned glass. Crystallization can be produced in this type by the effect of heat treatment. It can also be treated once more to form polycrystalline structure containing holes or some types of channels or any desired pattern. The products are used in electronics, and in some applications that include microchannels in optical fibers such as inkjet printer heads, substrates for pressure sensors and acoustic systems in headphones [6–8]. Common uses of other ceramic materials are as ceramic dishes, semiconductors, dental and magnetic materials and coating for rockets. The most successful commercial products in the glass-ceramics area include fire-rated glass, which is considered a specific type of material that is used to withstand risk of thermal shock damage. This material is characterized by its low coefficient of thermal expansion (CTE). The lower CTE of the glass-ceramics recommended them as successful materials for thermal applications [7, 8].

6.3.2 Protective glass-ceramic Protective glass-ceramic (armor) belongs to armor materials for the protection of people or equipment against high-speed projectiles or fragments. The ceramic armor materials contain Al2O3 and Zr2O3 species of high Young’s modulus and strength. It can be produced only at very high temperatures by costly manufacturing processes. But most glass-ceramics which are produced at relatively lower temperatures have lower hardness and Young’s modulus than the above-described ceramics and therefore cannot be used as armor materials. Another type of glass-ceramic is Schott’s

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Resistant. It can be transparent or opaque and is suitable as a base material for vehicular and personal armor systems.

6.3.3 Bioceramics Bioceramics, including biodental and bioactive glasses, are fabricated for special medical and dental applications. They can be used for repairing and reconstruction of damaged parts of the human tooth, e.g., porcelain crown and glass ionomer cement, or bone [9, 10]. Bioceramics should possess some specific requirements that make their properties compatible with those of the bone. This could be achieved by adjusting the composition and the structure of the material. In such a case the material used is considered as bioceramic. Based upon the nature of bond between the bioceramics and the bone, bioceramics are classified as: 1. bioinert, which can form a weak bioceramic/bone interface or 2. bioactive, that directly attaches to the bone and can form a strong interface. The bioinert ceramics include alumina (Al2O3) and zirconia (ZrO2) that are used for dental and orthopedic applications. Bioactive ceramics are mainly based on the formation and use of synthetic mineral phases with controlled properties such as compatibility and bioactivity. Such materials include calcium-deficient apatite and beta tricalcium phosphate phases [3, 4]. The good performance and withstanding long-term loading of bioceramics recommend them to be widely used as special materials for skeletal reconstruction. The great problem facing the use of bioceramics in the body is to replace diseased hard tissues such as bone, dentine, and enamel with a material that can function without failure. As is known, teeth (unlike bone) do not have the capacity to repair if they get damaged due to dental diseases. Therefore, there is an increasing demand for bioceramics for treatment and restoration. Bioceramics are well suited to meet this demand, which represents one of the fastest growing applications of biomaterials.

6.4 Preparation techniques Bioactive glasses and glass-ceramics can be obtained by various techniques. The conventional and most used method is to obtain the glass by cooling its melt. This method is adequate, especially for compositions that solidify without devitrification. It also represents the common method to obtain many types of glass-ceramics, either through melt cooling or by subjecting the glass to certain thermal treatments [11–16]. Nowadays, the sol-gel technique presents an alternative way to produce different types of glass and glass-ceramics. The process has been known since the 1840s and was rediscovered around forty years ago [17, 18]. The glasses are usually formed from organo-metallic precursors, such as tetramethoxysilane, tetraethyl orthosilicate

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and triethyl phosphate [19]. The sol-gel process is affected by different factors like the molar ratio of water (and/or solvent)-to-precursor, type of solvent, type and concentration of catalyst and the pH value of the emulsion [20]. This method is useful in producing different forms of glass, e.g., bulk glass, glass powder, thin film and fibers, and glass foam [21]. The main advantage of the sol-gel technique is the possibility to produce ceramic materials at relatively low temperatures. Saadaldin et al. [22] used the sol-gel technique to synthesize machinable bioactive glass-ceramics. Crystallization of a glass of the system SiO2–Al2O3–CaO–CaF2–K2O–B2O3–La2O3 indicated that miserite [KCa5(Si2O7)(Si6O15)(OH)F] is the dominant phase in the resulting glass-ceramic. The material is bioactive and can be used as an alternative to metallic titanium for dental implant applications. Recently, the hydrothermal technique has presented itself as a promising method for cheap and facile preparation of several materials, among which are different types of bioactive glasses and ceramics [23]. Formation of ceramic powder from the precursors’ solution could be assisted by microwave and ultrasonic irradiation [24]. Contrary to the sol-gel technique, simple and relatively cheap chemicals, such as Ca(NO3)2⋅4H2O, (NH4)2HPO4, NaOH and SiO2 are used as start materials in the hydrothermal technique. The obtained powders are usually calcinated at a temperature near Tg and sintered at higher temperature (1000–1100 °C) to get compact solid material. Bioactive glasses can be used as coatings for materials having good mechanical properties, but are bioinert. Examples of these materials are polycrystalline alumina Al2O3, ZrO2 and titanium metal and its alloys. As foreign bodies, implants of these materials are often surrounded by thin fibrotic encapsulation, preventing them from attaching to the organism. Coating these materials with a layer of bioactive glass allows forming composite materials that can bind to the living bone tissue and fulfill the required mechanical properties. A number of techniques can be applied to coat solid objects with a bioactive glass layer. Pulsed laser deposition is a well-known one [25–28], in which a pulsed laser beam is used for the ablation of bioglass target. The ablated flux forming the so called “plasma plume” is then condensated on the substrate. Deposition of a bioactive glass layer can be done by using sol-gel products. In this method, the substrate is immersed in the sol of glass. After withdrawal, the coating on the substrate must be dried, then heat treated at sufficiently high temperature (600–700 °C) to allow adhesion and calcination of the glass powder [29]. Radio-frequency magnetron sputtering is another way to coat substrates with bioactive glasses and glass-ceramics [30, 31]. Bioactive glass powder can be treated as an enamel to coat the substrate. Stanic et al. [32] coated yttria-stabilized tetragonal zirconia (YSTZ) by brushing an emulsion of the bioactive glass powder in a suitable binding polymer. A tightly adhered vitreous coating could be obtained by firing the dry coated samples at 1280 °C. Electrostatic spray deposition is another technique to coat bioinert substrates with bioactive glass powder. Leeuwenburgh et al. [33] deposited bioactive glass coatings on pure Ti discs. Precursor solutions with different Ca/P ratios were used to get different coatings.

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6.5 Structure of glass-ceramics

Intensity (a. u.)

Glass is generally defined as an amorphous material which lacks crystallization. If this glass is thermally treated to contain some different crystalline species, then the resultant material can be considered as a glass-ceramic. Therefore, glass-ceramics are considered polycrystalline solids produced by controlled crystallization of normal glasses. Controlled crystallization usually involves two stages, a nucleation stage and a crystallization stage. In the nucleation stage, small nuclei are formed within the parent glass. After the formation of stable nuclei, crystallization proceeds by growth of a new crystalline phase [4, 34, 35]. Fig. 6.1 shows the XRD patterns of heat-treated samples of 50SiO2–26Na2O– 21CaO–3P2O5 (mol%) glass at 620  °C (plot b) for 10 h and at 1000 °C (plot c) for 10 h, together with the spectrum of the as-prepared glass for comparison (plot a) [3]. It is realized from this figure that by subjecting the glass to a proper thermal treatment, it transforms to glass-ceramic, as featured by the peaks characteristic to apatite and wollastonite crystalline phases (spectra b and c) [3]. The amorphous hump around 2θ = 25–35˚ region is clearly observed in the asprepared glass. This indicates that the as-prepared 50SiO2–26Na2O–21CaO–3P2O5

2θ (degree) Fig. 6.1: XRD spectra of silicate glass of composition 50SiO2–26Na2O–21CaO–3P2O5 (mol%) for (a) as prepared glass, (b) heat treated sample at 620 °C, (c) heat treated sample at 1000 °C. The time of heat treatment is fixed for all treated samples at 10 hours [3].

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Heat treated at 1000 °C

Heat treated at 620 °C

As prepared

Degree of crytallinity (%)

glass has amorphous structure. On the other hand, the crystalline structure peaks become sharper and the amorphous hump nearly disappears after treating the glass at 1000 °C, resulting in crystallinity of 91 %. Increasing degree of crystallinity to 91 % with increasing temperature indicates that the initial amorphous matrix is well-crystallized by the effect of heat treatment. The relation between the temperature and crystallinity is shown in Fig. 6.2. It can be observed from this figure that the degree of crystallinity increases with increasing the temperature of treatment.

Temperature (°C) Fig. 6.2: Crystallinity-temperature dependence of composition 50SiO2–26Na2O–21CaO–3P2O5 (mol%) for as-prepared glass, heat treated sample at 620 °C, and heat treated sample at 1000  °C. The time of heat treatment is fixed for all treated samples at10 hours [3].

The features presented in Figs. 6.1 and 6.2 are in accordance with the SEM of this glass. Fig. 6.3 presents scanning electron micrographs of as-prepared sample (micrograph a) and a sample treated to 1000 °C for 10 h (micrograph b). It is shown that the increase of heat treatment temperature leads to an increase in the crystallinity, since needle-like crystals are obvious in graph b. Much attention has been focused on correlating the biocompatibility of glassceramics with the structural modifications. In this respect, the glass that solidified directly from the melt (usually contains amorphous apatite phases) should be treated, either by thermal treatment or by adding activator, to develop even a little concentration of crystalline or polycrystalline bioactive phases named apatite. The implant

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a

b

Fig. 6.3: SEM of silicate glass of composition 50SiO2–26Na2O–21CaO–3P2O5 (mol%) for (a) as prepared glass, (b) heat treated sample at 1000  °C for 10 hours [3].

material is generally bioactive if it contains apatite crystals, “the natural building substance of bone”. The most useful step therefore was to use controlled crystallization to segregate apatite crystals from the glass that solidified from the melt. It is important to determine nucleation and crystallization parameters, especially the concentration and species of nucleating agents in the glass-ceramics. These additives act as heterogeneous sites at which the nucleation of desired crystalline phases may take place. Therefore, by adjusting the type and concentration of the nucleating agents used, it is possible to control the crystallization process [3, 4, 36].

6.6 Crystallinity enhancement 6.6.1 By adding activator agents Glass-ceramics are normally produced in two steps. First, a glass is formed by any standard glass-forming process. Second, the glass sample is shaped, cooled and reheated above its glass transition temperature. The second step is sometimes repeated as a third step. In these heat-treatment processes, the sample partly crystallizes in the interior. In most cases, nucleating agents (e.g., noble metals, fluorides, ZrO2, TiO2, P2O5, Cr2O3 or Fe2O3) are added to the base glass composition to boost the nucleation process [4, 35, 36]. One of the most important commercial base glass systems is the Li2O–Al2O3–SiO2 (LAS). It may include additional components, such as CaO, MgO, ZnO, BaO, P2O5, Na2O and K2O, and some fining agents like As2O5 and SnO2 to meet special requirements. ZrO2 and TiO2 are the most commonly used nucleation agents. They can be used separately or added together to the glass batch. The very low CTE of the glassceramic obtained from this system is related to the crystalline β-quartz solid solution, which has a negative CTE [37].

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To obtain ceramic material with higher mechanical strength, a homogeneous finegrained structure is desired to be formed. So, it is important to control the nucleation and crystallization processes in terms of type of species and the amount of nucleating agents [38]. It was found that phase separation occurred on cooling from the melt in the glasses containing TiO2 as a nucleating agent and subsequent heat treatment caused the formation of large number of crystals approximately 5 nm in diameter. The formed crystals act as heterogeneous nucleation sites and allow crystallization in the remaining glass. In many glass-ceramic systems, more than one kind of nucleating agent was used to achieve optimum microstructure and properties. Müller [39] found that glasses containing TiO2 and ZrO2 as mixed nucleating agents have better nucleating efficiency. Finer grain size and higher strength could be obtained. Therefore, mixing more than one kind of nucleating agent was used to obtain optimum microstructure and properties. Apatite-containing glass-ceramics are very important for surgical implantation due to the high bioactivity; close crystallographic and chemical similarity to apatite of human bone tissue [40, 41]. The development of such materials is based on the controlled precipitation of crystalline phases from a certain base glass by the effect of nucleating agent or thermal heat treatment or both. In this situation, TiO2, Cr2O3, Al2O3 and Fe2O3 are added as individual agents in order to induce bulk crystallization of the phases. In this regard, some trials have been done to optimize and characterize the properties as well as the structure of glass-ceramics in terms of mixed (TiO2–Cr2O3), (TiO2– Al2O3) and (TiO2–Fe2O3) which served as mixed nucleating agents [4]. The introduction of nucleating agent to the main glass in the system 25NaF–21CaO–13P2O5–40B2O3 (mol%) can lower the crystallization temperature and time of crystallization. According to the DSC results and x-ray diffraction analyses [4], the sequence of phase precipitation and type of phases were defined. In the case of a composition without Cr2O3 additive, crystallization occurred at 735 °C. Crystallization temperature was found to decrease to 545 °C upon addition of one mol% Cr2O3 to the glass 25NaF– 21CaO–13P2O5–40B2O3 (mol%). Moreover, crystallization temperatures of glasses containing Fe2O3 or Al2O3 additive were found to be higher than that of the Cr2O3-containing glass. Also, Cr2O3 has a remarkable effect on the nucleation and the growth rates of apatite crystals [4], i.e. Cr2O3 is more effective than Al2O3 and Fe2O3. A larger amount of apatite particles is highly concentrated in glass containing Cr2O3 compared with that of the base glass and glasses containing Fe2O3 or Al2O3. A smaller variation in Cr2O3 content causes a larger effect than those arising from larger variation in the concentration of the other agents. So, the structural role played by (TiO2–Cr2O3) as mixed agents for crystallization of biodental glass-ceramics is considered to be more effective than that played by other mixed agents such as (TiO2–Al2O3) or (TiO–Fe2O3) [4]. Previously, it was found that the upper limit for TiO2 additives employed in 50SiO2–26Na2O–21CaO–3P2O5 (mol%) bioglasses is around 1 mol% [3]. Thus, it is recommended that bioglasses should include 1 mol% TiO2 and different amounts of Cr2O3

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as nucleating agents. The nucleation and crystallization processes were investigated in terms of changing Cr2O3 concentration. The properties and structure of the materials depend on the Cr2O3 content. For example, it is noted that thermal properties and hardness of the material are enhanced by increasing the concentration of Cr2O3.

6.6.2 By sintering process Glass-ceramics can also be produced by sintering of normal glass particles. In this case, crystallization starts at glass particle interfaces. A main advantage of the sintercrystallization process is that nucleating agents are not necessary, because the particle surfaces provide nucleation sites. A disadvantage of the sintering process is the presence of some residual porosity. However, concentration of the pores can sometimes be minimized or even eliminated by hot-pressing techniques. Sintering is also used to produce glass-ceramics from reluctant glass-forming compositions. Commercial applications of sintered glass-ceramics include devitrifying frit solder glasses for sealing TV tubes, marble-like floors and some bioactive glass-ceramics [42].

6.7 Dental glass-ceramics Generally, glass-ceramics are characterized by having several desirable properties compared to glasses and ceramics. The resistance to both surface damage and scratching is increased by controlling grain size. Also the tensile strength, resistance to wear and abrasion are enhanced. Nevertheless, the bioglass-ceramic has drawbacks such as its brittleness as in the case of other glasses and ceramics. Moreover, due to the restriction on the effective composition for biocompatibility or osteogenicity, the mechanical properties cannot be improved as for other glass-ceramics. Therefore, bioactive glass-ceramics cannot be used for making major load-bearing implants such as joint implants. However, they can be used as dental restorative composites or fillers for teeth. On the other hand, bioceramics used as dental restorations are biocompatible, have low thermal and electrical conductivities and are quite hard. These properties make them comfortable in the oral cavity. Moreover, these materials are extremely durable, relatively easy to manufacture and have superior aesthetics. Even dark tooth cores can be covered by glass-ceramic restorations. Glass-ceramic which produces highly esthetic results is ideal for fabricating teeth restorations. Dental glass-ceramics have a characteristic that they can be either pressed or machined to the desired shape in the dental laboratory. In addition, they have a natural appearance so that they look like normal teeth. As non-metal restorations, glass-ceramics have the privilege that they do not cause allergy [43].

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Various studies were concerned with the effect of combining metal implant and glass-ceramic on the efficiency of restoration. In this regard, Strunz et al. [44] coated cylinders from an alloy of CoCrMo with a layer of biocompatible adhesive enamel. The adhesive strength reached about 120 to 140 kp/cm2, showing that the method is useful in taking advantage of the stability of the metal implant and the tissue compatibility of the glass-ceramic. CoCrMo alloy coated with bioactive glass layer was also investigated by Cortés et al. [45]. They reported the formation of a bonelike apatite layer on the surface of the coating. The Ca/P ratio and the thickness of the apatite layer increases with increasing the time of immersion in simulated body fluid (SBF). Lopez-Esteban et al. [46] investigated coatings of SiO2–Na2O–K2O–CaO–MgO– P2O5 glasses on Ti-based alloys. They achieved excellent adhesion by controlling the glass composition and firing conditions. In vivo experiments on New Zealand white rabbits carried out by Drnovšek et al. [47] indicated that osseointegration and formation of bone could be enhanced on the surface of a porous Ti layer (on Ti-based implants) by filling the pores with nanoparticles of a bioactive glass.

