Design and Manufacturing of Composites

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Design and Manufacturing of Composites : Proceedings of the Second Joint Canada-Japan Workshop On Composites, Concordia University, Montreal, Quebec, Canada, August 1998 Hoa, S. V. Taylor & Francis Routledge 1566767083 9781566767088 9780585326672 English Composite materials--Congresses, Manufacturing processes-Congresses. 1998 TA418.9.C6J573 1998eb 620.1/18 Composite materials--Congresses, Manufacturing processes-Congresses. cover Page i

Design and Manufacturing of Composites Proceedings of the Second Joint Canada-Japan Workshop on Composites CONCORDIA UNIVERSITY MONTREAL, QUEBEC, CANADA AUGUST 1998 EDITED BY S. V. Hoa Concordia University H. Hamada Kyoto Institute of Technology

page_i Page ii Design and Manufacturing of Composites aTECHNOMIC®publication Technomic Publishing Company, Inc. 851 New Holland Avenue, Box 3535 Lancaster, Pennsylvania 17604 U.S.A. Copyright ã 1998 by Technomic Publishing Company, Inc. All rights reserved No part of this publication may be reproduced, stored in a retrieval system, or transmitted, in any form or by any means, electronic, mechanical, photocopying, recording, or otherwise, without the prior written permission of the publisher. Printed in the United States of America 10 9 8 7 6 5 4 3 2 1 Main entry under title: Design and Manufacturing of CompositesProceedings of the Second Joint Canada-Japan Workshop on Composites A Technomic Publishing Company book Bibliography: p. Includes index p. 317 Library of Congress Catalog Card No. 98-86276 ISBN No. 1-56676-708-3 HOW TO ORDER THIS BOOK BY PHONE: 800-233-9936 or 717-291-5609, 8AM-5PM Eastern Time BY FAX: 717-295-4538 BY MAIL: Order Department Technomic Publishing Company, Inc. 851 New Holland Avenue, Box 3535 Lancaster, PA 17604, U.S.A. BY CREDIT CARD: American Express, VISA, MasterCard BY WWW SITE: http://www.techpub.com PERMISSION TO PHOTOCOPY-POLICY STATEMENT Authorization to photocopy items for internal or personal use, or the internal or personal use of specific clients, is granted by Technomic Publishing Co., Inc. provided that the base fee of US $3.00 per copy, plus US $ .25 per page is paid directly to Copyright Clearance Center, 222 Rosewood Drive, Danvers, MA 01923, USA. For those organizations that have been granted a photocopy license by CCC, a separate system of payment has been arranged. The fee code for users of the Transactional Reporting Service is 1-56676/98 $5.00 + $ .25 page_ii Page iii Proceedings of the Second Joint Canada-Japan Workshop on Composites held in Montreal, Canada, August 1921, 1998, organized by the Concordia Centre for Composites and the Kyoto Institute of Technology. ORGANIZING COMMITTEE Co-Chairmen S.V. Hoa, CONCOM, Concordia University H. Hamada, Kyoto Institute of Technology

Committee Members F. Ellyin, University of Alberta R. Gauvin, Centre des matériaux composites de St-Jérôme J. Hansen, University of Toronto A. Kalamkarov, Dalhousie University J. Lo, CANMET A. Poursartip, University of British Columbia P. Lee-Sullivan, University of New Brunswick T. Vu-Khanh, Université de Sherbrooke W. Wallace, National Research Council Canada N. Amiji, Toshiba F. Baba, Mitsubishi Electronic M. Fujii, Seikou Chemical Machine N. Ikuta, Shonan Institute of Technology Y. Kagawa, University of Tokyo H. Kawada, Waseda University I. Kimpara, University of Tokyo T. Kimura, Fukui University I. Narisawa, Yamagata University N. Takada, Daiwa Seikou N. Takeda, University of Tokyo T. Wakui, Kawasaki Steel M. Zako, Osaka University LOCAL COMMITTEE Chairman S.V. Hoa Committee Members S. Amiouny, Concordia University K. Demirli, Concordia University R. Ganesan, Concordia University L. Lessard, McGill University A.D. Ngo, École de technologie supérieure F. Trochu, École Polytechnique Montréal X.R. Xiao, Concordia University WORKSHOP SECRETARIES Dr. A. Yokoyama, Mie University, Ms. S. Mérineau, Concordia University & Ms. A. Nakai, University of Tokyo ADVISORY COMMITTEE & SPONSORS M.M. Dumoulin, Industrial Materials Institute, NRC N. Esmail, Concordia University R. Fews, Bell Helicopter Textron D.J. Taddeo, Concordia University Foreign Affairs and International Trade Bombardier/Canadair Military Division Canadian Association for Composite Structures and Materials page_iii Page v

