Biomaterials In Asia: In Commemoration Of The 1st Asian Biomaterials Congress 9789812835758, 9789812835741

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Biomaterials in Asia In Commemoration of the 1st Asian Biomaterials Congress

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Biomaterials in Asia In Commemoration of the 1st Asian Biomaterials Congress Tsukuba, Japan

6 – 8 December 2007

edited by

Tetsuya Tateishi National Institute for Materials Science, Japan

World Scientific NEW JERSEY

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LONDON

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SINGAPORE

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BEIJING

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SHANGHAI

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HONG KONG

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TA I P E I

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CHENNAI

Published by World Scientific Publishing Co. Pte. Ltd. 5 Toh Tuck Link, Singapore 596224 USA office: 27 Warren Street, Suite 401-402, Hackensack, NJ 07601 UK office: 57 Shelton Street, Covent Garden, London WC2H 9HE

British Library Cataloguing-in-Publication Data A catalogue record for this book is available from the British Library.

BIOMATERIALS IN ASIA In Commemoration of the 1st Asian Biomaterials Congress Copyright © 2008 by World Scientific Publishing Co. Pte. Ltd. All rights reserved. This book, or parts thereof, may not be reproduced in any form or by any means, electronic or mechanical, including photocopying, recording or any information storage and retrieval system now known or to be invented, without written permission from the Publisher.

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ISBN-13 978-981-283-574-1 ISBN-10 981-283-574-1

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PREFACE History of the Asian Biomaterials Congress Tetsuya Tateishi, Chairman of the 1st ABMC Fellow Emeritus, Biomaterials Center, National Institute for Materials Science Biomaterials Center, National Institute for Materials Science, Tsukuba, Japan

The 1st Far-Eastern Symposium on Biomedical Materials was held October 5-8, 1993, at the Beijing Friendship Hotel in China. The scale of the event included 11 invited lectures, 58 general topics, and 38 poster presentations. This first monumental international conference on biomaterials in Asia was made possible by the efforts of two preeminent figures in the field, then-Prof. Ikada of Kyoto University in Japan and Prof. Zhang of Sichuan University in China. The setting was equally impressive, with the leaves of the Chinese parasol tree visible throughout Beijing, trembling soundlessly in the laminar flow wind characteristic of the Chinese continent, which was blowing gently through the slightly cloudy mid-autumn Beijing sky. Although this symposium was renamed the Asian Symposium on Biomedical Materials (ASBM) starting with the fourth meeting, a total of seven meetings were held up to 2006. As data on the symposium, Table 1 shows the transition in the number of presentation topics from the 1st Far-Eastern Symposium to the 7th ASBM. This table was based on materials provided by Prof. Tabata of Kyoto University, who served as Secretary-General of the symposium during this period. Around 1990, the amount of research on biomaterials grew rapidly in Japan as a result of the great efforts of our predecessors. This was a period when

Table 1. History of ASBM.

biomaterials research left behind the conventional paradigm, in which a smattering of bioscience was incorporated in existing fields of materials research. At the same time, universities began to produce professional researchers who committed their research lives to the full-scale study of biomaterials from the start, and the biomaterials R&D population increased substantially. In fact, progress on the road to comprehensive biomaterials research began to accelerate, as evidenced by the expansion and reorganization of Kyoto University’s Medical Macromolecular Research Center, which existed from 1980 to 1990, as the Biomedical Engineering Research Center. Prof. Ikada, who was the Director of the Biomedical Engineering Research Center at the time, in cooperation with Prof. Zhang, Director of the Biomaterials Engineering Center at Sichuan University, which was a representative presence in biomaterials research in China, planned the 1st Symposium on Biomedical Materials in order to secure a venue for research presentations, exchanges of information, and exchanges of researchers in East Asia. Given the general conditions in Asia, the time was ripe for holding this kind of symposium. An interesting story concerning the inception of another Asian biomaterials symposium was provided by Prof. Yui of the Japan Advanced Institute of Science and Technology (JAIST). When Prof. Yui was planning the introduction of JAIST, which was be established shortly, Prof. Akaike of the Tokyo Institute of Technology recommended holding an international

symposium. This provided the first impetus toward the creation of a new symposium. Prof. Yui then consulted with Prof. Okano of Tokyo Women’s Medical University, which led to concrete discussions about launching a regular international symposium, centered on Asia, to provide a forum for disseminating the results of research on biomaterials. Because Prof. Tsuruta had just joined JAIST as a Visiting Professor, this distinguished scientist was asked for guidance on various matters. As a result, the 1st Asian International Symposium on Biomaterials (AISB) was held at JAIST in 1997 (fifth year of Prof. Yui’s appointment at JAIST). At this time, researchers in the field of macromolecular biomaterials from Japan, Korea, and Taiwan were invited and the basic policy for holding future international symposiums was determined. This symposium was originally planned as a biennial event. However, because the year 2000 was also the 60th birthday of Prof. S. W. Kim of the University of Utah, the second meeting was held in 2000 rather than in 1999. Prof. Yui served as Secretary-General until the third meeting and, until then, the symposiums centered on Japan, Korea, and Taiwan. At each of the first three meetings, both proceedings and a commemorative book were issued. Japan was again the host country for the 4th Symposium, and Prof. Akashi, who was at Kagoshima University at the time, was to be in charge. However, just at that time, it was decided that he would move from Kagoshima University to Osaka University, and Prof. Tanaka, who was Managing Director of the Biomaterials Center at NIMS (presently professor at Tokyo Institute of Technology) was hastily placed in charge. Because this symposium was intended from the outset to be a public event of the Japan Society for Biomaterials, the fourth meeting was held under the sponsorship of that organization and, as a result, the scope was expanded to encompass not only macromolecular materials, but also fields related to metals and inorganic materials. Table 2 shows the transition in the number of presentation topics from the 1st AISB to the 5th AISB, based on data provided by Prof. Yui. From the 1st to the 4th meeting, the planning focused on invited lectures, and a large number of general topics were adopted for the first time at the fifth meeting. Until the 1990s, orthodox biomaterials research centering on macromolecular materials, ceramics, metals, and other composites and surface treatment techniques for improving biocompatibility represented the main direction in biomaterials. However, since around 2000, the application of biologically-derived substances and hybrid techniques incorporating cellular materials have become conspicuous As a result of the rapid progress of research

and development in cell scaffold materials as a basic technology supporting regenerative medicine, this trend has played a role in pushing tissue engineering and regenerative medicine into the spotlight in medical engineering. Actually, the fact that the number of topics began to increase rapidly at both of these two Asian biomaterials symposiums during this period truly tells the story of the conditions at that time. In addition, in 1998, the Biomedical Engineering Research Center at Kyoto University was reorganized as the Regenerative Medical Science Research Institute. On a personal note, on April 1, 2000, I transferred from a laboratory at the Agency of Industrial Science and Technology (former AIST), where I had worked for many years, to a professorship in the Department of Mechanical Engineering, School of Engineering, at The University of Tokyo, where I started a laboratory in bioengineering for regenerative medicine. The following year, on April 1, 2001, I became the concurrent managing director of the Tissue Table 2. History of ASBM.

Engineering Research Center at the National Institute of Advanced Industrial Science and Technology (new AIST; reorganized as an Independent Administrative Institution from the aforementioned Agency of Industrial Science and Technology). Thus, I also have many fond memories of efforts to keep up with the rapidly changing trends of the times. Recent research on biomaterials also includes nanobiotechnology, drug delivery systems (DDS), genetic engineering, molecular biology, and other

fields, in addition to bioengineering for regenerative medicine. This feels like a whole new world from the 1970s, when my own involvement in biomaterials began. In recent years, at both the ASBM and the AISB, the expansion in the range of presentation fields and the increase in the number of topics have been remarkable, reflecting this diversity. As a natural development, many persons have expressed hopes for the realization of an organization such as the Asian International Union of Societies for Biomaterials Science and Engineering, which attempts to create an organization of groups of specialists which can cope with this kind of diversification of research fields and overcome national and regional barriers. The same opinion has also been heard from influential persons in leadership positions in the biomaterials societies of the various Asian nations, and the Japanese Society for Biomaterials understands that negotiations have already begun, centering on directors with responsibilities for international matters. As described in this history, at least two general symposiums on biomaterials currently exist in Asia, and many researchers and scientists have continued to attend and present research topics at both. Given this situation, motions were made at the ASBM-7 steering committee meeting on Jeju Island and the AISB-5 steering meeting on Xiàmén in 2006 to hold the next round of symposiums in 2007 jointly, and the agreement of prominent persons leading Table 3. Number of papers submitted for 1st ABMC.

these two symposiums was obtained. As a result, the 1st Asian Biomaterials Congress (ABMC) was held in Tsukuba City, Ibaraki Pref., Japan (see Fig. 1). This Asian Biomaterials Congress was scheduled for December 6-8, 2007, at the Epochal Tsukuba International Congress Center, with a post-congress seminar and tours of laboratory facilities on December 8.

Figure 1. Location of Tsukuba City in Japan.

The Tsukuba area is about 60 km from Tokyo. It was originally a rural farming district. In 1963, it was chosen by the Japanese Government as the site of a new city that was to be devoted to academic research. You can see the results of that decision here today. During the period of more than 30 years since Tsukuba Science City was constructed, Tsukuba has come to play the role of a showplace for Japanese science and technology. From the viewpoint that Tsukuba was constructed as a large-scale science city led by the relocation of national research institutes and universities, mostly from Tokyo and the surrounding area, it can also be said this was Japan’s first “experimental city.” Originally, 45 national research institutes, 4 universities, and more than 200 private company laboratories formed the nucleus of the city; its population now exceeds 220,000. In the past, however, transportation was inconvenient, making Tsukuba seem like an island isolated from the mainland, and it was necessary to overcome this serious handicap when hosting international conferences. In August 2005, a high-speed rail line called the Tsukuba Express began service, reducing travel time from Akihabara in Tokyo to Tsukuba to only 45 minutes. As a result, Tsukuba is an increasingly popular spot in the northern Kanto area,

and its population is also increasing rapidly. Because the Epochal Tsukuba International Congress Center is less than a 10-minute walk from Tsukuba Station, which is the terminus of the Tsukuba Express, and is part of a complex that includes the Hotel Okura, this is a very suitable venue for the 1st Asian Biomaterials Congress. While the attendance at the first convention in 1993 did not exceed 110, for this present meeting we received about 400 applications from individuals in 15 countries desiring to attend this congress (Table 3). It has given us great pleasure to welcome the many participants, and I feel that we have had a series of extremely interesting and significant discussions. I hope that the exchange we had here will provide the basis for the development of new ideas and principles in the field of biomaterials and stimulate active future collaboration, while at the same time contributing to international friendship and understanding. In order to realize the immense potential of biomaterials and tissue engineering, an intensive international effort will be required to provide the basic structure-function relationships from the molecular to the tissue level and to develop the engineering systems analysis needed to produce functional tissue replacements including key technologies, such as biomolecular factors, cells, biomaterials, engineering design, bioinformatics. and cell-based technologies. I think the Congress will be held every other year but, in any case, what’s important is to create an association of the biomaterials societies in all countries. This will be a topic at the upcoming Congress, and I expect that we might see the birth of an Association of Asian Biomaterials Societies. If we can set the stage for this, holding this Congress will be a significant event indeed. At present, the executive committee for the international congress, an advisory committee, international committee, and domestic committee for the Asian Biomaterials Congress have been organized. Relative to the subject matter of the congress, the objective was to hold a congress with a strong allAsia flavor by adopting the widest possible range of biomaterials and related fields throughout Asia. The advice and cooperation of all those concerned is still sought. In closing, I would like to express my sincere appreciation to Prof. Zhang, Prof. Ikada, Prof. Okano and other representatives from each of their respective countries for contributing so much to make the Congress successful.

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The First Asian Biomaterials Congress Tsukuba, Japan December 6-8, 2007

CHAIRMEN TATEISHI, Tetsuya, Japan USHIDA, Takashi, Japan

Executive Committee IKADA, Yoshito, Japan OKANO, Teruo, Japan ZHANG, Xingdong, China XI, Tingfei, China LENG, Yang, Hong Kong, China KIM, Young Ha, Korea HSIUE, Ging-Ho, Taiwan, China TEOH, Swee-Hin, Singapore

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CONTENTS

Preface

v

Organizing Committee

xiii

Contributor Contact Details

xxi

PART I. General 1. Biomaterials and Mechanical Stimulation in Tissue Engineering

3

Tetsuya Tateishi, Guoping Chen, Naoki Kawazoe, Makoto Kawanishi, and Takashi Ushida

2. Osteoinduction and Ostogenic Genes Expression Regulated by Ca-P Bioceramics

24

Xingdong Zhang

3. Why is the Clinical Application of Tissue Engineering So Slow?

35

Yoshito Ikada

4. Temperature-responsive Cell Culture Surfaces for Cell Sheet Tissue Engineering Zhonglan Tang, Yoshikatsu Akiyama, and Teruo Okano

55

PART II. Tissue Engineering 5. Adipose Tissue-derived Mesenchymal Stem Cells have Lower Osteogenic Potential than Bone Marrow-derived Mesenchymal Stem Cells

73

Yoshihiro Katsube, Ousuke Hayashi, Motohiro Hirose, and Hajime Ohgushi

6. Amelogenin Overexpression in Tooth Development

87

Akiyoshi Taniguchi and Liming Xu

7. Mechano-active Tissue Engineering

98

Sang-Heon Kim, Youngmee Jung, Soo Hyun Kim, and Young Ha Kim

8. Biomaterials for In Vitro Expansion of Stem Cells

116

Yoshihiro Ito

PART III. Polymeric Biomaterials 9. Design of Artificial Extracellular Matrices and Their Application for Regenerative Medicine

143

Xiaoshan Yue, Masato Nagaoka, and Akaike Toshihiro

10. Nano-to-macro Architectures Polycaprolactone-based Biomaterials in Tissue Engineering

159

Teoh Swee Hin, Bina Rai, Tiaw Kay Siang, Chong Seow Khoon Mark, Zhang Zhiyong, and Teo Yiling Erin

11. Nanofiber-based Scaffolds for Tissue Engineering Hisatoshi Kobayashi, Yoshiro Yokoyama, Chiaki Yoshikawa, Satoshi Igarashi, Shinya Hattori, Takako Honda, Hiroyuki Koyama, and Tsuyoshi Takato

182

12. Electrospun Composite Nanofibrous Scaffolds for Tissue Engineering

194

Inn-Kyu Kang and Jung Chul Kim

13. Synthetic/Natural Hybrid Scaffold for Cartilage and Disc Regenerations

207

Gilson Khang, Soon Hee Kim, John M. Rhee, Munirah Sha’Ban, and Ruszymah Bt Hj Idrus

14. Novel Hydrogel Systems as Injectable Scaffolds for Tissue Engineering

222

Yoon Ki Joung and Ki Dong Park

15. Manipulation of Stem Cell Functions on Grafted Polymer Surfaces

234

Naoki Kawazoe, Likun Guo, Guoping Chen, and Tetsuya Tateishi

16. Gamma-ray Irradiated Poly(L-lactide) for Bone Repair

254

Kazuo Isama and Toshie Tsuchiya

PART IV. Metallic Biomaterials 17. Titanium Alloys with High Biological and Mechanical Biocompatibility

269

Mitsuo Niinomi

18. Biofunctionalization of Metals

292

Takao Hanawa, Yuta Tanaka, and Harumi Tsutsumi

19. Research on Biological Characteristic of Silver Nanoparticle Jinglong Tang, Ling Xiong, and Tingfei Xi

303

PART V. Bioceramics 20. Hydrothermal Synthesis of Hydroxyapatite Ceramics for Medical Application

317

Koji Ioku and Masanobu Kamitakahara

21. Bone Regeneration Materials Based on Calcium Phosphate Ceramics

327

Masanori Kikuchi, Yasushi Suetsugu, Yoshihisa Koyama, Shinichi Sotome, Soichiro Itoh, Kazuo Takakuda, Kenichi Shinomiya, Kazuya Edamura, Katsuyoshi Nagaoka, and Shigeo Tanaka

22. Silicate-based Bioactive Materials for Bone Regeneration

343

Jiang Chang and Lei Chen

23. Biomimetics: Bio-inspired Engineering of Human Tissue Scaffolding for Regenerative Medicine

364

David W Green and Besim Ben-Nissan

PART VI. Drug Delivery System 24. Nanobiomaterials Taking Aim at Drug and Gene Delivery

389

Zhongwei Gu and Yujiang Fan

25. Delivery Technology of Growth Factors to Realize Tissue Regeneration Therapy

402

Yasuhiko Tabata

26. Biostable Gradient Coating Implants with Drug Release

416

Shinn-Jyh Ding, Chun-Cheng Chen, and Chia-Che Ho

27. Design of Supramolecular Polyrotaxanes for DNA Delivery Atsushi Yamashita and Nobuhiko Yui

431

PART VII. Biomechanics 28. Biomechanical Aspects of Natural Articular Cartilage and Regenerated Cartilage

445

Teruo Murakami, Nobuo Sakai, Yoshinori Sawae, Itaru Ishikawa, Natsuko Hosoda, Emiko Suzuki, and Jun Honda

29. Wear Characteristics of a Monopivot Centrifugal Blood Pump for Circulatory Support

457

Takashi Yamane, Katsunobu Nonaka, Osamu Maruyama, Masahiro Nishida, Ryo Kosaka, Yoshiyuki Sankai, and Tatsuo Tsutsui

PART VIII. Evaluation and Standardization 30. In Vitro Biodegradation of Poly(Lactic-co-Glycolic Acid) Porous Scaffolds

467

Guoping Chen, Taiyo Yoshioka, Naoki Kawazoe, and Tetsuya Tateishi

31. Non-invasive Evaluation Technique for Cartilage Tissue Engineering

482

Shogo Miyata, Kazuhiro Homma, Tomokazu Numano, Takashi Ushida, and Tetsuya Tateishi

32. Analytical TEM Study of Biomineral Phases

494

Yang Leng and Renlong Xin

33. High-throughput Cytometry Using Antibody Arrays

506

Koichi Kato, In-Kap Ko, Toshinari Ishimuro, Mitsuaki Toda, Yusuke Arima, Isao Hirata, and Hiroo Iwata

34. Estimation of Endurance Performance of Hip Stems Using Finite Element Analysis Yu-Bong Kang, Duk-Young Jung, and Sadami Tsutsumi

517

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Contributor Contact Details Editor Tetsuya Tateishi Biomaterials Center, National Institute for Materials Science, 1-1 Namiki, Tsukuba, Ibaraki 305-0044, Japan Email: [email protected] Chapter 1 Tetsuya Tateishi Biomaterials Center, National Institute for Materials Science, 1-1 Namiki, Tsukuba, Ibaraki 305-0044, Japan Email: [email protected] Chapter 2 Xingdong Zhang National Engineering Research Center for Biomaterials, Sichuan University, 29 Wangjiang Road, Chengdu 610064, China Email: [email protected] Chapter 3 Yoshito Ikada Nara Medical University, Nara, Japan Email: [email protected]

Chapter 4 Teruo Okano Institute of Advanced Biomedical Engineering and Science, Tokyo Women’s Medical University, Japan Email: [email protected] Chapter 5 Hajime Ohgushi Research Institute for Cell Engineering, National Institute of Advanced Industrial Science and Technology, Amagasaki, Japan Email: [email protected] Chapter 6 Akiyoshi Taniguchi Biomaterials Center, National Institute for Materials Science, 1-1 Namiki, Tsukuba, Ibaraki 305-0044, Japan Email: [email protected] Chapter 7 Young Ha Kim Department of Materials Science & Engineering, Gwangju Institute of Science and Technology, Gwangju, Korea Email: [email protected] Chapter 8 Yoshihiro Ito Nano Medical Engineering Laboratory, RIKEN, Saitama, Japan Email: [email protected]

Chapter 9 Akaike Toshihiro Department of Biomolecular Engineering, Graduate School of Bioscience and Biotechnology, Tokyo Institute of Technology, 4259-B-57 Nagatsuta-cho, Midori-ku, Yokohama 226-8501, Japan Email: [email protected] Chapter 10 Swee-Hin Teoh Centre for Biomedical Materials Applications and Technology (BIOMAT), Department of Mechanical Engineering, National University of Singapore, 9 Engineering Drive 1 Singapore 117576 Email: [email protected] Chapter 11 Hisatoshi Kobayashi Biomaterials Center, National Institute for Materials Science, Tsukuba, Japan Email: [email protected] Chapter 12 Inn-Kyu Kang Department of Polymer Science, Kyungpook National University, Taegu 702-701, Korea Email: [email protected]

Chapter 13 Gilson Khang Department of Polymer.Nano Science and Technology, Chonbuk National University, 664-14, Dukjin Dong 1 Ga, Dukjin Ku, Jeonju 561-756, Korea Email: [email protected] Chapter 14 Ki Dong Park Department of Molecular Science and Technology, Ajou University, Suwon, Korea Email: [email protected] Chapter 15 Naoki Kawazoe Biomaterials Center, National Institute for Materials Science, 1-1 Namiki, Tsukuba, Ibaraki 305-0044, Japan Email: [email protected] Chapter 16 Kazuo Isama National Institute of Health Sciences, Tokyo, Japan Email: [email protected] Chapter 17 Mitsuo Niinomi Institute for Materials Research, Tohoku University, Sendai 980-8577, Japan Email: [email protected]

Chapter 18 Takao Hanawa Institute of Biomaterials and Bioengineering, Tokyo Medical and Dental University, 2-3-10 Kanda-Surugadai, Chiyoda-ku, Tokyo 101-0062, Japan Email: [email protected] Chapter 19 Tingfei Xi Center of Medical devices, National Institute for the Control of Pharmaceutical & Biological Products, Beijing, China Email: [email protected] Chapter 20 Koji Ioku and Masanobu Kamitakahara Graduate School of Environmental Studies, Tohoku University, Sendai, Japan Email: [email protected] Chapter 21 Masanori Kikuchi Biomaterials Center, National Institute for Materials Science, Tsukuba, Ibaraki, Japan Email: [email protected] Chapter 22 Jiang Chang Biomaterials and Tissue Engineering Research Center, Shanghai Institute of Ceramics, Chinese Academy of Sciences, Shanghai, 200050, China Email: [email protected]

Chapter 23 Besim Ben-Nissan Department of Physics and Advanced Materials, University of Technology, Sydney, Australia Email: [email protected] Chapter 24 Zhongwei Gu National Engineering Research Center for Biomaterials, Sichuan University, 29 Wangjiang Road, Chengdu 610064, China

Email: [email protected] Chapter 25 Yasuhiko Tabata Department of Biomaterials, Field of Tissue Engineering, Institute for Frontier Medical Sciences, Kyoto University, Kyoto, Japan Email: [email protected] Chapter 26 Shinn-Jyh Ding Institute of Oral Materials Science, Chung-Shan Medical University, Taichung 402, Taiwan, Republic of China Email: [email protected] Chapter 27 Nobuhiko Yui School of Materials Science, Japan Advanced Institute of Science and Technology, Ishikawa, Japan Email: [email protected]

Chapter 28 Teruo Murakami, Department of Mechanical Engineering, Kyushu University, Fukuoka, Japan Email: [email protected] Chapter 29 Takashi Yamane National Institute of Advanced Industrial Science and Technology, Tsukuba, Japan Email: [email protected] Chapter 30 Guoping Chen Biomaterials Center, National Institute for Materials Science, 1-1 Namiki, Tsukuba, Ibaraki 305-0044, Japan Email: [email protected] Chapter 31 Shogo Miyata Faculty of Science and Engineering, Keio University, Yokohama, Japan Email: [email protected] Chapter 32 Yang Leng Department of Mechanical Engineering, Hong Kong University of Science and Technology, Hong Kong, China Email: [email protected]

Chapter 33 Hiroo Iwata Institute for Frontier Medical Sciences, Kyoto University, Kyoto, Japan Email: [email protected] Chapter 34 Sadami Tsutsumi School of Dentistry, Nihon University, Tokyo, Japan Email: [email protected]

PART I

General

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Chapter 1 Biomaterials and Mechanical Stimulation in Tissue Engineering Tetsuya Tateishi1, Guoping Chen1, Naoki Kawazoe1, Makoto Kawanishi2 and Takashi Ushida2 1. Biomaterials Center, National Institute for Materials Science, Tsukuba, Japan 2. Division of Biomedical Materials and Systems Center for Disease Biology and Integrative Medicine, School of Medicine, The University of Tokyo7-3-1 Hongo, 113-0033 Tokyo, Japan

1

Introduction

As a rapidly emerging technology, tissue engineering holds the potential of a new approach to the repair and reconstruction of tissue and organs damaged by disease and accident. There are several factors required for tissue engineering. They include cell source, growth factor, biodegradable porous scaffolds and mechanical stress (Fig. 1). Biodegradable porous scaffolds play an important role in tissue engineering of cartilage as a physical support and also as an adhesive substrate for the isolated chondrocytes. Ideally, the scaffolds used in this application should meet several design criteria. For instance, they should permit cell adhesion, promote cell growth and allow retention of differentiated cell function, biocompatibility, biodegradability, high porosity, mechanical strength, and also malleable into desired shapes. Generally, three-dimensional biodegradable porous scaffolds can be fabricated from two kinds of biodegradable polymers. One is synthetic and the other is naturally derived polymer.

Figure 1. Seven factors for tissue engineering.

In the tissue engineering, the biological system is regarded as a multi-functional molecular machine, and it is the target of the tissue engineering to establish the principle for design and manufacture of artificial system simulating the cell and tissue structures. The ultimate goal of the application includes the development of bio-compatible substitutes of tissues and organs, such as artificial bone, muscle and blood vessel. The biological system has various functions superior to those of artificial materials and machines created by the human. For instance, the biological system is equipped with self-organizing capability to build up part of its body by itself. When it suffers from a damage, it can repair and reinforce through metabolism and controlled growth. Moreover, these functions can be implemented autonomously at various levels of hierarchical organization, such as tissues, cells and molecular assemblies, at normal temperatures and pressures through the reactions friendly to the environment. If the principle for design and manufacture of artificial system simulating the biological system became available, it would be possible to create artificial organs of improved bio-compatibility and retarded functional aging. In case of artificial bone materials, the bio-compatibility has been improved at a great stride, but still far from perfect. An artificial bone surgically implanted at an age of 60 was required to last up to 70, but owing to the extension of longevity, it has to remain intact up to 80. The long-term effects of metal ions dissolving out of

implanted materials are not known clearly. In the 'harsh' environment within the body, damages such as abrasion and corrosion may accumulate resulting in functional failure. These troubles are caused by the use of non-biological materials in the artificial organs. The reason for advancing the studies of tissue engineering is to create the truly bio-compatible devices by using living human cells. These efforts will be linked to environment-friendly science, technology and industrial production. There will be three parts in this paper. At first, we would like to introduce a method for the scaffold preparation developed by my group by using ice particulates as porogen materials. Secondly, we would like to talk about some new composite scaffolds developed by our group. And thirdly, we would show you some results of articular cartilage engineering by using the composite scaffolds. The porous three-dimensional temporary scaffolds play an important role as a physical support and an adhesive substrate for the isolated cells guiding the new tissue growth and organization.

2

Biomaterials for Scaffold [1-9]

Except β tricalcium phosphate (TCP) and hydroxyapatite, almost all the biodegradable scaffolds used in tissue engineering have been made from biodegradable polymers. There are two kinds of biodegradable polymer materials: synthetic polymer and naturally derived polymers. Among these biodegradable polymers, the synthetic polyesters such as poly(glycolic acid) (PGA), poly(lactic acid) (PLA), and their copolymer of poly(lactic-co-glycolic acid) (PLGA) and the naturally derived polymer collagen have been most attractive and most frequently used for tissue engineering. To prepare porous scaffolds of these polymer materials, several methods including phase separation, emulsion freeze-drying, gas foaming, fiber bonding and porogen leaching techniques have been developed. These methods have their respective advantages and also drawbacks. Phase separation technique has the difficulty to control pore structure. Emulsion freeze-drying technique and expansion technique often result in a closed cellular structure in the scaffold. Scaffolds prepared by the fiber bonding technique is devoid of structural stability. Compared to these methods, the porogen leaching method provides easy control of pore structure and has been well established in the preparation of

3-dimensional scaffolds for tissue engineering. However, this method needs airdrying at room temperature, which makes it unsuitable for protein incorporation. However proteins such as growth factors and cytokines are very important in promoting cell proliferation and differentiation into specific tissues. We have developed a modified porogen-leaching method by using preprepared ice particulates as the porogen material. At first, ice particulates were prepared by spraying cold water into liquid nitrogen. By controlling the spraying condition, ice particulates of 240 µm mean diameter were prepared. Ice particulates were then mixed with pre-cooled polymer solution (PLLA or PLGA in chloroform), the dispersion was mixed and freeze-dried, then polymer (PLLA or PLGA) foams were formed. The SEM photomicrographs of the cross-sections of the PLLA foams prepared with 80% and 90% ice particulates shows that the foams were highly porous with evenly distributed and interconnected pore structures. The pore shapes were almost the same as those of ice particulates. The foams became more interconnected as the weight fraction of ice particulates increased. Therefore, the pore structures of the porous scaffolds could be manipulated by varying the weight fraction, size of ice particulates (and polymer concentration). Because all the process was conducted at low temperature, this method would be suitable for protein incorporation. The scaffolds prepared by this method should be useful for tissue engineering.

3

Biodegradable Polymer Composite Sponge

Synthetic polymers such as PLA, PGA and PLGA, and naturally derived polymers have their respective advantages and drawbacks. The synthetic polymers can be easily processed into desired shapes with good mechanical strength and their degradation periods can also be manipulated. Despite these advantages, PGA, PLA and PLGA-derived scaffolds lack cell-recognition signals, and their hydrophobic property hinders successful cell seeding. On the other hand, naturally derived biodegradable polymer, such as collagen, offers the advantage of specific cell interaction and hydrophilicity, but scaffolds constructed entirely of collagen have poor mechanical strength and not easy to handle. So we thought of hybridizing these two kinds of biodegradable polymers to develop a novel kind of composite scaffold which will combine their advantages, but without the drawbacks.

The composite scaffolds were prepared as follows. At first, synthetic polymer sponge or mesh were prepared. The mechanically strong synthetic polymer sponge or mesh serves as a skeleton. And then we introduced the mechanically weak but highly cell compatible collagen microsponges were introduced into the pores of the sponge or interstices of the mesh to form the composite sponge or composite mesh (Fig. 2).

Figure 2. Hybridization of biodegradable synthetic polymers with collagen.

The composite structure of PLGA-collagen composite sponge was further confirmed by SEM- EPMA analysis. Collagen microsponges were formed in the pores of PLGA sponge and that the pore surfaces were also coated with collagen (Fig. 3). The wettability, mechanical strength and cell-interaction were investigated to confirm if the composite sponge combined the advantages of both PLGA and collagen. The wettability of a scaffold is considered to be very important for successful cell seeding into a scaffold. The contact angle of PLGA sponge with water was about 76 degrees, which indicated that the PLGA sponge was relatively hydrophobic. To seed cells in PLGA sponge, it is necessary to pretreat the scaffold by prewetting or surface hydrolysis. However, after hybridization with collagen, the contact angle decreased to about 31 degrees, thus the wettability increased. Cell seeding becomes very easy. The PLGA-

collagen composite sponge showed higher dynamic, static compression and tensile strength than PLGA and collagen sponges both in dry and wet states. When used for culture of mouse fibroblast cells, more cells adhered to the composite sponge than PLGA sponge. The cells adhered to the collagen microsponges and distributed throughout the composite sponge.

Figure 3. SEM photomicrographs of cross section of PLGA sponge and PLGA-collagen composite sponge.

After 5 days in culture, they spread and proliferated to cover all the available surfaces of the sponge. This result suggests good cell interaction of the PLGAcollagen composite sponge. From these results, we can conclude that the synthetic PLGA sponge serving as a skeleton, facilitated formation of the composite sponge into the desired shapes and reinforced the composite sponge, while collagen microsponges contributed good cell interaction and hydrophilicity, and therefore easy cell seeding.

4

Biodegradable Polymer Composite Sponge

The hybridization was achieved by forming collagen sponges between the interstices of synthetic biodegradable polymer mesh. In the case of the composite

mesh of a PLGA knitted mesh and collagen, collagen microsponges with interconnected microporous structures were formed in the interstices of the synthetic polymer mesh. The fiber bundles of polymer mesh and the collagen sponges were alternately chained (Fig. 4). The composite mesh possessed similar mechanical property to that of the polymer mesh, 500 times higher than that of the collagen sponge alone.

Figure 4. SEM photomicrographs of PLGA-collagen hybrid mesh.

When used for cell culture, more human skin fibroblasts adhered and proliferated quicker on the composite mesh than on the PLGA mesh. The cells adhered and spread well on the surfaces of the collagen microsponge of the composite mesh after being cultured for 5 days, and became layer structured after 2 weeks. These results suggest again that the composite mesh combined the advantages of both synthetic polymer mesh and collagen sponge.

5

Composite Sponge of Synthetic Polymer, Collagen and Hydroxyapatite

As collagen and hydroxyapatite are the two primary components of bone extracellular matrix and has been reported showing good osteoconductivity. We

hybridized synthetic polymer with collagen and hydroxyapatite to prepared a porous composite scaffold for tissue engineering of hard tissues.The sponge was prepared by depositing hydroxyapatite particulates on the collagen microsponge surfaces of synthetic polymer-collagen composite. The deposition of hydroxyapatite particulates was accomplished by alternate immersion of the PLGA-collagen sponge in CaCl2 and Na2HPO4 aqueous solutions and centrifugation. SEM photomicrographs show that hydroxyapatite particulates were formed on the surfaces of the collagen microsponges of the PLGA-collagen sponge. After one immersion cycle, the deposited particulates were scarce and small. They became denser and grew larger as the number of alternate immersion cycles increased. The surfaces of the collagen microsponges were completely covered with hydroxyapatite particulates after three cycles of alternate immersion (Fig. 5). And their level of crystallinity increased as the number of immersion cycles increased. Such synthetic polymer-collagen-hydroxyapatite composite sponges would prove useful as three-dimensional porous scaffolds for bone tissue engineering.

Figure 5. SEM photomicrographs of cross sections of PLGA-collagen-hydroxyapatite composite sponge after 3 and 7 cycles immersion.

6

Tissue Engineered Articular Cartilage

Bovine articular chondrocytes were isolated from the shoulder articular cartilage of a four-week old calf by collagenase digestion and subcultured in DMEM containing 10% FBS to obtain sufficient numbers of cells. The subcultured chondrocytes were collected and seeded onto the PLGA-collagen composite mesh, and cultured in vitro (Fig. 6). The chondrocytes adhered onto the composite mesh and continued to proliferate filling the spaces in the composite mesh.The cell/mesh constructs were overlapped and sutured together to control the thickness of implants (Fig. 7).

Figure 6. Tissue engineering of cartilage using PLGA-collagen composite mesh.

After one week in culture, the chondrocyte/scaffold constructs were implanted subcutaneously in the dorsum of nude mice. Here shows the nude mice immediately after implantation and after 8 weeks implantation. The implants were harvested after 2, 4, 8, and 12 weeks. Here shows the gross appearance of the implants after 8 and 12 weeks (Fig. 8). All the implants preserved their original shapes for all implantation periods and appeared glistening white. To see whether the engineered cartilage is articular hyaline cartilage or fibrocartilage, we examined implant samples by HE staining, and safranin O/fast green staining, immunohistological staining and gene expression study. We also measured the mechanical properties of the engineered cartilage.

Figure 7. SEM photomicrographs of extra cellular matrix of chondrocytes.

The histology of 2 and 4 weeks implants shows that the chondrocytes showed their natural round morphology. The PLGA fibers could still be seen after four weeks.GAGs were detected in the extracellular matrices. The histology of 8 and 12 weeks implants shows that the chondrocytes showed their natural lucuna structure, and were evenly distributed throughtout the scaffold for all the implants. The PLGA fibers gradually degraded and disappeared after 12 weeks. Homogeneous safranion O-positive GAGs were detected for all the implants Immunohistological staining with an antibody to type II collagen showed a homogeneous extracellular staining for type II collagen. Gene expression of articular cartilageneous matrices of collagen type II and aggrecan was detected and upregulated when the chondrocytes were cultured on the composite mesh. From these results, it can be concluded that the engineered cartilages are articular cartilage-like tissue, and the composite mesh can serve as an effective biodegradable scaffold for cartilage tissue engineering. We obtained good results concerning the mechanical properties of the engineered cartilage implants. The mechanical properties of the engineered implant after 12 weeks and bovine native articular cartilage were evaluated by a

dynamic compression mechanical test using a viscoelastic spectrometer. The dynamic complex modulus (E*) reached 37.8%, structural stiffness 57.0% and phase lag (tanδ) 86.3% of those of native bovine articular cartilage, respectively. Very high mechanical strength of engineered cartilage as we know. (The results measured at 35 Hz showed a similar trend to those measured at 11 Hz).

Figure 8. Gross appearance of tissue engineered articular cartilage.

7

Gas-controlled Cyclic Hydrostatic Pressure System for Bioreactor

Cells and tissues are known to be stimulated mechanically by diverse mechanical loading. For example, femur is loaded with compressive or tensile stress. In blood vessel, endothelial cells are loaded by shear stress with blood flow, smooth muscle cells are loaded by tensile stress with blood pulse. Chondrocytes in articular cartilage are thought to be loaded with hydrostatic pressure caused by compressive loading (Fig. 9). Autologous chondrocyte transplantation (ACT) using tissue engineering technique has undergone rapid development for treating articular cartilage defects10-12. Therefore the range that we can treat is only focal defect. More many cells are necessary to treat a large defect. Cells must be proliferated in vitro by a monolayer culture because of the very limited amount of donor biopsy

cartilage available. However, monolayer culture proliferation leads to dedifferentiation.

Figure 9. Various effects of mechanical stress on cultured cells.

Figure 10. System for cyclic hydrostatic pressure application. (A) Schematic drawing of the system. (B) Sealing of polyacetal cassette containing pellets with culture medium in plyolefin bag.

This is a process during which chondrocytes lose their ability to express articular cartilage-specific extracellular matrices (ECM) such as type II collagen and aggrecan, and produce fibroblast-specific ECM, type I collagen[13, 14]. To be used for clinical applications, dedifferentiated cells must be redifferentiated. Different methods have been tried to induce redifferentiation of dedifferentiated articular chondrocytes to regain their cartilaginous features. Application of growth factors, mechanical forces and creation of an environment supporting a spherical morphology such as in pellets, alginate beads or polymer gels have often been used to fulfill this purpose[15-22]. Hydrostatic pressure is one of the most frequently used mechanical factors for chondrocyte experiment. A variety of hydrostatic pressure loading devices have been reported for the cartilage cell experiment [23-28]. However, a gas exchange system did not exist at the pressure load except low pressure level [27]. But articular cartilage is exposed to intermittent mechanical compression with peak load amplitudes of about 3.5-18 MPa in vivo [29]. Recently Toyoda et al. also reported the difficulty of gas control at such pressure load [30]. Carver’s system [31] and our previous system [32] had medium circulation system not only hydrostatic pressure loading. Furthermore our system had a risk of contamination in long term culture. So we used the polyolefin bag from which the gas infiltration was confirmed. The bag is already used for a selfblood save and a blood transfusion in a medical field. And it was also confirmed to prevent bacteria and the virus. Medium exchange is possible in a part of a nozzle, too. And we used a polyacetal cassette so that each pellet was not contact together. This cassette is also already used for a clinical biopsy. New system was constructed for the application of hydrostatic pressure to the pellet cultured in bags. The system was composed of a cylinder pump for high performance liquid chromatography (SSC-3471; Senshu Scientific co., ltd.), a gas exchange tray, a pressure sensor (Yokohama system laboratory), a PC controlled air valve, and a pressure chamber (Fig. 10A). After 3 days tube culture, 6 pellets were put in each cassette made of a polyacetal (Biopsy six embedding cassettes; Sakura Finetek Japan co., ltd.), and then stuffed in a polyolefin bag (KBP-1000CP; Kawasumi Laboratories, inc.) (Fig. 10B). The bag was filled with serum free medium (15 ml). These 8 bags were put in a pressure chamber and filled with the RO water that had finished a gas exchange in an incubator 5% CO2 atmosphere at 37˚C beforehand. The pressure chamber

was placed in an incubator to maintain culture temperatures at 37 °C. Cyclic hydrostatic pressure was applied at magnitudes of 5 MPa at 0.5 Hz for durations of 4 hours per day for four days. The magnitude and frequency of the applied pressure were chosen to be within the physiological range normally occurring in the hip joint during daily activity [29]. Frequency and magnitude of the hydrostatic pressure were manipulated by the PC controlled air valve and flowing quantity of the pump. A pressure sensor was installed in the line to continuously monitor the internal pressure of the chamber. An identical chamber placed in the same incubator was used for the control pellets, but no loading was applied. Chondrocytes were extracted from the bovine normal knee joint cartilage. After 3 passages, dedifferentiated chondrocytes were applied to form a pellet (2.5 × 105 cells/0.5 ml). These pellets were cultured in chemically defined serum free medium with ITS+Premix for 3 days. Because a serum had a many factors that influence chondrocyte metabolism, we had to remove a serum. Especially thyroxine is one of a hormone in a serum. This hormone has wideranging effect, including involvement in the processes of hypertorophy and endochondral ossification at the growth plate [36]. This effect is not desirable for the chondrogenic differentiation. So, we used a serum only at dedifferentiation phase. Our serum free medium was a modification of the protocol outlined by Mackay et al [33]. The main change point was to have excluded dexamethasone and TGF-β. These two elements have been reported to affect chondrogenic differentiation in vitro [34, 35]. Semiquantitative reverse transcriptase–polymerase chain reaction showed a 5-fold increase in levels of aggrecan mRNA by cyclic hydrostatic pressure load (p < 0.01) (Fig. 11A, B). Type II collagen mRNA levels also upregulated to 4fold by cyclic hydrostatic pressure load (p < 0.01) (Fig. 11A, C). Type I collagen mRNA levels similarly reduced by cyclic hydrostatic pressure load and control groups (Fig. 11A, D). Type I collagen mRNA levels seemed to have decreased by three-dimensional pellet culture only. And redifferentiation of dedifferentiated bovine articular chondrocytes was enhanced by cyclic hydrostatic pressure load. Partial oxygen pressure (PO2) and partial carbon dioxide pressure (PCO2) of the medium in the bag reached equilibrium in 24 hours, and no eminent change

Figure 11. Effect of cyclic hydrostatic pressure on aggrecan, type I collagen, and type II collagen mRNA levels. (A) Ethidium bromide-stained agarose gel analysis of RT-PCR samples in each group. P3 means after passage three chondrocytes. Ratio of aggrecan (B), type II collagen (C), and type I collagen (D) to GAPDH PCR product normalized to the control mean in each group. Values are the mean ± SEM (n = 4). * p < 0.01.

Figure 12. Gas analysis of culture medium and RO water. Values are the mean ± SEM (n = 8).

was observed for three days afterwards. PO2 and PCO2 excellently controlled (Fig. 12).

Figure 13. Production of extracellular matrix. (A) The safraninO/fast green staining for glycosaminoglycan of the loaded pellets showed good stains than control pellets. (B) Sulphated proteoglycans were detected by Alcian-blue stainng. Metachromatic staining was strongly found in the loaded pellets than control pellets. Three magnifications (×10, 20 and 40) at each staining were used for microscopic observation. Scale bar, 200 mm.

In histological analysis, Safranin O/fast green staining for GAG of the loaded pellets showed good stains than control pellets (Fig. 13A). In the pellets, sulphated proteoglycans were detected by Alcian-blue stainng. Metachromatic staining was strongly found in the loaded pellets than control pellets (Fig. 13B). ECM expression of loaded pellets was higher than control pellet. These results suggest that gas-controlled cyclic hydrostatic pressure enhance the cartilaginous matrix formation of dedifferentiated cells differentiated in vitro. Sensitivity of chondrocytes to hydrostatic pressure is well established [24, 25, 28, 32, 37 ]. However there was no report of hydrostatic pressure effect on three-dimensional bodies which made only in dedifferentiated articular chondrocytes. Our results suggest that gas-controlled cyclic hydrostatic pressure within physiological range enhance the cartilaginous matrix formation of dedifferentiated cells differentiated in vitro.

In the field of cell-based therapy, ACT is most used for joint cartilage defect. In ACT, culture-expanded chondrocytes are transplanted under a cover of periosteum [10]. But culture-expanded cells lead dedifferentiation. Therefore the range that we can treat is only focal defect. More many cells are necessary to treat a large defect. This is one of the problems of ACT. Tallheden et al. reported that osteoarthritis chondrocytes have a good proliferation potential and are able to redifferentiate in a three-dimensional pellet model [38]. P. Angele et al. showed that cyclic hydrostatic pressure enhances the cartilaginous matrix formation of mesenchymal progenitor cells differentiated in vitro [39]. These kinds of the cell sources and our gas-controlled cyclic hydrostatic pressure loading system may solve this problem of ACT bone tissue engineering.

Figure 14. Outlook of tissue engineering using novel scaffolds.

8

Conclusions

A new method for the preparation of biodegradable porous scaffolds has been developed by using pre-prepared ice particulates as porogen material. A novel kind of composite biodegradable porous scaffold has been developed by forming collagen microsponges in the pores or interstices of a synthetic polymer sponge

or mesh. A composite sponge of synthetic polymer, collagen and hydroxyapatite has been developed for hard tissue engineering. Bovine articular cartilage-like tissue has been engineered by culturing chondrocytes in the PLGA-collagen scaffolds. The application of hydrostatic pressure with exchanged gas flow can be used for in vitro engineered tissue for clinical application. Enhancement of in vitro matrix synthesis by physical factors such as hydrostatic pressure and medium flow may realize in vitro reconstruction of cartilage and with requisite properties for repair of cartilaginous defect. Finally, we would like to say the scaffolds developed by our group can not only used for cartilage engineering, also usable for tissue engineering of other tissues and organs such as bone, skin, tendon, ligament, muscle, nerve, blood vessels, and etc. (Fig. 14).

Acknowledgments This work was supported in part by grants from the New Energy and Industrial Technology Development Organization.

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Chapter 2 Osteoinduction and Ostogenic Genes Expression Regulated by Ca-P Bioceramics Xingdong Zhang National Engineering Researcher Center for Biomaterials, Sichuan University, China

1

Introduction

Previous reports have shown that certain porous Ca-P bioceramics without any additional growth factor or living cell could induce bone formation, and the osteoinduction was material- and biological-dependent [1–4]. But the understanding of mechanism for biomaterial osteoinduction is not clear. It is known that a normal bone formation undergoes a series process, including cells, cell proliferation and differentiation and bone matrix secretion, which are regulated by a cascade of initiated gene activations in the cells under the guidance of environmental influences and molecular signals. Obviously, in osteogenesis integrated with biomaterials, the material factors are very important environmental influences. In order to explore the osteoinductive mechanism of Ca-P bioceramics, it has to be understood whether materials are involved in the initiation and regulation of a cascade of gene activities and material factors are the keys for the initiation and regulation processes.

2.1 The Molecular Events in Osteogenic Process There is a series of molecular events occurs in the process of the normal bone formation which regulates cells to proliferate and differentiate along the osteogenic pathway. The formation of bone that is required to integrate with biomaterials has to have the existence of undifferentiated mesenchymal stem cells (MSCs) with the ability to proliferate and differentiate along osteogenic pathway in osseous sites. It has been proved that in vivo fibro-connective tissues could grow in the pores of Ca-P bioceramics and provide undifferentiated MSCs after the ceramics were ectopically implanted in animals. But the undifferentiated MSCs coming from the fibro-connective tissues mainly are multi-potent MSCs which may differentiate into myogenic, adipogenic, fibroblastic, as well as chondrogenic and osteogenic cells dependent on environmental influences and signal molecules. It has been recognized that bone morphological proteins (BMPs) are very important osteoinductive signal molecules. Under the effects of BMPs, the multipotents MSCs could become bipotent osteogenic precursor cells which may differentiate into chondrogenic cells and osteoblasts. The differentiation pathway of the precursor cells is controlled by “master genes”. In osteogenic differentiation the core binding factor 1 (cbfa1) and osteoblast specific factor (osf2) are important regulatory genes (cbfa1/osf2), which regulate the osteogenic precursor cells to differentiate along osteogenic pathway. The cells will differentiate along chondrogenic pathway if no-expressions of cbfa1/osf2 in the MSCs. The cbfa1/osf2 in the undifferentiated cells is upregulated due to combined effects of BMPs and other factors in osteogenic differentiation. Further progress of the osteogenic differentiation is characterized by the expressions of bone-matrix proteins which are thought as osteogenic markers. The bone-matrix proteins mainly include collagen and non-collagenous specific proteins, the latter are alkaline phosphate (ALP), osteopontin (OPN), bone sailoprotein (BSP), osteocalcin (OCN), osteonectin et al. The proteins are produced by osteoblast and regulate the formation and growth of bone-mineral crystals of hydroxapatite. They are orderly expressed in preosteoblasts and osteoblasts under the effects of a cascade of activated genes, including cbfa1/osf2, TGF-β and so on. This paper mainly discusses the effects of material factors of Ca-P bioceramics to the expressions of BMP, cbfa1/osf2 and bone-matrix proteins in cells. The results show that in osteoinductive process of Ca-P bioceramics the material factors have key effects to activate and regulate the expression of genes under the synergism of biological environment [5].

2.2 The Enrichment of Materials to Endogenous Signal Molecules There are many osteoinductive signal molecules, for examples, BMPs, TGF-βs, PDGF, IGF-1, et al. But the in vivo concentration is very low which does not reach the threshold of the initiation of MSCs differentiation. The osteoinductive molecules are some proteins with low molecular weights. In vivo tests show that Ca-P bioceramics may be selectively absorbed in vivo proteins. The absorption ability and types absorbed proteins vary with different Ca-P ceramics. Figure 1 was the patterns of absorbed proteins by porous HA and TCP ceramics. The amount of absorbed proteins in HA ceramics is more than that in TCP ceramics, but the absorption for the proteins with low molecular weights is TCP＞HA[6].

(a)

(b)

Figure 1. Silver-strained PAGE profiles of proteins absorbed (a) in vitro and (b) in vivo. (a) Porous HA and TCP ceramics were incubated in fetal calf serum at 37℃ and 5% CO2 for one week; (b) Porous HA and TCP ceramics were put in diffusion chambers which were implanted in subcuts of wister rats for one week.

2.2.1 BMP BMPs are the most important osteoinductive molecules for the early osteogenic differentiation. BMPs act on multi-potent MSCs and lead the MSCs to become bi-potent MSCs. Thus, the formation of both bone and cartilage are promoted. In osteogenesis, BMPs upregulate the expression of cbfa1/osf2 in the bi-potent MSCs, thus regulate the cells to differentiate into osteoblasts. Fig. 2 is the immuno-histochemical staining sections of porous TCP/HA (BCP) ceramics ectopically implanted in pigs and dogs separately. The sections showed that the BCP could enrich intrinsic BMPs to the walls of ceramic pores

and stimulate the fibro-connective tissue cells ingrowth in the pores to secret BMPs[7].

BCP, in muscle of pig, 30d, BMP-2 enriched on the ceramic pore walls.

BCP in muscle of dog 15d. Positive reaction near capillary tissues.

Figure 2. Histological sections by immuno-histochemical staining.

The inhibition tests of BCP to enrich bBMP by the antibody of bBMP appeared that the enriched intrinsic bBMP was necessary in osteogenic diffentiation[8]. The expression of receptor of γBMP2 (γBMPR2) indicates that there are about 16.6% polyangular cells in the pores implanted ceramics existed γBMPR2 [Fig. 3], which means that a part of the fibro-connective tissue cells grew into the ceramic pores may accept γBMP2, thus to differentiate into osteogenic lineage.

(a)

(b)

Figure 3. γBMPR2 was expressed in some polymorphous cells ingrowth in the pores of BCP ceramics implanted in the muscles of dogs for 15 (a) and 30 (b) days.

The analyses of TEM super-structure and function of the cells showed that the MSCs may differentiate into preosteoblast and osteoblasts (Fig. 4).

*active cells; + endotheliocytes; ◊ HA; RER rough endoplasmic reticulum; col, collagen fibils .

Figure 4. Analysis of TEM super-structure and function of the cells.

In vitro absorption test of Ca-P ceramics to BMPs appeared that the amount of absorbed BMPs varied with the composition of the ceramics and the amount of absorbed BMPs is HA＞TCP/HA＞TCP. But the ability of the ceramics to induce bone formation was TCP/HA＞HA＞TCP [Fig. 5]. This implied that there were other osteoinductive molecules in the osteoinductive process.

Relation of the adsorbed bBMPs

Schematic of Ca-P ceramics to adsorb bBMPs

Figure 5. In vitro adsorption of porous Ca-P ceramics to bBMPs. The Ca-P ceramics have a strong absorption which is material-dependent.

2.2.2 TGF-β TGF-βs also are the important bone growth factors and their mainly effects are to stimulate the expression of bone matrix proteins. In vitro test of Ca-P ceramics to absorb TGF-βs in serum proteins of rats showed that the Ca-P ceramics may absorb in vivo TGF-βs and the amount of absorbed TGF-βs is material-dependent and is BCP＞TCP＞HA [Fig. 6].

and

Serum protein

Figure 6. In vitro absorption of Ca-P ceramics to TGF-βs in the rat serum.

2.3 The Regulation of Ca-P Ceramics to the Expression of Osteogenic “Master Genes” It has been identified for an osteoblastic differentiation that cbfa1/osf2 is ‘master genes’ which regulate bi-potent MSCs to differentiate into osteoblasts and the expression of specific bone matrix proteins associated with other osteogenic factors. The expression of cbfa1/osf2 in undifferentiated MSCs is upregulated by BMPs associated with other factors in which material factors have astrong effects. In situ hybridization studies have detected cbfa1/osf2 mRNA of bone marrow stromal cells cultured in vitro, and found that the cbfal/osf2 mRNA was strongly affected by materials composition and structure (Fig. 7). The expression of cbfa1/osf2 mRNA on BCP was much higher than that on HA [9]. Figure 7 (B) was the expression of osteocalcin(OCN) mRNA in the C2C12 cells, which indicated that the expression of OCN mRNA in C2C12 cells also were stimulated with the upregulation of cbfa1/osf2, and the expression level of OCN mRNA was similar to that of cbfa1/osf2, i.e. in BCP＞＞in HA.

20

18.40

(A)

10

5

6

(B) 5.85

2d 5d

13.88

13.52

fold change

fold change

15

2d 5d

4

3.27 2.38

2 3.59

1.02 1.00

2.53 1.50

0.52

1.03 1.00

0.19

0.20

HT1200

HT1250

0

0 HT1100

HT1200

HT1250

HA1250

HT1100

HA1250

Figure 7. Expressions of Cbfa1 and osteocalcin (OCN) in C2C12 cells cultured on HA and BCP for 2 days and 5 days By real-time RT-PCR analyses.

Figure 8 was the sections of immuno-histochemical staining for cbfa1, which showed that the cbfa1 expression presented on osteoblasts and osteocytes in the porous BCP ceramics implanted in the muscles of pigs for 3 months.

HA/TCP(pig,3M)

Immunohistochemistry

Figure 8. Cbfa1 expressed in osteoblasts and osteocytes.

The test results of the expressions of cbfa1/osf2 and OCN were in accord with the ability of Ca-P ceramics to induce bone formation, i,e the osteoinduction of BCP was stronger than that of HA.

2.4 The Regulation of Ca-P Ceramics to the Expressions of Bone

Matrix Proteins Bone-matrix proteins have been thought as osteogenic markers and have very important effects for the later stage of osteogenic differentiation. Bone-matrix proteins include collagen and non-collagenous proteins. Non-collagenous proteins mainly are alkaline phosphatase (ALP), osteopontin (OPN), bone sailoprotein (BSP), osteocalcin (OCN) and osteoconectin. Collagen is the major organic component of biological bone (80-90%), in which collagen type I occupies about 95%. Collagen endows bone with structural integrity and the formation of collagen substrutum triggers the differentiation of pre-osteoblasts into osteoblasts in association with other osteogenic factors. In the osteogenic differentiation, the expressions of collagen and osteoconectin occur early. The osteoconectin have the effects to regulate cell adhesion and mineral binding and inhibit HA crystal growth. The expression of osteopontin fluctuates throughout the osteogniec process. Early expression level of OCN is relatively low, which concerned in the “cement-forming” cells and may be a prerequisite for subsequent osteogenesis. Before osteoblastic differentiation, the expression of OPN declines and then is expressed at a higher level with OCN after mineralization has been initiated. The expression of BSP is presented when the formation of osteoblast. BSP has the function to initiate HA crystal formation. The expression of OCN is elevated following BCP expression and concerned in bone-mineral maturation and bone remodelling[5]. From the discussion mentioned above, various bone-matrix proteins have different functions and are expressed at different osteogenic stages. The bonematrix proteins are orderly produced by osteoblasts under the guiding of a cascade of activated genes. The initiation of gene activations is the results of effects of osteoinductive signal molecules and environmental influences. Previous reports have demonstrated the expressions of bone-matrix proteins in osteoinductive process of Ca-P ceramics (Figure 9) [10]. Further studies indicated that the expressions of bone-matrix proteins were material-dependent (Figure 10) [11]. Materials factors may involve in the regulation of Ca-P ceramics to expressions of bone-matrix proteins.

BCP, in muscles of dogs.

Figure 9. The expression of bone matrix specific proteins.

Expression time course of noncollagenous proteins in ectopicly implanted hydroxyapatite/tricalcium phosphate osteopontin osteonectin osteocalcin

150 100 50 0 0.25 0.5

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osteopontin osteonectin osteocalcin

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Figure 10. The time course and level of expressions of bone matrix proteins vary with the material variance.

2

Summary

In the osteoinductive process of Ca-P bioceramics, the materials may selectively absorb and stimulate cells to secrete endogenous osteoinductive signal molecules. The expression of osteogenic genes is regulated in association with the effects of biological environment. The same materials guide MSCs to proliferate and differentiate along osteogenic pathway, ultimately induce bone formation. The regulation of materials to the molecular events occurred in osteoinductive process may be considered as an elucidation for the osteoinducive mechanism. It may also be regarded as a starting point to design osteoinductive biomaterials by the regulation of materials composition and structure to the expressions of genes in cells, further to design the inductive materials for nonosseous tissues. The expression and regulation of genes in osteogenesis are very complex and also are not clear. The presentation only indicates that certain Ca-P bioceramics may intervene in the regulation of the expressions of cellular genes and are important factors for the regulation.

References 1. Zhang X., Zhou P., Zhang J., Chen W., Wu C., 1991. A study of porous block HA ceramics and its osteogenesis, Bioceramics and the Human Body, eds: A. Ravaglioli and A. Krajewski, Elsevier Science, 408-415. 2. Yang Z., Yuan H., Tong W., Zhou P., Chen W., Zhang X., 1996. Osteogenesis in extraskeletally implanted porous calcium phophate ceramics: Variability among differernt kinds of animals, Biomaterials 17, 2131-2137. 3. Yuan H., Yang Z., Li Y., Zhang X., J.D. de Bruijin K. De Groot 1998.Osteoinduction by calcium phosphate biomaterials, J.Mater. Sci. Meter. In Med. 9, 723-726. 4. Yuan H., Yang Z., J.D. de Bruijin K. De Groot, Zhang X., 2001. Material dependent bone induction by calcium phosphate ceramics: a 2.5-year study in dog, Biomaterials 22, 2617-2623. 5. Sodek J., Cheifetz S., 2000. Molecular regulation of osteogenesis, Bone Engineering, ed. T. E. Davies, Rainbow Graphic and printing Ltd, Hongkong, 31-43. 6. Rzeszutek K., Guo L., Davies J.E., 2002. Importance of including proteins in experimenteal models aimed at studying reactive surface apatite layers on calcium phosphate, Key Eng. Mater., 218-220, 829-830.

7. Yuan H., Zhou P., J.D. de Bruijin, Yang Z., K. De Groot, Zhang, X., 1998. Bone morphogenetic protein and ceramic-induced osteogenesis, J. Mater. Sci; Mater. Med. 9, 717-721. 8. Yuan H., 2001. Inhibition of calcium phosphate-induced bone formation by a monoclonal antibody of bovine bone morphogenetic proteins (McAb-bPMA). PhD thesis, University of Leiden, the Netherlants. 103-120. 9. Tan Y., Hong S., Wang X., 2007. Expression of core binding factor 1 (Cbfa1) and osteoblastic markers in C2C12 cells induced by calcium phosphate ceramics in vitro. J Mater Sci: MaterMed 18, 2237-2241. 10. Qu S., Guo X., Weng J., Cheng J.C.Y., Feng B., Yeung H.Y., Zhang X., 2004. Evaluation of the expression of collagen type I in porous calcium phophate ceramics implanted in an extra-osseous site, Biomaterials 25, 659-667. 11. Tan Y., Zhang L., He X., 2005. Gene expression of C2C12 cells on calcium phosphate ceramic in vitro, Key Engineering Materials, 288-289, 265-268.

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Chapter 3 Why is the Clinical Application of Tissue Engineering So Slow? Yoshito Ikada Nara Medical University, Nara, Japan

1

Introduction

The problems with donor organs, such as their limited availability and immunogenicity as well as the possible transmission of viral diseases, have encouraged many people to consider tissue-engineering methods as an alternative strategy to treat patients who need organ or tissue replacement as a result of accidents or diseases. The main purpose of tissue engineering is to regenerate crucial organs or tissues by using the patient’s own cells, in combination with scaffolds and growth factors [1]. Although growth factors are often required to promote tissue regeneration, bolus injection of growth factors in solution is not adequate in many cases due to the short half-life of factors administerted in the body and their rapid dissipation. Therefore, a drug delivery system is required to locally sustain growth factors and protect their bioactivity. This tissue engineering looks so promising for reconstructing diseased or missing tissues and organs that large numbers of studies have been conducted by biomedical engineers and medical people. In spite of their great efforts, it does not seem that both preclinical and clinical trials related to the tissue engineering have been increasingly reported to date. On the other hand, stem cells including embryonic stem (ES) and induced pluripotent stem (iPS) cells have attracted much attention of researchers of not only basic medicine but also clinical medicine and biomedical engineering. It is not surprising that focus has been shifted on the stem cells because of their high potentialities in tissue engineering. Apparently, such an increasing interest has retarded the clinical application of

tissue engineering, but it seems likely that this is not the only reason for the stagnant clinical trials in the field of tissue engineering. This presentation will overview the problems associated with the current tissue engineering from the practical point of view.

2

Surgery Before Tissue Engineering

Until recently, one of the major purposes of surgery was to resect diseased tissues from the body of patients to prevent the spreading of the disease to adjacent healthy tissues unless any effective treatments were clinically available. This kind of surgery is called “resective surgery”. Later, new surgery, called “reconstructive surgery” emerged, when biomaterials and artificial organs had been developed by collaborative efforts of biomedical engineers. Thanks to this novel technology, not only diseased but also lost tissues could be reconstructed, offering much benefit to patients. Among well-known techniques are total hip joint replacement, artificial vascular graft, and dental implant. This evolution of surgical interventions is shown in Fig. 1. Although the performance of biomaterials and artificial organs has been significantly improved within several decades, some medical devices applied to the reconstructive surgery have not always satisfied surgeons and patients. The major shortcomings of current artificial organs involve low capability compared with their natural counterparts, poor biocompatibility such as insufficient antithrombogenicity, and low mechanical durability, especially in artificial joints. Resective Surgery Reconstructive Surgery Regenerative Surgery Figure 1. Shift of surgical treatments.

An alternative compromise to overcome these disadvantages in current applications of man-made materials is to make use of patient’s own tissues or cadaver’s. Apparently, the patient’s tissues are much better than the cadaver’s, because the latter allogeneic tissues might invoke immune troubles such as rejection and pathogen transmission. Fig. 2 represents three illustrations for the reconstruction of a lost tissue with use of patient’s tissue. In Fig. 2(A), a defect of dura mater is covered with a fascial patch harvested from the patient. Fig. 2(B) shows the scheme of reconstruction of a fibula defect using the patient’s

periosteum remaining. It is clinically well-known that new bone will be formed from the mesenchymal stem cells present in the periosteum. When the tissue defect is large in size, scaffolding is necessary in addition to the periosteum to support bone cells generated from the periosteum, as shown in Fig. 2(C). In this case patient’s bone is used as a scaffold after making many small pores in the bone harvested from the patient’s body. As described above, our body, even of adult, stores the stem cells somewhere in the tissues, that are able to proliferate and differentiate into definite tissues, once stimulated. Fibroblast, osteoblast, and chondrocyte also can generate tissues, although these cells do not belong to the stem cell group. It should be, however, emphasized that tissue generation does not generally proceed unless the environment in the vicinity of the cells is adequate for generating tissues. The critical condition for such environments is to ensure an open space, where disturbing cells cannot infiltrate, but nutrients as well as oxygen for the cells are supplied, and a matrix, similar to the natural extracellular matrix (ECM), to which cells can attach, proliferate, and differentiate. In Fig. 2(B) and (C), the periosteal membrane and the perforated bone tissue provide the space and the matrix required for the new tissue generation, respectively. However, the surgical operation is discouraged by such problems that a tissue with an adequate size is not always available to harvest, while it is time-consuming to take out the tissue even if the size is acceptable for the tissue repair.

3

Emergence of Tissue Engineering

Clearly, it is more beneficial if an off-the-shelf biomaterial is available as an alternative of the patient’s own tissue to be used for making the environment optimal for the tissue reconstruction. In principle, there are no such natural tissues that are meaningless to patients. It is “the scaffold” that has been developed to replace this patient’s living tissue for new tissue formation. The scaffold is composed of biomaterials. The history of the man-made scaffold dates back to the 1980s. Probably, the first research group that employed such a man-made scaffold for tissue generation was that of J.P.Vacanti and his co-workers. They published a paper entitled “Selective Cell Transplantation Using Bioabsorbable Artificial Polymers as Matrices” in J.Pediatric Surgery in 1988 [2]. Fig. 3, represented in the paper, schematically demonstrates the principle of their study. Following is the quotation from the paper: ….. We report attaching parenchymal cells from liver, intestine, and pancreas onto bioerodable artificial polymers and then implanting these polymer-cell scaffolds into animals as a novel method of cell transplantation. ….. We reasoned that if we provide an organized scaffolding to which the cells were already attached, we could increase the total number of implanted cells.

(A) Dura mater reconstruction with patient’s fascia

(B) Bone reconstruction with patient’s periosteum

(C) Mandibular reconstruction with the patient’s vascularized bone and metal miniplates Figure 2. Reconstruction of tissues with the help of patient’s own tissues.

Artificial biodegradable polymers were chosen for several reasons. We could engineer configuration, manageability, tensile strength, and rate of degradation to a great degree with a man-made plastic and also might be able to modify the inflammatory response by modifying the material. The polymer material could be coated with other cell types or with attachment factors to increase cell attachment. In addition, we have considerable experience in placing biologically active molecules such as growth factors directly into the polymer, and allowing slow release of these agents in a controlled and predictable way. Finally, by the use of biodegradable matrices, we provide only a temporary scaffold, which eventually is reabsorbed, leaving structural support of the mass to mesenchymal elements supplied by the host, and modified by the implanted cells. ….. This technology allowed us to design the polymers to meet the biologic needs of the system we wished to create. The configuration of the polymer scaffold must have enough surface area for the cells to be nourished by diffusion while neovascularization occurs. The new blood vessels must interdigitate with the implanted parenchymal elements to continue to support their growth, organization, and function. Polymer discs seeded with a monolayer of cells and branching fiber networks both satisfy these needs. The branching fibers are based on the same principles that nature has used to solve the problem of increasing surface area proportionate to volume increases. …..Seeding this configuration with cells and implanting the structure as fibers allows us to implant large numbers of cells, each of which is exposed to the environment of the host. Therefore, free exchange of nutrients and waste can occur while neovascularization is achieved. ….. The results of this study demonstrate that cells from liver, intestine, and pancreas can be successfully harvested, and will attach to artificial biodegradable polymers. They will survive in culture in this configuration and can then be implanted into a host in a variety of locations. An inflammatory response that is mediated by both the wound and the nature of the polymer will occur. Successful engraftment of small clusters of hepatocytes and intestinal cells has been demonstrated. However, we do not yet have evidence of cell function in the new environment. Further studies to define the optimal characteristics of the polymer, attachment parameters, growth criteria, and function, need to be performed. ….. It is surprising that the basic idea of the current tissue engineering is almost entirely included in their paper. Unfortunately, the authors were not successful in obtaining a piece of regenerated tissue, but they could have observed evidence of tissue regeneration if they had chosen connective tissues such as bone and cartilage as the target of tissue regeneration instead of liver and pancreas.

Figure 3. Technique of chimeric neomorphogenesis. Appropriate parenchymal cells are harvested and dispersed. They are seeded onto the polymer matrix in cell culture, where attachment and growth occur. Partial hepatectomy is performed to stimulate growth of the transplant. The polymer-cell scaffold is then implanted into the recipient animal where neovascularization, cell growth, and reabsorption of the polymer matrix occurs.

4

Principle of the Current Tissue Engineering

Although two decades have just passed since the publication of the above-cited, pioneering work by Vacanti et al, the progress of the tissue engineering has not gone as smoothly as anticipated, especially in terms of clinical application. Before discussing the reasons for this slow advance, let us briefly summarize the framework of the current tissue engineering. It is widely accepted that the cell, scaffold, and growth factor are three basic pillars for the tissue engineering that is a constituent of the regenerative medicine, together with the other constituent, the cellular therapy, as represented in Fig. 4. In principle, the cellular therapy (or cell transplantation), like bone marrow transplantation, does not need any other materials than cells for the medical treatment, in contrast with the tissue engineering that demands a scaffold, that is, an artificial ECM, and in some cases, growth factors in addition to cells. It should be pointed out that even the cellular therapy has not yet achieved remarkable advances in the clinical application except for bone marrow transplantation, although this treatment can be performed only with cells. As shown in Fig. 4, the tissue engineering requires the cell, scaffold, and/or growth factor. Which material is the most responsible for the delay in tissue engineering?

5

Reasons for Slow Clinical Application of Tissue Engineering

5.1 Cells Since allogeneic cells are associated with the problem of immune rejection when implanted, their use is generally limited to experiments with scid animals. It would be ideal if a large amount of cells are clinically available as a result of proliferation and differentiation of patient’s own pluripotent stem cells. However, it seems likely that much further progress is needed for cell biology until to reach such a clinically applicable level. Currently, autogeneic, differentiated or progenitor cells are employed for the clinical application of tissue engineering. These cells include fibroblast, osteoblast, chondrocyte, and bone marrow cells. The most attractive cell source among them is the bone marrow, as it can differentiate into mesenchymal cells that will eventually produce bone, cartilage, and blood vessels among others. The largest problem involved in the bone marrow cells is the difficulty in multiplication of cell number maintaining the progenitor state. If an effective method for the cell multiplication is developed, it would significantly contribute to the clinical application of tissue engineering.

Cellular Therapy (Cell Transplantation)

Cell

Regenerative Medicine Tissue Engineering

Cell Scaffold Growth factor

Figure 4. Classification of regenerative medicine.

5.2 Scaffold When a fertilized ovum grows into a fetus in the maternal uterus, the ECM which is indispensable for the cell to generate the corresponding tissue is formed at the required timing and location according to the genetically predetermined schedule. However, when we lose part of our three-dimensional tissues by cancer or do not have a specific tissue because of congenital defect, it is almost impossible to regenerate the missing tissue with the use of the corresponding cells alone, because the ECM serving as scaffolding for the cells is also missing at the required site. This is the rationale for necessitating the artificial ECM, that is, the scaffold for the tissue engineering. This man-made scaffold should work only temporarily, because new natural ECM will be gradually formed along with the tissue regeneration. This means that the biomaterial used for the scaffold should disappear through bioabsorption without concurrent production of cytotoxic, biodegradation by-products. Another basic requirement for the scaffold is to have the structure that guides the tissue formation from cells and allows the access of oxygen and nutrients to the cells residing inside the scaffold. There is another important property that is essential to satisfy the demand of the scaffold, but has often been ignored by biomaterials researchers. That is the mechanical strength that should be high enough to resist the pressure from the inside and the loading from the outside and to allow firm fixation of the scaffold to the adjacent host tissue, for instance, with suturing. This is very critical as the property which the scaffold should meet, especially in the case of clinical applications of tissue engineering. If the animal model chosen for tissue engineering research is so small, such as mice and rat, most of the biomaterials employed for the fabrication of scaffold may satisfy the mechanical properties required for the scaffold. However, the mechanical strength necessary for the scaffold to be used for large animals and humans is much higher than that for small animals. The porous scaffolds prepared from collagen and glycolide-lactide copolymers may be too weak to allow suturing. Poly-L-lactide(PLLA) and poly(ε-caprolactone) can produce mechanically strong scaffolds, but remain in

the body for so long that they will disturb the smooth tissue regeneration. One exception is woven or non-woven fabrics from polyglycolide(PGA), but this polymer is generally too soft and high in degradation rate to be used as a biomaterial for scaffold fabrication. Therefore, it is very likely that the difficulty in fabricating a scaffold with appropriate mechanical strength and degradation rate is an obstacle to the progress of the clinical tissue engineering. A good example for that is the development of a bioabsorbable stent which is very effective in preventing tubular tissues from stricture which takes place during regeneration, as shown in Fig. 5 [3]. A special technique is necessary to fabricate a strong, bioabsorbable stent, even if the structure is as simple as shown in Fig. 5.

Figure 5. Protection of engineered tubular tissue by inserted, resorbable stent from stricture.

5.3 Growth Factors Growth factors are proteins that play a crucial role in cell proliferation and differentiation. Bone morphogenetic proteins (BMPs) are well-known in the tissue engineering, because these proteins induce bone regeneration in the body even at an ectopic site without any addition of osteogenic cells. Another example is basic fibroblast growth factor (bFGF) which gives rise to neovascularization. However, such valuable, biological effects of growth factors would disappear within a short period of time, if they are administered to the target site in bolus solution. To sustain the effects for much longer duration, the growth factors need an assist of a carrier to make the slow release of the proteins possible. It is, however, not easy to develop such a biomaterial that meets the demand as protein drug carrier. The fact that few biomaterials researchers are currently engaged in

the development of such carriers for protein drugs may be one of the reasons for the slow progress of tissue engineering.