6.8 Bioactive glass-ceramics Bioactive ceramics refer to materials that if implanted within the human body can encourage the formation of an interfacial bond with the bone. An example of a bioactive ceramic is synthetic hydroxyapatite, which encourages the formation of new bone on the surface of the implant. The physicochemical bonding between the implanted biomaterial and the bone is capable of withstanding stress caused by tension. The synthetic hydroxyapatite in the implanted glass-ceramic is useful due to its clinical use as a replacement for minor osseous parts [3]. Different principles are considered to control the bioactivity of glass-ceramics. These considerations are discussed on the basis of surface chemical studies of glassceramics containing crystalline apatite and wollastonite (A–W). The most useful principle is that the apatite species in the glass-ceramics does not play the main role in forming chemical bonds between the glass-ceramic and the bone. On the other hand, an apatite sub-layer formed on the surface of the glass-ceramic in vivo is responsible for bonding formation, since the bond between apatite crystals in the body environment can be a quite strong bond. The surface apatite layer was shown to be formed by a chemical reaction of calcium and silicate ions dissolved from the glass-ceramic with the surrounding body fluid. A P2O5-free CaO⋅SiO2 glass also formed the surface apatite and bonded to the bone [1, 3]. The same surface apatite was formed even on the surfaces of various kinds of ceramics, metals and polymers, when they were implanted near dissolving calcium and silicate ions in simulated body fluid. These results show that bioactivity of such materials is mainly based essentially on the simple components of CaO and SiO2 [1, 2].

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Heat treatment of Na2O–Ca2O–P2O5–SiO2, Na2O–Ca2O–P2O5–TeO2 and MgO–CaO– SiO2–P2O5 glasses gave glass-ceramics containing crystalline apatite (Ca10(PO4)6OHF) and beta-wollastonite (CaSiO3) in the host a glassy matrix. The bioactivity of these glass-ceramics was attributed to apatite formation on the surface of the implant in the body [1–4]. Dissolution of calcium and silicate ions from the glass is considered to play an important role in forming the surface apatite layer. It was reported that bioactive materials can be developed from (CaO, SiO2)-based glasses. For example, bioactivity of ceramics, metals and organic polymers coated with crystalline apatite was induced when such materials were placed in the vicinity of a (CaO, SiO2)-based glass in SBF. A bioactive bone cement, which is hardened within 4 min and bonded to living bone, forming an apatite, was obtained by mixing a (CaO, SiO2)-based glass powder with a neutral ammonium phosphate solution. Numerous clinical applications have shown intergrowth between glass-ceramics and human bone. Successful implants have already been made using some types of bioactive materials. Bioverits are machinable glass-ceramics that are very useful, because they can be easily modified during clinical procedures [6, 7]. A different type of low-density and high-bioactive response was prepared and studied. The glass in the system Na–Ca–Si–P–O has a Young’s modulus closer to that of cortical bone and much higher bioactivity than that of the other types of bioactive glass-ceramics. This glass-ceramic is about 30 to 50 percent crystalline, and its main phase is Na2O⋅2CaO⋅3SiO2 [1–4]. Recently, Zanotto and colleagues [48–51] have developed a new glass-ceramic, known as (Biosilicate®), that is based on the same Na–Ca–Si–P–O system. This glassceramic is characterized by greater than 99.5 percent crystallinity and is as bioactive as Hench’s “gold standard” bioglass 45S5 [52]. Another interesting class of bioactive ceramics is heat-generating bioactive or biocompatible glass-ceramics intended for use for hyperthermic treatment of tumors. For instance, the glass Fe2O3–CaO–SiO2–B2O3–P2O5 is easily crystallized to high level. This type of glass-ceramic should contain magnetite and wollastonite crystals and is characterized by its high-saturation magnetization. This glass-ceramic formed a calciumand phosphorous-rich layer on its surface and tightly bonded with bone within about eight weeks of implantation. In contrast, the parent glass did not efficiently form the calcium- and phosphorous-rich layer and could only bond with bone after 25 weeks. Under an external magnetic field, granules of this glass-ceramic filled-in bone have the ability to induce heat distribution to the surrounding bone. This induced heating process plays an important role in tumor treatment. Then, heat-generating bioactive or biocompatible glass-ceramics may be considered as promising materials intended for hyperthermic treatment of tumors [36].

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6.9 In vitro and in vivo test for bioactivity Calcium phosphate ceramic such as hydroxyapatite (HA) is a good candidate for bone substitutes due to its chemical and structural similarity to bone minerals. The bone mineral consists of tiny hydroxyapatite crystals in the nanoregime. Nanostructured hydroxyapatite is also expected to have better bioactivity than coarser crystals. In a recent study, glass-ceramics obtained from glasses in the silicate and tellurate systems were investigated [3, 34, 52]. Fourier transform infrared (FTIR) spectroscopy was used to identify the functional groups. X-ray diffraction (XRD) analysis was carried out to study the phase composition, crystallinity and the crystallite size of hydroxyapatite nanopowders that were sintered at different temperatures. The in vitro test was performed in a simulated body fluid (SBF) medium. The dissolution of calcium ions in SBF medium was determined by an atomic absorption spectrometer (AAS). FTIR results combined with the x-ray diffraction exhibited a single phase of hydroxyapatite with carbonate peaks in the FTIR spectrum. Most studies [1–4] indicate that increasing the sintering temperature increases the crystallinity and the crystallite size of hydroxyapatite nanopowders. The dissolution rate of hydroxyapatite nanopowders was higher than that of conventional hydroxyapatite powders and closer to biological apatite due to its nanostructure dimensions. It was concluded that sol-gel prepared hydroxyapatite nanopowders had superior bioresorption and similar chemical and crystal structure to natural bone apatite [2–4]. This can be indicated by the decreasing intensity of peaks representing crystalline phases, especially those at 2θ = 25–35˚. The decrease in XRD peaks after soaking in SBF is a common feature of bioactive glasses and glass-ceramics. Fig. 6.4 shows the effect of soaking time on the intensity of peaks for the glass 25NaF–21CaO–12P2O5–40B2O3–1TiO2–1Cr2O3 (mol%) [4]. It was previously inferred from mechanical tests that HA and ß-tricalcium phosphate (B-TCP) could become chemically tightly bonded to bone [53]. Like any implant, HA implants are not in direct contact with the host bone. Extensive bone remodeling takes place soon after implantation. Immature bone first fills the space between implant and bone, and is later replaced by mature bone. Thus there are two interfaces between implant and host bone. The first interface is the one between the old and the new bone tissues (called cement line). It can be described as an amorphous dense line and is composed of a mineralized, collagen-free matrix. The second interface is the one between the newly-formed bone and the implant material [53]. There are two types of bonds that should be formed between the implant and the natural bone. Firstly, various chemical bonds such as the hydrogen bond and the electrostatic ionic bond were proposed for bone bonding to ceramics, or glass including a direct ionic bond such as epitaxial HA growth on a substrate. Secondly, biological bonding is considered another bone bonding mechanism. In this respect, organic fibers key a bone / implant interface by inserting into the surface layer of the bioactive material. The interface is strengthened and toughened by the organic fibers. It was reported that the biological bonds appear as layers with amorphous polymerized

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Intensity (a. u.)

gel-like material. Moreover, the bonding was established through the incorporation of collagen fibers into the gel which itself formed a part of the implant surface. This proposed biological bonding mechanism has potential significance for strengthening the interface of bone and bioactive materials.

2θ (degree) Fig. 6.4: X-ray diffraction patterns of the glass 25NaF–21CaO–12P2O5–40B2O3–1TiO2–1Cr2O3 (mol%); (a) before immersion in SBF, (b) after 3 days in SBF and (c) after 7 days in SBF at 37 °C [4].

Although the basic chemical and biological mechanisms involved in forming a bioactive bond to bone tissue have been documented [3, 53], details of the processes involved at the atomic or molecular levels remain unclear. In this respect, an in vitro test is the one issue that determines the mechanism for formation of the bone apatite layer prior to bone bonding. The importance of the apatite is that it can derive the biological profile of the bioactive materials. The biological bonds can be considered or established through the in vivo test; through the incorporation of collagen fibers into the gel which itself formed a part of the implant surface. Therefore both in vitro and in vivo tests are important to study and follow the progress of chemical and biological bonds which might be involved in forming a bone bonding. The in vivo test of biomaterials ensured their biocompatibility and bioactivity with animal tissues, such as rabbit, since the osteocytes (bone forming cells) grew on the material-tissues interface. Moreover, the interface of bonding between tissues and implant is narrow and zigzag-like, showing an intense bonding with bone and rapid

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ingrowth of bone tissues. Figs. 6.5 and 6.6 show bone section stained with hematoxylin and eosin. The figures present the last stage of bone healing, in which the accumulation of osteoblasts is clearly visible and the cells were in a remodeling state [3].

Fig. 6.5: Bone section stained by hematoxylin and eosin after 30 days of implantation of bioactive silicate glass of composition 50SiO2–26Na2O–21CaO–3P2O5 (mol%) in trabecular bone of a New Zealand rabbit [3].

Fig. 6.6: Bone section stained by hematoxylin and eosin after 4 weeks of implantation of bioactive tellurate glass of composition 50TeO2–26Na2O–21CaO–3P2O5 (mol%) in trabecular bone of a New Zealand rabbit [3].

From the above figure, it can be noted that the in vivo test of the studied biotellurate glass 50TeO2–26Na2O–21CaO–3P2O5 (mol%) ensured their bioactivity with animal tissues interface. The interface of bonding with bone after implantation (4 weeks) showed rapid growth of bone tissues with a rate higher than that of 50SiO2–26Na2O– 21CaO–3P2O5 glass.

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[48] Roriz VM, Rosa AL, Peitl O, Zanotto ED, Panzeri H, de Oliveira PT. Efficacy of a bioactive glassceramic (Biosilicate) in the maintenance of alveolar ridges and in osseo integration of titanium implants. Clin. Oral Implants Res 2010, 21, 148–55. [49] Granito RN, Ribeiro DA, Rennó ACM et al. Effects of biosilicate and bioglass 45S5 on tibial bone consolidation on rats: A biomechanical and a histological study. J Mater Sci: Mater Med 2009, 20, 2521–26. [50] Moura J, Teixeira LN, Ravagnani C et al. In vitro osteogenesis on a highly bioactive glassceramic (Biosilicate®). J Biomed Mater Res A 2007, 82, 545–57. [51] Tirapelli C, Panzeri H, Lara EHG, Soares RG, Peitl O, Zanotto ED. The effect of a novel crystallised bioactive glass-ceramic powder on dentine hypersensitivity: a long-term clinical study. J Oral Rehabil 2011, 38, 253–62. [52] Kamal H. Effect of Nucleating Agents on the Structure and Properties of Some Bioactive Glasses: MSc thesis. Mansoura University, Egypt, 2005. [53] Jones JR, Hench LL. Factors affecting the structure and properties of bioactive foam scaffolds for tissue engineering. J Biomed Mater Res B: Appl Biomater 2004, 68, 36–44.

M. Shamshi Hassan, Touseef Amna, Mohamed Bououdina and Myung-Seob Khil

7 Metal oxide-based one-dimensional titania nanostructures via electrospinning: Characterization and antimicrobial applications 7.1 Introduction Nanostructured materials such as nanoparticles, nanofibers, nanorods, nanospheres, nanotubes and nanowires present new features and opportunities for enhanced performance in many promising applications (Costa LL et al., 2009; Seo MH et al., 2009; Wang BX et al., 2008). Nanotextured materials are of special significance owing to their excellent physicochemical properties such as catalytic, electronic, magnetic, mechanical and optical (Poudel B et al., 2005). They are widely applied in air and water purification technologies, photocatalysis, gas sensors, high effect solar cells and microelectronic devices (Nakahira A et al., 2004; Yu JG et al., 2006). Nanostructured materials with diverse morphologies have varying specific properties and, hence, the hybrid or/new applications of such materials are related to the shape and size of the nanotextured materials (Wang FM et al., 2007). The synthesis of nanostructured materials with specific shape and size, as well as the understanding of their formation mechanism are extremely important research aspects of materials science (Wang FM et al., 2007). Titanium dioxide (TiO2) or titania is a very well researched and celebrated material due to the stability of its chemical structure, biocompatibility, physical, optical and electrical properties. It exists in four mineral forms (Gianluca LP et al., 2008), viz: anatase, rutile, brookite and titanium dioxide (B) or TiO2 (B). Anatase type TiO2 has a crystalline structure that corresponds to the tetragonal system (with dipyramidal habit) and is used mainly as a photocatalyst under UV irradiation. Rutile type TiO2 also has a tetragonal crystal structure (with prismatic habit) and mainly used as white pigment in paints. Brookite type TiO2 has an orthorhombic crystalline structure. TiO2 (B) is a monoclinic mineral and is considered as a newcomer to the titania family. TiO2, therefore is a versatile material that has applications in various products such as paint pigments, sunscreen lotions, as a food coloring agent, electrochemical electrodes, capacitors, solar cells and in toothpastes (Meacock G et al., 1997). TiO2, especially its anatase phase, has attracted much attention for its potential application in degradation of various environmental pollutants, both gaseous and liquid (Ao CH et al., 2004). Moreover, TiO2 is one of the most investigated material for antibacterial applications. However, its shortcomings include a large band gap (~3.2 eV) which causes most of the solar spectrum to remain unutilized. To extend the optical absorp-

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tion of TiO2 region, various dopants (such as monometallic and bimetallic) have been added to the oxide to improve its photocatalytic and solar efficiency (Asahi R et al., 2001; Burda C et al., 2003; De Vos et al., 2002; Martyanov IN et al., 2004; Sakthivel S et al., 2003). Many previous investigations focused on the transition metals as dopants (Sun B et al., 2005; Wu JCS et al., 2004; Zhang W et al., 2004). Rare earth metals having incompletely occupied 4f and empty 5d orbitals have been found to promote catalysis, that is, the photocatalytic activity of TiO2 (Baiju KV et al., 2005; Xu AW et al., 2002). Wang et al., 2005 and Zhang et al., 1998 indicated that introducing two or more appropriate elements into nanocrystalline TiO2 particles could improve the photocatalytic effect of TiO2, showing the salutary effect of co-doping with transition and rare earth metals into nanocrystalline TiO2. Furthermore, nanowires, nanotubes and nanofibres of TiO2 have been successfully prepared by various methods such as electrochemical synthesis, template-based synthesis as well as the chemical-based route and the effectiveness and practical use of some of these materials for various applications are still being evaluated. Though several TiO2 nanocomposites with antimicrobial capabilities have been reported, there are still barriers in their antibacterial application under dark conditions. In the TiO2 nanocomposites, biocidal efficiencies depend on their light absorbance under UV and visible light. Therefore, most TiO2 compounds exhibit no antimicrobial performances in the dark condition. For this reason, some approaches with metal-doped TiO2, i.e., silver, palladium, and copper coated titania compounds, have been reported to provide efficient antibacterial activities without light irradiation (Mahltig B et al., 2007; Page K et al., 2007; Sunada K et al., 2003; Zhang H et al., 2009). In this chapter, we discuss the fabrication of one-dimensional titanium dioxide and metal oxide-doped titania nanofibers via electrospinning technique. The chapter is comprised of different sections and each section has been designated by a particular heading/title. We briefly describe nanostructured titania and its versatile properties. Furthermore, the doping and the utilization of TiO2 and doped titania nanofibers are discussed. Here, we mainly focus on metal oxide semiconductors such as zinc oxide, cobalt oxide, cerium oxide, nickel oxide and copper oxide nanoparticles which are generally used as dopants of titania to boost up the efficiency of titania. Additionally, an overview of the chemistry and biology of the aforementioned nanoparticles (NPs) is described. However, the focus is primarily on the antimicrobial applications. The antimicrobial activities of manufactured nanofibers are briefly presented. It has been observed that the preparation methods, doping with metal oxides (such as zinc oxide, cobalt oxide, cerium oxide, nickel oxide and copper oxide), morphologies, and the size of nanofibers show a significant impact on the toxic properties toward microorganisms. Nevertheless, this chapter provides an insight on the enhanced bactericidal effect of inorganic metal oxides-doped titania.

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7.2 General routes/procedures for the synthesis of nanofibers Presently, a number of processing techniques such as drawing (Ondarcuhu T et al., 1998), template synthesis (Feng L et al., 2002; Martin CR, 1996), phase separation (Ma PX et al., 1999), self-assembly (Liu GJ et al., 1999; Whitesides GM et al., 2002), electrospinning (Deitzel JM et al., 2001; Fong H et al., 2001), and so forth are used to synthesize polymeric nanofibers. Here in this chapter we discuss the electrospinning process. Moreover, the advantages of the electrospinning technique for the fabrication of nanofibers will also be addressed.