CONTENTS Preface

xi

Thermoplastic Composites I Structures and Mechanical Properties of Injection Molded CF/LCP Composites A. Fujita and F. Baba, Advanced Technology R&D Center, Mitsubishi Electric Corporation, 8-1-1, Tsukaguchi-honmachi, Amagasaki, Hyogo 661-8661, Japan

3

Prediction of Residual Stresses in Continuous Glass Fiber/Polypropylene Composites Y. Youssef and J. Denault, Industrial Materials Institute, National Research Council, 75 de Mortagne, Boucherville (Québec), Canada J4B 6Y4

7

Impact Properties of Stampable-Sheet Made of Glass Fiber and Polypropylene K. Nagayama, Chemical Research Laboratory Technical Research Laboratories, Kawasaki Steel Corporation, 1 Kawasaki-cho, Chuo-ku, Chiba-shi 260, Japan K. Fujiwara, Polymer Mechanics Laboratory, Faculty of Textile Science, Kyoto Institute of Technology, Matsugasaki, Sakyo-ku, Kyoto 606, Japan

15

Manufacturing I Special Techniques Used to Manufacture Conductive Composite Laminate Structure for Unique Air-Borne Geophysical System M. Artus and V. Kohli, Avcorp Industries Inc., Composites Materials Division, 840 Vadnais St., Granby, Quebec, Canada J2J 1A7

21

Simulation of Resin Injection in Parts with Variable Porosity in Liquid Composite Molding F. Trochu, X.-T. Pham, Y. Benoit, J. Breard, J.-F. Remacle and R. Gauvin, Ecole Polytechnique, Montreal, Quebec, Canada

22

Design Considerations of an All FRP Highway Cargo Tank 30 A. Chatillon, Tankcon FRP Inc., 4250 Marcel Lacasse, Boisbriand, Quebec, Canada J7H 1N3 Metal Matrix and Smart Composites Wear Characteristics of Alumina Particulate Reinforced Aluminum Based Composites 41 J. Lo, CANMET, Dept. of Natural Resources Canada, 568 Booth St., Ottawa, Ontario K1A 0G1, Canada J. Li and M. Phaneuf, Fibics Inc., 568 Booth St., Ottawa, Ontario K1A 0G1, Canada T. Murayama, IMRA American Inc., 1044 Woodridge Avenue, Ann Arbor, MI 48105-9774, U.S.A. Design and Fabrication of Smart Composites for Static Shape Control H. Wang and C. K. Jen, Industrial Materials Institute, National Research Council Canada, 75 de Mortagne Blvd., Boucherville, Québec, Canada J4B 6Y4 M. Giray and S. Kalaycioglu, Canadian Space Agency, 6767 Route de I'Aéroport, Saint-Hubert, Québec, Canada J3Y 8Y9 S. E. Prasad, Sensor Technology Ltd., 20 Stewart Road, Collingwood, Ontario, Canada L9Y 3Z4

49

page_v Page vi Fracture Behavior of Adhesively Bonded Composite-to-Metal Lap Joints with Thick Adherends J. F. P. Owens, Boeing Canada Technology Inc., 99 Murray Park Road, Winnipeg, Manitoba, Canada R3J 3M6 P. Lee-Sullivan, Department of Mechanical Engineering, University of New Brunswick, P.O. Box 4400, Fredericton, New Brunswick, Canada E3B 5A3

57

Fatigue and Dynamic Failure Prediction of Tensile Fatigue Life for GFRP/Metal Adhesive Joints M. Nakada, T. Imai and Y. Miyano, Materials System Research Laboratory, Kanazawa Institute of Technology, Yatsukaho, Matto, Ishikawa 924-0838, Japan S. Sihn and S. W. Tsai, Department of Aeronautics & Astronautics, Stanford University, Stanford, California 94305-4035, U.S.A.

67

Prediction of the Fatigue Behavior of Graphite-Epoxy Laminates Using Artificial Neural Network A. D. Ngo and Y. O. Abdesslam, Ecole de Technologie Supérieure, Université de Québec, Montréal (Québec), Canada

75

Impact Compressive Failure of GFRP Unidirectional Composites 81 J. Yuan and N. Takeda, Center for Collaborative Research (CCR), The University of Tokyo, 4-6-1 Komaba, Meguro-ku, Tokyo 153-8904, Japan A. M. Waas, Department of Aerospace Engineering, University of Michigan, FXB Bldg., Ann Arbor, MI 48109-2118, U.S.A.