6

Ex Vivo and In Situ Tissue Engineering

The tissue engineering has frequently been defined as a biomedical engineering to regenerate living tissues using cells outside the body. This is ex vivo or in vitro tissue engineering. Indeed, skin tissues like epidermis and dermis have been attempted to generate and success has been accomplished to some extent. For these thin tissues, oxygen and nutrients will be relatively readily supplied by diffusion to the cells confined to the generated tissue from the host tissue of patient when the thin constructed tissue is placed on the surface of the host tissue. The articular cartilage is another target that has attracted much attention of tissue engineers because of it’s unique avascular structure, but surgeons have not yet widely applied the articular cartilage reconstructed with the ex vivo tissue engineering. This is probably because the cartilage obtained by the ex vivo tissue engineering is too small in size to be clinically applied or too difficult to fix to the host cartilage. It should be emphasized that these skin tissues and articular cartilage are rather rare tissues that have been reconstructed with ex vivo tissue engineering. Most of other tissues and organs are thicker and bigger in size, and hence require assured vascularization. Such a tissue with vasculature is virtually impossible to regenerate with ex vivo tissue engineering, but only possible with in situ (or in vivo) tissue engineering. In the latter case, vasculature will be formed in the cell-scaffold composite with the help of the host tissue. This strongly suggests that much more attention should be focused on the in situ tissue engineering in order to promote the clinical application of tissue engineering. It should be kept in mind that living cells demand oxygen and nutrients, which can be supplied most efficiently from the vasculature originating from the host tissue, where the cell-seeded scaffold is implanted together with or without growth factors.

7

Clinical Trials of Tissue Engineering

Although tremendous amounts of studies have been conducted from the advent of tissue engineering, this is field still in its infancy from the application point of view. Notwithstanding, some defective tissues have been clinically treated with the tissue engineering. Below are shown some examples of the clinical applications.

7.1 Skin Skin is composed of the epidermis and dermis. Epidermis is formed by stratified keratinocytes, while dermis is composed of an ECM such as collagen fibers and protects inside organs. A bilayered artificial skin without cultured cells was first developed by Yannas et al. [4]. It is composed of an inner sponge layer of collagen and chondroitin-6-sulfate, a glycosaminoglycan, and an outer silicone layer. When placed on wounds, the inner sponge layer characteristically turns into dermis-like connective tissue. This kind of bilayered artificial skins, both of wet and dry types, is commercially available and patients with massive deep burns have successfully been treated with the material.

7.2 Articular Cartilage Natural articular cartilage is composed of a charged solid matrix phase consisting of charged proteoglycan macromolecules and collagen fibers, an interstitial fluid phase, and an ion phase. The proteoglycan matrix is mainly responsible for the equilibrium compressive stiffness of cartilage, the collagen fibers influence the instantaneous compressive and tensile response of the tissue, and the interstitial fluid phase affects the transient response of cartilage during compression. The ability of cartilage to regenerate or “heal” itself decreases with age, because, with age, chondrocyte numbers decrease, structural organization of cartilage becomes altered, and a well-defined calcified zone arises in cartilage. Wakitani et al. applied cell transplantation to repair human articular cartilage defects in osteoarthritis (OA) knee joints [5]. The study group comprised 24 knees of 24 patients with knee OA who underwent a high tibial osteotomy. Adherent cells in bone-marrow aspirates were expanded by culture, embedded in collagen gel, transplanted into the articular cartilage defect in the medial femoral condyle, and covered with autologous periosteum at the time of 12 high tibial osteotomies. The other 12 subjects served as cell-free controls. In the cell-transplanted group, as early as 6.3 weeks after transplantation the defects were covered with white to pink soft tissue, in which metachromasia was partially observed. Forty-two weeks after transplantation, the defects were covered with white soft tissue, in which metachromasia was partially observed for almost all area of the sampled tissue and hyaline cartilage-like tissue was partially observed.

7.3 Bone To prevent failure of the bone-prosthesis interface, Ohgushi et al. attempted a tissue engineering approach using mesenchymal cells of patients [6]. They collected a small amount of fresh marrow cells from the patient’s iliac crest and expanded the number of mesenchymal cells. They applied the cells to a ceramic ankle prosthesis and cultured them to form an osteoblasts/bone matrix on the

prosthesis, as shown in Fig. 6. They used the tissue-engineered prostheses on three patients suffering from ankle arthritis and followed their progress for at least 2 years. X-ray examinations revealed early radiodense appearance (bone formation) around the cell-seeded areas of the prostheses about 2 months after operation, after which a stable host bone-prosthesis interface was established. All patients showed high clinical scores after operation and did not exhibit inflammatory reactions.

Figure 6. Schematic presentation of tibia alumina component. Some surface areas consist of clusters of ceramic beads (represented in gray). Thus, the surface is porous and cells can be easily applied.

7.4 Mandible The loss of the mandibular arch after tumor excisions or traumatic injuries leaves the patient not only with a remarkable deformity of the face but also with a mandible that does not function for chewing, swallowing, or speaking. In such a case, reconstruction has often been performed by bone graft or implantation of prosthesis. Reconstruction by vascular or nonvascular autologous bone grafts has been considered to be reliable, but this treatment has disadvantages including donor site morbidity and insufficient supply. Autologous particulate cancellous bone and marrow (PCBM) that is rich in osteogenic progenitor cells and bone matrices has excellent properties as bone graft because it has full bone formation ability, and, in addition, spontaneous regeneration of donor sites is possible. However, PCBM does not, by itself, have structural strength and the ability to hold its desired shape and is not able to support the newly formed bone while it acquires enough strength to withstand external force. Furthermore, this framework should be preferably bioabsorbable and disappear from the implanted site after completion of bone repair. To address this issue, Kinoshita et al. developed a mesh manufactured from PLLA that can maintain high mechanical strength for a long period of time compared with other bioabsorbable materials [7]. The PLLA monofilaments with diameter of 0.3 or 0.6 mm were woven into mesh sheets, which could be cut with scissors and was easily molded by heating up to

70℃. Fig. 7 shows the PLLA mesh and tray used for mandible regeneration. After extensive preclinical studies with adult dogs, they started clinical studies using the PLLA sheet/tray and autologous PCBM since 1995 to establish a procedure for mandibular reconstruction engineering, as illustrated in Fig. 8. In 8 hospitals in Japan, 62 cases underwent mandibular reconstruction between 1995 and 2001 [8]. Mesh trays were used in 28 cases and mesh sheets in 6 cases. The PCBM was harvested from the iliac bone of patients and 10-40g of PCBM was transplanted to individual patients. The clinical results were evaluated as excellent when the area of osteogenesis, based on X-ray films 6 months after the surgery, was over two-thirds in comparison to right after operation. Results were evaluated as good when osteogenesis was less than two-thirds with no reconstruction required. All other results were graded as poor. And 40 cases (56%) were judged as being excellent, 17 cases (27%) as being good and 10 cases (16%) as being poor.

Figure 7. PLLA mesh sheet and mandibular mesh tray.

7.5 Orbital Floor Yamazaki et al. treated 9 patients with orbital floor fractures using PLLA meshes shaped mostly to ellipses ranging from 22 × 27 to 10 × 15 mm2 in size [9]. The largest amount of PCBM (4 g) was needed for the patient to whom a mesh sheet of 22 × 27 mm2 was applied. The necessary amount of PCBM was taken

readily using a bone curet after simply making a skin incision of 2-cm length in the waist and making a small hole of 7-8 mm diameter in the cortical bone of the iliac bone. The collected PCBM was filled in the gaps of the PLLA mesh. Postoperative symptoms were extremely mild, and few patients complained of pain. The CT images showed new bone formation at the orbital floor 3 months after operation. Diplopia was resolved in all patients. Examination of the visual field using the Hess screen showed clear improvements in all patients and showed almost symmetrical visual fields. No displacement of globes recurred for the longest follow-up of 61 months.

Figure 8. Schematic drawing of mandibular reconstruction by means of PLLA mesh and particulate cancellous bone and marrow (PCBM). The PLLA mesh tray was adjusted to the shape and size of the bone defect with cutting by scissors and warming at about 70℃. The PLLA mesh tray was fixed to the residual bone with stainless steel wires and filled with PCBM taken from the iliac bone.

7.6 Vascular Tissue It was shown in 2003 that bone marrow cells (BMCs) seeded onto a biodegradable scaffold contribute to the histogenesis of tissue-engineered vascular autografts (TEVAs) in canine in vivo models [10]. However, as no reports had described the function of endothelial cells or the mechanical properties and characteristics of TEVAs, Matsumura et al. evaluated the endothelial function and mechanical strength of TEVAs constructed with autologous mononuclear bone marrow cells (MN-BMCs) and a

l-lactide-ε-caprolactone copolymer, P(LA/CL), scaffold using a canine inferior vena cava (IVC) model, as shown in Fig. 9 [11]. Fig. 10 shows no statistical differences in strength among IVCs of dog (shown as a control) and 6- and 12-month TEVAs.

Figure 9. Tensile strength of biodegradable scaffold in vitro. The biodegradable scaffold used in this study diminished sequentially within 1 month. Data represent the mean ± standard error of 5 samples at each time point. *p < 0.01, **p < 0.001 vs. week 0.

After the successful results of the supplementary examination in a dog IVC replacement model, Shin’oka et al. reconstructed peripheral pulmonary in a 4-year-old girl with the patient’s own venous cells [12]. After that, 3 patients underwent tissue-engineered graft implantation with cultured autologous venous cells. However, since cell culturing was time-consuming and xenoserum had to be used, they began to use BMCs, readily available on the day of surgery, as a cell source. A 5-mL/kg specimen of bone marrow was aspirated with the patient under general anesthesia before skin incision. A P(LA/CL) tube serving as a scaffold for the cells was the same as described above. Twenty-three tissue-engineered conduits for extracardiac total cavopulmonary connection (TCPC) and 19 tissue-engineered patches were used for the repair of congenital heart defects. The patients’ages ranged from 1 to 24 years [13]. Mean follow-up after surgery was 490 ± 276 days. There were no complications such as thrombosis, stenosis, or obstruction of the tissue-engineered autografts. As shown in Fig. 11, the maximal trans-sectional area was calculated and compared with the implanted size in the TCPC group. There was no evidence of aneurysm formation or calcification on cineangiography or computed tomography. All tube grafts were patent and the

diameter of the tube graft increased with time (110 ± 7% of the implanted size), suggesting that these vascular structure may have the potential for growth, repair, and remodeling, and provide an important alternative to the use of prosthetic materials in the field of pediatric cardiovascular surgery.

Figure 10. Increase in mechanical properties of tissue-engineered vascular autografts Dunnett’s post hoc analysis (TEVAs). (A and B) Tensile strength (mN) and stiffness (mN/mm) of TEVAs (1, 3, 6 and 12 months after implantation). Kruska-Wallis test; p < 0.001. ***p < 0.001 vs. inferior vena cava, ††p < 0.01 vs. TEVAs at 1 month, ††† p < 0.001 vs. TEVAs at 1 month, ‡p < 0.05 vs. TEVAs at 3 months, ‡‡p < 0.01 vs. TEVAs at 3 months. (B) Stiffness/width of TEVAs increased with time. Stiffness/width was calculated as stiffness (mN/mm) = elastic modules (mN/mm2)×wall thickness (mm). Analysis of variance (ANOVA) p < 0.01. *p < 0.05 vs. TEVAs at 1 month. Data from dog inferior vena cava are shown as control. Data represent mean ± standard error.

7.7 Pericardial Tissue Patients with congenital heart defects require re-sternotomy during reoperation and staged-operation. Re-sternotomy and dissections are risks of the reoperations as mediastinal adhesion formation may result in the iatrogenic injury of the heart,

Patient 20: 12 months after operation

Figure. 11 Three-dimensional computed tomographic image 12 months after implantation of TCPC graft. Arrow indicates location of tissue- engineered conduit.

artery, and veins. Moreover, mediastinal adhesion makes surgeons difficult to identify anatomical feature, prolongs the operating time for dissection and might requie more blood transfusion. Several attempts have been made to resolve the problem related to pericardial adhesions after cardiac surgery by using different types of pericardial substitutes, including expanded PTFE patches, silicone rubber, and glutaraldehyde-treated bovine or porcine pericardium xenografts. Recently, Matsumura et al. reported the result of clinical trial performed using the gelatin sheet reinforced with PGA [14]. Five patients who were anticipated for further repair of congenital defects or for repeated cardiac surgery had received the anti-adhesive sheet. There was neither mortality in this trial nor morbidity such as prolonged inflammatory response, infectious complication, signs of tamponade, or early repeated thoracotomy due to the material. As shown in Tab.1, re-sternotomy took 19 to 34 minutes from the beginning of the operation to the Table 1. Effect of the Gelatin Sheet on Adhesions and Operative Times. Degree of adhesion at re-sternotomy

Time

Case

Gelatin sheet

Other site in mediastinum

op start to re-sternotomy (min)

re-sternotomy to CPB(min)

1

slightly

severe

20

261

2

slightly

severe

34

196

3

slightly

moderate

26

265

4 5

slightly slightly

severe moderate

24 19

176 160

sternum reopening. In each case, re-sternotomy was easily performed and there were no complications attributed to its use. All the doctors who joined the operation scored the anti-adhesive effectiveness as “good”.

7.8 Dura Mater To replace the dura mater, Yamada et al. fabricated a bioabsorbable sheet with sandwich structure as illustrated in Fig. 12 [15]. A non-woven PGA fabric was used to reinforce two P(LA/CL) sheets. This composite sheet displayed good mechanical properties and was completely absorbed 24 weeks after implantation in the back of rats. And 2 weeks to 26 months after implantation into 31 rabbits

Figure 12. Preparation of a three-layered sheet from P(LA/CL) and PGA for dura mater regeneration.

with dural defects, infection, cerebrospinal fluid leakage, evidence of convulsive of disorders, significant adhesion to underlying cortex, or calcification were not noticed in any case. The regenerated dura-like tissue had a high pressure-resistant strength 2 weeks after implantation. Yamada et al. applied this absorbable sheet to 53 patients during neurosurgical procedures [16]. The average follow-up was 35.5 months. The handling properties and biocompatibility of this sheet were satisfactory without any significant complication. In patients who underwent a second surgery performed more than 18 months after the initial operation, this sheet was found to have been replaced by autologous collagenous tissue, and chronic foreign-body reactions to this material were negligible.

8

Conclusions

Finally, let us make some remarks on pluripotent stem cells including ES and iPS cells. It is extremely important for biology, especially embryology and cellular biology to study these pluripotent cells, because they may eventually differentiate into all kinds of mature cells. However, it will take decades until any cells clinically applicable to the tissue engineering can be provided from these pluripotent cells after their selective differentiation and proliferation. If the cells differentiated from the patient’s stem cells are of endocrine type such as insulin-secreting β-cells in the Langerhans ilets and dopamine-secreting cells in the substantia nigra, these mature cells may be applicable to the clinical treatment of diabetes mellitus and Parkinson’s disease, respectively. However, for other non-endocrine diseases, a scaffold is mostly indispensable for the treatment in addition to the cells differentiated from the pluripotent cells. It may be concluded that for promoting the clinical tissue engineering now is the fine for us to come back to the starting point marked two decades ago.

References 1. 2.

3. 4.

Ikada Y., 2006. Tissue Engineering: Fundamentals and Applications, Academic Press, UK. Vacanti J.P., Morse M.A., Saltzman W.M., Domb A.J., Perez-Atayde A., Langer R., 1988. Selective Cell Transplantation Using Bioabsorbable Artificial Polymers as Matrices, J. Pediatr. Surg., 23, 3-9. Ikada Y., 2006. Challenges in tissue engineering, J. R. Soc. Interface, 3, 589-601. Yannas I.V., Burke J.F., Orgill D.P. Skrabut E.M., 1981. Wound tissue can utilize a polymeric template to synthesize a functional extension of skin, Science, 215, 174-176.

5.

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12. 13.

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16.

Wakitani S., Imoto K., Yamamoto T. Saito M., Murata N., Yoneda M., 2002. Human autologous culture expanded bone marrow mesenchymal cell transplantation for repair of cartilage defects in osteoarthritic knees, Osteoarthritis and Cartilage, 10, 199-206. Ohgushi H., Kotobuki N., Funaoka H., Machida H., Hirose M., Tanaka Y., Takakura Y., 2005. Tissue engineered ceramic artificial joint-ex vivo osteogenic differentiation of patient mesenchymal cells on total ankle joints for treatment of osteoarthritis, Biomaterials, 26, 4654-4661. Kinoshita Y., Kobayashi M., Fukuoka S. Yokoya S., Ikada Y., 1996. Functional reconstruction of the jawbones using poly(L-lactide) mesh and autogenic particulate cancellous bone and marrow, Tissue Engineering, 2, 327-341. Kinoshita Y., Yokoya S., Amagasa T. et al., 2003. Reconstruction of jawbones using poly(L-lactic acid) mesh and transplantation of particulate cancellous bone and marrow: Long-term observation of 40 cases, Int. J, Oral Maxillofac. Surg., 32(suppl.1), 117. Yamazaki T., Kinoshita Y., Ikada Y., 2001. Repair of human orbital floor defects using poly(L-lactide) mesh filled with particulate cancellous bone and marrow, Int. J. Oral Maxillofac. Surg. 30(suppl.) 92. Matsumura G., Miyagawa-Tomita S., Shin’oka T., Ikada Y., Kurosawa H., 2003. First evidence that bone marrow cells contribute to the construction of tissue-engineered vascular autografts in vivo, Circulation, 108, 1729-1734. Matsumura G., Ishihara Y., Miyagawa-Tomita S., Ikada Y., Matsuda S., Kurosawa H., Shin’oka T., 2006. Evaluation of Tissue-Engineered Vascular Autografts, Tissue Engineering, 12, 3075-3083. Shin’oka T., Imai Y., Ikada Y., 2001. Transplantation of a tissue-engineered pulmonary artery, N. Engl. J. Med., 344, 532-533. Shin’oka T., Matsumura G., Hibino N., Naito Y., Watanabe M., Konuma T., Sakamoto T., Nagatsu M., Kurosawa H., 2005. Midterm clinical result of tissue-engineered vascular autografts seeded with autologous bone marrow cells, J. Thorac. Cardiovasc. Surg., 129, 1330-1338. Matsumura G.., Shin’oka T., Ikada Y., Sakamoto T., Kurosawa K., 2008. Novel Anti-Adhesive Pericardial Substitute for Multistage Cardiac. Surgery, Asian Cardiovasc. thorac. Ann., 16, 309-312. Yamada K., Miyamoto S., Nagata I. Kikuchi H., Ikada Y., Iwata H., Yamamoto K., 1997. Development of a dura substitute from synthetic bioabsorbable polymers, J. Neurosurg., 86, 1012-1017. Yamada K., Miyamoto S., Takayama M. Nagata I., Hashimoto N., Ikada Y., Kikuchi H., 2002. Clinical application of a new bioabsorbable artificial dura mater, J. Neurosurg., 96, 731-735.

Chapter 4 Temperature-responsive Cell Culture Surfaces for Cell Sheet Tissue Engineering Zhonglan Tang, Yoshikatsu Akiyama, and Teruo Okano Institute of Advanced Biomedical Engineering and Science, Tokyo Women’s Medical University

1

Introduction

In the past two decades, cell-based therapies, regenerative medicine, and tissue engineering have progressed rapidly[1–8]. However, these conventional methods have limitations and present problems in clinical applications. When injecting single cells suspension, it is difficult to control the size, shape, and location of the cells in the host tissue. The migration of the injected cells suspension results in inefficient cell anchorage at the target host tissue site. In order to overcome these problems, tissue engineering techniques utilizing biodegradable scaffolds have been applied to the fabrication of three-dimensional tissue-like grafts[4–8]. During degradation of the scaffold, cells are expected to proliferate and migrate to replace the scaffold. However, the places previously occupied by the polymer scaffold are filled with a larger amount of extracellular matrix (ECM). These results also suggest that biodegradable scaffolds may not be suitable for regenerating cell-dense tissues such as the heart and the liver. Cells at the center of larger constructs become necrotic, while the cells on the periphery are unimpaired. The delivery of oxygen and nutrients and the removal of metabolic waste are thought to be limited by passive diffusion. Moreover, strong

inflammatory responses are often observed upon biodegradation of the scaffolds, which could damage the seeded cells, even though the biodegradable scaffolds are non-toxic and mechanically non-invasive, resulting in the failure of the engineered tissues[9,10]. In order to overcome the limitations of conventional methods, we have developed cell sheet engineering as an innovative approach for regenerative medicine[11–15]. In this technology, cells could be harvested as an intact monolayer cell sheet, including the deposited ECM, by a “temperatureresponsive culture dish,” which enables reversible cell adhesion to and detachment from the dish surface in response to temperature. In this chapter, we explain the features of thermo-responsive culture dishes and the mechanism of cells adhesion and detachment in responsive temperature. Furthermore, we have summarized novel temperature-responsive cell culture surfaces using functional monomers.

2

Temperature-Responsive Cell Culture Surfaces for Reversible Cellular Adhesion and Detachment

2.1 Ultrathin Poly(N-isopropylacrylamide) Gel Grafted on Cell Culture Surfaces Poly(N-isopropylacrylamide) (PIPAAm) is a well-known intelligent polymer, which exhibits temperature-responsive soluble/insoluble changes in an aqueous solution and has a lower critical solution temperature (LCST) at 32°C[16]. Thermoresponsive cell culture dishes were proposed by our laboratory 10 years ago. PIPAAm was introduced on the surfaces of commercially available tissue culture polystyrene (TCPS) dishes[17–19]. First, IPAAm monomers were dissolved in 2-propanol solutions. The solutions were then uniformly spread over TCPS dish surfaces. IPAAm monomers are polymerized and simultaneously covalently grafted onto the TCPS surface with a 0.3 MGy dose of electron beam (EB) irradiation. Owing to hydration of the grafted PIPAAm chains, PIPAAm grafted TCPS (PIPAAm-TCPS) surfaces exhibit hydrophilic and cell repellent surfaces at temperatures below 32°C. In contrast, PIPAAm-TCPS surfaces altered from a hydrophilic to hydrophobic state due to dehydration of the grafted PIPAAm chains, as the temperature rose above 32°C. Consequently, cells attached to, adhered to, and proliferated on PIPAAm-TCPS at temperatures above 32°C. Furthermore, the adhered cells were spontaneously detached from PIPAAm-TCPS surfaces when the temperature was reduced below 32°C.

Table 1. Characterizations of PIPAAm grafted surfaces. Amount of grafted PIPAAm (μg/cm2)a Thickness of PIPAAm layer (nm)b Contact angle (cosθ)c Cell adhesion (37oC) o

Cell detachment (20 C)

o

37 C 20oC

1.4 ± 0.1

2.9 ± 0.1

1,080

15.5 ± 7.2

29.3 ± 8.4

5,000

0.20 0.42

0.35 0.50

0.65 (40oC) 0.98 (10oC)

positive

negative

negative

positive

d

—

—d

a: n = 4, mean ± SD; b: n = 4, mean ± SD; c: n = 3; d: Not determined.

Such cell adhesion-detachment modulation on surfaces grafted with PIPAAm is a novel concept given that no enzymes or chelators are required. In order to achieve cell adhesion and detachment on PIPAAm-TCPS in response to temperature changes, a key requirement is that the PIPAAm layer grafted on the TCPS surface should have a thickness of 20 nm [19]. As shown in Table 1, two types of PIPAAm-TCPS show hydrophobic and hydrophilic property alterations from 37°C to 20°C. One PIPAAm-TCPS has a thinner PIPAAm layer (approximately 20 nm thick) and a lower density of grafted PIPAAm, while the other has a thicker layer (approximately 30 nm) and a larger PIPAAm density. However, a thick and large amount of grafted PIPAAm produces low aqueous contact angles (more hydrophobic) at both temperatures. Temperature-responsive cell adhesion and detachment were observed on a PIPAAm-TCPS with a thin and low-density PIPAAm layer (the density ranged from 1.4 to 2.0 μg/cm2). No cell adhesion occurred on the PIPAAm-grafted layer with a polymer density greater than 2.0 μg/cm2 and a PIPAAm thickness of more than 30 nm. Significantly, the contact angle reduced with an increase in the polymer density and thickness [20]. As reported in previous studies, the configuration of the graft chain and its mobility affect the wettability of the PIPAAm-grafted surface with varying temperatures. Single terminal graft of PIPAAm with liberated mobile chains exhibited more hydrophilic surfaces. In contrast, loopedchain configuration followed by multi-point graft of PIPAAm chains has more restricted polymer chains, showing more hydrophobic surfaces. Comparison of the swelling ratio of single-side fixed/cross-linked PIPAAm hydrogel, nonfixed/cross-linked PIPAAm gel revealed higher swelling ratio was significantly restricted with the single-side fixed gel. Further comparison of single side

fixed/cross-linked PIPAAm hydrogels with different thicknesses revealed that the swelling ratio for the thinner gels was less than half that of the gels with double thickness. These results suggested that the mobility of polymer chains in PIPAAm gels fixed on a covered glass surface was limited at the interface of the basal cover slips and produced extensive hydrophobic aggregates and limited hydration of the polymer chain. The limitations of polymer chain mobility progressively limited general polymer network responses through cross-linked points inside the polymer gel. PIPAAm chains had more free mobility and a greater ability to respond to stimuli and hydration in the outer regions than in the inner regions. This consideration is extended to PIPAAm-TCPS systems with an ultra thin PIPAAm layer.

Figure 1. Schematic illustration of the influence of the molecular mobility of grafted PIPAAm chains on cell adhesion characteristics at 37°C. The grafted PIPAAm gels are thin (left side) and thick (right side). The darker chains denote more restricted molecular motion, while brighter chains indicate greater mobility.

Fig. 1 schematically illustrates the possible chain motilities for PIPAAmTCPS with thin and thick PIPAAm layers. At the interface of the PIPAAm and polystyrene (PS) substrate, the PIPAAm chains were highly aggregated and dehydrated due to hydrophobic basal PS surfaces (aggregated region). Such aggregation and dehydration progressively extended to the outermost region, and was accompanied by limitations in chain mobility and an impact on the restricted layer, as illustrated in Fig. 1. For the thinner PIPAAm layer, the limited mobility of the polymer chains in the restricted layer promoted dehydration of the PIPAAm chains; the outermost surfaces were sufficiently hydrophobic to show the adhesive properties of the cells at 37°C. The restricted layer region is considered to be between 15 and 20 nm. Such promotion of dehydration and the limitations of the grafted polymer chains result in modulation of the cell attachment and detachment in response to temperature changes. In contrast, for a thicker PIPAAm layer, dehydration of the PIPAAm chains at the outermost surfaces (hydrated region) are not promoted even above the LCST, and do not suffer limitations from the polymer chains and dehydration generated at the basal PS surfaces. It is known that cells do not adhere on PIAAm hydrogel even at 37°C. The PIPAAm chains in the dehydrated layer are still dehydrated but are less restricted, like a free gel. As a result, cells could not adhere onto the surface

of the PIPAAm-TCPS with a thick PIPAAm layer. Therefore, controlling the PIPAAm surface graft density is a critical issue to reliably produce temperatureresponsive cell adhesion/detachment behavior. The thickness dependency of cell adhesion/detachment behavior was also proven in another study, where PIPAAm was grafted onto a glass cover slip using EB irradiation[21]. No cell adhesion occurred on the glass cover slip with a grafted PIPAAm of more than 1.28 μm/cm2 density and 7.4 nm thickness.

2.2 Cell Sheet Engineering Generally, enzymatic treatment using trypsin, pronase, collagenase, etc. is required for recovery of the cultured cells from cell culture surfaces such as tissue culture polystyrene. [22] The enzymatic treatment inflicts damage to cell membranes by hydrolyzing various membrane- associated proteins and cuts off the cell-cell junction, resulting in cell function impairment[22–24]. In contrast, for temperature-responsive cell culture surfaces, PIPAAm-TCPS, cells are spontaneously detached as a single contiguous sheet by reducing temperature after proliferating to confluence, due to the alteration from a hydrophobic to a hydrophilic surface. There is no significant difference of expression for phenotypic markers of cells between PIPAAm-TCPS and commercially available TCPS. Given mild temperature changes and no requirement for enzymatic treatment, cells are not subjected to damage such as destruction of cell-cell junction and cell membrane proteins during the recovery process. Thus, the recovered cell sheet maintains deposited ECM at the basal surface and cellcell junctions [25]. The deposited ECM allows easy adhesion of the recovered cell sheet to various surfaces such as the culture dish, other cells sheets, and the host tissue. Optimal temperature conditions for efficient cell recovery vary by cell species (20°C for cultured endothelial cells and 10°C for hepatocytes) [26]. The cell detachment process involves intracellular signal transduction and reorganization of the cytoskeleton, accompanied by the consumption of ATP molecules. In order to recover and manipulate the cell sheet, supporting membranes composed of poly(ethylene terephthalate), poly(vinylidene- difluoride), chitin, or parchment paper sheets were used[15, 25, 27]. This membrane allows the direct transfer of cell sheets onto native tissue or other cell sheets for preparation of the layered cell sheets. In addition to these membranes, we recently developed a polyion complex that can be used as a novel support[28]. We refer to such new cell manipulation technology as cell sheet engineering; we have utilized the technique to prepare a variety of cell sheets. In addition, many types of tissues have been successfully reconstructed using cell sheet engineering technology,

including skin[29], corneal epithelium [30,31], bladders [32], and cardiac tissue [33]. Tissue reconstruction using cell sheet engineering has several advantages over direct cell injection or tissue reconstruction with biodegradable scaffolds. Only a mild low temperature treatment for cell detachment (with no enzyme treatment) allows differentiated cellular form and complete functions. Cell sheets can be attached to host tissue and even to wound sites via the deposited ECM, with minimal cell loss. In contrast, there is often significant cell loss associated with direct injection of single cell suspensions. There are also fewer inflammatory responses, which are typically observed upon biodegradation of scaffolds. Bovine aortic endothelial cells (ECs) were cultured on ungrafted TCPS dishes and on PIPAAm-grafted dishes[25]. Cells adhered, spread, and proliferated to confluence in both kinds of dishes. As shown in Fig. 2a and 2b, cell morphologies exhibited in two kinds of dishes were similar. Moreover, immunofluorescence microscopy revealed that F-actin and fibronectin (FN) were deposited and accumulated on the grafted surfaces during culture. Cells maintained these proteins after cell sheet detachment (Figs. 2c and 2d). Since the actin filament exerted contractile forces sufficient to shrink the entire cell sheets, the detached cell sheets shrank and folded (Fig. 2c).

Figure 2. Bovine aortic endothelial cell sheets were detached from an ungrafted TCPS dish (a) and PIPAAm-grafted dish (b) after low-temperature treatment. Detached cell sheets were fixed and stained with rhodamine-phalloidin (c) or double-stained with an anti-FN antibody and a fluorescent dye for nuclei (d). FN matrix was recovered with the cell sheet and no remnants were observed on the surface where the cell sheet was detached (right bottom corner in (d)).

3

Functionalization of Temperature-Responsive Cell Culture Surfaces

3.1 Incorporation of Poly(n-butyl Methacrylate) to Intelligent Surfaces It is well known that transition temperatures of thermoresponsive polymers in water were modulated by copolymerization of IPAAm with a hydrophobic or hydrophilic monomer[34–37]. To control cell attachment and detachment on tissue culture surfaces at definite temperatures, IPAAm is copolymerized with a hydrophobic monomer, n-butyl methacrylate (BMA), and grafted to a TCPS surface using EB irradiation[38]. P(IPAAm-co-BMA)-grafted TCPS surfaces reduced transition temperature as BMA content increased, as determined by the contact angle. Cells adhered and proliferated to confluence on P(IPAAm-co- BMA)grafted TCPS as well as on PIPAAm-TCPS at 37°C. Although after the confluence the cells were successfully detached as a contiguous cell sheet from both PIPAAm-TCPS and P(IPAAm-co-BMA)-TCPS by reducing the temperature (20°C), a longer incubation period was required for complete cell detachment with increasing BMA content in the copolymers. At 28°C, detachment of cell monolayers was observed from only PIPAAm-TCPS and P(IPAAm-co-BMA)-TCPS with 1 mol% BMA introduced-grafted copolymer (P(IPAAm-co-BMA)-TCPS-IB1), while no cell monolayer detached from either P(IPAAm-co-BMA)-TCPS-IB3 or P(IPAAm-co-BMA)-TCPS-IB5 surfaces. P(IPAAm-co-BMA)-TCPS- IB3 and P(IPAAm-co-BMA)-TCPS-IB5 surfaces remained hydrophobic at this temperature due to the reduction in the transition temperature of the surface as a result of the incorporation of the BMA unit. In sharp contrast, complete recovery of cells sheet from P(IPAAm-co-BMA)TCPS-IB3 and P(IPAAm-co-BMA)-TCPS-IB5 surfaces were observed at 25°C and 20°C. These results indicate that cell attachment and detachment could be controlled to an arbitrary temperature by varying the BMA content incorporated into the PIPAAm-TCPS.

3.2 Co-grafting of Poly(N-isopropylacrylamide) and Poly(n-butyl Methacrylate) to Intelligent Surfaces Heterotypic cell-to-cell interactions are crucial to achieve and maintain specific functions in many tissues and organs[39]. To mimic such heterotypic cellular interactions in vitro, co-culture of different cell types has been carried out on various types of patterned surfaces. We developed a patterned surfaces modification technique to enable the co- culture of heterotypic cells and the

recovery of patterned co-cultured cell sheets. This new technique can be applied in tissue engineering, incorporating poly(n-butyl methacrylate) (BMA) into limited surface areas of PIPAAm-TCPS using EB irradiation and a metal mask with 1 mm holes[40]. The resulting surface consists of PIPAAm-grafted and P(IPAAm–BMA) co-grafted domain, exhibiting patterned dual thermoresponsive properties. The transition temperature of P(IPAAm–BMA) decreased with the increasing BMA content in the feed, as determined by the contact angle. On patterned surfaces, domain- selective adhesion and growth of rat primary hepatocytes (HCs) and bovine carotid endothelial cells (ECs) allowed patterned co-culture, exploiting hydrophobic/hydrophilic surface chemistry in response to the sole variable of culture temperature. At 27°C, seeded HCs adhered selectively onto hydrophobic, dehydrated P(IPAAm– BMA) domains (1-mm+ area), but not onto neighboring hydrated PIPAAm domains (Fig. 3a). Sequentially seeded ECs then adhered exclusively to hydrophobised PIPAAm domains with increasing culture temperature to 37°C, achieving patterned co-cultures (Fig. 3b). Alternation of polymer-grafted domains to hydrophilic property allows the release of co-cultured, patterned cell monolayers as continuous cell sheets with heterotypic cell interactions, by lowering the temperature to 20°C (Fig. 3d).

Figure 3. Patterned dual temperature-responsive surface for the co-culture of HCs and ECs. (a) HCs adhered and proliferated on P(IPAAm–BMA) co-grafted domain at 27°C; (b) ECs adhered and proliferated on PIPAAm-grafted domain at 37°C and were cocultured with HCs; (c) magnified view of the periphery of patterned co-cultures (square region in (b)); (d) co-cultured cell sheet detached from the patterned culture surface. Scale bars: (a), (b), and (d): 0.5 mm and (c): 0.2 mm.

Figure 4. Schematic illustration of the immobilization of biomolecules to a carboxylfunctionalized temperature-responsive polymers.

3.3 Biofunctionalized Intelligent Surfaces with 2-Carboxy-isopropylAcrylamide To conjugate various bioactive molecules with PIPAAm, reactive moieties such as acrylic acid (AAc) were introduced into PIPAAm[41]. However, such an introduction not only moves the phase transition shift but also dampens the sharp phase transition. Newly synthesized 2-carboxyisopropylacrylamide (CIPAAm) has a side chain structure that is similar to that of IPAAm and a functional carboxylate group; therefore, it overcomes these shortcomings[41.42]. In actuality, the P(IPAAm-co- CIPAAm) solution shows steep phase transition in response to temperature, and the LCST is almost the same as that of PIPAAm. Utilizing the advantage of CIPAAm, we prepared temperature- responsive poly(N-isopropylacrylamide-co-2-carboxyisopropylacryl- amide) (P(IPAAmco-CIPAAm)) copolymer grafted onto TCPS dishes and immobilized RGDS peptide(43), found in FN, type I collagen, and other ECM proteins, onto the surfaces using the reactive carboxyl groups[42]. These surfaces facilitate the spread of human umbilical vein endothelial cells (HUVECs) without serum, depending on the RGDS surface content at temperatures above 37°C (the LCST of the copolymer) (Fig. 4). Moreover, cells spread on RGDS-immobilized surfaces at 37°C detached spontaneously by lowering the culture temperature below the LCST, as hydrated, grafted copolymer chains dissociate immobilized RGDS from cell integrins. The cell-lifting impacts on hydration are similar to the results achieved using soluble RGDS in a culture, as a competitive substitution for immobilized ligands. Binding of cell integrins to immobilized RGDS on cell

culture substrates can be reversed spontaneously using mild environmental stimulation, such as temperature, without enzymatic or chemical treatment. These findings are important to control the specific interactions between proteins and cells, and subsequent “on-off” regulation of their function. [42, 44– 46] Furthermore, since the method allows for a serum-free cell culture and a trypsin-free cell harvest, it should be an attractive cell culture method that does not require mammalian-sourced components.

4

Rapid Recovery of Cell Sheets

4.1 Poly(N-isoproplacrylamide) Grafted on Porous Cell Culture Membranes Although cell sheets were recovered from PIPAAm surfaces together with their deposited FN matrix with a low temperature treatment, gradual hydration of the grafted PIPAAm chains from periphery towards the interior cause slow detachment process of cell sheets (Fig. 5a). As a result, a long incubation period is required to completely recover an intact cell sheet at reduced temperatures. The rapid recovery of cell sheets is important for maintaining the biological function and viability of recovered cell sheets. In addition, it also contributes reduction for time practical assembly of tissue structures and for patient burden in clinical step.

Figure 5. Schematic illustration of cell sheet detachment with varying water supply: (a) PIPAAm-TCPS, (b) PIPAAm-PM, and (c) P(IPAAm-co-PEG)-PM.

Figure 6. The average attached areas of cell sheets recovered from PIPAAm-TCPS (filled squares), PIPAAm-PM (filled circles) and P(IPAAm-co-PEG)-PM (blank circles) at 20°C.

A highly water permeable substrate is exploited to efficiently supply water molecules to interface between the cell sheets and the polymer surface and to accelerate the hydration of the hydrophobic PIPAAm chains beneath the cell sheet. We covalently grafted PIPAAm onto porous membranes (PMs) (PIPAAmPM) by EB irradiation [47]. At a cell culture temperature of 37°C, cells were attached and spread on this membrane as well as on a PIPAAm-grafted surface. Since the diameter of the pore was smaller than that of a single cell, the cell could adhere and proliferate on the porous membrane without influence. By reducing temperature below LCST (20°C), cell sheets successfully were detached from the PIPAAm-grafted PMs. Significantly, cell sheets detached from the PMs more rapidly than that from the TCPS surface (Fig. 6). Compared to the TCPS surface, hydration of the PIPAAm chains were promoted by water penetration from not only the cell sheet periphery, but also from the pore beneath the attached cell sheet (Fig. 5b) and achieved rapid detachment of the cell sheet.

4.2 Poly(Ethylene Glycol ) Co-grafted with Poly(N-Isoproplacrylamide) onto Porous Cell Culture Membranes Several years ago, our group reported on comb-type grafted PIPAAm hydrogels, where poly(ethylene glycol) (PEG) graft chains were introduced into a PIPAAm cross-linked network by co- polymerization. [48, 49] Such PIPAAm hydrogels exhibited rapid-responsive swelling/de-swelling characteristics while

maintaining the polymer transition temperature as observed for homopolymer IPAAm. Grafted hydrophilic PEG chains forming channels for water molecules enhanced the rate of the hydration of the grafted PIPAAm chains and, consequently, showed rapid swelling/de-swelling of the gels. Thus, we expected that introducing such PEG chains into PIPAAm-grafted surfaces could also achieve rapid cell detachment followed by rapid hydration of the grafted PIPAAm gel. PIPAAm was grafted with PEG chains onto porous membranes (P(IPAAmco-PEG)-PMs) by using EB irradiation[50]. For comparison, PIPAAm-PMs were prepared in a same way. Both polymer grafted MPs surfaces exhibited hydropholic/hydrophobic alternation by temperature change. At 37°C, contact angles of P(IPAAm-co-PEG)-PMs Grafted with 0.1 mol% PEG in seed (P(IPAAm-co-PEG0.1)-PM) and 0.5 mol% PEG in seed (P(IPAAm-coPEG0.5)-PM) (cosθ values = 0.535 ± 0.016) were almost equivalent to that of PIPAAm-PM (cosθ values = 0.530 ± 0.048) . In contrast, contact angles of P(IPAAm-co-PEG)-PMs decreased with an increase in concentration of PEG in feed at 20°C. These results suggested that grafted PEG chains scarcely exist in the outermost area of the membranes surface due to hydrophobically contracted PIPAAm chains at 37°C. Below LCST, embedded PEG chains interior of the dehydrated PIPAAm network with limited mobility are more mobile through hydration of the PIPAAm chains by lowering temperature, leading to rearrangement of PEG chains. Cells adhered and grew to confluence on a P(IPAAm-co-PEG)-PM as well as on a PIPAAm-PM at 37°C. However, they did not proliferate to confluency on (P(IPAAm-co-PEG1.0) surfaces because of more hydrophilic property. By reducing the temperature to 20°C, cell sheets detached from the P(IPAAm-co-PEG0.1)-PM more quickly than the PIPAAmTCPS and the PIPAA-PM (Fig. 6). These results suggest that both the pore beneath the cell sheet and the PEG chains enhanced water diffusion to hydrate the PIPAAm surface, resulting in particularly rapid cell sheet detachment (Fig.5c).

5

Conclusion

We have proposed a cell sheet engineering technology for use in tissue engineering by utilizing a thermoresponsive cell culture surface. An intact cell sheet is successfully recovered for fabricating functional tissue without the need for enzyme treatment. Moreover, modifications of temperature-responsive cell culture surfaces were newly carried out to promote the application in cell sheet

tissue engineering, enabling cell attachment/detachment at arbitrary temperature, co-culture of different species of cells on pattered temperature-responsive surfaces, and serum-free cultures on temperature-responsive surfaces immobilized with bioactive molecules. Cell sheet detachment behavior was also improved by using a porous membrane and incorporation hydrophilic moiety into PIPAAm component. Further modification can be expected to realize more complex tissue structure and functions by cell sheet engineering.

References 1.

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13. Yamato M., Akiyama Y., Kobayashi J., Yang J., Kikuchi A., Okano T., 2007. Temperature-responsive cell culture surfaces for regenerative medicine with cell sheet engineering, Pro. Polym. Sci., 32, 1123-1133. 14. Shimizu T., Yamato M., Kikuchi A., Okano T., 2001. Two-dimensional manipulation of cardiac myocyte sheets utilizing temperature-responsive culture dishes augments the pulsatile amplitude, Tissue Eng., 7, 141-151. 15. Kikuchi A., Okuhara M., Karikusa F., Sakurai Y., Okano T., 1998. Twodimensional manipulation of confluently cultured vascular endothelial cells using temperature-responsive poly(N-isopropylacrylamide)-grafted surfaces, J. Biomater. Sci. Polym. Ed., 9, 1331-1348. 16. Heskins M., Guillent J.E., James E., 1968. J. Macromol. Sci. Chem., A2, 1441. 17. Yamada N., Okano T., Sakai H., Karikusa F., Sawasaki Y., Sakurai Y., 1990. Makromol. Chem. Rapid Commun., 11, 571. 18. Okano T., Yamada N., Sakai H., Sakurai Y., 1993. A novel recovery system for cultured cells using plasma-treated polystyrene dishes grafted with poly(Nisopropylacrylamide), J. Biomed. Mater. Res., 27, 1243-1251. 19. Akiyama Y., Kikuchi A., Yamato M., Okano T., 2004. Ultrathin poly(Nisopropylacrylamide) grafted layer on polystyrene surfaces for cell adhesion/detachment control, Langmuir, 20, 5506-5511. 20. Kikuchi A., Okano T., 2005. Nanostructured designs of biomedical materials: Applications of cell sheet engineering to functional regenerative tissues and organs, J. Control. Release, 101, 69-84. 21. Fukumori K., Akiyama Y., Yamato M., Kobayashi J., Sakai K., Okano T., 2008. Acta Biomater., in press. 22. Waymouth C., 1974. To disaggregate or not to disaggregate, injury and cell disaggregation, transient or permanent? In vitro, 10, 97-111. 23. Osunkoya B. O., Mottram F. C., Isoun M. J., 1969. Synthesis and fate of immunological surface receptors on cultured Burkitt lymphoma cells, Int. J. Cancer, 4, 159-165. 24. Revel J. P., Hoch P., Ho D., 1974. Adhesion of culture cells to their substratum, Exp. Cell Res., 84, 207-218. 25. Kushida A., Yamato M., Konno C., Kikuchi A., Sakurai Y., Okano T., 1999. Decrease in culture temperature releases monolayer endothelial cell sheets together with deposited fibronectin matrix from temperature-responsive culture surfaces, J. Biomed. Mater. Res., 45, 355-362. 26. Okano T., Yamada N., Okuhara M., sakai H., Sakurai Y., 1995. Mechanism of cell detachment from temperature-modulated, hydrophilic-hydrophobic polymer surfaces, Biomaterials, 16, 297-303. 27. Hirose M., Kwon O.H., Yamato M., Kikuchi A., Okano T., 2000. Creation of designed shape cell sheets that are noninvasively harvested and moved onto another surface, Biomacromolecules, 1, 377-381.

28. Tang Z., Kikuchi A., Akiyama Y., Okano T., 2007. Novel cell sheet carriers using polyion complex gel modified membranes for tissue engineering technology for cell sheet manipulation and transplantation, React. Func. Polym., 67, 1388-1397. 29. Yamato M., Utsumi M., Kushida A., Konno C., Kikuchi A., Okano T., 2001. Thermo-responsive culture dishes allow the intact harvest of multilayered keratinocyte sheets without dispase by reducing temperature, Tissue Eng., 7, 473480. 30. Nishida K., Yamato M., Hayashida Y., Watanabe H., Okano T., Tano Y., 2004. Corneal reconstruction with tissue-engineered cell sheets composed of autologous oral mucosal epithelium, N. Engl. J. Med., 351, 1187-1196. 31. Nishida K., Yamato M., Hayashida Y., Watanabe K., Maeda N., Watanabe H., Nagai S., Kikuchi A., Tano Y., Okano T., 2004. Functional bioengineered corneal epithellial sheet grafts from corneal stem cells expanded ex vivo on a temperatureresponsive cell culture surface, Transplantation, 77, 379-385. 32. Shiroyanagi Y., Yamato M., Yamazaki Y., Toma H., Okano T., 2004. Urothelium regeneration using viable cultured urothelial cell sheets grafted on demucosalized gastric flaps, BJU Int., 93, 1069-1075. 33. Shimizu T., Yamato M., Isoi Y., Akutsu T., Setomaru T., Abe K., Kikuchi A., Okano T., 2003. Circ. Res. 2002, 90, e40. 34. Bae Y. H., Okano T., Kim S. W., 1990. J. Polym. Sci. Polym. Phys. 1990, 28, 923. 35. Takei Y. G., Aoki T., Sanui K., Ogata N., Okano T., Sakurai Y., 1993. Temperatureresponsive bioconjugates. 2. Molecular design for temperature-modulated bioseparations, Bioconjugate Chem., 4, 341-346. 36. Iwata H., Oodate M., Uyama Y., Amemiya H., Ikada Y., 1991. Preparation of temperature-sensitive membranes by graft polymerization onto a porous membrane, J. Membr. Sci., 55, 119-130. 37. Feil H., Bae Y. H., Feije J. n, Kim S. W., 1993. Effect of comonomer hydrophilicity and ionization on the lower critical solution temperature of N-isopropylacrylamide copolymers, Macromolecules, 26, 2496-2500. 38. Tsuda Y., Kikuchi A., Yamato M., Okano T., 2004. Control of cell adhesion and detachment using-temperature and thermoresponsive copolymer grafted culture surfaces, J. Biomed. Mater. Res. 2004, 69A, 70-78. 39. Bhatis S., Balis U., Yarmush M., Toner M., 1999. FASEB J., 13, 1883. 40. Tsuda Y., Kikuchi A., Umezu M., Okano T., 2005. The use of patterned dual thermoresponsive surfaces for the collective recovery as co-cultured cell sheets, Biomaterials, 26, 1885-1893. 41. Ebara M., Yamato M., Naga S. i, Sakai K., Okano T., 2004. Incorporation of new carboxylate functionalized co-monomers to temperature-responsive polymer-grafted cell culture surfaces, Surf. Sci., 570, 134-141. 42. Ebara M., Yamato M., Okano T., 2004. Temperature-responsive cell culture surfaces enable "on-off" affinity control between cell integrins and RGDS ligands, Biomacromolecules, 5, 505-510.

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recovery by co-grafting of PEG with PIPAAm onto porous cell culture membranes. Biomaterials, 24, 1223-1232.

PART II

Tissue Engineering

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Chapter 5 Adipose Tissue-derived Mesenchymal Stem Cells have Lower Osteogenic Potential than Bone Marrow-derived Mesenchymal Stem Cells Yoshihiro Katsube, Ousuke Hayashi, Motohiro Hirose, and Hajime Ohgushi Research Institute for Cell Engineering, National Institute of Advanced Industrial Science and Technology, Amagasaki, Japan

1

Introduction

Mesenchymal stem cells (MSCs) are able to differentiate into various functional cell types of mesodermal tissues [1–3]. With rapid advancements in tissue engineering, a multitude of applications for cultured MSCs in the construction of regenerative tissues have been demonstrated [3–5]. We have used human MSCs from patients’ bone marrow for treating various diseases centering on the bone diseases since 2001 [6–8]. MSCs with or without various scaffolds have been transplanted into more than 80 patients. Recently, multipotent mesenchymal cells were also found in adipose tissue [9]. These adipose tissue– derived MSCs (AMSCs; also referred to as ADSCs) have been shown to differentiate into many lineages, including chondrogenic [10–12], osteogenic [13–16], myogenic [17, 18], cardiomyogenic [19, 20], neurogenic [21–23], angiogenic [24–27], and hepatic [28] lineages. The chief advantage of adipose tissue includes the huge amount of fatty tissue available for cosmetic surgery, such as liposuction. It is expected to be a practical cell source for tissue

engineering and regenerative medicine [29]. However, it is not well investigated whether AMSCs have a similar potential for osteogenesis as do MSCs from other types of tissue, especially bone marrow–derived MSCs (BMSCs). In our previous study, we found the excellent bone forming capability of BMSCs compared with AMSCs [30]. In this review, we discuss in vitro as well as in vivo bone-forming potential of BMSCs and AMSCs on the basis of our latest study.

F344 rat

Adipose tissue

Collagenase digestion

Femur

Centrifugation

Adipocytes

Stromal-vascular fraction

Bone Marrow

7days culture Figure 1. Isolation methods of BMSCs and AMSCs.

2

Isolation and Culture of Stem Cells

Adipose tissue consists of mature adipocytes, preadipocytes, stem cells, (AMSCs), fibroblasts, vascular smooth muscle cells, endothelial cells and blood cells [31]. In generally, to isolate AMSCs, adipose tissue was first treated with collagenase and centrifuged to separate mature adipocytes and other cells. These cells without mature adipocytes are referred to as the stromal and vascular cell fraction (SVF). AMSCs are assumed to exist in the SVF and thought to be a heterogeneous cell population. To compare the bone differentiation capabilities of BMSCs and AMSCs, we isolated MSCs according to the generalized method, as illustrated in Fig. 1. We obtained femora, inguinal subcutaneous adipose tissue, and dermis from Fischer 344 7-week-old male rats. Bone marrow was flushed out using 10 mL of culture medium (minimal essential medium [MEM] with Earle’s salts and LGln) containing 15% FBS and antibiotics (100 U/mL penicillin G, 100 μg/mL streptomycin sulfate, and 0.25 μg/mL amphotericin B) from the femora and then dispersed through an 18-gauge needle. The culture medium including the bone marrow was filtered through a sterile 40 μm filter and centrifuged at 400 × g for 5 minutes. After centrifugation, the supernatant portion of the medium was discarded and the pellet was used as a source of BMSCs. Adipose tissue and dermis were each treated with 10 mL of PBS containing 3 mg/mL collagenase at 37°C for 1 hour. They were then centrifuged at 400 × g for 5 minutes after filtration by 40 μm filter. After centrifugation, the supernatants of the adipose tissue and dermis were discarded and their pellets were used as cell sources of AMSCs and fibroblasts, respectively. Previously, we reported that dermal fibroblasts did not exhibit mineralization and osteogenic ability [32]. Therefore, we used fibroblasts as a negative control for all further analyses. Three kinds of cell sources were placed into culture dishes for primary culture for 7 days. All cultures were maintained in a humidified atmosphere of 95% air and 5% CO2 at 37°C; the culture medium was renewed three times per week. The cell surface antigens of these primary cultured cells were analyzed by using a flow cytometer (FACS Calibur, BD, N.J., USA). These cells were positive for mesenchymal markers (CD29, -90) and negative for hematopoietic marker (CD45), though BMSCs contained some fraction of CD45 positive hematopoietic cells (Fig. 2). These hematopoietic cells express CD11b and MHC class II (data not shown) and are thought to be adherent hematopoietic cells, like macrophages. These results agree with previous reports that cell surface antigen expression of AMSCs is similar to that of other MSCs [31].

CD29

CD45

CD90

BMSCs

AMSCs

Fibroblasts

Figure 2. Cell surface antigen expressions of MSCs. Cells were stained with FITClabeled antibodies against the indicated antigens (black areas). Black lines indicate histograms of control antibody-stained cells.

3

Colony-Forming Capabilities of MSCs

To examine the frequency of stem cells in each cell source, we observed the colony-forming capabilities of MSCs. Cell sources of BMSCs were seeded at a rate of 1.0 × 104, 105, or 106 nucleated cells/cm2 on 20 cm2 dishes. Cell sources of AMSCs and fibroblasts were seeded at a rate of 1.0 × 102, 103, or 104 cells/cm2 on 20 cm2 dishes. After 6 days of culture, the colonies from each dish were counted and CFU frequencies (colony-forming units per seeded cell number) of cell sources were calculated. The observations of culture dishes are presented in Fig. 3. It was not possible to count the colonies of BMSCs when seeded at more than 1.0 × 106 cells/cm2 and those of the others when seeded at more than 1.0 × 104 cells/cm2 because they proliferated to near confluence. The number of CFUs of bone marrow was significantly lower compared with those of adipose tissue and dermis. We calculated that the frequency of CFU of the bone marrow was 100 times less than in adipose tissue and dermis [30].

1×102 cells/cm2

1×103 cells/cm2

1×104 cells/cm2

1×105 cells/cm2

1×106 cells/cm2

Bone marrow (BMSCs)

Adipose tissue (AMSCs)

Dermis (Fibroblasts) Figure 3. Colony-forming capabilities of MSCs. Photographs of various cell sources seeded at each cell densities are shown. Representative culture dishes were stained with Giemsa solution. Blue dots indicate colonies.

Sakaguchi et al. demonstrated that the colony number of human BMSCs in primary culture was the lowest among the various MSCs [33]. The differences in the colony-forming cells might be caused by the variety of cell types residing in each tissue. AMSCs and fibroblasts were derived from solid tissue. In contrast, bone marrow consists of a huge amount of floating cells including hematopoietic cells and a small amount of adherent cells on the bone surface.

4

Osteogenic Differentiation of MSCs

4.1 In Vitro Analysis We determined that 1.0 × 106 cells/cm2 dish for cell sources of BMSCs and 1.0 × 104 cells/cm2 dish for those of AMSCs and fibroblasts were optimal cell densities of primary culture. To evaluate the osteogenic differentiation capability of MSCs, MSCs were subcultured in the presence or absence of dexamethasone (Dex). As we previously reported, Dex induces mesenchymal cells to differentiate into osteoblasts and stimulates the development of mineralized matrices on a variety of materials [34, 35]. After primary culture, MSCs were harvested and seeded in

tissue culture plates at a cell density of 1.0 × 104/cm2. Under undifferentiated conditions, medium supplemented with 10 mM β-glycerophosphate was renewed three times per week. Under osteogenic differentiation conditions, the medium was further supplemented with 0.28 mM ascorbic acid-2-phosphate and 10 nM Dex. To visualize the mineralized extracellular matrix of subcultured cells, 1 μg/mL of calcein was added to the culture medium according to our previous report [32]. We could observe mineralized deposition under phase contrast microscopy after 2-week subculture. In the presence of Dex (osteogenic differentiation condition), BMSCs showed mineral deposition around the cells to form nodular aggregates [30]. Previously, we reported a novel quantitative method for detecting mineralization by cultured cells utilizing a calcium-binding fluorescent dye, calcein. Calcein has been found to be specifically incorporated and deposited into extracellular bone matrix, as evidenced by costaining with alizarin red S [36]. BMSCs showed strong calcein fluorescence in the matrix under fluorescence microscopy. On the other hand, AMSCs did not show any fluorescence. Fluorescence of calcein was determined using a Typhoon 8600 image analyzer (GE Healthcare, Little Chalfont, UK) (Fig. 4). To quantitation of the matrix, fluorescence intensities of each culture were measured. The intensity of BMSCs was 300 times as high as those of AMSCs and fibroblasts [30].

BMSCs

AMSCs

Fibroblasts

Calcein (+)

Calcein (-)

Dex (-)

Calcein (+)

Calcein (-)

Dex (+)

Figure 4. Laser-scanning images of MSCs under the osteogenic differentiation condition. Black dots indicate fluorescence of calcium-deposited calcein.

Dex (-)

Dex (+)

BMSCs

Dex (-)

Dex (+)

AMSCs

Fibroblasts Figure 5. ALP staining of MSCs under the osteogenic differentiation condition. Phase contrast microscopic images show ALP positive cells (red-stained cells). Scale bar = 200μm.

Alkaline phosphatase (ALP) activity is recognized as an early osteoblastic marker. ALP activities of the BMSCs for 2 weeks were significantly greater than those of AMSCs and fibroblasts [30]. ALP staining revealed that the cells around the mineralized matrix are strongly ALP positive (Fig. 5). Osteocalcin is a bone-specific protein and has been used as a late marker of osteogenic differentiation. The osteocalcin content of BMSCs was significantly greater than those of AMSCs and fibroblasts [30]. We also conducted real-time quantitative polymerase chain reaction (PCR) of ALP (Fig. 6A) and osteocalcin mRNAs (Fig. 6B). Both mRNA expression levels of the BMSCs under osteogenic conditions were greater than those of others. These in vitro results indicate that BMSCs have high osteogenic differentiation capability compared to AMSCs. Referring to previous reports that investigated the comparison of osteogenic ability between BMSCs and AMSCs, De Ugarte et al.reported no significant difference in osteogenic ability between human BMSCs and AMSCs [37]. In their report, quantification of ALP activity of 3-week subcultured BMSCs and AMSCs showed no significant difference. However, the ALP activity in their report was below 0.35 nmol p-nitrophenol release/μg DNA/minute. On the other hund, our previous study detected 100 times more activity of subcultured BMSCs (36 nmol p-nitrophenol release/μg DNA/minute) [30]. In contrast, Gun-II et al. reported that human AMSCs have inferior osteogenic capabilities compared to BMSCs [38]. They quantified the percentage of ALP-positive

cells under the osteogenic differentiation condition for both 2 and 3 weeks. The ALP-positive cells of the BMSCs were greater than those of the AMSCs.

ALP/GAPDH

A

1.2 1 0.8 0.6 0.4 0.2 0

Osteocalcin/GAPDH

B

BM SCs

AM SCs

Fibroblasts

BM SCs

AM SCs

Fibroblasts

1.2 1 0.8 0.6 0.4 0.2 0

Figure 6. ALP (A) and osteocalcin (B) gene expression analysis by real-time quantitative PCR. Each mRNA value was normalized to that of the housekeeping gene, GAPDH. Each result was expressed as relative value for BMSC sample as 1.

Sugiura et al. reported that the osteogenic ability of rat MSCs decreases with passaging [39]. Their results showed that ALP activity of rat BMSCs at passage 1 was almost five times that of rat BMSCs at passage 3 under the osteogenic differentiation condition. Therefore, when comparing the cellular activity of different cell sources, we should regard the passage time. The previous reports of low osteogenic ability of BMSCs used more passaged cells (passage 3 or more) than our study (passage 1). Furthermore, these previous reports dealt only with in vitro culture experiments. Therefore, we did in vivo implantation experiments, which confirmed the bone-forming ability of both BMSCs and AMSCs.

4.2 In Vivo Analysis To observe and quantify the in vivo bone-forming capability of primary cultured MSCs, we implanted MSCs/hydroxyapatite (HA) composites. We previously reported that BMSCs/HA composites can show a high level of in vivo boneforming ability [40–42]. According to the methods used in these reports, we prepared the composites. We used a HA disk (5 mm in diameter and 2 mm thick with 50% porosity, Cellyard; Pentax, Tokyo, Japan) as scaffold. Primary cultured BMSCs, AMSCs and fibroblasts were harvested and suspended with culture medium at a cell density of 1.0 × 107/mL. Each cell suspension was poured into a well of a 96-well tissue culture plate. A HA disk was soaked into each cell suspension and maintained in a humidified atmosphere of 5% CO2 at 37°C overnight for cell adhesion. We implanted subcutaneously into the back of the syngeneic recipient rat and the implants were recovered after 6 weeks. The capability of the composites for new bone formation was evaluated by micro-computed tomography. There were more massive areas of newly formed bone in the BMSCs/HA composite than in the other HA composites [30]. After the micro-CT analysis, the composites were histologically analyzed. They were decalcified, cut into sections and stained with hematoxylin and eosin. The histology of implants demonstrated that the BMSCs/HA composites showed newly formed bone together with active osteoblasts lining the pore areas of the composites (Fig. 7). In contrast, no bone formation or osteoblasts were seen in other composites. As seen with in vitro culture experiments, the in vivo ability was not observed in AMSCs.

BMSCs

AMSCs

Fibloblasts

Figure 7. Histological sections (hematoxylin and eosin stain) of MSCs/HA composites after 6-week implantation. Bone formation is indicated by arrowheads (red-stained area). Scale bar = 200μm.

Few reports analyzed the in vivo osteogenic differentiation capability of AMSCs. Lee et al. reported that osteogenic induced AMSCs in vitro could form

newly bone in vivo, however bone formation ability of AMSCs was not compared to that of BMSCs [43]. Cowan et al. compared osteogenic potential of mouse BMSCs and AMSCs by using a mouse calvarial defects model [44]. They reported that osteogenic capability of AMSCs was equivalent to that of BMSCs. However, this model is orthotopic implantation model and different from ectopic implantation model used in our study. Therefore, bone formation by recipient cells can not be negated, though the factor that promotes boneformation seems to be released from donor AMSCs.

5

Conclusions

Regenerative medicine and tissue engineering have been rapidly developing for practical and global use. Recently, adipose tissue–derived MSCs (AMSCs) have been shown to differentiate into many lineages, and they are considered as a potent cell source for tissue regeneration. However, our comparative study of the osteogenic potential using MSCs from bone marrow and adipose tissue demonstrated the superior bone forming capability of bone marrow than adipose tissue. It is well known that MSCs reside in many types of tissue and are able to differentiate into various functional cells. However, the potentials of each MSCs from different origin are thought to be different from one another. To achieve advances in regenerative medicine, knowledge of the potentials of various MSCs towards various cell/tissue regeneration is essential. Results of our present study indicate that bone marrow-derived MSCs (BMSCs) could be ideal candidates for their utilization in practical bone tissue engineering.

Acknowledgments This work was supported by grants from the New Energy and Industrial Technology Development Organization.

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17. Mizuno H., Zuk P.A., Zhu M., Lorenz H.P., Benhaim P., Hedrick M.H., 2002. Myogenic differentiation of human processed lipoaspirate cells, Plast Reconstr Surg 109, 199–209. 18. Rodriguez A.M., Pisani D., Dechesne C.A., Turc-Carel C., Kurzenne J.Y., Wdziekonski B., Villageosis A., Bagnis C., Breittmayer J.P., Groux H., Aihaud G., Dani C., 2005. Transplantation of a multipotent cell population from human adipose tissue induces dystrofin expression in the immunocompetent mdx mouse, J Exp Med 201, 1397–1405. 19. Planat-Bernard V., Menard C., Andre M., Puceat M., Perez A., Garcia-Veedugo J.M., Penicaud L., Casteilla L., 2004. Spontaneous cardiomyocyte differentiation from adipose tissue stroma cells, Circ Res 94, 223–229. 20. Strem B.M., Zhu M., Alfonso Z., Daniels E.J., Schreiber R., Begyui R., Maclellan W.R., Hendrick M.H., Fraser J.K., 2005. Expression of cardiomyocytic markers on adipose tissue-derived cells in a murine model of acute myocardial injury, Cytotherapy 7, 282–291. 21. Safford K.M., Hiock K.C., Safford S.D., Halvorsen Y.D., Wilkinson W.O., Gimble J.M., Rice H.E., 2002. Neurogenic differentiation of murine and human adiposederived stromal cells, Biochem Biophys Res Commun 294, 371–379. 22. Ashjian P.H., Elbarbary A.S., Edmonds B., De Ugarte D., Zhu M., Zuk P.A., Lorenz H.P., Benhaim P., Hedrick M.H. 2003. In vitro differentiation of human processed lipoaspirate cells into early neural progenitors, Plast Reconstr Surg 111, 1922–1931. 23. Kang S.K., Lee D.H., Bae Y.C., Kim H.K., Baik S.Y., Jung J.S., 2003. Improvement of neurological deficits by intracerebral transplantation of human adipose tissue–derived stromal cells after cerebral ischemia in rats, Exp Neurol 183, 355–366. 24. Miranville A., Heeschen C., Sengenes C., Curat C.A., Busse R., Bouloumie A., 2004. Improvement of postnatal neovascularization by human adipose tissue– derived stem cells, Circulation 110, 349–355. 25. Rehman J., Traktuev D., Li J., Merfeld-Clauss S., Temm-Grove C.J., Bovenkerk J.E., Pell C.L., Johnstone B.H., Considine R.V., March K.L., 2004. Secretion of angiogenic and antiapoptotic factors by human adipose stromal cells, Circulation 109, 1292–1298. 26. Planat-Benard V., Silvestre J.S., Cousin B., Andre M., Nibbelink M., Tamarat R., Clergue M., Manneville C., Saillan-Barreau C., Duriez M., Tedgui A., Levy B., Penicaud L., Casteilla L., 2004. Plasticity of human adipose lineage cells toward endothelial cells: physiological and therapeutic perspectives, Circulation 109, 656– 663. 27. Cao Y., Shu Z., Liao L., Meng Y., Han Q., Zhao R.C., 2005. Human adipose tissue– derived stem cells differentiate into endothelial cells in vitro and improve potential neovascularization in vivo, Biochem Biophys Res Commun 332, 370–379.

28. Seo M.J., Suh S.Y., Bae Y.C., Jung J.S., 2005. Differentiation of human adipose stromal cells into hepatic lineage in vitro and in vivo, Biochem Biophys Res Commun 328, 258–264. 29. Zuk P.A., Zhu M., Ashjian P., De Ugarte D.A., Huang J.I., Mizuno H., Alfonso Z.C., Fraser J.K., Benhaim P., Hedrick M.H., 2002. Human adipose tissue is a source of multipotent stem cells, Mol Cell Biol 13, 4279–4295. 30. Hayashi O., Katsube Y., Hirose M., Ohgushi H., Ito H., 2008. Comparison of osteogenic ability of rat mesenchymal stem cells from bone marrow, periosteum, and adipose tissue, Calcif Tissue Int. 82, 238–247. 31. Schäffler A., Büchler C., 2007. Concise review: adipose tissue-derived stromal cells–Basic and clinical implications for novel cell-based therapies, Stem cells 25, 818–827. 32. Uchimura E., Machida H., Kotobuki N., Kihara T., Kitamura S., Ikeuchi M., Hirose M., Miyake J., Ohgushi H., 2003. In-situ visualization and quantification of mineralization of cultured osteogenic cells, Calcif Tissue Int 73, 575–583. 33. Sakaguchi Y., Sekiya I., Yagishita K., Muneta T., 2005. Comparison of human stem cells derived from various mesenchymal tissues: superiority of synovium as a cell source, Arthritis Rheum 52, 2521–2529. 34. Ohgushi H., Dohi Y., Katsuda T., Tamai S., Tabata S., Suwa Y., 1996. In vitro bone formation by rat marrow cell culture, J Biomed Mater Res 32, 333–340. 35. Ohgushi H., Dohi Y., Yoshikawa T., Tamai S., Tabata S., Okunaga K., Shibuya T., 1996. Osteogenic differentiation of cultured marrow stromal stem cells on the surface of bioactive glass ceramics, J Biomed Mater Res 32, 341–348. 36. Hirose M., Kotobuki N., Machida H., Uchimura E., Ohgushi H., 2003. Quantitative monitoring of in vitro mineralization process using fluorescent dyes, Key Eng Mater 240, 715–718. 37. De Ugarte D.A., Morizono K., Elabarbary A., Alfonso Z., Zuk P.A., Zhu M., Dragoo J.L., Ashjian P., Thomas B., Benhaim P., Chen I., Fraser J., Hedrick M.H., 2003. Comparison of multi-lineage cells from human adipose tissue and bone marrow, Cells Tissues Organs 174, 101–109. 38. Gun-II I., Shin Y.W., Lee K.B., 2005. Do adipose tissue–derived mesenchymal stem cells have the same osteogenic and chondrogenic potential as bone marrow–derived cells? Osteoarthritis Cartilage 13, 845–853. 39. Sugiura F., Kitoh H., Ishiguro N., 2004. Osteogenic potential of rat mesenchymal stem cells after several passages, Biochem Biophys Res Commun 316, 233–239. 40. Ohgushi H., Goldberg V.M., Caplan A.I., 1989. Heterotopic osteogenesis in porous ceramics induced by marrow cells, J Orthop Res 7, 568–578. 41. Shimaoka H., Dohi Y., Ohgushi H., Ikeuchi M., Okamoto M., Kudo A., Kirita T., Yonemasu K., 2004. Recombinant growth/differentiation factor-5 (GDF-5) stimulates osteogenic differentiation of marrow mesenchymal stem cells in porous hydroxyapatite ceramic, J Biomed Mater Res A 68, 168–176.

42. Noshi T., Yoshikawa T., Ikeuchi M., Dohi Y., Ohgushi H., Horiuchi K., Sugimura M., Ichijima K., Yonemasu K., 2000. Enhancement of the in vivo osteogenic potential of marrow/hydroxyapatite composites by bovine bone morphogenetic protein, J Biomed Mater Res 52, 621–630. 43. Lee J.A., Parrett B.M., Conejero J.A., Laser J., Chen J., Kogon A.J., Nanda D., Grant R.T., Breitbart A.S., 2003. Biological alchemy: engineering bone and fat from fat-derived stem cells, Ann Plast Surg. 50, 610–617. 44. Cowan C.M., Shi Y.Y., Aalami O.O., Chou Y.F., Mari C., Thomas R., Quarto N., Contag C.H., Wu B., Longaker M.T., 2004. Adipose-derived adult stromal cells heal critical-size mouse calvarial defects, Nat Biotechnol. 22, 560–567.

Chapter 6 Amelogenin Overexpression in Tooth Development Akiyoshi Taniguchi and Liming Xu Biomaterials Center, National Institute for Materials Science, Tsukuba, Japan

1

Introduction

The Japanese population has aged rapidly for the past several years. According to population statistic estimation, the percentages of the aged population in Japan reached to 20% (2005), and 65 years and older will be more than double in 2030. Under such the societal condition, the maintenance of quality of life (QOL) of aged population is an important subject for Japanese government. Oral health is very important factor for QOL of the aged population. The relationship between the status of tooth preservation and QOL had been reported by Yoshida, et al. (1, 2). In their survey, they focused on the relationship among the number of existing teeth and life environment, health status, activities of daily living. The results indicated that the presence of teeth was very closely related to one’s daily activities. Periodontal disease likes other lifestyle-related disease, such as diabetes or hypertension, its development relate to lifestyle and with advancing age. And periodontal disease is one of the factors resulted in tooth loss for aged population. Recent development of tissue engineering technology will be possible to regenerate periodontal tissue or tooth. However, it is very difficult to generate complete periodontal tissue or tooth with reproducibility. The investigation of molecular mechanism of tooth development is very important for tooth regeneration. In this study, I focused on amelogenin, which is one of enamel matrix protein. Amelogenin proteins are produced by ameloblasts in the early stage of tooth

development. The outermost layer is the enamel, the bulk of the tooth is dentin, and there is a coating of cementum on the outer layer of the roots of tooth beneath the periodontal tissue. The enamel develops from dental epithelial cells, and the dentin develops from mesenchymal cells. The interactions between dental epithelial and mesenchymal cells are very important for the construction of tooth. Tooth development in human begins from the 6th week of intrauterine life. First, oral epithelium hollows toward to mesenchymal tissue, and becomes thick. At the 8th to 9th week of intrauterine life, the mesenchyme tissue nearby the thickened epithelium begins proliferation, is called bud stage. At the 9th to10th week of intrauterine life, the epithelium continues to proliferate to form enamel organ and becomes cap shape, is called cap stage. And at the 14th week of intrauterine life, the enamel organ begins to separate into inner enamel epithelium, enamel marrow and outer enamel epithelium, is called bell stage. By following epithelial-mesenchymal interactions, dental papilla cells differentiate into pre-odentoblasts, odontoblasts that secrete dentin matrix proteins (major constituents are collagen), and inner enamel epithelium differentiates into preameloblasts, ameloblasts that secrete enamel matrix proteins (major constituents are amelogenin). Subsequently, following matrix-mediated mineralization, the dentin and enamel were formed. Amelogenin is one of the major enamel matrix proteins. In the early stages of amelogenesis, enamel consists of 20-30% proteins. Amelogenins are the most abundant protein constituents that account for more than 90% of enamel matrix proteins in developing enamel [1]. It was well known that amelogenins form nanospheres by self-assembly under physiological conditions (physiological temperature and pH) both in vivo and in vitro [2-7]. The nanosphere structures of amelogenin have important role for matrix-mediated enamel mineralization. Many investigations have demonstrated that amelogenins play a major role in the structural organization of the mineral within the developing enamel and might also regulate the nucleation and growth pattern of the enamel hydroxyapatite crystals [8-13]. Immature enamel contains a complex mixture of amelogenin polypeptides, primarily due to the combined effects of alternative RNA splicing [14-20] and proteolytic processing [21]. Multi-type alternative spliced mRNAs generated by the amelogenin gene have been identified in various species. In mice, nine splicing products of amelogenin have been identified and detected as protein. Most of mouse amelogenins contain exon 7 [14, 15, 22] in addition to the recently identified exon 8/9 [23]. In rat, the amelogenin mRNAs containing exon 7 or exon 8 have been identified [18, 19, 24 ]. Full-length amelogenin, which contains exon 8, and leucine-rich amelogenin peptide or LRAP generated by spliced out of exon 6ABC are major splicing products (as shown in Fig. 1).