7.3 Electrospinning process Electrospinning is an inexpensive, effective and straightforward method of producing nonwoven nanofibrous mats. It (e-spinning) has been recognized as an efficient technique for the fabrication of polymeric nanofibers, organic-inorganic hybrid, inorganic-polymer composites, organic-organic composites, bimetallic-polymer composite nanofibers and ceramic nanofibers, etc. During recent years a variety of polymers have been successfully electrospun into ultrafine fibers, generally in solvent solution and some in melt form. In the past decade, electrospinning has attracted great attention due to an increased interest in nanoscale characteristics and technologies (Fridrikh SV et al., 2003; Li D et al., 2005; Sun XY et al., 2007). This technique has become a versatile and valuable route in the production of exceptionally long polymeric fibers with uniform diameters (ranging from nanometers to micrometers) (Huang ZM et al., 2003; Theron SA et al., 2004). The necessary components of an electrospinning apparatus include a high power voltage supply, a capillary tube with a needle or pipet, and a collector that is usually composed of a conducting material. The collector is held at a relatively short distance from the capillary tube, which contains a polymeric solution connected to the high power supply. Fig. 7.1 shows the electrospinning setup. In electrospinning, typically, a polymer solution or melt is extruded through a capillary tube to form a small droplet at the spinneret tip. The application of voltage between the tip of the spinneret and the collector generates surface charges in the polymer droplet, and when the applied voltage is above a critical value, electrostatic forces overcome the surface tension in the polymer drop causing it to stretch to form a Taylor cone and eventually a liquid jet that is accelerated toward the grounded collector. During the journey of the polymeric jet, it experiences stretching movement and whipping motion leading to continuous ultrathin randomly oriented fibers in the form of a nonwoven mat (Fig. 7.2). It is generally recognized that various factors control fiber formation (such as beaded/or beadfree) and diameter. Among other aspects, the parameters which greatly influence the fiber morphology were recognized as solution viscosity, surface tension, and electrical conductivity, etc. (Fong H et al., 1999). It has been observed that higher viscosity

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Negative terminal

and conductivity as well as lower surface tension tend to yield bead-free fibers, although extreme values can lead to practical difficulties in electrospinning (Kim B et al., 2005). In summary, the electrospinning technique has succeeded within the last years in establishing itself as an internationally recognized, highly versatile, enabling method granting access to a broad range of one dimensional functional nano-objects (such as

Cu wire

Syringe Sol-gel solution Charged Jet Rotating collector

Positive terminal Voltage supplier

Fig. 7.1: Schematic diagram for the utilized electrospinning setup.

Fig. 7.2: Representative digital image of nonwoven mat.

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nanofibers, nanorods as well as nanotubes). Significant progress has recently been made not only in the fundamental understanding and modeling of the details of the complex processes governing fiber formation but also in the development of a broad range of technical spinning devices. Nevertheless, the research fields which have gained abundant benefits from such developments are biomedical and life sciences with tissue engineering, drug delivery, wound healing, antimicrobial applications being characteristic targets. Electrospinning can without any doubt be expected to contribute significantly in diverse fields of nanoscience in near future.

7.4 General applications of electrospun nanofibers Electrospinning is a cost effective method for producing nonwoven nanofibrous mats, which intrinsically have 103 times larger surface to volume ratios, increased flexibility in surface functionalities, improved mechanical properties, and smaller pores than fibers produced using traditional methods. Potential applications based on such fibers, specifically their use as reinforcement in nanocomposite development have been realized. Although electrospun nanofibrous polymers are widely recognized to have potential applications in filtration, catalysis, and sensing (Frenot and Chronakis, 2003), they may also be useful in a wide variety of biomedical applications (Bhattarai N et al., 2006; Zhang SG, 2004; Zhang and Ma, 2000). In particular, polymer nanofibers can serve as tissue scaffolds, providing mechanical support for cellular activities and growth (Rho KS et al., 2006; Lim and Mai, 2009; Li WJ et al., 2009), because they resemble the body’s natural extracellular matrix (ECM) (Pham QP et al., 2006). Previous studies show that scaffolds with nanoscale features better support cell attachment and proliferation when compared to scaffolds with micrometer size structures due to increased cellular attachment (Bhattarai N et al., 2006; Pattison MA et al., 2005; Tuzlakoglu K et al., 2005; Li WJ et al., 2003; Li and Xia, 2004; Geng XY et al., 2005). Furthermore, electrospun nanofiber membranes have excellent porosity, which is essential to allow for better cell adhesion, cell ingrowth, and nutrient exchange during in vivo or in vitro cell culture (Pham QP et a., 2006; Toh YC et al., 2003). Moreover, nanotextured materials with open porosity, such as electrospun nanofiber mats, possess specific surface area (in the range 10–100 m2/g), which makes them attractive candidates for nanoparticle supports in water purification processes (Tuzlakoglu K et al., 2005) [99]. In addition, nanofibers can be used as filters since pores with sizes in the range ~1–10 μm can catch pollutants more efficiently than standard filters (Li WJ et al., 2003). In addition to filtration, nanofibers are also shown to be effective for wound dressing (Li and Xia, 2004; Geng XY et al., 2005). Besides to the aforementioned applications, the nanofibers can also be exploited for antimicrobial purpose. In this chapter we primarily address the antimicrobial functionalities of pristine TiO2 and metal oxide-based TiO2 nanofibers.

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7.5 Antimicrobial applications of metal oxide-based nanotextured materials/nanofibers Nanofibers with antimicrobial functionalities can facilitate development of very efficient materials such as membranes, fabric, filters, etc. which can be used for a wide variety of applications. In this respect, TiO2 and silver nanoparticles present themselves as the two best possible candidates because TiO2 can perform in the presence of UV light, whereas silver is active without UV light due to its intrinsic antimicrobial capability. Enhancement of the photocatalytic abilities of TiO2 can be achieved through modification of titania with metal deposited on its surface. Deposited metal can improve the photocatalytic reaction by acting as an electron trap to prevent electron-hole recombination, increasing the availability of electrons and holes for redox reactions. The presence of metal also shifts the absorption in the visible range, allowing for utilization of visible light. Antibacterial agents are of great importance in several industries, e.g., water disinfection, textiles, packaging, construction, medicine, and food (Yamamoto O, 2001; Li QL et al., 2001). The use of metal oxides as antimicrobial agents has the advantage of improved safety and stability, as compared to organic antimicrobial agents (Brayner R et al., 2006). Since most of the biological processes take place at the nanoscale level, a combined application of nanotechnology and biology can possibly solve important biomedical problems (Sandstead HH, 2006). Nanosized metal oxide particles are receiving ever-increasing attention for a large variety of applications. Titanium dioxide (TiO2) and zinc oxide (ZnO) nanoparticles are included in toothpaste, beauty products, sunscreens, and textiles. Ceramic nanopowders of metal oxides such as ZnO have been found to exhibit a marked antibacterial activity (Zhang L et al., 2006; Sawai J, 2006; Seven O et al., 2006). Nanotechnology has the potential to revolutionize the global food system. Nanomaterials that can be used in biological systems are required to be biocompatible (Shin YJ et al., 2007). The bactericidal effectiveness of metal nanoparticles has been suggested to be due to both their size and high surface-to-volume ratio. Such characteristics should allow them to interact closely with bacterial membranes, rather than the effect being solely due to the release of metal ions (Morones JR et al., 2005). The mixing of polymers and nanoparticles is opening new avenues for engineering flexible composites that exhibit desired and beneficial characteristics such as optical, mechanical properties and so forth (Lin Y et al., 2005). Additionally, metal nanoparticles alone or nanoparticle-polymers composites or coated onto surfaces are known to have a variety of potential antimicrobial applications. It is already well known that TiO2 photocatalytically disinfects bacteria but under UV light (Alrousan MAD et al., 2009; Yeung KL et al., 2009). Pd/TiO2 and Pd/SnO2 (Erkan A et al., 2006) thin films also require UV-A light for photo-disinfection. Antimicrobial nanoparticles that have been synthesized and tested for applications in antimicrobial packaging and food storage boxes include silver oxide nanoparticles (Sondi and Salopek-Sondi, 2004), zinc oxide, magnesium oxide nanoparticles (Jones N et al., 2007) and nisin particles produced

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from the fermenation of bacteria (Gadang VP et al., 2008). The use of metal oxides as antimicrobial agents has the advantage of improved safety and stability as compared to organic antimicrobial agents (Brayner R et al., 2006).

7.6 Concept of doping and composite nanofibers Modern trends in nanomaterials synthesis deal with the preparation of hybrid nanoconstructs/functional nano-objects with a variety of architectures such as nanowires, nanofibers, core shell, nanoflowers, and so on through various synthetic routes. A doping agent, commonly referred to as a dopant, is a trace impurity that is inserted into a substance (in very low concentrations) in order to alter/or enhance the basic properties of the substance. The controlled addition of known chemical species results in dramatic changes of various properties such as magnetic, electrical, optical, electronics, sensing and structural properties of basic material. Bimetallic nanomaterials, composed of two unlike metal elements, are of greater interest than the monometallic materials because of their improved characteristics. Generally speaking, great efforts have been devoted to the synthesis and development of metal-based nanocomposite materials. Especially, bimetallic nanoparticles composed of two different metal elements are of greater interest than the monometallic ones for the enhancement of the desired specific properties. This is because bimetallization can improve the properties of the original single metal and create a novel hybrid property, which may not be achieved by monometallic materials. The chemical composition as well as physical and chemical properties of the electrospun nanofibers can be tailored by selecting suitable polymers and other organic and/or inorganic components. The preparation of metal oxide nanofiber composites is challenging but can potentially lead to novel structures with the combined attributes of the constituent materials. This can be achieved by varying precursor ratios and tailoring the crystal structures during calcination. These mixed metal oxide ceramic nanofiber structures can be flexible and withstand high temperatures and are, therefore, attractive materials for use in a variety of applications. Electrospinning has been exploited for the preparation of metals, metal oxide semiconductor nanofibers, and of the corresponding polymer composites either by combining it with sol-gel processes or by following a polymer-based precursor route. Precursors of the target materials such as metal salts are introduced for this purpose into spinning solutions containing polymers. The precursor materials are, in a subsequent step, reduced to the target metals or metal oxides via thermal treatment or treatment in the presence of reducing agents such as, for instance, hydrogen. The results are polymer composite nanofibers. In a further step, the polymer can be removed completely, in general again by thermal treatment to yield nanofibers from the pure metal, metal oxide materials. A wide variety of nanofibers with specific properties were prepared via electrospinning based on polymer composite fibers with dispersed nanoparticles, for example NiO/TiO2 composite nanofibers (Amna T et al.,

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2013), cobalt-doped titanium dioxide nanofibers (Amna T et al., 2013c), Fe3O4/TiO2 hybrid nanofibers (Amna T et al., 2013a), Ce2O3/TiO2 composite nanofibers (Hassan MS et al., 2012), CuO/TiO2 composite nanorods (Hassan MS et al., 2012), other metal oxides or also based on composite materials with highly complex compositions (such as bimetallic and trimetallic nanocomposites) (Hassan MS et al., 2013a). Nevertheless, intensive research study has focused on designing the nanostructures of TiO2 and its composites. In recent years, a number of dopants have been tried in order to enhance the photocatalytic performance of titanium. Hassan et al. (2012a) carried out the comparative study of photocatalytic behavior of TiO2 nanofibers doped with rare earth elements. The experience with dopants has been varied because of their different roles in trapping electrons and/or holes on the surface. However, in this chapter we concentrate on the development of pristine TiO2 and its composite nanofibers fabricated mainly by electrospinning technique. The fabrication and characterization of metal oxide-based TiO2 nanocomposites is briefly described. The dopants of titania which are addressed in this chapter include zinc oxide, copper oxide, nickel oxide, cobalt oxide and cerium oxide along with their applications (Fig. 7.3). The pristine TiO2 nanorods fabricated through facile electrospinning method are described below.

Cerium Food packaging Cobalt

Titania

Nickel

Antimicrobial Applications

Water purification Textile fabrication

Copper

Zinc

Fig. 7.3: Graphical representation of titania dopants and their applications.

7.7 Development of pristine TiO2 nanofibers via electrospinning technique The synthesis of titania nanorods was carried out by electrospinning technique. In a typical procedure, polyvinyl acetate (PVAc, Mw=500,000 Aldrich, USA) solution (18 wt%) was prepared by dissolving PVAc in N, N-dimethylformamide (DMF, 99.5 assay, Showa Co. Japan) under magnetic stirring for 8 h at room temperature. 5 g of titanium isopropoxide (TIPP, 98.0 assay, Junsei Co. Ltd., Japan) was taken in a separate bottle and then a few drops of acetic acid were added until the solution became transparent. Then 6 g of PVAc solution was mixed slowly into the solution under vigorous stirring. The obtained solution was put in a 10 ml syringe with a stainless steel

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   115

needle. A copper pin connected to a high voltage generator was inserted in the solution as a positive terminal whereas a ground iron drum covered by a polyethylene sheet served as counter electrode. The solution was kept in the capillary by adjusting the inclination angle. A voltage of 20 kV was applied to this solution. The distance between the syringe needle tip and collector was fixed at 18 cm. The as-synthesized mat of composite fiber was initially dried at 80 °C for 24 h under vacuum and then calcined at 600 °C for 2 h in air atmosphere with a heating rate of 2 °C/min. Titanium (IV) isopropoxide is well known as a feasible parent material for the synthesis of pure titanium dioxide. PVAc, was chosen as a template because of its easy availability, hydrophobicity and no designed crosslinks. A new hybrid organic-inorganic solution was obtained by the chemical crosslinking reaction between TiP and PVAc. This process was carried out in an N,N-dimethylformamide solution, via a controlled crosslinking process that yielded macroscopically to a homogeneous and transparent solution. In hybrid solution system, transesterification of PVAc occurs and interchain links are formed. Titanium isopropoxide has been used as catalyst for transesterification reaction of PVAc. Furthermore, titanium is one of the most active metals for catalysis of ester polycondensation. Some strong interactions existed between PVAc and TiP phase, avoiding a macroscopic phase separation. However, it is difficult to get nanofibers by electrospinning highly crosslinked solution because of their high viscosity and surface tension. Therefore, acetic acid was used as a chelating agent of titanium isopropoxide, to decrease the crosslinking degree of TiP and PVAc. As a result of this, the crosslinking of PVAc through transesterification reactions with TiP were controlled to get a solution with suitable viscosity for electrospinnig. After calcinations of the hybrid nanorods, all organic groups were eliminated and titanium dioxide was formed as nanorods. Fig. 7.4 shows the XRD pattern of titania nanorods annealed at 600 °C and commercially available titania (P25). The XRD of the titania nanofibers shows (Fig. 7.4b) anatase (A) crystalline phase of TiO2 without any other phase or impurity (JCPDS No. 89-4921) (Sorapong P et al., 2005). Whereas P25 (Fig. 7.4a) is showing the presence of rutile (R) phase although anatase phase is the major phase. From the (101) peak of anatase TiO2, the average size of crystallite (D) was calculated using the Scherer equation: D = Kλ/βcosθ where β is the full width half maximum (FWHM) of the 2θ peak, K is the shape factor of the particles (it equals to 0.95), θ and λ are the incident angle and the wavelength of the X-rays, respectively. The average crystallite size of titania nanorods determined is ca. 3.5 nm as compared with ca. 4.5 nm of P25. Fig. 7.5a shows the FESEM image of nanorods having the uniform diameter size. Fig. 7.5b shows TEM image of nanorods having diameter in the range of 200–300 nm. TEM image shows smooth surface of nanorods without any structural defects. The individual TiO2 nanorods were comprised of small densely-packed nanocrystals. The ring patterns of selected area electron diffraction patterns demonstrate the polycrystallinity of TiO2 nanorods (inset of Fig. 7.5b). The aspect ratio of a shape is defined as the length of the major axis divided by the width of the minor axis (Catherine and Nikhil, 2002). Thus,

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Arbitrary unit (a. u.)

A

a

R A

b

2θ (degree) Fig. 7.4: XRD patterns of (a) commercially available titania P25 and (b) titania nanorods. Reprinted with permission from Hassan et al. (2012). Copyright 2012, American Scientific Publishers.

a

b

100 nm c

d

Fig. 7.5: (a) FESEM (b) TEM image (c) EDX spectra of titania nanorods (d) and thermal gravimetric analysis in an argon atmosphere and the corresponding first derivative of titania/PVAc composite. Reprinted with permission from Hassan et al. (2012). Copyright 2012, American Scientific Publishers.

spheres have an aspect ratio of 1. Herein, TiO2 nanorods have width of ~200–300 nm and aspect ratios between 15~25. Fig. 7.5c shows the energy-dispersive X-ray spectroscopy (EDX) analysis of the titania nanorods. It has been observed that the sample contains Ti and O; no other element impurity is detected, indicating the final product is free of impurity. TGA thermograms of titania/PVAc composite were shown in Fig. 7.5d. As pointed out by TGA, the composite sample showed a total of ~ 65 % weight loss. The first step of weight loss (~ 5 %) was due to the evaporation of water and solvent in the range of 25–200 °C. After 200 °C, the weight loss became more than 52 % due to the degradation of PVAc

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polymer up to 400 °C. The (~ 8 %) weight loss in the temperature range of 400–500 °C was was attributed to the complete decomposition of PVAc and the crystallization of amorphous TiO2 to anatase TiO2. On the first derivative curve, a main exothermic peak was observed at ~ 310 °C, suggesting the thermal events related to the decomposition of titanium tetra-isopropoxide along with the degradation of PVAc by dehydration on the polymer chain, which was confirmed by a dramatic weight loss in TG curve at the corresponding temperature range (200–400 °C). However, to improve the photocatalytic performance of titanium dioxide (TiO2), various approaches have been developed such as metal/nonmetal-doped TiO2, metal/semiconductor-coupled TiO2, and polymer– TiO2 composites. Deposited metal can improve the photocatalytic reaction by acting as an electron trap to prevent electron-hole recombination, increasing the availability of electrons and holes for redox reactions (Qin and Wang, 2006; Doshi and Reneker, 1995). The presence of metal also shifts the absorption into the visible range, allowing for utilization of visible light. At this point, some selected metal-based nanocomposites of TiO2 will be summarized and also their antibacterial performance reviewed briefly. For instance, zinc-doped titania nanofibers are discussed concisely below.