Thermoplastic Composites II Progressive Crushing of Compression-Molded Thermoplastic Composite Tubes H. Kawada, Department of Mechanical Engineering, Waseda University T. Honda, M. Takashima and H. Satoh, Graduate School of Waseda University, 3-4-1, Okubo, Shinjuku, Tokyo 169-8555, Japan

87

Development of "Fibro-Composites"Morphology of PBT/Polyolefin Blend K. Kitagawa, Kyoto Municipal Institute of Industrial Research, Chudoji, Simogyo-ku, Kyoto 600-8813, Japan H. Hamada and T. Semba, Kyoto Institute of Technology, Matsugasaki, Sakyo-ku, Kyoto 606-8585, Japan

91

Mechanism of Fatigue Fracture of Glass Fiber Reinforced Nylon 66 99 S. Odaka, T. Kuriyama, M. Kotaki and I. Narisawa, Department of Materials Science & Engineering, College of Engineering, Yamagata University, Jonan, Yonezawa City 992-8510, Japan Manufacturing II Thermal and Electron Beam Curing of Polymer CompositesA Comparison J. Raghavan and M. R. Baillie, Department of Mechanical and Industrial Engineering, University of Manitoba, Winnipeg, MB R3T 5V6, Canada V. J. Lopata, AECL, Whiteshell Laboratories, Pinawa, MB R0E 1L0, Canada

105

An Investigation of Autoclave Convective Heat Transfer 106 A. Johnston, P. Hubert, R. Vaziri and A. Poursartip, Composites Group, Department of Metals and Materials Engineering, The University of British Columbia, Vancouver, B.C., V6T 1Z4, Canada On the Processing and Testing of FRP Composites Incorporating Fiber Optic Sensors 114 A. L. Kalamkarov, S. B. Fitzgerald, D. O. MacDonald and A. V. Georgiades, Department of Mechanical Engineering, Dalhousie University, P.O. Box 1000, Halifax, Nova Scotia, B3J 2X4, Canada Textile Composites Predicting Shrinkage in Polyester Reinforced by Glass Fabrics V. Do-Thanh and T. Vu-Khanh, Université de Sherbrooke, Faculté des Sciences Appliquées, 2500 Boul. de I'Université, Sherbrooke, Québec, Canada J1K 2R1

125

133 Experimental and Numerical Analysis of Micro-Fracture Behavior in Textile Composites A. Nakai, Graduate School of Interdisciplinary Engineering Science, The University of Tokyo, page_vi Page vii 4-6-1 Komaba, Meguro-ku, Tokyo 153-8904, Japan H. Hamada, Faculty of Textile Science, Kyoto Institute of Technology, Matsugasaki, Sakyo-ku, Kyoto 606, Japan N. Takeda, Center for Collaborative Research (CCR), The University of Tokyo In situ Observation of Micro-Damage under Tensile Load: Single Fiber, Fiber Bundle 141 and Woven Fabric K. Nishiyabu and M. Zako, Osaka University, Graduate School of Eng., 2-1 Yamada-oka, Suita, Osaka 565-0871, Japan Impact Resistance of Multi-Reciprocal Braided Composites E. Kwan, X. R. Xiao and S. V. Hoa, Concordia Center for Composites, Concordia University, Montreal, Canada H. Wang, Industrial Materials Institute, National Research Council, Montreal, Canada H. Hamada, Kyoto Institute of Technology, Kyoto, Japan Analysis and Modeling

149

Free Vibration Analysis of Cantilevered Laminated Trapezoidal Plates 153 K. Hosokawa, J. Xie and T. Sakata, Department of Mechanical Engineering, Chubu University, 1200 Matsumotocho, Kasugai, Aichi 487-8501, Japan Mechanical Behavior of Sandwich-Type Composites with Waste of Fibrous Material 161 As Core Layer T. Kimura, Advanced Fibro-Science, Kyoto Institute of Technology, Matsugasaki, Sakyo-ku, Kyoto 606-8585, Japan Y. Kataoka, Department of Mechanical Engineering, Fukui University, 9-1 Bunkyo 3-chome, Fukui 910-8507, Japan Numerical Modeling Method of GFRP Laminate with Flexural Interphase T. Nishiwaki, ASICS Corp., 6-2-1 Takatsukadai, Nishi-ku, Kobe 651-2271, Japan S. Hayasaki and H. Hamada, Kyoto Inst. of Tech., Matsugasaki, Sakyo-ku, Kyoto K. Kitagawa, Kyoto Municipal Inst. of Indus. Res., Chudoji, Shimogyo-ku, Kyoto