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Figure 1. Schematic representations of two major amelogenin alternative spliced patterns in rat. The illustrations are based on the ref.19 with some modifications. The intron-exon structure of the chromosomal copy of the human amelogenin gene is shown at the top of the Figure. The exons are numbered above, while the number of nucleotides per exon is indicated below the bar. The alternatively spliced pattern for each rat amelogenin cDNA is illustrated below the human sequence.

The expression of amelogenin gene is regulated at the transcriptional level with some transcriptional activators and repressors. In the mouse amelogenin promoter analysis, it is suggested that C/EBPα plays a key role in the developmentally regulated expression of the amelogenin gene (25). The other investigation indicated that Msx2 interfered with binding of C/EBPα to its cognate site on the mouse amelogenin minimal promoter by protein-protein interaction (26). Xu. Y., et al. demonstrated that co-transfection of C/EBPα and NF-Y synergistically increased the promoter activity, suggesting that C/EBPα and NF-Y synergistically activate the mouse amelogenin gene and can contribute to its physiological regulation during amelogenesis (27). However, it is difficult to understand that the tissue and stage specific expressions of amelogenin gene were controlled by above described transcriptional regulations alone. The biological functions of amelogenin in processes of cell-differentiation have become widely appreciated. Previous studies suggested that in epithelialmesenchymal interactions, amelogenin exhibits specific biological effects on the mesenchyme (22, 24). However, effects of amelogenin on dental epithelial cells during differentiation have not been investigated. To understand the effects of amelogenin on dental epithelial cells and the regulation mechanisms of amelogenin gene expression in tooth development, Full-length mouse amelogenin protein was expressed by using baculovirus-insect cell expression system. The effects of amelogenin protein on differentiation marker gene expressions in dental epithelial cell lines were evaluated. The regulation mechanisms of amelogenin gene expression in tooth development

were investigated. The results from this study are helpful to better understand the biological functions and gene regulation of amelogenin in tooth development.

2

The Regulation Mechanisms of Amelogenin mRNA

In this study, amelogenin protein biological activity on the dental epithelial cells and the regulation mechanisms of amelogenin mRNA were investigated. The full-length mouse amelogenin protein (M180) was expressed and the biological activity of amelogenin on ameloblast like cell (HAT-7) differentiation was examined. The results showed that recombinant mouse amelogenin enhanced the expression level of endogenous amelogenin mRNA in HAT-7 cells (Fig. 2) (28). Next, the effects of recombinant amelogenin protein on amelogenin promoter activity and amelogenin mRNA stabilization were analyzed. The results indicated that amelogenin protein increased amelogenin mRNA stabilization, without any changes in amelogenin promoter activity. Furthermore, the result showed that exogenous FITC-labeled amelogenin protein was takenup by dental epithelial cells and co-localized with ER structure in the cytoplasm. Northwestern analysis showed that amelogenin proteins bind directly to themselves mRNA in vitro, suggesting that amelogenin protein acts as a transacting protein that specifically binds to mRNA.

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Figure 2. Effects of rmAmelogenin on mRNA expression in HAT-7 cells. Three days after treatment of HAT-7 cells with 3 μg/ml amelogenin, induction of mRNA for amelogenin (Ameg), ameloblastin (Ameb), Keratin 14, Fibroblast growth factor receptor 2b (FGFR2b), alkaline phosphatase (ALP), osteopontin (OPN), bone sialoprotein (BSP) and bone morphogenetic protein 2 (BMP-2), were determined by real-time PCR. The values are fold-induction compare to no-rmAmelogenin treatment control. The results were reproduced in three separate experiments. Error bars indicate the standard deviation of mean changes.

3

The Molecular Mechanism of Enhanced Amelogenin mRNA Stabilization

The molecular mechanism of enhanced amelogenin mRNA stabilization was examined. To identify the cis-regulatory region within amelogenin mRNA, various reporter systems using a deletion series of reporter plasmids were tested. A deletion at exon 6ABC of amelogenin mRNA resulted in a 2.5-fold increase in the amelogenin mRNA expression level, compared with that of full-length mRNA. The half-life analysis of mRNA indicated that the half-life of exon 6ABC deletion mutant is longer over 2-fold compared with that of full-length amelogenin. These results indicated that a cis-element exists in exon 6ABC of amelogenin mRNA. Moreover, recombinant mouse amelogenin protein extended the half-life of full-length amelogenin mRNA, but did not significantly alter the half-life of exon 6ABC-deletion mutant mRNA. These results suggested that exon 6ABC of amelogenin mRNA is important involved in amelogenin mRNA destability and amelogenin protein-mediated mRNA stabilization (Fig. 3) (29). The splice products produced by alternative spliced out of exon 6ABC are known as leucine-rich amelogenin peptides or LRAP, and have signaling effects on cells. This finding also suggests that the regulation of LRAP expression differs from the regulation of full-length amelogenin expression.

4

The Effects of LAMP1 and LAMP3 on M180 Amelogenin Uptake, Localization and Amelogenin mRNA Induction by Amelogenin Protein

We demonstrated that the uptake of M180 amelogenin protein in HAT-7 results in increased levels of amelogenin mRNA through enhanced mRNA stabilization. To determine the processes involved in the uptake of extracellular M180 amelogenin by cells and in amelogenin intracellular trafficking in the amelogenin protein-mediated amelogenin mRNA expression pathway, we investigated the effects of LAMP1 and LAMP3, which are candidate M180 amelogenin receptors, on M180 amelogenin uptake, localization and amelogenin mRNA induction by amelogenin protein, using anti-LAMP-1 and anti-LAMP-3 antibodies and siRNA analysis. The results indicate that LAMP3 blocking by anti-LAMP-3 decreases M180 amelogenin uptake, but does not affect amelogenin mRNA induction by amelogenin protein, suggesting that LAMP3 is related to amelogenin degradation.

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Figure 3. The destabilizing sequence located in the ORF. (Left) Schematic representation of reporter plasmids. Transcription was driven by the constitutive CMV promoter upstream of the amelogenin cDNA. Partial luciferase sequence (Luc, 309 bp ; represented as darkened box) was ligated between the 5’-UTR and ORF (pL-FL), ORF (pL-dORF), exon 2,3,5,6ABC (pL-d2,3,5,6ABC), exon6ABC(pL-d6ABC) and exon 6D(pL-d6D) were deleted. All plasmids included the amelogenin poly-adenylation signal (788–825 bp). (Right) Luciferase reporter mRNA levels in each reporter plasmid. HAT-7 cells were transiently transfected with each reporter plasmid separately, co-transfected with the renilla luciferase reporter plasmid (pRL-CMV) and cultured for 24 hr. The levels of luciferase reporter mRNAs were determined by quantitative real-time PCR and normalized against the pRL-CMV level. Results are representative of six experiments. Error bars indicate the standard deviation of mean changes. **p < 0.001 versus pL-FL.

Down-regulation by siRNA of LAMP1, which is the receptor for LRAP, does not affect M180 amelogenin uptake, localization or amelogenin mRNA induction by amelogenin protein. Thus, while LAMP1 is the specific receptor for LRAP, it is not a receptor for M180 amelogenin. These findings will aid further research into the understanding of M180 amelogenin function and expression (30). In this study, a unique auto-regulation mechanism of amelogenin was demonstrated using in vitro experimental systems. Figure 4 shows the summary of this work. Before exogenous recombinant amelogenin protein addition, endogenous amelogenin expression levels of both mRNA and protein are lower (Fig. 4, left panel) due to the instability of amelogenin mRNA. After recombinant amelogenin addition, extrocellular amelogenin proteins are re-taken up by cells (Fig. 4. right panel). In the cytoplasm, re-uptaken amelogenin proteins enhanced mRNA expression through directly binding to exon 6ABC region of mRNA and increasing amelogenin mRNA stabilization. This unique auto-regulation

mechanism is helpful for ameloblasts to secrete large amount of amelogenin proteins that occupied 90% of enamel matrix proteins. Before exogenous amelogenin addition

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Figure 4. Views of secretion and retakeup of amelogenin proteins. The left panel represents the transcription, translation and secretion processing of amelogenin before exogenous amelogenin protein addition. The right panel represents a positive feedback of amelogenin proteins after exogenous amelogenin protein addition. N: nucleus; C: cytoplasm.

5

Discussions

Ameloblasts, which develop from a layer of short epithelial cells, are known as inner enamel epithelium (IEE). Inner enamel epithelium interacts with an adjacent layer of ectomesenchyme cells from the dental papilla. Dental papilla differentiates into odontoblasts and secretes dentin layer that underlies enamel. Following intercellular signaling, inner enamel epithelium express amelogenin mRNA and secrete weakly amelogenin protein. Following continue differentiate, the overlying epithelial cells were elongated, and develop into preameloblasts. The preameloblasts secrete moderately amelogenin. A further functional differentiation of the ameloblasts implies with-drawal from the cell cycle, elongation and polarization, and initiation and/or up-regulation of the synthesis of enamel matrix components (e.g. amelogenin) at both of mRNA expression and protein secretion levels (31).

The data from this study indicated that ameloblasts are able to dramatically increase production of amelogenin in an autocrine fashion. Indeed, ameloblasts must secrete a large amount of amelogenin for enamel formation, considerably more than the other enamel matrix protein, during the short periods of tooth development. Amelogenin production in IEE cells was much less than that in differentiated ameloblasts during mouse incisor development. When dentin matrix is formed between the inner enamel epithelium and mesenchymal cells, amelogenin accumulates at the proximal side of the inner enamel epithelium. The deposition of amelogenin helps IEE cells reuptake amelogenin into the cytoplasm. Consequently, IEE cells rapidly accelerate production of amelogenin (Fig. 5).

Inner Enamel Epithelium (IEE)

Over Expression Enamel Dentin

Mesenchymal cells Figure 5. The illustration of early stages of tooth secretory cell development. Enamel organs contain inner enamel epithelium facing the dental papilla. Following signal interaction, inner enamel epithelium develops into preameloblasts. Preameloblasts differentiate into secretory ameloblasts, which mainly produce amelogenin. The expression of amelogenin mRNA and protein are indicated at the top of the figure. In inner enamel epithelium and preameloblasts stages, amelogenins are weakly expressed, and they easily diffused to the neighborhood. However, when dentin formation start, as show in odontoblasts stage, amelogenin proteins were accumulated nearby ameloblasts and were easily retaken-up by the ameloblasts.

Several laboratories have shown that the amelogenin was detected in the cytoplasm lysosome and stippled material-like substance adjacent to the coated pits of plasma membrane in presecretory ameloblasts and odentoblasts, suggesting that it is likely that amelogenin might be taken up by an endocytic pathway (32-34). Recently, amelogenin binding proteins have been cloned (35, 36). Wang et al. have identified integral membrane protein, CD63 that interact with amelogenin (35). Tompkins et al. also have shown LAMP-1 interacts with

LRAP (35). These results suggested that the amelogenin proteins are taken up by endocytic pathway of amelogenin into the HAT-7 cells via these binding molecules. However, our findings suggest that LAMP1 is the specific receptor for LRAP, it is not a receptor for M180 amelogenin (30). The data have shown exogenous FITC-labeled amelogenin co-localized with ER structure, suggesting that the re-takenup extrocellular amelogenins were released from endosome and were transported to the area near the ER structure, and in there amelogenin protein directly bind to amelogenin mRNA. The data from this study is helpful to better understand the biological functions and gene regulation of amelogenin in dental development, and provide important information for studying molecule-based tooth tissue regeneration.

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11. Bouropoulos N., Moradian-Oldak J., 2003. Analysis of hydroxyapatite surface coverage by amelogenin nanospheres following the langmuir model for protein adsorption, Calcif. Tissue. Int., 72, 599-603. 12. Wen H.B,, Moradian-Oldak J., Fincham A.G., 1999. Modulation of apatite crystal growth on Bioglass by recombinant amelogenin, Biomaterials, 20, 1717-1725. 13. Wen H.B., Moradian-Oldak J., Zhong J.P., Greenspan D.C., Fincham A.G., 2000. Effects of amelogenin on the transforming surface microstructures of Bioglass in a calcifying solution, J Biomed. Mater. Res., 52, 762-773. 14. Lau E.C., Simmer J.P., Bringas P.Jr., Hsu D.D., Hu C.C., Zeichner-David M., Thiemann F., Snead M.L., Slavkin H.C., Fincham A.G., 1992. Alternative splicing of the mouse amelogenin primary RNA transcript contributes to amelogenin heterogeneity, Biochem. Biophys. Res. Commun., 188, 1253-1260. 15. Simmer J.P., Hu C.C., Lau E.C., Sarte P., Slavkin H.C., Fincham A.G., 1994. Alternative Splicing of the Mouse Amelogenin Primary RNA Transcript, Calcif. Tissue Int., 55, 302-310. 16. Brookes S.J., Robinson C., Kirkham J. Bonass W.A.,1995. Biochemistry and molecular biology of amelogenin proteins of developing dental enamel, Arch. Oral Biol., 40, 1-14. 17. Simmer J.P.,1995. Alternative splicing of amelogenins, Connect Tissue Res., 32, 131-136. 18. Bonass W.A., Kirkham J., Brookes S.J., Shore R.C., Robinson C., 1994. Isolation and characterization of an alternative-spliced rat amelogenin cDNA: LRAP-a highly conserved, functional alternatively-spliced amelogenin? Biochim. Biophys. Acta., 1219, 690-692. 19. Li R., Li W., DenBesten P.K., 1995. Alternative splicing of amelogenin mRNA from rat incisor ameloblasts, J Dent. Res., 74, 1880-1885. 20. Hu C.C., Bartlett J.D., Zhang C.H., Qian Q., Ryu O.H., Simmer J.P., 1996. Cloning cDNA sequence and alternative splicing of porcine amelogenin mRNAs, J Dent. Res., 75, 1735-1741. 21. Moradian-Oldak J., Leung W., Simmer J.P., Zeichner-David M., Fincham A.G., 1996. Identification of a novel proteinase (ameloprotease-1) responsible for the complete degradation of amelogenin during enamel maturation, Biochem. J., 318, 1015-1021. 22. Veis A.,2003. Amelogenin gene splice products: potential signaling molecules, Cell Mol. Life Sci., 60, 38–55. 23. Papagerakis P., Ibarra J.M., Inozentseva N., DenBesten P., MacDougall M., 2005. Mouse amelogenin exon 8 and 9: sequence analysis and protein distribution, J Dent. Res., 84, 613–617. 24. Veis A., Tompkins K., Alvares K., Wei K., Wang L., Wang X.S., Brownell A.G., Jengh S-M., Healy K.E., 2000. Specific amelogenin gene splice products have signaling effects on cells in culture and in implants in vivo, J Biol. Chem., 275, 41263-41272. 25. Zhou, Y.L., and Snead, M.L. Identification of CCAAT/enhancer-bending protein a as a transactivator of the mouse amelogenin gene. J Biol Chem 275, 12273-12280, 2000. 26. Zhou Y.L., Lei Y., Snead M.L., 2000. Functional antagonism between Msx2 and CCAAT/enhancer-binding protein a in regulating the mouse amelogenin gene

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Chapter 7 Mechano-active Tissue Engineering Sang-Heon Kim1, Youngmee Jung1, Soo Hyun Kim1, and Young Ha Kim2 1. Biomaterials Research Center, Korea Institute of Science and Technology, Seoul, Korea 2. Department of Materials Science & Engineering, Gwangju Institute of Science and Technology, Gwangju, Korea

1

Introduction

It is well known that mechanical stimulation regulates the specialized structures and functions of mammalian cells, tissues, and organs. Mechanical stresses primarily originate either from tension that is caused by cells themselves or from extracellular matrix (ECM) through cell adhesion. Significant recent progress has been achieved through studies of cardiovascular tissues and skeletal tissues, such as bone and articular cartilage, in mechanical, stress-related cell biology [1-3]. The field of tissue engineering has progressed to develop means of regenerating damaged tissues and organs using cells and scaffolds. The fundamental role of a scaffold is not only to provide a temporary substrate on which transplanted cells can adhere but also to maintain mechanical integrity during the healing process and to deliver appropriate mechanical signals to adherent cells that ultimately comprise dynamic physiological systems [4-7]. Within tissue engineering, strategies concerning the mechanical environment of cells or tissues have emerged, and the resulting field has been termed “mechanoactive tissue engineering.” Specific examples include the development of elastic, mechano-active scaffolds that transmit mechanical stimulation to cells or protect

cells from mechanical forces for use in vascular and cartilage tissue engineering. The mechanical properties of native blood vessels and of cartilage provide researchers with key design elements to apply in properly developing a scaffold that will function under conditions similar to those of the native tissues. In this chapter, we review some of the recent advances in mechano-active tissue engineering, focusing on blood vessel and cartilage regeneration.

2

Mechano-biology

Biological processes such as apoptosis, proliferation, differentiation, and migration are regulated through both outside-in and inside-out signaling in the form of mechanical stimuli in addition to biochemical interactions. Numerous reports have shown that molecules such as growth factors receptors, cell adhesion receptors, G proteins, and ion channels act to transduce mechanical stimuli into biochemical signals within cells, although specific mechanoreceptors still have not been found on cell membranes. The process by which mechanical stimuli are sensed and transmitted to the nucleus to induce changes in cell morphology and phenotype is also not clearly understood. Nevertheless, tissue engineering research takes mechanical stimuli into consideration, particularly in efforts to engineer components of the cardiovascular system and articular cartilage. Blood vessels are dynamic tissues with high elasticity and strength suited to withstand both the flow of blood and the associated pressure. Blood vessel walls are comprised of three layers: the tunica intima, the tunica media, and the tunica adventitia. The intima consists of a single layer of endothelial cells (ECs) that are in direct contact with the blood flow and are supported by a subendothelial layer containing collagen fibers. Between the intima and the media is the internal elastic lamina, a layer of cross-linked elastin fibers [8]. The media contains smooth muscle cells (SMCs) that are subjected to the pulsatile load experienced by blood vessel walls. Numerous studies have reported that SMC phenotype and alignment are regulated significantly by mechanical stimulation, such as cyclic strain and pulsatile flow, in two-dimensional or three-dimensional culture systems [9-11]. Under mechanical stimulation, confluent SMCs orient perpendicular to applied strain and highly express SMC markers such as smooth muscle (SM) α-actin, myosin heavy chain, and caldesmon. In efforts to engineer blood vessels using bioreactor systems, interactions between ECs and SMCs (e.g., EC adhesion to and lining of the inner lumen in contact with cultured SMCs) are improved under the proper mechanical stimuli. In

addition, cyclic strain has been shown to induce stem cell differentiation into SMCs [12]. Native articular cartilage is influenced by a complex mechanical load that includes both compressive and shearing forces. In normal cartilage, the extracellular matrix exists as a highly organized composite of specialized macromolecules that distributes loads at the bony ends. Proper mechanical stresses are necessary to promote chondrogenesis. However, excessive stimulation damages chondral tissues and can lead to cartilage ossification and fibrous tissue formation. Articular cartilage defects result in inappropriate mechanical loading of the joint, which causes osteoarthritis. In articular cartilage, chondrocytes maintain a rounded morphology with limited cell-cell contact and are surrounded by ECM components such as collagen and glycosaminoglycan (GAG). This rounded morphology plays a pivotal role in maintaining and inducing the chondrocyte phenotype. Numerous reports have demonstrated that appropriate mechanical stimuli, including cyclic compressive strain and osmotic stresses, act as important factors to maintain or induce the chondrocyte phenotype with the rounded cell shape in three-dimensional culture [13, 14]. Hence, mimicking the mechanical stimulation associated with normal physiological function is crucial in properly regenerating articular cartilage.

3

Mechano-active Scaffolds

Mechano-active scaffolds have employed elastic materials in vascular and cartilage tissue engineering. Natural polymers such as collagen have been studied as an elastic scaffold created by cross-linking these polymers with chemicals such as glutaldehyde, carbodiimide, and divinylsulfone [15, 16]. Although chemical cross-linking increases the elasticity of the scaffolds, these chemicals also may be cytotoxic [17, 18]. On the other hand, polyglycolic acid (PGA) fibers were cross-linked with poly(L-lactic acid) (PLA) to form a synthetic, biodegradable polymer scaffold for mechano-active tissue engineering [19]. However, this synthetic polymer scaffold exhibited significant, permanent deformation under cyclic mechanical strain conditions [20-23]. A mechano-active polymer, poly-(L-lactide-co-ε-caprolactone) (PLCL) copolymer, was synthesized by the ring opening polymerization of L-lactide and ε-caprolactone in the presence of Sn(Oct)2 (Fig. 1). PLCL is composed of a soft matrix of ε-caprolactone moieties and hard domains of L-lactide units. The monomers used in this system differ greatly in mechanical properties and time to reach complete mass loss; however, once physically cross-linked in specific

Figure 1. Chemical synthesis and 1H-NMR spectrum of PLCL (50:50).

monomer ratios, the copolymer system exhibits a rubber-like elasticity [24]. The structure of PLCL was identified by 1H NMR spectra in CDCl3 (Fig. 1). The methine protons of the lactide unit appear as two singlets at δ 5.1–5.2 (a, a’) as a result of the sequence distribution of the lactyl and caproyl units, while the methylene protons of the caproyl unit adjacent to the ester group appear at δ 4.0–4.2 (c, c’) and δ 2.3–2.5 (g, g’). The copolymer compositions synthesized by these relative intensities almost identically correspond to the initial feed compositions. PLCL prepared from 50 wt % of L-lactide and 50 wt % of εcaprolactone is highly elastic and has been fabricated for microporous scaffolds using a variety of techniques such as extrusion-particulate leaching, gel-spinning, gel-pressing, freeze-drying, and electrospinning. The mechanical properties of PLCL scaffolds were measured and compared to those of PLGA scaffolds; PLCL scaffolds fabricated with 60% porosity exhibit a strain of 500%, and scaffolds of 90% porosity can be extended to 200%. The elastic properties of PLCL scaffolds were evaluated by measuring recovery after stretching, as

(A)

(F) PLGA

(B)

(D)

(C)

(E)

Figure 2. Elastic behaviors of PLCL scaffolds. PLCL scaffolds (A) were extended at 250% of initial length (B) with 3MPa for 5 sec and recovered (C) by releasing the load. The scaffolds were twisted (D) and folded (E) via a cyclic strain apparatus. PLGA (F) was broken at 20% strain.

shown in Figure 2. PLCL scaffolds with 90% porosity show 100% recovery at near 100% strain. In addition, PLCL scaffolds can be easily twisted and bent. In contrast, PLGA scaffolds largely deform and are broken even at strains as low as 20%. These data indicate that PLCL scaffolds are flexible and highly elastic, whereas PLGA scaffolds are stiff and brittle. To further examine the elastic properties of PLCL scaffolds, scaffolds with varying porosity were subjected to cyclic strain at 10% amplitude and 1 Hz frequency for 27 days in culture medium [25]. PLCL scaffolds of all tested porosities maintain excellent elasticity even in the hydrolytic medium over a 27-day experimental time course. Degradation tests show that the degradation rate of PLCL scaffolds is somewhat faster in vivo than in vitro, and this may be explained by enzymatic degradation possibly playing a role in degradation in the body. In addition, the CL moieties degrade faster than the LA units in PLCL scaffolds, although their hydrophilicities are in opposite order. This behavior appeared more prominently in vivo, probably indicating that amorphous regions composed of primarily CL units are first to be attacked by water, which can penetrate into the amorphous regions easier than into the hard domains that are composed primarily of LA units.

Figure 3. Fabrication of extruded tubular PLCL scaffolds. (A) Preparation of PLCL/NaCl mixture; (B) extrusion mold; (C) piston extrusion tool; (D) tubular scaffold extruded (I.D.: 4 mm, O.D.: 5 mm).

Tubular PLCL scaffolds were fabricated by a particulate-leaching/extrusion method (Fig. 3) and a gel-spinning technique (Fig. 4) for vascular tissue engineering. In conduit PLCL scaffolds that are fabricated by the particulate leaching/extrusion, SEM micrographs reveal extensive pores that are roughly spherical in shape. The pores appear to form interconnected networks and the porosities and pore sizes of the scaffolds can be varied independently by adjusting the fractions and sizes of particles used (Fig. 5A, B). We have used tubular PLCL scaffolds fabricated by the extrusion-particulate leaching technique as mechano-active scaffolds for the application of cultured, smalldiameter blood vessels. However, these extruded PLCL scaffolds are limited in mechanical strength, and vessel tensile properties are fundamental factors with which to evaluate scaffolds for vascular tissue engineering applications. In particular, the tensile strength of scaffolds is an important element for a successful vascular graft, because vascular grafts must have adequate strength to resist rupture or excessive dilation when subjected to pulsatile pressures during implantation. Therefore, a novel gel-spinning molding technique was used to fabricate a tubular, microporous, fibrous PLCL scaffold to overcome the limited mechanical strength of extruded PLCL scaffolds [26]. The cross-sectional image

Figure 4. Morphology (B) of tubular fibrous PLCL scaffolds prepared by a new gel-spinning technique (A).

and surface topography of fibrous PLCL scaffolds are shown in Figure 5C, D. Individual fibers are found crossed and/or fused together with neighboring fibers. These cross-linked fibrous networks cause the spun PLCL scaffolds to be open-pore structures and well-interconnected between pores. The fibrous, tubular PLCL scaffolds show good mechanical strength and cell adhesion and proliferation efficiencies, as compared to PLCL tubular scaffolds fabricated by the extrusion/particulate-leaching method. A sheet-form PLCL scaffold with 80% porosity and pores ranging from 300–500 μm in size was fabricated for cartilage tissue engineering using a gelpressing method [27]. The surface topology and cross-sectional images obtained by SEM reveal that the PLCL scaffolds have a homogeneously interconnected and open pore structure without a skin layer. These properties will improve cell seeding efficiency, cell in-growth, and, eventually, cartilage regeneration. We showed that tensile properties of sheet-form PLCL scaffolds can be controlled by changing the initial monomer composition of PLCL. Tensile tests show that

scaffolds that are fabricated from PLCL-1 (50:50 in LA:CA) have a lower tensile modulus (0.31 MPa) and a higher elongation at break (520%) than scaffolds that are fabricated from PLCL-2 (60:40 in LA:CA) which possess a modulus of 0.73 MPa and elongation of 136%. Both PLCL films and scaffolds exhibit a completely rubber-like elasticity and show almost complete recovery (over 94%) up to the tensile strain before a break occurs. These elastic properties suggest that the PLCL scaffolds are suitable for use in cartilage tissue engineering. Because the mechanical properties of PLCL scaffolds can be controlled by changing the monomer content, these materials can be applied in a range of applications related to cartilage tissue engineering that depend on specific mechanical properties for in-growth and eventual cartilage regeneration. PLCL-1 and PLCL-2 films have very high strains at break (1100% and 800%, respectively), low stresses at break (6.7 MPa and 7.5 MPa, respectively), and very low tensile moduli (0.015 MPa and 0.014 MPa, respectively).

A

B

C

D

Figure 5. SEM images of tubular PLCL scaffolds prepared by a particulate/extrusion method (A, B) and a new gel-spinning technique (C, D). A and C, outer surface; B and D, cross-section.

4

Mechano-active Vascular Tissue Engineering

In vivo vascular SMCs typically reside in mechanically dynamic environments, align in a specific direction, and exist in a contractile, differentiated phenotype, which is critical for the contractile functions of SMCs [25, 28]. SM tissues that

are engineered in vitro using conventional tissue engineering techniques may not be functional, because SMCs cultured in vitro usually do not align and usually do revert from a contractile, differentiated phenotype to a synthetic, non-differentiated phenotype. As previously mentioned, a number of studies have shown that mechanical stimuli regulate the characteristics of SMCs in culture systems that supply cyclic strain and pulsatile flow. Vascular SM reconstructed under such mechanically active conditions exhibits enhanced mechanical strength, collagen and elastin production, or a good patency after implantation. For these mechano-active systems, scaffolds must be able to deliver mechanical stresses to adhered SMCs. A non-woven PGA mesh has been used as a scaffold in the reconstruction of a variety of tissues including blood vessels. In previous studies, PGA mesh scaffolds had to be sewn with suture and inserted by a highly distensible silicone tubing that transduced pulsatile pressure indirectly into the scaffolds. Thus, the PGA material was limited as a mechano-active scaffold due to its non-elastic properties. Due to the need for elastic properties, we have developed PLCL as a mechano-active scaffold material for vascular tissue engineering. Specificallydesigned PLCL scaffolds were seeded with SMCs labeled with CM-DiI and

Figure 6. In vivo evaluation of SMCs-seeded PLCL scaffolds. CM-DiI-tagged SMCs seeded PLCL scaffolds were subcutaneously implanted in nude mice for 2, 5, and 8 weeks; (A)-(C), immunohistochemical staining for SM a-actin; (D)-(F), CM-DiI detection.

implanted into nude mice to investigate tissue compatibility, SMC growth, and in vivo scaffold degradation behavior [29, 30]. Immunohistochemical analyses of extracted scaffolds demonstrated that SM α-actin gradually increased in the implanted scaffolds (Fig. 6A-C). Fluorescent cells were clearly detectable under UV microscopy and easily distinguished from their unlabeled neighbors. The fluorescence was concentrated in each cell that was seeded in the early implantation period, but also extended and dispersed into the surrounding areas as the seeded cells extended and/or cells from surrounding areas grew into the implanted scaffolds (Fig. 6D-F). The implanted PLCL scaffolds degraded at a proper rate, while SM tissues were regenerated in the scaffolds. Thus, PLCL scaffolds exhibit good biocompatibility for SMCs and a proper in vivo degradation rate. To investigate the effects of mechanical stimulation on the proliferation and phenotype of SMCs adhered to PLCL scaffolds, an extruded tubular scaffold was utilized as a three-dimensional cell culture substrate for SMCs under pulsatile strain and shear stress conditions [24]. We hypothesized that a radial distention would induce the phenotype of SMCs in vitro to be similar to that of SMCs in vivo.

Figure 7. SEM images and H&E staining of tissue engineered vessels; SEM images of SMCs grown statically (A) or under a pulsatile flow (B) for 8 weeks on the scaffolds surface; H&E staining of the cross-section of the scaffolds cultured statically (C) or under a pulsatile flow (D). The arrow indicates the direction of pulsatile flow through the inside of tubular scaffolds.

Aortic SMCs were seeded onto PLCL scaffolds and subjected to pulsatile strain and shear stress in culture in pulsatile perfusion bioreactors. The tubular scaffolds maintain their shape without permanent deformation during the culture period. SMCs proliferate and eventually cover the surface of the scaffolds over 8 weeks in pulsatile perfusion bioreactors. In control, static culture, cells grow much slower and the scaffold surface was not completely covered at 8 weeks. Pulsatile strain and shear stress enhance SMC proliferation and collagen production. In additional to collagen, elastin is a major ECM component secreted by functional SMCs in native blood vessels. The elastin content of transplanted PLCL scaffolds was examined as a measure of SMC differentiation in a mechanically active culture system. Elastin expression is greater in the scaffolds exposed to mechanical stimuli relative to that in static constructs [31]. Furthermore, significant cell alignment in a direction radial to the distending direction, which is similar to that of native SM tissues, is observed in SM tissues that are exposed to radial distention, whereas SMCs in SM tissues reconstructed under static conditions align randomly (Fig. 7). Western blot analysis demonstrates that the expression of SM α-actin is upregulated by 2.5-fold in SM tissues reconstructed under the mechano-active conditions, as compared to that in tissues grown in static conditions [24]. This study demonstrates that tissue engineering of SM tissues in vitro using pulsatile perfusion bioreactors and elastic PLCL scaffolds leads to enhanced tissue development and differentiated SMC phenotype. In a different strategy, we used a gel-spinning molding technique to fabricate fibrous tubular scaffolds with high mechanical strength and elasticity. Our aim was to develop a seamless tubular PLCL scaffold that would resist rupture or leakage under high pressure to be applied for the implantation of blood vessels, such as the aorta, rather than for use in in vitro culture for tissue reconstruction. A double-layered tubular scaffold was fabricated by a novel gelspinning technique, implanted into a canine abdominal aorta, and characterized through examination of mechanical and biological properties for blood vessel reconstruction (unpublished data).

5

Mechano-active Cartilage Tissue Engineering

Due to the lack of self-healing capacity of avascular and aneural tissues, even minor cartilage defects can result in mechanical joint instability and progressive damage, and cartilage damage is notoriously difficult to treat and cure. Although many studies have explored cartilage regeneration using techniques associated

with tissue engineering and cell therapy (e.g., autologous chondrocyte implantation and marrow stimulation), these approaches have not adequately addressed certain problems in cartilage regeneration, such as fibrocartilage formation and lower mechanical strength of implants relative to that of native cartilage [32, 33]. Articular hyaline cartilage is subjected to particularly complex loads that affect its development and maintenance in the body [34]. Hence, mechanical stimulation associated with normal physiological movement is crucial in properly re-forming articular cartilage through tissue engineering. The successful generation of functional, engineered cartilage will require a mechano-active scaffold that can transmit mechanical signals to adherent cells in the physiologically dynamic in vivo environment. To deliver the required mechanical signals associated with the surrounding biological environment of cartilage, we fabricated a sponge-type microporous scaffold from the elastic PLCL copolymer by a gel-pressing method. Xie et al. evaluated microporous PLCL scaffolds in an ex-vivo system for mechano-active, scaffold-based cartilage tissue engineering. A new cell-seeding technique was devised to improve chondrocyte distribution into and viability in microporous PLCL scaffolds under compression force-induced suction. They successfully delivered cells to microporous PLCL scaffolds with high cell seeding efficiency, cell viability, and homogeneous cell distribution under multiple cycles of suction. Xie et al. focused on how chondrocytes respond bio-mechanically to various modes of compressive loading, specifically loading frequency, loading duration per cycle, loading period, and continuous or intermittent compression, in mechano-active PLCL scaffolds. mRNA expression of ECM molecules associated with cartilage is observed in chondrocyte-containing PLCL scaffolds cultured at three different loading frequencies: 0.05, 0.1, and 0.5 Hz. Type II collagen expression is the highest in scaffolds cultured at 0.1Hz, whereas little difference is observed in aggrecan mRNA levels at these frequencies. Continuous dynamic compression results in decreased mRNA levels of the cartilage-dependent ECM components, aggrecan and type II collagen. An intermittent loading (24-h cycle of loading and unloading) program maintains high levels of ECM component mRNA expression. Continuous dynamic compression also causes the release of sulfated glycosaminoglycan (S-GAG) from the chondrocyte-containing PLCL scaffolds into the surrounding medium [35, 36]. Thus, an excessive mechanical stress is not favorable for mRNA expression and protein accumulation, both of which influence cartilage formation. An appropriate mechanical stimulation program is required for the construction of functional cartilage using mechano-active scaffolds in vitro.

Figure 8. A bioreactor in compressive mode (A) and the units of the cradle of scaffolds and the plunger for compressive forces (B).

To evaluate the suitability of microporous PLCL scaffolds for mechanoactive cartilage tissue engineering, chondrocyte-seeded PLCL scaffolds were cultured for 10 days or 25 days under continuous compressive deformation of 5% strain at 0.1 Hz using a compressive-mode bioreactor (Fig. 8) or under control static conditions, and subsequently implanted subcutaneously into nude mice [37]. The collagen and GAG content of mechanically stimulated scaffolds increases significantly over 10 days in culture compared to that in static-cultured scaffolds. Histological analysis shows that mechanically stimulated implants formed mature and well-developed cartilaginous tissue, as evidenced by the presence of chondrocytes within lacunae and the abundant accumulation of SGAGs (Fig. 9). However, more unhealthy lacunae shapes and hypertrophic forms are observed in the implants that have been mechanically stimulated for 24 days as compared with those stimulated for 10 days. This result suggest that the proper periodical application of dynamic compression can encourage the maintenance of chondrocyte phenotype and enhanced GAG production, which will improve the function of cartilaginous tissue constructed both in vitro and in vivo. We compared the mechanical properties of elastic PLCL scaffolds to those of conventional rigid polymers scaffolds, such as PLA and PLGA scaffolds. The PLCL scaffolds possess a completely rubber-like elasticity, are easily twisted and bent, and exhibit an almost complete (over 97%) recovery from applied

Figure 9. Histological studies of implants at 8 weeks. The sections were stained with H&E (A–E), Masson’s Trichrome (F–J), Alcian Blue (K–O), or Safranin O (P–T). Articular cartilages of rabbits were stained as the positive controls (A, F, K, P). The images are of the constructs stimulated mechanically for 10 days (B, G, L, Q) or 24 days (D, I, N, S) and of the constructs not stimulated for 10 (C, H, M, R) or 24 days (E, J, O, T).

strain (up to 500%), while the control PLA scaffolds show little recovery after strain. We evaluated their abilities to promote cartilaginous tissue formation and cartilage regeneration in vitro, in nude mice, and in a rabbit cartilage defect model. We seeded scaffolds with rabbit chondrocytes, cultured them in vitro, and subcutaneously implanted them into nude mice for up to 8 weeks. In vitro and in vivo accumulation of extracellular matrix on the cell-PLCL constructs demonstrates that these scaffolds can not only sustain but significantly enhance chondrogenic differentiation. Moreover, the mechanical stimulation of the dynamic in vivo environment promotes deposition of the chondral extracellular matrix onto implanted PLCL scaffolds. In contrast, on the PLA scaffolds, most of the chondrocytes de-differentiate and form fibrous tissues. In the rabbit defect model, animals implanted with PLCL scaffolds exhibit significantly enhanced cartilage regeneration compared to those that receive an empty control or PLGA scaffold. These results indicated that the mechano-active PLCL scaffolds

effectively deliver mechanical signals associated with biological environments to adherent chondrocytes, suggesting that these elastic PLCL scaffolds can successfully be used for cartilage regeneration.

6

Conclusions

Mammalian cells and tissues reside in mechanically dynamic microenvironments in the body. Inappropriate physical loads upon tissues result in tissue deformation. Numerous studies have confirmed that proper mechanical stimuli applied to cells or tissues are involved in maintaining cell/tissue morphology and inducing specialized functions. Mechanical stimuli have recently been applied to regenerate functional tissues, in particular, tissues in the cardiovascular system and articular cartilage. To reconstruct functionally active SM tissues and chondrogenic tissues that are comparable to the native tissues, the in vitro re-creation of the in vivo mechano-active microenvironments may be necessary in the tissue engineering process. Tissue engineering scaffolds play a crucial role in delivering mechanical stresses from the extracellular environment to the cells responsible for tissue formation. PLCL is a rubber-like, elastic, biodegradable polymer that has developed as a key material for mechano-active tissue engineering. In both in vitro and in vivo studies, PLCL scaffolds have been beneficial in sustaining SMC phenotype and chondrogenic differentiation in vascular tissue engineering and cartilage tissue engineering, respectively, presumably through effectively transmitting mechanical signals through dynamic microenvironments. PLCL will be an excellent candidate material for mechano-active scaffolds that deliver mechanical stresses to cells and tissues via a transducer.

Acknowledgments This work was supported in part by a grant of the Korea Health 21 R&D Project, Ministry of Health & Welfare, Republic of Korea (A050082).

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Chapter 8 Biomaterials for In Vitro Expansion of Stem Cells Yoshihiro Ito Nano Medical Engineering Laboratory, RIKEN, Saitama, Japan

1

Introduction

For achievement of regenerative medicine, in vitro expansion of stem cells which can be differentiated to various cells is very important. However, it is very difficult to efficiently and safely culture some stem cells such as hematopoietic stem cells in cord blood (CB) or human embryonic stem (hES) cells. For example, hES cells cultured on mouse feeder cells expressed an immunogenic nonhuman sialic acid [1]. Human iPS cell developed by Yamanaka’s group is cultured on MatrigelTM which is a solubilized basement membrane extracted from the Engelbreth-Holm-Swarm mouse sarcoma, a tumor rich in extracellular matrix proteins [2]. Therefore, the culture systems are investigated by many researchers. Here our strategies for development of biomaterials for expansion of hematopoietic stem cells in CB and ES cells are discussed (Figure 1).

2

Culture of Human Umbilical Cord Blood Cells

Human umbilical CB is regarded as a very attractive medical source for transplantation therapy and regenerative medicine, because its abundance in primitive hematopoietic progenitors and stem cells, immature immunoresponsiveness, and easiness of procurement from the CB bank.

However, the major limitation to use of CB exists in its limited absolute cell number for widespread medical applications. Many investigators have explored in vitro expansion methods to increase the numbers of hematopoietic cells [3-5]. Indeed, successful CB expansion by the combination of hematopoietic growth factors (GFs) was proposed by several groups [6-8], but the expansion efficiency in GF combination culture was lower. Dexter-type long term bone marrow (BM) culture which was established by mimicking BM environment was excellent culture system to expand stem or progenitor cells [9]. Coculture with BM stromal cells was demonstrated as a promising system for the maintenance and expansion of hematopoietic primitive progenitors [10-12]. However, numerous number of culture systems have been reported mainly by using murine cell lines and stromal medium composed with calf and horse serum. These systems were considered as xenotransplantation in human hematopoietic cell expansion. To reduce the risk such as infectious disease, culture system without animal-derived stromal cells is desired.

Figure 1. There are two types of stem cells for resource of tissue engineering or regenerative medicine. A red frame indicates the target process in this article.

Prevously it was reported that human CB cells were greatly expanded in the presence of stem cell factor (SCF), thorombopoietin (TPO), and Flk-2/Flt-3 ligand (FL) by coculture with primary [13] or immortalized [14] human BM stromal cells. However, it is troublesome to provide cells for each expansion. Therefore, we attempted to chemically fix the stromal cells (Figure 2) [15].

Morphology of stromal layers treated with glutaraldehyde (GA) was almost the same as that of non-treated. Furthermore, GA treated layers maintained their morphology even after 4 weeks, though partially formed interstitial spaces. In contrast to the GA treatment, ethanol treatment induced shrinkage of cells and the layer formed flat and very superficial structure. Para formaldehyde (PFA) treatment did not induce shrinkage of the cells, however, some cell fragments were removed during the process of the preservation.

Figure 2. Chemically fixed feeder cells were prepared as a new technology for culture of stem cells.

To assess the surface structure of stromal cell layers in a more direct fashion, scanning electron microscopy (SEM) of the cell was performed. When treated with GA, stromal cell layers formed multi-laminate structure consisted of amoeboid shape cells with processes on the surface. Stromal cell layers treated with ethanol were largely plain. Many small granules and nucleus structure also could be seen in the surface, revealing that the plasma membrane structure was mostly lost by the ethanol treatment. Morphology of PFA treated layers are similar to that of ethanol treated, indicating that the PFA fixation was not enough to maintain the cell surface structure during the process of SEM measurement. Hoetelmans et al. [16] reported plasma membrane integrity was poorly preserved after fixation in acetone or methanol by SEM observation.

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Figure 3. Effect of various culture conditions on expansion of human cord blood progenitors. Stromal cells for hematopoietic cell coculture were pretreated with glutaraldehyde (GA), paraformaldehyde (PFA), or ethanol. Two thousand CD34+ cells were cultured with or without stromal cells in 2ml of α-MEM supplemented with 20% FCS, 1% BSA, each 10 ng/ml of SCF, TPO, and FL. After 2 weeks of coculture, hematopoietic cells were collected and estimated (a) Fold increase of total cell number, (b) Fold increase of CD34+ cell number, and (c) Fold increase of CFU-C. n=3.

Our SEM observation of ethanol treated stromal cells strongly suggested this phenomena, that is, plasma membrane and intracellular structure was damaged by ethanol treatment and many small granules and nucleus structure could be seen in the surface. On the other hand, GA treatment relatively preserved the stromal cell surface integrity than ethanol treatment as shown by SEM.

2.1 Effect of Fixation Methods of Stromal Cells on the Ex Vivo Expansion of Human CB Progenitors We next examined hematopoiesis supporting ability of the fixed stromal cells by comparing with that of the viable stromal cells (Figure 3). Two thousand of human CB CD34+ cells were cultured in serum containing medium supplemented with SCF, TPO, and FL. After 2 weeks, hematopoietic cells were collected and were analyzed by counting of number of total cells, CD34+ cells, and CFU-C. The number of total cells, CD34+ cells, and CFU-C after 2 weeks of stroma-free culture increased by 110-, 9-, and 15-times at the initial input cell number, respectively. This expansion degree in bovine serum culture was slightly higher than that previously estimated in AB-serum containing medium [17,18] or serum free medium [13]. As compared with stroma-free culture, GA fixed stromal cells enhanced the number of total cells, CD34+ cells, and CFU-C by 3.6-, 3.7-, and 3.5-fold, respectively, compared to 2 weeks of stroma-free culture value. Even though the expansion efficacy of fixed stromal cells was lower than that of viable stromal cells, GA-fixed stromal layers have the ability to enhance the CB cell expansion. On the other hand, enhancement of the number of total cells, CD34+ cells, and CFU-C by PFA or ethanol fixed stromal cells was approximately 1.2- or 0.6-fold, respectively, compared to 2 weeks of stroma-free culture value. These results reveal that the hematopoiesis supporting ability depended on the fixation methods of stromal cells.

2.2 Stromal Cell Contact Promoted the Expansion of Progenitors To estimate whether inhibition of direct cell-cell contact between stromal cells and hematopoietic progenitor cells reduces the expansion efficacy, CD34+ cells was cultured on 0.4 μm microporous membrane separately with stromal cells. Figure 4 shows the fold increase in the number of total cells, CD34+ cells, and CFU-C after 2 weeks of culture with stromal cells compared with that without stromal cells using 3 different CB donors. In viable stromal cells, contact inhibition brought the reduction in the number of CD34+ cells and CFU-C. In GA-fixed stromal cells, contact inhibition also observed and the number of expanded cells was almost the same as that in stroma free culture. The result reveals the hematopoiesis supporting ability of GA fixed stromal cells was induced by the contact between their surface and hematopoietic cells.

Relative increase in cell number

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Figure 4. Effect of direct cell-cell contacts between hematopoietic cells and stromal cells on expansion of human cord blood progenitors in two weeks culture. ‘Relative increase’ means the fold increase compared with basic values in cytokine combination culture. ‘Well’ represent the separation culture of hematopoietic cells and stromal cells by use of a cell culture transwell insert. (a) Relative increase of total cell number, (b) Relative increase of CD34+ cell number, and (c) Relative increase of CFU-C. n=3.

Hurley et al. [19] reported that GA treatment did not destroy the antibody recognition of cell surface adhesion molecules and extracellular matrix proteins such as fibronectin, thrombospondin, and collagen. Recently, Bhatia et al. [20] reported that enhancement of primitive progenitor preservation by coculture with glutaraldehyde (GA) fixed stromal cells. Inhibition of direct cell-cell contact reduced the expansion of hematopoietic progenitor cells by stromal cells [13] although there is controversy with the requirement of stromal cells for the in vitro maintenance and expansion of hematopoietic cells [21, 22]. The present investigation demonstrated that GA fixed stromal cell would have the capacity to expand CB cells as a result of preserved cell surface integrity. Because fixed stromal cells are in metabolically inactivated state and lost the ability to secrete diffusible factors, hematopoiesis supporting ability was mainly depended on the direct cell-cell contact mechanism. Considering that integrin family [23], cell adhesion molecules [24,25], CD44 [26], and selectin family [27] were known as mediators of cell-cell contact to regulate the hematopoietic proliferation and that membrane protein such as Notch ligand, jagged-1 [28,29] and mKirre [30] was found on the BM stromal cell membrane as a hemotopoietic stem cell supporting protein, these mediators would be candidate functional molecules on GA-fixed stromal cells.

2.3 GA-fixed Stromal Cells Stored for Several Weeks Maintained the Ability to Expand Human CB Progenitors Hematopoiesis supporting ability of stromal cells was investigated to estimate whether refrigeration preservation of the GA fixed stromal cells was possible, and the result was shown in Figure 5. Because viable stromal cells detached from the plastic support several hours after incubation in PBS, ex vivo expansion by viable stromal cells 1 or 4 weeks after refrigeration preservation was not conducted. The result revealed that hematopoietic supporting ability of GA-fixed stromal cells maintained at least for 4 weeks. Previously Higashiyama et al, [31] used the chemically fixed cells for investigation of cell membrane-associated growth factors such as heparinbinding epidermal growth factor (HB-EGF). They demonstrated that HB-EGF worked on another cells within the cell membrane by chemically fixation methods. The present study demonstrated that chemical fixation technique is useful for not only academic research but also cell culture methodology.

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Figure 5. Expansion of human cord blood progenitors on refrigeration preserved glutaraldehyde-treated stromal cells. Because viable stromal cells detached from the plastic support several hours after incubation in PBS, ex vivo expansion by viable stromal cells 1 or 4 weeks after refrigeration preservation was not tested (N.T.). Expansion data of 0 week was the same as in (A). (a) Fold increase of total cell number, (b) Fold increase of CD34+ cell number, and (c) Fold increase of CFU-C. n=3.

3

Culture of ES Cells

ES cells were first derived from the inner cell mass of mouse blastocysts in the early 1980s [32]. ES cells have the unique ability to give rise to any type of somatic cell lineage. Therefore, ES cells are important for analyzing the molecular mechanisms of self-renewal and embryogenesis, and are frequently used to generate transgenic mice. In addition, the implementation of cell-based or regenerative therapies requires ES cells as a renewable source of cells. Efficient in vitro expansion of ES cells is therefore important for the development of ES cell technologies. When ES cells are cultured, the environment in which they are propagated has the potential to influence their capacity to act as therapeutic agents in tissue engineering. It is therefore important to understand the response of ES cells to synthetic or biological matrices to ascertain their usefulness as implant materials. Previous studies demonstrate that mouse embryonic fibroblasts (MEFs) support the continued propagation of ES cells in the primitive undifferentiated state while retaining their pluripotency. In the absence of MEFs, gelatin can similarly support the growth and propagation of mouse ES cells in the presence of cytokine leukemia inhibitory factor (LIF), as recently described by Hamazaki et al. [33]. In the case of human ES cells, recently bFGF is reported to be useful for growth of human ES cells with keeping the undifferentiated state [34-37]. However, these replacements cannot be generalized for all lines of ES cells. For culture of primate ES cells including human ES cells, some types of human nurse cells were reported [38-40]. Among them human amniotic endothelial (HAE) cell, which was reported by Miyamoto et al. [41], are useful by taking into consideration the easiness of acquisition. However, there are two disadvantages of usage of these nurse cells. One is troublesome for preparation before each ES cell culture. In addition, the nurse cells cannot be completely removed for the next culture after trypsinization.

3.1 Chemically Fixed Cells for ES Cells As described in the section of CB cells, chemically fixed nurse cells supported the growth of hematopoietic stem cells. Therefore, the nurse cells were chemically fixed for supporting the growth of ES cells with keeping the undifferentiated state [42]. GA- or PFA-fixed MEF cells and HAE cells were prepared for culture of murine and primate ES cells, respectively. HAE was immortalized by infection with hTert cDNA as described in expansion of CB cells. Figure 6 shows the micrographs of murine ES cells cultured on several materials. Although the colony size of murine ES cells cultured on chemically

fixed MEF cells was a little smaller than that on MEF which was treated with mitomycine C, they have clear outlines. This result indicated that the murine ES cells grew with keeping the undifferentiated state. Freeze-dried chemically fixed MEF cells also supported colony formation of murine ES cells. On the other hand, on gelatin-coated surface, colony formation was not sufficient and ES cells spread out from the colonies. This indicates that the differentiation beginning on the gelatin-coated surface. On freeze-dried non-fixed MEF cells ES cells formed colonies but spread out a little.

Figure 6. Phase contrast micrographs of mouse ES cells cultured on different materials for 4 days. A: mitomycine-C-treated MEF, B: GA-fixed MEF, C: PFA-fixed MEF, D: freeze-dried MEF, E: freeze-dried GA-fixed MEF, F: freeze-dried PFA-fixed MEF, G: Gelatin. Reproduced with permission [42].

To confirm the undifferentiation state, the activity of alkaline phosphatase of murine ES cells was measured by staining (Figure 7). The cells on chemicallyfixed cells were stained stronger, even before and after freeze-drying, than that cultured on MEF treated with mitomycine-C. On the other hand, the cells cultured only on the gelatin-coated dish were not stained so much. Immunostaining of SSEA-1 and SSEA-4, which are known to be the markers of undifferentiation and differentiation of mouse ES cells, respectively, were performed and the unidifferentiation of mouse ES cells on chemically

fixed cells. Expression of transcription factor Oct-3/4 of murine ES cells were investigated by RT-PCR. On all materials mouse ES cells expressed SSEA-1 and Oct-3/4, but not SSEA-4. Considering LIF was added into all culture media, the expression was considered to be natural. However, although on gelatin surface, less expression was observed by RT-PCR.

Figure 7. Staining for Alkaline phosphatase of mouse ES cells cultured on different materials for 6 days. A: mitomycine-C-treated MEF, B: GA-fixed MEF, C: PFA-fixed MEF, D: freeze-dried MEF, E: freeze-dried GA-fixed MEF, F: freeze-dried PFA-fixed MEF, G: gelatin. Reproduced with permission [42].

Growth rates of mouse ES cells on different materials were investigated and no significant difference among mouse ES cells on the surfaces was observed. These results indicate that the chemically-fixed MEF cells supported the growth of murine ES cells with keeping undifferentiated state in comparison to native MEF cells even before and after freeze-drying. The freeze-drying was not influenced on both non-fixed MEF and chemically-fixed MEF. However, in the case of non-fixed MEF, after one time culture of mouse ES cells and their detachment it cannot support the growth of new ES cells. Since trypsin treatment is required for detachment of cultured ES cells, it decomposed the non-fixed MEF. On the other hand, chemically-fixed MEF can be repeatedly employed and effective on undifferentiated growth of mouse ES cells. The fixed cells were considered to be stable for trypsinization.

Figure 8. Phase contrast micrographs of monkey ES cells cultured on mitomycine-Ctreated MEF (A), GA-treated HAE (B), PFA-treated HAE (C), freeze-dried HAE (D), freeze-dried GA-fixed HAE (E), freeze-dried PFA-fixed HAE (F), and gelatin (G) for 4 days. Reproduced with permission [42].

Although the monkey ES cells spread out on gelatin-coated dish, the cells on MEF treated with mitomycine C or on chemically fixed HAE cells grew with forming colonies (Figure 8). On the MEF cells the colonies were all round, but on the fixed and immortalized HAE the shape of colonies were spindle. It seemed that the monkey ES cells interacted with HAE more than MEF. In this investigation, because the immortalized HAE did not survive in the presence of mitomycine C, the cells were not prepared as a control. Therefore, MEF treated with mitomycine C was employed as the control in the present study. When the cells were stained by using the activity of alkaline phosphatase, all of the colonies were stained except for the cells on gelatin-coated dish as shown in Figure 8. Staining of SSEA-1 and SSEA-4 was performed and the result indicates the cells kept the undifferentiated state on these materials. This was also supported from the data on Oct-4 expression. Freeze-drying did not reduce the supporting activity. Monkey ES cells also grew on chemically fixed HAE cells. Neither chemically fixation nor freeze drying significantly affected the supporting activity, although the growth rates were different from each other. Considering

fixation reagent, growth supporting ability of PFA-fixed HAE was a little high than that of GA-fixed one even before and after freeze-drying. On gelatincoated surface monkey ES cells grew with differentiation.

Figure 9. Staining for Alkaline phosphatase of monkey ES cells cultured on mitomycine-C-treated MEF (A), GA-treated HAE (B), PFA-treated HAE (C), and freeze-dried HAE (D) for 3 days, and freeze-dried GA-fixed HAE (E), freeze-dried PFAfixed HAE (F), and gelatin (G) for 8 days. Reproduced with permission [42].

3.2 Synthetic Biomateirals for ES Cells Many researchers have reported on the importance of the culture substrates for ES cells [43-55]. Harrison et al. [44,45] reported that fluoride-containing hydroxyapatite supported mouse ES cell growth and the biodegradable substrate affected the pluripotency. Horak et al. [46] cultured mouse ES cells on poly(2hydroxyethyl methacrylate)-based slabs. Langer’s group [47,48] cultured human ES cells in the presence of various materials. A Germany group [49] also investigated the combination of biomaterials with stem cells. Considering the importance of immobilized biosignal molecules [50], Makino et al. [51] and Nagaoka et al. [52,53] reported that immobilized LIF and cadherin supports the

growth of undifferentiated mouse ES cells, respectively. On the other hand, Kurosawa et al. [54] and Konno et al. [55] employed polypropylene tubes and phospholipid polymers, respectively, for the formation of embryoid bodies.

Figure 10. Chemical structures of biomaterials used for ES cells in this investigation. Reproduced with permission [56]

We prepared different types of polymers and covalently immobilized on conventional polystyrene tissue culture dishes to investigate the effects of polymer surface properties on the culture of mouse ES cells (Figure 10) [56]. Four types of photoreactive polymer were prepared and photoimmobilized: azidophenyl-derivatized poly(acrylic acid) (Az-PAc) as an anionic polymer, azidophenyl-derivatized polyallylamine (Az-PAllAm) as a cationic polymer, azidophenyl-derivatized poly(2-methacryloyloxyethyl phosphorylcholine-comethacrylic acid) (Az-PMAc50) as a zwitterionic polymer, and azidophenylderivatized gelatin (Az-gelatin) as a biological polymer. Phase contrast microscopic images of ES cells after three days of culture on the photoimmobilized polymers are shown in Figure 10. The ES cells immediately and spontaneously aggregated on the Az-PMAc50 surface, which comprised concentrated phospholipid polar groups. The Az-PAc surface, which was anionically charged, also induced ES cell aggregation. The ES cells appeared to form an embryoid body (EB)-like cell aggregation that did not adhere to the Az-PMAc50 surface or the Az-PAc surface. The cell aggregates

moved randomly when the Az-PMAc50 and the Az-PAc culture plates were shaken. The ES cell aggregates that formed on the Az-PMAc50 surface were larger than the cell aggregates that formed on the Az-PAc surface. In contrast, the ES cells strongly adhered to the Az-PAllAm surface, which was positively charged, but did not form aggregates. On the Az-gelatin surface, the ES cells adhered and formed colonies. The cells did not move when the Az-PAllAm and Az-gelatin culture plates were shaken.

Figure 11. Murine ES cells cultured for 3 days on synthetic biomaterials. Reproduced with permission [56]

Figure 12 shows the results of staining for ALP, a marker of the undifferentiated state of the cells, after four days in culture. On the Az-PAc and Az-PMAc50 surfaces, the cell aggregates that formed readily stained for ALP, although on the Az-PAllAm it was difficult to detect the staining of the colonies of cells because the cells did not form colonies. On the Az-gelatin surface, ES cell colonies were not densely stained for ALP, which indicates partial differentiation. It is considered that this result related to the cell spreading from colonies or the deformation of colonies, as shown in Figure 11.

Figure 12. ALP staining of murine ES cells cultured on polymers for 4 d. Reproduced with permission [56].

Figure 13 shows the increase in cell numbers on the photoimmobilized polymer surfaces after five days. The highest number of cells occurred on the Az-gelatin surface. The growth of cells on the Az-PMAc50 and Az-PAc surfaces, to which the ES cells did not adhere, was lower than the growth of cells on the Az-gelatin surface. The lowest growth rate was observed on the AzPAllAm surface, although cells adhered to this surface. These results indicate that the biological interaction of gelatin is important for cell growth, rather than the simple adhesive interaction. The LIF/stat3 signaling cascade is a mechanism of signal transduction in ES cells [57]. The cytoplasmic protein stat3 is activated by the binding of LIF to the LIF and gp130 receptors, which both reside on the ES cell membrane. Activated stat3 (phosphorylated stat3) induces the expression of the transcription factor Oct3/4. In this study, phosphorylated stat3 (p-stat3) and total stat3 (t-stat3) in cells on the immobilized photoreactive polymer surfaces were detected using western blotting. The relative intensity of p-stat3 to t-stat3 (p-stat3/t-stat3),

which indicates the activation of signal transduction, was investigated. As a result, phosphorylation of stat3 was observed on all polymer surfaces, and no significant difference between the surfaces was found. Although ALP staining was not observed in the cells on the Az-PAllAm surface by microscopy, at the molecular level of detection, the undifferentiated state of the cells was revealed. Expression of the transcription factor Oct3/4 is necessary for ES cells to maintain an undifferentiated state [58]. Phosphorylation of stat3 induces the expression of Oct3/4. We confirmed the expression of Oct3/4 by RT-PCR. When the relative expression of Oct3/4 to G3PDH (HKG) was calculated using image analysis, the same level of expression was observed in cells on all polymer surfaces (Figure 13).

Figure 13. Growth of ES cells on polymers after 5 d culture (n = 5). Reproduced with permission [56].

The expression of GATA4, which indicates the early endodermal state or differentiation to the EB [59,60], was also measured. Using the data in Figure 14, the relative intensity of GATA4 to G3PDH was calculated. The values were 0.4, 0.2, 0.18, and 0.02 on Az-PMAc, Az-PAc, Az-gelatin, and Az-PAllAm, respectively. High expression of GATA4 occurred in cells on the Az-PMAc50 surface. This indicates that the Az-PMAc50 surface effectively induced differentiation of ES cells to the EB. Strong expression of GATA4 was not observed in the cells on the Az-PAc surface, although the cells aggregated on this surface. Similar expression was observed in the cells on the Az-gelatin, but, in this case, the colonies were spread out. In contrast, cells grown on the AzPAllAm surface did not express GATA4, which we considered was due to the lack of cell aggregation.

Many researchers have investigated the behavior of ES cells cultured on various biomaterials. Properties such as hydrophilicity/hydrophobicity, electrostatic charge, and topographical roughness affect cell adhesion, growth, and differentiation. In this study, the effect of electrostatic charge on mouse ES cells was investigated in the presence of LIF. Considering that it is known that LIF activates stat 3 [57,61], and expression of Oct3/4 [58], and ALP [62], it is reasonable that the undifferentiated state of cells was maintained on all surfaces in the presence of LIF.

Figure 14. RT-PCR detection of G3PDH, Oct3/4, and GATA4 from mouse ES cells cultured on polymers for 6 d. Reproduced with permission [56].

However, the morphology, growth, and differentiation of ES cells to EBs depended on polymer surface properties. Neither Az-PAc nor Az-PMc50 surfaces induced adhesion of ES cells, which led to EB formation. When an EB did form, the growth of cells was reduced. Negative charges or zwitterions reduced the interaction of ES cells with the polymer surface and the cells aggregated to form an EB. EB size and expression of GATA4 (an indicator of EB formation) in cells grown on the Az-PMAc50 surface were greater compared with cells grown on the Az-PAc surface. Recently, we reported that EB formation was efficiently induced on a phosphorylcholine-derived polymercoated surface [55]. Therefore, the reduction of cell-surface interactions with this polymer would enhance the cell–cell interactions required to form an EB, and thus explain the larger size of the EBs and stronger expression of GATA4 on the Az-PMAc50 surface.

ES cells adhere to cationic or gelatin surfaces. However, mouse ES cells formed colonies on the gelatin surface, but not on the Az-PAllAm surface. Some biological interactions are required to induce ES cells to aggregate and adhere to a surface. The growth of ES cells was significantly enhanced on the Az-gelatin surface, and GATA4 expression was slightly induced. These phenomena are the same as those for normal gelatin-coated surfaces. Recently the group of Chen [63] investigated chondrogenic differentiation of human mesenchymal stem cells (MSCs) on photoreactive polymer –modified surfaces. They reported that Az-PAllAm-modified surface supported cell adhesion and proliferation and also promoted chondrogenic differentiation of the MSCs. The Az-Pac-modified and polystyrene surfaces supported cell adhesion and differentiation, but not differentiation. Polyethylene glycol (PEG)modified surface did not cell adhesion, but promoted chondrogenic differentiation.

4

Conclusion

Chemically fixed human stromal, MEF, and HAE cells have the supporting ability in in vitro expansion of human CB progenitor cells, murine ES cells, and primate ES cells, respectively. Considering that the chemically fixed cells reduce the biological contamination of nurse or feeder cells and can be maintained in refrigeration preservation or freeze-dried, the system is very convenient for culture stem cells. However, ideally a defined artificial substrate without any biologically derived components is desired for stem cell culture. Therefore, as a fundamental research, some different surfaces were prepared and biological interactions between ES cells were investigated. Not only biological properties but also physic-chemical properties significantly affected the behavior of stem cells. In future, by consideration on these results, new biomaterials will be designed and developed for stem cell culture.

Acknowledgments This work was supported by Kanagawa Academy of Science and Technology (KAST) and the Ministry of Education, Culture, Sports, Science and Technology (MEXT) of Japan.

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PART III

Polymeric Biomaterials

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Chapter 9 Design of Artificial Extracellular Matrices and Their Application for Regenerative Medicine Xiaoshan Yue1, Masato Nagaoka2, Akaike Toshihiro1 1. Department of Biomolecular Engineering, Graduate School of Bioscience and Biotechnology, Tokyo Institute of Technology, 4259-B-57 Nagatsuta-cho, Midori-ku, Yokohama 226-8501, Japan 2. Department of Cell Biology, Neurobiology and Anatomy, Medical College of Wisconsin, 8701 Watertown Plank Road, Milwaukee, WI 53226, USA

1

Introduction

Most of the cell functions are regulated through attachment to extracellular matrices, or through intercellular interactions, while cytokines and growth factors were also well known to regulate cell functions as well. In the developmental process, all these complicated events together controlled the formation of certain organs. With the establishment of embryonic stem (ES) cells, it has become no more a dream for researchers to study on developmental process at cellular levels. As ES cells have provided a powerful tool for the research of regenerative medicine, the studies on ES cells has turned out to be a more and more attractive field. In order to actualize regenerative medicine, the research of extracellular matrices, which plays a central role for cell attachment, has been focused by a lot of researchers. Here, the “extracellular matrices” indicates not only the normal matrices existing around cells in vivo, but also includes artificial matrices designed with improved functions. In this review, we are going to have a brief discussion on the research work related to extracellular

Figure 1. Generation of stem/progenitor cells and their application.

matrices, especially focusing on the artificial matrices incorporated with intercellular adhesion molecules, cytokines, and growth factors, for controlling cell functions.

2

Extracellular Matrices for Regenerative Medicine

2.1 Regenerative Medicine by Using Stem Cells Since several years ago, with the increasing requirement for tissue therapy and treatment for organ damages, the research on regenerative medicine has turned out to attract more and more researchers’ attention. However, the lack of donors and the immune rejection problems have restricted their broadened clinical application. In order to regenerate certain organs and to develop transplantation therapy, lots of efforts have been made on artificial organs, and the establishment of ES cells has shed light on researches on regenerative medicine [1-3]. ES cells as well as progenitor cells are with self-replication ability and the potential to differentiate into required cell types under certain induction conditions. ES cells were established from the inner cell mass of blastocyst that could develop into all tissue of embryo, while progenitor cells are found in certain organs and tissues of the body with a restricted potency to differentiate into multi-lineage

cells (Fig. 1). Lots of researchers had devoted themselves trying to induce ES cells or progenitor cells differentiating into certain cell types with high efficiency. By using these cells, the research of regenerative medicine has developed to a new height, and with the establishment of human ES cells and induced pluripotent stem (iPS) cells, new approaches have been made in the application of modern medical treatment. However, up to now, the mechanism which controls the differentiation and proliferation process of ES cells and progenitor cells are still not well understood, while the mechanism underling the maintenance of pluripotency is also not well studied.

Figure 2. Environmental signals regulating ES cell functions.

According to recent researches, ES cell functions are regulated by environmental signals mainly from four types of process or stimuli: (1) stimulation from growth factors and cytokines [4, 5]; (2) uptake of molecules through endocytosis or transporters [6]; (3) cadherin-mediated cell-cell adhesion [7]; (4) integrin-mediated cell-matrix interactions [8] (Fig. 2). Later in this review, we will introduce the artificial extracellular matrices for regulating ES cell functions, depending on several kinds of the stimuli mentioned above.

2.2 Extracellular Matrices for Controlling ES Cell Differentiation Extracellular matrices are cell attachment materials and essential for the maintenance of tissue architectures. In addition, extracellular matrices regulate differentiation and proliferation of cells during developmental process. The extracellular matrices are secreted by cells that exist in a given tissue or organ, so that the composition and ultrastructure of the ECM are determined by factors that influence the phenotype of these cells [9, 10]. In turn, different composition of extracellular matrices has different modulating effects for cells [10, 11]. For example, it has been reported that the direction of cell differentiation can be affected through different ways by changing compositions of extracellular matrices [12-14]. Also, with the application of certain types of extracellular matrices, ES cells could be induced particularly into hepatocytes and neural cells [15, 16]. Therefore, it is considered as a very important task for us to study the effect of extracellular matrices and also the cellular molecular changes induced by changing the components of extracellular matrices.

2.3 Regulation of ES Cells Differentiation by Other Signal Induction As mentioned previously, cell functions are regulated not only by extracellular matrices, but also by cell-cell interactions, growth factors and cytokines as well. For cell-cell interaction, signal transduction through the interactions between ligands and their receptors such as cadherin/cadherin, Eph/ephrin, Notch/Delta, and so on, can also execute important roles. However, the signaling pathways from these factors are very complicated, and it is still difficult to analyze each signaling pathways separately. Furthermore, in vivo, cytokines and growth factors not only exist in soluble form, but also these factors are immobilized on the cell surface polysaccharide side chains or onto extracellular matrices [17, 18]. Furthermore, these two different forms, the soluble form and the immobilized form, might activate different signal pathways in different time span and with different strength [18]. These researches indicated that with the spatial-temporal control of these factors, it might be possible to control cell functions in a better exactitude. Also, since the differentiation processes is regulated by these extracellular factors and intercellular interactions as well, more precise analysis is required to clarify the effects of cell-cell interactions and extracellular factors in controlling differentiation process.

2.4 Application of Artificial Materials to Regenerative Medicine Extracellular matrices are known as cell attachment materials. However, nowadays the researches of extracellular matrices are no more restricted to the traditional studies on natural matrices such as collagen, fibronectin, and laminin. Recently, many studies concerning developing artificial materials with the application of genetic technology and molecular manipulation were also reported. In other words, with molecular technology and macromolecular synthesis methods, we could produce specific artificial matrices based on chemically synthesized materials with specific physical or chemical properties, and endow them with cell recognizable characteristics and abilities to regulate cell functions. For example, cell sheet engineering by using temperaturesensitive materials [19] is applied for the nondestructive harvesting of differentiated cells which could be used in transplantation therapy. Also, by patterning or spatial-temporally organizing artificial molecules, we can bring these artificial matrices into different applications especially for regulating cell functions. Up to now, we have designed and produced several kinds of biomaterials. Hepatocytes specific material PVLA (poly N-p-vinylbenzyl-O-　-dgalactopyranosyl-[1→4]-d-gluconamide), taken as an example, was firstly synthesized by Professor Kobayashi from Nagoya University, and was found to be able to function as an artificial matrix for hepatocyte culture [20, 21]. Actually, several kinds of glycan-carrying polymers were synthesized at first, and among these polymers, we found that only PVLA which carried a galactose part could be specifically recognized by the cell surface receptor (asialoglycoprotein receptor: ASGR) of hepatocytes (Fig. 3). By controlling coating concentration of PVLA, we can control cell shape and differentiation state of hepatocytes. On the surface coated with high concentration of PVLA, bile acid release as a typical differentiated function of hepatocytes was maintained at higher levels. Furthermore, since the expression of ASGR in the progenitor cells of hepatocytes is kept at a low level, they would attach to PVLA-coated surface at low efficiency and so that could be separated from the whole liver by using PVLA-coated surface [21-23]. This might provide us a useful method for getting the progenitor cells for regenerative application to injured liver. Besides PVLA, we also reported the effect of using alginate/galactosylated chitosan/heparin scaffold as a culture system for hepatocytes. A highly porous hydrogel (sponge-like) scaffold, 150-200 m pore size in diameter, was fabricated with alginate, galactosylated chitosan and

heparin, through electrostatic interaction, and applied to culturing hepatocytes. It was confirmed that hepatocyte spheroids were more rapidly formed in this culture system, and these stable spheroids might enhance liver-specific functions, so that this might provide us a new synthetic ECM for designing bioartificial liver devices [24].

Figure 3. Molecular structures of several glycopolymers.

3

Design and Application of Engineered Matrices by Genetic Technology

3.1 Design and Application of Fusion Protein for the Regulation of Cell Function As described previously, cell functions are regulated by the crosstalk between signals from extracellular environment; therefore, the analysis of the signaling pathways, especially cell-cell interaction, was considered to be complicated and difficult. In addition, the effect of immobilized form of growth factors is still unclear.

In our study, we tried to design several cell recognizable fusion proteins with molecular technology and gene manipulation. These fusion proteins consist of the functional domain and the immobilizable domain such as the Fc domain of mouse immunoglobulin (IgG), which can be adsorbed to polystyrene surface with a hydrophobic interaction. We used intercellular adhesion proteins including E-cadherin and N-cadherin; and growth factors including EGF, HGF; cytokines including LIF, Activin, BMP, and Wnt. These fusion proteins are applied to cell culture and to the induction of differentiation of stem cells. Furthermore, these proteins are expected for the application in analysis of signaling pathway activated by target molecules and for controlling the direction of ES cell differentiation. Here, we are going to take these fusion proteins for examples to discuss applications of artificial extracellular matrices to the research of regenerative medicine.