7.8 Doping of titania with metal oxide 7.8.1 Doping of titania with zinc The remarkable properties of zinc oxide (ZnO) have attracted the interest of many researchers in the past few years. ZnO is a direct band gap semiconductor with band gap energy of 3.36 eV at room temperature, high exciton binding energy of 60 meV and high dielectric constant (Singh S et al., 2007). Hence, the luminescent properties of ZnO have attracted considerable attention due to its potential application in ultraviolet light emitting devices. This band gap semiconductor has numerous potential applications, particularly in the form of thin films, nanowires, nanorods or nanoparticles and can be introduced to optoelectronic and electronic devices (Starowicz and Stypula, 2008). They also can be used in the production of chemical sensors and solar cells (Singh S et al., 2007). One of the most remarkable commercial applications of ZnO nanoparticles is their use in the production of sunscreens and cosmetics, due to their property of blocking broad UV-A and UV-B rays (Huang Z et al., 2008). ZnO nanoparticles are believed to be a non-toxic, biosafe and biocompatible nanomaterial (Zhou J et al., 2006), although few reports have described some toxicological activity of ZnO nanoparticles. There are few reports on the antimicrobial properties of ZnO nanoparticles against Staphylococcus aureus and Escherichia coli (Jones N et al., 2008; Brayner R et al., 2006). However, the interaction of nanoparticles with micro-organisms and biomolecules is an extensive field of research which is largely unexplored. In this regards, Amna et al. have reported the successful fabrication of zinc-doped titania by electrospinning technique (Amna T et al., 2012). Additionally, the authors have described the possible interaction

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Intensity (a. u.)

mechanism for enhanced antibacterial activity. Zn-doped titania nanofibers were synthesized by using precursor Zn (NO3)2⋅6H2O and titanium isopropoxide via electrospinning. Fig. 7.6 shows the XRD pattern of Zn-doped titania nanofibers annealed at 600 °C, which reveals the formation of TiO2 anatase crystalline phase without any other phase or impurity (PDF-89-4921). The low Zn content and the large difference between the ionic radii of Zn2+ (0.88 Å) and Ti4+ (0.745 Å) make it difficult for Zn2+ to substitute Ti4+. Therefore, it is possible that Zn ions exist mainly in the form of ZnO clusters and dispersed on TiO2 crystallite surfaces. The average crystallite size (ca. 3.17 nm) of Zndoped titania powder was determined from the broadening of the anatase (101) peak (2θ = 25.2). Fig. 7.7 shows FE-SEM micrographs of Zn-doped titania nanofibers at low and high magnifications, respectively. Fig. 7.7a shows the presence of nanofibers having uniform diameter size, whereas Fig. 7.7b shows a higher magnification image of nanofibers having a diameter in the range of 200–300 nm. Fig. 7.8a shows the energy disper-

2θ (degree) Fig. 7.6: XRD pattern of the Zn doped titania nanofibers calcined at 600 °C. Reprinted with permission from Amna et al. (2012). Copyright 2012, Springer.

a

b

Fig. 7.7: FE-SEM image of the Zn doped titania nanofibers at (a) low and (b) high magnification. Reprinted with permission from Amna et al. (2012). Copyright 2012, Springer.

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sive X-ray spectroscopy (EDX) analysis of Zn-doped titania nanofibers. It has been observed that the sample contains Ti, Zn, and O; no other element impurity is detected, indicating the final product is pure and free of impurity and composed of ZnO and TiO2. Figs. 7.8b and 7.8c show EPMA of Zn-doped titania nanofibers. The presence of Zn particles in the titania is confirmed from the EPMA image. The EPMA image clearly shows that Ti is the main element, and Zn is also uniformly dispersed on the surface of the titania. However, in this study, ZnO phase could not be detected by XRD because its content may be below the XRD detection limit. According to the crystal lattice parameters, ZnO and TiO2 cannot combine in a single crystal. Accordingly, a main TiO2 nanofiber doped with ZnO nanoparticles is the expected structure of the obtained product. To affirm this hypothesis, TEM analysis has been conducted (Fig. 7.9). As shown in Fig.  7.9a, tiny black particles distributed along the nanofibers can be observed. As shown in the high resolution (HR)-TEM image (Fig. 7.9b), the black nanoparticles do have different crystal lattice parameters than the matrix as the crystal plane is closer, so one can claim that these nanoparticles represent ZnO. The antimicrobial activity against E. coli ATCC 52922 (E. coli) and S. aureus ATCC 29231 (S. aureus) was assessed in vitro, and an attempt was made to find the minimum inhibitory concentration of the

TiKa,8

a

ZnLa1, 4

b

c

Fig. 7.8: (a) EDX spectra of the zinc doped titania nanofiber (b, c) EPMA mapping result of zinc doped titania nanofiber. Reprinted with permission from Amna et al. (2012). Copyright 2012, Springer.

a

200 nm

b

5 nm

Fig. 7.9: TEM image of the prepared ZnO-doped TiO2 nanofibers at (a) low resolution and (b) HR-TEM image, the red circle refers to a ZnO nanoparticle. Reprinted with permission from Amna et al. (2012). Copyright 2012, Springer.

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nanofibers capable of inhibiting the growth of the above-mentioned pathogenic strains. To understand the antibacterial activity and acting mechanism of synthesized Zndoped titania nanofibers deeply, the structural and morphological changes of E. coli ATCC 52922 and S. aureus ATCC 29231 bacterial cells were studied after exposure by nanofibers. The antibacterial activity of nanofibers against the above-mentioned bacteria were investigated by calculation of minimum inhibitory concentration (MIC) and analyzing the morphology of the bacterial cells following the treatment with nanofiber solution. The lowest concentration of Zn-doped titania nanofiber solution inhibiting the growth of S. aureus ATCC 29231 and E. coli ATCC 52922 strains was found to be 0.4 μg/ ml and 1.6 μg/ml, respectively (Fig. 7.10). Furthermore, Bio-TEM analysis (Fig. 7.10) dem-

a

b

1.0 μm

200 nm

c

1.0 μm

d

200 nm

Fig. 7.10: (a) Representative bio-transmission electron microscope images of native Escherichia coli (b) attachment of nanofibers with cells (red arrow and circle) (c) cleavage inside the cells by nanofibers (red arrows), (d) disintegration of E. coli (red arrows). Reprinted with permission from Amna et al. (2012). Copyright 2012, Springer.

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onstrated that exposure of the selected microbial strains to the nanofibers led to disruption of the cell membranes and leakage of the cytoplasm. In conclusion, the combined results suggested that doping promotes the antimicrobial effect; synthesized nanofibers possess a very large surface-to-volume ratio and may damage the structure of the bacterial cell membrane, also depressing the activity of the membranous enzymes which causes bacteria to die in due course. The nanocomposites of titania with copper oxide will be discussed in the next section of this chapter.

7.8.2 Doping of titania with copper The antimicrobial properties of both silver (Sondi I and Salopek-Sondi, 2004; Sheikh FA et al., 2010) and copper nanoparticles (Cioffi N et al., 2005) have been previously reported, and both have been coated onto or incorporated into various materials (Li Z et al., 2006). However, copper has attracted particular attention because it is the simplest member of the family of copper compounds and exhibits a range of potentially useful physical properties such as high temperature superconductivity, electron correlation effects and spin dynamics (Tranquada JM et al., 1995). As an important p-type semiconductor, CuO has found many diverse applications such as in gas sensors, catalysis, batteries, high temperature superconductors, solar energy conversion and field emission emitters. In the energy-saving area, energy transferring fluids filled with nano CuO particles can improve fluid viscosity and enhance thermal conductivity (Kwak and Kim, 2005). CuO crystal structures possess a narrow band gap, giving useful photocatalytic or photovoltaic properties as well as photoconductive functionalities (Xu JF et al., 1999). Moreover, CuO is cheaper than silver, easily mixed with polymers and relatively stable in terms of both chemical and physical properties. Highly ionic nanoparticulate metal oxides, such as CuO, may be particularly valuable antimicrobial agents as they can be prepared with extremely high surface area and unusual crystal morphologies (Gao F et al., 2009). However, limited information on the possible antimicrobial activity of nano CuO is available. Therefore, the development of nanostructured CuO/TiO2 rods with antimicrobial properties is of considerable interest. In this direction, Hassan et al. discussed the preparation of novel CuO/ TiO2 composite nanorods using precursor Cu(NO3)2.3H2O and titanium isopropoxide by sol-gel electrospinning technique (Hassan MS et al., 2013). In accordance with the crystal lattice parameters, CuO and TiO2 cannot combine in a single crystal. Accordingly, main TiO2 nanorods doped with CuO nanoparticles is the expected structure of the obtained product. To affirm this hypothesis, TEM analysis has been conducted (Fig. 7.11). As shown in Fig. 7.11a, tiny black particles distributed along the nanorods can be observed. The TEM image again confirms the diameter size of the CuO/TiO2 nanorods is 100 nm. As shown in the HR-TEM image (Fig. 7.11b), the black nanoparticles do have different crystal lattice parameters than the matrix as the crystal plane is closer, so one can claim that these nanoparticles represent CuO.

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a

50 nm

b

2 nm

Fig. 7.11: (a) TEM (b) HR-TEM image (green circles represent the CuO particle). Reprinted with permission from Hassan et al. (2013). Copyright 2013, Elsevier.

Nevertheless, although there are numerous publications on the bactericidal effect of TiO2 materials upon irradiation, to the best of the authors’ knowledge there has been no thorough report on the antibacterial effect and mechanism of CuO/TiO2 composite nanorods. Accordingly, Hassan et al. have investigated the mechanism of disinfection by electrospun CuO/TiO2 composite against common pathogenic bacteria under visible light. For this purpose, they chose Gram-positive S. aureus and Gram-negative E. coli strains supplemented with nanorod solution and an attempt was made to find the minimum inhibitory concentration of the nanorods capable of inhibiting the growth of the above mentioned pathogenic strains. The growth curves of E. coli and S. aureus treated with CuO/TiO2 composite nanorods and pristine TiO2 nanorods are shown in Fig. 7.12 (a, b) and insets of Fig. 7.12 (a, b) respectively, by measuring optical density at 600 nm. In the presence of 0, 5, 15, 25, 35 and 45 μg/ml of pristine TiO2 and CuO/TiO2 composite nanorods, the growth curves of testing strains included three phases: lag phase, exponential phase, and stabilization phase. However, decline phases in each growth curve could not be revealed because we only assayed the total numbers of bacteria, including live and dead ones, based on the value of OD600. In the absence of CuO/TiO2 composite nanorods, both the organisms reached exponential phase rapidly. But exposed to the above-mentioned concentrations of nanorods, E. coli cells and S. aureus were lagged to 4 h. With increasing concentration of synthesized nanorods, the delay was more evident. The observed MIC of the TiO2 and CuO/TiO2 composite nanorod solution was found to be different. The TiO2 nanorods slightly inhibited bacterial cell growth under light irradiation as the contact time increased. The nanosized TiO2 can penetrate the cell wall and partially hinder bacterial growth. In contrast, the CuO/TiO2 composite nanorods exhibit excellent antibacterial performance against both E. coli and S. aureus. In comparison to the pristine TiO2 (insets of Fig. 7.12a, b) nanorods, the CuO/TiO2 composite nanorods effectively killed both E. coli (Fig. 7.12a) and S. aureus (Fig. 7.12b). However, in the case of both microbial strains, it has been seen that with an increase in the concentration of TiO2 and CuO/TiO2 composite, the inhibition has also increased. A noticeable difference in the growth rate was observed for S. aureus and E. coli between 3–8 h of incubation. The highest concentration (45 μg/ml) of the CuO/TiO2 composite solution was found to exhibit dramatic toxicity against both the tested pathogenic strains.

7.8 Doping of titania with metal oxide   

2.0

1.4

1.0 OD600nm

1.6

0.8

5.0 μg/ml 15 μg/ml 25 μg/ml 35 μg/ml

0.6

45 μg/ml

0.4 0.2

1.2 OD600nm

0.0 μg/ml

0.0 μg/ml 5.0 μg/ml 15 μg/ml 25 μg/ml 35 μg/ml 45 μg/ml

1.2 1.8

   123

0.0

2 4 6 8 10 12 14 16 18 20 Time (h)

1.0 0.8 0.6 0.4 0.2 0.0 2

1.2

2.2

1.0

2.0 1.8

OD600nm

1.6 1.4

OD600nm

2.4

4

6

8

10 Time (h)

12

14

16

20

0.0 μg/ml

0.0 μg/ml 5.0 μg/ml 15 μg/ml 25 μg/ml 35 μg/ml 45 μg/ml

0.8

18

5.0 μg/ml 15 μg/ml 25 μg/ml 35 μg/ml

0.6

45 μg/ml

0.4 0.2 0.0

2

4

1.2

6

8 10 12 14 16 18 20 Time (h)

1.0 0.8 0.6 0.4 0.2 0.0 2

4

6

8

10 Time (h)

12

14

16

18

20

Fig. 7.12: Growth curves of (a) E. coli and (b) S. aureus exposed to different concentrations of CuO/ TiO2 composite nanorods (insets (a) and (b) show the growth curves of E. coli and S. aureus exposed to titania nanorods respectively). Reprinted with permission from Hassan et al. (2013). Copyright 2013, Elsevier.

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The nanocomposites of titania with nickel will be discussed in the next section of this chapter. Also the broad-spectrum of antimicrobial activity of nickel-doped titania nanofibers will be presented.

7.8.3 Doping of titania with nickel Nickel oxide is a material of interest for a variety of practical applications such as an electrochromic material for smart windows (Niklasson and Granqvist, 2007; Wu and Yang, 2007), gas sensing (Arshak KI et al., 2011), photovoltaics (Kim SS et al., 2006; He JJ et al., 1999), catalysis (Li J et al., 2008), light-emitting devices (Xi YY et al., 2008), and electrochemical capacitors (Zhao DD et al., 2008). Although many applications and properties of nickel oxide based materials have been investigated, there are however very few reports on their antibacterial activity. Nevertheless, Amna et al. inspected the mechanism of disinfection by electrospun NiO/TiO2 composite nanofibers against foodborne pathogens (Amna T et al., 2013). The utilized NiO/ TiO2 composite nanofibers were prepared by electrospinning of a sol-gel composed of nickel nitrate hexahydrate, titanium isopropoxide and poly(vinyl acetate). The obtained electrospun nanofiberous mat was vacuum dried at 80 °C and then calcined at 600 °C in air for 2 h. The FE-SEM images of calcined samples at different magnifications reveal good surface morphology made up of nanosize particles (Fig. 7.13). The diameter size of the composite nanofibers is in the range of 70–90 nm. The nanosize building block morphology as depicted by the SEM images is beneficial for numerous applications. The EDX spectra (Fig. 7.13d) of the NiO/TiO2 composite nanofibers contain Ti, Ni and O; no other element impurity is detected, indicating that the final product is pure and free of impurity and composed of NiO and TiO2. Fig. 7.14 shows the TEM and high resolution (HR)-TEM images along with selected area electron diffraction (SAED) and fast Fourier transform (FFT) micrographs of the NiO/TiO2 composite nanofibers obtained after firing (or annealing) at 600°C. The TEM image (Fig. 7.14a) further confirms the fiber diameter in the range of 70–90 nm with nanosize building block particles. The SAED pattern (inset Fig. 7.14a) shows a set of bright rings, each ring indicating a different lattice plane of the anatase TiO2. Furthermore, HR-TEM (Fig. 7.14b) shows the parallel crystalline planes, which are confirmed by FFT micrograph (inset Fig. 7.14b). Tiny black particles distributed along the nanofibers can be observed in TEM which are more apparent in HR-TEM. As depicted in the HR-TEM, the black nanoparticles do have a different crystal lattice to the matrix, so one can claim that these nanoparticles represent NiO. The antibacterial activity was tested against four common foodborne pathogenic bacteria viz., S. aureus, E. coli, S. typhimurium and K. pneumoniae by MIC method taking five different concentrations (5–45 μg/ml). In the present study the lowest concentration of NiO/TiO2 composite solution inhibiting the growth of tested strains was found to be 5 μg/ml. To sum up, novel NiO/ TiO2 composite nanofibers which possess large surface-to-volume ratio with excel-

7.8 Doping of titania with metal oxide   

   125

lent antimicrobial activity were fabricated, which can be used to inhibit the microbial growth associated with food stuffs. Furthermore in continuance, the fabrication and antibacterial potential of nanocomposites of titania with cobalt will be summarized in the next section of the chapter.

a

b

c

d

Fig. 7.13: FESEM image of the NiO/TiO2 composite nanofibers at different magnifications (a, b, c) and (d) EDX spectra. Reprinted with permission from Amna et al. (2013). Copyright 2013, Springer.

a

50 nm

b

2 nm

Fig. 7.14: TEM image of the prepared NiO/TiO2 composite nanofibers at (a) low resolution and (b) HR-TEM image, the blue circle refers to a NiO nanoparticle. The inset in (a) shows the SAED pattern and inset in (b) shows the FFT micrograph. Reprinted with permission from Amna et al. (2013). Copyright 2013, Springer.