165

Ceramic/Metals/Polymer Hybrid Composites Design and Applications of Metal/FRP Hybrid Structures 175 P. Kim, Shonan Institute of Technology, 1-1-25 Tsujido Nishikaigan, Fujisawa 251, Japan Mechanical Forming of Aluminum Matrix Composites H. J. McQueen Mech. Eng., Concordia Univ., Montreal, H3G 1M8, Canada E. Evangelista, Mechanics, Univ. of Ancona, 1-60131, Italy

183

Effect of Plasma Treatment on Surface of Glass Fiber for Plastic Based Composites 184 A. Nakahira, H. Akamizu and K. Kijima, Dept. of Chem. and Materials Tech., Kyoto Institute of Technology, Gosho Kaido-cho, Matsugasaki, Sakyo-ku, Kyoto 606-8585, Japan Y. Suzuki, National Industrial Research Institute of Nagoya, 1-1 Hirate-cho, Kita-ku, Nagoya 462, Japan S. Ueno and S. Nishijima, ISIR, Osaka University, 8-1 Mihogaoka, Ibaraki, Osaka 567, Japan Design and Applications Hierarchical Layerwise Higher-Order Finite Elements for Laminated Composite I. Kimpara, K. Kageyama and K. Suzuki, Department of Naval Architecture and Ocean Engineering, The University of Tokyo, 7-3-1 Hongo, Bunkyo-ku, Tokyo 113-8656, Japan

191

Design of an Implant and External Fixation for the Treatment of Bone Fracture in 199 Consideration of Mechanical Properties of Cortical Bone T. Hirai, INTEC-HIRAI Ltd., Miyamaki-nanasegawa, Kyotaname, Kyoto 610-0313, Japan Y. Watanabe, Kyoto Prefectural University of Medicine, Kawaramachi-Hirokoji, Kamigyo-ku, Kyoto 620, Japan A. Yokoyama, Kyoto Institute of Technology, Matsugasaki, Sakyo-ku, Kyoto 606-8585, Japan Design of a Polymer-Based Composite Container for the Long-Term Storage of 205 Radioactive Materials H. W. Bonin, V. T. Bui, P. J. Bates, J.-F. Legault and J. Y. S. D. Pagé, Department of Chemistry & Chemical Engineering, Royal Military College of CanadaCollège Militaire Royal du Canada, P.O. Box 17000, Station Forces, Kingston, Ontario K7K 7B4 page_vii Page viii Glass Fiber from Canada Resist in Acid Condition Y. Fujii, Seikow Chemical Engineering & Machinery, Ltd., Kukuchi 3-13-33, Amagasaki, 661-0977, Hyogo, Japan

212

Thermoplastic Lined FRP Dual-Laminate Composites for Corrosive ApplicationsAn 216 Overview P. Habib, C.P.F. Dualam Inc., 11750 J. J. Joubert, Montreal, Quebec H1E 7E7 Prediction of Failure in Unsaturated Polyester Reinforced by Plain Weave Glass Fabric H. Nguyen-Hoa and T. Vu-Khanh, Université de Sherbrooke, Faculté des Sciences Appliquées, Département de Genie Mécanique, 2500 Boul. de I'Université, Sherbrooke, Québec, Canada J1K 2R1

224

Poster Session Optimum Design of Composites with Functional Properties by Genetic Algorithm 235 A. Goto, Osaka Sangyo University, Faculty of Engineering, 3-1-1 Nakagaito, Daito, Osaka, Japan A. Yokoyama, Mie University, Faculty of Education, Kamihamacho, Tsu, Mie, Japan Curing of Thick Angle-Bend Thermoset Composite Part: Curing Cycle Effect on Thickness and Fiber Volume Fraction Variation M. I. Naji and S. V. Hoa, Concordia Center for Composites, Concordia University Montreal, Quebec, H3G 1M8, Canada

241

In-situ Cure Monitoring of Graphite/Epoxy Composites Using Fiber Optics and 249 Ultrasonics J.-Y. Chen and S. V. Hoa, Dept. of Mech. Eng., Concordia Univ., 1455 de Maisonneuve Blvd. W., Montreal, Quebec H3G 1M8 C.-K. Jen and H. Wang, Industrial Materials Institute, NRC, 75 de Mortagne Blvd., Boucherville, Quebec J4B 6Y4 Influence of Reinforcing Continuous Graphite Fibers, Environment and Physical Aging on the Visco-Elastic Properties and Fracture of a Thermoset Polymer Matrix J. Raghavan and C. I. Viswanathan, Department of Mechanical and Industrial Engineering, University of Manitoba, Winnipeg, MB R3T 5V6, Canada