3.2 Application of Artificial Extracellular Matrices to ES Cell Culture In order to establish the new culture system for ES cells, immobilized model proteins are designed as novel artificial extracellular matrices. Mouse ES cells are conventionally cultured on feeder cells of mouse embryonic fibroblast or on gelatinized surface, which result in the aggregated colony formation. Several reports suggested that the aggregated colony formation induce the heterogeneity of cell population in the colonies. Therefore, the novel method that could maintain ES cells in the homogeneous environment should be useful for the maintenance of ES cell feature and for the single cell-based analysis of ES cells. E-cadherin is a well-known intercellular adhesion molecule that is important for the early developmental process and cell sorting, and this molecule is highly expressed in ES cells; therefore, we tried to apply the fusion protein of Ecadherin and Fc domain of IgG to improve the method of mouse ES cell culture. On an E-cad-Fc-immobilized surface, ES cells proliferated without formation of colonies but show a scattering morphology, which is quite different behavior from conventional culture systems [25]. However, since up to now no one has ever observed the phenomena that ES cells proliferate while keeping a scattering morphology, it was questioned whether the characteristics of ES cells cultured on E-cad-Fc surfaces are changed or not. In order to clarify this issue, we examined two of the most important characteristics which distinguish ES cells from other types of cells: the pluripotency and self renewal ability. As a result, we found that with the

Figure 4. Maintenance of undifferentiated state of ES cells on E-cad-Fc surface. (A) Immunostaining results of undifferentiated cells with Oct-3/4; (B) RT-PCR results for checking the expression of undifferentiated marker genes.

addition of LIF, ES cells cultured on E-cad-Fc-coated surface could maintain the expression of marker genes for undifferentiated state (oct-3/4, rex-1, and nanog) as same level as those on gelatin-coated surfaces (Fig. 4). This indicates that the undifferentiated state is well maintained on E-cad-Fc-coated surfaces. Next, we checked the pluripotency of ES cells cultured on E-cad-Fc-coated surfaces. To induce multiple differentiations, ES cells are induced to form an

aggregated structure of embryoid body (EB), which mimics developmental process to differentiate into three germ layer-derivatives. Here in order to confirm the maintenance of the ability to differentiate into multiple cells, ES cells cultured on E-cad-Fc-coated surface or gelatin-coated surface were induced to form embryoid bodies for the induction of differentiation. ES cells cultured on E-cad-Fc-coated surface revealed a decrease of undifferentiated marker genes such as oct-3/4, rex-1, and nanog, while the expression of maker genes for the three germ layers (neurod3 for ectoderm, gata-1, T/brachyury, flk-1, hbb for mesoderm, and -fetoprotein, transthyretin, vitronectin for endoderm) increased in a similar pattern as that on gelatinized surface. Furthermore, we confirmed the maintenance of pluripotency by the formation of teratoma and chimera mouse, and the ability of germline transmission was also confirmed (Fig. 5). From these results, we indicated that when cultured on E-cad-Fccoated surface, ES cells showed a unique morphology without losing their characteristics as the stem cells [25].

Figure 5. A. RT-PCR results of expression of marker genes for three germ layers. 1. Undifferentiated ES cells; 2. ES cells cultured on gelatin coated surfaces; 3. ES cells cultured on E-cad-Fc coated surfaces. B. Germline transmission.

In addition, the proliferative activity of ES cells was enhanced when cultured on E-cad-Fc-immobilized surfaces (Fig. 6A, B). Since the proliferation of cells was inhibited when ES cells are cultured in high density even on E-cadFc-coated surface, the intercellular interaction might inhibit cell proliferation.

Furthermore, transfection efficiency was also enhanced on E-cad-Fc surface (Fig. 6C). With all these merits, E-cad-Fc immobilized system can be expected to be applied as a novel culturing system for mouse ES cell culture [25].

Figure 6. A. Proliferative ability of ES cells cultured on gelatinized (open square) or on E-cad-Fc-coated surface (filled square). B. The DNA synthesis activity was measured by the incorporation of BrdU. C. Transfection efficiency. ** P < 0.01,§P < 0.001.

ES cells form colonies in traditional culturing system; however, we can not say that colony formation is a prerequisite or the best condition for ES cell culturing. Firstly, with the formation of colonies, heterogeneity might be generated in the microenvironment around the cells. Cells located at the basal position of colonies could directly interact with basal matrices so that the signals from basal matrices might be enriched, while cells located at the apical position of colonies might be easier to accept the stimulation of growth factors or cytokines which exist around the colonies. On the other hand, at the inner part of colonies, intercellular interaction might be strengthened. As a result, every

single cell is maintained in a distinct microenvironment. This may cause the heterogeneous signal transduction depending on the place of colonies where the cell exists. Actually, recently several studies suggested that heterogeneity might be caused as a result of colony formation both in mouse ES cells [26-28] and in human ES cells [29, 30]. For example, colony formation was reported to induce the heterogeneous expression of Nanog and SSEA-1 [26]. Also, differential responsiveness to gp130 ligands within colonies was found to be organized in a radial manner [27]. On the contrary, when cultured on E-cad-Fc coated surface, ES cells can be maintained in a homogeneous condition so that the heterogeneous environment caused by colony formation in the conventional culture system could be eliminated. Furthermore, by using E-cad-Fc-coated surface, the single cell-based assay should be possible. In this way, E-cad-Fc immobilized system could be expected for its application in mechanism analysis of ES cells and also for the establishment of differentiation induction system for ES cells. All these approaches might be applied broadly in the field of basic research and even for clinical achievement in regenerative medicine. In addition, we also designed the immobilizable molecule of LIF to improve the ES cell culturing system. LIF is a pleiotropic cytokine that inhibits the differentiation of mouse ES cells. Several reports indicated that LIF could be immobilized on the extracellular matrices and that the signaling pathways are different; therefore, we designed the immobilizable model protein of LIF as same procedure as E-cad-Fc and applied for the culture of mouse ES cells. On the surface co-immobilized with LIF-Fc and E-cad-Fc, ES cells can be maintained at undifferentiated state without additional soluble LIF supplementation, and activation of LIF signal can be sustained [31]. In this way, with the idea to design novel artificial extracellular matrices by corporately immobilizing E-cad-Fc and other cytokines such as LIF-Fc, it might provide a new method for ES cell culture and also for the analysis of signaling pathways activated by cytokines or growth factors immobilized in the extracellular matrices.

3.3 Application of Artificial Extracellular Matrices for Controlling Cell Functions In order to control stem cell function, there are several studies that design artificial matrices with the technology to immobilize functional molecules. For example, Beckstead et al. focused on Notch signal pathway which is activated through intercellular interaction, and they tried to induce ES cell differentiation

on surfaces immobilized with Notch’s ligand, Jagged [32]. Here, we focused on N-cadherin molecule which is highly expressed in cardiomyocytes and neural cells and was though to be an important molecule during the differentiation into neural or cardiac cells. To analyze the direct effect of N-cadherin mediated interaction during the differentiation process, we designed an immobilizable Ncadherin model molecule (N-cad-Fc) with the same method mentioned above and applied this model protein to the research of the differentiation process into cardiomyocytes or neural cells. We used P19 cells as a model cell, which could be induced to differentiate into neural cells or cardiomyocytes/skeletal muscle cells with the supplementation of retinoic acid or DMSO, respectively [33]. With traditional differentiation inducing methods, cells are first cultured in a suspending condition and induced to form aggregates (embryoid bodies). Then these embryoid bodies are transferred to new dishes to help them attach to ECM and differentiation will occur successively [33]. As discussed above, the formation of aggregates might cause heterogeneity for every single cell. In contrast, on an N-cad-Fc-coated surface, P19 cells were maintained as single cells with the same morphology as those cultured on E-cad-Fc-coated surface, and the differentiation into neural cells was induced in single cell level without the requirement for EB formation (data not shown). However, when neural differentiation was induced on the surface coated with high-concentration of Ncad-Fc, the expression of neural marker genes such as nestin, map2, neurod3 was repressed comparing to the traditional gelatin-coated surface, while the expression of cardiac marker genes such as cardiac actin, gata-4 was promoted (data not shown). These results indicated that the direction of differentiation into neural cells or cardiomyocytes might be regulated by the strength of interaction with N-cadherin molecules. By immobilizing cell adhesion molecules and growth factors related to cell differentiation process, it might become possible to analyze signaling pathways from a single target molecule. Also, by immobilizing molecules with proper proportion, the direction of differentiation of stem cells might also be controlled. As a supplement, besides physical immobilization methods, chemical immobilization methods were also studied. For example, the immobilization of erythropoietin was reported for culturing erythropoietin-dependent human leukemia cell lines [34]. In addition, the immobilization of growth factors such as EGF was also reported for their application in analyzing signal pathways from the immobilized form of growth factors [35, 36]. The chemical immobilization, however, might cause the decrease of activity, and it is still difficult to control a site-specific immobilization. The method we used is simple

and easy; however, we could not control the orientation of immobilized protein under current conditions. Therefore, the improvement of technology to immobilize the target molecules should be needed for further application.

4

Conclusions

Most of cells, including stem cells, keep their functions through cell-matrix interaction and also cell-cell interaction, and they can response to a variety of inductions caused by soluble and immobilized factors existing in surrounding environment. However, as a result of the complexity of signaling transduction caused by the variety of factors and also their crosstalks, there are still a lot of unsolved problems left in the research of molecular mechanisms of every single factor. Therefore, researching signaling pathways focused only on the target molecules was regarded as very important work. This might be realized by using the immobilizing systems. In addition, the design of growth factors and other molecules related to cell response might be very important for the control of cell fate. Furthermore, the induction of spatial-temporal controlled signals and the construction of more elaborate artificial extracellular matrices could be realized by using micro-patterning technology. Since the recombinant proteins are easily modified, we can improve functions of these natural molecules and even spatial control of their activity. As a result of all the development in biomaterials, it can be expected that in the future, more and more artificial extracellular matrices would be made out and be applied into researches or even clinical trials for regenerative medicine.

Acknowledgments This work was supported by Grant-in- Aid for Scientific Research S (15100008) and A (19200038), and a Grant of the 21st Century COE Program.

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Chapter 10 Nano-to-macro Architectures Polycaprolactone-based Biomaterials in Tissue Engineering Teoh Swee Hin, Bina Rai, Tiaw Kay Siang, Chong Seow Khoon Mark, Zhang Zhiyong and Teo Yiling Erin Centre for Biomedical Materials Applications and Technology, Department of Mechanical Engineering, National University of Singapore

1

Introduction

Current applications of biomaterials for tissue engineering (TE) involve the combination of a scaffold with cells and/or biomolecules that promote the repair and regeneration of tissues. These TE constructs are under vigorous exploration and have resulted in the incessant emergence of diverse strategies. However, despite intense efforts, the ideal tissue engineered construct has yet to be discovered. From an engineering perspective, significant advances in scaffold design and fabrication have evolved in the past decade. Our interdisciplinary group at the National University of Singapore have evaluated the parameters necessary to process polycaprolactone (PCL) and PCL-based composites by fused deposition modeling (FDM). Other cutting-edge technologies such as biaxial-stretching and electrospinning, have also been utilized to manufacture PCL-based biomaterials with unique macro and nano-architectures that are useful for drug, growth factor and cell delivery applications. The present chapter serves to introduce the developments achieved by our group with respect to the extensive applications of PCL-based biomaterials investigated, with special emphasis on bone TE and the importance of structure-property relationship in biomaterial design.

2

Development of “Cell/PCL Film” Sheet Technology

Functional soft tissues are anisotropic and heterogeneous, comprising of mixed populations of cells organized in defined orientations. This spatial organization is crucial for proper tissue function. In the heart, for example, this organization is required to effect functional contraction [1]. Several investigators apply cell sheet approaches to address this issue of tissue organization (Table 1). However, cell sheets are mechanically weak, making handling difficult and unsuitable for direct implantation into high stress applications. The generation of a functional small diameter blood vessel, for example, will require up to four months of culture before maturity [2]. Cell sheets can be used in conjunction with scaffolds, to provide mechanical support in the interim period prior to tissue maturation [3-6]. However, these scaffolds are invariantly isotropic, with no concessions for the formation of compartmentalized tissue. Table 1. Summary of cell sheet-based research. Institution Tokyo Women’s Medical and Dental University, Japan Cytograft Tissue Engineering, USA National University of Singapore, Singapore Tokyo Women’s Medical and Dental University, Japan Harvard Medical School, USA

Target tissue Trachea, myocardium, liver, cornea Blood vessel

Ref 7-10 2, 11

Bone, skin

5, 6

Skin

3

Genitourinary

4

Remarks First demonstration of cell sheet concept, pre-clinical stage Whole blood vessel formed, pre-clinical stage Cell sheet used in conjunction with scaffold Cell sheets cultured on collagen membranes Multilayer cell sheets used in conjunction with scaffold

A possible solution would be to seed the cells on bioresorbable films, and subsequently assemble the tissue construct layer-by-layer using these cell-film constructs (Figure 1). This configuration further allows specific functional compartments to be generated, which best approximates native architecture. The films will provide sufficient mechanical strength for a defined period of time, allowing earlier implantation and further maturation within the native microenvironment. We have, in our laboratory, developed unique biaxially-stretched PCL films for use as scaffolds in the creation of cell-film sheet constructs [12-16]. These films are pliable and facilitate mechano-transduction during the remodeling

phase in vivo. The film thickness can be varied to achieve different stiffness, which has been shown to impact cell responses [17]. In addition, the slow degradation rate of PCL allows for retention of mechanical properties during remodeling, as well as reduced acidosis arising from the accumulation of degradation products [18].

Stack

PCL film

Nanofibers

Perivascular Cell Compartment

Endothelial Cell Compartment Shear

Heparin

Blood Flow Figure 1. Schematic diagram of stratified tissue created using layer-by-layer construction to produce an anisotropic vessel graft fabricated from polycaprolactone.

Various groups have studied the use of solvent cast PCL films for tissue engineering applications, and have shown the material to be compatible with many cell types, including osteoblasts, smooth muscle cells and endothelial cells [19-20]. PCL can also be electrospun to form nanofibrous mats for vascular tissue engineering applications [21]. However, such PCL films have low mechanical strength, as compared to other scaffold materials, such as PLLA and PGA. Thus, biaxial stretching was carried out to improve the mechanical properties of the films [22]. Following solvent casting and heat pressing, the films are stretched in orthogonal directions, resulting in a doubling of tensile strength, accompanied by a reduction in thickness of films by ten-fold. The use of roll milling as a solvent-free approach to generate these films, and the use of spin casting techniques to fabricate sub-micron thick films were reported recently [16]. In assembling the constructs, diffusion of nutrients, particularly oxygen is crucial to allow stacking of multiple layers. It was previously demonstrated that

the stretched films permitted the permeability of oxygen by Knudsen mechanisms [13]. The use of needle perforators and laser drilling techniques further improved the permeability, with a 50% increase in water vapor transmission rate [23-24]. Aside from altering bulk physical properties, surface modifications are often employed to confer bioactivity to the otherwise inert synthetic materials. Cell responses are directly mediated by interactions with biomaterial surfaces [25]. Although PCL films are hydrophobic and not optimal for cell attachment, the surface can be readily modified with sodium hydroxide to improve wettability [13]. These films have been shown to support the functional adhesion of various cell types, namely periosteal cells, dermal fibroblasts and keratinocytes [12, 13, 26]. Wet chemistry methods such as UV engraftment can be employed to functionalize the surface as well [15, 27]. Biomolecules, such as collagen can then be immobilized onto the film surface to facilitate improved cell attachment and proliferation of various cell types, including myoblasts and smooth muscle cells [15, 28]. Cold discharge plasma, which is less penetrative, has been demonstrated for use in the immobilization of hydrogel layers onto polymeric films [29]. In our lab, this method has been applied to stretched PCL films with similar effect, and reduced degradation of mechanical properties. Aside from surface chemistry, morphology and topography exert strong influences on cell responses [30-31]. Nanotopography was shown to reduce mesenchymal stem cells’ adhesion and increase motility [32]. Taken into the context of tissue engineering, this may favor more rapid tissue ingrowth and integration. We have recently demonstrated the use of electrospinning to introduce nanotopography onto flat film surfaces, resulting in increased surface roughness and wettability, suggesting the possible use of this method to modulate cell responses to tissue engineered scaffolds [33].

3 PCL Films with Laser-Modified Surfaces The use of lasers in the surface modification of polymers offers many advantages over chemical and physical methods. Laser processing allows precise modification of the surface and keep the bulk properties of the materials intact while ensuring that the resulting modified surfaces are free from contaminants [34].The maintenance of the bulk properties of the polymer is essential in preserving the biological and structural properties of the

biomaterials which allow it to repair, maintain or enhance the bodily structure and function [35]. There has been much research on the effect of laser modifications on the surface of polymer films. Excimer lasers, such as KrF (λ = 248 nm), ArF (λ = 198 nm) and ultra-fast femtosecond lasers, were widely reported to be efficient tools for precision micro-structuring of polymer materials [36-39]. Excimer lasers make use of the photo-chemical effect [40] whereby polymer chains are broken down by the high energy photon absorption while the femtosecond laser causes material breakdown through multi-photon absorption (MPA) [41-42]. These two ablation mechanisms would result in minimal photo-thermal effect, and are suitable for precision laser micro-fabrication of materials. Studies had also been conducted to improve the surface wettability of polymer films. Tiaw et. al. [34] found that different surface morphologies and sizes of perforation on PCL films (Figure 2) through different laser parameters can reduce the water contact angle, hence improving the wettability of the film. Revez et. al. [43] reported that the photochemical treatment of PTFE using excimer lasers caused the surface to become hydrophilic. Others have also reported on the effects of irradiation wavelength and atmospheric gases on the polymers. Lippert et. al. [44] studied the effects of irradiation wavelength of triazenopolymer and discovered that chemical modifications and micro-structural changes are dependent on both wavelength and threshold fluence of the polymer. Bityurin et. al. [45] demonstrated that both modification and etching kinetics of PMMA were strongly influenced by ambient atmosphere. CO2 laser has also been used to treat the surface of polydimethysiloxane (PDMS) membrane to create a super-hydrophobic polymer while keeping the bulk properties of the substrate intact [46], and on polyethylene terephthalate (PET) membrane to induce morphology changes that subsequently affects the contact angle at different pulses [47]. In the biomedical field, modifying the surfaces of biocompatible polymers can lead to improved interactions between the polymer surface and the surface of cells. It is thus a technique that is gaining much recognition in recent times to engineer the surface of biomaterials suitable for specific tissue engineering applications [48].

(A)

(B)

(C)

Figure 2. Micro-perforations produced on PCL films using (a) Femtosecond, (b) Nd:YAG and (c) KrF excimer laser.

4

PCL Scaffolds for Drug Delivery Applications

An important requirement on advanced tissue engineering scaffold is the accurate controlled delivery of bioactive factors and drugs to enhance and guide tissue regeneration for therapeutic purposes [49]. Extensive studies were conducted on the release kinetics of polymeric materials for drug delivery applications and it was found to depend generally on the properties of polymers, drug characteristics and environmental conditions [50-60]. These polymer properties include hydrophilicity, bioresorbability and bioerodibility, degradation rates, swelling behavior and general permeability of the polymers. Drug characteristics also affect the delivery mechanism, specifically the molecular weight of the drugs and its solubility. Drug release is based on 2 main types of release mechanisms, or a combination of them, namely [60-66], diffusion and dissolution/degradation. Diffusion-controlled drug release mechanism works on the basis of either (a) drug release through rate-controlling membrane from a reservoir, resulting in

zero-order release or (b) through drugs dispersed in polymer matrix without a rate controlling barrier, resulting in non-zero order of release. In the dissolution/degradation paradigm, drug release is dependent on the degradation or dissolution of the polymer matrix. For example, PCL is generally investigated for long term control release of drugs due to its slow degradation rate via bulk hydrolysis of its ester bonds [66-72]. In most cases, drug delivery occurs as a combination of type 1 and type 2 diffusion mechanisms, whereby diffusion is accompanied by simultaneous polymer degradation. Surface area is a controlling factor for both type 1 and type 2 delivery mechanisms, thus influencing the rate of drug release. Previous studies have shown that an increase in the surface area to volume ratio (SA: VR) of a bioresorbable polymeric carrier resulted in an increase in the rate of drug delivery [73]. The surface area of the drug delivery scaffold can be controlled by its architecture. In recent advances of fabrication technology, customized 3D scaffold with honeycomb architectures can be obtained by fused deposition modeling [74]. For example, PCL can be fabricated into honeycomb scaffolds of controlled porosity using micro-extrusion and nanofibrous mesh can be generated by electrospinning, as shown in Figure 3. (A)

(B)

Figure 3. Scanning electron micrograph of (A) Scaffold fabricated using micro-extruder and (B) Film fabricated via electrospinning.

PCL has a relatively slow degradation rate, and the short term drug delivery mechanism prior to its degradation is largely diffusion-limited. It is therefore vital to control the fabrication parameters, such that the surface area of the scaffold can be engineered for optimal type 1 release. Using micro-extrusion technology, the porosity of the drug delivery scaffold can be controlled to vary surface area, and consequently, the SA: VR. This relationship is clearly shown in Figure 4. One can then customize the rate of release based on the surface area by choosing a specific porosity for the delivery scaffold.

Using empirical analysis, comparing PCL delivery scaffolds of similar dimension (50x50x0.5mm), the scaffold (50% porosity) fabricated using microextruder method with laydown pattern of 00/900 is able to be achieve a SA: VR of 7.94 as compared to 4.08 for non-porous film fabricated by heat press or solvent casting method. The micro-extruded scaffolds have almost double the VR value. It is therefore expected that the rate of drug delivery is higher for micro-extruded scaffold as compared to non-porous films. In conclusion, the architecture, specifically the SA, of the drug delivery scaffold is a crucial criterion to be considered when engineering a specific rate of drug delivery. Consequently, the diffusion and degradation kinetics for a drug-polymer system can be designed to cater for the drug release profile desired.

Surface area to vol ratio (mm^-1)

18 16

Relationship between porosity and surface area to vol ratio

14 12 10 8 6 4 2 0 -2 0

20

40

60 Porosity (%)

80

100

120

Figure 4. Relationship for porosity of scaffold against total surface area for drug delivery.

5

PCL-TCP Scaffolds for Growth Factor-based BTE

The goal of bone repair grafts is to restore vascular and bone architecture with biological and mechanical properties analogous to intact bone. To achieve this, our group introduced a bone regenerative strategy consisting of PCL - 20% tricalcium phosphate (PCL-TCP) composite scaffolds loaded with platelet-rich plasma (PRP). Synthetic scaffolds when implanted into bone defects are generally encapsulated by fibrous tissue and become isolated from the surrounding bone, limiting their use in bone repair [75]. This evoked the development of bioactive composite materials comprising of a biodegradable polymeric phase and a

bioactive inorganic phase that can spontaneously bond to and integrate with living bone. The inorganic phase may provide a pH buffering effect on the polymer degradation behavior that in polyesters is strongly acidic [76-77]. Ceramics exhibit a high binding affinity for proteins mainly via electrostatic interactions and are widely used as carriers for growth factors [78]. For example, β - TCP particles are common carriers for recombinant bone morphogenetic protein-2 and this combination augmented bone repair [79-81]. Finally, ceramics may enhance the mechanical strength of the biodegradable polymer [76-77]. The above-described factors instigated our research team to design and fabricate novel, three-dimensional bioactive composite scaffolds consisting of PCL physically blended with 20 % of TCP microparticles (Figure 5). The PCLTCP mixture is converted to filament form, and then fabricated to scaffolds by a form of rapid prototyping technology called fused deposition modeling (FDM) [75, 83]. The scaffolds produced have a unique honeycomb, fully interconnected matrix architecture [84-85]. The inclusion of 20% TCP provided the composites with bioactivity by nucleating a hydroxyapatite layer on its surface [86]. These scaffolds were also proven as viable in vitro delivery systems for rhBMP-2 [87-88] and PRP [89], thus conferring an osteoinductive nature.

Figure 5. Scanning electron micrographs of 3D (A) PCL and (B) PCL-TCP scaffold structure produced by fused deposition modeling and having a 0/60/120° lay down pattern. Both pictures represent the top view, displaying a typical array of equilateral triangles. TCP was evident as particles on the walls of the PCL scaffold. [73]

A rat nonunion femoral defect model together with contrast enhanced microcomputed tomography (micro-CT) was adopted to test this bone regenerative approach [90]. Each rat received a PRP loaded PCL-TCP scaffold in one leg with the contralateral defect receiving a plain PCL-TCP scaffold. This properly

controlled 3-month study provided solid evidence of PRP’s beneficial effect in stimulating vascularisation and extracellular matrix production when combined with PCL-TCP scaffolds (Figure 6A). Our study revealed that eighty-three percent bone unions across critical-sized femoral defects were achieved when treated with PRP delivered by PCL-TCP scaffolds after 12 weeks of implantation. The micro-CT images revealed that control femurs had random bone islands throughout the defect site, while PRP-treated femurs achieved bone unions with callus formation at 12 weeks (Figure 6B). This concurred with our mechanical tests, with higher stiffness demonstrated by PRP-treated femurs. These results suggest strongly that our proposed strategy, that is, PRP delivery by PCL-TCP scaffolds, were capable of stimulating angiogenesis, and subsequent osteogenesis in a rat critical-sized femoral defect model. This is on par with the body’s natural healing response where a vascular network must be established first to provide nutrients and gaseous exchange to facilitate the initial deposition and preceded by maintenance of bone repair and remodeling processes. The scaffolds architecture and porosity allowed for the infiltration of blood vessels as well as minerals.

(A)

(B)

Figure 6. Micro-CT images of (A) Vascularisation at 3 weeks and (B) Bone union, with callus formation after 3-months treatment with PRP-loaded PCL-TCP scaffolds. [75]

In a separate study, PRP loaded PCL-TCP scaffolds tested in a dog dental implant model, accelerated wound healing and stimulated mandibular bone regeneration in 8 mongrels for 6 and 9 months [91]. PRP-treated defects exhibited significantly higher bone volume fractions (BVF) compared to controls at both time points, but the difference was not as pronounced as it was at six months. An increase in BVF was detected for PRP-treated and control groups from six to nine months. Compared to the previous rat study, the BVF

values obtained were higher as was anticipated due to the longer implantation times. Histological analysis showed that PCL-TCP scaffolds experienced 33% degradation from six to nine months, finally occupying only 46.9% of the crosssectional area. Looking at the innermost section of defect site, it was noted that the entire area surrounding the implants was mostly covered with new bone, indicating bony ingrowth into the center of the scaffold. Hence, PCL-TCP scaffolds together with PRP facilitated the placement and subsequent stability of the dental implants for up to nine months.

6 PCL-TCP Scaffolds for Cell-based BTE In combination with the cutting-edge PCL-TCP scaffolds, we are exploring suitable cellular sources for cell-based BTE. The characteristics of an appropriate cellular source would include: (1) no or low immunogenicity, (2) no tumorigenicity, (3) immediate availability, (4) availability in pertinent quantities, (5) rapid cell proliferation rate, (6) predictable and consistent osteogenic potential as well as (7) good integration into the surrounding tissues [92]. Currently, several cell-based BTE strategies have been proposed, including the use of fresh bone marrow (BM), differentiated osteoblasts and mesenchymal stem cells (MSCs) (Table 2). BM has been tested in nonunion bone defects since the 1980s [93-98], because they are easy to harvest and contain osteogenic progenitors, which can enhance bone regeneration. However, their effectiveness in older patients or patients after chemotherapy is rather limited because of the significant decrease in osteogenic progenitors [96]. Allogenic transplantations are rarely applicable, because BM contains T lymphocytes that may cause multisystem graft-versus-host syndrome [97]. The use of differentiated osteoblasts has been shown to enhance the rate and extent of bone regeneration [98-99], but has been limited by issues of expansion in vitro and donor variability [96, 99]. MSCs, also known as marrow stromal cells or colony forming unit – fibroblast (CFU-F), was first identified from adult bone marrow by Friedenstein et al in the late 1960s [100-101]. After nearly half a century of research, MSCs were found to play an essential role for bone fracture healing. They are the progenitors for osteoblasts, osteocytes and bone lining cells [102-103]; and are involved in bone remodeling via control of the opposing activities of osteoblasts and osteoclasts [104]. Furthermore, MSCs help to form bony callus by differentiating into the osteoblasts and chondroblasts during bone fracture healing and create a regenerative micro-environment by expression of a large spectrum of bioactive molecules [94, 104-105]. MSCs represent a promising

cell source for BTE for the following reasons: (1) Easy isolation by plastic adhesion method or other antibody selection techniques [106]; (2) Well-defined osteogenic differentiation pathway and MSCs have been shown to generate greater bone formation than fresh bone marrow in preclinical studies [96, 107108]; (3) Non-immunogenicity, hence suitability for allogenic application; and (4) Osteogenic potential not affected by cryostorage [109]. Drawbacks do exist and clinical applications using MSCs from adult BM is hampered by the low existing frequency, high cellular senescence, limited proliferation capacity [110-111] and osteogenic potential [112-114]. As a result, lots of effort has been directed at searching for an alternative source of MSCs, leading to the isolation of MSCs with osteogenic potential from a diverse range of tissue types and ontogeny, including postnatal including adipose [115] and perinatal such as the umbilical cord [116], umbilical cord blood [117-118], amniotic fluid [119-120], as well as prenatal such as fetal blood, bone marrow and liver [121-124]. Table 2. Comparison of different cellular sources for bone tissue engineering. Collection

Expansion

Storage

Allogenic application

Osteogenic potential

Fresh BM

+++

N.A.

N.A.

N.A.*

+

Osteoblast

+

+

+

N.A.*

+++

MSC from adult BM

++

++

+++

++

++

MSC from fetal BM

++

+++

+++

+++

+++

N.A.*: allogenic application probably can be achieved by proper HLA matching

We performed a study to systematically compare human MSCs from different ontological and anatomical origins, including fetal bone marrow, umbilical cord, adult bone marrow, and adult adipose tissue. In vitro comparative studies were done in both monolayer culture systems and on 3D PCL-TCP bioactive scaffolds to investigate their proliferative capacity, osteogenic differentiation and mineralization. In vivo studies were carried out by investigating the ectopic bone formation capacities of different MSCs

incorporated into PCL-TCP scaffolds. The ontological and anatomical origin of MSC has a profound influence on the proliferative and osteogenic capacity of MSC. Human fetal MSC (hfMSC), derived from a unique origin, were found to be a good cell source (Table 2). They express embryonic stem cell markers such as Oct-4 and Nanog, indicating their primitive nature, and had higher proliferative capacities with 30 hours’ doubling time and more than 70 population doublings before reaching cellular senescence. Moreover, they have significantly greater osteogenic potential and achieved more ectopic bone formation when compared to other MSC sources. (Unpublished data) The hfMSCs have also been shown to have lower immunogenicity than adult BM derived MSCs [125-126]. A challenging issue for the clinical translation of hfMSC incorporated PCLTCP scaffolds is the maintenance of viability and pre-differentiation prior to implantation. In vitro, in the absence of a vasculature network, the static culture cannot provide sufficient nutrient and oxygen supply for large 3D scaffolds because of the limited diffusion capacities and high demand for nutrient and gas exchange created by high cellular density. Studies showed that 3D tissue in the static culture is limited to a depth of few hundred microns [127-128]. Bioreactors have been introduced to increase mass transfer to mitigate the diffusion limitation of 3D scaffolds and provide necessary mechanical stimulation to trigger the mechanotransduction signaling pathway for better cell differentiation. 129-131 Rotating wall vessels (RWV) bioreactor, an example of bioreactors used in TE, have the characteristic of low-shear, threedimensionality and high mass transfer and thus provide dynamic flow culture conditions to promote tissue synthesis [132-136]. We recently developed a novel biaxial-rotating RWV bioreactor, which can rotate simultaneously in two orthogonal axes. Computational simulation has shown that it can achieve manifold increases in fluid velocity with significant improvements in fluid transport through the scaffolds compared to conventional uni-axial rotating bioreactor and static culture [137]. This bi-axial rotating bioreactor also provided suitable preclinical culture environment for the hfMSC incorporated PCL-TCP scaffolds by facilitating high viability, proliferation, homogenous cell and ECM distribution and osteogenic differentiation/mineralization. After just one month of culture in the bi-axial bioreactor, hfMSC incorporated PCL-TCP scaffolds were found to become bone-like tissues, with differentiated osteoblast trapped inside the highly mineralized ECM. (Unpublished data)

7 Conclusions This chapter has served to introduce the various applications of PCL-based biomaterials for tissue engineering, with emphasis on BTE and the importance of structure-property relationships. Unique biaxially-stretched PCL films with superior mechanical properties have been developed for use as layer-by-layer cell-film constructs. PCL macroporous scaffolds and nanofibrous films are being tested for drug delivery. PCL-TCP composite scaffolds with honeycomblike, fully-interconnected architecture have been proven suitable for delivery of growth factors in the repair of bone defects. These scaffolds have also been explored as carriers for several stem cell types. Hence, the biomedical applications for PCL-based biomaterials are still in its infancy and represents exciting future research.

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Chapter 11 Nanofiber-based Scaffolds for Tissue Engineering Hisatoshi Kobayashi1, Yoshiro Yokoyama1, Chiaki Yoshikawa1, Satoshi Igarashi1, Shinya Hattori1, Takako Honda1, Hiroyuki Koyama2, and Tsuyoshi Takato2 1. Biomaterials Center, National Institute for Materials Science, Tsukuba, Japan 2. Tisuss engineering Dept. The University of Tokyo Hospital, Tokyo Japan

1

Introduction

Recently nanotechnology is blooming through numerous researches and one of the important targets is to use in medical fields. Several researchers have reported novel nanoparticles and nanofibers and their possibilities in the fields of at cutting edge medical prosthesis, drug delivery, pharmaceutical and tissue engineering scaffolds, etc. Various applications of nanomaterials are well summarized in the other review papers [1, 2]. As this reasons, nano-sized materials are similar size range with various proteins, extracellular matrix (ECM), and physiological factors, that mimic surrounded natural tissues and cells. And nanomaterials such as nanoparticles and nanofibers have large relative surface area against volume. Considering that the material size and feature can substantially affect the morphology, functionality and cell-cell interactions of cells grown on ECM, cells would show good attachment and proliferation in micro and nanostructured materials. Moreover if it is degraded with cell or tissue growth, and the shape is enhanced the permeability of the nutrition and metabolites compared with other shaped scaffolds. Due to statement above, nanofibers made from bioabsorbable and biocompatible polymers are promising the application to tissue engineering. In fact, nanofiber made from biodegradable polymer such as collagen (gelatin) [3,

4], chitosan [5], poly(glycolide), (PGA) [6, 7], poly(lactide) (PLA) [8], poly (ε-caprolactone) [9] and their random copolymers [10, 11] were reported. The compatibilities with various cells such as human embryonic palatal mesenchymal cells [4], smooth muscle cells [10], endothelial cells [10], human bone marrow stromal cell [12] and neural stem cells [13] are demonstrated. When the engineered tissue implanted into the living body, the host tissue, for instance blood vessels, should be invaded into the scaffold rapidly to maintain the seeded cells and immature tissue. And these nanofiber scaffolds come in contact with living bodies such as blood or body fluid, and it is used, and therefore a boundary surface also carries an important role to decide the performances as well as the 3D-structure such as the porosities, fiber spacing, alignment, etc. To control these parameters, various processing methods are progressing. Here we developed new technique to give surface micro-structure on the nanofiber mats consist of PGA to improve. To prove the importance of 3D structure, new sponge form scaffold was developed and comparison between nanofiber sheet and nanofiber sponge was carried out in rat in vivo model. And we also developed the new nanocomposites fibers, which consisted of PGA and collagen(PGA/Col), and investigated the nanostructure of the the PGA/Col and degradation pattern of the composites. The angiogesesis activities were evaluated in vitro.

2

Nanofiber Construct Preparation by Electrospinning

In the early 20th century, the principal of the electrospinning was developed [14] and revived at the end of 20th century together with the progress of nanotechnology. In general, electrospinning is performed by the simple set-up as Disposable syringe and Syringe pump Stainless needle

Collector: Metal plate High Voltage Power Supply

a) Conventional system

Figure 1. Schematic illustration of conventional electrospinning set-up.

shown in Fig. 1. Mainly just three components, high voltage suppliers, glass syringe with small nozzle, and metal-made collector, are required. The parameters affecting elecrospininng is summarized as below. A) Polymer solution properties, for example, viscosity, conductivity, and surface tension, etc. B) Set-up related factor, for example, flow rate of the polymer solution, nozzle shape, applied electric potential, and distance between nozzle and collector, etc. C) Environment, for example, humidity, temperature, a convection of gas.

However, in some polymer case, nanofibers via conventional electrospinning method showed macroscopically two-dimensional structure and very dense structure preventing from going into inside of biomaterials by several cells. Therefore we need to modify the electrospinning process to creat 3D structure and the density control.

2.1 Micro-structured Nanofiber Mat with Template Collector System Surface topology is one of the important factors to give functionality of materials. In general, micro-replication technology is sometimes utilized to give microstructure on the surface. But in case of nanofiber mat cannot use the imprinting system after the creation of nano mat. When the hot press process would use the replication process, the outer surface of the nanofiber mat create filmy skin layer by the heating process, and lose the nanofiber structure, and then the nanofibereffect would be lost. From that point of view, we developed combination process with the micro-imprinting system in the electrospininng process [15]. The electrospinning with template collector set-up was shown in Fig. 2 a). On the conventional electrospinning method, in principal, nanofibers are strongly drawn to the collector by high voltage charge, are then sticked and compressed to the collector surface. Therefore when the template is placed as the collector part, the replication is achieved at the same time. Disposable syringe and Syringe pump

Disposable syringe and Syringe pump

Stainless needle

Stainless needle

Collector: Metal template High Voltage Power Supply

a) Template collector

Collector: Metal bath filled with liquid High Voltage Power Supply

b) Bath collector

Figure 2. Schematic illustration of modified electrospining. a) Template collector for micro-structured nanfiber mats manufacturing. b) Metal bath collector for nanofiber sponge manufacturing.

We used the micro-patterned template instead of flat collector in electrospinning equipment. “Micro-patterned nanofibrous material” demonstrated regular and same sized dents sizing 200 micrometer diameter, 37 micrometer depth and 250 micrometer space between two dents. PGA was dissolved with 1,1,1,3,3,3-hexafluoro-2-propanol. Micro-patterned templates made from acrylic resin were coated by platinum and then used as collector of electrospinning equipment. Electrospinning are performed based on previous reported condition by Boland et al [16, 17] with slight modification. Applied potential is 23kV, 67mg/mL of PGA solution was used and its flow rate was controlled at 10mL/hour. The distance between nozzle and collector was 25cm. By using micro-patterned template instead of collector, nanofibrous non-woven mats were formed on it. This material was dried overnight in a vacuum at room temperature. After drying, materials were removed from template and used for observation of Scanning electron microscopy (JSM-5600LV, JEOL, Tokyo Japan) and laser microscopy(VK-8510, KEYENCE, Osaka Japan). The distribution of fiber diameter is 100-1600 nanometer and averaged one is about 650 nanometer. As a result, nanofibrous materials formed on it and the products possesses regular micro-pattern. As shown in Figure 3, micro-patterned nanofibrous material possesses regular and same sized dents on its surface. In case of column shape of the template, the sizes of dents were 200 micrometer diameter, 37 micrometer depth and 250 micrometer distance between two dents. These sizes were almost similar to template (200 micrometer diameter, 37 micrometer depth, 250 micrometer distance). Similarly, we had also succeeded in a trapezoidalpatterned nanofibrous materials. This achievement indicated that we can control the size, shape and space by the designs of templates.

2.2 Sponge Form Preparation with Wet Bath Collector System On the electrospinning method, nanofibers are strongly drawn to the fiber collector by high voltage charge. The flying charged-fiber is landed on the collector surface, then sticked and strongly compressed to the collector surface. Therefore, it is difficult to control the density of nanofibers in the direction of the thickness and the control of 3D structure in the conventional electrospinning method. This caused very much restriction for the scaffold application. On that point of view, novel wet electrospinning system was developed as shown in Fig. 2 b). The 3D structure of nanofiber fabric was controlled by combining electrospinning with metal bath collector, filled with low surface tension liquid which is poor solvent for the fiber. In this system, the charged fibers are also strongly drawn to the collector and nanofibers are created. But the flying charged fiber is landed on the collector, liquid which has low surface tension, become a kind of cushion and the created fibers does not stick and compressed on the collector surface.

200 μ 250 μ

37 μ

200 μ

37 μ

μ

Template

200 μ 200 μ

Dents on nano fiber mat

Figure 3. SEM and laser microscopy photomicrographs of Micro-structured nanofiber mats and the template collectors.

Spongiform nanofiber 3-dimensional fabric controlled the fiber density was successfully formed with simplicity using the novel wet electrospinning system [18,19]. PGA was used as the specimen. PGA was dissolved in 1, 1, 1, 3, 3, 3 hexafluoro-2-propanol at concentration of 90 mg/ml. The positive output lead of a high voltage supply, set to 25 kV, was attached to a 21 gauge nozzle. A grounded stainless bath collector filled with 99 % t-BuOH. After fabricating nanofibers in the spinning stainless bath via wet electrospinning system, nanofiber fabric was freeze with 99 % t-BuOH in the freezer for more than 2 hours. And then nanofiber fabric was vacuum-freeze dried using a freeze dryer for 1 day in order to keep the form of fabrics. The SEM images and digital camera photographs of PGA nanofibers fabricated by conventional electrospinning system and novel wet electrospinning system with 99% t-BuOH were shown in Table 1. Each nanofibers consisted of fibers with diameters ranging from 200 nm to 1400 nm. The average diameter of PGA fiber prepared by conventional electrospinning system is 676 ± 355 nm. And the average diameter of fiber fabricated by wet electrospinning system is 636 ± 204 nm. Both nanofibers have the same average diameter and deviation.

The spongiform fabric had large spaces between the nanofibers in comparison with the non-woven mat. Comparing with usual nanofiber non-woven mat, spongiform nanofiber fabric has lower apparent density and higher porosity. When the water is used as the solvent, fibers float on the water due to its high surface tension. Therefore, fibers are collected on the plane like a nonwoven mat prepared by conventional system. t-BuOH has low surface tension. By using such solution, fibers can sink in the solvent, and have a degree of directional freedom for 3-dimension. Then spongiform fabric with the remarkable low bulk density is made in the wet spinning bath. In this study, PGA spongiform nanofiber 3-dimensional fabric was successfully formed using novel wet electrospinning system with low surface tension solvent. Spongiform nanofiber fabric had lower bulk density than nanofiber non-woven mats made by conventional electrospinning system. Table 1. Comparison between PGA nanofibers fabricated by conventional electrospinning system and novel wet electrospinning system with 99% t-BuOH.

The novel nanofiber fabrication system could be very useful for the structure control of the nanofiber fabric. And the resultant nanofiber fabrics would be useful as various applications.

3

In Vitro and In Vivo Evaluation for the PGA Nanofiber Fabrics

3.1 In Vitro Evaluation of the PGA Spongiform Nanofiber Fabric Considering that this material is applied to the vascular invasion bed, interaction of the capillary cells with the PGA nanofiber fabrics must be checked. In this study, the human umbilical vein endothelial cells (HUVEC) were used and

cultured in Endothelial Cell Growth Medium containing growth factor in 5% CO2 atmosphere at 37oC. HUVEC at passage 3-4 were seeded to materials at a density of 1.5x104 cells/cm2 for cell attachment assay. After one day culture, the samples were fixed with 4% glutaraldehyde solution and the cell attachment was evaluated by SEM observation. And the cells migrations into the scaffolds were evaluated by cross section of the samples stained with Giemsa staining. As shown in Fig. 4, HUVEC were well adhered on the surface of both PGA non-woven mat and Spongiform PGA. But in case of the PGA non-woven mat, the seeded cells just stay near the surface, not migrated into the central part of the substrate during the time period (see the cross sectional photo of PGA nonwoven mat; blue color, i.e. cells, distributed only the outer surface). This seems to be caused by too much fiber density. On the contrary, in case of spongiform PGA, HUVEC migrated into the deep inside even in one day after the seeding (see the cross sectional photo of spongiform PGA; blue color, i.e. cells, distributed not only the outer surface but also the inside of the sponge). These results suggested the importance of the 3D structure for the cell distribution. For the gene analysis, HUVEC at passage 3-4 were seeded to materials at a density of 9.54x104 cells/cm2. Samples are obtained on 3 and 7days by trypsin treatment and centrifugation. mRNA extraction and cDNA synthesis are performed. These cDNAs are amplified by PCR reaction (95oC for 5min, 94 oC for 1min–60 oC for 1min x43cycles and 60 oC for 10min) and amplified samples are applied electrophoresis gel and detected by UV lump. The results were shown in Table 2. Gene expressions of anginogenesis marker, especially Integrinαv and VEGF receptor, were clearly detected on the PGA nanofiber fabrics on 7days after seeding compared with polystylene culture dish. Table 2. Angiogenesis related gene expression of the HUVEC on PGA nanofiber fabric.

Collage TypeIα Chain

ECM

3 days

7 days

Detected

Detected

CD-44

Angiogenesis marker

Detected

Detected

ED-B domain of fibronectin

Angiogenesis marker

Detected

Detected

Integrin αv

Angiogenesis marker

Detected

Clearly detected

VEGF receptor

Angiogenesis marker

Detected

Clearly detected

GAPDH

Control

Detected

Detected

This result suggested that the material (i.e. PGA) has some potential to induce angiogenesis. In this experiment, no much difference of gene expression was observed in both non-woven and spongiform PGAs. Therefore it seems that the phenomena were caused by the degradation products such as oligo-glycolic acid or glycolic acid.

Figure 4. Human Umbilical Vein Endothelial Cells adhesion on the PGA nanofiber fabrics.

3.2 In Vivo Evaluation of the PGA Spongiform Nanofiber Fabric To prove the importance of 3D structure for tissue engineering scaffolds, comparison between nanofiber sheet and nanofiber sponge was carried out in rat in vivo model. The PGA spongiform nanofiber fabrics and nanofiber sheets were implanted in rat sub-aponeurosis glutea for 7days and histological evaluations were carried out. As shown in Figure 5, in case of the PGA non-wovon mat, no cell and no tissue ingrowth was observed during one week. On the other hand, cells and tissue was starting to migrate in the spongiform PGA nanofiber fabric with newly formed capillaries within 7days after implantation. This fact supported the in-vitro cell culture results as well as the results of gene expression assay. PGA is biodegradable polymer, and it would be degraded within 3 weeks to several months, depend on the molecular weight and the crystallinity. Within

this experiment periods, the degradation process by hydrolysis is just started, the tenacity of PGA is still kept high enough to keep the bulk shape. So, the PGA non-woven mat, made by conventional electrospining method, has too much nanofiber density agaisnt tissue migration in the scaffold in the early implantation stage. On the contrary, the spongiform PGA nanofiber fabric had a spongy structure with large spaces between the constructed nanofibers. Therefore, the host cells and tissue would be migrated easier into the scaffold, even if degradation of PGA is not enough. Even in case of the PGA non-woven mat, neovascularization were observed in outskirts of the PGA non-woven mat. This result was coincident with the results of HUVEC gene expression. This result clearly demonstrated the PGA nanofiber itself has angiogenesis activity. The cells and tissue migration into the spongiform PGA was faster than that into the PGA non-woven sheet. But the tissue migration was not finished yet even in case of spongiform PGA mats within this periods. Very few cells and tissue was observed in the central part of the sample at 7 days after implantation. These results suggested that when we need to achieve the large defect regeneration with more quick tissue ingrowth, we need to combine somehow the molecules which accelerate wound repair.

Figure 5. In vivo comparison of spongiform PGA with conventional PGA non-woven mat.

On this point of view, now we are trying to develop the PGA/Col composite nanofiber sponge and combination of the growth factor such as b-FGF and VEGF. As fundamental research for the PGA/Col composite nanofiber, we developed the new quantitative evaluation methods for cell-nanofiber interaction [20]. From the results of the analysis, we have found out optimum ratio of the PGA and collagen mixture (4:6) and cell-fiber interaction affected by the fiber diameters and strongly enhanced in the nanofibers compared with microfibers. From the TEM observation of the PGA/Col nanocomposite fiber, phase separation of the PGA and Collagen was observed. And the PGA/Col composite nanofibers were degraded and deformed within 5days in PBS(-) at 37oC. Based on the results, preliminary implantation of the spongiform PGA/Col composite fabric was carried out. The complete set of the experiment has not finished yet. But in some case, inside of the spongiform PGA/Col composite nanofiber was occupied the newly formed tissue with neovascular just within 5days after implantation. This result suggested the angiogenesis and the tissue migration activities of PGA/Col nano-composites and spongiform helps the tissue ingrowth into the nanofibers construct made by electrospinning. Further studies, such as the detailed characterization of the nanostructure, degradation profiles, cellrecognition of the nanoscale structure, etc., are still necessary.

4

Conclusions

When the nanofiber based scaffold made from bioabsorbable and biocompatible polymers would use tissue engineering application, the 3D-structure such as the porosities, fiber spacing, alignment, etc. is one of the most important factor to get good results. And vasculogenesis acceleration produced by scaffold materials, is another important factor to achieve the large area tissue regeneration. In the series of this study, we developed new technique to construct 3D structure by simple set-up of the electrospinning method. And we demonstrated that PGA and PGA/Col have some potential to accelerate neovascularization of the implanted materials. Combination of the spongiform nanofiber with the PGA/Col had a great potential to achieve the quick repair of the tissue defect and to generate higher functional tissue and organs.

Acknowledgments This work was supported by the Ministry of Education, Culture, Sports, Science and Technology and the Ministry of Health, Labour and Welfare of Japan. The animal study was supported

References 1. Huang Z.-M., Zhang Y.-Z., Kotaki M., Ramakrishna S., 2003. A review on polymer nanofibers by electrospinning and their applications in nanocomposites, Composites Science and Technology. 63, 2223-2253. 2. Ashammakhi N, Ndreu A., Piras A., Nikkola L., Sindelelar T., Ylikauppila H., Harlin A., Chiellini E., Hasirci V., Redl H., 2006. Biodegradable Nanomats Produced by Electrospinning: Expanding Multifunctionality and Potential for Tissue Engineeering, J. Nanosci. Nanotechnol. 5, 2693-2711. 3. Matthews J.A., Wnek G.E., Simpson D.G., Bowlin G.L, 2002. Electrospinning of collagen nanofibers, Biomacromolecules. 3, 232-238. 4. Li M., Mondrinos M.J., Gandhi M.R., Ko F.K., Weiss A.S., Lelkes P.I., 2005. Electrospun protein fibers as matrices for tissue engieneering, Biomaterials. 26, 5999-6008. 5. Ohkawa K., Cha D., Kim H., Nishida A., Yamamoto H., 2004. Electrospinning of chitosan, Macromol. Rapid Commun. 25, 1600-1605. 6. Boland E.D., Wnek G.E., Simpson D.G., Pawlowski K.J., Bowlin G.L., 2001. Tailoring tissue engineering scaffolds using electrostatic processing techniques: a study of poly (glycolic acid) electrospinning, J. Macromol. Sci Pure. Appl. Chem. A38, 1231-1243. 7. Boland E.D., Telemeco T.A., Simpson D.G., Wnek G.E., Bowlin G.L., 2004. Utilizing acid pretreatment and electrospinning to improve biocompatibility of poly (glycolic acid) for tissue engineering, J. Biomed. Mater. Res. Part B. Appl Biomater. 71B, 144-152. 8. Zong X., Kim K., Fang D., Ran S., Hsiao B.S., Chu B., 2002. Structure and process relationship of electrospun bioabsorbable nanofiber membranes, Polymer. 43, 4403-4412. 9. Reneker D. H., Kataphinan W., Theron A., Zussman E., Yarin A.L., 2002. Nanofiber garlands of polycaprolactone by electrospinning, Polymer. 43, 6785-6794. 10. Li W.J., Laurencin C.T., Caterson E.J., Tuan R.S., Ko F.K., 2002. Electrospun nanofibrous structure: A novel scaffold for tissue engineering, J. Biomed. Mater. Res. 60, 613-621. 11. Mo X.M., Xu C.Y., Kotaki M., Ramakrishna S., 2004. Electrospun P(LLA-CL) nanofiber: a biomimetic extracellular matrix for smooth muscle cell and endothelial cell proliferation, Biomaterials. 25, 1883-1890. 12. Jin H.J., Chen J., Karageorgiou V., Altman G. H., Kaplan D.L., 2004. Human bone marrow stromal cell responses on electrospun silk fibroin mat, Biomaterials. 25, 1039-1047. 13. Yang F., Murugan R., Wang S., Ramakrishna S., 2005. Electrospinning of nano/micro scale poly (L-lactic acid) aligned fibers and their potential in neural tissue engineering, Biomaterials. 26, 2603-2610.

14. Formhals A., 1934. Process and apparatus for preparating artificial threads, US Patent 1975504. 15. Igarashi S., Tanaka J., Kobayashi H., 2007. Micro-PatternecNanofibrous Biomaterials, J. Nanosci. Nanotechnol. 7, 814-817. 16. Ohkawa K., Cha D., Kim H., Nishida A., Yamamoto H., 2004, Electrospinning of chitosan, Macromol. Rapid Commun. 25, 1600-1605. 17. Boland E.D., Wnek G.E., Simpson D.G., Pawlowski K.J., Bowlin G.L., 2001. Tailoring tissue engineering scaffolds using electrostatic processing techniques: a study of poly (glycolic acid) electrospinning, J. Macromol. Sci Pure. Appl. Chem. A38, 1231-1243. 18. Yokoyama Y., Hattori S., Yoshikawa C., Yasuda Y., Koyama H., Takato T., Kobayashi H., submitted. Novel wet electrospinning system for fabrication of spongiform nanofiber 3-dimensional fabric, materials letter. 19. Kobayashi H., Yokoyama Y., Takato T., Koyama H., Ichioka S., 2007 submitted. Japanese patent 2007-103201. 20. Tian F., Hosseinkhani H., Hosseinkhani M., Khademhosseini A., Yokoyama Y., Estrada G. G., Kobayashi H., 2008. Quantitative analysis of cell adhesion on aligned micro- and nanofibers, J. Biomed. Mater. Res. Part A. 84A, 291-299.

Chapter 12 Electrospun Composite Nanofibrous Scaffolds for Tissue Engineering

Inn-Kyu Kang1 and Jung Chul Kim2 1. Department of Polymer Science and Engineering, Kyungpook National University, Daegu 702-701, South Korea 2. Department of Immunology, School of Medicine, Kyungpook National University, Daegu 700-422, South Korea

Abstract Electrospinning has recently emerged as a leading technique for the formation of nanofibrous structures made of synthetic and natural extracellular matrix component. In this study, nanofibrous scaffolds were obtained by coelectrospinning poly(3-hydroxybutyrate-co-3-hydroxyvalerate) (PHBV) and collagen Type І in hexafluoro-isopropanol (HIFP). The resulting fiber diameters were in the range between 300 and 600 nm. The nanofiber surfaces were characterized by attenuated total reflection Fourier transformed infrared spectroscopy, electron spectroscopy for chemical analysis, and atomic force microscopy. The PHBV and collagen components of the PHBV-Col nanofibrous scaffold were biodegraded by PHB depolymerase and a collagenase Type I aqueous solution, respectively. It was found, from the cell-culture experiments, that the PHBV-Col nanofibrous scaffold accelerated the adhesion of the NIH 3T3 cell compared to the PHBV nanofibrous scaffold, thus showing a good tissue engineering scaffold.

1

Introduction

Collagen is a natural ECM component of many tissues, such as skin, bone, tendon, ligament, and other connective tissues. Therefore, it has a more native surface, relative to synthetic polymers, which favors cellular attachment as well as being chemotactic to cells. It is well known that collagen plays an essential role in providing a scaffold for cellular support and thereby affecting cell attachment, migration, proliferation, differentiation and survival. On the other hand, poly(3-hydroxybutyric acid-co-3-hydroxyvaleric acid) (PHBV) is a wellknown biodegradable, biocompatible, non-toxic, and thermoplastic polyester that is produced by bacteria [1]. PHBV emerges as a new generation of PHBbased materials with properties that are adjustable via changing the hydroxyvaleric acid (HV) content [2, 3]. A highly porous structure is desirable to allow for cell seeding or migration. Throughout the material, pore size plays a critical role in both cell adhesion and in the exchange of nutrient and metabolic waste. In biomedical applications, for example, the electrospinning technique can be used to construct wound dressings [4], drug delivery platform [5], and tissue engineering scaffolds [6,7]. The nanofibrous collagens easily lose their fibrous form in an aqueous medium. Accordingly, they need to be crosslinked for keeping a proper mechanical stability. [8,9]. In this study, the blended polymer solution of PHBV and collagen Type І was used to fabricate an electrospun nanofibrous scaffold. Characteristics of the nanofibrous scaffolds were investigated using ATR-FTIR spectroscopy and electron spectroscopy for chemical analysis. The behavior of the fibroblasts on the nanofibrous scaffolds was also investigated.

2

Materials and Methods

2.1 Electrospinning & Surface Characterization Hexafluoro-isopropanol (HFIP) was chosen to dissolve both PHBV and collagen, based on the previous reports [10]. Transparent polymer solution for electrospinning was obtained by dissolving both collagen and PHBV in HFIP with sufficient stirring under room temperature. To examine the effect of collagen content on fiber morphology, 2 wt% of polymer solution was prepared using the different ratios of PHBV and collagen. The presence of collagen in the nanofibrous scaffolds was confirmed using a FT-IR spectrometer. The surface

chemical composition was analyzed using electron spectroscopy for chemical analysis.

2.2 Biodegradation & Cell Behaviors In order to examine the enzymatic degradation of nanofibrous scaffolds, samples were placed in a PBS containing collagenase type І (10 mg/ml). The proliferation of NIH 3T3 cells seeded on the nanofibrous scaffolds were determined using a colorimetric immunoassay based on the measurements of 5bromo-2'-deoxyuridine (BrdU), that was incorporated during DNA synthesis [11]. The production of extracellular materials (ECM) in cultured matrices was assayed by immunohistochemical staining with α-SMA and cytokeratin 8.

2.3 Open Wound Healing Test Athymic nude mice were purchased from the Oriental Company (Busan Korea) and used at 6-7 weekd old. Three mice were used for each of four experimental groups (PHBV and PHBV-Col) and one untreated control group, and at least six sets of total experiments were carried out. Mice were anesthetized with Forane solution, and full thickness skin wounds of 1 cm diameter were prepared. Then matrices cultured for 3-5days were grafted, cell layer down, onto the wound surfaces. Matrices were changed every 3-4 days after grafting. When changing matrices, macroscopic photographs were obtained, and images of the decreasing wound areas were copied into transparent films. The areas on the copied films were calculated and normalized using the Image J Program.

3

Results

3.1 Morphology of Electrospun Nanofibrous Scaffolds We have explored the effect of collagen content on fiber morphology using polymer blends with different ratio of PHBV and collagen (7:3, 5:5, 3:7). When the ratio of collagen is high, the as-spun fibers could not keep their morphology because the collagen is readily swellable or partially soluble in water. In addition, the nanofiber containing high amount of collagen has poor mechanical property. Based on these preliminary results, we thought the most proper ratio of PHBV and collagen is 7:3 and used this ratio for the preparation of

electrospun nanofibrous scaffolds. Moreover, we have electrospun by changing the fluid and processing parameters. Figure 1 shows the SEM images of electrospun nanofibrous scaffolds that were obtained under maximized condition. The images showed continuous fiber morphology and the fibers did not contain beads that were independent of the kind of polymer. The PHBV fiber diameters were in the range between 300 and 600 nm and it decreased after the incorporation of collagen (PHBV-Col).

Figure 1. SEM micrographs of nanofibrous scaffolds. (a) PHBV, (b) PHBV-Col, (c) collagen

3.2 Characterization of Nanofibrous Scaffolds Figure 2 shows the ATR-FTIR spectra of the collagen (a) and PHBV-Col (b) nanofibrous scaffolds. Sachlos et al. [12] examined the effect of processing on the tertiary structure of collagen by monitering the wavenumber of the N-H strctching vibration peak. They reported that the position of the N-H peak was almost not shifted by ethanol treatment and critical point drying while significantly shifted with increasing the temperature over 40oC. As shown in Figure 2, the N-H peak of collagen nanofibrous scaffold (a) was almost same as that of PHBV-Col nanofibrous scaffold. In addition, the amide I (1651cm-1) and amide II (1540 cm-1) of collagen nanofibrous scaffold (a) were also same as those of PHBV-Col (b). These results indicate that the collagen type I preserve its conformation after blending with PHBV. Changes in the chemical structure of nanofibrous scaffolds were investigated using ESCA. Figure 3 shows the ESCA survey scans of the nanofibrous scaffold surfaces. The collagen nanofiber scaffold (Figure 3 (c)) showed three peaks corresponding to C1s (binding energy, 285eV), N1s (binding energy, 400eV) and O1s (binding energy, 532eV), while the PHBV nanofiber scaffold (Figure 3 (a)) showed two peaks corresponding to C1s and O1s.

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3.3 In Vitro Biodegradation Figure 4 illustrates the morphological changes of the nanofibrous scaffold surfaces before (a, b, c) and after (d, e, f ) incubation with a Type I collagenase aqueous solution. For the results, the surface morphology of the PHBV (a) and PHBV-Col (b) nanofiber scaffolds did not change after they were dipped in a PBS solution for 12h. The surface of the collagen nanofiber (Figure 4, c), however, lost its original fiber morphology, probably due to swelling caused by water. On the other hand, the collagen nanofiber scaffold was partially biodegraded by the treatment of a Type I collagenase aqueous solution, as shown in Figure 4 (f ).

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Figure 4 SEM images of nanofibrous scaffolds incubated for 12h in PBS solution (a,b,c) and in collagenase Type І solution (d,e,f): a, d = PHBV; b, e = PHBV-Col; c,f = collagen.

3.4 Cell-scaffold Interaction Collagen has an Arg-Gly-Asp sequence that can be recognized by one of the receptors of cell membrane. Figure 5 shows the SEM images of the NIH 3T3 fibroblasts that adhered to the nanofibrous scaffolds when cultured in a Dulbecco’s modified eagle medium containing 10% fetal bovine serum for 4 h. The cells are more adhered to the PHBV-Col than to the PHBV scaffold. In the case of the collagen nanofibrous scaffold, the cells spread largely and formed clusters with a stratified appearance to the cell layers (Figure 5 (c)). The result of Figure 5 showed that the order of cell adhesion is collagen > PHBV-collagen > PHBV. This result suggests that the introduction of collagen could improve the adhesion of fibroblasts. Figure 6 shows the proliferation of cells on the nanofibrous scaffolds when cultured in a DMEM with 10% serum for 68 h. Cell proliferation on the PHBV nanofibrous scaffold was significantly accelerated by the introduction of collagen Type I (PHBV-Col) (P < 0.01). It is considered that the high cell proliferation on the collagen and collagen-containing PHBV may probably be due to the high cell adhesion on the same scaffold (Figure 5). Figure 7 shows the biological expression of cells cultured on the PHBV matrices. MAbs of human cytokeratin 8 and α-SMA were colocalized to identify extracellular materials (ECM) production in matrices cocultured for 5 days. Stronger signals of α-SMA fibers were produced in PHBV-Col than in PHBV, while both matrices showed similar expression of cytokeratin 8 on round cytoplasm of ORS cells. This result was correlated with faster cell attachment and growth of cells in PHBV-Col over PHBV (Figs. 5 and 6). Wound healing effects of cocultured matrices are shown in Figure 8. Until the second day after grafting with matrices, all experimental sets showed similar healing progress. However, on the fourth day, PHBV-Col set produced better wound closure than PHBV and the control set did.

4

Discussion

Unlike a conventional fiber fabrication process, electrospinning provides a straightforward way to fabricate fibrous scaffolds with fiber diameters in the tens of nanometers [13]. In the present study, we have studied cell behavior on nanofibrous scaffolds which was brought about by electrospinning using biodegradable PHBV and Type Ι collagen. Wagner et al. [14] have fabricated

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Figure 8. Wound healing progress of PHBV matrices cocultured with ORSand DS cells or matrices alone.

elecrospun polyetherurethaneurea (PEUU)/collagen scaffolds by combining PEUU with Type I collagen at various ratios. In their results, the SEM images revealed continuous fiber morphology at all of the examined ratios. The diameter of co-electrospun fibers was in the range between 100 and 900 nm. The results, from the cell experiments, showed that the cells adhered well to the fibers with diameter smaller than the size of the cells [15]. These findings has led to a necessity of mimicking nanoscale structures that are found in natural ECM when creating a polymeric synthetic ECM scaffold for tissue engineering [16]. In order to illustrate the presence of collagen in the co-electrospun scaffolds, a PHBV-Col scaffold was subjected to ATR-FTIR measurements. The resulting spectrum yielded peaks that were characteristic of collagen, as reported previously by other researchers [17]. Figure 2(a) illustrated common bands of collagen appeared at approximately 1651 cm-1 (amide I) and 1540 cm-1 (amide II), corresponding to the stretching vibration of C=O bond and the bending vibration of N–H bond, respectively. When blended and electrospun PHBVwith

collagen, the peak of amide I and N-H stretching vibration did not shift. This result indicates that the conformation of native protein is preserved in the PHBV-Col scaffold. The enzymatic degradation of PHBV films by PHB depolymerase has been reported previously [17]. Park et al. [18] have carried out biodegradation tests of PHBV in the form of fibrous structures or film, by using a simulated municipal solid waste aerobic composting method. They reported that the degradation of the PHBV non-woven structures was faster than the PHBV film. In this study, the surfaces of the PHBV and PHBV-Col nanofiber mats were severely eroded by the PHB depolymerase treatment (data not shown). On the other hand, the PHBV-Col nanofibers were nearly not biodegraded by the treatment of collagenase solution, and their cylindrical morphology was preserved. We have redone this experiment using high collagenase concentration of 50 mg/ml for 24 h, but the morphology of PHBVCol did not change. It should be notice that after 24h degradation, the collagen fiber nearly biodegraded. These results indicate that PHBV was not sensitive to collagenase at all. Recently, a great deal of research has focused on the influence of scaffold microarchitecture on cell behavior [13, 19, 20]. In this research, the PHBV-Col nanofibrous scaffold accelerated the adhesion of NIH 3T3 cells as compared with the PHBV nanofibrous scaffold. This is probably because nanofiber can provide microarchitecture for cell adhesion and collagen can interact with the cells. From the biological standpoint, almost all of the tissue and organs are deposited in nanofibrous structures. Highly porous nanofiber mat with high surface area offers a bio-mimicking structure during the process of tissue regeneration, more structural space for accommodation and attachment of cells, and enables the efficient exchange of nutrient and metabolic waste. Collagen has an Arg-Gly-Asp sequence that can be recognized by one of the receptors of cell membranes. As shown in Figure 5, the cells are more adhered to the PHBVCol than to the PHBV. It is considered that high number of cells adhered to the collagen and PHBV-Col is probably due to the interaction of a cell membrane receptor with Arg-Gly-Asp sequence of collagen. Pores in a tissue-engineered scaffold make up the space in which cells reside. Pore properties such as porosity, dimension and volume are parameters directly related to the success of a scaffold. High porosity provides more structural space for cell accommodation and makes the exchange of nutrient and metabolic waste between a scaffold and environment more efficient. These characteristics are fundamental criteria for a successful tissue-engineered scaffold. Nanofibrous scaffolds are very porous, but the pores formed are much

smaller than the normal cell size of a few to tens of micrometers, which would inhibit cell migration. As such, the phenomenon of cell ingrowth into a nanofiber structure has been doubted by many researchers and which leading to some reasonable explanations. Pores in an electrospun structure are formed by differently oriented fibers lying loosely upon each other. When cells perform amoeboid movement to migrate through the pores, they can push the surrounding fibers aside to expand the hole as the small fibers offer little resistance to cell movement. This dynamic architecture provides cells with an opportunity to optimally adjust the pore diameter and grow into the scaffold even though some pores are relatively small. Therefore, some pores with smaller diameters in this structure may not hinder cell migration. Such a hypothesis of dynamic cell–scaffold interaction needs to be further investigated. In this study, nanofibrous scaffolds were obtained by co-electrospinning poly(3-hydroxybutyrate-co-3-hydroxyvalerate) (PHBV) and collagen Type І in hexafluoro-isopropanol (HIFP). The resulting fiber diameters were in the range between 300 and 600 nm. We have investigated a biological wound dressing that improves early-stage wound healing and a technique that reduces the time between preparation and patient use. To achieve efficient biodressings that contain proliferative cells, we cocultured ORS and DS cells on PHBV-based nanofiber matrices of varying hydrophilicities. We found that cocultured PHBV-Col had the most positive effect on wound closure.

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Hocking P.J., Marchessault R.T., 1994. Chemistry and technology of biodegradable polymers, New York: Blackie Academic & Professional 48-96. Kim B.S., Lee S.C., Lee S.Y., Chang H.N., Chang Y.K., Woo S.I., 1994. Production of poly(3-hydroxybutyric acid) by fed-batch culture of Alcaligenes eutrophus with glucose concentration control, Biotechnol Bioeng. 43, 892-898. Juni K., Nakano M., 1987. Poly(hydroxyl acids) in drug delivery, CRC Crit. Rev. Ther. Drug Carrier Syst. 3, 209-232. Khil M.S., Cha D.I., Kim H.Y., Kim I.S., Bhattarai N., 2003. Electrospun nanofibrous polyurethane membrane as wound dressing, J. Biomed. Mater. Res. Appl. Biomater 67, 675-679. Zeng J., Xu X., Chen X., Liang Q.Z., Bian X.C., Yang L.X., Jing X.B., 2003. Biodegradable electrospun fibers for drug delivery, J. Controlled Release 92, 227231. Jin H.J., Fridrikh S.V., Rutledge G.C., Kaplan D.L.,2002. Electrospinning Bombyx mori silk with poly(ethylene oxide), Biomacromolecules 3, 1233-1239.

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Fertala A., Han W.B., Ko F.K., 2001. Mapping critical sites in collagen П for rational design of gene-engineered proteins for cell-supporting materials, J. Biomed. Mater. Res., 57, 48-58. Matthews J.A., Wnek G.E., Simpson D.G., Bowlin G.L., 2002. Electrospinning of collagen nanofibers, Biomacromolecules 3, 232-238. Buttafoco L., Kolkman N.G., Engbers-Buijtenhuijs P., Poot A.A., Dijkstra P.J., Vermes I., Feijen J., 2006. Electrospinning of collagen and elastin for tissue engineering applications, Biomaterials 27, 724-734. Ito Y., Hasuda H., Kamitakahra M., Ohtsuki C., Tanihara M., Kang I.K., Kwon O.H., 2005. A Composite of Hydroxyapatite with Electrospun Biodegradable Nanofibers as a Tissue Engineering Material, J. Biosci. Bioeng. 100, 43-49. Maghni K., Nicolescu O.M., Martin J.G., 1999. Suitability of cell metabolic colorimetric assays for assessment of CD+ T cell proliferation: comparison to 5bromo-2-deoxyuridine (BrdU) ELISA, J Immunol. Methods 223, 185-194. Sachlos E., Reis N., Ainsley C., Derby B., Czernuszka J.T., 2003. Novel collagen scaffolds with predefined internal morphology made by solid freeform fabrication, Biomaterials 24, 1487-1497. Zhang Y.Z., Ouyang H., Lim C.T., Ramakrishna S., Huang Z.M., 2005. Electrospinning of gelatin fibers and gelatin/PCL composite fibrous scaffolds, J. Biomed. Mater. Res. Part B Appl. Biomater. 72, 156-165. Stankus J.J., Guan J.J., Wagner W.R., 2004. Fabrication of biodegradable elastomeric scaffolds with sub-micron morphologies, 2004. J. Biomed. Mater. Res. Part A 70, 603-614. Xu C.Y., Inai R., Kotaki M., Ramakrishna S., 2004. Aligned biodegradable nanofibrous structure: a potential scaffold for blood vessel engineering, Biomaterials 25, 877-886. Camacho N.P., West P., Torzilli P.A., Mendelsohn R., 2001. FTIR spectroscopic imaging of collagen and proteoglycan in bovine cartilage, Biopolymers 62, 1-8. Kang I.K., Choi S.H., Shin D.S., Yoon S.C., 2001. Surface modification of polyhydroxyalkanoate films and their interaction with human fibroblasts, Int. J. Biol. Macromol. 28, 205-212. Choi J.S., Lee S.W., Jeong L., Bae S.H., Min B.C., Youk J.H., Park W.H., 2004. Effect of organosoluble salts on the nanofibrous structure of electrospun poly(3hydroxybutyrate-co-3-hydroxyvalerate), Int. J. Biol. Macromol. 34, 249-256. Li M.Y., Mondrinos M.J., Gandhi M.R., Ko F.K., Weiss A.S., Lelkes P.I., 2005. Electrospun protein fibers as matrices for tissue engineering, Biomaterials 26,59996008. Riboldi S.A., Sampaolesi M., Neuenschwander P., Cossu G., Mantero S., 2005. Electrospun degradable polyesterurethane membranes: potential scaffolds for skeletal muscle tissue engineering, Biomaterials 26, 4606-4615.

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Chapter 13 Synthetic/Natural Hybrid Scaffold for Cartilage and Disc Regenerations Gilson Khang1, Soon Hee Kim1, John M. Rhee1, Munirah Sha’Ban2, and Ruszymah Bt Hj Idrus2 1. Department of Polymer Nano Science and Technology, Chonbuk National University, 664-14, Dukjin Dong 1 Ga, Dukjin Ku, Jeonju 561-756, Korea 2. Department of Physiology, Faculty of Medicine, University Kebangsaan Malaysia, Jalan Raja Muda Abdul Aziz, 50300 Kuala Lumpur, Malaysia

1

Introduction

It has been recognized that tissue engineering offers an alternative technique to whole organ and tissue transplantation for diseased, failed or malfunctioned organs. Millions of patients have been suffered by end-stage organ failure or tissue loss annually. In the United State alone, at least eight million surgical operations had been carried out each year, requiring a total national health care cost exceeding $400 billion annually [1-4]. In order to avoid the shortage of donor organ and these problems, a new hybridized method combined with cell and biomaterials has been introduced as tissue engineering very recently. To reconstruct a new tissue by tissue engineering, triad components such as (i) cells which are harvested and dissociated from the donor tissue including nerve, liver, pancreas, cartilage and bone as well as embryonic stem, adult stem or precursor cell, (ii) biomaterials as scaffold substrates which cells are attached and cultured resulting in the implantation at the desired site of the functioning tissue and (iii) growth factors which are promoting and/or preventing cell adhesion, proliferation, migration and differentiation by up-regulating or down-regulating the synthesis of protein, growth factors and receptors must be needed.