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7.8.4 Doping of titania with cobalt Cobalt oxides (CoO and Co3O4) materials possess remarkable optical, electrical and magnetic properties and are therefore commonly used for photocatalysis and electromagnetic applications (Wang RM et al., 2004; Li T et al., 2004). Despite a variety of interesting properties as aforementioned, cobalt oxides also possess antibacterial potential. Amna et al. explored the electrospinning technique to successfully fabricate Co-doped TiO2 and the manufactured nanofibers were investigated for antibacterial effect (Amna T et al., 2013). Specifically, this study examined the effects of Co-doped TiO2 concentration and exposure time on the growth and viability of the bacteria. The fabrication of Co-doped TiO2 nanofibers was carried out by facile electrospinning. Briefly, Poly(vinyl acetate) (18 wt%) solution was prepared by dissolving PVAc in N, N-dimethylformamide under magnetic stirring at room temperature. 5 g of titanium isopropoxide was taken in a separate bottle and a few drops of acetic acid were added to it until the solution became transparent. The acetic acid was added to avoid precipitation of the solution. 5 wt% of alcoholic solution (ethyl alcohol) of Co(NO3)3. 6H2O was added into the TIP solution. Finally the obtained sol-gel was subjected to electrospinning. Fig. 7.15 (a, b) shows SEM images of as-synthesized nanofibers at different magnifications. The nanofibers have a continuous, uniform and smooth surface with an average diameter of 350–450 nm. The presence of cobalt oxide in the doped nanofibers was further confirmed by EDX analysis. The EDX spectrum (Fig. 7.15c) contains Ti, Co and O; no other element impurity is detected, indicating the final product is free of impurity and composed of cobalt oxide and TiO2 only. Fig. 7.16 shows the TEM image

a

10 μm

b

1 μm c

Fig. 7.15: (a–b) SEM images at different magnifications and (c) EDX spectrum of the electrospun Co-doped TiO2 nanofibers. Reprinted with permission from Amna et al. (2013). Copyright 2013, Elsevier.

   127

7.8 Doping of titania with metal oxide   

along with an SAED pattern and FFT micrograph of the doped nanofibers. The TEM image (Fig. 7.16a) further confirmed the fiber diameter ~400 nm. The dark black particles on the surface of the nanofiber may be due to the presence of cobalt oxide in the composite nanofiber. The FFT micrograph shows the parallel crystalline planes which confirm the high crystallinity of the sample (Fig. 7.16b). The SAED pattern is composed of some bright spots which supports the polycrystalline nature of electrospun nanofibers (inset Fig. 7.16b).

a

b

2 nm

200 nm

Fig. 7.16: TEM image (a) FFT micrograph (b) of the electrospun Co-doped TiO2 nanofibers, the inset represents the SAED pattern. Reprinted with permission from Amna et al. (2013). Copyright 2013, Elsevier.

1.4

0.0 μg/ml

1.2

a

0.0 μg/ml

1.2

Pristine TiO2 10 μg/ml

20 μg/ml

OD600nm

OD600nm

1.0

Pristine TiO2 10 μg/ml

1.0

b

5.0 μg/ml

5.0 μg/ml

0.8 0.6

0.8

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0.6 0.4

0.4

0.2

0.2

0.0

0.0 0

4

8 12 Time (h)

16

20

0

4

8 12 Time (h)

16

20

Fig. 7.17: Growth curves of (a) S. aureus and (b) S. typhimurium exposed to different concentrations of Co-doped TiO2 nanofibers. Data are the average from triplicate experiments. Reprinted with permission from Amna et al. (2013). Copyright 2013, Elsevier.

To examine the antibacterial effect, the double dilution method was used. Different concentrations (5, 10 and 20 μg/ml) of Co-doped TiO2 nanofibers (Fig. 7.17a, b) were used in the present study. The obtained results indicate that the inhibition is both a time and concentration-dependent matter. Moreover the doping of Co in TiO2 promotes the bactericidal effect against two common foodborne pathogenic microorganisms (Staphylococcus aureus and Salmonella typhimurium). The Co ions interacted with the negatively charged bacterial cells and adhered to the bacterial cell walls.

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The bacteria and the Co-doped TiO2 developed electrostatic forces and we consider this electrostatic attraction may be the reason for their adhesion and bioactivity. The current work suggests that the enhanced antibacterial effect of the nanofibers might be explained by a synergistic effect of the components. Moreover, the synthesis of Cerium-doped titania via electrospinning process will be discussed at the end of this section. Additionally, the antimicrobial potential of synthesized Co-doped TiO2 nanofibers has also been discussed in a nutshell in this chapter.

7.8.5 Doping of titania with cerium CeO2 can exist as both Ce(III) and Ce(IV) forms (Robinson RD et al., 2002) and the ratio of the two oxidation states appears to be size dependent (Zhang F et al., 2002), with increasing Ce(III) at the lower size. In recent years, nanocrystalline cerium oxide particles have attracted great attention and been extensively studied. Cerium oxide has numerous potential applications such as UV-blockers and filters (Tsunekawa S et al., 2000; Yamashita M et al., 2002), gas sensors (Garzon FH et al., 2000), high refractive index material (Mogensen M et al., 2000) as well as luminescent material (Yu XJ et al., 2001). Cerium oxide also has applications as catalysts in fuel cell technology (Logothetidis S et al., 2003), catalytic wet oxidation (Larachi F et al., 2002), engine exhaust catalysts (Dario and Bachiorrini, 1999), and photocatalytic oxidation of water (Bamwenda and Arakawa, 2000). Moreover, cerium oxide has been used as a polishing agent for glass mirrors, plate glass, television tubes, ophthalmic lenses, and precision optics. Recently, it has also been demonstrated that cerium oxide nanoparticles exhibit antioxidant activity at physiological pH and therefore may be useful in biomedical applications for protecting cells against radiation damage, oxidative stress, or inflammation (Perez JM et al., 2008; Tarnuzzer RW et al., 2005). Generally speaking, relatively few studies have focused on the effects of cerium oxide and other metal oxide nanoparticles on bacterial systems (Ju-Nam and Lead, 2008). Thill et al. pointed out the toxicity of commercial cerium oxide nanoparticles to Escherichia coli and showed that the nanoparticles adsorbed to the bacterial cell surface (Thill et al., 2006). On the other hand, Hassan et al. have utilized the electrospinning technique to successfully fabricate Ce(III) oxide–titania composite nanofibers (Hassan MS et al., 2012). Fig. 7.18 shows the XRD pattern of the electrospun nanofibers before and after calcination. The XRD of the dried electrospun nanofibrous mat prepared from cerium nitrate/TIP/PVAc (Fig. 7.18a) showed no distinguished peak due to the amorphous nature of the sample. All the reflections after firing (or annealing) at 600 °C (Fig. 7.18b) could be indexed based on the anatase titania phase (PDF-89-4921). The peak at 30.6° is showing the presence of Ce2O3 (PCPDF-78-0484) in the sample. From SEM images of precursor nanofibers (Fig. 7.19a), it can be seen that the prepared Ce (NO3)3/TIP/PVAc composite fibers have a uniform and smooth surface with diameter ~ 500 nm. However, when the

7.8 Doping of titania with metal oxide   

   129

Intensity (a. u.)

T -TiO2 C -Ce2O3

2θ (degree) Fig. 7.18: XRD pattern of the electrospun nanofibers (a) without calcinations and (b) with calcinations at 600 °C. Reprinted with permission from Hassan et al. (2012). Copyright 2012, Elsevier.

a

b

5 μm

2 μm

d

c

500 nm Fig. 7.19: (a–c) FESEM image of the Ce2O3/TiO2 composite nanofibers at different magnifications and (d) EDX spectra. Reprinted with permission from Hassan et al. (2012). Copyright 2012, Elsevier.

precursor fibers were heated to 600 °C for 2 h in air atmosphere, the PVAc polymer was selectively decomposed but the nanofibers remained as an unbroken structure with reduced diameter size (~ 300 nm) (Fig. 7.19b). The decrease in nanofiber diameter after calcination is due to the loss of PVAc from the precursor fiber and the crystallization of the oxides. The SEM image of the calcined sample at high magnifications showed the fascinating porous morphology on the surface which is beneficial for numerous applications (Fig. 7.19c). Presence of cerium oxide in the composite fibers was further confirmed by EDX analysis. The EDX spectrum (Fig. 7.19d) of the composite nanofibers contains Ti, Ce and O; no other element impurity is detected, indicating the

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final product is pure and free of impurity and composed of cerium oxide and titania. The presence of Ce particles in titania was again confirmed by the EPMA image of composite nanofibers (Fig. 7.20). The EPMA image clearly shows that Ti is the main element, and Ce is also uniformly dispersed on the surface of the titania nanofibers. Fig. 7.21 shows TEM and high resolution (HR)-TEM images along with the selected area electron diffraction (SAED) pattern of the composite nanofibers obtained after firing (or annealing) at 600 °C. TEM images at low and high magnifications (Fig. 7.21a and b) further confirmed the fiber diameter in the range of about ~ 300 nm with unique porous morphology. In HR-TEM (Fig. 7.21c), nanoparticles showed two different crystal lattices in the matrix which confirms the presence of cerium oxide in titania. The SAED pattern (Fig. 7.21d) is composed of some bright spots and rings which support the polycrystalline nature of the as-spun composite nanofibers.

SE,255

a

TiKa,2

b

CeLb1,1

c

Fig. 7.20: EPMA mapping result of Ce2O3/TiO2 composite nanofibers, blue color square represents the selected area. Reprinted with permission from Hassan et al. (2012). Copyright 2012, Elsevier.

a

200 nm

c

b

100 nm

d

5 nm Fig. 7.21: TEM image of the prepared Ce2O3/TiO2 composite nanofibers (a, b) at different resolutions (c) HR-TEM image and (d) the SAED pattern. Reprinted with permission from Hassan et al. (2012). Copyright 2012, Elsevier.

7.9 Plausible antibacterial mechanism of TiO2 / doped-TiO2 nanostructures    

   131

Specifically, this study scrutinized the effects of synthesized Ce2O3–TiO2 nanofibers’ concentration and exposure time on the growth and viability of S. aureus and S. typhimurium. From this study it was given to understand that Ce2O3–TiO2 nanofibers are toxic to food pathogens. This study also demonstrated the fabrication of novel Ce2O3–TiO2 composite nanofibers with increased surface-to-volume ratio using a simple electrospinning process and offered supportive evidence to indicate that Ce2O3–TiO2 nanofibers can inhibit S. aureus and S. typhimurium growth and even kill the cells through destroying the bacterial membranous configuration (Fig. 7.22).

a

b

c

1 μm

d

0,5 μm

Fig. 7.22: Growth curves of (a) S. aureus and (b) S. typhimurium exposed to different concentrations of Ce2O3/TiO2 composite nanofibers. Data are average from duplicate experiments. Representative transmission electron microscope images of (c) S. aureus (inset native S. aureus) and (d) S. typhimurium (inset native S. typhimurium) with composite Ce2O3/TiO2 nanofibers. Reprinted with permission from Hassan et al. (2012). Copyright 2012, Elsevier.

7.9 Plausible antibacterial mechanism of TiO2 / doped-TiO2 nanostructures Oxidative stress induced by reactive oxygen species (ROS) is one of the most important antibacterial mechanisms of engineered nanoparticles. It has already been established that most of the semiconductor on band gap illumination produces reactive oxygen species like HO•−, O2•−−, HO2•− and H2O2, (Hu et al. 2010; Gaya and Abdullah

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2008) which may react with bacteria. In the presence of moisture, O2•−− generates reactive oxidizing species like HO•−, HO2•− and H2O2. Based on the antibacterial results discussed in the previous section, the probable antibacterial mechanism has been suggested as follows. Perhaps cell lysis may be comprised of two stages: an initial fast and passive process involving physical adsorption or ion exchange at cell surfaces, and a slower transport of metal ions into bacterial cells as represented schematically in Fig. 7.23. More precisely, when the pristine titania nanorods solution was irradiated under UV light, the production of reactive oxygen species (ROS) by aqueous suspensions occurred. UV light induces a separation of charge, generating a hole (h+) in the valence band and an electron in the conduction band. At the surface of the excited particle, the valence band holes abstract electrons from water and/or hydroxyl ions, generating hydroxyl radicals (OH•). Electrons reduce O2 to produce the superoxide anion O2−•. ROS generated by UV-irradiated TiO2 nanorods might have interacted with the outer wall of the cell in the beginning and, further, the generated free radicals entered into the inner wall of the cell leading to the disruption of the internal contents of the cell and, as a result, the cells were deformed leading to disorganization and leakage (Hassan MS et al., 2012). In contrast, the doped titania also depicted excellent antimicrobial activity in visible light (Amna T et al., 2013; Hassan MS et al., 2013). In the case of doped TiO2 nanostructures, UV irradiation was not found to be essential.

Interaction with bacteria

O2–* S RO – * hʋ O2 RO OH S *

O2

eCB–

UV Light Separation of charges

hVB+ H2O or OH–

OH*

Titania nanorods in broth

Plasma membrane Peptidoglycan Periplasmie space Nucleoid

Gram Positive Bacteria

Rupture

Peptidoglycan Plasma membrane Nucleoid Periplasmie space Outer membrane (Lipopolysaecharide Rupture and Protein) Gram Negative Bacteria

Disintegration and death

Disintegration and death

Fig. 7.23: Possible acting mechanism of fabricated nanorods on tested pathogenic strains. Reprinted with permission from Hassan et al. (2012). Copyright 2012, American Scientific Publishers

7.10 Concluding remarks   

   133

In conclusion, it was demonstrated that the exposure of the microbial strains to various nanomaterials led to the disruption of the cell membranes and leakage of the cytoplasm which causes bacteria to die eventually. The results pointed out the oxidative attack from the exterior to the interior of the bacteria by hydroxyl radicals as the primary mechanism of photocatalytic inactivation.

7.10 Concluding remarks The investigation of interactions between emerging nanostructures and biological systems is of utmost importance in order to develop safe nanotechnology for biomedical applications. It is a well-known saying that it is better to be safe than sorry. Thus inorganic nanoparticles/metal oxide-based nanotextured materials manufactured for biomedical applications should be evaluated before they are utilized. Titanium is widely used as a material for permanent implants in orthopedic and dental applications. It is well known that Ti shows a mechanically stable interface towards bone (osseointegration). The good biological properties are due to the beneficial properties of the native oxide (TiO2) that forms on Ti when exposed to oxygen. Properties of nanocrystals are strongly influenced by their size. Nanomaterials exhibit novel properties owing to their unique geometry. They can be used for a wide variety of applications. The use of metal oxides as antimicrobial agents has the advantage of improved safety and stability, as compared to organic antimicrobial agents. The bactericidal effect of titania, silver, copper, and zinc oxide nanoparticles on the microorganisms is very well known. This chapter focused on the fabrication, characterization and antibacterial potential of one-dimensional pristine TiO2 and metal oxide-doped TiO2 (in particular zinc oxide, cobalt oxide, cerium oxide, nickel oxide and copper oxide) nanofibers by facile sol-gel electrospinning technique. The physicochemical properties of the synthesized nanomaterials were determined by various sophisticated techniques such as X-ray diffraction pattern (XRD), field emission scanning electron microscopy (FESEM), energy-dispersive X-ray spectroscopy (EDX), electron probe microanalysis (EPMA) and transmission electron microscopy (TEM). Additionally the antibacterial potential of metal oxides and metal oxide-doped TiO2 has been reviewed. In this regard the microorganisms which were mostly tested and discussed include S. aureus, E. coli, S. typhimurium and K. pneumonia. The antibacterial potential of the aforementioned nanotextured materials was evaluated by effective MIC method. Moreover, the various fabrication methods were discussed briefly, however the electrospinning technique was discussed somewhat in detail. Furthermore, the concept of doping and potential effect of doping in enhancing the antibacterial activity has been revealed. Finally, this chapter sheds light on the applications of metal oxide, titania, metal oxide-doped titania as potential antibacterial agents as well as proposes new directions of research

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on interactions between emerging nanomaterials and biological systems to develop safe nanotechnology for biomedical applications.

Acknowledgment This chapter was supported by research funds of Chonbuk National University.

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Luisa Russo, Sabrina Zaccaria, Maria Assunta Autiello and Assunta Borzacchiello

8 Hydrogels for biomedical applications 8.1 Hydrogels: Classification and basic structure Hydrogel materials consisting of water-swollen hydrophilic polymer networks exhibit a large number of specific properties highly attractive for a wide range of applications (Fig. 8.1). In particular their ability to contain water, stability in aqueous media and softness make hydrogels compatible with biological systems and thus means they are suitable for different biomedical applications such as tissue engineering, wound healing, controlled drug release, contact lenses, to name a few [1, 2]. Hydrogels can be classified on the basis of different properties such as method of synthesis, ionic charges and physical structure [3]. As regards the first type of classification, the following properties can be considered: homo-polymer hydrogels or networks (they are composed of one type of hydrophilic monomer), copolymer hydrogels (if made of two types of monomers, at least one hydrophilic), multi-polymer hydrogels

MC

Fig. 8.1: Schematic representation of the hydrogel structure.