2

Biomaterials for Tissue Engineering

2.1 Importance of Scaffold Matrices in Tissue Engineering Scaffolds might be played a very critical role in tissue engineering. The function of scaffolds is to direct the growth of cells seeded within the porous structure of the scaffold or of cells migrating from surrounding tissue. The majority of mammalian cell types are anchorage-dependent resulting in dying if an adhesion substrate is not provided. Scaffold matrices can be used to achieve cell delivery with high loading and efficiency to specific sites. Therefore, the scaffold must provide a suitable substrate for cell attachment, cell proliferation, differentiated function and cell migration. The prerequisite physicochemical properties of scaffolds are (i) to support and deliver for cells, (ii) to induce, differentiate and conduit tissue growth, (iii) to target cell-adhesion substrate, (iv) to stimulate cellular response, (v) wound healing barrier, (vi) biocompatible and biodegradable, (vii) relatively easy processability and malleability into desired shapes, (viii) highly porous with large surface/volume, (ix) mechanical strength and dimensional stability, (x) sterilizability, and so on.

2.2 Natural Polymers Many naturally occurring scaffolds can be observed as biomaterials for tissue engineering purposes. One of the typical examples is the extracellular matrix (ECM) that is very complex biomaterials and controls cell function. For the ECM of tissue engineering, natural and synthetic scaffolds are designed to mimic specific function. The natural polymers are alginate, proteins, collagens (gelatin), fibrins, albumin, gluten, elastin, fibroin, hyarulonic acid, cellulose, starch, chitosan (chitin), sclerolucan, elsinan, pectin (pectinic acid), galactan, curdlan, gellan, levan, emulsan, dextran, pullulan, heparin, silk, chondroitin 6sulfate, small intestine submucosa (SIS), acellular dermis, polyhydroxyalkanoates, and so on. Much of the interest in these natural polymers comes from their biocompatibility, relatively abundance and commercial availability, and ease of processing [5].

2.3 Synthetic Polymers and Poly(α-hydroxy ester)s One of the most significant shortages of natural polymers is typical expensive, suffering from batch-to-batch variation, and the possibility of cross contamination from unknown virus or unwanted disease due to the isolation

from plant, animal, and human tissue. On the contrary, synthetic polymeric biomaterials might be easily controlled physicochemical properties and quality and no immunogenecity. Also, it can be processed with various techniques and supplied consistently in large quantity. In order to adjust the physical and mechanical properties of tissue engineered scaffold at desired place in the human body, the molecular structure, molecular weight and so on are easily adjusted during the synthetic process. There are largely divided two categories such as (i) biodegradable and (ii) nonbiodegradable. Some nondegradable polymers include polyvinylalcohol, poly(hydroxylethylmethacryalte) (PHEMA), and poly(N-isopropylacryamide). Some synthetic degradable polymers are the family of poly(α-hydroxy ester)s such as polyglycolide (PGA), polylactide (PLA) and its copolymer poly(lactide-co-glycolide) (PLGA), polyphosphazene, polyanhydride, poly(propylene fumarate), polycyanoacrylate, polycaprolactone, polydioxanone, biodegradable polyurethanes and so on [1, 5]. The family of poly(α-hydroxy acid)s such as PGA, PLA and its copolymer PLGA that are among the few synthetic polymers approved for human clinical use by US Food and Drug Administration (FDA) are extensively used or tested for the scaffold materials as a bioerodible material. It has been used for three decades as suture of PGA, bone plate, screw and reinforced materials for PLA, and drug delivery devices of PLGA in surgical operation and whose safety has been proved in many medical applications. The synthetic methods and physicochemical properties such as melting temperature, glass transition temperature, tensile strength, Young’s modulus, and elongation were reviewed elsewhere.

2.4 Bioceramic Scaffolds Bioceramic is a term introduced for biomaterials that are produced by sintering or melting inorganic raw materials to create an amorphous or a crystalline solid body that can be used as an implant. Porous final products have been mainly used scaffolds. The components of ceramics are calcium, silica, phosphorous, magnesium, potassium, and sodium. Bioceramic used in the fabrication for the tissue engineering might be classified as nonresorbable (relatively inert), bioactive or surface active (semi-inert), and biodegradable or resorbable (noninert). Alumina, zirconia, silicone nitride and carbons are inert bioceramics. Certain glass ceramics are dense hydroxyapatites (HA, 9CaO·Ca(OH)2·3P2O5) semi-inert (bioactive), and calcium phosphates, aluminum-calcium-phosphates, coralline, tricalcium phosphates (3CaO·P2O5), zinc-calcium-phosphorous

oxides, zinc-sulfate-calcium-phosphates, ferric-calcium-phosphorous-oxides and calcium aluminates are resorbable ceramics. Among of these bioceramics, synthetic apatite and calcium phosphate minerals, coral-derived apatite, bioactive glass and demineralized bone particle (DBP) since they are widely used in hard tissue engineering area [2].

3

Natural/synthetic Hybrid Scaffolds for Tissue Engineering

3.1 PLGA Hybrid Scaffolds Although a poly(α-hydroxyacid) family have been extensively tested as scaffolding materials for tissue engineering due to relatively good mechanical properties, low toxicity and predictable biodegradation kinetics, its poor mechanical strength, small pore size, releasing of acidic degradables, induction of thick fibrotic capsules and hydrophobic surface properties have limited its usage. In order to solve these problems, several techniques as a surface treatment, an introduction of bioactive molecules, a development of manufacturing method for porous structure, a hybrid with bioactive materials and so on have been developed. Khang et al. reviews four categories on the focus of hybrid, composite and complex biomaterials for the application of hybrid scaffolds such as (i) poly(α-hydroxyester) family with natural polymer and bioceramics, (ii) bioceramic scaffolds with other biomaterials, (iii) natural polymer with other biomaterials, and (iv) miscellaneous in order to approach to a more natural three dimensional environment and support biological signals for tissue growth and reorganization [5]. In our previous works, natural/synthetic hybrid scaffolds as DBP/PLGA[6], SIS/PLGA[7], hyaruronic acid/PLGA[8], fibrin/PLGA[9-11] and so on have been developed for the application of the regeneration of bone, spinal cord, cartilage and spinal cord injury. Among of these works, fibrin/PLGA hybrid scaffolds are mainly discussed on this chapter.

3.2 Fibrin/PLGA Hybrid Scaffolds for Cartilage Regeneration In Vivo and In Vitro Articular cartilage has a limited potential to repair, consequently, damaged articular cartilage will further degenerate and eventually turn into osteoarthritis.

Autologous chondrocytes implantation was first published by Brittberg et al [12]. Recently, articular cartilage repair has been given much intention in orthopaedic tissue engineering. Usually scaffolds are designed as a highly porous three-dimensional (3D) structure to allow cells to accommodate and grow inside, as well as to organize cells into a 3D tissue. Many trials have successfully cultured articular chondrocytes [13], formed neocartilage tissue [14] and transplanted autologous neocartilage to the defect, so biocompatible scaffolds that afford the proliferation of cartilage and accumulate the matrix have been widely investigated [15]. As we discussed earlier, numerous attempts have been made for successful tissue reconstruction using PLGA-based scaffold either by PLGA itself or in combination with natural polymers such as collagen [16], and ECM scaffolds, i.e. SIS [7, 17] as well as DBP [18] in our previous studies. The incorporation of bioactive molecules in PLGA is believed to mediate cells behavior, e.g. proliferation, differentiation and function. To minimize cells lost during in vitro seeding procedure, we used fibrin to immobilize cells and to provide homogenous cells distribution in PLGA scaffolds. Fibrin has been widely used for cartilage reconstruction purposes. [9~11] We hypothesized that fibrin would be an ideal cell carrier/transplantation matrix and to enhance in vitro chondrogenesis of rabbit articular chondrocytes by mean of morphological, histological, biochemical and phenotypically similar to the normal hyaline cartilage. Articular cartilage was aseptically isolated from the femoral condyles and patellae of 6 weeks-old New Zealand White rabbits and isolated chondrocytes were cultured in the mixture of equal volume of F12/DMEM. PLGA (mole ratio 50:50, molecular weight 33,000 g/mole, Resomer RG 503 H) was purchased from Boehringer Ingelheim Pharma GmbH (Ingelheim, Germany). Microporous 3D PLGA scaffolds (0.2% w/v) were fabricated by the solvent casting/salt leaching technique using sodium chloride as porogen. Each sample was assigned into two experimental groups - cultured chondrocytes were seeded into (1) PLGA scaffolds with fibrin (fibrin/PLGA) and (2) PLGA scaffolds without fibrin (PLGA). One million cells per scaffold were incorporated and resuspended with the (1) commercially available fibrin glue kit from Greenplast® (Green Cross P. D. Co., Yongin, Korea) and (2) culture medium. PLGA scaffolds were pre-treated with calcium chloride solution for PLGA/fibrin group prior to cells seeding. The ‘chondrocytes-fibrin’ admixture was seeded into PLGA scaffolds and was allowed to polymerize within 5 min. Chondrocytes suspension in culture medium was seeded directly into PLGA scaffolds and incubated for 5 min before adding culture medium. All constructs

were implanted at the dorsum of athymic nude mice. Resulted in vivo constructs were harvested at 1, 2 and 4 weeks post-implantation. The morphology of cells and distribution of cartilaginous ECM in the fibrin/PLGA and PLGA were examined via histological staining Figure 1. Before implantation, H&E staining showed that cells and ECM filled the space of the fibrin/PLGA. A significant number of round morphological chondrocytes cluster and cartilaginous ECM formation were observed in the fibrin/PLGA hybrid construct than in the PLGA construct. The in vitro fibrin/PLGA constructs exhibited superior histo-architectural characteristics of cartilage-like tissue compared to control. The closely-packed cartilage-isolated cells were homogeneously distributed in the ECM and exhibited rounded morphology with lacunae embedded in basophilic ground substance. The pericellular and interterritorial matrix was strongly stained by the characteristic red of Safranin O, to indicate the presence of proteoglycan-rich matrix in the construct. Cartilaginous ECM deposition was further demonstrated by positive Alcian blue staining to confirm the presence of accumulated GAG. The formation of cartilaginous tissue in fibrin/PLGA constructs was remarkably evident at each time point of 1, 2 and 4 weeks after in vivo implantation. As shown in Figure 1, in fibrin/PLGA construct, an increase in the implantation period resulted in lower cells to matrix ration similar to that of native hyaline cartilage. Cartilage-isolated cells with lacunae were sparsely distributed within the homogenous ECM in concert with the presence of specific histochemicals property of accumulated proteoglycan-rich matrix and GAG. Whereas, in PLGA, presence of chondrocytes in their natural round morphology increased with an increase in the implantation period; resulted in higher cells to matrix ration. This phenomenon indicated the progress of cartilaginous tissue formation in the PLGA group. The difference between the fibrin/PLGA hybrid construct and the PLGA group was clearly visible in term of the overall cartilaginous tissue formation, cells organization and ECM distribution in the specimens. At the end of the in vivo experiment, the fibrin/PLGA constructs exhibited good quality cartilage-like tissue compared to the PLGA group. We analyzed collagen type II and type I immunolocalization of the fibrin/PLGA hybrid construct. Collagen type II was detected in all specimen. Collagen type II exhibited strong immunopositivity at the specific region of the in vitro fibrin/PLGA hybrid construct, mainly localized surrounding the pericellular and inter-territorial matrix. After in vivo implantation, the fibrin/ PLGA hybrid construct showed more homogeneous distribution of collagen type II than in the PLGA scaffold at each time point as shown in Figure 2. Analysis of collagen type I in fibrin/PLGA hybrid constructs showed moderate

H&E

Safranin-O

Alcian Blue

Fibrin/PLGA 1 wk PLGA Fibrin/PLGA 2 wks PLGA Fibrin/PLGA 4 wks

PLGA

Figure 1. The morphology of cells and distribution of cartilaginous ECM in the fibrin/PLGA hybrid scaffold and PLGA were examined via haematoxylin and eosin (H&E), Safranin-O staining and Alcian blue staining in in vivo.

Collagen type II

Collagen type I

Fibrin/PLG 1 wk PLG Fibrin/PLG 2 wks PLG Fibrin/PLG 4 wks

PLG

Figure 2. Collagen type II was detected in all in vivo specimens and at each time point of 1, 2 and 4 weeks post-implantation.

positive immunoreactivity throughout the ECM of specimens at each time point of 1 and 2 weeks post-implantation. Interestingly, after 4 weeks implantation, both fibrin/PLGA hybrid construct and the PLGA group showed no collagen type I expression, however in some cases, both groups showed very weak collagen type I expression. Fibrin/PLGA hybrid scaffold promotes chondrogenesis of rabbit articular chondrocytes proven by means of morphology, histology, immunohistochemistry, chondrogenic gene expression and sGAG production. This study suggests that fibrin/PLGA hybrid scaffold may serve as a potential cell delivery vehicle and a structural basis for in vivo development of tissueengineered articular cartilage construct. Future studies using appropriate animal model for autologous in vivo system are necessary to further validate the feasibility of fibrin/PLGA hybrid constructs for articular cartilage restoration.

3.3 Fibrin/PLGA Hybrid Scaffolds for IVD In Vitro The intervertebral disc (IVD) is the cornerstone of the joint complex comprises the spinal motion segment. The IVD functions to permit limited motion and flexibility, while maintaining segmental stability by absorbing and distributing external loads. The structure of the normal IVD includes nucleus pulposus (NP) composed primarily of proteoglycans and collagen type II with a capacity to absorb and distribute load and the outer annulus fibrosus (AF) with wellorganized layers consist of collagen type II and collagen type I serve to stabilize the motion segment. The structure and function of the IVD may be altered by processes including normal physiological aging, mechanical factors i.e. trauma and repetitive stress, segmental instability of the spine, inflammatory and biochemical factors. Patient of degenerative disc disease forms the majority of patients with back pain of spinal origin. The IVD undergoes extrinsic morphological changes during their lifetime. In fact, most of the adult population will have a degenerative disc by the 6th decade of life. The current treatment options range from NSAIDS to invasive procedures including spinal fusion and arthroplasty. Unfortunately these treatments are not solving the root of the problem which is the degeneration of IVD itself. As tissue engineering and regenerative research becomes more advance, there is a trend toward reversing the aetiology of disease and much effort has been put into regeneration of IVD [19]. Six weeks-old New Zealand White rabbits were euthanized, lumbar discs were obtained using an osteotome by means of an anterior approach in an en bloc fashion from L1–L2 to L7–S1 IVD. The AF and NP tissue were aseptically

dissected from the lumbar discs. AF and NP cells were cultured separately in the mixture of equal volume of F12/DMEM supplemented with 10% FBS and HEPES buffer 1M. PLGA and fibrin/PLGA hybrid scaffolds were used as same one in section 3.2. Total sGAG production was normalized by the dried weight of each sample and represented as relative sGAG content in percentage (%). After 1, 2 and 3 weeks of in vitro culture AF cells cultured in fibrin/PLGA exhibited 0.290±0.009, 0.341±0.004 and 0.443±0.014 relative sGAG content, respectively. Whilst PLGA exhibited 0.255±0.025, 0.319±0.45 and 0.360±0.007 relative sGAG content, respectively after 1, 2 and 3 weeks. Apparently, sGAG production was higher in fibrin/PLGA than in PLGA and the differences between sGAG production magnitudes were significantly distinguished by week 2 (1.07-fold, p=0.04) and week 3 (1.23-fold, p=0.0003) as shown in Figure 3A. As shown in Figure 3B, after 1, 2 and 3 weeks of in vitro culture NP cells cultured in fibrin/PLGA exhibited 0.169±0.021, 0.276±0.007 and 0.277±0.018 relative sGAG content, respectively. Whilst PLGA exhibited 0.109±0.017, 0.208±0.0207 and 0.230±0.005 relative sGAG content, respectively after 1, 2 and 3 weeks. The sGAG production in fibrin/PLGA was superior to PLGA throughout in vitro culture. The magnitudes were significantly higher by 1.55fold, 1.33-fold and 1.2-fold after 1 week (p=0.008), 2 weeks (p=0.009) and 3 weeks (p=0.038), respectively. Approximately 1x105 cells per scaffold were cultured in the fibrin/PLGA and PLGA scaffolds. Seeded-cells adhered onto scaffold, proliferated and produced matrices filling in the void spaces of the scaffolds. No sign of cartilaginous tissue formation in fibrin/PLGA and PLGA was observed after 1 week of in vitro 3D culture of AF and NP cells. At 2 weeks culture, minimal cartilaginous tissue formation was observed in the fibrin/PLGA seeded with AF cells with several cells clusters filling up the void spaces of the scaffold, Figure 4. The newly synthesized ECM was strongly stained by the characteristic red of Safranin-O, indicating presence of the proteoglycan-rich matrix corroborated with positive Alcian Blue staining confirming GAG accumulation. The formation of cartilaginous tissue in fibrin/PLGA was remarkably evident by 3 weeks of in vitro culture. The closely-packed cells were homogeneously distributed in the basophilic ECM in concert with the presence of specific histochemicals property of proteoglycan-rich matrix and GAG. Fibrin/PLGA developed predominantly superior histoarchitecture when compared to PLGA group. In PLGA, few rounded cells cluster was spotted throughout the specimen. For NP specimen, as shown in Figure 5, cartilaginous tissue formation was also more prominent in fibrin/PLGA than PLGA. Clearly, the

Figure 3. Total sGAG production was normalized by the dried weight of each sample and represented as relative sGAG content in percentage (%). The sGAG production was higher in fibrin/PLGA seeded with AF cells (A) and NP cells (B) than in the PLGA group.

Figure 4. Safranin-O indicating presence of the proteoglycan-rich matrix corroborated with positive and Alcian Blue staining confirming GAG accumulation of the fibrin/PLGA and PLGA group seeded with AF cells at 2 weeks culture.

formation of cartilaginous tissue was evident by the third week of in vitro culture in the fibrin/PLGA. The differences between the fibrin/PLGA and PLGA group both in AF and NP cases were distinguishable in term of overall cartilaginous tissue formation, cells organization and ECM distribution in all specimens. The presence of accumulated proteoglycan-rich matrix and GAG at the core region was significant and was intensely stained at 2 weeks and greatest at 3 weeks. Conversely, AF and NP cells cultured in PLGA without fibrin demonstrated slower progress of cartilaginous tissue formation in vitro.

Figure 5. At 2 weeks cartilaginous tissue formation was more prominent in fibrin/PLGA seeded with NP cells when compared with PLGA. Clearly, the formation of cartilaginous tissue was evident by the third week of in vitro culture in the fibrin/PLGA. The presence of accumulated proteoglycan-rich matrix and GAG at the core region was significant and was intensely stained at 2 weeks and greatest at 3 weeks. Conversely, NP cells cultured in PLGA without fibrin demonstrated slower progress of cartilaginous tissue formation in vitro. The differences between the fibrin/PLGA and PLGA group were distinguishable in term of overall cartilaginous tissue formation, cells organization and ECM distribution in all specimens.

Fibrin supports higher cells proliferation, maintains phenotypic expression and promotes greater sGAG production and cartilaginous tissue formation of AF and NP cells cultured in PLGA. This study suggests that fibrin/PLGA hybrid scaffold may serve as a potential cell delivery vehicle and a structural

basis for in vitro tissue-engineered IVD. This in vitro study revealed promising results; hence, future studies utilizing the in vivo system are necessary to further validate the development of tissue-engineered IVD using fibrin and PLGA composite.

4

Conclusions

Tissue engineering including regenerative medicine shows tremendous potential as a revolutionary research push. Also, many successful results have been reported the potential for regenerating tissues and organs such as skin, bone, cartilage, nerve of peripheral and central, tendon, muscle, corneal, bladder and urethra, and liver as well as composite systems like a human phalanx and joint on the basis of scaffold biomaterials from polymers, ceramic, metal, composites and its hybrids. As previously emphasized, scaffold materials must contain the site of cellular and molecular induction and adhesion and must allow for the migration and proliferation of cell through porosity. It should also maintain strength, flexibility, biostability and biocompatibility to mimic a more natural, three dimensional environments. From this point of view, the control over precise biochemical signal must be needed by the combination of scaffold matrix and bioactive molecules including genes, peptide molecules and cytokines. Moreover, the combination of the cells and redesigned bioactive scaffolds has attempted to expand to a tissue level of hierarchy. In order to achieve this goal, the novel hybrid scaffold biomaterials, the novel scaffolds fabrication methods and the novel characterization methods must be developed.

Acknowledgments This work was supported by grants from KMOWH(0405-BO01-0204-0006) and SCRC (SC4110).

References 1.

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Khang G, Kim S. H., Kim M. S., Lee H. B., 2008. Hybrid, Composite, and Complex Biomaterials for Scaffolds, In: Principles of Regenerative Medicine, edited by A. Atala, R. Lanza, J. A. Thomson, and R. M. Nerem, Elsevier, San Diego, 636-655. Khang G, Kim S. H., Kim M. S., Lee H. B., 2007. A Mannual for the Fabrication of Tissue Engineered Scaffolds, World Scientific Publishing Co.

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Chapter 14 Novel Hydrogel Systems as Injectable Scaffolds for Tissue Engineering Yoon Ki Joung and Ki Dong Park Department of Molecular Science and Technology, Ajou University, Suwon, Korea

1

Introduction

To date a variety of biomaterials have been developed in many biomedical and pharmaceutical fields. Hydrogels are considered as important biomaterials in various biomedical fields on account of their properties, such as highly hydrophilic nature, flexible structure and good biocompatibility [1,2]. From a clinical perspective, the use of these hydrogels as injectable scaffolds (e.g. sponge, mesh etc.) is more attractive than implantable scaffolds with surgical procedures because it can minimize drawbacks, such as patient’s discomfort, risk of infection, scar formation, and cost of the operation [3]. Injectable hydrogels are based on the in situ hydrogel formation triggered by external stimuli, mild chemical reactions and stereocomplexation, etc. In situ forming hydrogel is also advantageous for biomedical applications because the hydrogel formation do not require toxic crosslinking agents or organic solvents. The hydrogel that can induce significant and reversible phase transition in response to environmental stimuli (pH, ionic strength, temperature and electric field etc.) is known as stimuli-responsive hydrogels [4]. Among such hydrogels, thermo-responsive hydrogels are one of the most commonly studied classes of environmentally sensitive polymer systems [5]. Some polymers cause a phase transition upon heating or cooling of the aqueous solution. In such polymers, the temperature range for sol-gel transition can be formed near the body temperature by modulating the balance between hydrophilic and hydrophobic moieties of

block copolymers, realizing to inject scaffolds into the body [6,7]. Pluronic, triblock copolymers of poly(ethylene oxide) and poly(propylene oxide), and their 4-arm derivatives, Tetronic, have been focused on by our group to utilize in thermo-responsive hydrogel. The conjugation of these to natural polymers, such as chitosan and heparin, allow to improving potential as scaffold. External stimuli such as temperature [8-10] and pH [11-13] permit to form a physical hydrogel of which the gelation is reversible and the mechanical properties are largely poor. On the other hands, light [14-17] and mild chemical reactions such as Michael addition [18-20] and enzymatic reaction [21,22] permit to form a chemical hydrogel of which the sol-gel transition is irreversible and the mechanical properties are relatively good. Exceptionally, stereocomplex formation of certain polymer allows forming a physical hydrogel, but the sol-gel transition and physical properties are similar to chemical hydrogels [23]. For biodegradable hydrogels, enantiomeric block copolymers of poly(lactic acid) (PLA) have been investigated as potential in situ hydrogel formation. The copolymers contained enantiomeric PEG-PLA block copolymers with triblock [24,25] and 8-arm (star-shape) [26-28]. In our group, examples of in situ forming hydrogels by Michael addition and stereocomplexation will be introduced. Our group has focused on growth factor delivery by protein-binding affinity of heparin during tissue regeneration using hydrogels. Heparin, the highly sulfated polysaccharide belonging to the family of glycosaminoglycans (GAG), has numerous important biological activities, associated with its interaction with diverse proteins [29]. Heparin is the most popular clinical anticoagulant that has been widely applied to reducing thrombus formation at blood-biomaterial interfaces [30-32]. Over the last few decades, it has been reported that heparin interacts with a number of biologically important proteins including enzymes (i.e., thrombin), lipoproteins, growth factors, chemokines, viral coat proteins and extracellular matrix (ECM) proteins, and so on. [33-35]. Such properties of heparin would be utilized for entrapping, and transporting or retaining the heparin-binding proteins for therapeutic uses. In our group, two kinds of injectable hydrogel systems combined with heparin have been developed and evaluated. In this chapter, four kinds of novel hydrogel systems as injectable scaffolds will be introduced. First example is thermo-responsive chitosan-Pluronic hydrogel preserving biocompatible nature. Second one is another thermoresponsive Tetronic-PLA-heparin hydrogel, which is specialized by biodegradable nature and controlled release of certain proteins. Third one is heparin-conjugated PLGA-PEG-PLGA tri-block copolymer hydrogel, showing

simultaneously in situ hydrogel formation and heparin conjugation. Final example is an in situ forming hydrogel formed by different driving force for gelation, stereocomplexation between enantiomeric PLA chains. It is expected that all examples can be applied to injectable biomaterials in type of hydrogel as cell delivery carriers or bulking materials for tissue engineering and other biomedical fields.

2

Chitosan-Pluronic Hydrogels

The conjugation of chitosan to Pluronic, PEO-PPO-PEO block copolymer, can be considered as an approach of developing biocompatible and injectable scaffold for tissue regeneration. In addition, the conjugation of GRGDS, cell adhesive peptide, is expected to promote the cell adhesion and spreading. A RGD (Arg-Gly-Asp)-conjugated chitosan-based hydrogel was investigated as injectable scaffold for articular cartilage regeneration. RGD-conjugated CP (RGD-CP) copolymer was prepared by coupling carboxyl group in the peptide with residual amine group in CP copolymer. The concentration of conjugated RGD was quantified by amino acid analysis (AAA) and the rheological experiment of RGD-CP hydrogel was investigated. The amount of bound RGD was measured to be 0.135 µg per 1 mg of CP copolymer. Viscoelastic parameters of RGD-CP hydrogel showed thermo-sensitivity and suitable mechanical strength at body temperature for cell scaffolds (over 100 kPa of storage modulus). Viability of bovine chondrocyte and the amount of expressed glycosaminoglycans (GAGs) in RGD-CP hydrogels were evaluated together with alginate hydrogels as a control during 14 days. The both results revealed that RGD-CP hydrogel is more superior to alginate hydrogel as shown in Fig. 1. From these results, it is demonstrated that conjugating RGD to CP hydrogels promotes to improve cell viability and proliferation including extra cellular matrix (ECM) expression. Therefore, obtained results suggest that RGDconjugated CP hydrogels are very suitable for chondrocyte culture and possible its application to tissue engineering of articular cartilage tissue.

3

Tetronic-PLA-Heparin Hydrogels

Heparin-conjugated biomaterials have a lot of potential as the delivery vehicle of heparin-binding proteins and biocompatible materials for tissue regeneration. Such concepts leads to a novel hydrogel which comprised of a thermosensitive

Figure 1. The viability of bovine chondrocyte (upper) and the synthesized GAG amounts (lower) in the RGD-CP hydrogel and alginate hydrogel for 14 days culture.

polymer (Tetronic), heparin and a biodegradable linkage (PLA chain) to utilize as a versatile cell-supporting scaffold for tissue regeneration. A novel scaffold, Tetronic-PLA-heparin (TLH) hydrogel, was prepared for improving tissue

regeneration by coupling heparin to polymerized Tetronic-PLA (oligolactide). Tetronic is tetra-functional block copolymer synthesized by the sequential addition of propylene oxide and ethylene oxide to ethylenediamine [X]. Tetronic® 1307 was found to gel at 25 °C, which is similar with Pluronic® F127. OLA was introduced to the conjugate in order to control hydrophilic/ hydrophobic ratio and give biodegradable property. Conjugating heparin to the synthetic Tetronic-OLA can provide another function as a matrix for delivering heparin-binding proteins. Aqueous TLH solutions showed thermo-sensitive behavior, demonstrating potential for injectable hydrogel. The content and the activity of conjugated heparin were determined to be 0.61 g per gram of total polymer and 67.2% of intact heparin activity, respectively. The basic fibroblast growth factor (bFGF) binding assay presented relatively high bFGF affinity of TLH hydrogel, which indicates applicability for growth factor delivery. Finally, chondrocyte culture on hydrogels revealed that the cell viability and the amount of synthesized glycosaminoglycan (GAG) for TLH hydrogel were higher than those for alginate gel.

4

Heparin-Containing PLGA-PEG-PLGA Hydrogels

For many years, our group has investigated on heparin-conjugated materials and thermo-responsive hydrogels. Heparin has been conjugated to linear PLA [X], star-shaped PLA [X] and Tetronic-PCL [X,X] to prepare blood compatible materials and long-term growth factor delivery carriers, demonstrating that heparin-conjugated materials could be applied to blood compatible biodegradable polymers and polymeric micelles releasing growth factors. We have also studied some thermo-sensitive hydrogels of Pluronic-grafted chitosan [X,X]. Our previous results demonstrated that the hydrogels were thermosensitive and possible to be utilized as injectable materials for tissue regeneration. Recently our group designed a novel system for in situ crosslinkable hydrogel composed of heparin and PLGA-PEG-PLGA as an injectable scaffold for tissue regeneration. The thermo-sensitive hydrogel was prepared through Michael-type addition between thiolated heparin and diacrylated PLGA-PEG-PLGA. Characteristic results demonstrate that the hydrogel improved the bioactivity as well as the thermal and physical stability of the hydrogel. Novel heparin-conjugated PLGA-PEG-PLGA hydrogels were prepared via Michael-type addition between thiolated heparin and PLGA-PEG-PLGA diacrylate. Firstly, thiolated heparin (HP-SH) was prepared by conjugation of thiolacid dihydrazide and following reduction. The structure and the thiol

determination of obtained HP-SH were characterized by 1H NMR and Ellman method. Anticoagulant activity and pKa value of the obtained HP-SH were determined by aPTT test and UV absorbance measurement, resulting in 79.3% and 10.5, respectively. Then, PLGA-PEG-PLGA diacrylate was synthesized by bulk ring-opening polymerization of D,L-lactide (DLLA) and glycolide (GA) with PEG and stannous 2-ethylhexanoate, followed by the acrylation of terminal groups, which was characterized by 1H NMR and GPC. Finally, synthesized HP-SH was conjugated to PLGA-PEG-PLGA diacrylate by Michael-type addition. Phase diagram of hydrogels was recorded by vial tilting method and release test of heparin from the hydrogels was performed, confirming temperature dependent sol-gel transition behaviors. Obtained results demonstrated that our heparin-conjugated hydrogels can be utilized as a novel injectable and tissue-compatible scaffold because of its ability of in situ gel formation, thermo-sensitivity and growth factor binding.

Cumulative released (%)

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80 60 40

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20 0 0

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Time (Hour) Figure 2. In vitro release profiles of bFGF measured by enzyme-linked immunosorbent assay (ELISA) from the (A) TL and (B) TLH hydrogel using supplemented PBS release buffer (0.01 M, pH 7.4) at 37 °C.

5

Four-Arm PEG-PLA Stereo-Complex Systems

The formation of PLA stereo-complexes provides advantages such as easy injection, the high density of crosslinking and fast gelation. In our group, 4-arm

PEG-PDLA (4-PPD) and 4-arm PEG-PLLA (4-PPL) block copolymers were prepared to give a stereo-complexed hydrogel for in situ forming injectable system. The hydrogel could be formed in situ by simply mixing aqueous solutions of 4-PPD and 4-PPL block copolymers, showing fast gelation and high mechanical strength. In situ gel formation through stereocomplex formation is a promising candidate for injectable hydrogels. Enantiomeric fourarm PEG-PLA block copolymers and their stereocomplexed hydrogels were prepared by bulk ring-opening polymerization of D-lactide and L-lactide, respectively, with stannous octoate as a catalyst. 7 Alginate gel

GAG amount (㎍)

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Time (day)

Figure 3. Viability of bovine chondrocyte in TLH hydrogel (white) and alginate hydrogel as a control (gray). On each time points (1, 3, 7 and 14 days), viable cell number was quantified by MTS assay. The GAG amount at each time points (1, 3, 7, and 14 days) produced by bovine chondrocyte encapsulated in TLH hydrogel (white) and alginate hydrogel control (gray).

The prepared polymers were characterized by 1H NMR, FT IR, GPC and TGA, confirming the tailored structure and chain lengths. The swelling and degradation behavior of the hydrogels formed from a selected copolymer series are observed in different concentrations. The results showed that a larger portion of the polymers in a solution induced a slower degradation rate.

Temperature(°C)

(a)

Concentration (wt%) (b) Figure 4. Photograph showing temperature-responsive sol-gel transition of PLGA-PEGPLGA triblock copolymers and phase diagram of PLGA-PEG-PLGA in situ conjugated with thiolated heparin.

The rheological result indicated that the prepared hydrogel underwent in situ gelation and had favorable mechanical strength. In addition, its feasibility as an injectable scaffold was evaluated using a media dependence test for cell culture.

It was found that a Tris solution was more favorable for in situ gel formation than PBS and DMEM solutions. These results suggest that in situ forming hydrogel through the formation of a stereocomplex with enantiomeric 4-arm PEG-PLA copolymers was accomplished. Overall, enantiomeric 4-arm PEGPLA copolymers are new species of stereocomplexed hydrogels that are suitable for further research into injectable hydrogels.

6

Conclusions

A variety type of injectable hydrogels were developed and evaluated as scaffolds for tissue regeneration. The conjugation of RGD and chitosan lead to enhanced cell viability and GAG expression. The conjugation of heparin to hydrogels grants more advantageous properties mechanical properties. In another heparin-conjugated hydrogel, the release test of heparin demonstrated that heparin was covalently connected with PLGA-PEG-PLGA via Michael-

Figure 5. (a) Photographs showing in situ hydrogel formation by mixing 20 wt% solutions of 4-arm PEG-PDLA (4-PPD) and 4-arm PEG-PLLA (4-PPL) and (b) Swelling ratios of stereocomplex hydrogels formed from 18, 20 and 22 wt% solutions at room temperature.

type addition, accordingly was slowly released from PLGA-PEG-PLGA hydrogel matrix. Obtained results suggest that this hydrogel is a promising injectable scaffold for tissue regeneration by the incorporation of growth factors. In PEG-PLA stereo-complexed hydrogels, the swelling/degradation and rheological results revealed that in situ hydrogelation and mechanical strength favorable for injectable hydrogels. In addition, evaluations of the feasibility as an injectable scaffold concluded that the Tris buffer solution is favorable for in situ hydrogel formation but acidification of the hydrogel matrix can limit in vitro evaluations.

Acknowledgments This work was supported by grants from the Korea Health 21 R&D Project, Ministry of Health & Welfare (02-PJ3-PG6-EV11-0002), and NanoBiotechnology Project (Regenomics), Ministry of Science & Technology (B020214), Republic of Korea.

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Chapter 15 Manipulation of Stem Cell Functions On Grafted Polymer Surfaces Naoki Kawazoe, Likun Guo, Guoping Chen, Tetsuya Tateishi Biomaterials Center, National Institute for Materials Science, Tsukuba, Japan

1

Introduction

The surface properties of biomaterials and scaffolds such as chemical composition, nano- or microstructured morphology, wettability, and electrostatic property are quite important for cell behaviors such as cell adhesion, proliferation, extracellular matrix (ECM) secretion, and differentiation [1-5]. To elucidate the effects of surface properties, especially surface chemistry on cell functions, various methods have been reported to present the surfaces with different functional groups. Surfaces with different chemical groups such as methyl, hydroxyl, carboxyl, and amino groups have been used for the culture of mesenchymal stem cells, fibroblasts, chondrocytes, osteoblasts, myoblasts, and neural cells to study the effects on cell functions [6-11]. The effects of surface chemistry on cell functions depend not only on surface composition but also on culture conditions and cell type. Self-assembled monolayers (SAMs) of alkanethiols on gold are a useful model system to investigate the effects of surface chemistry on cell functions systematically. Many researchers have used this system to study the functional surface groups on the adsorption and conformational change of proteins and the resulting effects on cell adhesion, spread, alignment, and proliferation. Methyl, hydroxyl, carboxyl, and amino groups have been presented on gold surfaces by SAMs to study their effects on cell functions in different culture conditions.

The effects depend not only on surface composition, but also on culture condition and cell type. Silane-modified glasses presenting methyl (–CH3), hydroxyl (–OH), carboxyl (–COOH), amino (–NH2), and sulfoxyl (–SH) have also been used. Curran et al. [12, 13] reported the effects of surfaces presenting these functional groups on the differentiation of human mesenchymal stem cells (MSCs). MSCs are a prospective cell source for tissue engineering because they are relatively easy to obtain from a small aspirate of bone marrow, and are multipotent to differentiate into different cell lineages such as osteoblasts, chondrocytes, adipocytes, and neural cells [14]. The glass control and –CH3 surfaces maintained the multipotent phenotype of MSCs, the –NH2- and –SHmodified surfaces promoted osteogenesis, and the –OH- and –COOH-modified surfaces promoted chondrogenesis. The SAM method and silane-modification are only applicable to gold and glass substrates. Matsuda and Ito have developed a method of photochemical modification that can be used to introduce functional groups to the surfaces of any organic substrate [15-17]. The introduced groups are covalently bound to the surface and remain stable during long-term cell culture. This chapter first describes the chondrogenic differentiation of MSCs on photoreactive polymer-grafted surfaces. Then it discusses the adipogenic differentiation of MSCs on micropatterned polymer surfaces. Photochemical modification was used to introduce functional groups to cell culture polystyrene plate surfaces. This method can be used for the surface modification of any organic substrate, and the grafted polymers are stable because of covalent immobilization.

2

Chondrogenic Differentiation of Human Mesenchymal Stem Cells on Polymer-Grafted Surfaces

This section examines the effects of functional groups on mesenchymal stem cell functions. The functional groups were introduced on the surface of polystyrene cell culture plates and their effects on the adhesion, proliferation, and chondrogenic differentiation of MSCs were investigated. Three kinds of surfaces with different functional groups/electrostatic properties were designed and prepared by the photochemical method, namely a NH2/positively charged surface, which was grafted by photoreactive PAAm, a COOH/negatively charged surface, which was grafted by photoreactive PAAc, and a neutral surface, which was grafted by photoreactive PEG.

2.1 Surface Grafting of Polyallylamine, Poly(acrylic acid ), and Poly(ethylene glycol ) Photoreactive azidophenyl-derivatized PAAc (AzPhPAAc, 1) conjugate was synthesized by coupling PAAc with 4-azidoaniline, as shown in Fig. 1a. Azidophenyl-derivatized PAAm (AzPhPAAm, 3) was synthesized by coupling PAAm with 4-azidobenzoic acid (Fig. 1b and c). N-(4-azidobenzoyloxy) succinimide (2) was first synthesized (Fig. 1b). The AzPhPAAm was then synthesized by coupling PAAm with N-(4-azidobenzoyloxy) succinimide (Fig. 1c). Azidophenyl-derived PEG (AzPhPEG, 4) was synthesized by the reaction of bis-amino PEG (PEG) and N- (4-azidobenzoyloxy) succinimide (Fig. 1d). The amounts of the azidophenyl groups in the polymers were determined by 1HNMR. The percentages of the carboxylic groups in the PAAc and the amino groups in the PAAm and bis-amino PEG coupled with the azidophenyl groups were 6.2%, 8.6%, and 100.0%, respectively.

Figure 1. Synthetic scheme of azidophenyl-derivatized poly(acrylic acid) (AzPhPAAc) (a), N-(4-azidobenzoyloxy) succinimide (b), azidophenyl-derivatized polyallylamine (AzPhPAAm) (c) and azidophenyl-derivatized poly(ethylene glycol) (AzPhPEG) (d).

AzPhPAAm, AzPhPAAc, and AzPhPEG were then dissolved in water, placed in the wells of 6-well polystyrene cell culture plates, and air-dried at room temperature in the dark. The plates were irradiated with ultraviolet light at an intensity of 105μJ/cm2. After irradiation, the irradiated plates were immersed in diluted hydrochloric acid (pH 4), alkaline solution (pH 10), and ultrapure water, respectively and then sonicated to completely remove any unreacted polymers. Observation of the grafted surfaces under a phase-contrast microscope also showed no evidence of defects of the grafted surfaces. These results indicate that the polystyrene surfaces were homogenously grafted by the photoreactive polymers. The plates were sterilized with 70% ethanol aqueous solution and used for cell culture.

2.2 Cell Adhesion, Proliferation, and Differentiation 2.2.1 Cell adhesion and aggregation Human bone marrow-derived MSCs (7.0 × 105 cells/mL) at passage 4 in Dulbecco's Modified Eagle Medium (DMEM) serum-free medium were seeded onto each well of the PAAm-, PAAc-, PEG-grafted, and non-grafted 6-well cell culture plates (7.0×104/well). Each well was supplemented with chondrogenic induction medium and cultured for another 2 weeks under static conditions. The chondrogenic induction medium consisted of serum-free DMEM containing glucose, L-glutamine, penicillin, streptomycin, non-essential amino acids, proline, ascorbic acid, dexamethasone, and transforming growth factor-beta 3 (TGF-β3).

Figure 2. Phase-contrast micrographs of MSCs cultured on PAAm-, PAAc-, PEGmodified, and polystyrene (PS) surfaces in chondrogenic induction medium.

The cells on the each polymer-grafted surface were observed with optical microscope (Fig. 2). The cells adhered to the PAAm-grafted surface and spread after 30 min culture, spread more after 3 h, and proliferated to confluence after 3 days. After reaching confluence, the cells gradually aggregated and detached to form pellets. The PAAc-grafted, and polystyrene surfaces also supported MSCs adhesion, but there were some differences among them. The MSCs adhered more rapidly to the PAAm-grafted surface than to the PAAc-grafted and polystyrene surfaces. This difference might be caused by the different electronic properties of these surfaces. The PAAm-grafted surface promotes cell adhesion through the electrostatic attractive interaction between the positively charged surface and the negatively charged cells. The negatively charged PAAc-grafted and polystyrene surfaces do not provide such attractive interaction for cell adhesion. The PEG-grafted surface did not have any electrostatic attractive interaction between surface and cells. It also did not support protein adsorption. Therefore, the cells did not adhere to the PEG-grafted surface and aggregate directly after cell seeding. Although the cells cultured on all the surfaces formed pellets, the effects of the PAAm- and PEG-grafted surfaces were more evident than were those of the PAAc and polystyrene surfaces. Cells on PEG-grafted surface began to aggregate immediately after cell seeding and formed pellets after 24 h. The rapid formation of the pellets on the PEG-grafted surface indicates that cell–cell interaction is stronger than cell–surface interaction. The cells on the other surfaces also formed pellets after becoming confluent. Shrinkage of the confluent cells may result in the detachment of the cell sheet from the surface and in cell aggregation. The cells cultured on the PAAm-grafted surface in the serum medium also detached and formed pellets, but more slowly than those in the chondrogenic induction medium. Culture in the chondrogenic medium facilitated the cell detachment from the surface. 2.2.2 Cell proliferation Cell proliferation was determined by the WST-1 assay. This method is a colorimetric assay for the investigation of cell viability and proliferation based on the cleavage reaction of a tetrazolium salt (WST-1) by mitochondrial dehydrogenases in viable cells. Increased enzyme activity leads to an increase in the amount of formazan dye, which is measured with a spectrophotometer. 96well cell culture plates were used. The three kinds of photoreactive polymers were grafted on the surfaces of the wells in the same manner as the grafting procedure described above by changing the volume of the eluted aqueous solution. The MSC suspension solution in DMEM serum-free medium was added to each well and cultured for 3 h or 1 day.

Figure 3. Proliferation rates of MSCs cultured on PAAm-, PAAc-, PEG-modified, and polystyrene (PS) surfaces in serum-free DMEM. Data represent the average±SD of six samples.

After incubation, the culture medium was aspirated and DMEM supplemented with fetal bovine serum (FBS) was added along with WST-1. The plates were then incubated at 37 °C. After incubation, the absorbance of the samples against the background control on a microtiter plate reader was obtained at a wavelength of 440 nm. The cells proliferated on the PAAm- and PAAc-grafted surfaces at almost the same rate as those on the polystyrene cell culture plates (Fig.3). Almost no cells were detected on the PEG-grafted surface because very few cells attached to the surface.

2.3 Chondrogenic Differentiation 2.3.1 Histological staining The cell pellets formed after 2 weeks culture were harvested for histological examination. The pellets formed during cell culture in the chondrogenic differentiation medium were fixed in formalin, embedded in paraffin, and sectioned. The pellet sections were stained with hematoxylin and eosin (H/E) for the nucleus and stained with safranin O/fast green and toluidine blue to visualize the extracellular glycosaminoglycan (GAGs) (Fig.4). The results indicates that the cells cultured on the PAAm- and PEG-grafted surfaces had a round morphology, whereas those on the PAAc and polystyrene surfaces had a spindle, fibroblast-like morphology. The bright safranin O-positive stain indicated that GAGs were abundant and homogeneously distributed around the cells cultured on the PAAm- and PEG-grafted surfaces. Toluidine blue staining revealed the typical metachromasia (purple color) of articular cartilage, coinciding with the

Figure 4. Hematoxylin/eosin, safranin-O/fast green, and toluidine blue staining of pellets that formed on PAAm-, PEG-, PAAc-modified, and polystyrene (PS) surfaces after culture in chondrogenic induction medium for 2 weeks.

results of safranin O staining. However, the cells on the PAAc-grafted and polystyrene surfaces were not positively stained by safranin O and toluidine blue. 2.3.2 Immunohistological staining Type I, type II collagens, and cartilage proteoglycan were immunohistologically stained using rabbit anti-human type I collagen antibody, mouse anti-human type II collagen monoclonal antibody, and mouse anti-human cartilage proteoglycan monoclonal antibody. The stainings indicate that the pellets formed on the PAAm- and PEG-grafted surfaces were positively stained with type I collagen, type II collagen, and cartilage proteoglycan, while the pellets formed on the PAAc-grafted and polystyrene surfaces did not show any obvious positive staining for type II collagen and cartilage proteoglycan (Fig.5). These results indicate that the PAAm- and PEG-grafted surfaces provided microenvironments for MSCs to change to a round morphology and produce cartilaginous ECM.

Figure 5. Type I collagen, type II collagen, and cartilage proteoglycan staining of pellets that formed on PAAm-, PEG-, PAAc-modified, and polystyrene (PS) surfaces after culture in chondrogenic induction medium for 2 weeks.

2.3.3 Gene expression The gene expressions of type I, II, and X collagens, sox9, and aggrecan in the pellets were analyzed by real-time PCR (Fig.6). The MSCs cultured in the wells in the chondrogenic induction medium for 2 weeks were washed with PBS. The cell pellets were taken from the culture plates and frozen in liquid nitrogen. The frozen pellets were crushed into powder. Total RNA was isolated from the powder using a commercially available RNA extraction reagent and was converted to cDNA by reverse transcriptase. Real-time PCR was performed for GAPDH, types I, II, and X collagens, sox9, and aggrecan. The level of expression of each target gene was normalized to GAPDH. The primer and probe sequences followed those of Martin et al. [18] and Schaefer et al. [19].

Figure 6. Real-time PCR results of mRNA expression of type I collagen (a), type II collagen (b), type X collagen (c), sox 9 (d), and aggrecan (e) of the MSCs cultured on PAAm-, PEG-, PAAc-modified and control surfaces in chondrogenic induction medium for 2 weeks. The data are normalized to GAPDH. MSCs/P4 is the cells seeded onto the surfaces.

The MSCs cultured on the PAAm- and PEG grafted surfaces expressed all these genes. The expression of genes encoding type II and type X collagen, sox9, and aggrecan were increased. The cells cultured on the PAAc- and polystyrene surfaces expressed genes encoding type I collagen, a low level of sox9, and almost no genes encoding type II and X collagens and aggrecan. The gene expression results coincided with the histological, immunohistochemical and biochemical results, which indicated that the PAAm- and PEG-grafted

surfaces promoted the chondrogenic differentiation of the MSCs, but that the PAAc-grafted and polystyrene surfaces did not.

2.4 The Effects of the Polymer Charges on Chondrogenesis The electrostatic properties of a biomaterial surface often affect cell functions such as cell adhesion, proliferation, and differentiation. In this study, the positively charged PAAm-grafted surface supported cell adhesion, proliferation, and the chondrogenic differentiation of MSCs. The negatively charged and polystyrene surfaces supported cell adhesion and proliferation, but not chondrogenic differentiation. The neutral PEG-grafted surface supported neither cell adhesion nor proliferation, but did promote chondrogenic differentiation. Although both the PAAm and PEG-grafted surfaces promoted chondrogenic differentiation of MSCs, the PAAm-grafted surface supported cell adhesion whereas the PEG-grafted surface did not. The positively charged PAAm-grafted surface is more appealing for tissue engineering because, at first, it can support cell adhesion, and then switch to differentiation of the proliferated cells. Thus, cell proliferation and differentiation could occur on the same surface at different times. The 2 week culture results using PAAm, PAAc, and PEG clearly demonstrated the early effects of chargeable polymer grafted surfaces on the chondrogenic differentiation of MSCs, and the results might be applicable to other chargeable polymers. The cells were monolayer and then formed pellets after a few days of culture. The pellets can provide a three-dimensional microenvironment that facilitates chondrogenic differentiation. When cultured on the surface of culture plate for a few passages, primary chondrocytes may change from their original round morphology to a spindle, fibroblast-like shape and lose their ability to express articular cartilage-specific ECM proteins such as type II collagen and aggrecan. Instead, they express and produce fibroblastspecific ECM, type I collagen [20, 21]. The chondrocytes dedifferentiate and change their phenotypes. Three-dimensional microenvironments are necessary to promote cell differentiation [22, 23]. The formation of pellets on PAAm-, PAAc-, PEGgrafted, and polystyrene surfaces in the present study should provide additional information regarding the chondrogenic differentiation of MSCs. The PAAmand PEG-grafted surfaces promoted the chondrogenic differentiation of the MSCs. However, the pellets formed on the PAAc-grafted and polystyrene surfaces did not display any evidence of chondrogenesis, indicating that these surfaces did not support the chondrogenic differentiation of MSCs.

3

Adipogenic Differentiation of Mesenchymal Stem Cells on Micropatterned Polymer Surfaces

The effect of the surface properties on stem cell differentiation has been widely studied because the manipulation of stem cell differentiation remains a great challenge in tissue engineering. However, all these methods compare the effects of diverse surfaces on cell functions by culturing the cells separately on different surfaces. It is desirable to culture the cells on a surface having domains of different chemical groups to compare their effects more directly. The patterning technique provides a powerful tool for creating different patterned domains of chemical groups on a surface [24-27]. Patterning technology provides a useful tool for cell culture to compare the effects of diverse molecules on cell functions under the same culture conditions. Ito et al. [28-33] used this method to micropattern immobilization of growth factors and cytokines to investigate their effects on cell proliferation and differentiation. The present study demonstrated the application of this method for the investigation of surface properties on stem cell differentiation. We micropattern grafted polymers of PAAm, PAAc, and PVA on the surfaces of polystyrene plates and used them for cell culture of MSCs. Their effects on the adipogenic differentiation of MSCs were studied.

3.1 Preparation of Micropatterned Polymer Surfaces The preparation scheme of a micropatterned surface is shown in Fig.7. An aqueous solution of a photoreactive polymer was coated on a polystyrene plate and air-dried. The cast plate was covered with a photomask and UV- irradiated. The photoreactive polymer in irradiated areas should be intermolecularly and intramolecularly crosslinked and grafted to the polystyrene surface. In unirradiated areas, the polymer is not crosslinked but removed by washing with

Figure 7. Preparation scheme of micropatterned polymer surface.

an acidic or alkaline solution or ultrapure water. After washing, the polymer surface having the same pattern as that of the photomask is obtained. Stripe patterns of PAAm, PAAc, and PVA was prepared as follows. First, photoreactive AzPhPAAm and AzPhPAAc were synthesized as described in the previous section. Azidophenyl-derivatized (AAzPhPVA) was synthesized by coupling poly(vinyl alcohol) with 4-azidobenzoic acid. The solutions of AzPhPAAm, AzPhPAAc, and AzPhPVA were was then placed on the tissue culture polystyrene plate and air-dried at room temperature in the dark. The plate was covered with a patterned quartz photomask having a 200-μm-wide stripe and irradiated with ultraviolet light at an intensity of 5.0×103 μJ/cm2 from a distance of 15 cm for 25 s. The 200-μm-wide stripe pattern was used because it was suitable for cell patterning and microscopic observation. After irradiation, the plate was immersed in ultrapure water or dilute hydrochloric acid (pH 3.0) and an alkaline solution (pH 10) and then sonicated to completely remove any unreacted polymer from the unirradiated areas. After complete washing, The PAAm, PAAc-, and PVA-patterned surfaces were obtained. Observation with an optical microscope demonstrated that the PAAm, PAAc, or PVA formed the same striped pattern as that of the photomask (Fig. 8).

Figure 8. Photomicrographs of photomask (a), PAAm- (b), PAAc- (c), and PVApatterned (d) polystyrene surfaces.

3.2 Cell Adhesion on Micropatterned Polymer Surfaces 3.2.1 Cell culture The micropatterned polystyrene plates were put in a cell-culture dish and a glass cylinder was placed over each micropatterned polymer-grafted polystyrene plate. Cell suspension solution was added into the glass cylinder. The hMSCs were cultured in the serum medium for 1 week and allowed to reach confluence. To induce adipogenic differentiation, the hMSCs were cultured in adipogenic differentiation medium [34]. The adipogenic differentiation medium was DMEM serum medium supplemented with 1 μM dexamethasone and 0.5 mM methyl-isobutylxanthine, insulin (10 μg/mL), and 100 μM indomethacin. The cells on the micropatterned surfaces were incubated in the differentiation medium for 1, 2, and 3 weeks and used for cell staining. Alternatively, for analyzing gene expression, the hMSCs were cultured on six-well polystyrene plates homogenously grafted with PAAm, PAAc, and PVA in the adipogenic differentiation medium and serum medium (control) for the same periods. 3.2.2 Cell adhesion and proliferation MSCs were cultured on PAAm-, PAAc-, and PVA-patterned surfaces in serum medium. Immediately after cell seeding, the cells distributed evenly on the polymer-patterned surfaces (Fig. 9). The cells on the PAAm- and PAAcpatterned surfaces adhered to both the PAAm- and PAAc-grafted areas and the non-grafted polystyrene areas after culture for 4 hours, and proliferated to cover the whole patterned surface after culture for 7 days. No cell pattern was observed. The results suggest that the PAAm-grafted, PAAc-grafted, and polystyrene surfaces supported cell adhesion and proliferation of MSCs in serum medium. However, the MSCs cultured on the PVA-patterned surface formed a cell pattern following the substrate pattern. The cells adhered to the non-grafted polystyrene areas, but not on the PVA-grafted areas. The cells on the PVA-grafted areas moved to the non-grafted areas and proliferated to form the same pattern as the striped pattern of the nongrafted area. The PAAm and PAAc-grafted surfaces might affect cell adhesion through electrostatic interaction and adsorbed proteins. Therefore, these surfaces supported cell adhesion. PVA has been reported to inhibit cell adhesion [35]. The cells on the micro-patterned surfaces spread and aligned in parallel with the polymer pattern that is consistent with the results of many investigations showing that cells align with the long axis of grooved surfaces [3]. The striped patterns of all the polymers guided the cell alignment.

Figure 9. Phase-contrast photomicrographs of MSCs cultured on PAAm-, PAAc-, and PVA-patterned polystyrene surfaces in serum medium immediately after cell seeding, and after culture for 4 hours and 1 week.

3.3 Adipogenic Differentiation on Micropatterned Polymer Surfaces Adipogenic differentiation of MSCs on the polymer-patterned surfaces was investigated by the formation of lipid vacuoles and gene expression analysis. The formation of lipid vacuoles and staining with Oil Red-O demonstrates that the MSCs differentiate to adipocytes when cultured in the adipogenic differentiation medium. 3.3.1 Formation of lipid vacuoles After the cells were cultured on the PAAm, PAAc, and PVA-patterned surfaces in serum medium for 1 week, the culture medium was changed to adipogenic differentiation medium and the cells were further cultured for 1, 2, and 3 weeks. The cells were also further cultured in the serum medium as control.

Figure 10. Phase-contrast photomicrographs of MSCs cultured on PAAm-, PAAc-, and PVA-patterned polystyrene surfaces in adipogenic differentiation, and control mediums for 2 week with and without Oil Red-O staining.

Lipid vacuoles were observed after culture for 1 week. The lipid vacuoles were scarce after 1 week culture, and increased after 2 and 3 weeks culture (Fig.10 for 2 weeks culture). The lipid vacuoles distributed evenly on the PAAm- and PAAc-patterned surfaces, while they followed the cell pattern on the PVApatterned surface. The lipid vacuoles were observed only in the areas where the cells existed. The lipid vacuoles were further stained with Oil Red-O[36]. The Oil Red-Opositive cells were observed after culture in adipogenic differentiation medium for 1, 2, and 3 weeks. The Oil Red-O positive cells showed an even distribution on the PAAm and PAAc-patterned surfaces. There was no significant difference between the striped areas of PAAm, PAAc, and polystyrene. However, the Oil Red-O-positive cells showed a pattern on the PVA-micropatterned surface. The PVA grafted areas did not support cell adhesion. Therefore, no Oil Red-Opositive cells were detected on the PVA-grafted areas. The fraction of Oil Red-

O-positive cells increased with culture time. No lipid vacuoles were detected on the PAAm-, PAAc-, and PVA-patterned surfaces in the control medium. The MSCs cultured on the micropattern of the polymers in the control medium did not differentiate into adipocytes, indicating that the induction of adipogenic differentiation of the PAAc-grafted, PAAm-grafted, and polystyrene surfaces require the synergistic effect of the adipogenic differentiation factors. 3.2.2 Gene expression analysis The effects of PAAm, PAAc, and non grafted polystyrene surfaces on adipogenic differentiation of MSCs were further demonstrated by the results of gene expression analysis. The gene expressions of MSCs cultured on the PAAm grafted, PAAc-grafted, and non-grafted polystyrene surfaces were analyzed by real-time PCR (Fig.11). The adipogenesis marker genes encoding peroxisome proliferator γ2 (PPARγ2), lipoprotein lipase (LPL), and fatty acid binding protein 4 (FABP4) were analyzed. PPARγ2 is a transcription factor, and LPL and FABP4 are late markers for adipogenesis.

Figure 11. Real-time PCR results of mRNA expression of PPARγ2 (a), LPL (b), and FABP4 (c) genes of MSCs cultured on PAAm-, PAAc-, and non-grafted polystyrene (PS) surfaces in adipogenic differentiation and control mediums for 2 weeks.

All these genes were detected in the cells when cultured on the PAAm- and PAAm-grafted and polystyrene surfaces. The expression level of these genes increased a small amount with culture time. The cells cultured on these surfaces in the control medium did not express these genes. The gene expression and Oil Red-O staining results indicate that the cells on the PAAm-, and PAAc grafted and non-grafted polystyrene surfaces differentiated to adipocytes. The PAAm-, and PAAc-grafted and non-treated polystyrene surfaces support the adipogenic differentiation of MSCs.

3.4 The Effects of the Grafted Polymer on Adipogenesis The PAAm, PAAc, and PVA were micropattern-grafted on the surface of cellculture polystyrene plate and their effects on the adipogenic differentiation of MSCs were compared. The MSCs adhered to the PAAm- and PAAc-patterned surfaces evenly, while they selectively adhered to the non-grafted areas of the PVA patterned surface. We demonstrated that the PAAm-, PAAc-grafted, and polystyrene surfaces supported adipogenic differentiation of MSCs, while PVAgrafted surfaces did not support the adhesion and adipogenic differentiation of MSCs. The differentiated cells distributed homogeneously on the PAAm- and PAAm-patterned surfaces. No pattern of differentiated cells formed on these two patterned surfaces. The MSCs differentiated to adipocytes on the PVApatterned surface following the micropattern of the surface. There were no cells on PVA-grafted striped areas. These results suggest that the PAAm, PAAc, and polystyrene surfaces promoted adipogenic differentiation of MSCs while the PVA surface did not support adipogenic differentiation.

4

Conclusion

The adhesion, proliferation, and differentiation of mesenchymal stem cells were controlled by changing the electrostatic property using photo-grafted polymer surfaces or micropatterned grafted polymer surfaces. The results in this chapter will provide important information for the design of scaffolds for tissue engineering.

Acknowledgments This work was supported in part by grants from the New Energy and Industrial Technology Development Organization (NEDO).

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Chapter 16 Gamma-ray Irradiated Poly(L-lactide) for Bone Repair Kazuo Isama and Toshie Tsuchiya National Institute of Health Sciences, Tokyo, Japan

1

Introduction

Poly(l-lactide) (PLLA) with a high molecular weight is used as biodegradable screws, pins and plates for internal bone fixation in the orthopedics. PLLA has been well reported on a good osteocompatibility in vivo and in vitro. The γ-ray irradiated PLLA sample was implanted in vivo, and newly bone was formed around the PLLA implant [1]. It was not clear whether there was the effect of γ-irradiation on the formation of newly bone in this result. However, it was the fact that γ-irradiation decreased the molecular weight and mechanical strength of PLLA. If the satisfied mechanical property was maintained, the γ-irradiation was suitable for PLLA devices. We performed the wear test of the γ-irradiated PLLA sheets and measured the particle size distribution of wear debris. On the other hand, PLLA fibers formed bone-like apatite in a simulated body fluid [2]. It was reported that the apatite layer formed on the bioactive glass increased the attachment and initial proliferation of osteoblasts [3]. If the apatite-forming ability of PLLA is increased by γ-irradiation, there may be a good influence on osteoblasts cultured on the irradiated PLLA. We clarified the effects of the γ-irradiated PLLA sheet on the osteoblasts and apatite formation in vitro.

2

Wear Characteristic of the Gamma-ray Irradiated PLLA Sheets

2.1 Gamma-ray Irradiation of the PLLA Sheets The PLLA sheets made of high molecular weight PLLA with thickness of 0.3 mm were obtained from Shimadzu Co. (Kyoto, Japan). The PLLA sheets were γ-ray irradiated at the dose of 10, 25 or 50 kGy using 60Co as the radiation source. The γ-irradiated PLLA sheets were preserved in the silica gel desiccator until next measurement.

2.2 Molecular Weight of the PLLA The molecular weight of the γ-ray irradiated PLLA was determined by gel permeation chromatography. The polydispersity index was calculated as the ratio of the weight average molecular weight (Mw) to the number average molecular weight (Mn). The Mw of the γ-irradiated PLLA extremely decreased with the increasing irradiation dose. The Mw of 271,000 of the unirradiated PLLA was decreased to 95,000 by irradiation at 50 kGy. In contrast, the polydispersity index of the γ-irradiated PLLA was confined to the slight increase with the increasing irradiation dose, compatible with a random cleavage in the degradation mechanism [4–6]. Yoshioka et al. reported γ-irradiation of PLLA caused random cleavage of molecular chain with hydrolysis of ester bonds. In addition, they detected decomposition products having a molecular weight higher than lactic acid in alkali hydrolysis products of irradiated PLLA, and they suggested crosslinkage of molecular chain also occurred [7, 8]. We also analyzed of the γ-irradiated PLLA by high performance liquid chromatography after alkali hydrolysis. However, the quantity of decomposition products having a molecular weight higher than lactic acid was extremely slight. Otto et al. also observed that the molecular weight of PLLA was decreased from 160,000 to 35,200 by γ-irradiation at 25 kGy [9]. Thus, γ-irradiation caused cleavage for molecular chain and decreased the molecular weight of PLLA.

2.3 Wear Test The PLLA sheets were cut out in the disk with the 14.0 mm diameter, and glass column of 11.0 mm diameter and 2.5 g weight was bonded on each PLLA disk. Then, the PLLA specimen was put in the cylindrical vessel of the 30.0 mm inside

diameter, in which bottom plane was #400 waterproof abrasive paper. Five milliliter of balanced electrolyte solution was added in the cylindrical vessel, and the whole vessel was gyrated of 15 mm radius at 200 rpm for 1 hour using a rotatory shaker.

0 kGy

10 kGy

25 kGy

50 kGy

Figure 1. The particle size distribution of wear debris derived from the γ-irradiated PLLA sheet. The center line showed the mean, and the vertical width showed the mean ± 2SD (n = 9).

2.4 Coulter Counter Analysis of Wear Debris The particle size of PLLA wear debris in balanced electrolyte solution obtained by wear test was measured using the Coulter counter. The orifice tube with nominal aperture diameter of 100 μm was used and the particle diameter was measured in the range of 2–60 μm. The particle size distribution was obtained from mean number of each particle diameter. The mean particle diameter of PLLA wear debris was calculated from 9 times experiment.

2.5 Particle Size Distribution of Wear Debris from the PLLA Sheets The particle size distributions of wear debris derived from the γ-ray irradiated PLLA sheets were shown in Fig.1. With the increasing irradiation dose, the particle size distribution of wear debris derived from irradiated PLLA shifted toward the smaller diameter size. The relationship between the irradiation dose of PLLA and the mean diameter of PLLA wear debris was shown in Fig.2. The mean diameter of PLLA wear debris was decreased 9.3% by irradiation at 50 kGy. The mean diameter of PLLA wear debris significantly decreased (P < 0.0001 by ANOVA) with the increasing irradiation dose. The tensile strength of irradiated PLLA also decreased with the increasing irradiation dose [4]. When the abrasive wore the PLLA specimen, the surface of PLLA would be easily cracked, because the tensile strength was lower. In fact, the minute crack had been observed on the surface of the 50 kGy irradiated PLLA disk,

Particle diameter (µm)

5.8 5.6 5.4 *

5.2

* *†

5.0 4.8 0

10

25

50

Irradiation dose (kGy)

Figure 2. The effect of γ-ray irradiation on particle size distribution of PLLA wear debris. *Significant difference compared with unirradiated PLLA at P < 0.01. †Significant difference compared with 10 kGy irradiated PLLA at P < 0.01.

microscopically. Therefore, the decrease of tensile strength of PLLA by the γ-irradiation caused the decrease in particle diameter of PLLA wear debris.

3

Apatite Formation on the Gamma-ray Irradiated PLLA Sheets

3.1 Soaking in the Acellular Medium The PLLA sheet was cut into 14.0 mm diameter disk and laid in a 24-well dish. The complete medium of 1 ml was added without the cells. Then, the dish was stored in a 37°C humidified atmosphere of 5% CO2, and the complete medium was changed three times a week. After soaking for 2 weeks, the PLLA disk was washed in deionized water five times quickly and dried in a silica gel desiccator.

3.2 Surface Analysis The surface of the PLLA sheet before and after soaking in the complete medium without the cells was characterized by scanning electron microscope (SEM), energy dispersive X-ray analysis (EDX), Fourier transform infrared spectroscopy (FT-IR) and X-ray photoelectron spectroscopy (XPS) according to the conventional methods.

Figure 3. The SEM image (a) and the EDX spectrum (b) of the PLLA sheet after soaking in the acellular medium for 2 weeks.

3.3 Apatite Formation on the PLLA Sheets in the Acellular Medium The SEM micrograph exhibited crystal particles on the surface of the PLLA sheet after soaking in the complete medium without the cells (Fig.3a). The crystal particles were identified with hydroxyapatite by EDX, FT-IR and XPS spectra (Fig.3b). The phosphate band in attenuated total reflection (ATR)/FT-IR spectra became strong with irradiation dose (Fig.4a). Moreover, the element rations of calcium and phosphorus increased but that of carbon decreased with irradiation dose, in XPS analysis (Fig.4b). The amount of hydroxyapatite formed on the γ-ray irradiated PLLA sheet increased with irradiation dose [10].

3.4 Surface Carboxyl Group on the PLLA Sheets The surface carboxyl group on the γ-ray irradiated PLLA sheets was determined by XPS in combination with chemical derivatization. The carboxyl group content of the PLLA surface was increased by γ-irradiation. The γ-irradiation increased the apatite-forming ability of the PLLA sheet. Tanahashi and Matsuda reported that some negatively charged groups such as phosphate and carboxyl group strongly induced apatite formation in a simulated body fluid. They described that the apatite formation was initiated via calcium ion-absorption upon complexation with a negative surface-charged group [11]. In our study, the molecular weight of PLLA decreased with hydrolysis of ester bonds by γ-irradiation [4]. Therefore, the surface density of carboxyl group of the γ-irradiated PLLA sheets increased with irradiation dose, and the carboxyl group would promote the apatite-forming ability of the PLLA sheet. a

b 50 Element ratio [mass%]

5

Intensity

4 3 kGy 50 ↑ 2 0 kGy 1 0

C

40 30

Ca

20 10

P

0

1200

1000

800 -1

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0

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Irradiation dose [kGy]

Figure 4. The phosphate band (a) and the element ratios of calcium, phosphorus and carbon of the γ-irradiated PLLA sheet after soaking in the medium.

Figure 5. The surface density of carboxyl group of the γ-irradiated PLLA sheets by XPS with chemical derivatization (a) and the relation between the surface density of carboxyl group and the molecular weight of the γ-irradiated PLLA sheet (b).

4

Osteoblast Differentiation on the Gamma-ray Irradiated PLLA Sheets

4.1 Micromass Culture of Osteoblasts Mouse osteoblast-like MC3T3-E1 cells (RIKEN Cell Bank, Japan) and normal human osteoblast NHOst cells (Clonetics Corporation, MD, USA) were grown in alpha minimum essential medium (α-MEM) supplemented with 20% fetal bovine serum. The PLLA sheet was cut into 14.0 mm diameter disk and laid in a 24-well dish. The 20 μl of cell suspension (2×106 cells/ml) was delivered on the disk. After the cells were attached on the disk, 1 ml of the complete medium that contained 10 mM disodium β-glycerophosphate in the culture medium was added. The complete medium was changed three times a week, and the cells cultured for 2 weeks in a 37°C humidified atmosphere of 5% CO2.

4.2 Assay of Proliferation and Differentiation of Osteoblasts The cell proliferation was estimated with the cell number, the protein and DNA content. The number of the cells cultured on the PLLA sheet was determined by WST-8 assay. The protein and DNA contents of the cell lysate were measured by the Lowry method and the fluorescence assay using Hoechst 33258 dye, respectively [12, 13].

The osteoblastic differentiation was estimated with the calcification, the collagen synthesis and the alkaline phosphatase (ALP) activity. The calcium depositions of the cell cultures were stained by alizarin red S, and the areas stained dark-red were measured. The calcification was calculated as the normalized area in the cell number. Moreover, the collagen synthesis was evaluated by the hydroxyproline content of the cell lysate, and the ALP activity of the cells was measured using p-nitrophenylphosphate as a substrate [12, 13]. The osteoprogenitor cells first differentiate into immatute osteoblasts characterized by the expression of ALP and then into mature osteoblasts characterized by the expression of osteocalcin and calcification [14].

4.3 Osteoblasts Cultured on the PLLA Sheets The cell number of MC3T3-E1 cells cultured on the PLLA sheet did not change with increasing irradiation dose (Fig.6a). The protein and DNA contents of the cells also did not change. The other side, the cell number, protein and DNA contents of NHOst cells cultured on the PLLA sheet slightly decreased with irradiation dose (Fig.6b). The calcification of MC3T3-E1 cells (Fig.7a) and

200

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150

Cell number [%]

Cell number [%]

200

100 50 0

(b)

150 *

100 50 0

0

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25

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Irradiation dose [kGy]

0

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Irradiation dose [kGy]

Figure 6. The cell numbers of (a) MC3T3-E1 and (b) NHOst cells cultured on the γ-irradiated PLLA sheet.

NHOst cells (Fig.7b) remarkably increased with irradiation dose. The collagen synthesis and ALP activity of MC3T3-E1 and NHOst cells also increased as same as the calcification, respectively [12, 15]. The γ-ray irradiated PLLA remarkably promoted the differentiation of osteoblasts. The γ-irradiated PLLA

hardly affected the proliferation but remarkably promoted the differentiation of osteoblasts. It was expected that the low molecular weight PLLA eluted to the medium, because the molecular weight of PLLA decreased by γ-irradiation. Otto et al. also reported when mouse osteoblastic cells were cultured with γ-irradiated PLLA wire for 48 hours, DNA content did not change, but ALP activity increased by 28% [16].

200

(a) **

150

**

100 50 0

Calcification [%]

Calcification [%]

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(b)

150

** *

**

100 50 0

0

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Irradiation dose [kGy]

0

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Irradiation dose [kGy]

Figure 7. The calcifications of (a) MC3T3-E1 and (b) NHOst cells cultured on the γ-irradiated PLLA sheet.

Ikarashi et al. reported that heat treatment of PLLA did not affect the proliferation of MC3T3-E1 cells cultured on heat treated PLLA, but the differentiation of MC3T3-E1 cells was increased. They described that lower change in the molecular weight of PLLA was a cause of stimulation of MC3T3E1 cells cultured on the heat treated PLLA [17]. Moreover, they reported that the low molecular weight PDLLA did not affect the proliferation, but increased the differentiation of MC3T3-E1 cells. They also indicated that low molecular weight PDLLA stimulated the differentiation of MC3T3-E1 cells [18]. In our recent studies, the low molecular weight PLLA enhanced the differentiation of MC3T3-E1 cells but inhibited that of NHOst cells [14, 19]. The present results, which the differentiations of MC3T3-E1 and NHOst cells both increased on the γ-irradiated PLLA sheet, would not be caused by the low molecular weight PLLA. The surface of the γ-irradiated PLLA should good influence on the differentiation of osteoblasts.

Fujibayashi et al. compared in vivo bone ingrowth and in vitro apatite formation on Na2O-CaO-SiO2 glasses. The quantities of newly bone formed on the glasses correlated with their apatite-forming abilities in simulated body fluid. They propose to evaluate the apatite-forming ability in order to confirm the in vivo bioactibity of biomaterials [20]. In our present study, the γ-irradiation enhanced the apatite-forming ability of the PLLA sheet, and then the γ-irradiated PLLA sheet promoted the differentiation of osteoblasts. The osteoblast differentiation should connect with the apatite formation on the γ-irradiated PLLA sheet.

5

Conclusions

The molecular weight of PLLA decreased with the increasing irradiation dose. In addition, the particle size distribution of PLLA wear debris shifted toward the smaller diameter size, and the mean diameter of PLLA wear debris significantly decreased, with the increasing irradiation dose. It was indicated that the lowering of the molecular weight by γ-ray irradiation caused the decrease in tensile strength of irradiated PLLA and the particle size of PLLA wear debris derived from irradiated PLLA. The hydroxyapatite was formed on the PLLA sheet in the acellular medium, and the γ-irradiation enhanced apatite-forming ability of the PLLA. On the other hand, the γ-irradiated PLLA hardly affected the proliferation but promoted the differentiation of osteoblasts with increasing irradiation dose. It was suggested that the connection between the apatite formation and the osteoblast differentiation on the γ-irradiated PLLA sheets.

Acknowledgments A part of this study was financially supported by the Budget for Nuclear Research of the Ministry of Education, Culture, Sports, Science and Technology, based on the screening counseling by the Atomic Energy Commission, and supported by Health Labour Sciences Research Grants from the Ministry of Health Labour and Welfare.

References 1. Otto T.E., Patka P., Haarman H.J.Th.M., Klein C.P.A.T., Vriesde R., 1994. Intramedullary bone formation after polylactic acid wire implantation, J. Mater. Sci.: Mater. Med. 5, 407–410. 2. Yuan X., Mak A.F.T., Li J., 2001. Formation of bone-like apatite on poly(L-lactic acid) fibers by a biomimetic process, J. Biomed. Mater. Res. 57, 140–150. 3. Olmo N., Martin A.I., Salinas A.J., Turnay J., Vallet-Regi M., Lizarbe M.A., 2003. Bioactive sol-gel glasses with and without a hydroxycarbonate apatite layer as substrates for osteoblast cell adhesion and proliferation, Biomaterials 24, 3383– 3393. 4. Isama K., Tsuchiya T., 2001. Change in the particle size distribution of poly(Llactide) wear debris by γ-ray irradiation, Bull. Natl. Inst. Health Sci. 119, 61–64. 5. Reich G., 1998. Ultrasound-induced degradation of PLA and PLGA during microsphere processing: influence of formulation variables, Eur. J. Pharm. Biopharm. 45, 165–171. 6. Mohr D., Wolff M., Kissel T., 1999. Gamma irradiation for terminal sterilization of 17beta-estradiol loaded poly-(D,L-lactide-co-glycolide) microparticles, J. Control. Release. 61, 203–217. 7. Yoshioka S., Aso Y., Otsuka T., Kojima S., 1995. The effect of γ–irradiation on drug release from poly(lactide) microspheres, Radiat. Phys. Chem. 46, 281–285. 8. Yoshioka S., Aso Y., Kojima S., 1995. Drug release from poly(dl-lactide) microspheres controlled by γ-irradiation, J. Control. Release. 37, 263–267. 9. Otto T.E., Patka P., Haarman H.J.Th.M., Klein C.P.A.T., Vriesde R., 1994. Intramedullary bone formation after polylactic acid wire implantation, J. Mater. Sci. Mater. Med. 5, 407–410. 10. Isama K., Tsuchiya T., 2005. Osteoblast differentiation and apatite formation on gamma-irradiated PLLA sheets, Key Eng. Mater. 288–289, 409–412. 11. Tanahashi M., Matsuda T., 1997. Surface functional group dependence on apatite formation on self-assembled monolayers in a simulated body fluid, J. Biomed. Mater. Res. 34, 305–315. 12. Isama K., Tsuchiya T., 2002. Effect of γ-ray irradiated poly(L-lactide) on the differentiation of mouse osteoblast-like MC3T3-E1 cells, J. Biomater. Sci. Polym. Ed. 13, 153–166. 13. Isama K., Matsuoka A., Haishima Y., Tsuchiya T., 2002. Proliferation and differentiation of normal human osteoblasts on dental Au-Ag-Pd casting alloy: Comparison with cytotoxicity to fibroblast L929 and V79 cells, Mater. Trans. 43, 3155–3159. 14. Isama K., Ikarashi Y., Tsuchiya T., 2002. Surface analysis and osteoblast functuon on gamma-ray irradiated poly(L-lactide), BIO INDUSTRY 19, 21–29.

15. Thompson G.J., Puleo D.A., Ti-6Al-4V ion solution inhibition of osteogenic cell phenotype as a function of differentiation timecourse in vitro, Biomaterials 17, 1949–1954. 16. Otto T.E., Nulend J.K., Patka P., Burger E.H., Haarman H.J.Th.M., 1996. Effect of (poly)- L-lactic acid on the proliferation and differentiation of primary bone cells in vitro, J. Biomed. Mater. Res. 32, 513–518. 17. Ikarashi Y., Tsuchiya T., Nakamura A., 2000. Effect of heat treatment of poly(Llactide) on the response of osteoblast-like MC3T3-E1 cells, Biomaterials 21, 1259– 1267. 18. Ikarashi Y., Tsuchiya T., Kaniwa M., Nakamura A., 2000. Activation of osteoblastlike MC3T3-E1 cell responses by poly(lactide), Biol. Pharm. Bull. 23, 1470–1476. 19. Isama K., Tsuchiya T., 2003. Enhancing effect of poly(L-lactide) on the differentiation of mouse osteoblast-like MC3T3-E1 cells, Biomaterials 24, 3303– 3309. 20. Fujibayashi S., Neo M., Kim H.M., Kokubo T., Nakamura T., 2003. A comparative study between in vivo bone ingrowth and in vitro apatite formation on Na2O-CaOSiO2 glasses, Biomaterials 24, 1349–1356.

PART IV

Metallic Biomaterials

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Chapter 17 Titanium Alloys with High Biological and Mechanical Biocompatibility Mitsuo Niinomi Institute for Materials Research, Tohoku University, Sendai 980-8577, Japan

1

Introduction

The main metallic biomaterials are stainless steels, cobalt (Co) alloys, and titanium (Ti) and its alloys. Among these biomaterials, the biocompatibility of Ti and its alloys is the highest. Because Ti alloys exhibit excellent biocompatibility and have high corrosion resistance and specific strength, which is the ratio of density to strength, strength/density, the demand for Ti alloys as biomaterials has increased, and extensive research and development on the use of Ti alloys for biomedical applications is being carried out. Among the current practical Ti alloys available for biomedical applications, Ti-6Al-4V ELI is the most widely used. Ti-6Al-4V ELI was initially used for aerospace applications and then for surgical applications. It was found that vanadium (V) present in Ti6Al-4V ELI was toxic for surgical applications; however, no problems have been encountered. Therefore, Ti-6Al-7Nb and Ti-5Al-2.5Fe where V in Ti-6Al4V ELI is replaced with Nb or Fe, which are nontoxic elements that act as βstabilizing elements, similar to V, have been developed. Among these alloys, the demand for Ti-6Al-7Nb is gradually increasing. Ti-15Sn-Nb-Ta-Pd and Ti15Zr-Nb-Ta-Pd have also been developed [1]. Finally, many low-modulus β-type Ti alloys, whose Young’s moduli are relatively close to that of the cortical bone (10–30 GPa), have been developed. Since most of these alloys are designed considering both biological and mechanical biocompatibility, they are composed of nontoxic and allergy-free elements. Recently, researchers have attempted to develop Ti alloys that have functionality as well as biological and mechanical

biocompatibility for biomedical applications. For example, the authors have developed Ti-29Nb-13Ta-4.6Zr (TNTZ) and they are investigating the properties of this alloy to use it in practice [2-5]. The research and development of functional β-type Ti alloys with a low modulus for biomedical applications are described in this paper, focusing on TNTZ.

2

Biologically and Mechanically Biocompatible Alloy Design

In order to design biologically and mechanically biocompatible Ti alloys, nontoxic and allergy-free alloying elements are selected. In this case, the reported data on cell viability, corrosion resistance, biocompatibility, and rate of metallic allergy of various pure metals and representative metallic biomaterials are used [6, 7]. Consequently, Nb, tantalum (Ta), and zirconium (Zr) are selected as the most harmless alloying elements for Ti. In particular, the addition of nickel (Ni) is avoided. Since Ni is added in very few Ti alloy systems, its hazardous effects are rarely encountered in most of the Ti alloys. Next, the alloy design for reducing the Young’s modulus, which is the most important factor among various mechanical biocompatibility factors, is carried out. It is advantageous to develop β-type Ti alloys because they have a lower modulus than α- and (α + β)-type alloys. The α phase has a hexagonal closed packed structure (HCP), whose atomic density is the highest, and it is the main constituent phase of α-type Ti alloys. In contrast, the β phase has a body centered cubic structure (BCC), which is the main constituent phase of β-type Ti alloys, and its atomic density is relatively lower than that of the α phase. Therefore, large amounts of Nb and Ta are added to Ti in order to fabricate βtype Ti alloys. A small amount of Zr is generally added because it dissolves in both the α and β phases and increases the strength of the resulting alloy On the basis of this concept, the chemical composition of low-modulus Ti-Nb-Ta-Zr system alloys is determined. In this case, it is convenient to use d-electron alloy design based on the DVα-X cluster method [8] because it requires a minimum number of experimental samples to determine the chemical composition of the alloy developed. The authors have developed Ti-29Nb-13Ta-4.6Zr using the delectron alloying method carrying out further research and development on this alloy. Rack et al. [9] have simultaneously developed Ti-35Nb-7Zr-5Ta. Ti36Nb-2Zr-3Ta-O has also been developed [10]5). Even though this alloy has been developed for consumer applications, it is expected to be used in medical field.

3

Biocompability

The contact micro radiogram (CMR) image of the boundaries of bone and lowrigidity Ti-29Nb-13Ta-4.6Zr (in this case, as-solutionized conditions), Ti-6Al4V ELI, or SUS 316L stainless steel implanted into lateral femoral condyles of rabbits, is shown in Fig. 1 [11]. Each specimen is surrounded by a newly formed bone; the bone tissue is partly in direct contact with the specimen. However, the extent of the direct contact is greater in Ti-29Nb-13Ta-4.6Zr than in Ti-6Al-4V ELI and SUS 316L stainless steel. Therefore, it can be concluded that the biocompatibility of Ti-29Nb-13Ta-4.6Zr with bone is excellent.

Observation of tissue reaction and bone formation

Direct Ti-29Nb-13Ta-4.6Zr

Direct Ti-6Al-4V ELI

Lucent line SUS 316L stainless steel

Figure 1. CMR image of boundary of each specimen and bone at 8 weeks after implantation.

4

Mechanical Biocompatibility

4.1 Low Modulus A low modulus has been proved to be effective in accelerating the healing of bone fracture, inhibiting bone absorption, and good bone remodeling through experimental implantation of an intramedullar rod made of TNTZ into a rabbit tibia fracture model [12]7). The bending strengths of tibia, in which the intramedullar rod is implanted, and control tibia have been reported [13]8)．A comparison of three-point bending strengths between a healed tibia fracture model, from which the intramedullar rod made of TNTZ, Ti-6Al-4V ELI, or SUS 316L stainless steel has been removed 16–18 months after implantation,

and the control tibia is shown in Fig. 2 [13]8). It is observed that the three-point bending strength of the healed tibia fracture model is less than that of the control tibia in the cases of Ti-6Al-4V ELI and SUS 316L stainless steel intramedullar rods. On the other hand, the three-point bending strength of the healed tibia fracture model is greater than that of the control tibia in the case of the TNTZ intramedullar rod. Further, the bone structure has been reported to be abnormal in the case of the SUS 316L stainless steel intramedullar rod [12]7).

kgf 80

TNTZ (n=3)

Ti-6Al-4V ELI (n=3)

SUS316L (n=3)

+19.4%

60

-15.0% -0.2% +13.2%

40

-12.5%

+9.2% -8.5%

-9.1%

※died by +10.5% disease at 43 weeks

20 Control side

Intramedullary rod side

Control side

Intramedullary Control rod side side

Intramedullary rod side

Figure 2. Three-point bend strength: fracture load.

The fact that the low Young’s modulus significantly affects bone remodeling has also been reported by observing the healing processes of fracture models of tibia in which bone plates made of TNTZ, Ti-6Al-4V ELI, and SUS 316L stainless steel [14]9) have been implanted. Figs. 3, 4, and 5 [14]9) show the X-ray follow-up of bone fracture, healing from 4 through 18 weeks after the bone plate made of SUS 316L stainless steel, Ti-6Al-4V ELI, or TNTZ was implanted into the rabbit tibia fracture model. Fracture healing was almost the same in the cases of SUS316L stainless steel, Ti-6Al-4V ELI, and TNTZ. Initially, callus formation was observed 2 weeks after implantation, which increased to 3 weeks; this was in agreement with the normal fracture healing time. Bone union occurred 4 weeks after implantation, and the fracture line was barely visible approximately 8 weeks after implantation. The experimental fracture trace completely disappeared 16–20 weeks after implantation. However, bone atrophy (thinning of the cortical bone) was observed under the bone plate,

which occurred at different time intervals in the cases of SUS 316L stainless steel, Ti-6Al-4V ELI, and TNTZ. In the case of SUS316 stainless steel (Fig. 3), the atrophy of the cortical bone began 7 weeks after implantation, and the bone almost disappeared 12 weeks after implantation. In the case of Ti-6Al-4V ELI (Fig. 4), the atrophy of the bone began 7 weeks after implantation, and the bone almost disappeared 14 weeks after implantation. In the case of TNTZ (Fig. 5), the atrophy of the bone began 10 weeks after implantation, and the bone almost disappeared 18 weeks after implantation. Therefore, since the period from the beginning of bone atrophy to disappearance of the bone is the longest in the case of TNTZ, a low Young’s modulus is required to inhibit bone atrophy.

4ｗ

8ｗ

12ｗ

14ｗ

18ｗ

Figure 3. X-ray follow-up 4 to 18 weeks after implantation for SUS 316L stainless steel.

4ｗ

8ｗ

12ｗ

14ｗ

18ｗ

Figure 4. X-ray follow-up 4 to 18 weeks after implantation for Ti-6Al-4V ELI.

4ｗ

8ｗ

12ｗ

18ｗ

14ｗ

Figure 5. X-ray follow-up 4 to 18 weeks after implantation for TNTZ.

A

A B

(a)

(b) Middle

A B (a)

(b) Distal

B (c)

A B (c)

Figure 6. CMR images of cross sections of middle and distal parts of rabbit tibia fracture model, in which a TNTZ bone plate was implanted for 44 weeks, and the control tibia: (a) cross section of fracture model, (b) high-magnification CMR image of branched part of outer and inner formed bone, and (c) cross section of control tibia.

Figure 6 shows the CMR images of the cross sections of the middle and distal parts of the rabbit tibia fracture model, in which a TNTZ bone plate was

implanted for 44 weeks, and the control tibia. An increase in the tibia diameter can be observed in both the middle and distal parts. With regard to the increase in the tibia diameter in the case of TNTZ, a double-wall structure with different X-P densities and a clear boundary line in the middle and distal parts is observed, where the shape of the inner wall is similar to that of the original cortical bone. Therefore, it seems that the outer cortical bone is newly formed, and the intramedullar bone tissue is formed from the remains of the old cortical bone, which is a possible result of bone remodeling with the low-rigidity bone plate. This can be attributed to the fact that the increase in the tibia diameter increases the bending rigidity of the tibia, which may reduce the shear stress around the point of fixation.

4.2 Fatigue Strength Since TNTZ is a β-type Ti alloy, its fatigue strength can be improved by second-phase precipitation hardening through aging treatment. Figure 7 [15]10 ) shows the aging curves of TNTZ aged at 573 K, 673 K, and 723 K after solution treatment. The hardness increases with the aging time at all aging temperatures, and at the same aging time, the hardness is greater with aging temperature. 400

PA

Vickers Hardness, HV

573 K

PA

673 K

350

OA

723 K OA

300

UA

Under Aging

PA

Peak Aging

UA UA

OA Over Aging PA

250

OA UA

200

150

As-ST

1

10

102

103

104

Time, t / ks Figure 7. Relationship between Vickershardness and aging time of TNTZ subjected to aging at 573 K, 673 K and 723 K after solution treatment.

.

The results of fatigue life evaluations (S-N curves) of TNTZ in under-aged (UA), peak-aged (PA), and over-aged (OA) conditions at the above-mentioned aging temperatures reveal that the fatigue life of untreated TNTZ is greater than that of TNTZ under solutionized conditions. The S-N curves of TNTZ in UA, PA, and OA conditions at the aging temperatures of 573 K and 723 K are shown in Figs. 8 and 9 [15]10) as the representative ones. A comparison of the fatigue limit range between TNTZ and Ti-6Al-4V ELI, which is the representative Ti alloy for biomedical applications, is also shown in these figures. The fatigue limit in all the cases is within the fatigue limit range of Ti-6Al-4V ELI. Under all conditions, the fatigue limits of TNTZ at 723 K are greater than those of TNTZ at 573 K, and they lie between the central and upper fatigue limits of Ti6Al-4V ELI. Fatigue Limit

Tensile Strength

Fatigue Ratio

AST

320 MPa

550 MPa

0.58

UA573 K

505 MPa

880 MPa

0.57

PA573 K

548 MPa

899 MPa

0.61

OA573 K

530 MPa

805 MPa

0.67

Maximum Cyclic Stress, σmax / MPa

1000 800

Fatigue Limit Range of Ti-6Al-4V ELI

600 400 200 0

104

105

106

107

Number of Cycles to Failure, Nf Figure 8. S-N curves of AST (as-solutionized), UA573 K, PA573 K, and OA573 K obtained from fatigue tests in air.

Maximum Cyclic Stress, σmax / MPa

Fatigue Limit

Tensile Strength

Fatigue Ratio

AST

320 MPa

550 MPa

0.58

UA723 K

590 MPa

982 MPa

0.60

PA723 K

680 MPa

1059 MPa

0.64

OA723 K

637 MPa

1006 MPa

0.63

1000 800 600

Fatigue Limit Range of Ti-6Al-4V ELI

400 200 0

104

105

106

107

Number of Cycles to Failure, Nf Figure 9. S-N curves of AST, UA723 K, PA723 K, and OA723 K obtained from fatigue tests in air.

The tensile characteristics under the same aging conditions are shown in Fig. 10 [15]10). The tensile strength increases, but elongation decreases significantly, except in the case of the UA condition at 573 K. It has been reported that the increase in the tensile strength is significant, but the elongation is fairly low at an aging temperature of 673 K, although its S-N curve is not represented. At an aging temperature of 723 K, both the tensile strength and elongation increase, leading to a good balance between strength and elongation. Among all the three conditions, the balance between strength and elongation is found to be the best under PA conditions. Figure 11 [15]10) shows Young’s moduli of TNTZ subjected to various heat treatments. It can be observed that the Young’s modulus increases after aging treatment, and it is approximately 80 GPa under PA conditions. However, this value is fairly small. The microstructures of

TNTZ aged at 573 K and 723 K are shown in Figs. 12 and 13[15]10), respectively. The ω phase is precipitated in the former case, and the α phase is precipitated in the β matrix phase in the latter case. The ω phase enhances the trend of brittleness.

1400

30 0.2% Proof Stress Tensile Strength

25

Elongation

1000

20

800 15 600 10

400 200

5

0

0

Elongation ( %)

0.2 % Proof Stress, σ0.2 / MPa

Tensile Strength, σB / MPa

1200

A O

PA

UA

A O

PA

UA

3 72

3 72

3 72

K

K

K

K

K

3 67

3 67

K

K

K

K

3 67

73

73

73

A5 O

5 PA

5 UA

T

AS

Figure 10. Tensile properties of TNTZ subjected to solution treatment at various aging temperatures.

OA 723 K

PA 723 K

UA 723 K

OA 673 K

PA 673 K

OA573 K

UA 673 K

40

PA573 K

UA573 K

80

AST

Modulus of Elasticity / GPa

120

0 Figure 11. Modulus of elasticity of TNTZ subjected to various heat treatments.

(a) UA573 K

: β phase 200 nm

: ω phase

(b) PA573 K

: β phase : ω phase

(c) OA573 K

: β phase : ω phase

Figure 12. TEM micrographs and key diagrams of (a) UA573 K, (b) PA573 K, and (c) OA573 K. Beam direction is parallel to [110].

(a) UA723 K Grain Boundary

500nm

: β phase : α phase

(b) PA723 K

Grain Boundary

: β phase : α phase

(c) OA723 K

Grain Boundary

: β phase : α phase

Figure 13. TEM micrographs and key diagrams of (a) UA723 K, (b) PA723 K, and (c) OA723 K. Beam direction is parallel to [110].

900 Fatigue Limit Range of Ti-6Al-4V ELI

Maximum Cyclic Stress,σmax/MPa

800

700 TNTZCR TNTZCR aged at 673K

600

TNTZCR aged at 723 K

500

400

300 103

104

105

106

107

108

Number of Cycles to Failure，Nf Figure 14. S-N curves of TNTZ in as-solutionized conditions (TNTZST) and as-coldrolled conditions (TNTZCR) and TNTZST and TNTZCR subjected to aging at 598 K, 673 K, and 723 K for 259.2 ks with fatigue limit range of Ti-6Al-4V ELI in air.

βtr: 1013K

ST : 1063K, 3.6ks

WQ

Aging : 723 K

TNTZCR CRR : 87. 5%

673 K

, 259.2ks

WQ

Figure 15. Schematic representation of thermomechanical processing of TNTZ. ST, WQ, and CRR indicate solution treatment, water quenching, and cold rolling ratio, respectively.

The fatigue life of TNTZ can be increased further by thermomechanical treatments. The fatigue life of TNTZ can be found in the upper fatigue limit range of Ti-6Al-4V ELI (Fig. 14 [16]11)) by thermomechanical treatment together with solution treatment, cold sever rolling, and aging treatment, which is schematically shown in Fig.15 [16]11). In addition, the balance between strength and ductility is excellent, although the Young’s modulus increases to approximately 80 GPa; however, this value is still considerably lower than that of Ti-6AL-4V ELI. Therefore, α phases with lengths of several hundred nanometers are precipitated in the β matrix phase.

4.3 Fretting Fatigue

Maximum Cyclic Stress, σmax /MPa

When fretting and fatigue occur simultaneously, the fatigue strength of Ti alloys used for biomedical applications decreases significantly. This may occur in the cases of bone plate and screw and hip joint stem and bone, where two bodies remain in contact with each other under cyclic loading conditions. This is termed fretting fatigue. Fretting fatigue is important for biomaterials. Figure 16 [17]12) shows the S-N curves of TNTZ subjected to solution treatment. These curves were obtained from plain and fretting fatigue tests

600 Plain fatigue in air Plain fatigue in Ringer’s solution Fretting fatigue in air Fretting fatigue in Ringer’s solution

500 400 300 200 100 0

R=0.1 f=10Hz

104

105 106 Number of Cycles to Failure, Nf

107

Figure 16. S-N curves of TNTZ subjected to solution treatment obtained from plain fatigue and fretting fatigue tests in air and Ringer’s solution.

Modulus of Elasticity, E/ GPa

performed in air and in a simulated body fluid (Ringer’s solution) with N2 gas bubbling. The plain fatigue strength in Ringer’s solution is nearly equal to that in air in both low- and high-cycle fatigue life regions. However, the trend in the fretting fatigue strength in air and in Ringer’s solution is different from that of the plain fatigue strength and between low- and high-cycle fatigue life regions. Further, the fretting fatigue strength in Ringer’s solution is greater than that in air in low-cycle fatigue life region; however, in the high-cycle fatigue life region, the fretting fatigue strength in Ringer’s solution is less than that in air. This occurs due to a significantly high lubrication effect caused by Ringer’s solution in the low-cycle fatigue life region and a significantly high corrosion effect caused by Ringer’s solution in the high-cycle fatigue life region. It has been reported that the fatigue strength due to fretting decreases with the Young’s modulus of the material, as shown in Fig.17 [17]12). Therefore, TNTZ is advantageous for biomedical applications in terms of stress shielding and fretting fatigue.

90

Ti-15Mo-5Zr-3Al (Annealing)

80

Ti-29Nb-13Ta-4.6Zr (STA) （ STA）

70

60

Ti-29Nb-13Ta-4.6Zr （ ST） (ST)

50 1.0

1.5

2.0

2.5

3.0

3.5

4.0

Pf/Ff Figure 17. Relationships between fretting damage ratio Pf/Ff, modulus of elasticity, or Vickers hardness; Pf: plain fatigue limit and Ff: fretting fatigue limit.

4.4 Wear The wear resistance of Ti alloys is generally inferior to that of other biomaterials such as Co-Cr alloys and alumina. There is a need to improve the wear resistance of Ti alloys used in biomedical applications. Surface hardening is one of the most effective methods for improving the wear resistance of Ti alloys. Several techniques such as oxidation, nitriding, electroplating, PVD/CVD coating, and thermal spray have been investigated. Among these techniques, oxidation and gas nitriding have relatively simple operation processes. When Ti alloys are oxidized in air, a hard oxide layer of TiO2 is formed in addition to the α-phase; that is, an oxygen-rich region is formed near the surface due to the diffusion of oxygen. In the α-case, the precipitation of the α-phase is enhanced because oxygen is an α-stabilizing element; the precipitation of the αphase increases the hardness of Ti alloys. Fig. 18 [18]13) shows the weight losses of TNTZ and Ti-6Al-4V ELI after the friction wear tests in Ringer’s solution. After solution treatment, each alloy is subjected to oxidation at 773 K and 1073 K for various durations indicated in the figure. The wear resistances of both TNTZ and Ti-6Al-4V ELI are improved by oxidation, although the wear resistance of Ti-6Al-4V ELI decreases due to oxidation at 773 K. The wear resistance of Ti-29Nb-13Ta can be improved at a relatively low temperature.

as ST 43.2 ks oxidized

3.6 ks oxidized 86.4 ks oxidized

Weight Loss / g

0.006 0.005 0.004 0.003 0.002 0.001 0

at 773K

at 1073K

Figure 18. Weight loss of TNTZ after friction wear tests in Ringer's solution. Each alloy was subjected to oxidation at 773 K and 1073 K for various time intervals after solution treatment.

TNTZ1223NP

TNTZ1123NP

TNTZ1073NP

10

TNTZ1023NP

20

TNTZST

Weight Loss / mg

30

0 Figure 19. Weight loss of TNTZ subjected to solution treatment and nitriding obtained from friction wear tests in Ringer’s solution.

Figure 20. Cell viabilities of TNTZ and Ti64 subjected to solution treatment and nitriding, CP Ti, and control (cell disc).

Fig. 19 [19]14) shows the weight losses of Ti-6Al-4V ELI (Ti64) subjected to solution treatment and gas nitriding after friction wear tests in air and Ringer’s solution. It was observed that the wear resistance of Ti-6Al-4V ELI improved after subjecting it to gas-nitriding at temperatures in both air and Ringer’s

solution In this case, hard nitrides such as TiN and TiN2 are formed on the surface of Ti alloys, the amounts of which increase with temperature. As the gas-nitriding temperature increases, the ratio increase of TiNi is greater than that of TiN2. Furthermore, the cytotoxicity of Ti-6Al-4V ELI subjected to gas nitriding also improves, as shown in Fig. 20 [19] 15).

5

Superelastic Function

TiNi is the only shape memory alloy that has been put into practical use, and it is receiving considerable attention for use in stents, guide wires of catheters, orthodontic wires, etc. However, it contains a large amount of Ni, which is a high-risk element causing metal allergies. Therefore, there is a need to develop shape memory Ti alloys that do not contain Ni. In other words, functionalities such as superelasticity and shape memory are required in addition to a low Young’s modulus for biomedical Ti alloys. Superelastic and/or shape memory β-type Ti alloys without toxic and allergic elements are being currently under development. As shown in Fig. 21 [20]15), after appropriate thermomechanical treatment, TNTZ β-type Ti alloys exhibit superelastic behavior; however, it does not

1200

Stress, MPa

1000

Maximum elastic strain ≒2.8%

800 600 400 200 0

0

φ1.0mm W.R.=99.89%

0.5

1.0

1.5

2.0

2.5

3.0

3.5

4.0

Strain, % Figure 21. Tensile loading-unloading S-N curves of as-cold-drawn TNTZ (f1.0 mm). WR indicates work ratio.

exhibit a shape memory effect. The reason for the superelastic behavior of the alloy is not yet clearly understood. The maximum elastic strain is observed to approximately 3% after thermomechanical treatment. In the Ti-Nb-Ta-Zr system alloy, a change in the chemical composition can induce a shape memory effect, which is caused by the deformation-induced martensitic transformation and its reversion. The following Ti-based alloys have been developed: Ti-Nb-Sn system alloys, Ti-Mo-Ga system alloys, Ti-Nb-Al system alloys, Ti-Mo-Al system alloys, TiTa system alloys, Ti-Nb system alloys, Ti-Sc-Mo system alloys, Ti-Mo-Ag system alloys, Ti-Mo-Sn system alloys, and Ti-Nb-Ta-Zr system alloys [21, 22]16) 17). Since shape memory β-type Ti alloys are pseudo-elastic in nature, they exhibit a superelastic behavior. It has been reported that by varying the composition of the elements in the alloys, approximately 6% elastic strain can be obtained in one of the Ti-Nb-Al system alloys [23]18). The shape memory effect of these alloys can be attributed to the deformation-induced martensitic transformation and its reversion. Ni-free superelastic and/or shape memory β-type Ti alloys are expected to be used in stents, guide wires for catheters, orthodontic wires, etc.

6

Bioactive Surface Modification

Bioactive surface modification is required for Ti and its alloys in order to increase their biofunctionality. Hydroxyapatite surface modification is widely investigated in order to achieve direct bonding with bone. In this case, there is a possibility that hydroxyapatite may be either exfoliated from the Ti alloy substrate or fractured by loading. In particular, a surface-modified layer such as a hydroxyapatite layer is exposed to mechanically severe conditions due to cyclic stress. The authors are currently investigating the hydroxyapatite surface modification of TNTZ using a calcium phosphate glass coating method [24]19). The firing process at a temperature over solution treatment temperature is included in that case, and then the strength deceases, there is a case where the aging treatment is carried out when the increase in the strength is required. Subsequently, the severity of the mechanical conditions increases. Fig. 22 [25]20) shows the tensile bonding strength of 20 μm and 5 μm thick calcium phosphate glass ceramic layers coated on TNTZ subjected to aging treatments at 723 K for different time intervals. It is observed that the decrease in the tensile bonding strength of the coating layer can be inhibited using the 5 μm. This implies that exfoliation or cracking of the coating layer occurs when its thickness is 20 μm, while this does not happen not when the thickness is 5 μm.

Tensile Bonding Strength, σb / MPa

50 : DC20

40

: DC5 30 20 10 0

1

BA

10

102

103

Aging Tme, ta/ks

Figure 22. Relationship between tensile bonding strength of calcium phosphate glassceramic coating layer with a thickness of 5 μm (DC5) or 20 μm (DC20), and aging time in Ti-29Nb-13Ta-4.6Zr. BA indicates before.

Maximum Cyclic Stress (MPa)

The S-N curves obtained from the fatigue tests of TNTZ with a 5 μm thick coating layer subjected to various aging treatments are shown in Fig. 23 [26]21). The fatigue strength of TNTZ is significantly improved by subjecting it to aging treatment after coating treatment. An SEM (scanning electron microscopy) 1000 TNTZ TNTZ1 TNTZ2

800

f=10 Hz

600

400

200 103

104

105 106 107 Number of Cycle to Failure

Figure 23. S-N curves of as-solutionized TNTZ, calcium phosphate-glass ceramic coated TNTZ (TNTZ1) and aged calcium phosphate-glass ceramic coated TNTZ (TNTZ2).

fractograph of the TNTZ fatigue fracture surface is shown in Fig. 24 [26]21). Exfoliation or cracking of the coating layer is not observed. The strong bonding of the coating layer can be maintained by controlling its thickness. Simple bioactive surface modification processes such as alkali treatment [27, 28]22) 23) and electrochemical treatment24) are applicable to TNTZ. However, in these processes, the formability of bioactive hydroxyapatite is less than that on the case of pure Ti.

Calcium phosphate glass-ceramic coated layer -

Matrix

5μm

Figure 24. SEM fractograph of calcium phosphate glass-ceramic coated TNTZ (number of cycles N = 557,862).

7

Conclusions

In order to decrease the Young’s modulus, porous materials such as porous pure Ti and its alloys are energetically investigated, although this is not mentioned in our study. In terms of functionality, the further high ordering of biofunctionality of biomaterials through harmonizing metals, ceramics and polymer is progressed. For example, investigations on combining metals and polymer are energetically carried out. In these cases, the main base metallic biomaterials are Ti and its alloys. There is a need to improve the performance and functionality of these alloys. Further research and development of new harmless Ti alloys are expected.

Acknowledgments The author would like to thank Professors T. Hattori with Meijo University, Nagoya, japan and T. Kasuga with Nagoya Institute of Technology, Nagoya, Japan, and Dr. K. Morikawa with Aichi Medical University, Nagakute, Japan for their coraborate works. The author also would like to thank Dido Steel co. Ltd. and GCOE at Tohoku University for their support to complete this paper.

References 1. Okazaki Y., Rao S., Tateishi T., Ito Y., 1998. Cytocompatibility of Various Metal and Development of New Titanium Alloys for Medical Implants, Mat. Sci. Eng. A A243, 250-256. 2. Niinomi M., Kuroda D., Morinaga M., Kato Y., Yashiro T., 1998. New β Type Titanium Alloys with High Biocompatibility, Non-Aerospace Application of Titanium and Its Alloys, Eds. Froes, F. H., Allen, P. G., Niinomi, M., TMS, 217223. 3. Niinomi M., 2003. Cyto-toxicity and Fatigue Performance of Low Rigidity Titanium Alloy, Ti-29Nb-13Ta-4.6Zr, for Biomedical Applications, Biomaterials 24, 26732683. 4. Niinomi M., Hattori T., Niwa S., 2004. Biomaterials in Orthopedics, Material Characteristics and Biocompatibility of Low Rigidity Titanium Alloys for Biomedical Applications, 2004. Eds. Yaszemski, M. J., Trantolo, D. J., Lewandrowski, K. U., Hasirci, V., Altobelli, D. E., Wise, D. L., Marcel Dekker, INC, 41-62. 5. Niinomi M., Akahori T., Makai M., 2007. In-Situ X-ray Analysis of Mechanism of Nonlinear Super Elastic Behavior of Ti-Nb-Ta-Zr System Beta-type Titanium Alloy for Biomedical Applications, Mater. Sci. Eng. C in press,. 6. Kawara H., 1992. Cytotoxicity of Implantable Metals and Alloys, Bulletin of Japan Institute of Metals 31, 1033–1039. 7. Steinemann S. G. 1980. Corrosion of Surgical Implants-in vivo and in vitro Tests”, Evaluation of Biomaterials, eds. G. D. Winter, J. L. Leray, K. De Groot, and John Wiley & Sons Ltd.1–34. 8. Kuroda D., Niinomi M., Morinaga M., Kato Y., Yashiro T., 1998. : Design and Mechanical Properties of New-type Titanium Alloys for Implant, Mater. Sci. and Eng. A A243, 244–249. 9. Rack H. J., Qazi J. I., 2005. Titanium Alloys for Biomedical Applications, Mater. Sci. Eng. C 26, 269–1277. 10. Saito T., Furuta T., Hwang J. H., Kuramoto S., Nishino K., Suzuki N., Chen R., Sakuma T., 2003. Multifunctional Alloys Obtained via Dislocation-free Plastic Deformation Mechanism, Science 300, 464–467.

11. Niinomi M., Hattori T., Morikawa K., Kasuga T., Suzuki A., Fukui H., Niwa S., 2002. Development of Low Rigidity 　type Titanium Alloy for Biomedical Applications, Mater. Trans. 43, 2970–2877. 12. Niinomi M., Hattori T., Niwa S., 2004. Biomaterials in Orthopedics, Material Characteristics and Biocompatibility of Low-rigidity Titanium Alloys for Biomedical Applications, eds. M. J. Yaszemski, D. J. Trantolo, K. U. Lewandrowski, V. Hasirci, D. E. Altobelli, and D. L. Wise, Marcel Dekker Inc., 41–62. 13. Niinom, M., Akahori T., Nakai M., Hattori T., 2007. Low-modulus Multifunctional Titanium Alloys for Biomedical Applications, Kinzoku 77, 128–134. 14. Sumitomo N., Noritake K., Hattori T., Morikawa K., Niwa S., Sato K., Niinomi M., 2008. Experiment Study on Fracture Fixation with Low-rigidity Titanium Alloy— Plate Fixation of Tibia Fracture Model in Rabbit, J. Mater. Sci.: Mater. Med. 19, 1581-1586. 15. Akahori T., Niinomi M., Noda A., Toda H., Fukui H., Ogawa M., 2006, Effect of Aging Treatment on Mechanical Properties of Ti-29Nb-13Ta-4.6Zr Alloy for Biomedical Applications, J. Japan Institute of Metals 70, 295–303. 16. Akahori T., Niinomi M., Ishimizu K., Fukui H., Suzuki A., 2003. Effect of Thermomechanical Treatment on Fatigue Characteristics of Ti-29Nb-13Ta-4.6Zr, J. Japan Institute of Metals 67, 652–660. 17. Niinomi M., Akahori T., Yabunaka T., Fukui H., Suzuki A., 2002. Fretting Fatigue Characteristics of Newly Developed 　-type Titanium Alloy for Biomedical Applications in Air and Simulated Body Environmen, Tetsu-to-Hagane 88, 553–560. 18. Niinomi M., Akahori T., Nakamura S., Fukui H., Suzuki A., 2002. Friction Wear of Surface Oxidized Newly Developed 　-type Titanium Alloy for Biomedical Applications in Simulated Body Environment, Tetsu-to-Hagane 88, 567-574. 19. Akahori T., Niinomi M., Nakai M., Nishimura H., Takei Y., Fukui H., Ogawa M., 2008. Wear and Mechanical Properties, and Cell Viability of Gas-Nitrided 　-type Ti-Nb-Ta-Zr System Alloy for Biomedical Application, Mater. Trans. 40, 166–174. 20. Niinomi M., Akahori T., Katsura S., Yamauchi K., Ogawa M., 2007. Mechanical Characteristics and Microstructure of Drawn Wire of Ti-29Nb-13Ta-4.6Zr for Biomedical Applications, Mater. Sci. and Eng., C 27, 154–161. 21. Nitta K., Watanabe S., Masahashi N., Hosoda H., Hanada S., 2001. Ni-free Ti-NbSn Shape Memory Alloys”, Structural Biomaterials for the 21st Century, eds. M. Niinomi, T. Okabe, E. M. Taleff, D. R. Lesure, and H. E. Lippard, 25–34. 22. Niinomi M., 2003. Recent Research and Development in Titanium Alloys for Biomedical Applications and Healthcare Goods, Science and Technology for Advanced Materials 4, 445–454. 23. Hosoda H., 2004. Superelastic Titanium Alloys for Biomedical Applications”, Front Line of Research and Development of Titanium Alloys, Japan Institute of Metals, 9– 13.

24. Kasuga T., Nogami M., Niinomi M., Hattori T., 2003. Bioactive Calcium Phosphate Invert Glass-Ceramic Coating on 　-type Ti-29Nb-13Ta-4.6Zr, Biomaterials 24, 283–290. 25. Akahori T., Niinomi M., Koyanagi S., Kasuga T., Toda T., Fukui H., Ogawa M., 2006. Aging Treatment and Mechanical Properties of Calcium Phosphate Glassceramic Coated Ti-29Nb-13Ta-4.6Zr for Biomedical Applications, J. Japan Institute of Metals 70, 314–321. 26. Li S. J., Niinomi M., Akahori T., Kasuga T., Yang R., Hao Y. L., 2004. Fatigue Characteristics of Bioactive Glass-ceramic Coated Ti-29Nb-13Ta-4.6Zr for Biomedical Application, Biomaterials 25, 3341–3349. 27. Kim H. M., Miyaji F., Kokubo T., 1997. Effect of Heat Treatment on Apatiteforming Ability of Ti Metal Induced by Alkali Treatment, J. Materials Science: Materials in Medicine 8, 341–347. 28. Akahori T., Niinomi M., Nakai M., Fukuda H., Fukui H., Ogawa M., 2007. Bioactive Ceramic Surface Modification of -Type Ti-Nb-Ta-Zr System Alloy by Alkai Solution Treatment, Mater. Trans. 48, 380–384.

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Chapter 18 Biofunctionalization of Metals Takao Hanawa, Yuta Tanaka, and Harumi Tsutsumi Institute of Biomaterials and Bioengineering, Tokyo Medical and Dental University, Tokyo, Japan

1

Introduction

Metals are typically artificial materials and have no biofunction that leads low attraction of metals as biomaterials, while this viewpoint is shortsighted and reflects misunderstandings. On the other hand, abrupt technological evolution on ceramics and polymers make possible to apply these materials to medical devices during the last three decades; in fact many devices consisting of metals have been substituted by those consisting of ceramics and polymers. In spite of this event, over 70% of implant devices still consist of metals and this share is currently maintained, because of their high strength, toughness, and durability. Metallic biomaterials cannot be replaced with ceramics or polymers at present. In addition, completion of regenerative medicine requires at least another few decades. In other words, artificial materials such as metals will survive as biomaterials in future. When a metallic material is implanted into a human body, immediate reaction occurs between its surface and the living tissues. In other words, immediate reaction at this initial stage straightaway determines and defines a metallic material’s tissue compatibility. Surface modification is a process that changes a material’s surface composition, structure and morphology, leaving the bulk mechanical properties intact. With surface modification, tissue compatibility of surface layer could be improved. Dry-process (using ion beam) and wet-process

(which is performed in aqueous solutions) are conventionally predominant surface modification techniques to accelerate bone formation. In addition, metals with biofunctions have been required in the recent past. For example, stents are placed at stenotic blood vessels for dilatation, and blood compatibility or prevention of adhesion of platelets is necessary. In guide wires and guiding cathetels, lubrication in the blood vessels is important for proper sliding and insertion. If metals are used as sensing devices, the control of cell adhesion is necessary. Infection due to biofilm formation on implant devices must be inhibited. For these purposes, the fundamental property is to control the adsorption of proteins, adsorption of cells, platelets, and bacteria. In this chapter, endeavor for biofunctionalization of metals using some techniques will be introduced.

2

Immobilization of Functional Molecules

The immobilization of biofunctional polymers on a noble metals such as gold is usually conducted by using the bonding –SH or –SS– group (Fig. 1); however, this technique can only be used for noble metals. The adhesion of platelets and adsorption of proteins, peptides, antibodies, and DNA is controlled by modifications of the above technique. On the other hand, poly(ethylene glycol), PEG, is a biofuctional molecule on which adsorption of proteins is inhibited. Therefore, immobilization of PEG to metal surface is an important event to biofunctionalize the metal surface. A class of copolymers based on poly (L-lysine)g-poly (ethylene glycol), PLL-g-PEG, has been found to spontaneously adsorb from aqueous solutions onto TiO2, Si0.4Ti0.6O2, and Nb2O5 to develop bloodcontacting materials and biosensors [1,2]. In another case, TiO2 and Au surfaces are functionalized by the attachment of poly(ethylene glycol)-poly(DLlactic acid), PEG-PLA, copolymeric micelles. The micelle layer can enhance the protein resistance of the surfaces up to 70% [3]. Peptides with terminal cysteine residues were immobilized on maleimide-activated oxides [4-6]. A surface of stainless steel was firstly modified by a silane-coupling agent, SCA, (3mercaptopropyl)trimethoxysilane. The silanized stainless steel, SCA-SS, surface was subsequently activated by argon plasma and then subjected to UV-induced graft polymerization of poly(ethylene glycol)methacrylate, PEGMA. The PEGMA graft-polymerized stainless-steel coupon, PEGMA-g-SCA-SS, with a high graft concentration and, thus, a high PEG content was found to be very effective to prevent the absorption of bovine serum albumin and γ-globulin [7]. These processes require several steps but are effective for immobilization;

OH O=C

O-N- Protein O=C H

S

S Au

Figure 1. Immobilization of biofunctional polymers on a noble metals such as gold is usually conducted by using the bonding –SH or –SS– group.

however, no promising technique for the immobilization of PEG to a metal surface has been so far developed.

3

Immobilization of PEG to Metals with Electrodeposition

No successful one-stage technique for the immobilization of PEG to base metals has ever been developed. In this section, immobilization of PEG modified both terminals or one terminal with amine bases onto titanium surface using electrodeposition will be explained. Both terminals of PEG were terminated with –NH2 (B-PEG; PEG1000 Diamine, NOF Corporation, Japan), and only one terminal was terminated with –NH2 (O-PEG; SUNBRIGHT MEPA-10H, NOF Corporation, Japan). The chemical structures of the PEGs are shown in Fig. 2. The molecular weights of all PEGs were about 1000. These terminated PEGs were dissolved in a 0.3-mol L-1 NaCl solution with a concentration of 2mass%. In the solution, the –NH2 terminal was dissociated and charged as –NH3+. The pH of the solution with BPEG was 11.2, and that of the solution with O-PEG was 11.0. The resultant solution was used as an electrolyte for electrodeposition at 310 K. A commercially pure titanium disk with grade 2 was metallographically polished and ultrasonically rinsed in acetone and deionized water. The titanium disk was fixed in a polytetrafluoroethylene holder that was insulated from the electrolyte except for an opening made for electrodeposition. The cathodic potential was charged from open circuit potential to -0.5V vs. SCE with a sweep rate of 0.1 Vs-1 and maintained at this potential for 300s. During charging, the terminated PEGs were electrically migrated to the titanium cathode and deposited on it as shown in Fig. 2. For comparison, titanium was immersed in the electrolyte containing B-PEG for 2h and 24h without any electric charge at 310 K. After

electrodeposition, specimens were rinsed in deionized water and dried with a stream of nitrogen gas (99.9%).

Terminated with -NH2 H H H O C C O H H

H n

H H H H H H H H C O C C O C C C N nH H H H H H H

－

＋ OH-

O－ O2

Ti

H+ CP Ti

Both terminalterminated PEG

One terminalterminated PEG

PEG without termination

H2

H H H H H H H H H H N C C C O C C O C C C N nH H H H H H H H H H

NH3+

Pt

2 mass% PEG + 0.3 mol L-1 NaCl （pH11） pH11）

+ O－ NH3 O－ + NH3 － O O－ + NH3 O－ O－ PEG

Figure 2. PEG molecules were terminated with amine bases at one terminal or both terminals. Amine bases dissociate and are positively charged in aqueous solution and electrically attracted to titanium surface with cathodic charge, and eventually PEG molecules are immobilized.

Fig. 3 shows the thicknesses of the PEG deposition layers determined by ellipsometry. These thicknesses are measured in air; therefore, the real thickness in solutions is larger than these values. The thickness of the deposition layer, in other words, the amount of deposited PEG, is the largest in this order: 24himmersion B-PEG, electrodeposition of B-PEG for 300s, electrodeposition of O-PEG for 300s, and 2h-immersion B-PEG. This indicated that electrodeposition was more effective than immersion for the deposition of PEG on the titanium surface. However, the PEG layer increased after a 24-h immersion, indicating that the charged terminals of PEG attracted the electrostatically titanium surface that is covered by titanium oxide with a large number of hydroxyl groups. In electrodeposition, the thickness of the B-PEG deposition layer was larger than that of the O-PEG deposition layer. This does not necessarily indicate that more B-PEG than O-PEG was deposited because ellipsometry was conducted in air and the PEG molecules collapsed on titanium.

Thickness, d / nm

5 4 3 2 1 0 B-PEG

O-PEG

B-PEG B-PEG 2h 24h Immersion

300s

Electrodepsition

Figure 3. Thickness of the PEG layer deposited on titanium by electrodeposition and immersion.

The B-PEG has more density per molecule after deposition on titanium surface than the O-PEG because both terminals are attracted to the titanium surface in the B-PEG. Therefore, the apparent thickness of the B-PEG in air was larger than that of the O-PEG.

-NH oriented to titanium NHNH-O Stable bonding

Random direction

OneOne-PEG Electrodeposition

BothBoth-PEG Immersion NH3+

NH3+ NH3+

O-

NH3+

O-

NH3+

NH3+ NH3+

NH3+

O-

ElectroElectrodeposition

NH3+…OUnstable bonding

Immersion

NH3+

NH3+ NH3+

NH3+

O-

BothBoth-PEG Electrodeposition

O-

+ NH3+ NH3

O-

O-

NH

NH

NH

O

O

O O

NH

NH3+

O

NH

NH

NH

NH

NH

NH

NH

NH

NH

O

O

O

O

O

O

O

O

O

Ti

Ti

Ti

Random

BrudhBrudh-shape

U-shape

Figure 4. Schematic model of iimobilized manners of PEG to titanium surface with immersion and electrodeposition.

The bonding manner of PEG to titanium surface is significant to design PEGimmobilized materials, while characterization techniques for the determination of immobilization manner of PEG are a little. Immobilization manner of PEG was characterized using X-ray photoelectron spectroscopy (XPS) with angleresolved technique and glow discharge optical emission spectroscopy (GD-OES). As a result, not only electrodeposition but also immersion led to the immobilization of PEG onto titanium surface. However, more terminated amines combined with titanium oxide as an ionic NH-O by electrodeposition, while more amines randomly existed as NH3+ in the PEG molecule by immersion. Moreover, the difference of amine termination led to different bonding manner, U-shape in PEG terminated both terminals and brush in PEG terminated one terminal. Schematic illustration of immobilization manners of PEG molecules are shown in Fig. 4. Characterization with XPS and GD-OES is useful to determine immobilization mode of PEG to solid surface [8, 9].

(A)

(B)

5 μm

Figure 5. Human blood from a healthy volunteer was drawn into a syringe with 1 mL of 3.8% sodium citrate solution used as an anticoagulant at a ratio of 9 parts blood to 1 part citrate. Plate-rich plasma (PRP), 1 x 105 platelets 　L-1, was obtained from a freshly citrated blood. A 0.25 mol L-1 CaCl2 solution was added to PRP. Ti and PEGelectrodeposited Ti, which was incubated at 310 K in advance, were immersed into PRP at 310 K for 5 min. Thereafter, Ti was rinsed with PBS(-), fixed with 2% glutaraldehyde, dehydrated, and observed through a scanning electron microscope, SEM. Platelet adhesion is inhibited on PEG-electrodeposited Ti surface (A), while platelets adhered on untreated Ti surface and fibrin network is formed on it (B).

The concentrations of hydroxyl groups located inside and on the surface oxide films of a commercially pure titanium, cp-Ti, a type 316L austenitic stainless steel, SS, and a cobalt-chromium-molybdenum alloy, Co-Cr-Mo, were evaluated using X-ray photoelectron spectroscopy, XPS, and a zinc-complex

substitution technique. As a result, the concentration of the active hydroxyl groups on Co-Cr-Mo was significantly larger than those on cp-Ti and SS: The immobilized amount of PEG to Co-Cr-Mo alloy was the largest. The amounts of the PEG layer immobilized on the metals were governed by the concentrations of the active hydroxyl groups on each surface oxide in the case of electrodeposition; it was governed by the relative permittivity of the surface oxide in the case of immersion [10]. PEG-immobilized surface inhibited the adsorption of proteins, adhesion of platelets (Fig. 5) and bacteria, therefore, this electrodeposition technique is useful and universal biofinctionalize metal surfaces.

4

Metal-Polymer Composites

Polymers are widely used as biomaterials because of their high degree of flexibility, biocompatibility, and technologic properties, while the polymers show insufficient strength and long-term durability for some purposes because of their structures. On the other hand, metals have good mechanical properties, especially toughness, and long-term durability. However, the biocompatibility of metals is generally inferior to that of polymers and ceramics because no biofunction is added to the metals during the manufacturing process. If a polymer and a metal could be bonded and used as a composite material, a new material having good biocompatibility and high mechanical strength could be created. The unequivocal relationship between the shear bonding strength and the chemical structure at the bonding interface of a metal-polymer composite through a silane coupling agent (3-(trimethoxysilyl) propyl methacrylate (γ-MPS)) was investigated [11]. As the base materials for the composite, Ti and a segmentated polyurethane (SPU) were employed. The chemical structure of the Ti/γ-MPS/SPU interface is illustrated in Fig. 6. According to glow discharge optical emission spectroscopy (GD-OES), the intensity of S in the γ-MPS layer increased with the increase of the concentrations of the γ-MPS solution and immersion times. In other words, the number of the γ-MPS molecular unit and the thickness of the γ-MPS layer increased with the concentration of the γ-MPS solution and the immersion time. Shear bonding stress of Ti/γ-MPS/SPU interface increased with the increase of the concentration of the γ-MPS solution only in the case of 1-min immersion. On the other hand, the shear bonding stress of the Ti/γ-MPS/SPU interface formed from 1.0 and 2.0% γ-MPS solutions significantly increased with immersion times.

Metal Figure 6. Metal-polymer composite through a silane coupling agent.

If the number of molecular units per Ti surface area is small, each molecular unit keeps a distance from its neighbors and consequently falls down to the Ti surface. Also, -Si-O-Si- bonding network among molecular units is not sufficiently formed. On the other hand, if the number of molecular units per Ti surface area is large, the molecular units will crowd at the interface and stand perpendicular to the Ti surface. As a result, the thickness of the γ-MPS layer will increase. In this case, the -Si-O-Si- bonding network is sufficiently formed. A thick γ-MPS layer would be attributable to the increase in the shear bonding strength because a thick γ-MPS has more molecular units containing S-H groups bonded to SPU. The Ti-SPU composite was fractured leaving the SPU component elements on the fractured surface, determined by XPS. However, more residual SPU on the fractured surface of the Ti-SPU composite with the γ-MPS layer existed than that without a γ-MPS layer. The SPU elements remained on the fractured surface as a result of the presence of the γ-MPS layer. The thicker the γ-MPS layer was, the larger the SPU area fraction on the fractured surface was (Fig. 7). According to the above results, the thickness of the γ-MPS layer is controlled by the concentration of the γ-MPS solution and the immersion time. The shear bonding stress of the Ti/γ-MPS/SPU interface increased with the increase in the

Fractured surface Thin SH γ-MPS

SH

O Si O O

SH Si O Si O O O

SH

O Si O O

SH Si O Si O O O

SH

O Si O O

Si O Si O O O

Fractured surface Thick γ-MPS

SH

SH

SH

SH

Si O Si O Si O Si O O O O O

SH

SH

SH

SH

Si O Si O Si O Si O O O O O

SH

SH

SH

SH

Si O Si O Si O Si O O O O O

Figure 7. Schematic model of the fractured region before and after the shear bonding test in the case of a thin γ-MPS layer (top) and a thick γ-MPS layer (bottom).

thickness of the γ-MPS layer. The Ti-SPU composite was fractured inside the SPU, and the shear bonding stress of the Ti/γ-MPS/SPU interface increased with the increase in the SPU area fraction. The γ-MPS is very useful to improve the bonding strength of the Ti-SPU composite. The factor governing the shear bonding strength between Ti and SPU is the thickness of the γ-MPS layer. This study should lead to enhancements in the creation of metal-polymer composites for artificial organs [11]. On the other hand, the shear bond strength of the Ti/SPU interface increased with the UV irradiation, according to the increase of crosslinkage in SPU. UV irradiation to a Ti-SPU composite is clearly one of the causes governing the shear bond strength of the Ti/SPU interface [12]. Also, active hydroxyl groups on the surface oxide film are clearly one of the causes governing the shear bond strength [13].

5

Conclusions

Metallic materials are widely used in medicine not only for orthopedic implants, but also as cardiovascular devices and for other purposes. Biomaterials are always used in close contact with living tissues. Therefore, interactions between material surfaces and living tissues must be well controlled. Metal surface maybe biofunctionalize by various techniques such as immobilization of biofunctional

molecules and creation of composite with biopolymers. These techniques make it possible to apply metals to a scaffold in tissue engineering.

References 1. Kenausis G.L., Vörös J., Elbert D.L., Huang N., Hofer R., Ruiz-Taylor L., Textor M., Hubbell J.A., Spencer N.D., 2000. Poly(L-lysine)-g-poly(ethylene glycol) layers on metal oxide surfaces: Attachment mechanism and effects of polymer architecture on resistance to protein adsorption, J. Phys. Chem. B104, 3298-3309. 2. Huang N.P., Michel R., Vörös J., Textor M., Hofer R., Rossi A., Elbert D.L., Hubbell J.A., Spencer N.D., 2001. Poly(L-lysine)-g-poly(ethylene glycol) layers on metal oxide surfaces:surface-analytical characterization and resistance to serum and fibrinogen adsorption, Langmuir 17, 489-498. 3. Huang N.P., Csucs G., Emoto K., Nagasaki Y., Kataoka K., Textor M., Spencer N.D., 2002. Covalent attachment of novel poly(ethylene glycol)-poly(DL-lactic acid) copolymeric micelles to TiO2 surfaces, Langmuir 18, 252-258. 4. Xiao S.J., Textor M., Spencer N.D., Sigrist H., 1998. Covalent attachment of celladhesive, (Arg-Gly-Asp)-containing peptides to titanium surfaces, Langmuir 114, 5507-5516. 5. Xiao S.J., Textor M., Spencer N.D., Wieland M., Keller B., Sigrist H., 1997. Immobilization of the cell-adhesive peptide Arg-Gly-Asp-Cys (RGDC) on titanium surfaces by covalent chemical attachment, J. Mater. Sci. Mater. Med. 8, 867-872. 6. Rezania A., Johnson R., Lefkow A.R., Healy K.E., 1999. Bioactivation of metal oxide surfaces. 1. Surface characterization and cell response Langmuir 15, 69316939. 7. Zhang F., Kang E.T., Neoh K.G., Wang P., Tan K.L., 2001. Surface modification of stainless steel by grafting of poly(ethylene glycol) for reduction in protein adsorption, Biomaterials 22, 1541-1548. 8. Tanaka Y., Doi H., Iwasaki Y., Hiromoto S., Yoneyama T., Asami K., Imai H., Hanawa T., 2007. Electrodeposition of amine-terminated poly(ethylene glycol) to titanium surface, Mater. Sci. Eng. C27, 206-212. 9. Tanaka Y., Doi H., Kobayashi E., Yoneyama T., Hanawa T., 2007. Determination of immobilization manner of amine-terminated poly(ethylene glycol) electrodeposited to titanium surface with XPS and GD-OES, Mater. Trans. 48, 287-292. 10. Tanaka Y., Saito H., Tsutsumi Y., Doi H., Imai H., Hanawa T., 2008. Active hydroxyl groups on surface oxide film of titanium, 316L stainless steel, and cobaltchromium-molybdenum alloy and its effect on the Immobilization of poly(ethylene glycol), Mater. Trans. 49, 805-811. 11. Sakamoto H., Doi H., Kobayashi E., Yoneyama T., Suzuki Y., Hanawa T., 2007. Structure and strength at the bonding interface between a titanium-segmentated

polyurethane composite through 3-(trimethoxysilyl) propyl methacrylate for artificial organs. J. Biomed. Mater. Res. 82A, 52-61. 12. Sakamoto H., Hirohashi Y., Doi H., Tsutsumi Y., Suzuki Y., Noda K., Hanawa T., 2008. Effect of UV irradiation on the shear bond strength of titanium with segmented polyurethane through γ-mercapt propyl trimethoxysilane, Dent. Mater. J. 27, 124-132. 13. Sakamoto H., Hirohashi Y., Saito H., Doi H., Tsutsumi Y., Suzuki Y., Noda K., Hanawa T., 2008. Effect of active hydroxyl groups on the interfacial bond strength of titanium with segmented polyurethane through γ-mercapt propyl trimethoxysilane. Dent. Mater. J. 27, 81-92.

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Chapter 19 Research on Biological Characteristic of Silver Nanoparticle Jinglong Tang, Ling Xiong, and Tingfei Xi Center of Medical devices, National Institute for the Control of Pharmaceutical & Biological Products, Beijing, China

1

Introduction

Recent years, silver nanoparticles (SNPs), with stable physiochemical properties have been widely used in the field of medicine, due to their superiority in antibacterial activity over common silver [1, 2]. Numerous studies have demonstrated that nano-materials can possess distinct biological characteristics, with some nano-materials being specifically distributed to the targeted organs [3–5], while some nano-materials are prone to affecting normal cellular activity [6, 7].Did SNPs possess these biological characteristics? It is well known that silver ions (Ag+)can enter the blood circulation, accumulate in tissues and organs, can induce toxicity in the liver and kidney, and even cause human death [8, 9]. SNPs are only 1-2 magnitudes larger than silver ions (silver ionic radius 0.126 nm), so it’s possible that SNPs can cross some barriers between blood and tissue. It is, therefore, possible that SNPs produce similar toxicity to Ag+. One study has revealed that SNPs can move into the blood circulation system via the blood-pulmonary barrier, and can show a systemic distribution [10]. These suggested SNPs could distribute throughout the body. However, the distribution, translocation and accumulation pathways of medically administered SNPs, which do not access the body via the respiratory system, may be different. And the state of SNPs (particle or silver

ion) in the body and how they cross the physical barrier is largely unknown. A greater understanding of the translocation, distribution and accumulation of SNPs in target organs is required. On the other hand, some studies have demonstrated that SNPs could generate cytoxicity [11, 12]. But the cytoxicity mechanism of SNPs is largely unknown. Some researchers suggested that SNPs could get into cells in a short time by cell fusionand endocytosis probably and may affecting normal cellular activity [13]. But this must be validated. In this chaper, we will introduce recent developments of biological characteristic of SNPs according to our group’s research.

2

Distribution, Translocation and Accumulation of Silver Nanoparticles in Rats

Our study was designed to investigate the distribution and accumulation of SNPs in rats with subcutaneous injection. Ninety Wistar female rats (120g±5g) were randomly divided into three groups: control group, SNPs group, and SMPs group. Each group was treated with its corresponding suspension by a subcutaneous injection, at a dose of 62.8 mg/kg silver in a volume of 1 mL. The main organs of the experimental animals were harvested for ultrastructural analysis by transmission electron microscopy (TEM) and for silver content analysis by inductively coupled plasma mass spectrometry (ICP-MS) at 2, 4, 8, 12, 18, and 24 weeks. 0.25% SNPs

0.20%

SMPs

0.15% 0.10% 0.05% 0.00% 2wk

4wk

8wk

12wk

18wk

24wk

Figure 1. Percentage of silver content in different organs to total silver content.

Results indicated that SNPs translocated to the blood circulation and distributed throughout the main organs, especially in the kidney, liver, spleen, brain and lung. SMPs, however, could not invade the blood stream, or organ tissues (Fig. 1, Fig. 2). Ultrastructural observations indicate that those SNPs that had accumulated in organs could enter different kinds of cells, such as renal tubular epithelial cells and hepatic cells in the form of particles (Fig. 3). 6

liver

4

brain

1

spleen lung

2

0

spleen lung kidney liver

kidney

3

brain

silver content (μg)

5

SNPs

SMPs

brain

heart

spleen

lung

kidney

liver

thighbone

uterus and ovary

adrenal gland

Figure 2. Silver particles distribution in different organs at 4th week.

According to the results, we could conclude that SMPs never translocate into the blood, while the injected SNPs do translocate into kidney, liver, spleen, brain and lung via blood circulation, which would explain the significantly different distributions of the two in vivo. Why can SNPs enter the blood, whereas SMPs cannot? We presume that physical size determines the different distributions of the two particles. Blood vessels and lymphatic tissues are plentiful in subcutaneous tissue and dermis, especially capillary vessels. After subcutaneous injection, dispersed SNPs may be phagocytozed into vascular endothelial cells, and then enter the blood by exocytosis. Also the injection may result in local inflammation of the subcutaneous tissues, increasing vasopermeability and

a

b

300nm

100nm

c

d

300nm

300nm

Figure 3. TEM photomicrograph of different organs. SNPs in the renal tubular endothelial cells (a), SNPs in hepatocytes (b), SNPs in spleen lymphocytes (c) and SNPs in normal neurons (d).

contracting the endothelial cells. The scattered SNPs can then enter the blood via 0.5-1.0 μm crevices between the endothelial cells, which are flat, with a width of 10-15 μm and a length of 25-50 μm. But SMPs or agglomerated SNPs could not pass the crevices or be phagocytozed by endothelial cells due to their bigger size [14].

SNPs in vivo may translocate throughout the body by two approaches. Firstly, SNPs dissolved in body fluid may produce systemically distributed Ag+. Secondly, SNPs could interact with some proteins and be distributed evenly in the body through protein metabolism [15]. Although the silver content in kidney, liver, spleen, brain and lung of the SNPs group is significantly higher than those of SMPs group, it is not sure that SNPs could translocate in vivo in the form of particles. Using only ICP-MS it was impossible to distinguish whether the silver in vivo exists as ions or as particles. So the results, that silver content in organs is higher in the SNPs group, may be because SNPs can dissolve more Ag+ due to its surface effect (SNPs are likely to possess a larger specific surface area, a higher Gibbs energy and an incomplete surface structure compared with SMPs at the same dose. These factors together may result in SNPs in vivo dissolving more Ag+ than SMPs). To determine the state of SNPs in vivo, the ultrastructural analysis was conducted by TEM. According to the ultrastructural pictures (Fig. 3). We can conclude that the subcutaneously injected SNPs distribute in the different organs as particles. This outcome authenticates the presumption by Takenaka et al, which maintains that silver is a bio-inert material, unreactive to the oxygen in body fluids, and could not dissolve completely in vivo. In conclusion, after subcutaneous injection in rats, monodisperse SNPs has at most 0.15% that translocated into blood circulation and distributed throughout the body, and accumulated in kidney, liver, spleen, brain, and lung in the form of particle. But SMPs and aggregates made up of SNPs could not translocate into blood circulation. After translocation to the different organs, monodisperse SNPs could enter into some kinds of cells such as renal tubular epithelial cell, hepatocyte, neuron, and disperse in the cytoplasm.

3

Effects of Silver Nanoparticles on L929 Cells In Vitro

Because it is proved that SNPs could enter into cells in experiment in vivo. An in vitro experiment was carried out to investigate the the cytoxicity and its mechanism of SNPs. There are three groups: control group, SNPs group, and SMPs group. L929 cells were cultured with cell culture medium, SNPs dillution of serials concentration (2.5, 5, 10, 25, 50, 100, 250, 500μg/ml) and SMPs dillution of same serials concentration for 24h respectively. Cellular morphology, ultrastructure changes, and mitochondrial function (MTT assay) were assessed.

The results of the MTT assay assay demonstrated that exposure to SNPs for 24 h resulted in concentration-dependent decrease of RGR (related growth rate). It did not produce cytotoxicity up to the concentration of 50μg/ml, but exhibited a signifficant (P 0.05). Statistical analysis using Scheffe’s multiple comparison test showed that the bond strength of 700oC-treatment coatings significantly declined by about 39%. The 600oCtreatment coating had greater bond strength than the as-sprayed coating, although there was not significantly different (p > 0.05). Post-test observations on the stud side indicated that the coatings mainly fractured in the HA-rich layer. The differences in bond strength can be explained in terms of stress relief and microstructural changes. When heat-treated at higher temperatures such as 700oC, the regions near the substrate may result in a larger amount of cracks caused by the large volume shrinkage during crystallization and phase transformation. In turn, it is detrimental to the bonding strength of the coating, although it possesses a higher degree of crystallinity.

3.4 Electrochemical Test The OCP-time plots of the coating samples heat-treated at different temperatures as a function of time along with the as-sprayed coating are shown in Fig. 2. It seems that the as-sprayed coating was in a steady state possibly due to the apatite precipitation, although with a more negative initial potential of -0.47 V.

Figure 2. Open circuit potential-time (top) and typical polarization (bottom) curves of graded coatings before (a) and after heat treatment at 400oC (b), 500oC (c), 600oC (d), and 700oC (e) in deaerated HBSS at 37 oC.

On the contrary, all heat-treated samples except the coating heat-treated at 400oC showed a higher initial OCP, tending to a decreased potential ranging from -0.10 V to -0.25 V after OCP scanning for 5 h. Figure 2 also shows the typical potentiodynamic polarization curves of the plasma-sprayed coatings without and with heat treatment in deaerated HBSS at 37oC. All electrochemical parameters, including corrosion potential (Ecorr), current density (icorr), and polarization resistance (Rp), of coating samples are also compiled in Table 1. Concerning the corrosion potential, there are significant differences (p < 0.05) among all test samples. After heat treatment at 400, 500, 600 and 700oC, corrosion potentials of the heat-treated samples were found to be -598, -518, -464 and -692 mV (vs. SCE), respectively. Our results confirmed that the 600oC-treatment coating exhibited a significantly better corrosion-resistance than all the other coatings by virtue of a more noble corrosion potential, although all polarization curves were characterized by a very similar trend. The corrosion current density and polarization resistance of the samples were determined from the potentiodynamic polarization curves using the Tafel extrapolation method. As for the current density, it was determined that the average values of the heat-treated coatings of between 83 and 153 nA/cm2, which were comparable to that of the as-sprayed coating (142 nA/cm2), were Table 1. Mean and standard derivation values of electrochemical parameters of the assprayed coating and heat-treated graded coatings at different temperatures after electrochemical test. Ecorr (mV)

icorr (nA·cm-2)

Rp (kΩ·cm2)

-576 ± 87a,b,c

142 ± 54a,b

72 ± 16a

400

-598 ± 60a,b

134 ± 35a,b

87 ± 29a

500

-518 ± 46b,c

95 ± 25a,b

162 ± 33b,c

600

-464 ± 58c

83 ± 23a

178 ± 49c

700

-692 ± 77a

153 ± 39b

103 ± 22a,b

Coating As-sprayed

Heat-treated (oC)

Number of samples is at least six in each subgroup. Mean values followed by the same superscript letter in the same column are not significantly different (p > 0.05) according to Scheffe’s post-hoc multiple comparisons.

dependent on treatment temperature, revealing the significant difference (p < 0.05). It is worthwhile to note that heat-treated coatings at 700oC indicated a lower corrosion potential and larger corrosion current density as compared to the 500oC- and 600oC-treated coatings, possibly due to higher temperatureinduced structure changes that led to the penetration of solution into the coating/substrate interface through perpendicular and/or parallel defects. In contrast to the current density, there was an increase in the Rp values by about a factor of two, illustrating that heat treatment endowed the coatings better corrosion resistance except for the treatment at 400oC. The critical factors influencing the corrosion behavior of HA-based coatings are the quality (crystallinity, purity, residual stress and ion substitution in the apatite lattice) and structure (porosity and cracks) [27, 28]. The porosity is a characteristic of plasma sprayed coatings and strongly affects their corrosionresistance. Generally speaking, the corrosion rate increases with increasing porosity of the coatings. The electrolyte infiltrates into the inner portion of the coating through structural imperfections such as pores and cracks or pinholes existing in the coating, and it comes into contact with the deeper portion of the coating [28], causing corrosion. The obtained polarization resistance (Rp) can be used to determine the porosity that corresponds simply to the ratio of the polarization resistance of the uncoated substrates and the coated samples [29]. Using a modified equation, the ratio of Ps/Phs = Rp,hs/Rp,s can be used to represent the change in porosity, where Ps and Phs are the porosity of the assprayed and heat-treated coatings, and Rp,s and Rp,hs are the polarization resistance of the as-sprayed and heat-treated coatings, respectively. Substituting the obtained Rp values into the above-mentioned equation, it is obvious that Ps/Phs is approximately two for the samples heat-treated at temperature greater than 400oC, indicating that there was a positive effect on the reduction of the porosity occurred in those samples. This was because the heat treatment apparently reduced plasma spray-induced layer defects, as described earlier in the morphology change, resulting in heat-treated coatings possessing higher corrosion-resistance. More importantly, the in vitro electrochemical test results indicated that treatment at 600oC had a more beneficial and desired effect on corrosion behavior than the as-sprayed and the other three heat-treated samples at 400, 500 and 700oC from the viewpoints of Rp and corrosion potential.

3.5 In Vitro Drug Release Gentamicin loading onto coatings is a clinically relevant concept in the context of total joint arthroplasty and dental surgery. We used antibiotic-soaked

coatings in an in vitro drug release study. The appearances of the heat-treated coating surfaces were similar to those of as-sprayed coatings, which had wellflattened splats and shiny glassy films and irregularly shaped particles, as shown in Fig. 3a. Randomly distributed pores of different sizes as well as microcracks were also observed. Gentamicin might incorporate into the pore within the plasma sprayed coating. In contrast to the image in Fig. 3a, the gentamicinloaded surface became much smoother and quite uniform, but fractures were visible possibly due to drying (Fig. 3b). After drug release, coating morphologies are similar to those without drug loading (Figs. 3c, d).

Figure 3. Scanning electron micrographs: as-sprayed coating (a) with gentamicin loading (b), the release of gentamicin for as-sprayed coating (c) and 600oC-treatment coating (d).

Gentamicin release profiles from the two coatings in PBS as a function of time are shown in Fig. 4. Generally, the release curves can be separated into an initial fast release, followed by a slow release pattern. The fast release is mainly caused by the dissolution of the drug that is physically adsorbed on coating implants, and the slow release may be attributed to chemically adsorbed drugs. It can be seen that during the fast release, the rate of the untreated coatings was larger than that of the heat-treated coatings. It is noted that, at the initial 1 h, the two coatings released almost entirely gentamicin. Osteoconductive coatings have the potential to serve as drug carriers to prevent infection in the setting of total joint arthroplasty and dental therapy [21].

Figure 4. The release profiles of gentamicin as a function of time from coatings without (a) and with heat treatment at 600oC (b). The insert is the short-term release profiles.

4

Conclusions

The interfacial bond strength and dissolution behavior of plasma sprayed coatings were related on their microstructural and chemical inhomogeneity. After heat treatment at 500–700oC, the graded coatings had an enhanced crystallinity by a factor of three but a temperature-sensitive bond strength. Among the treatment temperatures, 600oC seems to endow the graded coating with better corrosion-resistance with an increased polarization resistance value by approximately two times as compared to the as-sprayed samples. Improved corrosion resistance was due to a coating surface modification with higher degree of crystallinity and less dissoluble non-apatite phases (TCP), as well as a reduction in coating defects when plasma-sprayed coatings were subjected to post-deposition heat treatment. Gentamicin loading might provide effective release at early time points (up to 1 h).