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(they are composed of more than three types of monomers) or interpenetrating polymeric networks if they are prepared by swelling a network of polmer1 in monomer2, to make an intermeshing network of polymer1 and polymer2. Also the influence of ionic charges is important to obtain different hydrogels such as neutral hydrogels, anionic hydrogels, cationic hydrogels, or ampholytic hydrogels when both charges are present on the same hydrogel. Moreover, hydrogels can be classified on the basis of their physical structure. Their chains in fact can be randomly arranged to form amorphous hydrogels, present dense regions of ordered macromolecules, i.e. crystallites to form semi-crystalline hydrogels or also to form hydrogen-bonded hydrogels structures, supermolecular networks and hydrocolloidal aggregates [4]. Another important classification for hydrogels is done on the basis of the nature of their crosslink. They can be classified as physical and chemical gels; in particular the first are also called ‘reversible’ because the networks are held together by molecular entanglements, and/or secondary forces including ionic, H-bonding or hydrophobic forces while the chemical ones are known as ‘permanent’ because they are covalently-crosslinked networks. There are several crosslinking methods to obtain both chemical and physical gels. In the first case they are realized by radical polymerization, high energy irradiation, chemical reaction with a complementary group or using enzymes. In the second case hydrogels can be physically crosslinked by ion interactions, from amphiphilic block and graft copolymers or by crystallization. Both types of gels do not present a homogenous aspect; they usually contain regions of low water swelling and high crosslink density, called ‘clusters’, that are dispersed within regions of high swelling, and low crosslink density. This may be due to hydrophobic aggregation of crosslinking agents, leading to high crosslink density clusters. Depending on the solvent composition, temperature and solids concentration during gel formation, hydrogel structure can show a phase separation and water-filled ‘voids’ or ‘macropores’ can form. In chemical gels, free chain ends represent gel network ‘defects’ which do not contribute to the elasticity of the network. Other network defects are chain ‘loops’, which also do not contribute to the permanent network elasticity. Furthermore, the structure of hydrogel can be influenced by the external environment such as pH, ionic strength, temperature, electromagnetic radiation, etc. showing a swelling behavior dependent on physiological environmental stimuli [5]. Polimeric materials commonly used in the realization of hydrogels for biomedical applications are natural, synthetic or a combination of these [6]. The choice of material is important to obtain hydrogels with different properties for specific applications. For example, natural polymers offer advantageous properties such as nontoxicity and biocompatibility whereas by using synthetic material it can be possible to obtain hydrogels with a well-defined degradation kinetic and better mechanical properties.

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   143

8.1.1 In situ forming hydrogels For some applications it is advantageous if the hydrogels can be formed in situ (i.e. in vivo). Recently an increasing number of in situ-forming systems have been developed for various biomedical applications, including tissue repair, drug delivery, and cell encapsulation. These systems consist of injectable fluids that can be introduced into the body in a minimally invasive manner before gelling within the desired tissue, organ, or body cavity. These systems present numerous advantages, the main one being that they do not need a surgical procedure for placement. Both physical and chemical crosslinking strategies have been pursued to achieve in situ gelation.

Physical crosslinking methods A variety of environmental triggers, such as pH, temperature or ionic strength, can be used to achieve physical crosslinking of polymer chains. Different physicochemical interactions, such as hydrophobic or electrostatic interactions, hydrogen bonding or supramolecular chemistry, can be responsible for the formation of the gel. Regarding the hydrophobic interactions, polymers with hydrophobic domains can crosslink in aqueous media through reverse thermal gelation (sol-gel chemistry). In this kind of polymers, usually referred to as gelators, a hydrophobic segment is attached to a hydrophilic polymer segment by post-polymerization grafting or by directly synthesizing a block copolymer to create a polymer amphiphile. Such amphiphiles are water soluble at low temperature, but with increasing temperature, hydrophobic domains aggregate to minimize the hydrophobic surface area contacting the bulk water, reducing the amount of structured water surrounding the hydrophobic domains and maximizing the solvent entropy. The gelation process will occur at a temperature which depends on several factors: the concentration of the polymer, the length of the hydrophobic block, and the chemical structure of the polymer. The more hydrophobic the segment, the larger the entropic cost of water structuring, the larger the driving force for hydrophobic aggregation, and the lower the gelation temperature. The chemical structures of some common hydrophobic blocks are shown in Fig.  8.2. They can undergo reverse thermal gelation at a temperature close to the physiological one. The most widely-used reverse thermal gelation polymers are the poloxamers (or Pluronics), triblock copolymers of poly(ethylene oxide)–poly(propylene oxide)– poly(ethylene oxide) (PEO–PPO–PEO) [7]. Carbohydrate (i.e. hyaluronic acid) grafting to poloxamers reduces the critical gelation concentration and the in vivo dissolution rate of the networks because of the high viscosity of the carbohydrate grafts. Poloxamers can also be modified by adding an additional polymer block at each chain terminus, forming an ABCBA pentablock copolymer with improved properties for some biomedical applications. For example, pentablock PDMAEMA–PEO–PPO–PEO–

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   8 Hydrogels for biomedical applications

O

O *

n*

O n

m

* CH3

Poly(propylene oxide) PPO

n

*

O

CH3

*

O

Poly(lactide-co-glycolic acid) PLGA

O

HN H3 C

*

CH3

Poly(N-isopropylacrylamide) PNIPAM

Fig. 8.2: Chemical structures and abbreviations of common thermogelling hydrophobic blocks.

PDMAEMA forms free-flowing liquids at room temperature but forms elastic hydrogels at concentrations above 12% when heated [8]. Among the thermoresponsive polymers poly(N-isopropylacrylamide) (PNIPAM) is one of the most widely studied. PNIPAM–poly(phosphorylcholine)–PNIPAM triblock copolymers have been reported which form gels at 6–7 wt% when the phase transition temperature of the PNIPAM arms (~ 32 °C) is exceeded [9]. Grafting PNIPAM linear chains onto natural polymers can also convert those polymers into physically crosslinkable hydrogels. Moreover, several natural polymers also undergo reverse thermal gelation. For example, chitosan solutions containing glycerol-2-phosphate [10] or chitosan grafted with 40 wt% PEG, gel at a temperature close to 37 °C. Physical and pharmacokinetic properties of physically crosslinked gels can be tuned by changing the morphology of the gelator. For example, star diblock copolymers can be produced from multifunctional controlled radical polymerization initiators to form highly efficient thermally gelling polymers. Star diblocks are typically prepared by first polymerizing a water-soluble polymer in the center of the star (e.g. phosphorylcholines) followed by a thermosensitive (e.g. poly(propylene oxide) methacrylate) or pH-sensitive (e.g. 2-diisopropylaminoethyl methacrylate) gelator in the outer arms [11]. The advantage of such polymer architectures is that lower polymer concentrations are required to realize strongly gelled systems. Another way to significantly lower the polymer fraction required to form gels is to formulate thermogelling polymers as particles due to long-range electrostatic double layer interactions between the surfaces of the particle. For example, biocompatible interpenetrating polymer network microgels comprising poly(N-isopropylacrylamide) and polyacrylic acid undergo reversible gelation to form particle assemblies at 33 °C at weight concentrations above 2.5 wt%. This critical concentration is approximately one order of magnitude lower than that observed for conventional poloxamer systems [12]. Electrostatic interactions have been widely investigated for crosslinking in situ gelling polymers. One advantage of this approach is that crosslinking (or de-crosslinking) can be triggered by pH changes which ionize or protonate the ionic functional groups that cause gelation. Electrostatic interactions may occur between a polymer and a small molecule or between two polymers of opposite charge to form a

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hydrogel (Fig. 8.3). For example elastin-like polypeptides have been crosslinked via electrostatic interactions between their cationic lysine residues and anionic organophosphorus crosslinkers under physiological conditions [13].

add polymer of opposite charge

add small molecule cross-linker of opposite charge

Fig. 8.3: In situ physical gelation through electrostatic interactions with an oppositely-charged polymer or an oppositely-charged small-molecule crosslinker.

As examples of polymer–polymer crosslinking, ionic-complementary peptides with alternating positive and negative charge domains can self-assemble to form hydrogels in situ [14]. Electrostatic interactions can also be used to crosslink microparticle or nanoparticle gels. In this way it is possible to create three-dimensional particle assemblies as in the case of dextran microspheres coated with anionic and cationic polymers which exhibit spontaneous gelation upon mixing, due to ionic complex formation between the oppositely charged microparticles [15]. Regarding hydrogen bonding interactions, they can be used to formulate injectable hydrogels. Mixtures of two or more natural polymers can display rheological synergism, meaning that the viscoelastic properties of the polymer blends are more gel-like than those of the constituent polymers measured individually [16–17]. This is due to hydrogen bonding interactions between the polymer chains, in turn facilitated by the compatible geometries of the interacting polymers. Blends of natural polymers, such as gelatin–agar [18], starch–carboxymethyl cellulose [19], and hyaluronic acid–methylcellulose [20], form injectable physically crosslinked gel-like structures which exhibit

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excellent biocompatibility due to the absence of chemical crosslinkers and to the similarity to extracellular matrix polymers. However, in vivo hydrogen-bonded networks can dilute and disperse over a few hours due to influx of water, restricting their use to relatively short-acting drug release systems unless some other form of crosslinking is used. A latest approach to forming hydrogels in situ involves using specific molecular recognition motifs and/or supramolecular chemistry. The most common type of crosslinking interaction of this type is the formation of inclusion complexes between poly(alkylene oxide) polymers and cyclodextrins [21–23]. Naturally occurring macromolecules can also assemble into hydrogels in situ. For example, interactions between glycosaminoglycans and polymer-grafted peptide sequences can be used to rapidly form hydrogels (e.g. heparin and heparin-binding peptide) [24]. Synthetic thermally-associating polypeptides can also crosslink through hydrophobic domain interactions to yield hydrogels with high mechanical strength. In this case, assembly is controlled not only by the amphiphilic amino acid sequences but also by the chain conformations of the assembling polypeptides [25].

Covalent crosslinking strategies for forming hydrogels in situ One of the major advantages of the physically crosslinked hydrogels is that they do not need chemical modification or the addition of crosslinking agents for the in situ gelification. However, they also have some limitations because the strength of a physically crosslinked hydrogel is directly related to the chemical properties of the constituent gelators thus restricting the design flexibility of such hydrogels. Additionally, the tissue residence time of physically bonded hydrogels is often poor due to dilution. In contrast, covalent crosslinking prevents both dilution of the hydrogel matrix and diffusion of the polymer away from the site of injection. The most common synthetic strategies for in situ crosslinking hydrogels are based on the introduction of a crosslinker that can be added to reactive pre-polymers as small molecules or conjugated directly to them. As an example for small-molecule crosslinkers, hydrogen peroxide has been used to produce in situ crosslinked hydrogels as dextran–tyramine [26] and hyaluronic acid–tyramine [27]. These hydrogels have a controllable gelation time ranging from 5 s to 9 min according to the reactant concentrations used. The main drawback of the small-molecule crosslinking method is the potential toxicity of residual unreacted small-molecule crosslinkers. For example, glutaraldehyde, which is often used to form carbohydrate-based hydrogels [28], is also toxic because it is a tissue fixative. To overcome this problem it is possible to use polymers pre-functionalized with reactive functional groups as crosslinking agents. In this case, the main disadvantage is that significant polymer modification chemistry may be required to prepare the functionalized pre-polymers. Moreover, the pre-gel polymers are often themselves

8.2 Structure-properties relationship   

   147

cytotoxic, even when prepared from highly biocompatible polymer precursors, and can also form potentially tissue-reactive oligomers during the polymer degradation. Several types of linkages can be made depending on the desired speed of crosslinking and biodegradability of the resulting conjugates. An example is the formation of a hydrazone bond, through the reaction of an aldehyde and a hydrazide, that facilitates fast crosslinking of gel precursors [29]. Another widely investigated in situ crosslinking synthetic strategy is the Michael addition between a nucleophile (i.e. an amine or a thiol) and a vinyl group. This is particularly useful due to its rapid reaction time, its flexibility in forming multiple types of bonds, and the relative biological inertness of the polymeric precursors. As an example of this chemistry, a mixture of thiol-modified heparin and thiol-modified hyaluronic acid can be gelled with PEG diacrylate to form a hydrogel which can prolong the release of basic fibroblast growth factor in vivo [30].

8.2 Structure-properties relationship The structure and properties of a specific hydrogel are extremely important in selecting which materials are suitable for the specific application. In this section, the theory describing the mechanical properties and the swelling behavior of hydrogels is shown and the structure-properties relationship of hydrogels is discussed.

8.2.1 Hydrogel mechanical properties Hydrogels’ time dependent properties The dynamical mechanical analysis allows to obtain information on the viscoelastic properties of the hydrogels by measuring the response of a sample when it is deformed under periodic oscillation [31]. In a dynamic experiment the material is subjected to a sinusoidal shear strain (or stress): 𝛶 = 𝛶0 sin(𝜔𝑡)

(8.1)

where ϒ0 is the shear strain amplitude, ω is the oscillation frequency (also expressed as 2πf where f is the frequency in Hz) and t the time. The mechanical response, expressed as shear stress τ of viscoelastic materials, is intermediate between an ideal pure elastic solid (obeying Hooke’s law) and an ideal pure viscous fluid (obeying Newton’s law) and therefore is out of phase with respect to the imposed deformation as expressed by: 𝜏 = 𝐺 ∗ (𝜔) 𝛶0 sin(𝜔𝑡 + 𝛿)

(8.2)

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   8 Hydrogels for biomedical applications

From (2) it is possible to obtain the shear stress τ expressed as the sum of two contributions: 𝜏 = 𝐺 ∗ (𝜔) 𝛶0 sin(𝜔𝑡) cos(𝛿) + 𝐺 ∗ (𝜔) 𝛶0 cos(𝜔𝑡) sin(𝛿)

(8.3)

and if it is defined 𝐺󸀠 (𝜔) = 𝐺 ∗ cos(𝛿) 𝐺󸀠󸀠 (𝜔) = 𝐺 ∗ sin(𝛿)

(8.4)

It is obtained: 𝜏 = 𝐺󸀠 (𝜔) 𝛶0 sin(𝜔𝑡) + 𝐺󸀠󸀠 (𝜔) 𝛶0 cos(𝜔𝑡)

(8.5)

where G′(ω) is the shear storage (or elastic) modulus and G″(ω) is the shear loss (viscous) modulus. G′ gives information about the elasticity or the energy stored in the material during deformation, whereas G″ describes the viscous character or the energy dissipated as heat. The dependences of the elastic and viscous moduli upon frequency are usually defined as mechanical spectra. The combined viscous and elastic behavior is given by the absolute value of complex shear modulus G󵠥: 𝐺 ∗ (𝜔) = √𝐺󸀠2 + 𝐺󸀠󸀠2

(8.6)

or by the absolute value of complex viscosity η󵠥 defined as: 𝜂 ∗ (𝜔) =

√𝐺󸀠2 + 𝐺󸀠󸀠2 𝜔

(8.7)

which is usually compared with the steady shear viscosity in order to evaluate the effect of large deformations and shear rates on the material structure. The ratio between the viscous modulus and the elastic modulus is expressed by the loss tangent: tan 𝛿 =

𝐺󸀠󸀠 𝐺󸀠

(8.8)

where δ is the phase angle. The loss tangent is a measure of the ratio of energy lost to energy stored in the cyclic deformation [31]. The phase angle, δ, is equal to 90° for a purely viscous material, 0° for a pure elastic material, and 0°  G″), as reported in Fig. 8.4a. The limit between the two regions is represented by the crossover frequency, usually expressed as ωc. These observations can be explained by considering the following. At low frequency, the molecular chains can release stress by disentanglement and molecular rearrangement during the period of oscillation, and hence, the solution shows viscous behavior (G″ > G′). At high frequency, however, molecular chains cannot disentangle during this short period of oscillation, and therefore, they behave as a temporary crosslinked network, and the elastic behavior (G′ > G″) is prevalent [33–35]. As regards, indeed, the mechanical spectra of the gels the elastic modulus curve is higher than the viscous modulus one over the frequency range analyzed [36]. In particular, the strong gel spectrum (Fig. 8.4b) is characterized by G′ and G″ curves that are almost frequency independent; moreover G′ values are typically 1–2 orders of magnitude greater than G″ values (tan δ lower than 0.1) [37–42]. In the case of a weak gel, instead, G′ and G″ moduli (Fig. 8.4c) show a slight frequency dependence and the ratio G″/G′ is higher than 0.1 [43]. Such rheological behavior is typical of a biological gel such as protein and polysaccharides based network (collagen, hyaluronic acid) and soft tissues [44–45].

Stress strain behavior Most hydrogels in their swollen state can be considered as rubber, that is, lightly crosslinked networks with a rather large free volume that allows them to respond to external stresses with a rapid rearrangement of the polymer segments. When a hydrogel is in the region of rubber-like behavior, its mechanical behavior is dependent mainly on the architecture of the polymer network [46]. To derive the relationship between the network characteristic and the mechanical stress-strain behavior, clas-

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sical and statistical thermodynamics and phenomenological approaches have been used to develop an equation of state for rubber elasticity. Combining those theories, a constitutive equation expressing the modulus as a function of structural features such as molecular weight between crosslinking points can be expressed as: 𝐺=

2 𝜌𝑅𝑇 𝑟0

𝑀𝐶 𝑟2𝑓

(1 −

2𝑀𝐶 𝑀𝑛

)

(8.9)

where r is the density of the polymer, Mc is the number average molecular weight between crosslinks, r̅02/r̅f2, is the ratio of the end-to-end distance in a real network versus the end-to-end distance of the isolated chains; it is generally approximated as 1 when it is unknown.