Acknowledgments The author is indebted to Mr. S.D. Hsieh for DC plasma spraying facilities. The work was supported by National Science Council of the Republic of China under the contract No. NSC 95-2314-B-040-011-MY2.

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13. Ding S.J., Huang T.H., Kao C.T., 2003. Immersion behavior of plasma-sprayed modified hydroxyapatite coatings after heat treatment, Surf. Coat. Tech. 165, 248– 257. 14. Weng J., Cal T., Chen J., Zhang X. 1995. Significance of water promoting amorphous to crystalline conversion of apatite in plasma sprayed coatings, J. Mater. Sci. Lett. 14, 211–213. 15. Chen J., Tong W., Cao Y., Feng J., Zhang X., 1997. Effect of atmosphere on phase transformation in plasma-sprayed hydroxyapatite coatings during heat treatment, J. Biomed. Mater. Res. 34, 15–20. 16. Chen C.C., Ding S.J., 2006. Effect of heat treatment on characteristics of plasma sprayed hydroxyapatite coatings, Mater. Transact. 47, 935–940. 17. Li L.C., Deng J., Stephens D., 2002. Polyanhydride implant for antibiotic deliveryfrom the bench to the clinic, Adv. Drug. Delivery Rev. 54, 963–986. 18. Dion A., Berno B., Hall G., Filiaggi M.J., 2005. The effect of processing on the structural characteristics of vancomycin-loaded amorphous calcium phosphate matrices, Biomaterials 26, 4486–4494. 19. Akashi A., Matsuya Y., Unemori M., Akamine A., 2001. Release profile of antimicrobial agents from α-tricalcium phosphate cement, Biomaterials 22, 2713– 2717. 20. Schofield S.C., Berno B., Langman M., Hall G., Filiaggi M.J., 2006. Gelled calcium polyphosphate matrices delay antibiotic release, J. Dent. Res. 85, 643–647. 21. Radin S., Campbellt J.T., Ducheyne P., Cuckler J.M., 1997. Calcium phosphate ceramic coatings as carriers of vancomvcin, Biomaterials 18, 777–782. 22. Pourbaix M., 1984. Electrochemical corrosion of metallic biomaterials, Biomaterials 5, 122–134. 23. Arends J., Christoffersen J., Christoffersen M.R., Eckert H., Fowler B.O., Heughebaert J.C., Nancollas G.H., Yesinowski J.P., Zawacki S.J. 1987. A calcium hydroxyapatite precipitated from an aqueous solution, J. Crystal. Growth. 84, 515– 532. 24. Wen J., Leng Y., Chen J., Zhang C., 2000. Chemical gradient in plasma-sprayed HA coatings, Biomaterials 21, 1339–1343. 25. Sergo V., Sbaizero O., Clarke D.R., 1997. Mechanical and chemical consequences of the residual sresses in plasma sprayed hydroxyapatite coatings, Biomaterials 18, 477–482. 26. Milosevski M., Bossert J., Milosevski D., Gruevska N., 1999. Preparation and properties of dense and porous calcium phosphate, Ceram. Int. 25, 693–696. 27. Sridhar T.M., Kamachi Mudali U., Subbaiyan M., 2003. Preparation and characterisation of electrophoretically deposited hydroxyapatite coatings on type 316L stainless steel, Corros. Sci. 45, 237–252. 28. Cao Y., Weng J., Chen J., Feng J., Yang Z., Zhang X., 1996. Water vapour-treated hydroxyapatite coatings after plasma spraying and their characteristics, Biomaterials 17, 419–424.

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Chapter 27 Design of Supramolecular Polyrotaxanes for DNA Delivery Atsushi Yamashita and Nobuhiko Yui School of Materials Science, Japan Advanced Institute of Science and Technology, Ishikawa, Japan

1

Introduction

Gene delivery using virus-based vector systems is risky due to genomic integration concerns, and have low usefulness due to single use limitations [1]. Therefore, many researchers has paid their attention to invent the non-viral gene vector systems using cationic polymers and cationic liposome. Cationic polymers can form a polyplex in some buffer conditions through an electrostatic interaction with phosphate anions of pDNA. However, low transfection efficiency still remains a bottleneck, preventing the use of cationic polymers in clinical applications [2]. In order to achieve the effective gene delivery, the cationic polymers are required to overcome the following issues: (1) formation of a polyplex, which is stable against counter polyanion, (2) protection of DNA from nuclease attack in blood circulation and cytosol, (3) cellular uptake via endocytosis, (4) endosomal escape, (5) nuclear localization, and (6) DNA release and the interaction with gene expression machinery such as RNA polymerase in cell nucleus.[2] Biocompatibility is also a very important factor for cationic polymers in clinical applications. Linear polyethyleneimine (LPEI) is one of the effective gene carriers, because the secondary amines of LPEI are protonated under weakly acidic conditions in endosome/lysosome, and the buffering effects are thought to induce osmotic swelling of endosomal and lysosomal interior, resulting in the rupture of the endosome/lysosome and the subsequent DNA release into cytoplasm [3]. Many efforts have been made to

develop gene vectors using cationic polymers. It was assumed that the polyplex stability in the blood circulation as well as the pDNA release from the polyplex in target cells was an important factor to enhance the transfection activity of non-viral gene carriers. In addition, the low transfection efficiency of non-viral gene carrier principally arises from low transcription efficiency in nuclei rather than an intracellular trafficking [4,5]. Therefore, a novel strategy is needed to make a multi-functional gene carrier, which forms a stable polyplex in the extracellular and the endosome/lysosome interior, while induces the efficient DNA release from the polyplexes after their movement around the nuclear compartment [6-8]. However, it was expected that overcoming the various barriers on transfection was difficult in conventional polymers (e.g. polyethyleneimine, poly(L-lysine)). Therefore, strategies of using some supermolecules such as polyion micelle [9], multi-functional envelope type nano devise (MEND) [10], and polyrotaxane [11-13] have been examined as a multifunctional gene carrier. Supramolecular structures of polyion micelle and MEND have been revealed as effective gene carriers. On the other hand, the usefulness of polyrotaxane structure for gene carrier is not well known, although the polyrotaxanes provided many novel properties such as mobile motion of cyclic molecule threading onto linear polymer [14], multivalent interaction [15,16] and stimuli responsive dissociation. [17] In this review, we will introduce recent studies on supermolecule-based gene carrier using biodegradable polyrotaxanes after introduction of PEI-based gene carriers.

2

Polyethyleneimine (PEI)-Based Gene Carriers

Cationic polymers have been frequently used for DNA encapsulation and delivery included polylysine (PLL), chitosan, polyamidoamine dendrimers, and polyethylenimine (PEI). Particularly, linear PEI (LPEI) and fractured polyamidoamine dendrimers showed high transfections in various animal and cultured cells due to their proton sponge effect [3]. Although these polycations have been used as transfection agents, most of these cationic materials have relatively high cytotoxicity and difficulties in formulation to use as they are [18]. In addition, low transfection efficiency still remains a bottleneck, preventing the use of polycations in clinical applications. Therefore, many researchers have paid their attention to increasing the biocompatibility and the transfection activity of polycations. An inclusion complex (polypseudorotaxane) consisting of a linear polyethylenimine with Mw of 22,000 (LPEI) and γ-cyclodextrins (γCDs) (LPEI/γ-CD) has been examined as a gene carrier, which can improve the biocompatibility of polycation (LPEI) [19]. The LPEI22k/γ-CD significantly increased cell viability even at high N/P ratio, and the polyplex maintained the high transfection activity of LPEI, because the threading γ-CDs of LPEI22k/γCD could decrease the charge density of LPEI22k. The decreased cation charge

density of polyplex by γ-CD threading may prevent from the cell membrane damages, resulting in high cell viability even at higher N/P ratio. In addition, decreasing the charge density of LPEI22k was achieved without any covalent bounds, resulting in effective proton sponge effect of LPEI. However, LPEI/γCD could not improve the transfection activity of LPEI. On the other hand, some research groups introduced the disulfide linkages into the main chain of PEI (biodegradable PEI), because the cleavage of disulfide linkages by cytosolic glutathione decreases the molecular weight of LPEI to induce the pDNA release in the cytosolic milieu and/or nuclei, and to improve the cell viability. However, biodegradable PEI was not succeeded to enhance the transfection activity. In addition pDNA release from biodegradable PEI required high concentration of glutathione, because of an excess amount of disulfide linkages can result in over stabilization of the polyplex against the exchange reaction of counter polyanion.

3

Biodegradable Polyrotaxanes

In order to overcome the over stabilization due to excess amount of disulfide linkages, we focused on a biodegradable polyrotaxane structure. The biodegradable polyrotaxane, in which dimethylaminoethyl-modified αcyclodextrins (DMAE-α-CDs) are threaded onto a poly(ethylene glycol) (PEG) (Mw = 4000) chain capped with benzyloxycarbonyl tyrosine (Z-L-Tyr) via disulfide linkages that exist only at both termini of the PEG chain (DMAE-SSPRX), was examined as a multi-functional gene carrier[12,13]. It is expected that the polyrotaxane will enhance the pDNA release in the cytosol, because only two disulfide linkages can avoid the over stabilization. The cleavage will lead to triggering pDNA release via the dissociation of the inclusion complex between α-CDs and PEG (Fig. 1). In addition, tertiary amino groups, which were introduced into the threading α-CDs of DMAE-SS-PRX, may induce polyplex formation and endosomal escape due to low pKa value of DMAE groups. In deed, a polyplex was formed by mixing very small amounts of the DMAE-SS-PRX with pDNA (Fig. 2A). On the other hand, the DMAE-α-CD, one of the building blocks of the DMAE-SS-PRX, did not form any polyplexes under the N/P ratio of 10 (Fig. 2B). These results indicate that the polycationic nature of the DMAE-SS-PRX contributes to polyplex formation. In vitro pDNA release experiments in the presence of 10 mM dithiothreitol (DTT) confirmed that pDNA was released from the DMAE-SS-PRX polyplex (Fig. 2C). On the other hand, the polyplex of DMAE-introduced non-degradable polyrotaxane (DMAE-PRX), which has no disulfide linkages, was stable in the same condition (Fig. 2D). In addition, rapid endosomal escape of DMAE-SS-PRX polyplex and pDNA delivery to the nucleus were observed by confocal laser scanning microscopy. The biocompatibility and transfection activity of DMAESS-PRX were significantly higher than non-degradable PRX (DMAE-PRX)

Figure 1. Polyplex formation of DMAE-SS-PRX and its pDNA release by supramolecular dissociation.

Figure 2. Agarose gel electrophoretic images of pDNA polyplex with DMAE-SS-PRX (A) and DMAE-α-CD (B). pDNA release from the polyplex with the DMAE-SS-PRX (C) and the DMAE-PRX (D) (N/P = 5) in the presence of 10 mM DTT and dextran sulfate (Mn = 25 000) as a counter polyanion [12]. [Permission from American Chemical Society]

(Fig. 3 A, B). These results strongly support our strategy that the DMAE-SSPRX overcomes over stabilization of the polyplex, because only two disulfide linkages can avoid the over stabilization, resulting in high biocompatibility and high transfection activity.

Figure 3. Cells viability of the DMAE-SS-PRX, the DMAE-PRX and LPEI polyplexes evaluated by MTT assay (n=4) (A). Transfection activities with the DMAE-SSPRX, DMAE-PRX, and LPEI measured by luciferase assay (n = 7-9) (B) [12]. [Permission from American Chemical Society]

Recently, we prepared biocleavable polyrotaxanes with different numbers of threading α-CD and amino (DMAE) groups to optimize the molecular design of DMAE-SS-PRX. [20] The DMAE-SS-PRXs were prepared in the three steps. PEG-BA-SS was added to α-CD saturated water, resulting in forming the pseudopolyrotaxane. Two different mixture ratios of α-CD and PEG-SS-BA were employed to adjust roughly the number of α-CDs in the SS-PRX [16]. To prepare the SS-PRX, the respective pseudopolyrotaxanes were capped with ZTyr using a condensation agent in methanol, followed by purification to obtain α18-SS-PRX and α29-SS-PRX. The number of DMAE groups per α-CD molecule could be also controlled by varying mixture ratio of CDI and DMEDA to hydroxyl groups in α-CD to be 0.8-6.2, which was determined by comparing peak integrations for both C(1) protons in α-CD around 4.8 ppm and methyl protons on nitrogen around 2.0 ppm in the NMR spectrum measured in D2O containing NaOD. The PRXs without SS linkages (DMAE-PRXs) were prepared as a reference in a similar manner by employing ethylenediamine in place of cystamine for DMAE-SS-PRXs. The structural parameters of the obtained DMAE-(SS)-PRXs are summarized in Table 1. Figure 4 shows the pDNA release from these polyplexes in the reductive condition. To examine the

Table 1. Synthetic results for the DMAE-SS-PRXs and DMAE-PRXs. Sample code

# of α-CD[a]

# of DMAE /PRX[d]

Total Mw[e]

14DMAE-α18-SS-PRX

18

14

25,000

29DMAE-α18-SS-PRX

18

29

27,000

76DMAE-α18-SS-PRX

18

76

32,000

99DMAE-α18-SS-PRX

18

99

35,000

35DMAE-α29-SS-PRX

29

35

37,000

52DMAE-α29-SS-PRX

29

52

39,000

110DMAE-α29-SS-PRX

29

110

46,000

180DMAE-α29-SS-PRX

29

180

53,000

14DMAE-α16-PRX

16

14

23,000

29DMAE-α16-PRX

16

29

24,000

61DMAE-α16-PRX

16

61

28,000

96DMAE-α16-PRX

16

96

32,000

29DMAE-α29-PRX

29

29

36,000

71DMAE-α29-PRX

29

71

41,000

97DMAE-α29-PRX

29

97

44,000

164DMAE-α29-PRX

29

164

52,000

Mw of PEG = 3,500 – 4,500. [a] [c] The averaged threading number of α-CDs in a PRX and the averaged number of DMAE groups in an α-CD were determined from the peak integration of 1H NMR spectra. [b] The threading % of α-CDs in a PRX was calculated from [a]. [d] [e] The number of DMAE groups in a PRX and the molecular weight of DMAE-SS-PRXs were calculated from [a] and [b]. The numerical relative errors of the number of α-CDs in DMAE-α18-SS-PRXs, DMAE-α29-SS-PRXs, DMAE-α16-PRXs and DMAE-α29-PRXs, which are caused by the Mw distribution of PEG, are 32%, 20%, 36% and 20%, respectively.

ability of pDNA release, a dextran sulfate competitive displacement assay was performed in the presence of 10 mM dithiothreitol (DTT). The DTT concentration of 10 mM was comparable to intracellular reductive condition (approximately 10 mM glutathione). In the case of DMAE-SS-PRXs, free pDNA bands were observed at a lower concentration of dextran sulfate, compared with the non-degradable PRXs. These results indicate that the pDNA release from DMAE-SS-PRXs polyplexes was induced by the SS cleavage,

followed by the dissociation of the supramolecular structure of PRXs. In addition, the pDNA release from DMAE-SS-PRXs depended on the numbers of α-CD and DMAE groups. The pDNA release from 14, 29, 76 and 99DMAE18α-SS-PRXs was induced by the adding of dextran sulfate, in which concentrations were 0, 8.5, 35 and 50 μg mL-1, respectively.

Figure 4. Agarose gel electrophoretic images of the released pDNA from polyplexes (N/P = 5) of LPEI (A), DMAE-α18-SS-PRXs (B-E), DMAE-α29-SS-PRXs (F-I), DMAE-α16-PRXs (J-M) and DMAE-α29-PRXs (N-Q) in the reductive condition. The reductive condition was prepared by adding DTT (10 mM) and incubated for 60 min at 37 °C. Each sample was mixed with a solution of dextran sulfate (Mw = 25,000) and incubated at room temperature for 20 min, followed by agarose gel electrophoresis. Concentrations of dextran sulfate were 0, 8.5, 17, 35, 50, 70 and 100 μg mL-1.

Most of the DMAE-SS-PRX polyplexes were found to release the pDNA only in the presence of both 10 mM DTT and counter polyanion, except for 14DMAE-α18-SS-PRX, which released pDNA in the absence of dextran sulfate once DTT had been added to the polyplex solution. From the results of the ethidium bromide displacement assay (Fig. 5), it is considered that the relatively easy pDNA release from the DMAE-SS-PRXs with the low number of amino groups would be due to loosely packed polyplex. The transfection activity of DMAE-SS-PRXs seems to be related with the ability of pDNA

Figure 5. Relative fluorescence intensity (Fr) of ethidium bromide (EtBr) in solution with pDNA and DMAE-α18-SS-PRXs or DMAE-α29-SS-PRXs at various N/P ratios.

release from these polyplexes, which dependent on the numbers of α-CD and/or DMAE groups in the polyrotaxane carrier (Fig. 6). The DMAE-SS-PRXs with low numbers of α-CD and amino groups exhibited high transfection activity, except for the 14DMAE-α18-SS-PRX, which has the lowest number of amino groups. A 29DMAE-α18-SS-PRX, in which the numbers of α-CD and DMAE groups are 18 and 29 molecules, revealed high transfection activity compared with other DMAE-SS-PRXs. These results might be due to the polyplex stability against counter polyanion and the appropriate timing of pDNA release in the cytosol, which are controlled by the numbers of α-CD and amino groups in the DMAE-SS-PRXs. However, it is difficult to explain the reason why the transfection activity of 14DMAE-α18-SS-PRXs was significantly lower than those for the other DMAE-SS-PRXs, although an effective pDNA release was observed in the reductive condition (Fig. 4B).

Figure 6. Transfection activity in the polyplexes of DMAE-α18-SS-PRXs (A), DMAEα29-SS-PRXs (B) and DMAE-α16-PRXs (C). Luciferase activity in the NIH/3T3 cells was measured 48 h after adding the polyplexes. Results were expressed as relative light units (RLU) per mg of cell protein.

We expected that the low transfection activity of 14DMAE-α18-SS-PRX might be due to a premature pDNA release in the cytosol, presumably causing its degradation by cytosolic nucleases. In deed, the results of confocal laser scanning microscopic (CLSM) observation suggested that the specific result of 14DMAE-α18-SS-PRX might be due to a premature pDNA release from the most dissociative 14DMAE-α18-SS-PRX polyplex in the cytosol (Fig. 7). Therefore, it is suggested that the transfection activity would be related to an appropriate timing for pDNA release.

Figure 7. Confocal laser scanning microscopic (CLSM) images of NIH/3T3 cells 3 h after the transfection at an N/P ratio of 5; polyplexes of 14DMAE-α18-SS-PRX (A), 29DMAE-α18-SS-PRX (B), and 29DMAE-α16-PRX (C). The endosome/lysosome (green) was stained by LysoSensor DND-189 and the blue fluorescence indicates a Hoechst 33258-stained nucleus. Rhodamine-labeled pDNA shows red fluorescence.

4

Conclusion

A stable polyplex with positively charged surface was formed by mixing very small amounts of the DMAE-SS-PRX with pDNA. The pDNA release from the polyplex occurred through disulfide cleavage of the DMAE-SS-PRX and the subsequent interexchange with polyanions. Rapid endosomal escape and pDNA delivery to the nucleus were achieved by the DMAE-SS-PRX polyplex. In addition, we prepared various types of PRXs to examine how the numbers of αCD and amino groups affect polyplex formation, polyplex stability against counter polyanion, pDNA release in the reductive condition and transfection activity. All of the DMAE-SS-PRXs showed polyplex formation at low N/P ratio and a high stability of the polyplex against the counter polyanion. The pDNA release in the reductive condition increased with a decrease in the numbers of α-CDs and amino groups, resulting in high transfection except for the 14DMAE-α18-SS-PRX. These results suggest that the appropriate timing of pDNA release in the cytosol, which is controlled by the numbers of α-CD and amino groups in the DMAE-SS-PRXs, is an important parameter to enhance

the transfection activity of cationic polymers. These findings are believed to be an important knowledge of designing non-viral vectors and developing new gene therapy using vectors.

Acknowledgments We acknowledge Prof. Hideyoshi Harashima, Dr. Hidetaka Akita, Dr. Yuma Yamada (Hokkaido University, Japan), Prof. Kentaro Kogure (Kyoto Pharmaceutical University, Japan), Prof. Atsushi Maruyama, Dr. Arihiro Kano (Kyushu University, Japan), Dr. Ryo Katoono, Mr. Daizo Kanda, Mr. Atsuto Yoshihiro (Japan Advanced Institute of Science and Technology, Japan), Ass. Prof. Tooru Ooya (Kobe University, Japan), Dr. Hak Soo Choi (Harvard Medical School, Boston), Dr. Motoichi Kurisawa (Institute of Bioengineering and Nanotechnology, Singapore) for their collaboration. This work was supported by the Ministry of Education, Culture, Sports, Science and Technology of Japan.

References 1. Wolff J.A., 2002. The “grand” problem of synthetic delivery, Nat. Biotechnol. 20, 768-769. 2. De Smedt S.C., Demeester J., Hennink W. E., 2000. Cationic Polymer Based Gene Delivery Systems, Pham. Res. 17, 113-126. 3. Boussif O., Lezoualc’h F., Zanta M.A., Mergny M. D., Scherman D., Demeneix B., Behr J.P., 1995. A versatile vector for gene and oligonucleotide transfer into cells in culture and in vivo: Polyethylenimine, Proc. Natl. Acad. Sci. USA 92, 7297-7301. 4. Hama S., Akita H., Ito R., Mizuguchi H., Hayakawa T., Harashima H., 2006. Quantitative comparison of intracellular trafficking and nuclear transcription between adenoviral and lipoplex systems, Mol. Ther. 13, 786-794. 5. Hama S., Akita H., Iida S., Mizuguchi H., Harashima H., 2007. Quantitative and mechanism-based investigation of post-nuclear delivery events between adenovirus and lipoplex, Nucleic Acids Res. 35, 1533-1543. 6. Pichon C., LeCam E., Gue´rin B., Coulaud D., Delain E., Midoux P., 2002. Poly[Lys-(AEDTP)]: A Cationic Polymer That Allows Dissociation of pDNA/Cationic Polymer Complexes in a Reductive Medium and Enhances Polyfection, Bioconjugate Chem. 13, 76-82. 7. Gosselin M.A., Guo W., Lee R.J., 2001. Efficient gene transfer using reversibly cross-linked low molecular weight polyethylenimine, Bioconjugate Chem. 12, 989994. 8. Neu M., Germershaus O., Mao S., Voigt K.H., Behe M., Kissel T., 2007. Crosslinked nanocarriers based upon poly(ethylene imine) for systemic plasmid delivery: In vitro characterization and in vivo studies in mice, J. Control. Release 118, 370-380.

9. Miyata K., Kakizawa Y., Nishiyama N., Harada A., Yamasaki Y., Koyama H., Kataoka K., 2004. Block catiomer polyplexes with regulated densities of charge and disulfide cross-linking directed to enhance gene expression, J. Am. Chem. Soc. 126, 2355-2361. 10. Kogure K., Akita H., Yamada Y., Harashima H., 2008. Multifunctional envelopetype nano device (MEND) as a non-viral gene delivery system, Adv. Drug Delivery Reviews 60, 559–571. 11. Harada A., Li J., Kamachi M., 1992. The molecular necklace: a rotaxane containing many threaded α-cyclodextrins, Nature 356, 325-328. 12. Ooya T., Choi H.S., Yamashita A., Yui N., Sugaya Y., Kano A., Maruyama A., Akita H., Kogure K., Harashima H., 2006. Biocleavable polyrotaxane-lasmid DNA polyplex for enhanced gene delivery, J. Am. Chem. Soc. 128, 3852-3853. 13. Yamashita A., Yui N., Ooya T., Kano A., Maruyama A., Akita H., Kogure K., Harashima H., 2006. Synthesis of a biocleavable polyrotaxane-plasmid DNA (pDNA) polyplex and its use for the rapid nonviral delivery of pDNA to cell nuclei, Nature Protocols 1, 2861-2869. 14. Yui N., Ooya T., 2006. Molecular mobility of interlocked structures exploiting new functions of advanced biomaterials., Chemistry 12, 6730-6737. 15. Ooya T., Utsunomiya H., Eguchi M., Yui N., 2005. Rapid Binding of Concanavalin A and Maltose-Polyrotaxane Conjugates Due to Mobile Motion of α-Cyclodextrins Threaded onto a Poly(ethylene glycol), Bioconjugate Chem. 16, 62-69. 16. Ooya T., Eguchi M., Yui N., 2003. Supramolecular Design for Multivalent Interaction: Maltose Mobility along, J. Am. Chem. Soc. 125, 13016-13017. 17. Ooya T., Yui N., 1999. Synthesis of theophylline–polyrotaxane conjugates and their drug release via supramolecular dissociation, J. Control. Release 58, 251-269. 18. Moghimi S. M., Symonds P., Murray J. C., Hunter A. C., Debska G., Szewczyk A., 2005. A two-stage poly(ethyleneimine)-mediated cytotoxicity: implications for gene transfer/therapy, Mol. Ther. 11, 990-995. 19. Yamashita A., Choi H.S., Ooya T., Yui N., Akita H., Kogure K., Harashima H., 2006. Improved Cell Viability of Linear Polyethylenimine through γ-Cyclodextrin Inclusion for Effective Gene Delivery, Chembiochem. 7, 297-302. 20. Yamashita A., Yui N., Ooya T., Kano A., Maruyama A., Akita H., Kogure K., Harashima H., 2008. Synthesis of a biocleavable polyrotaxane-plasmid DNA (pDNA) polyplex and its use for the rapid nonviral delivery of pDNA to cell nuclei, J. Contrl. Release in press.

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PART VII

Biomechanics

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Chapter 28 Biomechanical Aspects of Natural Articular Cartilage and Regenerated Cartilage Teruo Murakami1, Nobuo Sakai1, Yoshinori Sawae1, Itaru Ishikawa2, Natsuko Hosoda3, Emiko Suzuki3and Jun Honda4 1. Department of Mechanical Engineering, Kyushu University, Fukuoka, Japan 2. National Institute of Health Sciences, Tokyo, Japan 3. Graduate School of Systems Life Sciences, Kyushu University, Fukuoka, Japan 4. Graduate School of Engineering, Kyushu University, Fukuoka, Japan

1

Introduction

The natural synovial joints have excellent load-carrying capacity and tribological performance with very low friction and low wear even under high load conditions for various daily activities. However, the deficiency of lubricating ability and the deterioration in load-carrying capacity in synovial joints particularly for aged people appear to induce the initiation of osteoarthritis. To maintain the excellent performance, the adaptive multimode lubrication mechanism [1,2] should operate under various activities in which elastohydrodynamic lubrication, weeping lubrication, hydration lubrication [3,4], biphasic lubrication [5,6], adsorbed film and/or gel film lubrication can become synergistically effective depending on the severity of operating conditions. In natural joints subjected to severe loading and rubbing, a local direct contact between rubbing cartilage surfaces occurs, and thus the adsorbed film is removed and the underlying gel layer also is damaged. It is important to repair the damaged surface layer by not only synovial fluid containing phospholipids, proteins, glycoproteins, hyaluronic acid etc. but also a supply of proteoglycan

from extracellular matrix produced by chondrocytes [7]. For supply of proteoglycans including lubricin to superficial zone and/or surface gel film, the chondrocytes in surface zone appear to play an important role. The chondrocyte and articular cartilage adapt to changing mechanical environments, but the detailed adaptive and/or restorative process has not yet been clarified. Furthermore, for tissue-engineered cartilage, the appropriate mechanical stimulation is expected to enhance the metabolic activity of chondrocytes [8]. To study biomechanical aspects in articular cartilage and regenerated cartilage, in situ observation of compressive behaviors of articular cartilage including chondrocytes was conducted [9-11]. Articular cartilage has high quantity of water of 70 to 80%. Therefore, the application of biphasic theory [12] is required to understand the complicated time-dependent mechanical behaviors accompanied by fluid pressurization, and thus the corresponding finite element analyses were conducted. Furthermore, the effectiveness of mechanical stimulation on chondrocyte-agarose constructs during culture tests was examined.

2

In Situ Observation of Compressive Behaviors of Articular Cartilage Including Chondrocytes

2.1 Materials and Methods The articular cartilage has a biphasic viscoelastic property based on high water content up to 80%. In this study, therefore, the time-dependent deformation of compressed articular cartilage was observed under unconfined compression in the compressive apparatus shown in Fig. 1 [9-11] located on the stage of confocal laser scanning microscope (CLSM). The articular cartilage specimens are prepared from the intact femoral condyle of porcine knee joints. The cylindrical specimen of 3 mm diameter was prepared by the punch. Then, the cylinder was sectioned in half with a scalpel. The semi-cylindrical specimen includes the subchondral bone of about 0.5 mm thickness. In order to observe the whole morphology and to visualize the cells in compression tests, the live chondrocytes were stained with calcein-AM, at 1 μL/mL and 37°C for 30 min. The compression tests of articular cartilage were carried out by means of the newly developed compression apparatus (Fig. 1) [9-11] with high precision within 0.2 μm for position control, which was allocated on the stage of CLSM.

Cartilage Compression 3mm 2～3mm Figure 1. Compressive apparatus on CLSM stage.

The compression speed can be adjusted from 1 μm/s to 4 mm/s by feed-back control of a DC servo-motor. In most of compressive tests, 10 to 15% total compressive strain was applied by moving impermeable alumina ceramic plate at constant speed. Based on these visualized images, the time-dependent and depth-dependent changes in local strain of articular cartilage were evaluated.

2.2 Experimental Results Figure 2 shows an example of change in compressive stress for articular cartilage in unconfined compression test at high compressive speed of 2 mm/s at definite total deformation. The occurrence of peak stress and stress relaxation are shown. The peak stress of about 2 MPa is supported by both pressurization in fluid phase and elastic stress in solid phase in biphasic articular cartilage. During the stress relaxation process, the fluid gradually flows according to pressure gradient and then finally the pressurization diminishes at equilibrium.

Stress MPa

2.5

Displacement

0.20

2.0

0.15

1.5 1.0

0.25

Peak Stress

Stress

0.10 0.05

0.5 0.2

0.4

0.6

0.8

0 1.0

Displacement mm

0.30

3.0

Time Timess Figure 2. Changes in stress and displacement for unconfined compression test of articular cartilage.

These phenomena appear to have some relation to changes in local strain. On the basis of evaluation of changes in local strain of articular cartilage in visualization test and biphasic finite element analyses, the time-dependence and depth-dependence for mechanical behaviors are discussed. The local strain in cartilage specimen in compression test is estimated by calculating the changes of the distance before and after compression in perpendicular direction to the cartilage surface between the definite stained chondrocytes as reported in the previous paper [9-11] in Fig. 3 as follows; Local strain ε = (a – b) / a Before compression

a

Immediately after compression

b

(a)

Compression

(b)

Equilibrium state

c

(c)

Figure 3. Fluorescence images for estimation of local strain in articular cartilage: (a) before compression, (b) immediately after compression, (c) at equilibrium.

The changes in local strain estimated from the above equation are depicted in Fig. 4. The hollow symbols correspond to the local strain immediately after compression and the solid symbols indicate equilibrium strain. Thus, the local strain of biphasic cartilage exhibited remarkable time-dependent behaviors. The deformation of the deeper zone was clearly recovered probably accompanied by the flow of fluid into the middle and deep zone. On the contrary, the surface zone was largely compressed than average strain during stress relaxation in unconfined compression. The large strain at equilibrium indicates that the excessive deformation of the surface zone can be prevented by fluid pressurization for immediately after compression and the elastic modulus of solid phase at equilibrium has clear depth-dependency, i.e., the elastic modulus of solid phase at surface zone is very lower than the deep zone.

S urface Surface

S ubchondralbone bone Subchondral

0.5 Im m ediatelyafter aftercompression com pression Immediately Equilibrium E quilibrium

L Local o c a l sstrain tra in

0.4 0.3 0.2 0.1 0 -0.1

0

0.2

0.4

0.6

0.8

1.0

Relativeposition position R elative

Figure 4. Changes in local strain in unconfined compression test [9].

2.3 Finite Element Analyses To elucidate the biphasic behaviors of articular cartilage in compression tests, the biphasic finite element methods were applied. Firstly, three-dimensional analysis based on the method by Vermilyea and Spilker [13] was applied for an axisymmetric cylindrical model with uniform elastic modulus [10,11]. The cylindrical model of 1.5 mm radius and 2.0 mm thickness is composed of tetrahedron elements with 10 nodes (including midpoints), where the bottom plane is fully fixed as the tidemark connected to subchondral bone without fluid flow. The upper surface was uniformly compressed to attain 15% total compression at a constant rate for 10 s and then kept as constant total compressive deflection under unconfined condition for outer cylindrical surface. In this study, the typical values for constituent volume fraction; solid φ s = 0.17, fluid φ f = 0.83, Lame constants; λ = 0.1 MPa, μ = 0.3 MPa were used. For the permeability, constant values in horizontal direction of 9.14 x 10 –15 m4/N•s were used and the formula proposed by Jurvelin et al. [14] was applied in the compressed direction. The influence of surface conditions on the compressive behavior was evaluated by changing the friction level and permeability. That is, the surface boundary conditions are as follows; Condition 1: Fixed, impermeable, Condition 2: Frictionless, impermeable, Condition 3: Fixed, permeable, Condition 4: Frictionless, permeable. From this analysis, time-dependent strain behaviors near the surface at 0.1 mm

(at 5% depth) from the surface under these surface conditions are shown in Fig. 5 for compression of articular cartilage with uniform elastic modulus. Condition 2 corresponds to intact articular cartilage with very low friction and very low permeability. Under condition 2, the strain attains a slightly higher value than the average but approaches asymptotically the average strain of 0.15 during stress relaxation process. This time-dependent behavior did not correspond to the larger strain in surface zone as shown in Fig. 4. Therefore, the depth dependence of elastic properties was taken into consideration.

Times Figure 5. Changes in strain near surface of cartilage model with uniform elastic modulus under four kinds of surface conditions [11].

In this study, two-dimensional biphasic FEM analysis [15] with depthdependent elastic modulus under the condition 2 for 10% compression for 0.2 s (1 mm/s) was applied to simulate actual strain behavior. The commercial software of finite element method ABAQUS (v.6.5) was applied to a simplified two-dimensional square model of 3 mm horizontal length and 2 mm thickness. The biphasic pore pressure plane strain elements (CPE8RP: 8-node biquadratic displacement, bilinear pore pressure, reduced integration) of 10 μm x 10 μm square were used. The total number of the elements was 60,000. In this FEM analysis, the void ratio is assumed as 4 (water content is 80%). The Poisson’s ratio of solid phase is 0.125 and the permeability is 2.0 x 10 –15 m4/N・s. The average Young’s modulus is 0.74 MPa,. The depth-dependency of elastic modulus was estimated from the strain distribution at equilibrium after stress relaxation in constant total compression test, as shown in Fig. 6.

Local strain in experiment

Strain %

40 30 20 10 0

0

20

40

60

80

Distance from surface %

(a) Strain distribution at equilibrium

100

Young's modulus MPa

50

2.5 2.0 1.5 1.0 0.5 0

0

20

40

60

80

Distance from surface %

100

(b) Estimated elastic modulus

Figure 6. Depth dependence of elastic property of articular cartilage [15].

By applying of depth-dependent elastic property to FEM analysis, as a matter of course, the analyzed depth-dependent strain distribution at equilibrium coincided with the measured distribution. Furthermore, the deformed profile of circumference of cartilage specimen under compression was confirmed by FEM analysis for the model of depth-dependent elastic modulus. In the compression tests, a remarkable lateral bulging of surface zone (Fig. 7(a)) was observed, but the FEM analysis for a model of uniform elastic modulus showed the barreled profile in Fig. 7(b). In contrast, the biphasic FEM model with depth-dependent elastic modulus exhibited the appropriate profile as shown in Fig. 7(c).

Surface Surface

a)

(b)

(a) Observed profile

(c)

(b) Uniform modulus

(c) Depth dependent modulus

Figure 7. Side profiles of cartilage specimen immediately after compression in observation and FEM analyses for models different in elastic modulus property (vertical direction is compressive direction).

As described above, the importance of the depth-dependence of Young’s modulus was confirmed for deformed profile. However, the load carrying average stress by FEM analysis immediately after compression is about 0.17

MPa, which is about one order lower than the experimental one of about 2 MPa. Therefore, we examined the effect by evaluating parameters such as permeability, void ratio and elastic modulus on stress level. On the various parametric evaluation [15], we could simulate the time-dependent load-carrying capacity of cartilage by applying time-dependency to instantaneous elastic modulus as shown in Fig. 8. It should be noticed in biphasic FEM analysis that about half of mean stress is supported by fluid pressurization in fluid phase [16].

Experimentalvalue value Expreimental Analyzed value Analytic value

2.0 2

Mean stress [MPa]

Mean stress MPa

2.5

1.5

1.0 1 0.5

0

0

5

0

10

10

15

Time [s]

20

20

25

30

30

Times Time s Figure 8. Comparison of stress relaxation behavior [15].

In the present FEM analysis, we used usual biphasic model, but the effectiveness of fibular spring elements in cartilage model is pointed out by Li et al [17]. For the next stage, we plan to report biphasic FEM analysis with spring element reflecting the property of collagen fiber network.

3

Evaluation of Compressive Stimulation on Chondrocyteagarose Construct as Regenerated Cartilage

3.1 Materials and Methods To improve the mechanical properties of tissue-engineered cartilage, the appropriate mechanical stimulation is expected to enhance the metabolic activity of chondrocytes. In this study, the influence of compressive stimulation was evaluated in culture tests [18]. Chondrocytes were isolated from metacarpal-phalangeal joints of steers using a sequential enzyme digestion process. The cylindrical chondrocyteagarose constructs with a diameter of 4 mm and a height of 2.5 mm were

prepared as 4 wt% constructs with Sigma Type IX-A agarose to give an initial cell density of 1 x 107 cells/mL, and cultured in sterile culture medium (DMEM + 20% FBS) within a humidified tissue culture incubator controlled at 37°C and 5% CO2. Cyclic 15% compressive deformation was applied to test specimens at 1 Hz for 6 hours a day during their culture period by a mechanical loading system mounted within the incubator (Fig. 9). Constructs in the control group were placed in the same CO2 incubator and cultured in a 24 well tissue culture plate. Their upper surface was in contact with a loading plate; however dynamic loading was not applied to them. Constructs in the static group were maintained under the free-swelling condition, and ECM tissue elaboration and mechanical properties were evaluated in comparison to the compressed and the control specimens. The constructs in all groups were cultured up to 22 days and all culture medium was changed every two to three days.

2.5 mm

0.25 mm = 10%

Figure 9. Compressive loading apparatus for chondrocyte-agarose construct.

Chondrocyte-agarose constructs have uniform elastic property and the deformed profile correspond to uniform lateral bulging. Fluid pressurization appears to stimulate chodrocytes as described above for cyclic compression at 1Hz.

3.2 Mechanical Characterization and GAG Assay The mechanical characteristics of cell-agarose specimens were evaluated by an unconfined compression test after culture periods of 1, 8, 15 and 22 days. A single specimen was put on a 35 mm culture dish mounted on a mechanical testing machine, and immersed in the culture medium. 10 % compressive strain was applied by an impermeable stainless steel plunger at a strain rate of

20 %/min. The tangent modulus of cell-agarose construct was calculated from the linear region of the stress-strain curve, from 5 % to 10 % strain. Glycosaminoglycan (GAG) content of cultured cell-agarose constructs was determined by using the well established dimethyl-methylene blue (DMMB) Assay [18]. The GAG content of each sample was then determined using a standard curve derived from absorbance of solutions containing chondroitin sulfate from bovine trachea.

3.3 Results and Discussion The changes in tangent modulus and GAG content are shown in Fig. 10. The cyclic compression upregulated the GAG biosynthesis and the tangent modulus at day 15 of the cultured construct. However, it was observed that the cyclic compression also enhanced the GAG release from constructs into the culture media by the pumping effect [18]. Therefore, the total GAG synthesis of the cultured constructs should be evaluated by the sum of GAG content in the constructs and in the culture media. In these constructs, the compression group had synthesized a significantly larger amount of GAG compared to the static and control groups at day 22. control control static static compressed compressed

40.0 40

20 20.0

0 0.0 day1 day1 day8 day8 day15 day15 day22 Culture Period (days) (n=8-10)

(a) Tangent modulus

control GAG content wet weight G A G content (% w et w% eight)

Tangent Modulus [kPa] Tangent modulus kPa

60.0 60

static compressed 0.35 0.3 3.0

C ontrol S tatic C om pression

0.25 0.2 2.0 0.15 0.1 1.0 0.05

00 -0.05

D ay1

D ay8

D ay15

D ay22

day1 day8 day15 day22 C ulture P eriod Culture Period (days) (n=3-5)

(b) GAG content

Figure 10. Changes in tangent modulus and GAG content of chondrocyte-agarose constructs. Error bars mean standard deviation.

To examine the morphological characteristics of elaborated extracellular matrix, immunofluorescence observation was conducted on separate specimens (18). By the CLSM images of immunofluorescently stained constructs, the

distribution of proteoglycan (keratan sulfate) and collagen (type II) molecules were visualized within the construct. At the present stage, the structure and mechanical properties of regenerated cartilage do not reach the levels of natural articular cartilage. However, further optimization of culture conditions with appropriate environment of optimal scaffold, nutrient supply, possible addition of growth factors and mechanical stimulation is expected to improve the properties of regenerated cartilage. The time-dependent and depth-dependent mechanical behaviors in natural articular cartilage described above will be able to be reflected in evaluation of improved regenerated cartilage in the future.

4

Conclusions

The visualized compression tests of articular cartilage and corresponding biphasic FEM analyses indicated that the consideration of depth-dependent elastic modulus controlled the time-dependent and depth-dependent strain behavior of cartilage. The effectiveness of mechanical stimulation was examined for the enhancement of the metabolic activity of chondrocytes in tissueengineered cartilage. The influence of cyclic compressive loading on the tangent modulus and GAG content of chondrocyte-agarose constructs during culture tests was confirmed.

References 1. Dowson D., 1966-67. Modes of Lubrication in Human Joints, Proc. Instn. Mech. Engrs., 181, Pt3J, 45-54. 2. Murakami T., Higaki H., Sawae Y., Ohtsuki, N., Moriyama, S., Nakanishi, Y., 1998. Adaptive multimode lubrication in natural synovial joints and artificial joints, Proc. Instn. Mech. Engrs., Part H, 212, 23-35. 3. Sasada T. 2000. Lubrication of Human Joints – Nature of Joint Friction and “Surface Gel Hydration Lubrication”. J. Japanese Society for Clinical Biomechanics (in Japanese), 21, 17-22. 4. Ikeuchi K. 2007. Origin and future of hydration lubrication, Proc. Instn. Mech. Engrs.,Part J: J. Engineering Tribology, 2007, 221, 301-305. 5. Ateshian G.A. 1997. Theoretical Formulation for Boundary Friction in Articular Cartilage, J. Biomech. Eng., 119(1), 81-86. 6. Forster H. Fisher J. 1999. The influence of continuous sliding and subsequent surface wear on the friction of articular cartilage. Proc. Instn. Mech. Engrs., Part H, 213(4), 329-345.

7. Murakami T., Sawae Y., Ihara M., 2003. The Protective Mechanism of Articular Cartilage to Severe Loading:Roles of Lubricants, Cartilage Surface Layer, Extracellular Matrix and Chondrocyte, JSME International Journal, Ser. C, 46, 594-603. 8. Temenoff J.S., Mikos A.G., 2000. Review: tissue engineering for regeneration of articular cartilage, Biomaterials, 21, 431-440. 9. Murakami T., Sakai N., Sawae Y., Tanaka K., Ihara M., 2004. Influence of Proteoglycan on Time-dependent Mechanical Behaviors of Articular Cartilage under Constant Total Compressive Deformation, JSME International J., Ser. C, 47, 10491055. 10. Murakami T., Sakai N., Sawae Y., Kurohara Y., Ishikawa I., Okamoto M., 2006. Time-dependent Mechanical Behaviors of Articular Cartilage and Chondrocytes under Constant Total Compressive Deformation, Biomechanics at Micro- and Nanoscale Levels, Vol. II, Ed. By Wada, H., World Scientific, 37-47. 11. Murakami T., Sakai N., Sawae Y., Okamoto M., Ishikawa I., Hosoda N., Suzuki E., 2007. Depth-dependent Compressive Behaviors of Articular Cartilage and Chondrocytes, Biomechanics at Micro- and Nanoscale Levels, Vol. IV, Ed. By Wada, H., World Scientific, 36-46. 12. Mow V.C., Kuei S.C., Lai W.M., Armstrong C.G., 1980. Biphasic creep and stress relaxation of articular cartilage in compression theory and experiment, Journal of Biomechanical Engineering, 102, 73-84. 13. Vermilyea M.E., Spilker R.L., 1992. A hybrid finite element formulation of the linear biphasic equations for hydrated soft tissue, International Journal for Numerical Methods in Engineering, 33, 567-593. 14. Jurvelin J.S., Buschmann M.D. Hunziker E.B., 2003. Mechanical anisotropy of the human knee articular cartilage in compression, Proc. Instn. Mech. Engrs., Part H, 217, 215-219. 15. Hosoda N., Sakai N., Sawae Y., Murakami T., 2008. Depth-Dependence and TimeDependence in Mechanical Behaviors of Articular Cartilage in Unconfined Copmression Test under Constant Total Deformation, J. Biomechanical Science and Engineering, 3, 209-220. 16. Murakami T., Nakashima K., Sawae Y., Sakai N., Hosoda N., 2008. Roles of adsorbed film and gel layer in hydration lubrication for articular cartilage, Proc. Instn. Mech. Engrs., Submitted. 17. Li L.P., Soulhat J., Buschmann M.D., A. Shirazi-Adl A., 1999. Nonlinear analysis of cartilage in unconfined ramp compression using a fibril reinforced poroelastic model, Clinical Biomechanics, 14, 673-682. 18. Sawae Y., Honda J., Suzuki E., Morita Y., Watanabe M., Sanada T. Murakami T., 2007, Experimental characterization of regenerated cartilage model cultured under cyclic compression, Proc. ATEM (Advanced Technology in Experimental Mechanics) ’07 (CD-ROM), JSME-MMD.

Chapter 29 Wear Characteristics of a Monopivot Centrifugal Blood Pump for Circulatory Support Takashi Yamane1, Katsunobu Nonaka1, Osamu Maruyama1, Masahiro Nishida1, Ryo Kosaka1, Yoshiyuki Sankai2, and Tatsuo Tsutsui2 1. National Institute of Advanced Industrial Science and Technology, Tsukuba, Japan 2. University of Tsukuba, Tsukuba, JAPAN

1

Introduction

Donors for heart transplantation can be found only four patients a year in Japan. The remaining patients on the waiting list should depend on artificial hearts. The first-generation ventricular assist devices (VADs) were mostly pulsatile devices and were used only in hospitals. The second-generation VADs (Micromed DeBakey VAD, Jarvik research Jarvik-2000, Thoratec HeartMate II, SunMedical EVAHERAT) are small implantable rotary devices and allow the patients to be discharged from hospital, although their durability is limited since the impeller is supported by mechanical bearings or mechanical seals. The thirdgeneration VADs are implantable rotary types with long durability utilizing non-contacting bearings to allow patients to be discharged (Terumo DuraHeart, Ventracor VentrAssist, BerlinHeart INCOR, Arrow CorAide)[1]. The purpose of VAD application is not only for “bridge” to transplantation but also for “bridge-to-bridge” from short term VAD to long term VAD. We are developing a single pivot (monopivot) centrifugal pump for short term circulatory support including the above bridge-to-bridge use.

In an earlier study, the material combinations of ceramics or polymers, and pivot radius, were compared through the rotating wear test [2-3]. The reports on the wear characteristics of pivot-bearing pumps have so far been limited to double-pivot centrifugal types[4-5]. In these studies, the wear rate was reported to be 0.64 μm/day and 0.32 μm/day for left and right ventricular assist devices (LVAD and RVAD), respectively. In the present study, the wear characteristics of pivot bearings, used in monopivot centrifugal blood pumps of second-generation VADs, were investigated. Regarding a monopivot centrifugal pump, the pivot is concentratively exposed to phenomenon of wear, hemolysis, and thrombus formation. The selection of close radii radius between male and female pivots was examined to be effective or not in rotating wear tests and in animal tests.

2

Materials and Methods

[Monopivot pump] We are developing a monopivot centrifugal pump for circulatory assist lasting more than two weeks (Fig. 1). A closed-type impeller, with 4 vanes of 50 mm in diameter, is supported by a pivot bearing at the bottom and by a magnetic bearing with permanent magnets at the top. The pivot is surrounded by a cylindrical space of 7mm in diameter, a so-called a washout hole, and is supported by two bridges. The impeller is driven by a brushless DC Closed Impeller

Magnetic Suspension

Gap size Monopivot

(a) M onopivot mechanism

(b) M onopivot centrifugal pump, M C105

Figure 1. Pivot mechanism and monopivot centrifugal pump.

motor through an axial-flux magnetic coupling to reduce the eddy current losses. The gap size around the male pivot behind the impeller is 2.6 mm. Pivot axial load applied in the wear tests was determined based on a measurement with a load cell set just under the female pivot of a monopivot pump (Fig. 2). Since the maximum hydrodynamic lift was 7.8N (800gf) at 3000rpm and 13.7N (1400gf) at 4000 rpm, the necessary magnetic force to keep the pivot contact was regarded to be 9.8N (1000gf) assuming pump pressure as 300 mmHg. 16

DD7 (33wt% Glycerol aq)

Axial Load (N)

14 12 10

Hy drody namic lift

8

4000rpm 3000rpm 2000rpm 1000rpm

6 4 2 0 0.0

5.0 10.0 Flow Rate (L/min)

15.0

Figure 2. Axial load measurement for a monopivot pump.

[Rotating wear test] We conducted rotating wear tests with test pieces dipped in saline. The apparatus was composed of an upper rotating spindle and a lower axial loading part with two pulleys and weights (Fig. 3). To prevent splash, four fins were located in the saline vessel, whose temperature was maintained at 37 °C. The saline was supplied one-way from a bag (flow rate: 17 mL/h). Although five spindles could have been used simultaneously, only one was used for the present study to avoid differences in vibration, the adjustment of the upper and the lower centers, etc. between five spindles. [Animal test] Animal tests with a monopivot pump, MC105, were conducted for 4-5 weeks. Regarding the MC105, the radii of both the male pivot (material: Carbon) and the female pivot (material: UHMWPE) were 1.5 mm to eliminate the space for thrombus formation. Institutional guidelines for the care and the use of laboratory animals were observed.

φ4xL6.5 Male Piv ot

φ3xL3.5

Female Piv ot

6x6x t4.9

Rotational speed: 2000 rpm Axial load: 9.8 N Liqud: NaCl aq solution, 　　　37’C, 17mL/h

Figure 3. Apparatus of rotating wear test for pivot bearings.

3

Results and Discussion

[Former study] In our former study it was found that ultra-high molecular-weight polyethylene (UHMWPE) was better for female pivot material than any other hard materials[2, 3]. It was also found that the axial wear was minimum when the radius of male pivot was 1.5mm in a case of combination of sphere male pivot and planer female pivot under the axial loading of 9.8N. The first durability test of a monopivot pump, DD3, was performed for 90 days in a closed circuit filled with glycerin water solution. The male pivot radius was 1.5 mm (material: Al2O3) and the female pivot radius was 5 mm (material: ultra-high molecularweight polyethylene [UHMWPE]). Measuring the impeller/casing gap with a microscope and a video camera, the wear rate was found to be 350μm over 90 days, namely 3.9 μm/day (Fig. 4). [Rotating wear test] To reduce the wear of pivot bearing a set consisting of a carbon male pivot (r = 1.5mm) and a UHMWPE female pivot (r = 1.6mm), with closer radii, was selected and tested for 36 days. The wear rate was reduced to 24 μm over this period, namely to a rate of 0.67 μm/day (Fig. 5). It was found that a larger contact area was effective to reduce the wear rate.

Figure 4. Gap size trend during a durability test of a monopivot pump.

Male: 1.5R (Carbon) Female: 1.6R (UHMWPE,depth: 0.40mm) Axial load: 9.8 N Rotational speed: 2000rpm Dipped in: 37’C NaCl(aq)

Figure 5. Axial wear for a combination of close male/female pivots radii.

[Performance test and hemolysis test] The overall pump efficiency for monopivot pump MC5 was measured and was found to be 14% and 18% at conditions of 100mmHg-5L/min and 300mmHg-5L/min, respectively, using 33% 22°C glycerol solution. Hemolysis test was conducted at 100mmHg5L/min with bovine blood whose hematcrit was adjusted to 30% with saline. The plasma free hemoglobin was measured over 4 hours, the obtained normalized index of hemolysis was NIH = 0.0013 g/100L, which was almost half that of a commercial centrifugal pump (Medtronic BP80).

[Animal test] Animal experiments were conducted with sheep at the University of Tsukuba. The pump was put in a harness on the back of the animal; the inlet cannula was inserted into the left ventricle through the left atrium, and the outlet cannula into the aorta through the left carotid artery in a reverse manner. The activated coagulation time (ACT) was kept at around 200 s using heparin, as would be the case in clinical applications. In the preliminary test, the thrombus observed behind the tongue area was successfully eliminated by modifying the tongue shape to round and removing stagnation. Animal tests with the MC105 pump were conducted with sheep over 3, 15, 29 and 35 days (Fig. 6). The assisted flow rate was 1-2 L/min, and the pump rotational speed was kept at 1900 rpm, namely 100 mmHg pumping pressure, except for a period of recovering from sucking.

15 days (#1)

3 days (#2)

29 days (#3)

35 days (#4)

Device: monopivot pump MC-105 coated with MPC polymer Position: between left ventricle and carotid artery Cause of No Days Pump rpm Pump W Pump FlowACT(s) Thrombus termination 1 15 1940 5.0 1.2-0.7 160-260 none bleeding inflow valve 2 3 1900 5.3 1.2-1.4 130-180 cannula unclosure 15003 29 1900 3.1-5.8 1.0-2.0 140-260 male pivot infection 4 35 1910 4.6-5.5 0.9-1.5 170-250 none alive Figure 6. Results of animal tests with monopivot pump MC105.

Only with the 29-day experiment, thrombus formation was found around the male pivot. This may have been because the rotational speed was reduced to 80% and the flow rate decreased to 1.0 L/min on the second day, to avoid sucking. As a result, it was found that the center of the pivot did not wear and

the section became shaped like the letter W (Fig. 7). It was also observed that the thrombus that formed caused a slight wear at the shoulder of the female pit. With the 3-, 15-, and 35-day experiments no thrombi were observed. With the 35-day experiment, without thrombus development, the shape of the female pivot was measured using laser shape microscope before and after the experiment (Fig. 7); the total wear rate was found to be 73.6μm in 35 days, namely 2.1μm/day.

(Depth of female pivot)

UHMWPE before 35-day use

(Depth of female pivot)

UHMWPE after 35-day use 1.9 2.8

(Depth of female pivot)

UHMWPE after 29-day use

Figure 7. Pivot wear after 35 days and 29 days use of MC105 pump in animal tests.

4

Conclusions

Regarding a monopivot centrifugal pump, the selection of close radii radius between male and female pivots was examined. A combination of close radii of male/female pivots resulted in the wear rate of 0.67 μm/day in wear tests and that of 2.1μm/day in an animal test. In general, close radii combination between male/female pivots would provide more durability to the pump. In an extreme case no gap combination would add an effect of no thrombus formation.

References 1. Takatani S., Matsuda H., Hanatani A., Nojiri C., Yamazaki K., Motomura T., Ohuchi K., Sakamoto T., Yamane T., 2005. Mechanical circulatory support devices (MCSD) in Japan: current status and future directions, Journal of Artificial Organs, 8-1, 13-17. 2. Yamane T., Maruyama O., Mizuhara K., Nishida M., Nonaka K., Tateishi T., 2001. Durability enhancement of monopivot magnetic suspension blood pump, Journal of Congestive Heart Failure and Circulatory Support, 1-4, 317-320. 3. Nonaka K. K., Maruyama O., Yamane T., Miyoshi H., Mizuhara K., 2004. Improvement of anti-wear function of pivot bearing for a centrifugal blood pump (part1): Wear mechanisms in spinning motion, Japanese Journal of Tribology, 48-3, 247-258. 4. Makinouchi K., Nakazawa T., Takami Y., Takatani S., Nose’ Y., 1996. Evaluation of the wear of the pivot bearing in the Gyro C1E3 Pump, Artificial Organs, 20-6, 523528. 5. Asai T., Watanabe K., Ito S., Tsujimura S., Motomura T., Shinohara T., Glueck JAQ, Nose’ Y, 2004. Real-time studies of the pivot bearing in the NEDO Gyro PI-710 centrifugal blood pump, Artificial Organs, 28-10, 899-903.

PART VIII

Evaluation and Standardization

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Chapter 30 In Vitro Biodegradation of Poly(Lacticco-Glycolic Acid) Porous Scaffolds Guoping Chen, Taiyo Yoshioka, Naoki Kawazoe, and Tetsuya Tateishi Biomaterials Center, National Institute for Materials Science, Tsukuba, Japan

Abstract A protocol in which the condition is close to the in vivo pH environment was established for in vitro degradation evaluation of biodegradable porous scaffolds. Degradation of PLGA sponges was evaluated with the protocol. The PLGA sponges degraded with incubation time. For the first 12 weeks, the weight loss increased gradually and then remarkably after 12 weeks. In contrast, the number-average molecular weight (Mn) decreased dramatically for the first 12 weeks, and then less markedly after 12 weeks. Thermal analysis showed that the glass transition temperatures (Tg) decreased rapidly for the first 12 weeks, and the change became less evident after 12 weeks. These results suggest that the degradation mechanism of PLGA sponges was dominated by autocatalyzed bulk degradation for the first 12 weeks, and then by surface degradation after 12 weeks. Physical aging was observed during incubation at 37°C. The heterogeneous structure caused by physical aging might be one of the driving forces that induced autocatalyzed bulk degradation. The protocol suppressed extreme changes of the pH and will be useful in the degradation evaluation of porous scaffolds for tissue engineering.

1

Introduction

Porous scaffolds play an important role in tissue engineering as templates to accommodate cells and guide new tissue formation [1-3]. During the regeneration of new tissues or organs, the scaffolds should degrade and eventually disappear. Porous scaffolds constructed from aliphatic biodegradable polyesters, such as poly(lactic acid) (PLA), poly(glycolic acid) (PGA), and poly(lactic-co-glycolic acid) (PLGA), have been frequently used for tissue engineering because of their versatile biodegradability and mechanical properties. The porous scaffolds should have a degree of biodegradability that matches the formation of the new tissues or organs to allow the templating role of the scaffolding to gradually be replaced by the extracellular matrices. Therefore, evaluation of the degradation of porous scaffolds provides important information for the design and selection of biodegradable polymers for tissue engineering. Many methods have been established by mimicking the in vivo environment to study the degradation of biodegradable porous scaffolds [4-7]. For these reported methods, the buffer solution remained unchanged or completely changed, which resulted in an extreme change of pH before and after the change of buffer solution. The physiological pH is 7.4 and its change is very mild. A protocol that is close to the in vivo pH environment is necessary to evaluate the degradation of porous scaffolds. In this study, we established a new protocol with a stable pH for in vitro degradation evaluation of PLGA sponges. Their degradation mechanism was discussed in detail by analyzing the changes in weight, molecular weight, thermal properties, morphology, and pH.

2

Materials and Methods

2.1 Fabrication of PLGA Sponges Poly(D,L-lactic-co-glycolic acid) (PLGA) with a copolymer ratio of 75/25 (lactic acid/glycolic acid) was purchased from Sigma-Aldrich, Inc., (St. Louis, MO). Its weight-average (Mw) and number-average molecular weights (Mn) and polydispersity index (PDI) measured by gel permeation chromatography (GPC) were 109,520 ± 1,670, 47,870 ±1,560, and 2.3 ±0.1, respectively. PLGA sponges were fabricated by the particulate-leaching technique [8] using sieved sodium chloride (NaCl) particulates with a diameter range from 355 to 425 μm. Briefly, the NaCl particulates were added to a PLGA solution in chloroform at a weight ratio of PLGA/NaCl of 9/1, and mixed well. The chloroform was allowed to evaporate by air drying in a draft for 1 day, followed by 3 days of vacuum

drying. To leach out the NaCl particulates, the dried PLGA/NaCl composite was immersed in deionized water that was changed every hour. The washing was continued until the weight of the dried sponge did not change. Finally, the perfect leaching out of NaCl was confirmed by a simple qualitative analysis with a silver nitrate (AgNO3) aqueous solution, by which the change of transmittance caused by silver chloride precipitates of an ionic-reaction-product was detected by measuring light transmission. After drying, the PLGA sponges were cut into pieces having a dimension of 13.5 x 13.5 x 5 mm. The cross section of PLGA scaffold was observed by a scanning electron microscope (SEM) (JSM-6400Fs, JEOL, Ltd., Tokyo, Japan) operated at a voltage of 20 kV. The PLGA sponge was cut with a razor blade and then coated with platinum using a sputter coater (Sanyu Denshi Co., Tokyo, Japan). The pore size of the PLGA was measured by analyzing the pore size of 10 randomly chosen SEM images of the cross sections of the PLGA sponges. The porosity of the PLGA sponges was determined by a mercury porosimeter (Autopore IV, Shimadzu, Kyoto, Japan).

2.2 In Vitro Degradation Test An in vitro degradation test of the PLGA sponges was conducted in phosphate buffer solution (PBS) under pH 7.4 at 37°C with mechanical shaking (60 shakes/min). PBS was prepared by mixing 18.2％ (v/v) of 1/15 mol/l KH2PO4 aqueous solution and 81.8％(v/v) of 1/15 mol/l KH2PO4 aqueous solution; the pH of the mixture solution was adjusted to 7.4. The PBS was autoclaved before use. Before the degradation test, the PLGA sponges were sterilized by 70% ethanol aqueous solution. This treatment served another important role as a prewetting treatment, which makes PBS permeate into all the pores of the sponges. After complete washing with sterile PBS, the sponges were immersed in 20 ml sterile PBS. 50 mL disposable sterile polypropylene centrifuge tubes were used as test vessels. To suppress the pH change during the degradation test to a minimum, only the upper three-fourths of the PBS was replaced with fresh PBS every week. The sponges were collected every four weeks, washed with deionized water, air-dried for 1 day, and vacuum-dried for another 3 days. The dried samples were used for various evaluations. The pH of the removed PBS was measured by a pH meter (Shimadzu, Co., Kyoto, Japan).

2.3 Evaluation of Degraded PLGA Sponges The weights of the PLGA sponges were measured with an electrical balance, AG 135 (Mettler-Toledo International Inc., NY, USA). Weight loss in % was calculated according to a simple equation:

Weight loss (%) =

(W0 − Wt ) W0

100

(1)

where W0 is the initial weight and Wt is the weight at a given time point. Both W0 and Wt were measured after vacuum drying for 3 days. Mn and PDI of the PLGA sponges before and after degradation were determined by GPC using a high-performance liquid chromatography system, HLC-8220GPC (Tosoh Co., Tokyo, Japan), with two TSK gel columns (GMHHR-M, Tosoh Co., Tokyo, Japan). Chloroform was used as the elution solvent at a flow rate of 1.0 ml/min at 40°C; TSK polystyrene standards (Tosoh Co., Tokyo, Japan) were used for calibration. Glass transition temperature (Tg) and enthalpy relaxation (ΔHg) of the sponges were determined by differential scanning calorimetry (DSC) using a DSC instrument, DSC8240 (Rigaku Co., Tokyo, Japan). Each DSC measurement consisted of the following three steps: (1) heating from -15 to 200°C; (2) cooling from 200 to -15°C; (3) heating from -15 to 200°C. All steps were scanned at a rate of 10°C/min under a nitrogen gas flow at a rate of 50 ml/min. Only the DSC curves obtained from heating steps 1 and 3 were recorded, which were called the 1st and 2nd scans, respectively. The instrument was calibrated with indium, tin, and lead. The morphology of the degraded PLGA sponges was observed by the SEM operated at a voltage of 20 kV. Ten samples at each time point were used for the weight and pH measurements. Three samples at each time point were used for GPC. The data were used to calculate the means and standard deviations.

3

Results

3.1 Pore Size and Porosity of PLGA Sponges The PLGA sponges were prepared by the particulate-leaching method. SEM observation demonstrated that the PLGA sponges had a highly porous structure with the average pore size of 388.3 ± 51.8 μm. The average pore size was highly consistent with that of the NaCl particulates. The porosity of the PLGA sponges was 89.9 ± 1.1% measured by mercury porosimetry.

3.2 Change of Weight and Molecular Weight The PLGA sponges were incubated in PBS for 24 weeks and their weight change was recorded. The PLGA sponges lost weight with an increase in

Weight Loss (%)

incubation time (Fig. 1). The weight loss was slow for the first 8 weeks, moderate between 8 and 12 weeks, and increased rapidly after 12 weeks. SEM observation demonstrated that new, small pores were formed inside the wall and the wall of PLGA sponges became more porous after 24 weeks incubation (Fig. 2). 100 90 80 70 60 50 40 30 20 10 0 0

4

8

12

16

20

24

Incubation Time (weeks)

Figure 1. Change of weight loss (%) as a function of incubation time. Data represent the average ± SD.

Figure 2. SEM photomicrograph of the cross section of PLGA sponge after degradation for 24 weeks. Scale bar indicates 10μm.

The Mn and PDI of the PLGA sponges were determined by GPC using a high-performance liquid chromatography system (Fig. 3 and 4). The Mn decreased

dramatically for the first 12 weeks and only slightly after 12 weeks. After 24 weeks, the Mn decreased to 3,030 ± 480. The PDI increased rapidly for the first 8 weeks, and then decreased after 8 weeks. The decrease became less evident after 12 weeks.

Number Average Molecular Weight

30000 y = 27294e -0.1614x R 2 = 0.999 y = 24373e -0.1331x R 2 = 0.956

20000

y = 19219e -0.0935x R 2 = 0.851

10000

0

0

4

8 12 16 20 Incubation time (weeks)

24

28

Figure 3. Change of number average molecular weight (Mn) as a function of incubation time. Blue, red, and green lines are fitting curves by using Eq. 5 for the plots of the first 12, 16, and 24 weeks, respectively. Approximate function equations and R 2 values estimated from each of the fitting curves are also shown. Data represent the average ± SD.

Polydispersity Index

6 5 4 3 2 1 0

4

8 12 16 20 Incubation time (weeks)

24

Figure 4. Change of polydispersity index (PDI) as a function of incubation time. Data represent the average ±SD.

3.2 Degradation Mechanism Degradation of biodegradable polyesters commonly proceeds by chemical hydrolysis reaction of the ester bonds in its backbone, and thus results in formation of carboxylic acid end groups that act as a catalyst in the reaction [9]. The hydrolytic degradation mechanism can be classified into surface and bulk degradation [10,11]. In the surface degradation mechanism, the polymers are degraded only on their surface, which results in a decrease of size and loss of weight (Fig. 5a). In contrast, in bulk degradation, the degradation proceeds uniformly in the vertical direction to the surface, which results in a decrease of molecular weight (Fig. 5b). As a special case, the acidic degradation products of the bulk degradation may accumulate if the degradation products cannot diffuse from the polymer matrix. Therefore, the internal (core) part of the polymer matrix can degrade faster than the outer part (Fig. 5c). This type of degradation mechanism is known as an autocatalyzed bulk degradation [12]. The autocatalyzed bulk degradation also results in a decrease of molecular weight. The hydrolysis reaction of polyesters is shown as Equation (2). R-COO-R ' + H2O

(a) surface degradation mechanism

R-COOH + HO-R '

(2)

(b) bulk degradation mechanism

(c) autocatalyzed bulk degradation mechanism Figure 5. Schematic illustration of the three degradation mechanisms in the hydrolysis of bulk polymers.

The kinetics of the hydrolysis without and with the autocatalyzed effect can be expressed with Equations (3) and (4), respectively: -

-

d[E] dt d[E] dt

=

=

d[COOH] dt d[COOH] dt

= kd[COOH][H2O]

(3)

= kd ' [COOH][H2O] [E]

(4)

where [H2O], [E], and [COOH] are concentrations of water, ester, and carboxyl end groups in the polymer matrix, respectively, and kd and kd’ are apparent rate constants in each reaction. Assuming that [H2O][E] is unchanged and [COOH]=1/M n, the relationship between Mn and incubation time (t) in the nonautocatalyzed reaction (Eq. (5)) and autocatalyzed reaction (Eq. (6)) can be derived from Equations (2) and (3), respectively [13]. Mn(0) is the initial Mn of the polymer. Mn(t)-1 - Mn(0)-1 = kdt

(5)

In Mn(t) = In Mn(0) – kd '

(6)

Curve fitting of the data from the first 12 weeks in Fig. 3 with Eq. (6) showed a high correlation coefficient (R 2 = 0.999). The Mn(0) (27,290) obtained from the curve fitting coincided well with that obtained by GPC measurement (28,060 ± 2,040). However, good correlation was not achieved by curve fitting with Eq. (5). These results suggest that the degradation mechanism for the first 12 weeks is mostly based on the autocatalyzed bulk degradation mechanism. After 12 weeks, the correlation coefficient gradually decreased and the values of M n(0) obtained from the curve fitting gradually deviated from the GPC data with an increase of incubation time. The values of R 2 and M n(0) obtained from the curve fitting with 16 and 24 weeks of data were 0.956 and 24373, and 0.851 and 19,220, respectively. These results suggest that the degradation after 12 weeks was not dominated by bulk degradation. The degradation mechanism changed from one dominated by bulk degradation to one dominated by surface degradation. The PDI data can also be explained by the above degradation mechanism. During the autocatalyzed stage dominated by bulk degradation, the internal parts of the PLGA sponges degraded quickly and produced some acid oligomers and monomers, which resulted in an increase of the PDI. Between 8 and 12 weeks, the pooled oligomers and monomers flood out, which resulted in a decrease of the PDI. During the stage dominated by surface degradation (after 12 weeks), the change in the PDI became less evident.

3.3 pH Change of PBS Three-fourths of the PBS was replaced every week. The pH of the PBS was 7.40 immediately after the PBS was replaced with three-fourths fresh PBS. The pH of the removed three-fourths PBS was measured every week when the PBS was replaced. The pH of the removed PBS reflected the variation of the pH during the preceding week of incubation. The pH of the removed PBS was in the range of 7.43 and 7.24 (Fig. 6). The range of variation was very narrow. This result indicates that partial change of the PBS could return the pH to 7.40 and also avoided extreme pH change during incubation, which might be close to physiological condition. Therefore, the present protocol provided an incubation condition similar to the in vivo pH environment. 7.45

pH

7.35

7.25

7.15 0

4

8

12

16

20

24

Incubation time (weeks) Figure 6. Change of pH of the removed PBS as a function of incubation time. Data represent the average ±SD.

3.4 Change of Tg and ΔHg The PLGA sponges used in this study were confirmed to be completely amorphous by DSC measurement. It is known that amorphous polymers below their glass transition temperatures (Tg) are in a non-equilibrium state, and annealing of the polymers below Tg long enough results in a relaxation of the thermodynamic quantity toward an equilibrium state [14]. This phenomenon is called physical aging, and the resultant relaxation of enthalpy is enthalpy relaxation. The quantity of enthalpy relaxation can be detected by DSC as ΔHg. The change of Tg and ΔHg of the PLGA sponges during incubation was

determined by DSC, and is shown in Fig. 7. The data represented by the circles were obtained from the 1st scan. The data of the triangles were from the 2nd scan. The ΔHg in the 1st scan (open circles) increased for the first 8 weeks and then decreased. The ΔHg data obtained from the 2nd scan (open triangles) were almost the same for all samples. The average ΔHg of the 2nd scan was 2.84 ± 0.40 J/g. The ΔHg data from the 1st scan were higher than were those from the 2nd scan. The nonzero ΔHg in the 2nd scan might be caused by reannealing during the cooling and heating processes. The ΔHg data from the 1st scan were higher than were those from the 2nd scan, which suggests the occurrence of enthalpy relaxation in all scaffolds. 55

15

10 45 40 5

⊿ Hg (J/g)

Tg(°C) Tg (°C )

50

35 30 0

4

8 12 16 Incubation time (weeks)

24

0

Figure 7. Change of glass transition temperature (Tg) (closed) and enthalpy relaxation (ΔHg) (open) as a function of incubation time. The circles were obtained from the 1st scan, and the triangles from the 2nd scan.

The Tg obtained from the 2nd scan is more accurate than is that from the 1st scan and the Tg from the 2nd scan was used [15]. The Tg decreased for the first 12 weeks and remained almost unchanged after 12 weeks. Tg is dependent on Mn at neither extremely low nor high Mn; the dependence can be described by the empirical equation: Tg (Mn) = Tg (∞) -

C

Mn

(7)

where Tg(∞) is a Tg in infinite Mn, and C is a constant [16]. The decrease of Tg for the first 12 weeks might be ascribed to the decrease of Mn because of bulk degradation. After 12 weeks, the degradation was dominated by surface hydrolysis and the change of Mn became less evident, as shown in Fig. 3. Therefore, there was no evidence of Tg change after 12 weeks. It should be pointed out that Tg was lower than the incubation temperature (37°C) after 8 weeks. The decrease of ΔHg in the 1st scan after 8 weeks might be due to an energetic recovery of PLGA sponges at a temperature above its Tg.