8.2.2 Hydrogel swelling The swelling behavior of biomedical hydrogels in biological fluids can be described by a variety of theoretical models aiming at the prediction of the swelling behavior. One theory that can be used with reasonable success is the Flory-Rehner analysis [47]. This thermodynamic theory states that a crosslinked polymer gel, which is immersed in a fluid and allowed to reach equilibrium with its surroundings, is subject only to two opposing forces, the thermodynamic force of mixing and the retractive force of the polymer chains. At equilibrium, these two forces are equal. The water chemical potential change can be calculated at constant temperature and pressure as follows: 𝜇1 − 𝜇1,0 = 𝛥𝜇mix + 𝛥𝜇elastic

(8.10)

Here, 𝜇1 is the chemical potential of water in the system, 𝜇1,0 is the chemical potential of pure swelling water, and 𝛥𝜇mix and 𝛥𝜇elastic are the mixing and elastic contributions to the total chemical potential change. The parameter 𝛥𝜇mix can be determined from the thermodynamics of the biomedical polymer-water mixing process and 𝛥𝜇elastic can be determined from the rubber elasticity theory as reported below: 2 𝛥𝜇mix = 𝑅𝑇[ ln(1 − 𝜈2,𝑠 ) + 𝜈2,𝑠 + 𝜒1 𝜈2,𝑠 ]

(8.11)

and 1/3 𝛥𝜇elastic = (𝑅𝑇𝑉1 /𝜈𝑀𝑐 )(1 − 2𝑀𝑐 /𝑀𝑛 )(𝜈2,𝑠 −

𝜈2,𝑠 ) 2

(8.12)

where, χ1is the biomedical polymer-water interaction parameter, V1 is the molar volume of water, 𝜈̅ is the specific volume of the biomedical polymer, 𝜈2,s is the volume fraction of the swollen gel, Mc is the number-average molecular weight between the

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   151

crosslinks, and Mn is the molecular weight of linear polymer chains prepared under the same conditions without crosslinking. Equation (12) is written for the hydrogels that were crosslinked in the absence of a solvent. Equations (10) through (12) lead to the expression for the true Mc of a nonionized hydrogel: 𝜐 𝑉1

2 1 = − 𝑀𝑐 𝑀𝑛

2 [ln (1 − 𝜈2,𝑠 ) + 𝜈2,𝑠 + 𝜒𝜈2,𝑠 ] 1/3 [𝜈2,𝑠 − 𝜈2,𝑠 ( 𝜑2 )]

.

(8.13)

For the case of bio-hydrogels crosslinked in the presence of water, equation (13) is modified to account for the water-induced elastic contributions swelling: 2 1 = − 𝑀𝑐 𝑀𝑛

𝜐 𝑉1

2 [ln (1 − 𝜈2,𝑠 ) + 𝜈2,𝑠 + 𝜒𝜈2,𝑠 ] 𝜈

𝜈2,𝑟 [( 𝜈2,𝑠 ) 2,𝑟

1/3

−

𝜙 2

𝜈

(8.14)

( 𝜈2,𝑠 )] 2,𝑟

where, 𝜈2,r is the volume fraction of the polymer in the relaxed state, i.e. immediately after crosslinking but prior to swelling/deswelling [48], and 𝜑 is the functionality of the crosslinking agent. The primary mechanism of release of many drugs from hydrogels is diffusion, occurring through the space available between macromolecular chains. This space is often regarded as the ‘pore’. A structural parameter that is often used in describing the size of the pores is the average distance between consecutive crosslinks, or mesh size, ξ [49]. From the knowledge of number average molecular weight between the crosslinks Mc it is possible to estimate the end-to-end distance of the unpertubated (solvent free) 1/2 state, (𝑟0−2 ) : 1/2

(𝑟0−2 )

= 𝑙 (2

𝑀𝑐 1/2 1/2 ) 𝐶𝑛 𝑀𝑟

(8.15)

where l is the bond length, Cn is the characteristic ratio of the polymer, Mr is the molecular weight of the repeating unit. Finally, the mesh size, ξ, can be calculated from: 𝜉 = (𝑟0−2 )

1/2

−1/3 𝑉2,𝑠 .

(8.16)

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8.3 Biomedical applications 8.3.1 Tissue engineering Recently hydrogels have become especially attractive to the new field of ‘tissue engineering’ as matrices for repairing and regenerating a wide variety of tissues and organs. In fact, numerous strategies are currently being introduced for the realization of hydrogels used as scaffolds, a synthetic extra cellular matrix (ECM) able to organize cells into a three-dimensional architecture and to present stimuli that direct the growth and formation of a desired tissue. Hydrogels possess features that are suitable to act as synthetic ECM; indeed they may contain pores large enough to accommodate living cells, or they may be designed to dissolve or degrade, releasing growth factors and creating pores into which living cells may penetrate and proliferate [50]. They have many different functions in this field being applied also as space filling agents (providing bulking, preventing adhesions, or functioning as bioadhesives) and as drug delivery vehicles [51]. Adequate scaffold design and material selection for each specific application depend on several variables, including physical properties (e.g. mechanics, degradation, gel formation), mass transport properties (e.g. diffusion), and biological properties (e.g. cell adhesion and signaling). For example, in the use as bulking material the most basic design requirements are the abilities to maintain a desired volume and structural integrity for the required time. So the success of the constructs is highly dependent on the design of the scaffold but also of the selected material. Over the past decade, a variety of naturally- and synthetically-derived materials have been utilized to form gels for tissue engineering applications [52]. Natural biomaterials present macromolecular properties similar to the natural ECMs so they present important advantages exhibiting excellent biocompatibility and bioactivity while synthetic polymers may lack informational structure for positive cell biological response. However, synthetic hydrogels have chemical and physical properties more controllable and reproducible than those of natural polymers and present the advantage of being reproducibly produced with specific block structures, molecular weights, and degradable linkages; so compared to natural ones, they can offer improved control of the matrix architecture and chemical composition. One approach to creating an ideal hydrogel for tissue engineering applications is to realize composite hydrogels incorporating bioactive elements into synthetic hydrogels for increased cellular bioactivity. A variety of synthetic and naturally derived materials may be used to form hydrogels for tissue engineering scaffolds. Example of synthetic materials shown in Tab. 1 are poly(ethylene oxide) (PEO), poly(vinyl alcohol) (PVA), poly(acrylic acid) (PAA), poly(propylene furmarate-co-ethylene glycol) (P(PF-co-EG)), and polypeptides whereas representative naturally-derived polymers include agarose, alginate, chitosan, collagen, fibrin, gelatin, and hyaluronic acid (HA). As regards the first class, PEO

8.3 Biomedical applications   

   153

and the chemically similar poly(ethylene glycol) (PEG) are currently the most commonly applied synthetic hydrogel polymers for tissue engineering. They are hydrophilic polymers that can be photocrosslinked by modifying each end of the polymer with either acrylates or methacrylates or mixed with the appropriate photoinitiator and crosslinked via UV exposure [53]. Moreover, block copolymers of PEO and poly(llactic acid) (PLLA) or PEG and PLLA are used to obtained thermally reversible hydrogels and degradable hydrogels containing hydrolytically degradable poly(lactic acid) (PLA) and enzyme specific cleavage sequences of oligopeptides [54]. Another synthetic hydrophilic polymer widely explored for use in space filling and drug delivery applications is PVA that can be physically crosslinked through repeated freezing/thawing methods or chemically crosslinked with glutaraldehyde [55], succinyl chloride, etc. Collagen and hyaluronic acid (HA) are the most attractive materials for applications in tissue engineering being the main components in connective tissues in human and other animals. The first presents a basic structure composed of three polypeptide chains, which wrap around one another to form a three-stranded rope structure [56] and its main advantage is a controlled degradation locally by cells present in the engineered tissue because it is naturally degraded by metalloproteases, specifically collagenase, and serine proteases [57–58]. The mechanical properties of a hydrogel realized with collagen can be enhanced by introducing various chemical crosslinkers (i.e., glutaraldehyde, genipin and carbodiimide). Hyaluronic acid, consisting of multiple repeating disaccharide units of N-acetylD-glucosamine and D-glucuronic acid, plays an essential role in many biological processes such as tissue hydration, nutrient diffusion, proteoglycan organization, and cell differentiation and the main properties that make the use of this material advantageous are its good biocompatibility and biodegradability. Moreover, hyaluronic acid hydrogels, formed by covalent crosslinking with hydrazide derivatives, esterification, and annealing [59], often include both collagen and alginate to form composite hydrogels with better properties [60]. Like HA, both alginate and chitosan are hydrophilic, linear polysaccharides. Alginate is a hydrophilic and linear polysaccharide composed of (1–4)-linked β-D-mannuronic acid (M) and α-L-guluronic acid (G) monomers, used in a variety of medical applications including cell encapsulation and drug stabilization and delivery, because it gels under gentle conditions, has low toxicity, and is readily available. Despite its advantageous features, alginate may not be an ideal candidate for tissue engineering because it does not specifically degrade. Chitosan is polycationic polysaccharide contains glucosamine and N-acetylglucosamine molecules, thus structurally similar to naturally occurring glycosaminoglycans (GAGs) and it is considered a biodegradable polysaccharide, which can be metabolized by human enzymes such as lysozyme [61]. Chitosan has been investigated for a variety of tissue engineering applications in recent years due to its biocompatibility, biodegradability, low immunogenicity and cationic nature, but unmodified

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chitosan can only be dissolved in acidic solutions due to its strong intermolecular hydrogen bonds and this limits its applications as an injectable hydrogel. As regards the specific applications of scaffold as space-filling agents, generally they are used as bulking material to treat conditions of vesicoureteral reflux, urinary incontinence, vocal fold augmentation [62–63] and above all for both plastic and reconstructive surgery. An example is the use of dextranomer/HA copolymer for vesicoureteral reflux [64]. Moreover, they are used to prevent post-operative adhesions; in particular in their use as anti-adhesives, synthetic materials are preferred because cells lack adhesion receptors to them. PEG hydrogels for example have been used in several applications to prevent post-operative adhesion and in some approaches it is also grafted to other materials such as the poly l-lysine to form copolymers that can adhere on one side and provide a non-adhesive brush-like surface on the other [65]. HA-based hydrogels are widely used as dermal fillers (DF) for the rejuvenation of the dermis. HA is an important structural element in the skin; its concentration in the dermis decreases with age promoting the formation of wrinkles. With the use of dermal fillers it is possible to replace them, achieving a natural and younger appearance. One of the most important aspects of the clinical use of DF is their persistence in the human body. The performance of dermal filler in vivo seems to be directly affected by physical materials properties. In particular, the clinical data appear to correlate with the concentration of the polymer and with the product between the concentration and the percent elasticity, so these should be crucial parameters for the clinical performance of DF [66]. The use of gels as hydrated three-dimensional networks of polymers that provide a place for cells to adhere, proliferate, and differentiate finds application in the engineering of a wide range of tissues: cartilage, bone, muscle, fat, and neurons. For example, weak gels consisting of HA containing alginate and HA alone, promote a differentiated chondrocyte phenotype and expression of type II collagen [67]. Also, alginate has been used more widely than other hydrogels to assess the in vivo potential of hydrogel scaffolds for cartilage engineering. In several applications it has been mixed with chondrocytes and injected into the site of interest [68]. Moreover, there are several examples of composite constructs such as semi-interpenetrating networks obtained interpenetrating HA in a fibrillar collagen scaffold [43] or also composite hydrogels that combine the advantages of both natural and synthetic polymers for engineering fibro-cartilaginous tissues such as porous three-dimensional composite scaffolds realized using hyaluronic acid derivatives and PCL to improve mechanical properties [69]. Thanks to its properties, PCL is also used in different applications for bone tissue engineering [70], but to exhibit improved tensile strength and moduli in hard tissue engineering PCL scaffolds are reinforced by inorganic nanofiller. Halloysite nanoclay for example has the potential to substantially improve both strength [71] and bioactivity and has been investigated in the processing of bone cements [72] and recently also used to prepare PCL/NC fibrous scaffolds via electrospinning technique

8.3 Biomedical applications   

   155

after intercalating nanoclay within PCL by solution intercalation method [73]. Moreover, hydrogels are being widely investigated to engineer nearly every tissue such as adipose tissue. In this field hyaluronic acid derivatives have shown the ability to favor adipocyte growth [74], blood vessel, skeletal muscle [75–76], intervertebral disk [77]. As an injectable carrier to engineer the nucleus pulposus, different hydrogels based on modified hyaluronic acid have been used [78].

8.3.2 Drug delivery Nowadays hydrogels play an important role in many biomedical applications, such as tissue engineering scaffolds, biosensor and drug carriers. Among these applications, hydrogel-based drug delivery devices have become a major area of research interest with several commercial products already developed  [79] because of their unique physical properties. The hydrogels’ porous structure can easily be tuned by controlling the density of crosslinks in the gel matrix and the affinity of the hydrogels for the aqueous environment in which they are swollen. Drugs can be loaded into the hydrogel pores and subsequently released at a rate dependent on the diffusion coefficient of the small molecule or macromolecule through the gel network. Using hydrogels it is possible to control pharmacokinetic: specifically that a depot formulation is created from which drugs slowly elute, maintaining a high local concentration of drug in the surrounding tissues over an extended period, although they can also be used for systemic delivery. Usually hydrogels are also highly biocompatible, therefore they can be successfully used in vivo. Biocompatibility is promoted by the high water content of hydrogels and the physiochemical similarity of hydrogels to the native extracellular matrix, both compositionally (particularly in the case of carbohydrate-based hydrogels) and mechanically. Biodegradability or dissolution can occur through enzymatic, hydrolytic, or environmental (e.g. pH, temperature, or electric field) pathways, but degradation is not always desirable depending on the time scale and location of the drug delivery device. Hydrogels are also rather deformable and can conform to the shape of the surface to which they are applied. Although hydrogels have so many advantageous properties, they also have several limitations. The low tensile strength of many hydrogels limits their use in load-bearing applications and can result in the premature dissolution or flow away of the hydrogel from a targeted local site. In addition, the quantity and homogeneity of drug loading into hydrogels may be limited, particularly in the case of hydrophobic drugs. The high water content and large pore sizes of most hydrogels often result in relatively rapid drug release. Ease of application can also be problematic; although some hydrogels are sufficiently deformable to be injectable, many are not, necessitating surgical implantation. Therefore, a successful drug delivery device relies on intelligent network design and on accurate a priori mathematical modeling of drug release profiles. Using an ordered polymer network, composed of macromers with

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well-understood chemistries, it is possible to obtain hydrogels with well-defined physicochemical properties and reproducible drug-release profiles. An aspect to be highlighted regarding the use of hydrogels in drug delivery is that they are excellent candidates for encapsulating biomacromolecules including proteins and DNA therapeutics due to their lack of hydrophobic interactions which can denature these fragile species [80].

8.3.2.1 Design criteria for hydrogels in drug delivery Typically the high water content of most hydrogels induces rapid release of drugs from the gel matrix, particularly in the case of hydrophilic drugs. For this reason, a range of strategies have been explored to reduce the release rate of drug from hydrogels. Numerous design criteria are crucial for drug delivery formulations and have to be evaluated before hydrogel fabrication and drug loading. Among transport properties, the most notable variable is the drug diffusion coefficient, which depends on the drug molecular size and is characteristic of the polymer network. Hydrogel crosslinking density has a great influence on diffusivity; in addition, physical properties of the hydrogel also affect drug release. Regarding the physical properties, polymer molecular weights, composition, and polymer/initiator concentrations influence hydrogel swelling and degradation. Eventually, the stimuli-responsiveness of a hydrogel network can also mediate the amount and rate of drug delivery. The understanding of transport and physical properties is especially crucial in modeling molecule release. However, even if a hydrogel delivery formulation is designed with the appropriate physical and transport properties, it may still fail to perform its therapeutic role when implanted in vivo due to a localized inflammatory response. Proper material selection, fabrication process, and surface texture of the drug delivery device are therefore always critical in designing biocompatible hydrogel formulations for controlled release. Both physical and chemical strategies can be employed to enhance the binding between a loaded drug and the hydrogel matrix to extend the duration of drug release, as illustrated schematically in Fig. 8.5. Electrostatic interactions between charged drugs and ionic polymers have often been employed to increase the strength of the interactions between the hydrogel and a target drug to delay drug release. Because of their multivalent anionic charge, phosphate-functionalized polymers are useful. As an example, the uptake of cationic lysozyme into N-isopropylacrylamide-based hydrogels functionalized with polyoxyethyl phosphate-containing comonomer is significantly enhanced compared to nonfunctionalized PNIPAM hydrogels [81]. Even carbohydrate-based polymers are useful to prolong the release of a charged drug, due to the presence of anionic and cationic functional groups [82]. Charge interaction has also been cited as one of the reasons for using hyaluronic acid as a delivery vehicle for local anesthetics [83], since it is anionic, while most local anesthetics are cationic in aqueous solution.

8.3 Biomedical applications   

a

b

drug

bound drug

cleavable cross-linker hydrogelbound charge

   157

free drug

water, enzyme polymer

Fig. 8.5: Strategies for enhancing the interaction between a loaded drug and a polymeric gel to slow drug release: (a) physical, (b) chemical.

A drug can also be conjugated to the hydrogel matrix through a covalent bond, in this case its release is mainly controlled by the rate of chemical or enzymatic cleavage of the polymer–drug bond. Alternately, drug release may be regulated via the hydrolysis of the polymer backbone, possibly inducing the release of a partially modified drug analogue. Alternatively, a crosslinker between the drug and the polymer can be engineered to give specific durations of release. For example, by changing the length of a sulfide-based crosslinker from three to four carbons, the time required to release bound paclitaxel from a hydrogel was increased from approximately 4 days to 2 weeks [84].

Incorporation of drugs Hydrogels have been used to deliver hydrophilic, small-molecule drugs, characterized by high solubilities in both the hydrophilic hydrogel matrix and the aqueous solvent swelling the hydrogel. In this case, it is fairly simple to load a high quantity of drug into a swollen hydrogel by simple partitioning from a concentrated aqueous drug solution and subsequently releasing the hydrophilic drug payload into an aqueous environment. Nevertheless, in the case of large macromolecular drugs (e.g. proteins, nucleic acids, etc.), this process is relatively ineffective, due to their diffusive limitations to their partitioning into a hydrogel phase. Similarly, for hydrophobic drugs, which are scarcely soluble in both the aqueous and the hydrogel phases, the process is quite inefficient. However, both of these classes of drugs are becoming increasingly important clinically as a result of improved understanding of the molecular mechanism of disease and the more frequent application of molecular design approaches for small-molecule drug design. Macromolecular drug uptake is typically limited by the diffusion through the hydrogel network. This problem can be overcome at least partially by engineering the pore size of hydrogels. Hydrogel-based hydrophobic drug delivery is in many aspects a more difficult problem because of the intrinsic incompatibility of the hydrophilic hydrogel network and the hydrophobic drug. Thus, for this kind of drug there are two main issues: how to load the hydrophobic drug into the gel matrix and, once present, how to effectively release the drug into the aqueous

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gel environment. Numerous strategies have been used to improve hydrophobic drug loading into hydrogels. One simple approach is to generate hydrophobic domains within hydrogels. A simple method to do this is the copolymerization with hydrophobic comonomers, introducing statistically distributed hydrophobic sites within the polymer networks. This strategy introduces binding sites for hydrophobic drugs and condenses the bulk dimensions of the gel, reducing the average pore size and slowing diffusion-limited release. For example, copolymerization of acrylic acid-2-ethylhexyl ester in a methacrylic acid-based hydrogel improves the loading of p-hydroxyanisole and extends its release from the hydrogel independently from pH [85]. More recently, molecular design approaches have been applied to maximize the affinity of a polymer-bound hydrophobic domain for a particular drug target while minimizing the non-specific binding of other hydrophobic compounds in the gel environment. Performing a screening of a range of different small molecules with hydrophobic binding properties it is possible to identify those which bind most strongly to a given drug. Alternately, hydrogel networks can be modified to generate hydrophobic domains. Hydrophobic side chains can be grafted onto the polymer precursors which can selfassemble to form hydrophobic domains within the bulk hydrogel network and bind hydrophobic drugs. Using grafting or copolymerization methods to render the hydrogel hydrophobic one should take into account that the hydrogel might de-swell and delocalized surface and bulk hydrophobicity, introduced into the hydrogel network, could potentially reduce the biocompatibility and/or the low protein binding properties of hydrogels. Cyclodextrins are of interest in this context given their hydrophilic exterior, which is useful for maintaining the bulk hydrophilicity and swelling state of the hydrogel, and their hydrophobic interior, which can facilitate the entrapment and controlled release of hydrophobic drugs. For example, grafting cyclodextrin to the hydrogel provides improved control over drug release kinetics.

8.3.2.2 Drugs release from hydrogels formulations To properly design a hydrogel system to be used as a drug delivery carrier the modeling of drug delivery phenomena and the accurate prediction of release profiles from complex hydrogel systems are pivotal. Creating a fundamental understanding of drug transport processes is the first step towards developing a suitable mathematical model. Mass transport governs the translocation of drug from the hydrogel’s inner part to the surrounding environments. A variety of factors affect the mass transport of encapsulated molecules, among them the network crosslinking density, extent of swelling, gel degradation, the size and charge of the encapsulated molecules, and the physical interactions these molecules exhibit for themselves and for the polymer matrix. If specific drug-binding moieties are present within the hydrogels, the kinetics and/or thermodynamics of drug–ligand binding need to be taken into account and quantified to predict the controlled release of the encapsulated molecules.

8.3 Biomedical applications   

   159

Dynamic hydrogels For degradable hydrogels the rates of matrix swelling and degradation govern the diffusion of encapsulated or tethered molecules. To accurately predict the unique molecule release profiles that occur with many degradable hydrogels, additional parameters, not taken into account for not biodegradable drug carriers, need to be considered. For example, in the case of enzymatically degradable hydrogels, one challenging issue is how to model the rate of enzyme (e.g. matrix metalloproteinases, MMPs) production by invading cells. Stimuli-sensitive hydrogels can sense changes in complex in vivo environments and utilize these triggers to modify drug release rates. For this system, it is critical to model the dynamic swelling response in order to predict solute release. In the case of ionic or pH-sensitive hydrogels at a fixed pH and salt concentration, the swelling of ionic hydrogels is balanced by the osmotic pressure and the relaxation of the polymer chains. Thermodynamically, the total free energy can be expressed as: 𝛥𝐺𝑇 = 𝛥𝐺𝑒 + 𝛥𝐺𝑚 + 𝛥𝐺𝑜 .

(8.17)

Here, 𝛥GT is the total Gibbs free energy, 𝛥Ge is the free energy contributed by the elastic force of the polymer chains, 𝛥Gm is the free energy of mixing, and 𝛥Go is the free energy due to osmotic pressure. When the swelling of an ionic hydrogel is in equilibrium (𝛥GT = 0), the decreased elastic free energy is balanced by the free energy of mixing and osmotic pressure. Based on this concept, [86–89], a continuum model has been developed to describe the macroscopic behaviors of pH-responsive poly(methacrylic acid) (PMAA) hydrogel membranes accounting for charge density, ionic strength, stress, strain, and electric field [87]. It was found that membrane swelling was slower than shrinking. Following the equilibrium, swelling of anionic pHresponsive hydrogels appears to be proportional to pH with a sharp increase around the pKa of the charge group. A mathematical simulation of drug release from thermo-sensitive hydrogels has to correlate gel swelling and diffusive controlled molecule release. The following equation (18) was used to assess glucose and insulin diffusivities in n-isopropylacrylamide gels [90]: 𝐷𝑒 (1 − 𝜙)3 = 𝐷0 (1 + 𝜙)2

(8.18)

where De and D0 are the effective molecule diffusivities in the gel and in pure solvent, respectively. 𝛷 is the polymer volume fraction of the gel. Using this equation, the effective diffusivities of molecules encapsulated within thermo-sensitive hydrogels can be estimated as a function of temperature. Once the molecule diffusivity is determined, a release profile can then be predicted using Fick’s law of diffusion [90].

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Composite hydrogels Modeling drug release from composite hydrogel systems has proven to be challenging due to the anisotropy of these systems, both in terms of spatial and transport properties. Composite hydrogels can be either multi-layer or multi-phase. These composite systems have great potential in delivering multiple protein therapeutics for tissue engineering applications where temporal and spatial control over drug delivery is desirable. The simultaneous delivery of multiple proteins is known to occur in vivo during angiogenesis, bone remodeling, and nerve regeneration. For example, several angiogenic proteins including vascular endothelial growth factor (VEGF), basic fibroblast growth factor (bFGF), transforming growth factor beta (TGF-β), platelet-derived growth factor (PDGF), and matrix metalloproteinases (MMPs) are involved in the angiogenesis process. Marui et al. discovered that the dual delivery of bFGF and hepatocyte growth factor (HGF) from collagen microspheres greatly increased blood vessel formation in an animal model [91]. Crosslinked hyaluronan (HA) hydrogels have been used to simultaneously deliver VEGF and keratinocyte growth factor (KGF) to enhance angiogenesis [92]. Alginate hydrogels have been used to deliver bone morphogenetic protein-2 (BMP2) and transforming growth factor-β (TGF-β3) and showed enhanced bone formation compared to delivery of either single protein [93]. In multi-layer systems, a basal polymer layer is fabricated, followed by lamination of subsequent layers. Different proteins can be encapsulated into each layer during fabrication and tunable multiple-protein release or unique single-protein release profiles are made possible by independently adjusting the crosslinking density of each layer. Many models have been developed for predicting drug release from multi-layer hydrogel composites [94–95]. In addition to multiple-protein delivery, multi-layer matrices can also be used to decrease the problematic burst release. A multi-laminated hydrogel system has been prepared by photopolymerization to obtain a zero-order release profile, through a non-uniform initial drug loading in multi-laminated hydrogels.[96–98]. Another strategy for multiple-protein delivery is multi-phase systems. In this approach, prefabricated microspheres containing one or more proteins are uniformly embedded within a hydrogel containing a second protein [99–101]. The release of the microsphere-encapsulated protein is delayed due to the combined diffusional resistances of the microsphere polymer and the surrounding gel. Hydrogels containing gelatin microspheres were fabricated to independently control the delivery of insulinlike growth factor-1 (IGF-1) and transforming growth factor-β1 (TGF-β1). Release profiles can be adjusted by varying the protein loading in each polymer phase [99]. 

Micro-nanoscale hydrogels Over the past few decades, polymeric microspheres and, more recently, nanoparticles have been widely used for sustained or targeted drug delivery  [102]  as well as cell

8.3 Biomedical applications   

   161

encapsulation [103–105].  Alternatively, micro/nanoparticles made from hydrophilic hydrogels are more suitable for encapsulating these fragile biomacromolecules, such as DNA. These miniaturized drug-containing vehicles can be fabricated in vitro and then administered via an oral [106–107] or nasal route [108–109] or injected into the patients in a minimally invasive manner to increase patient compliance. Two types of mathematical approaches have been used to predict molecule release from hydrogel microspheres: macroscopic diffusion models and microscopic Monte Carlo simulations. [110–111]. For macroscopic modeling, the most applicable models are still based on Fick’s second law of diffusion. One of the unique challenges facing microscaled matrix delivery systems is burst release due to the high surface-to-volume ratio of these particulate systems [112–113]. Burst release may cause a ‘dose-dumping’ effect and is potentially harmful to patients in clinical applications. Several methodologies have been developed in an attempt to decrease the degree of burst release. These include increasing crosslinking density of the matrix surface  [114–115], coating additional drug-free layers  [95,  112, 116], embedding the drug-containing particles within a bulk polymeric matrix  [99–101, 117], and loading drug unevenly with higher concentrations toward the center of the matrix [118–119]. The prediction of burst release is problematic as the exact mechanism has not been elucidated. Typically, diffusion-controlled release can be divided into two phases: a rapid burst phase and a prolonged diffusion-controlled phase. [112]

In situ forming hydrogel As previously said, injectable in situ forming hydrogels have been widely exploited for drug delivery applications. Several physical or chemical crosslinking mechanisms have been used for in situ network formation. The common disadvantage of physical crosslinking, however, is that the gels thus formed are unstable and may disintegrate rapidly and unpredictably. For long-term drug delivery applications, covalent crosslinking methods performed under physiological conditions, such as photopolymerization of multi-vinyl macromers, are more favorable compared to physical crosslinking methods as they produce relatively stable hydrogel networks with predictable degradation behaviors. Photopolymerization of degradable hydrogels has been applied in protein [120– 122]  and gene delivery [123–125] and permits in situ encapsulation of these species during network fabrication. When in situ forming hydrogels are used to deliver macromolecules such as DNA and protein, reduced or incomplete release of these biomolecules is commonly observed [122–125]. The factors influencing incomplete biomolecule release from these hydrogel carriers has commonly been attributed to the fabrication processes. When in situ forming gels are used to deliver proteins, irreversible interactions between the encapsulated proteins and polymerizing polymer chains decrease the

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efficacy of the therapeutic agent. The modification of hGH has been identified by reactive thiol macromers in a PEG-based hydrogel system prepared via Michael-type addition reaction  [126]. However, free radicals produced from the photoinitiation process can attack DNA molecules during UV irradiation, leading to DNA damage. Based on similar observations during protein encapsulation, Lin and Metters utilized a metal-ion-chelating molecule, iminodiacetic acid (IDA), to block undesirable protein–polymer conjugation reactions mediated by free radicals.

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Index antibacterial applications 107 apatite 94 apatite crystals 96 articular cartilage 41 – mechanical properties 43 – organization 42 articular replacement 44 β-tricalcium phosphate (β-TCP) 77 bioactive 90 bioactive ceramics 77, 99 bioactive glass-ceramics 99 bioactive glasses 90 bioactivity 17, 19, 21, 23, 24, 25, 27, 29, 78 bioceramics 92 biocompatibility 17, 19, 22, 23, 25, 40, 78, 95, 98 biocompatible material 89 biocomposites 17, 21 biodegradable polymers 70 biodental materials 90 biomaterial 37, 89 biomedical implants 69 bioresorbable 70, 79, 80 bioresorbable composite 69, 70, 80 bioresorbable membranes 80 bone repair 1

composite nanofibers 113 composites manufacturing methods 78 compression molding 78 copolymers 76 copper oxide (CuO) 121 crosslinked polymers 33 crosslinking 39 crystallinity enhancement 96 – by adding activator agents 96 – by sintering process 98 crystallization 96, 98 dental glass-ceramics 98 dental implant 93 dental restorations 98 doping 113, 117 drug delivery 143, 152, 153, 155, 156, 157, 158, 160, 161 electrospinning process 109 – setup 109 electrospun nanofibers 111 – antimicrobial applications 112 ethylene oxide 78 extrusion 78 foot and ankle 79

calcium phosphate glass fibers 78 calcium silicate 77, 78 cartilage 41 – frequent medical issues 44 cartilage replacement 33, 44 cemented total hip replacements 53 ceramic polymer composites 1 – requirements 2 cerium oxide (CeO2) 128 chitosan based composites 10 – biofunctionality 10 – derivatives 11 – disadvantages 10 – fabrication techniques 10 – mechanical properties 10 – tissue compatibility 10 clinical applications 79 cobalt oxides (CoO and Co3O4) 126 comparison of cemented and cementless THR 55 composite materials 93

gamma irradiation 78 GBR 80 glass-ceramics 90 glass-ceramics and bioactive glass-ceramics 90 – Bioceramics 92 guided bone regeneration (GBR) 76 HA 77 HAp-Ti/Ti6Al4V 17, 18, 27, 29 hard tissue applications 1 hip joint prosthesis 53 historical background 89 hydrogel 33, 141, 142, 143, 144, 145, 146, 147, 148, 149, 150, 151, 152, 153, 154, 155, 156, 157, 158, 159, 160, 161 – categories 34 – hydrogels in situ 145, 146 – in situ forming hydrogel 143, 161 – mechanical properties 36

170   

   Index

– properties 34 – swelling 142, 147, 150, 151, 156, 158, 159 hydrogels as biomaterials 37 hydroxyapatite (HA) 77, 99 hyperthermic treatment 100 injection molding 78 in vitro and in vivo test for bioactivity 101 knee 79 load-bearing 17, 21, 22, 25 mechanical properties 17, 19, 21, 22, 25, 26 melt molding 78 metallic implant materials 89 – Dental amalgam 90 metal oxide-based one-dimensional titania nanostructures 107 nickel oxide (NiO) 124 non-metal restorations 98 nucleating agents 96 nucleation 96, 98 orthognatic surgery 79 orthopaedic applications 73 patterned glass 91 PCL 69, 73, 76, 77, 78 PDLLA (poly-DL-lactide) 75 pediatric orthopedics 80 PGA 73, 75, 76, 80 phosphate-based glasses 77 PLA 73, 74, 75, 76, 77, 78, 80 plasma 17, 20, 21, 24, 25, 26, 28, 29 PLGA 73, 76, 78 PLGA copolymers 76, 78 PMMA composites 60 polyester based composites 7 – development potential 9 – mechanical characteristics 8 – tissue compatibility 9 polyethylene based composites 3 – advantages 3 – bioactivity 5 – mechanical properties 4 – production methods 4 polyglycolic acid 73

polylactic acid 73, 74 poly lactide-co-glycolide 73 polymer 73 polymer composites 53 polymethymethacrylate based composites 6 – mechanical properties 7 – toxic effects 6 poly(vinylalcohol) 33 polyvinyl alcohol (PVA) 38 polyvinyl alcohol (PVA) hydrogels 38 – biomedical applications 40 poly α-hydroxy acids 69, 73 poly ε-caprolactone (PCL) 73, 76 preparation techniques 92 resorbable membranes 76 self-reinforcing (SR) technique 79 shoulder 79 sinter-crystallization process 98 solid state drawing 78 spine 80 sterilization of bioresorbable composite 78 structure of glass-ceramics 94 TiO2 nanocomposites 108 – antibacterial mechanism 131 – applications 111 – development 114 – synthesis of nanofibers 109 tissue engineering 141, 152, 153, 154, 155, 160 titania 107 – dopants 108 titanium dioxide (TiO2) 107 – dopants 108 Total Hip Replacement (THR) 53 – economic and clinical factors 56 UHMWPE composites 57 wollastonite 94 zinc oxide (ZnO) 117 Zn-doped titania nanofibers 119