4

Discussion

There are a number of protocols that have been used for evaluating degradation of biodegradable polymers. In this study, we established a protocol by making some modifications. Only three fourths of PBS was changed every week and the incubation test was done with mild mechanical shaking to make the environment nearer to the in vivo pH environment. Through the entire incubation time the maximum acidic pH value was successfully suppressed to pH 7.24, which change is very small compared to the results of other protocols. The mild shaking facilitated the diffusion of the buffered solution to avoid extreme pH change in the degrading regions. The protocol in the present study provided conditions closer to in vivo pH environment and the evaluated degradation behavior of the biodegradable polymers might more realistically reflect the in vivo degradation of biodegradable polymer scaffolds after implantation. The degradability of PLGA sponges was evaluated by the new protocol. There were evident turning points in the curves of weight loss, Mn, and Tg at the 12th week of incubation time. The PLGA sponges lost less weight, while their Mn and Tg decreased rapidly within the initial 12 weeks. In contrast, after 12 weeks, the PLGA sponges lost significant weight, but changes in their Mn and Tg were less evident. These results suggest that the degradability of PLGA sponges during the initial 12 weeks was dominated by autocatalyzed bulk degradation and that after 12 weeks was dominated by surface degradation. The significant weight loss during 8 and 12 weeks might be mostly attributed to the flooding out of accumulated acidic products due to the autocatalyzed bulk degradation. The flooding out of pooled acidic products during the period resulted in the drop of pH and PDI. Until now, the degradation of PLGA bulk materials and porous scaffolds at pH 7.4 and 37°C has been reported to be dominated by autocatalyzed bulk degradation through the whole degradation period [5,17]. In this study, however, we found that the degradation of PLGA sponge was dominated by autocatalyzed bulk degradation only in the early

period. The walls of the sponges were porous and their thickness was from a few um to dozens of um. Autocatalyzed bulk degradation occurred in the thin walls of the sponges, but did not last as long as the bulk materials because of the easy collapse of the autocatalyzed degradation regions and flooding out of degraded products. The autocatalyzed bulk degradation generated many small pores in the walls and greatly increased the surface area, which made the surface degradation override the autocatalyzed bulk degradation. The mechanism was confirmed by fitting of the relationship between molecular weight and incubation time with the equation for theoretical autocatalyzed bulk degradation. However, it should be emphasized that the biodegradability and its mechanism of biodegradable polyesters may be affected by various factors [18-26]. For example, crystallinity has a strong influence on the degradation mechanism because of degradation selectivity between amorphous and crystalline regions [27]. The degradation mechanism proposed in this study has been based on only one system, i.e. a specific copolymer ratio, molecular weight, porosity, pore size, amorphous property, and etc. Its application to other biodegradable scaffolds prepared with different methods and having different intrinsic properties need to be verified by further experiment. The driving force to induce the autocatalyzed bulk degradation of PLGA sponges at the initial stage might be heterogeneous structural fluctuation, which can lead to a selective hydrolysis reaction in the lower density regions. There are various reasons to form density fluctuations in amorphous polymer materials. For example, the fluctuation can be caused by thermal fluctuation, convection due to solvent evaporation, and incorporation of air bubbles, which may be affected by scaffold preparation methods such as phase separation, melt molding, and particulate leaching [28]. Physical aging may be one of the main reasons. It has been reported that PLGA undergoes physical aging during treatment or storage at room temperature and the effect becomes more evident when PLGA is incubated at 37°C. Various physical and mechanical properties of amorphous materials have been reported to be affected by physical aging [14,29,30]. Takahara et al. have clarified that the density fluctuation temporally becomes larger at an early stage of physical aging [31]. These reports demonstrate that the structural rearrangement of molecular chains by physical aging proceeds via a heterogeneous structure. In this study, physical aging was demonstrated by thermal analysis. The physical aging was more evident when Tg was higher than the incubation temperature (37°C) and less evident when Tg was lower than 37°C. The strong physical aging during the initial 8 weeks might increase the heterogeneous structural fluctuation, and thus partially account for the initiation of autocatalyzed bulk degradation.

5

Conclusions

The degradation of PLGA sponges was evaluated by a new protocol under shaking and with three fourths of the PBS being changed every week. The PLGA sponges degraded with the increase of incubation time. The degradation mechanism during the first 12 weeks was mostly based on autocatalyzed bulk degradation and that after 12 weeks it was mainly based on surface degradation. The new protocol suppressed the extreme variation of pH and will be useful for evaluating the degradation of porous scaffolds for tissue engineering.

Acknowledgments This work was supported by the New Energy and Industrial Technology Development Organization (NEDO) of Japan.

References 1. Langer R., Vacanti, 1993. Tissue Engineering, Science, 260, 920-926. 2. Ikada Y., 2006. Tissue Engineering: Fundamentals and Applications, In: Interface Science and Technology vol.8, Academic Press/Elsevier. 3. Chen G., Ushida T., Tateishi T., 2002. Scaffold Design for Tissue Engineering. Macromol Biosci., 2, 67-77. 4. Lu L., Garcia C.A., Mikos A., 1999. In vitro degradation of thin poly(DL-lactic-coglycolic acid) films, J Biomed. Mater. Res., 46, 236-244. 5. Wu L., Ding J., 2004. In vitro degradation of three-dimensional porous poly(D,Llactide-co-glycolide) scaffolds for tissue engineering, Biomaterials, 25, 5821-5830. 6. Wu L., Ding J., 2005. Effects of porosity and pore size on in vitro degradation of three-dimensional porous poly(D,L-lactide-co-glycolide) scaffolds fro tissue engineering, J Biomed. Mater. Res., 75A, 767-777. 7. Wu L., Zhang J., Jing D., Ding J., 2006. Wet-state mechanical properties of threedimensional polyester porous scaffolds, J Biomed. Mater. Res., 76A, 264-271. 8. Mikos A.G., Thorsen A.J., Czerwonka L.A., Bao Y., Langer R., 1994. Preparation and characterization of poly(L-lactic acid) foams, Polymer, 35, 1068-1077. 9. Loo S.C.J., Ooi C.P., Wee S.H.E., Boey Y.C.F., 2005. Effect of isothermal annealing on the hydrolytic degradation rate of poly(lactide-co-glycolide)(PLGA), Biomaterials, 26, 2827-2833. 10. Gopferich A., 1996. Mechanisms of polymer degradation and erosion, Biomaterials, 17, 103-114.

11. Burkersroda F., Schedil L., Gopferich A., 2002. Why degradable polymers undergo surface erosion or bulk erosion, Biomaterials, 23, 4221-4231. 12. Monhammadi Y., Jabbari E., 2006. Monte carlo simulation of degradation of porous poly(lactide) scaffolds, 1: Effect of porosity on pH. Macromol Theory Simul, 15, 643-653. 13. Tsuji H., 2004. Biodegradable polymers. Corona Publishing Co., Ltd., 124-128. 14. Hutchinson J.M., 1995. Physical aging of polymers, Prog. Polym. Sci., 20, 703-760. 15. Deng M., Uhrich K.E., 2002. Effects of in vitro degradation on properties of poly(DL-lactide-co-glycolide) pertinent to its biological performance, J Mater. Sci. Mater. Med., 13, 1091-1096. 16. Fox T.G., 1954. Flory PJ. The glass temperature and related properties of polystyrene -influence of molecular weight, J Polym. Sci., 14, 315-319. 17. Agrawal C.M., McKinney J.S., Lanctot D., Althanasiou K.A., 2000. Effects of fluid flow on the in vitro degradation kinetics of biodegradable scaffolds for tissue engineering, Biomaterials, 21, 2443-2452. 18. Jung J.H., Ree M., Kim H., 2006. Acid- and base-catalyzed hydrolyses of aliphatic polycarbonate and polyesters, Catalysis Today, 15, 283-287. 19. Freed L.E., Marquis J.C., Nohria A., Emmanual J., Mikos A.G., Langer R., 1993. Neocartilage formation in vitro and in vivo using cells cultured on synthetic biodegradable polymers, J Biomed. Mater. Res., 27, 11-23. 20. Park T.G., 1995. Degradation of poly(lactic-co-glycolic acid) microspheres: effect of copolymer composition, Biomaterials, 16, 1123-1130. 21. Tsuji H., Tezuka Y., Yamada K., 2005. Alkaline and enzymatic degradation of Llactide copolymers. II. Crystallized films of poly(L-lactide-co-D-lactide) and poly(L-lactide) with similar crystallinities, J Polym. Sci. Polym. Phys., 43, 10641075. 22. Tsuji H., Ikeda Y., 1998. Properties and morphology of poly(L-lactide). II. Hydrolysis in alkaline solution, J Polym. Sci. Polym. Chem., 36, 59-66. 23. Lam K.H., 1994. Nieuwenhuis P, Molenaar I. Biodegradation of porous versus nonporous poly(L-lactic acid) films, J Mater. Sci. Mater. Med., 5, 181-189. 24. Holy C.E., Dang S.M., Davies J.E., Shoichet M.S., 1999. In vitro degradation of a novel poly(lactide-co-glycolide) 75/25 foam, Biomaterials, 20, 1177-1185. 25. Grayson A.C.M., Cima M.J., Langer R., 2005. Size and temperature effects on poly(lactic-co-glycolic acid) degradation and microreservoir device performance, Biomaterials, 26, 2137-2145. 26. Tsuji H., Tezuka Y., 2005. Alkaline and enzymatic degradation of L-lactide copolymers, 1: Amorphous made films of L-lactide copolymers with D-lactide, glycolide and ε-caprolactone, Macromol. Biosci., 5, 135-148. 27. Tsuji H., Ikada Y., 2000. Properties and morphology of poly(L-lactide) 4. Effects of structural parameters on long-term hydrolysis of poly(L-lactide) in phosphatebuffered solution, Polym. Degrad. Stabil., 67, 179-189.

28. Yang S., Leong K.F., Du Z., Chua C.K., 2001. The design of scaffolds for use in tissue engineering. Part I. Traditional factors,Tissue Eng., 7, 679-689. 29. Hutchinson J.M., Smith S., Horne B., Gourlay G.M., 1999. Physical aging of polycarbonate: enthalpy Relaxation, Creep Response, and Yielding Behavior, Macromolecules, 32, 5046-5061. 30. Pan P., Zhu B., Inoue Y., 2007. Enthalpy relaxation and embrittlement of poly(Llactide during physical aging, Macromolecules, 40, 9664-9671. 31. Takahara K., Saito H., Inoue T., 1999. Physical aging in poly(methyl methacrylate) glass: densification via density fluctuation, Polymer, 40, 3729-3733.

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Chapter 31 Non-invasive Evaluation Technique for Cartilage Tissue Engineering Shogo Miyata1, Kazuhiro Homma2, Tomokazu Numano3, Takashi Ushida4 and Tetsuya Tateishi5 1. Faculty of Science and Engineering, Keio University, Yokohama, Japan 2. National Institute of Advanced Industrial Science and Technology, Japan 3. Tokyo Metropolitan University, Japan 4. Graduate School of Medicine, University of Tokyo 5. National Institute for Material Science, Japan

1

Introduction

Articular cartilage is avascular tissue covering articulating surfaces of bones, and it functions to bear loads and reduce friction in diarthrodial joints. It is a porous gel of large proteoglycan aggregates containing high fixed charge density (FCD) embedded in a water-swollen network of collagen fibrils [1,2]. Although articular cartilage may function well over a lifetime, traumatic injury or the degenerative changes associated with osteoarthritis (OA) can significantly erode the articular layer, leading to joint pain and instability [3]. Because of its avascular nature, articular cartilage has a very limited capacity to regenerate and repair. Moreover, the natural response of articular cartilage to injury is said to be variable and, at best, unsatisfactory. Therefore, numerous studies have reported tissue-engineering approaches to restore degenerated cartilage and repair defects; these approaches involve culturing autologous chondrocytes in vitro to create three-dimensional tissue that is subsequently implanted [4–7]. In these tissue engineering approaches, it is important to assess the biophysical and biochemical properties of the engineered cartilage. These

material properties of the engineered constructs are detectable only via direct measurements that are invasive and require destructive treatments such as histological analysis, biochemical quantification, and mechanical indentation testing. However, the application and utilization of these tissue engineering approaches in a clinical setting requires a noninvasive method of assessing the maturity of the actual regenerated cartilage tissue for therapeutic use. Moreover, the method should be applicable to various aspects of cartilage regenerative medicine, including the characterization of the regenerated tissue during in vitro culture and in vivo evaluation after transplantation. Magnetic resonance imaging (MRI) of articular cartilage is well accepted [8–11] and has been applied in recent years. In addition, the highly promising area of MRI to assess cartilage biochemistry is under investigation [12,13]. Indeed, specific material properties of cartilage explants have already been correlated with specific MR imaging parameters. For instance, increased water content in degenerated cartilage has been correlated with increased selfdiffusion of water [14], and loss of proteoglycan has been correlated with MRIdetermined FCD (fixed charge density) using gadolinium diethylenetriaminepentaacetic acid (Gd-DTPA2-) as a contrast agent [15]. Although the noninvasive assessment of tissue maturity and the nondestructive evaluation of molecular structure are important, we believe no previous study has fully evaluated the relationships between the biochemical properties and MRI measurements of regenerated cartilage consisting of articular chondrocytes. Previous study has indicated that MR images of autologous chondrocyte transplants may show clinically significant variations [16]; neither biochemical properties nor the FCD of regenerated articular cartilage has been evaluated. In this paper, we introduce our evaluation technique for tissue-engineered cartilage using MRI. We tested the hypothesis that MRI measurements of tissue-engineered cartilage correlate with biochemical and biomechanical properties and that these novel approaches can be used to assess cartilaginous matrix material properties during tissue reconstruction.

2

Magnetic Resonance Imaging (MRI) of Tissue Engineered Cartilage

Quantitative MRI evaluations were performed on a 2.0-T Biospec 20/30 System with a B-GA20 Gradient System (Bruker, Karlsruhe, Germany) with a maximum gradient strength of 100 mT/m. The MRI data acquisition and reconstruction

were performed using the ParaVision (Bruker) software system. In all MRI experiments, three or four sheets of the disks were stacked in layers and placed into glass tubes containing phosphate buffered saline (PBS) (Fig. 1). The measured parameters included longitudinal (T1) and transverse (T2) relaxation time and water self-diffusion coefficient (Diff). A longitudinal relaxation time map (T1-map) was obtained with a short echo time (TE: 15 ms) spin-echo sequence with different repetition time values (TR: 100 ms to 15 s, 16 steps). A transverse relaxation time map (T2-map) was obtained with a long repetition time value (TR: 15 s) spin-echo sequence with different echo time values (TE: 30 ms to 450 ms, 29 steps). A diffusion coefficient map (Diff-map) was calculated from the images obtained using a conventional diffusion weighted spin-echo (SE-DWI, TR: 15 s, TE: 35 ms) sequence with different b values (0, 74, 275, 603, 1059 s/mm2). All sequence were performed with a field of view (FOV) of 50 × 50 mm2, matrix size 64 × 64, and slice thickness 3 mm. The values of the relaxation time (T1 and T2) and the relative diffusion coefficient (Diff*) were calculated as the average of the specimen from the obtained T1-, T2-, and Diff-maps. The value of Diff* (= DiffS/DiffP) was calculated by normalizing the diffusion coefficient of the sample (DiffS) by the diffusion coefficient of PBS (DiffP) around the sample. All MRI measurements were carried out with no contrast agent at room temperature (23°C).

Figure 1. Schematic diagram of MR Imaging [17].

We used agarose gel culture for tissue-engineered cartilage model, because agarose is a biocompatible, thermosensitive hydrogel that offers excellent homogeneity and stability for assessing both biophysical and biochemical properties during in vitro culture, and has been used widely in cartilage mechanobiology. Chondrocyte-seeded agarose gels were prepared as described previously [18,19]. Briefly, the bovine chondrocytes in the feed medium (DMEM/F12 + 20% FBS + 50 μg/ml L-ascorbic acid) were mixed with an equal volume of PBS containing low-melting temperature agarose at 37°C to obtain 1.5 × 107 cells/ml in 2% (wt/vol) agarose gel and cast into a custom-made mold. After gelling at 4°C for 25 min, approximately 50 disks of 8 mm in diameter and 1.5 mm in thickness were cored out from the large gel plate by using a biopsy

Figure 2. T1-maps of day 1 (a), day 7 (b), and day 28 (c) post-inoculation specimens [17].

Figure 3. T2-maps of day 1 (a), day 7 (b), and day 28 (c) post-inoculation specimens [17].

Figure 4. Diff-maps of day 1 (a), day 7 (b), and day 28 (c) post-inoculation specimens [17].

punch. The chondrocyte/agarose disks were fed 2.5 ml medium/disk every alternate day and maintained in 5% CO2 atmosphere at 37°C up to 28 days in culture. Figure 2, 3 and 4 show the MRI maps of the enginnered cartilage. At the first stage of the culture (day 3), T1 and Diff of the engineered cartilage showed values similar to those of the PBS around the cartilage; hence, it was difficult to distinguish the boundaries between the engineered cartilage and the bath solution (PBS) in the MRI maps (Fig. 2a and 4a). By the end of the culture (day 28), the boundaries were distinct in both T1- and Diff-maps (Fig. 2a−2c and 4a−4c). In contrast, the boundary between the specimen and the PBS remained clear in the T2-map during the culture time (Fig. 3a−3c). The T1, T2, and Diff* values of the engineered cartilage were averaged, and the results are summarized in Figure 5. T1 and Diff* of the tissue-engineered cartilage had decreased with an increase in the culture time (Fig. 5a and 5c). On the other hand, T2 of the engineered cartilage showed considerably lower values than those of the PBS in

(a)

(b)

(c) Figure 5. Longitudinal relaxation time (a), transverse relaxation time (b), and relative diffusion coefficient (c) of the tissue-engineered cartilage during the culture time [17]. The values represent mean +/– S.D. (n = 3).

the glass tube throughout the culture time (Fig. 3), and these values tended to increase slightly with the culture time (Fig. 5b).

3

Relationships Between Mechanical Property and Quantitative MRI Measurement

To confirm the correlations between the quantitative MRI measurements and the biophysical and biochemical properties of the tissue-engineered cartilage, we performed linear regression analyses among the MRI-derived parameters (T1, T2, and Diff), the biochemical composition (sGAG content), and the biophysical properties (Em) of the engineered cartilage. Simultaneously with MRI experiments, mechanical testing and biochemical analysis were carried out to determine equilibrium compressive modulus Em and sGAG content of the engineered cartilage. As described previously, the specimens were analyzed by performing a stress relaxation test (maximum compressive strain, 20%) in unconfined compression. Following to the mechanical testing, the sGAG content was measured using DMMB assay.

(a)

(b)

(c) Figure 6. Scatter plots for the relationship between the equilibrium compressive modulus Em and longitudinal relaxation time (a), transverse relaxation time (b), and relative diffusion coefficient (c) [17]. Solid line represents the linear regression line.

The Em of the engineered cartilage (Fig. 6a and 6c) showed a strong correlation with T1 and Diff but a weak correlation with T2 (Fig. 6b). Similarly, the tissue sGAG concentration (Fig. 7a and 7c) and were found to be strongly correlated with T1 and Diff.

(a)

(b)

(c) Figure 7. Scatter plots for the relationship between the sulfated glycosaminoglycan concentration and longitudinal relaxation time (a), transverse relaxation time (b), and relative diffusion coefficient (c) [17]. Solid line represents the linear regression line.

4

Evaluation of Fixed Charge Density of Tissue Engineered Cartilage

For ‘native’ articular cartilage, the gadolinium-diethylene triamine pentaacetic acid (Gd-DTPA2-) -enhanced T1 imaging technique has been used to predict the PG content [15] and spatial distribution [18]. Furthermore, the negative fixed charge density (nFCD) can be estimated from consecutive T1 relaxation time measurement using Gd-DTPA2--enhanced MRI and be related to the PG concentration. This MRI technique is already well known as the “delayed

gadolinium enhanced magnetic resonance imaging of cartilage” (dGEMRIC) technique. The technique is based on utilization of the two-negative charge of the MRI contrast agent (Gd-DTPA2-). Sulfated glycosaminoglycans (sGAG) of the PGs are negatively charged in the cartilage which is known as nFCD: The electric exclusion force between this nFCD and the negatively charged contrast agent result in the inverse distribution of the contrast agent to the PGs distribution in the tissue. Consequently, T1 relaxation time and the nFCD determined by Gd-DTPA2--enhanced MRI are correlated with the PG concentration. In our previous study, we determined the FCD of the tissueengineered cartilage by Gd-DTPA2--enhanced MRI technique. The MRI measurements were performed with a 2.0-Tesla Bruker Biospec 20/30 system using Gd-DTPA2- contrast agent. In all MRI measurements, the specimens were put into glass tubes filled with PBS (Fig. 1). The longitudinal relaxation time map, T1-map, was obtained with a short-echo time (TE: 15 ms), spin-echo sequence with different repetition time values (TR: 100 ms to 15 s, 16 steps). Subsequently, the specimens were balanced in PBS containing 1 mM Gd-DTPA2– (Magnevist®, Nihon Schering, Osaka, Japan) for 12 hours; the longitudinal relaxation time map in the contrast agent, T1Gd-map, was obtained again with a short-echo time (TE: 15 ms), spin-echo sequence with different repetition time values (TR: 30 ms to 5 s, 13 steps). Finally, using the relaxivity (R) value of Gd-DTPA2– in saline (5.24 in our MRI system), the concentration of the contrast agent was estimated using the formula [Gd-DTPA2–] = 1/R(1/T1Gd – 1/T1). The negative fixed charge density (FCD) was calculated as follows

FCD =

[ Na + ]b [Gd − DTPA 2− ]t [Gd − DTPA 2− ]b

−

[ Na + ]b [Gd − DTPA 2− ]b [Gd − DTPA 2− ]t

,

where subscript b stands for bath solution and subscript t stands for cartilaginous tissue [20]. All MRI measurements were performed at room temperature 23°C. In the gadolinium-enhanced MR imaging measurements, longitudinal relaxation time of the bulk PBS containing Gd-DTPA reagent showed 0.179 ± 0.06 seconds in our MRI system. The T1Gd of the cultured specimen increased as a function of tissue maturation (0.197 ± 0.001 to 0.222 ± 0.003 seconds). In the T1Gd-maps, the [Gd-DTPA2–] in the specimen decreased, and the boundary between the specimen and the PBS bath became clearer with increased time in culture (Fig. 8). Interestingly, the T1Gd tended to show higher value in the

Figure 8. Quantitative water proton T1 maps in the presence of Gd-DTPA2- at day 3 (A), day 7 (B), day 14 (C), day 21 (D), day 28 (E) [19].

circumferential area of the disk than in the internal area from day 14 to 28. The FCD calculated from [Gd-DTPA2–] increased according to the time in culture (17.7 ± 1.8 to 40.4 ± 2.2 mM). As time in culture lengthened, the gross appearance of the cultured disk became increasingly opaque. Typical Safranin O-stained sections of the cultured specimens are shown in Figure 9. Over the culture time, the chondrocytes in the agarose gel appeared round, similar to the “native” articular cartilage. Figure 9 shows that the chondrocytes synthesized a thin shell of pericellular matrix (~ day 10) and expanded the volume of the cartilaginous matrix (~ day 28). The DMMB assay revealed that the sGAG content of the chondrocyte/agarose disks increased as a function of tissue maturation (0.19 ± 0.27 to 13.2 ± 1.9 mg/mL-disk-vol). Finally, the sGAG content of the reconstructed cartilaginous disk reached approximately 20% of the “native” articular cartilage (data not shown).

Figure 9. Histological appearance of chondrocyte/agarose disk stained by Safranin O at day 1 (A), day 10 (B), day 21 (C), day 28 (D) [19].

Figure 10. Scatter plots relating the tissue fixed charge density (FCD) to the sulfated glycosaminoglycan (sGAG) content [19]. The correlation between the fixed charge density and the sGAG content can be clearly seen (r = 0.95, n = 30, p < 0.001).

To correlate gadolinium-enhanced MRI and biochemical properties, the sGAG content of the tissue was plotted as a function of the FCD. From the linear regression analysis, the FCD correlated significantly with the sGAG content (r = 0.95, n = 30, P < 0.001) (Fig. 10), and the tissue [Gd-DTPA2–] correlated with the sGAG content by r = 0.83, n = 30, P < 0.001.

5

Conclusions

In conclusion, we evaluated the changes in the quantitative MRI parameters and matrix FCD derived from the MRI measurement of tissue-engineered cartilage that consisted of articular chondrocytes and hydrogel scaffolds. We found significant linear correlations between the MRI measurements and the biophysical and biochemical properties of the engineered tissue. Finally, we suggest that the quantitative MRI approach could be a useful noninvasive method for assessing the material property of tissue-engineered cartilage during the in vitro culturing process.

Acknowledgments This research was supported in part by the Special Coordination Funds for Promoting Science and Technology, and by a Grant-in-Aid for Young Scientists (B) (No. 18700414) from the Ministry of Education, Science, Sports and Culture of Japan, and by a grant from Tateishi Science and Technology Foundation of Japan.

References 1. Mow V.C., Kuei S.C., Lai W.M., Armstrong C.G., 1980. Biphasic creep and stress relaxation of articular cartilage in compression? Theory and experiments, J. Biomech. Eng., 102, 73-84. 2. Lee R.C., Frank E.H., Grodzinsky A.J., Roylance D.K., 1981. Oscillatory compressional behavior of articular cartilage and its associated electromechanical properties, J. Biomech. Eng., 103, 280-292. 3. Hunziker E.B., 1999. Articular cartilage repair: are the intrinsic biological constraints undermining this process insuperable? Osteoarthritis Cartilage, 7, 15-28. 4. Langer R.S., Vacanti J.P., 1999. Tissue engineering: the challenges ahead, Sci. Am., 280, 86-89. 5. Wakitani S., Goto T., Young R.G., Mansour J.M., Goldberg V.M., Caplan A.I., 1998. Repair of large full-thickness articular cartilage defects with allograft articular chondrocytes embedded in a collagen gel, Tissue Eng., 4, 429-444. 6. Aoki H., Tomita N., Morita Y., Hattori K., Harada Y., Sonobe M., Wakitani S., Tamada Y., Culture of chondrocytes in fibroin-hydrogel sponge, 2003. Biomed. Mater. Eng., 13, 309-316.

7. Chen G., Sato T., Ushida T., Hirochika R., Tateishi T., 2003. Redifferentiation of dedifferentiated bovine chondrocytes when cultured in vitro in a PLGA-collagen hybrid mesh, FEBS Lett., 542, 95-99. 8. Burgkart R., Glaser C., Hyhlik-Durr A., Englmeier K.H., Reiser M., Eckstein F., 2001. Magnetic resonance imaging-based assessment of cartilage loss in severe osteoarthritis: accuracy, precision, and diagnostic value, Arthritis Rheum., 44, 20722077. 9. McCauley T.R., Disler D.G., 2001. Magnetic resonance imaging of articular cartilage of the knee. J. Am. Acad. Orthop. Surg., 9, 2-8. 10. Schiller J., Naji L., Huster D., Kaufmann J., Arnold K., 2001. 1H and 13C HR-MAS NMR Investigations on Native and Enzymatically Digested Bovine Nasal Cartilage, MAGMA 13, 19–27. 11. Schiller J., Huster D., Fuchs B., Naji L., Kaufmann J., Arnold K., 2004. Evaluation of Cartilage Composition and Degradation by High-Resolution Magic-angle Spinning Nuclear Magnetic Resonance, Methods. Mol. Med., 101, 267–285. 12. Potter K., Butler J.J., Horton W.E., Spencer R.G., 2000. Response of engineered cartilage tissue to biochemical agents as studied by proton magnetic resonance microscopy, Arthritis Rheum., 43, 1580-1590. 13. Gray, M.L., Burstein, D., Xia, Y., 2001. Biochemical (and functional) imaging of articular cartilage, Semin. Musculoskelet. Radiol. 5, 329-343. 14. Shapiro E.M., Borthakur A., Kaufman J.H., Leigh J.S., Reddy R., 2001. Water distribution patterns inside bovine articular cartilage as visualized by 1H magnetic resonance imaging, Osteoarthritis Cartilage, 9, 533-538. 15. Bashir A., Gray M.L., Burstein D., 1996. Gd-DTPA2- as a measure of cartilage degradation. Magn. Reson. Med., 36, 665-673. 16. Alparslan L., Minas T., Winalski C.S., 2001. Magnetic resonance imaging of autologous chondrocyte implantation, Semin. Ultrasound CT MR, 22, 341-351. 17. Miyata S., Homma K., Numano T., Furukawa K., Tateishi T., Ushida T., 2007. Feasibility of noninvasive evaluation of biophysical properties of tissue-engineered cartilage by using quantitative MRI, J. Biomech., 40, 2990-29985. 18. Bashir A., Gray M.L., Hartke J., Burstein D., 1999. Nondestructive imaging of human cartilage glycosaminoglycan concentration by MRI, 41, 857-865. 19. Miyata S., Homma K., Numano T., Furukawa K., Tateishi T., Ushida T., 2006. Assessment of Fixed Charge Density in Regenerated Cartilage by Gd-DTPA Enhanced MR Imaging, Magn. Reson. Med. Sci., 5, 73-78.

Chapter 32 Analytical TEM Study of Biomineral Phases Yang Leng1 and Renlong Xin1,2 1. Department of Mechanical Engineering, Hong Kong University of Science and Technology, Hong Kong, China 2. School of Materials Science and Engineering, Chongqing University, Chongqing, China

1

Introduction

Bone related mineralization in physiological systems includes two aspects: formation of biominerals on osteoconductive/inductive implanted materials and formation of bone minerals in bone structure. Understanding the structures and properties of biominerals are crucial for decoding mechanism of new bone formation and osteoconduction/induction of biomaterials. It is well known that bone related biominerals are calcium phosphates in nature. It is commonly believed that the calcium phosphate of biominerals exhibits crystal structure of hydroxyapatite (HA). A number of studies however suggest that the biominerals may include other calcium phosphates, such as octacalcium phosphate (OCP) [1-7]. Brown and co-workers proposed that octacalcium phosphate (OCP) was a precursor phase of hydroxyapatite crystals in tooth and bone minerals in 1980s [1, 8]. The plate-like morphology, the low calcium to phosphor molar ratio (Ca/P) (less than 1.67) and the presence of HPO42- in bone minerals have been considered as supporting evidence [3, 5]. Bodier-Houlle et al. reported evidence of OCP to HA transformation in formation of human dentine crystals [9]. Crane et al. presented Raman spectroscopic evidence for OCP deposition during intramembranous mineralization [4]. We found that OCP more likely precipitate on bioceramic surfaces than HA in simulated body fluid and animal model

[10, 11]; and our theoretical analysis indicates formation of OCP is more likely than that of HA in body fluid environments due to fast nucleation rates of OCP crystals [12]. Determining the crystal structure of biominerals often encounters technical challenges because of limitations in their amount, size and chemical stability. This chapter views our work on analyzing biominerals using transmission electron microscopy (TEM). It describes special techniques of biomineral sample preparations, analytical methods and discussed possible confusion in phase identifications and phase instability in TEM examinations.

2

TEM Experimental

Biomineral examinations were conducted in a high resolution TEM system (JEOL 2010F, Japan) equipped with a field emission gun (point resolution: 0.19 nm) and an X-ray energy dispersive spectrometer (EDS). The electron beam was accelerated with 200 kV and its intensity was kept as low as possible during TEM examinations to minimize possible damage of samples by incident electrons. Bright field and high-resolution TEM (HRTEM) images were digitally recorded with a Gantan 794 Slow-Scan CCD (Multiscan) camera. Analysis of HRTEM images with fast Fourier transform (FFT) was conducted with the DigitalMicrograph 3.7.1 (Gatan, Inc., Pleasanton, CA) software package. Electron diffraction patterns were recorded on negative photo films. The diffraction ring Au (111) of polycrystalline gold (sputtered on the carbon film) were used as an internal standard to calibrate the dimensions of diffraction patterns.

3

Biominerals on Implants

3.1 Sample Preparation TEM sample preparation is often a challenge task, because a TEM sample should have thickness about 100 nanometers to ensure its electronic transparency. Fortunately, the biominerals formed on implanted biomaterials in physiological environments do not need thinning processing because one of their dimensions is sufficiently thin for TEM examinations. For example, Fig. 1 shows morphology of the biominerals precipitated on porous bioceramic specimens implanted in muscle of animal models. The precipitates are either in shape of flake or rod. The sample preparation for such biomineral precipitates becomes relatively simple. We harvested the bioceramic implants after a period of implantation by sacrificing animals. The harvested implants with surrounding tissues then were

a

b

Figure 1. Biomineral precipitation on α-tricalcium phosphate (TCP) specimens implanted in muscle of a) rabbit and b) dog.

washed with phosphate-buffered saline (PBS). The attached tissues were cleaned with a mixture of PBS (90 wt%) and pepsin (10 wt%). The samples of the biomineral precipitates were separated from an implant in an ethanol solution with mechanical forces of ultrasound vibrations. The precipitates, separated from the implant and dispersed in the solution, were picked up with a TEM copper grid coated with a carbon film. After air drying, the copper grid was ready to be mounted in a TEM sample holder for examinations.

3.2 Results The samples of biomineral precipitates prepared by the aforementioned methods generated high quality bright field images (Figs. 2a and 3a) in TEM. Also, the samples can produce excellent single crystalline diffraction patterns (Figs. 2b and 3b). The flake-like biomineral sample extracted from an α-TCP ceramic implant in rabbit muscle (Fig. 2) was identified as OCP. The diffraction pattern revealed d-spacings of 9.4Å and 6.8Å, which matched well with 9.38Å of OCP – (110) plane spacing and 6.83Å of OCP (001) plane spacing, respectively. Also, the measured interplanar angle on Fig. 2b matched with the one between OCP – (110) and (001) (89.7°). The rod-like biomineral sample extracted from an α-TCP ceramic implant in dog muscle (Fig. 3) was identified as HA. The diffraction pattern revealed d-spacings of 4.7Å and 6.9Å, which matched well with 4.71Å of HA (110) plane spacing and 6.88Å of HA (001) plane spacing, respectively. The measured interplanar angle between HA (110) and (001) on Fig. 3b was exact 90°, matching with the hexagonal structure of HA. In fact, the hexagonal cross-sections of rod-like precipitates shown in SEM micrograph

a

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Figure 2. OCP flake precipitate on the α-TCP implanted in rabbit muscle: (a) TEM bright field image; (b) corresponding diffraction pattern, identified as that of OCP with [110] zone axis.

a

b

Figure 3. HA rod precipitate on the α-TCP implanted in dog muscle: (a) TEM bright field – image; (b) corresponding diffraction pattern, identified as that of HA with [1 10] zone axis.

(Fig. 1b) revealed their hexagonal symmetry of crystal structure as it was confirmed by electron diffraction in TEM (Fig. 3b). HRTEM images provided additional ultrafine-structural information of biomineral precipitates. For example, Fig. 4 shows the HRTEM analysis results

of a rod precipitate. The HRTEM fringes (Fig. 4b) revealed the differences in their crystalline structure from areas 1 and 2 of rod marked in Fig. 4a. Under the same amount of defocus, Figure 4b exhibits two distinct fringe patterns, corresponding to areas 1 and 2 in Fig. 4a. Fast Fourier transformation (FFT) enabled us to extract symmetrical information in the HRTEM fringes and generated the patterns similar to the diffraction ones shown in Figs. 2b and 3b. From the FFT patterns, we identified that area 1 was of the OCP structure and area 2 was of the HA structure. In addition, Fig 4b revealed the orientation – relation between OCP and HA in a single rod as OCP (110) //HA (110) and – OCP (001)//HA (001). Interestingly, this orientation relation is the same as that we have found in solid state transformation of synthetic OCP to HA [13].

a 1

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Figure 4. OCP/HA co-existence in biomineral precipitation. a) Bright field TEM image of rod precipitate; b) HRTEM image including areas 1 and 2 and their FFT patterns at the bottom reveal their OCP and HA structure, respectively.

4

Bone Minerals

4.1 Sample Preparation The bone mineral samples were collected from cortical bones of human tibiae belonged to a deceased 53-year old man. We extracted bone minerals with ambient temperature process to avoid any temperature-induced phase transformation in bone minerals during extraction. The fresh cortical bone samples were fixed in 4% paraformaldehyde (pH 7.4) for 24 hours and then they were immersed in a 10% (w/v) ethylenediamine tetraacetic acid (EDTA) solution (pH 7.4) for extractions of bone minerals at room temperature. The EDTA solution was prepared by dissolving EDTA disodium salt dehydrate (reagent quality) in distilled water and the pH value of solution was adjusted with sodium hydroxide. The mineral extraction in the EDTA solution was conducted with the assistance of ultrasonic waves generated by a conventional ultrasonic cleaner that was operated at a frequency of 42 kHz. The EDTA was changed every day during the three-day extraction processing. After extraction, 10 ml of the EDTA solution was transferred to a tube filled with 1 ml of double distilled water in order to separate the bone minerals from attached organic substances. The tube was centrifuged at 4,000 rpm for 5 minutes. After centrifuging, the top 10 ml of the solution in the tube was immediately poured away in order to obtain mineral-containing, organic-free solution at the tube bottom. The ultrasound vibration was also used to assist dispersing bone mineral particles in the solution, before the solution was dropped on a copper grid with carbon film for TEM observations.

4.2 Results The majority of the bone minerals extracted by the EDTA solution exhibited irregular and plate-like morphology. Figure 5 shows bright field TEM images of bone minerals in human cortical bone. The EDS examinations confirmed calcium and phosphorous contents in bone minerals (the inset in Fig. 5a). Mostly bone mineral particles showed very weak contrast in TEM because of their nanometer thickness. Occasionally, we observed large pieces of bone minerals as shown in Fig. 6a. The diffraction pattern of bone mineral piece from the area, marked by a solid circle in Fig. 6a, revealed the OCP structure (Figs. 6b). The diffraction pattern in Fig. 6b was identified as that of OCP with the [110] zone axis, which is identical to OCP pattern shown in Fig. 2b. The OCP crystal structure was found by electron diffraction in the area marked by the arrow in Fig. 6a.

a

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Figure 5. TEM bright-field images of minerals extracted from cortical bone of human tibia: (a) lower-magnification image; (b) higher-magnification image. The inset in (a) is an EDS spectrum of several bone minerals.

a

b 220 002

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Figure 6. (a) TEM bright-field image of several bone minerals from human cortical bone; (b) the diffraction pattern from the solid circle area in (a) identified as OCP with the [110] zone axis; (c) the diffraction pattern from the dashed circle are in (a) identified as – HA with the [1 10] zone axis.

We also identified the HA structure in the area marked by a dashed circle in Fig. 6a. HA in bone minerals is commonly considered as the ‘bone apatite’, because it is calcium-deficient and carbonate-contained. Note that the bone apatite shows little difference in diffraction pattern from stoichiometric HA. – The diffraction pattern in Fig. 6c was identified as that of HA with [110] zone axis, identical to the HA pattern in Fig. 3b. Interestingly, the diffraction patterns in Fig. 6b and Fig. 6c also revealed the OCP/apatite orientation relation in bone minerals is exactly as we observed on biomineral precipitates as shown in Fig. – – 4b, that is, OCP (110) //apatite (110) and OCP (001)//apatite (001).

(a) a

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5 nm Figure 7. OCP structure in bone minerals revealed by HRTEM. (a) bright-field image of a bone mineral sample; (b) an HRTEM image of the square area in (a); (c) the FFT pattern of (b), which is identified as OCP with the [110] zone axis.

It was rather lucky for us to obtain excellent diffraction patterns of bone mineral crystals shown in Figs. 6b and 6c. In fact, such diffraction patterns were extremely difficult to be found in bone minerals, because most of bone mineral samples were too small for a TEM aperture to single out their single crystalline

areas. Frequently, we had to rely on HRTEM techniques for their phase identification. For example, Fig. 7b shows an HRTEM image of the square area of a bone mineral sample in Fig. 7a. The corresponding FFT of HRTEM is shown in Fig. 7c. The HRTEM fringe image and corresponding FFT match with those of OCP. Note that FFT shown in Fig. 7c is geometrically identical to the diffraction pattern shown in Fig. 6b. The apatite structure was also frequently revealed by HRTEM in bone minerals. For example, Fig. 8 shows a typical HRTEM image of apatite crystals in bone minerals. The HRTEM image (Fig. 8b) and corresponding FFT (Fig. 8c) – matched well with those of HA with the [110] zone axis. The HRTEM image shown in Fig. 8 was obtained without tilting the plate-like bone mineral samples. – It implies that the flat surface of bone mineral is parallel to the apatite {110} plane.

a

50 nm

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b (110)

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001 110

2 nm Figure 8. Apatite structure in bone minerals revealed by HRTEM. (a) bright-field image of a bone mineral sample; (b) HRTEM images from the square area in Fig. 8a. (c) The – FFT pattern of an HRTEM image (b) showing the HA structure with the [1 10] zone axis.

5

Discussion

OCP has a triclinic crystal structure (a = 19.692Å, b = 9.523Å, c = 6.835Å, α = 90.15°, β = 92.540° γ = 108.65°) which is distinctively different from HA’s hexagonal crystal structure (a = 9.424 Å, c = 6.879 Å) [5]. However, their identities can be confused by their diffraction patterns. We found that the most frequently observed OCP diffraction pattern is the one with [110] zone axis that – looks similar to that of HA with the [110] zone axis. Figure 9 superimposes such two patterns in order to illustrate their similarity and differences. Although their high similarity, as illustrated in Fig. 9, two main differences enabled us to distinguish OCP from HA: (1) the d-spacing of (110) in the HA pattern is 4.71 Å, – while the d-spacing of (110) in the OCP pattern is 9.38 Å, approximately double of the d-spacing of HA (110); (2) the interplanar angle of (110) and (001) is – exactly 90° in the HA pattern, while the interplanar angle of (110) and (001) is about 89.7° in the OCP pattern. Note that the large d-spacing of 9.38 Å in OCP makes the OCP pattern unique in diffraction among all kinds of calcium phosphate structures. Thus, OCP can unmistakably be identified from the d-spacing of 9.38 Å in a diffraction pattern.

(110)HA

_

(110)OCP (001)HA

(001)OCP

– Figure 9. Schematic comparison the diffraction patterns of HA with B = [1 10] (marked with open circles) and OCP with B= [110] (marked with solid squares). Similarity of these two patterns can cause confusion in phase identification.

We have noted that OCP phase is vulnerable to structure changes under electron beam irradiation in TEM. The electron beam can raise temperature of OCP samples rapidly due to their poor thermal conductivity. Figure 10 shows the example of electron beam induced structure changes in an OCP sample.

After irradiated by a high intensity electron beam, the OCP sample generates bubble-like damages. In fact, the OCP has transformed to HA and produced a side-product of the phase transformation as bubbles [13]. Thus, the diffraction information and HRTEM images of OCP crystals presented in this articles were obtained by carefully lowing the intensity of electron beam to avoid irradiation damage. We believe that poor thermal instability is one of the main reasons that OCP was not found and often misidentified as apatite, particularly in bone minerals. Our findings might raise attention to the roles of OCP in biomineralization. Our studies provided evidence of OCP as a precursor phase during biomineralization processing and as an existing phase in bone minerals.

Figure 10. OCP is relatively instable under electron beam irradiation in TEM. The OCP can be transformed to HA during TEM observation and results in generation of “bubbles”, the side-products of transformation.

6

Conclusions

Octacalcium phosphate is identified as a crystalline phase existing in biominerals, including mineral precipitates in physiological conditions and bone minerals. Octacalcium phosphate and hydroxyapatite can co-exist in a simple – piece of biominerals with crystalline orientation relations as OCP (110) //HA – (110) and OCP (001)//HA (001).

Acknowledgments This work was financially supported by the Research Grants Council of Hong Kong and the internal research funds from the Hong Kong University of Science and Technology. The human bone sample was provided by Dr. Xia Guo at the Hong Kong Polytechnic University.

References 1. Brown W.E., Eidelman, N., and Tomazic, B., 1987. Octacalcium phosphate as a precursor in biomineral formation, Adv. Dent. Res. 1, 306-313. 2. Cho G., Wu Y., Ackerman J.L., 2003. Detection of hydroxyl ions in bone mineral by solid-state NMR spectroscopy, Science 300, 1123-1127. 3. Chow L.C. Eanes E.D., 2001. Octacalcium Phosphate. Basel: Karger. 4. Crane N.J., Popescu V., Morris M.D., Steenhuis P., Ignelzi Jr M.A., 2006. Raman spectroscopic evidence for octacalcium phosphate and other transient mineral species deposited during intramembranous mineralization, Bone 39, 434-442. 5. Elliott J., 1994. Structure and chemistry of the apatites and other calcium phosphates. Amsterdam: Elsevier. 6. Loong C.K., Rey C., Kuhn L.T., Combes C., Wu Y., Chen S.H., Glimcher M.J., 2000. Evidence of hydroxyl-ion deficiency in bone apatites: An inelastic neutronscattering study, Bone 26, 599-602. 7. Pasteris J.D., Wopenka B., Freeman J.J., Rogers K., Valsami-Jones E., Van Der Houwen J.A.M., Silva M.J., 2004. Lack of OH in nanocrystalline apatite as a function of degree of atomic order: Implications for bone and biomaterials, Biomaterials 25, 229-238. 8. Brown W.E. Smith J.P., 1962. octacalcium phosphate and hydroxyapatite, Nature 1962, 1050-1055. 9. Bodier-Houlle P., Steuer P., Voegel J.C., Cuisinier F.J.G., 1998. First experimental evidence for human dentine crystal formation involving conversion of octacalcium phosphate to hydroxyapatite, Acta Cryst. Sect. D 54, 1377-1381. 10. Xin R., Leng Y., Chen J., Zhang Q., 2005. A comparative study of calcium phosphate formation on bioceramics in vitro and in vivo, Biomaterials 26, 6477-6486. 11. Xin R., Leng Y., Wang N. Ultrastructure study of hydroxyapatite precipitation on ceramic surfaces in dog model, Mater. Sci. Eng. C. in press. 12. Lu X. Leng Y., 2005. Theoretical analysis of calcium phosphate precipitation in simulated body fluid, Biomaterials 26, 1097-1108. 13. Xin R., Leng Y., Wang N., 2006. In situ TEM examinations of octacalcium phosphate to hydroxyapatite transformation, J. Cryst. Growth 289, 339-344.

Chapter 33 High-throughput Cytometry Using Antibody Arrays Koichi Kato, In-Kap Ko, Toshinari Ishimuro, Mitsuaki Toda, Yusuke Arima, Isao Hirata, and Hiroo Iwata Institute for Frontier Medical Sciences, Kyoto University, Kyoto, Japan

1

Introduction

Information on specific surface markers provides various possibilities in the processing and analysis of living cells. One of the most representative cases is the immunophenotyping of blood-related cells. Various surface markers specific for differentiation stages are analyzed to characterize the population of hematopoietic stem cells [1]. A panel of cluster of differentiation (CD) antigens is clinically adopted for typing leukemia for diagnosis [2]. In the realm of stem cell-based regenerative medicine, the purification of specific cell types using information on surface markers will be necessarily crucial for the safe and efficient treatments of damaged tissues [3]. Conventionally, surface markers have been analyzed by flow cytometry (FCM) in which the reactivity of antibodies against surface markers is evaluated for individual cells in a population taking advantage of techniques for fluorescent labeling and their detection [4]. This analytical method has provided profound possibilities for the characterization of living cells. However, the throughput of analysis is still limited even with a state-of-the-art multicolor FCM apparatus. This limitation is critical when analysis is targeted to a large number of surface markers.

To circumvent this limitation, antibody arrays have been developed in our laboratory [5–8] as well as others [9,10]. This is the novel analytical technique that has been realized using advanced technologies for microfabrication and material surface engineering. The basic mechanism of the analytical method is to assess simultaneously the reactivity of surface antigens expressed on cells with surface-immobilized multiple antibodies (Fig. 1). Different antibodies are immobilized in site-addressable manners, and the expression pattern of surface markers is simply attained by inspecting cell adhesion on every spot with standard optical microscopes. The simplicity of this method makes it useful for first screening over a large number of surface markers. Although the method is similar to the analysis of protein–protein interactions using protein microarrays, the technique described here utilizes interactions of antibody with antigens carried by living cells. In this paper, we describe the outline of high-throughput cytometry using antibody arrays, based on our published data [5–8]. The first part of this paper introduces the method we used for the fabrication of antibody arrays. In the second section, general features of surface marker analysis using antibody arrays are described, focusing on the quantitative aspect of the analysis. Because the throughput of analysis largely relies on a detection method, the third section is devoted to discuss the techniques that allow to determine the number of cells bound to antibody arrays. Finally, the forth section demonstrates the feasibility Y Y

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Figure 1. Strategy to survey surface markers using an antibody array. Cells are captured on antibody-immobilized spots (ex. spots with antibodies 1 and 3), if cells express surface markers specific for immobilized antibodies. By inspecting captured cells for all the antibody-immobilized spots on a single array, one can attain the pattern of surface markers expressed in the population tested. Drawings are not to scale.

of the antibody array, showing the results of surface marker analysis conducted for neural stem cells (NSCs). In the case of the NSCs and many other cell types as well, the heterogeneity of cell populations is critical problem in the identification of surface markers expressed on a specific cell type.

2

Fabrication of Antibody Arrays

To date, various techniques have been developed for the microfabrication of organic and inorganic materials. Among them, the photo-assisted patterning of alkanethiol monolayers provides a simple and reproducible means to fabricate antibody arrays. Figure 2 shows the outline of array preparation that we are routinely employing for laboratory-scale experimentation. The method involves several steps of chemical reaction at the surface of glass-based substrates. First, a thin gold layer is deposited on a glass plate, and subsequently a selfassembled monolayer (SAM) of methyl-terminated alkanethiol (CH3-SAM) is formed on gold. The CH3-SAM is overlaid with a photomask, and the CH3SAM beneath circular windows of the mask is photolytically degraded to create spots presenting again a bare gold surface. Then, carboxylic acid-terminated SAM (COOH-SAM) is formed within the spots, and the carboxylic acid thus introduced is used to covalently immobilize antibodies via the standard carbodiimide activation method. Finally, the surface of an array is treated with blocking agents such as bovine serum albumin to minimize nonspecific cell adhesion during surface marker analysis. The site-addressable immobilization of antibodies requires to dispense antibody solutions correctly into the separate spots with a typical diameter of 1 mm. This can be manually done with a micropipette for laboratory-scale antibody arrays. The large contrast in hydrophilicity between the spot (activated COOH-SAM) and the surrounding area (CH3-SAM) makes it easy to locate a droplet (~100 nL) of antibody solution within a spot. For the preparation of arrays loaded with hundreds of antibodies, automated machines such as microarrayers may be suitably used. The size of antibody-immobilized spot determines the capacity of the spot for cell capturing. For instance, ca. 5000 lymphocytes (~10 μm in diameter) can be captured at most on a spot with a diameter of 1 mm. As described later, such capacity is required for the quantitative evaluation of cells that express a specific surface marker with sufficient accuracy.

CH3-SAM Gold Glass UV light Photomask

COOH-SAM

NHS ester

Antibody solution

Antibody Albumin

Figure 2. Scheme for the preparation of an antibody array through photo-assisted patterning of an alkanethiol monolayer and the subsequent site-addressable immobilization of a panel of antibodies. NHS: N-hydroxysuccineimide. (Reprinted with permission from [6]. Copyright 2005 Elsevier)

As described later, cell binding assays can be reproducibly performed by constructing a parallel plate chamber on an antibody array [8]. The chamber serves to prevent the inhomogeneous distribution and dislocation of cells, which is critically important for quantitative cell binding assays.

3

Quantitative Immunophenotyping

To evaluate the feasibility of antibody arrays for surface marker analysis, we tested three leukemia cell lines whose surface markers had been well

characterized. The cell lines included CCRF-CEM (T-cell acute lymphoblastic leukemia), Ramos (B-cell Burkitt’s lymphoma), and HL-60 (myelomonocytoidcell acute promyelocytic leukemia). Figure 3 shows the results of cell binding assays. The array displayed 22 antibodies against different CD antigens together with 3 nonspecific immunoglobulins as controls. As can be seen, the binding patterns are different between the three cell lines. Given the specificity of cell

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Figure 3. Parallel analysis of 22 surface markers expressed on leukemia cells by cell binding assays on antibody arrays. (A) Position of antibodies and control immunoglobulins (G1, G2a, and G2b) immobilized on an array. Antibodies are represented with their CD numbers or the name of a specific antigen (HLA-DR) for short. (B–D) Digital microscopic images of (B) CCRF-CEM, (C) Ramos, and (D) HL-60 cells bound to the antibody array. Bars: 1 mm. (Reprinted with permission from [7]. Copyright 2007 Elsevier)

CCRF-CEM determined (%)

binding, these patterns represent the profiles of surface antigens expressed on the tested cells. These patterns are in good agreement with the expression patterns of surface markers reported for three cell lines including CCRF-CEM, Raji (B-cell Burkitt’s lymphoma), and HL-60 [10]. Our data show that similar assays can also be performed for adherent cells such as mesenchymal stem cell lines [7]. The number of cells adhering to the spots are varied depending on the cells tested. For example, on the spots with anti-CD4 antibody (Fig. 3 B and D), the number of cells is larger with CCRF-CEM than HL-60. This is likely due to the difference in the content of CD4-expressing cells between the two cell lines. In practice, model experiments in which CCRF-CEM and HL-60 were physically mixed at various ratios were carried out to determine the number of cells bound to the spots with immobilized anti-CD5 antibody. CCRF-CEM is known to abundantly express CD5, whereas no HL-60 express CD5 antigen. As shown in Fig. 4, the number of bound cells (CCRF-CEM) is proportional to the content of CCRF-CEM in the mixed suspension. This result indicates that the content of a specific cell type can be determined by antibody array-based cytometry. The possibility of quantitative analysis is on account of the following mechanism: Initially, plated cells settle on the spot to generate a stationary cell monolayer of a similar composition to the bulk suspension. Cells that express an 120 100 80 60 40 20 0

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Figure 4. Cell binding to an anti-CD5 antibody spot from the mixed suspensions of fluorescently-stained CCRF-CEM and non-stained HL-60 at various ratios. The graph shows the fraction of CCRF-CEM bound to the anti-CD5 antibody spot from the mixture of CCRF-CEM and HL-60 at various compositions. Percent CCRF-CEM was determined by dividing the number of fluorescently active cells bound to the spot by total cells bound from the suspension of 100% CCRF-CEM. The data are expressed as mean ± standard deviation for five different spots. (Reprinted with permission from [7]. Copyright 2007 Elsevier)

antigen can be captured on the specific antibody spot, and these cells are detected as bound cells. This implies that the ratio of the number of bound cells to the total cells in the monolayer (defined by the seeding density and the spot size) is identical to the composition of the specific cells in suspension. It is worth comparing the array-based cytometry with conventional FCM with regard to their performance. First of all, the number of target antigens that can be analyzed on a single array far exceeds that by FCM. As demonstrated above, our laboratory-scale antibody arrays are designed to analyze 22 different surface antigens. In the case of antibody array analysis, a few thousands of test cells are required for the assessment of a single surface marker. This number is almost comparable to the requirement for FCM. The limitations of antibody arrays include the inability of detecting multiple surface markers co-expressed on an identical cell.

4

High-Throughput Detection

The throughput of array-based quantitative analysis relies on the speed of bound cell determination. Although standard optical microscopes can be used for this A spacer

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Figure 5. The preparation of an antibody array and SPR monitoring. (A) Diagram showing an antibody array equipped with a parallel-plate chamber. (B) Photograph of a parallel-plate antibody array. (C) Setup of a homemade SPR apparatus used for imaging. (D) Drawing for the monitoring of cell binding by SPR imaging. Diagrams are not to scale. (Reprinted with permission from [8]. Copyright 2007 American Chemical Society)

purpose, it is time-consuming to determine the number of cells for the entire spots on an array. To improve the throughput of the analysis, we have exploited the technique by which bound cells can be determined simultaneously for multiple surface markers in parallel [8]. The technique is based on surface plasmon resonance (SPR) imaging (Fig. 5). SPR imaging is a surface-sensitive optical technique that allows to probe refractive index changes near the surface of a metal substrate with 2-D spatial information [11]. As demonstrated in Fig. 6A, CCRF-CEM interacting with the immobilized antibody generates SPR signals. Specific cells binding to an array have significant contributions to the elevation of the refractive index changes, whereas nonspecific cells with round shapes have extremely minor contributions. A IgG CD5

CD5 IgG

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Figure 6. Real-time SPR monitoring of affinity cell binding to the antibody array. (A) Brightness images recorded with the SPR apparatus of the antibody array at 0, 10, 20, and 30 min after cell injection into the parallel-plate antibody array that displayed antiCD5 antibody and nonspecific immunoglobulin G (IgG). The identification of the spots are indicated in the image. (B) The intensity of reflected light as a function of time after cell injection. Data are shown for two similar spots of both anti-CD5 antibody and nonspecific IgG and for the BSA-blocked surrounding region. (C) Parallel immunophenotyping of CCRF-CEM cells by SPR imaging. Data are subtracted by those for background regions and expressed as mean ± standard deviation for three spots. (Reprinted with permission from [8]. Copyright 2007 American Chemical Society)

In right of such mechanism, the SPR-based detection method allows the parallel monitoring of cell binding on multiple spots in real-time without washing out unbound cells (Fig. 6B). Monitoring without the washing-out procedure has great advantage to the accurate determination of bound cells because shear forces generated by fluid flow would cause the partial detachment of bound cells. As shown in Fig. 6C, CCRF-CEM bound to the spots with antibodies against CD3, CD4, CD5, CD7, CD8, CD38, CD45, and CD71. The intensities of reflected light were positively correlated with the mean channel fluorescence determined by FCM at a similar level of sensitivity.

5

Surface Marker Identification

Antibody array-based analysis was carried out for surface markers expressed on NSCs [5,6]. NSCs are recognized as one of the potential cell sources for the cell transplantation therapy of neurodegenerative diseases and brain and spinal cord injuries. The standard culture method, called neurosphere culture, can be used to obtain NSCs. However, neurospheres are always obtained as a heterogeneous population containing differentiated cells as well as undifferentiated stem cells. Moreover, information on specific surface markers of NSCs is limited, which makes it difficult to analyze the population for the quality control of NSCs and their isolation. Therefore, we applied antibody arrays to the analysis of surface markers expressed on NSCs. Due to the heterogeneity of test cells, surface marker expression cannot be simply predicted from cell binding patterns on an antibody array. To overcome this problem, we combined the binding assays with the immunofluorescent staining of nestin, an intracellular marker of NSCs and analyzed the correlation between nestin and surface marker expression. Figure 7 shows the representative results for dissociated neurosphere-forming cells. The fraction of nestin-expressing cells is relatively low for CD56, moderate for CD15, and high for CD57. These results indicate that nestin-expressing NSCs are found to a high content in a CD57-expressing population. In contrast, CD56-expressing cells were found to co-express nestin to the lowest extent among other cell types. As demonstrated above, the antibody array provides a versatile analytical tool that enables to correlate the multiple surface markers expressed on NSCs to their characteristics including intracellular marker expression. The application of this method is not limited to NSCs, but the same approach may be applicable to other cell types by employing an appropriate panel of antibodies.

CD15 (84±2%)

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Figure 7. Immunofluorescent staining of nestin intracellularly expressed in neurosphereforming cells bound to the microarray. (a–c) Phase-contrast and (d–f) fluorescence images are shown for cells on the spot of (a, d) anti-CD15, (b, e) anti-CD56, and (c, f) anti-CD57 antibodies. The fraction of nestin-expressing cells is indicated above the corresponding images. Bars: 50 μm. (Reprinted with permission from [6]. Copyright 2005 Elsevier)

6

Summary and Prospective

This paper introduces the new platform that enables the high-throughput analysis of surface marker expression. Due to its feasibility and simplicity as well as cost effectiveness, the array-based analysis provides a complementary method to the conventional FCM. The versatility of the array-based method will open a wide variety of medical and biotechnological applications, for instance, the quality control of stem cells for regenerative medicine and clinical diagnostic of infectious diseases.

Acknowledgements This work was supported by Kobe Cluster, the Knowledge-Based Cluster Creation Project (MEXT), Core Research for Evolutional Science and Technology (CREST/JST), and Consortium R&D Projects for Regional Revitalization (METI).

References 1. Wognum A.W., Eaves A.C., Thomas T.E., 2003. Identification and isolation of hematopoietic stem cells, Arch. Med. Res. 34, 461–475. 2. Riley R.S., Massey D., Jackson-Cook C., Idowu M., Romagnoli G., 2002. Immunophenotypic analysis of acute lymphocytic leukemia. Hematol. Oncol. Clin. North Am. 16, 245–299. 3. Eiges, R., Schuldiner M., Drukker M., Yanuka O., Itskovitz-Eldor J., Benvenisty N., 2001. Establishment of human embryonic sten cell-transduced clones carrying a marker of undifferentiated cells, Curr. Biol. 11, 514–518. 4. Bonner W.A., Hulett H.R., Sweet R.G., Herzenberg L.A., 1972. Fluorescence activated cell sorting, Rev. Sci. Instrum. 43, 404–409. 5. Ko I.K., Kato K., Iwata H., 2005. Antibody microarray for correlating cell phenotype with surface marker, Biomaterials 26, 687–696. 6. Ko I.K., Kato K., Iwata H., 2005. Parallel analysis of multiple surface markers expressed on rat neural stem cells using antibody microarrays, Biomaterials 26, 4882–4891. 7. Kato K., Toda M., Iwata H., 2007. Antibody arrays for quantitative immunophenotyping, Biomaterials 28, 1289-1297. 8. Kato K., Ishimuro T., Hirata I., Arima Y., Iwata H., 2007. High-throughput immunophenotyping by surface plasmon resonance imaging, Anal. Chem. 79, 8616– 8623. 9. Fujii Y., Anderson J.M., Matsuda T., 2008. Antibody-bound cell microarray for immunophenotyping: Surface modification and lymphocyte subpopulations, J. Biomed. Mater. Res. B Appl. Biomater. in press. 10. Belov L., de la Vega O., dos Remedios C.G., Mulligan S.P., Christopherson R.I., 2001. Immunophenotyping of leukemias using a cluster of differentiation antibody microarray, Cancer Res. 61, 4483–4489. 11. Arima Y., Ishii R., Hirata I., Iwata H. 2006. Development of surface plasmon resonance imaging apparatus for high-throughput study on protein-surface interactions, e-J. Surf. Sci. Nanotech. 4, 201–207.

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Chapter 34 Estimation of Endurance Performance of Hip Stems Using Finite Element Analysis Yu-Bong Kang1, Duk-Young Jung2, and Sadami Tsutsumi1 1. School of Dentistry, Nihon University, Tokyo, Japan 2. Department of Medical Engineering & Cardiology Institute of Development, Aging and Cancer, Tohoku University, Sendai, Japan

1

Introduction

Total Hip Replacement (THR) is successfully applied to the patients affected by hip diseases. The reliability of hip joint prostheses has been greatly increased during the last decade [1]. Almost of THR operation is being success, however, THR prosthesis failure appeared after the operation. For example, from 1992 to 2003, 511 cases of failures were reported to U.S Food and Drug Administration (U.S FDA) in America. Also in Japan, 37 cases of failure were reported to Ministry of Health, Labour and Welfare in Japan in 2006. After the THR operation, during daily life, daily activity generates dynamic stresses varying in time that effected to the changing loading condition, resulting in the mechanical fatigue failure of prosthesis. Aseptic loosening and localized acetabular or femoral osteolysis was influenced fix conditions. These are considered as the reason for prosthesis failure [2-3]. Therefore, it is important to ensure the hip prosthesis on fatigue failure. For reduce the THR prosthesis failure, prosthetic designs have been improved under many effort, and material technology and the adoption of new surgical procedure is necessary.

Fatigue testing for applications, the final step in the development of THR, have been performed to validate the safety of components for fatigue breaking before using in clinical. In femoral stems, fatigue test is usually performed according to ISO 7206 series [4]. ISO7206 series consist of various parts, under the general title implants for surgery – Partial and total hip joint prostheses. Among them, Part 4 and 8 established the evaluation method which describes the test apparatus and procedure to assess the fatigue resistance with torsion. Both procedure s require the stem fixed in the apparatus using bone cement. The cement constrains the distal end of the stem at 0.4CT from the center of the femoral head; the stem is oriented at 10º in adduction and for the test with torsion also at 9º in flexion. The applied cyclic load waveform varies from 300N to 2300N for the test, with a frequency ranging from 4 to 30Hz. A fatigue reliable stem must withstand 5×105 load cycles. However, laboratory fatigue test may take long time and high cost. Furthermore, if THR prostheses design parameters or materials were changed, the laboratory fatigue test had to restart. These are demerits of laboratory fatigue test. In order to solve these problems, using Finite Element Analysis (FEA) is suggested as one of the effective method. By using FEA, we may reduce the time and cost taken in the test and available to more rapid and flexible response with reasonable cost. In addition, we can predict a fatigue property and check the design parameters of modified THR prostheses. The purpose of this study was to develop an efficient method to evaluate the endurance performance and mechanical stability of THR prostheses using FEA. FEA was performed with two aims: first, evaluation of fatigue strength of THR stems; second, estimation of influence on THR prostheses design parameters.

2

Finite Element Models

Three dimensional finite element models of THR stems were constructed to analyze the stress distribution and estimate the endurance performance. Two different stem models were prepared based on the commercially produced stems used in clinical. In addition, in order to investigate the effect of stem size on the mechanical performance, four models were made by changing the scale of original models proportionally. Fig. 1 shows the finite element models used in this study. Model-A-S1 and S2 was made by shrinking the model-B with the CT length and stem diameter, respectively. Model-B-L1 and L2 was made by enlarging the model-B with the model-A. Table 1 showed the representative dimensions of finite element models. The number of nodes and elements of all

C A

B D

Model-A

Model-A-S1

Model-A-S2

C A

B D

Model-B

Model-B-L1

Model-B-L2

Figure 1. Finite element models of hip stems used in this study.

Table 1. The dimension of hip stem FE models. Model

A:Neck angle [degree]

B:CT length [mm]

C:Offset Length [mm]

D:Diameter of fix location [mm]

A

45

163

38

19

A-S1

42

145

35

17

A-S2

42

125

32

15

B

40

145

28

14

B-L1

38

163

32

15

B-L2

38

182

36

19

Table 2. The number of nodes and elements of FE models. Model

Nodes

Elements

A

24860

16135

A-S1

11406

7114

A-S2

10860

6738

B

8571

5201

B-L1

13423

8365

B-L2

20081

12918

models were shown in Table 2. All of finite element models were constructed with 10 node tetrahedron element as a linear isotropic structure. Boundary condition was based on ISO7206 standard Part 4 and 8 at year of 2002. Fig. 2 shows the boundary conditions applied to the finite element models. According to ISO7206, the stems were embedded at location of 0.4CT (CT: the length between center of head and the distal tip of stem) from center of bone head to fix location. In addition, the stem was oriented 10° valgus /9° flexion. The 5.3kN loading forces applied to the stem surface in a directional perpendicular to fix location.

F

9°

0.4CT

Fix line

10°

Figure 2. Boundary condition based on ISO7206-4 applied to FE models.

Material properties applied to finite element models showed in Table 4. Ti15Mo-5Zr-3Al was considered as the material used to stem. Table.4 Material properties applied to FE models. Material

Young's modulus

Poisson ratio

Ultimate strength

Fatigue strength

Ti-15Mo-5Zr-3Al

80GPa

0.3

1075MPa

580MPa

In this study, fatigue life of hip stems was estimated with the results of stress analyses. Maximum principle stresses were obtained from FEA, and utilized to estimate the fatigue life in numerical simulation. Stress life (S/N) approach was used for estimating the fatigue life of the stems. All fatigue analyses were performed according to infinite criteria (N = 107 cycles). As the loading condition in ISO 7206/4, a cyclic load from 300N to 2300N was applied to femoral component in a direction perpendicular to fix line with a frequency between 4Hz and 30Hz. In this study, however, the loading force was directly applied to stem surface because femoral component was not made in finite

element models. The alternating stress versus number of cycles (S-N curve) for implant materials obtained from laboratory fatigue test as shown in Fig. 3. The advantages of this approach is represents both initiation and propagation of cracks in the aggressive environment [6, 7]. ANSYS Workbench Ver.10SP1 (CYBERNET SYSTEMS Co, Ltd., Japan) was used to make the finite element models and conduct the FEA and fatigue simulations.

Figure 3. Fatigue property of Ti-15Mo-5Zr-3Al obtained from laboratory fatigue test.

3

Results

The results showed the different tendency by the stem design. Fig. 4 showed the results of the maximum principle stresses occurred to neck region and fix location in model-A and B. In model-A, the stress occurred to fix location was lower than that to neck region. In model-B, conversely, the stress occurred to fix location was higher than that to neck region. In addition, the stress in fix location of model-B was over the ultimate strength of the material used for the stem. This indicates the possibility of breaking in this location. In results of analyzing models (model-A-S1 and S2) changed the size of model-A, the stress occurred both to neck region and fix location increased compared with original size (model-A) as shown in Fig. 5. Especially, the stress occurred to neck region of model-A-S2 was 1.75GPa, it increased 300% of the value occurred to original size. The stress occurred to fix location was 567MPa, it increased 205%

Figure 4. The comparison with maximum principle stress between model-A and B.

Figure 5. The results of maximum principle stresses in Model-A.

of the value occurred to original size. This indicates the neck region is more dangerous for breaking when the stem size shrink with the same design of model-A. Also, in results of fatigue simulation, fatigue life of neck region decreased drastically in model-A-S2 as shown in Fig. 6.

Figure 6. The results of fatigue simulation in Model-A.

In results of models enlarged model-B, the stresses occurred to the fix location decreased as shown in Fig. 7. the stress in fix location of model-B-L2 was 650MPa, it decreased 40% of model-B with original size. However, the stresses occurred to the neck region decreased slightly. The stress in neck region of model-B-L2 was 548MPa, it decreased 14% of model-B with the original size.

Figure 7. The results of maximum principle stresses in Model-B.

Figure 8. The results of fatigue simulation in Model-B.

Fig. 8 shows the results of fatigue simulation in model-B. Fatigue life improved by enlarging the size of stem. However, CT length of model-B-L2 was 182mm by enlarging the stem size of model-B, it was too long even compared with model-A. Then, it might be unsuitable for applying to patients with averaged body type in Japanese.

4

Discussion

Total Hip Replacement (THR) is successfully applied to patients affected by hip diseases. As the THR are used in a human body which has particular condition over a long time, it is required a performance for enduring to cyclic loading with 4~5 times of a body weight. Thus, it is very important to estimate the fatigue strength of THR components before implant. However, it is impossible to evaluate the fatigue life of THR components implanted in a human body. Thus the fatigue strength of an artificial hip stem has been estimated using the test established by ISO7206. In this method, however, the fatigues strength of stems is just estimated if the stems endure 107 cycles under the test condition. In this study, the fatigue strength of hip stems was evaluated using FEA which is well known as an effective tool to evaluate the strength of structural objects. FEA was conducted on two commercial commercially stems used in

clinical. In addition, the effect of changing the size was investigated. In previous studies [5-6], as the neck region and distal end of stems were reported as the critical location of fatigue breaking. By FEA, the maximum principle stresses concerned with the fatigue breaking occurred to these locations were evaluated. In results, the stresses occurred to stem neck and fix location showed the higher value as a whole in all models. This was highly suggestive the breaking could be resulted in these locations. FEA was used for evaluation of fatigue failure related in stem designs [7-9]. However, there are no studies which evaluated an effect of stem size concerned with a physical type. Thus, in order to consider for improvement of the stem strength, the effect of changing the size of stems was evaluated by FEA. This could investigate the size effect to strength for applying to different body type. In results of model-A-S1 and S2 which was shrunk the size of model-A, the stresses occurred to neck region drastically increased compared with that to fix location. Model-A-S2 was made by shrinking the size of model-A to fit the stem diameter of ix location of model-B. Then, the CT length of model-A-S2 shorten more than that of model-B, the stress occurred to fix location decreased compared with model-B. However, as the offset length of model-A-S2 was longer than that of model-B, this was considered that the stresses in neck region increased. In contrast, the results of model-B-L1 and L2 which was enlarged the size of model-B, the effect of enlarging the stem size could not show for decrease the stress in the neck region. Model-B-L1 was enlarged of stem size with the CT length of model-A. The offset length of model-B-L1 was shorter than that of model-A and the diameter of fix location was smaller than that of model-A. Thus, the stresses occurred to neck region could not decrease and rather increase in fix location compared with model-A. Also, the stresses occurred to neck region of model-B-L2 was same and could not decrease in the fix location compared with model-A. Model-B-L2 was enlarged the size of model-B to fit the stem diameter of fix location in model-A. The offset length of model-B-L2 was shorter than that of model-A, but the CT length of model-B-L2 was longer than that of model-A. Model-A have been used for patients in averaged body type of male in Japanese. In contrast, model-B was designed originally depending on a petit body type in Asian female. The results of FEA could show that it was insufficient for improving the strength just to change the size of stem with design. In addition, as the bone shape around hip joint is different between male and female, the hip stem with same design cannot apply even changing the stem size.

In this study, FEA was conducted for evaluating the effect of changing the size with the same design on the mechanical strength of THR stem component. In the results, it was insufficient to advance the strength even changing the size with same design. However, there are some limitations in this study. Fatigues strength is long term performance on cyclic loading to THR components. However, in this study, fatigue strength was evaluated using results of static analysis even with fatigue property of material used to the stem. Fatigue life of a hip stem was resulted in long term dynamic loading affected in daily life. In addition, loading force affected in daily life would change the degree and direction. As such a loading case could not be considered in FEA condition, it is difficult to predict the fatigue life accurately using FEA. Therefore, the maximum principle stress concerned with the fatigue strength was investigated for evaluation. However, even by fatigue test method such as ISO7206, the stresses occurred to stem can’t estimate quantitatively. By using FEA, the critical location for breaking will predict with estimation of stress values.

5

Conclusion

In order to evaluate the effect of changing stem sizes for improving the fatigue strength of total hip replacement stem components, finite element analyses were conducted to two type models used in clinical for different body type. As the results, the neck and fix location of hip stem was critically to break by stress concentration and it is insufficient even to change the scale of stem with same design for enhancement of fatigue strength. In case of enlarging the stem sizes used for petit body type, the effect of changing the scale was found to improve the fatigue life. However, these models might not apply as the stem scale was unmatched to bone shape around hip joint. In contrast, in case of shrinking the stem sizes used for patients averaged body type, the stress occurred to neck region increased drastically and also fatigue life decreased. Thus, this model design could not apply to patients with petit body size without modification.

Acknowledgments This work was supported by New Energy and Industrial Technology Development Organization (NEDO) for R&D project planning and formation, project management and post-project technology evaluation functions.

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