257

258 Impact Fatigue Fracture of Glass Fiber Reinforced Thermoplastics K. Itoh, T. Kuriyama, M. Kotaki and I. Narisawa, Department of Materials Science & Engineering, College of Engineering, Yamagata University, Jonan, Yonezawa City 992-8510, Japan Vibration Damping Properties of Adhesive Joints of CFRP Laminates Y. Tanimoto, A. Tange and Z. Maekawa, Kyoto Institute of Technology, Goshokaido-cho, Matsugasaki, Sakyo-ku, Kyoto 606, Japan T. Nishiwaki, ASICS Corporation, 6-2-1, Takatsukadai, Nishi-ku, Kobe 651-22, Japan

262

The Effect of the Chemical Metamerism Polyolefin on the Friction and Wear of Bronze Powder Filled High Density Polyethylene A. Saito, H. Takahashi and I. Um, College of Science and Technology, Nihon University, 7-24-1 Narashino-dai Funabashi-shi, Chiba 274-8501, Japan

266

Mechanical Behavior of Polyolefin Composites Using Wastes of Fibrous Material As 274 Matrix and Reinforcement T. Kimura, Advanced Fibro-Science, Kyoto Institute of Technology, Matsugasaki, Sakyo-ku, Kyoto 606-8585, Japan Y. Kataoka, Department of Mechanical Engineering, Fukui University, 9-1, Bunkyo 3-chome, Fukui 910-8507, Japan Y. Kondo, Industrial Technology Center of Fukui Prefecture, 61 Kawaiwashizuka-cho, Fukui 910-0102, Japan T. Takahashi, Japan Polyolefins Co. Ltd., 3-2, Yakou 2-chome, Kawasaki-ku, Kawasaki 210-8548, Japan Identifying Delamination in Composite Beams Using Response Surface Methodology 278 Y. Shimamura, A. Todoroki, H. Kobayashi, H. Nakamura and K.-I. Iwasaki, Tokyo Institute of Technology, O-okayama 2-12-1, Meguro-ku, Tokyo 152-8552, Japan Effects of Flexible Interphase on Mechanical Properties of Unidirectional Carbon Fibre Reinforced Composites page_viii

282

Page ix K. Kitagawa, Kyoto Municipal Institute of Industrial Research, Chudoji, Shimogyou-ku, Kyoto 600, Japan S. Hayasaki and H. Hamada, Kyoto Institute of Technology, Matsugasaki, Sakyo-ku, Kyoto 606, Japan Impact Properties of Braided Composites with Flexible Interphase 288 A. Nakai, Graduate School of Interdisciplinary Engineering Science, The University of Tokyo, 4-6-1 Komaba, Meguro-ku, Tokyo 153-8904, Japan H. Hamada, Faculty of Textile Science, Kyoto Institute of Technology, Matsugasaki, Sakyo-ku, Kyoto 606, Japan K. Kitagawa, Kyoto Municipal Institute of Industrial Research, Chudoji, Shimogyo-ku, Kyoto 606, Japan N. Takeda, Center for Collaborative Research (CCR), The University of Tokyo Evaluation of Delamination Energy Release Rates by Layerwise Higher-Order Finite 296 Element K. Suzuki, I. Kimpara and K. Kageyama, Department of Naval Architecture and Ocean Engineering, The University of Tokyo, 7-3-1 Hongo, Bunkyo-ku, Tokyo 113-8656, Japan CAE for SMC Molding H. Hamada, T. Hasegawa and E. Tanigaki, Kyoto Institute of Technology, Matsugasaki, Sakyo-ku, Kyoto 606-8585, Japan H. Naito, Sekisui Chemical Co., Ltd., Minami-ku, Kyoto 601, Japan

300

Energy Absorption Capability of Braided Composites 304 H. Hamada and K. Kameo, Kyoto Institute of Technology, Matsugasaki, Sakyo-ku, Kyoto 606, Japan A. Nakai and N. Takeda, The University of Tokyo, 4-6-1 Komaba, Meguro-ku, Tokyo 153, Japan Stochastic Characteristics of Interlaminar Shear Strengths of Laminated Composites 308 C. Zhang, R. Ganesan and S. V. Hoa, Concordia Centre for Composites, Department of Mechanical Engineering, Concordia University, Montreal, Quebec, Canada H3G 1M8 Author Index

317 page_ix Page xi

PREFACE The second joint Canada-Japan workshop on Composites follows the success of the first Canada-Japan workshop on Composites held in Kyoto, Japan in August 1996. The second workshop held in Montreal in August 1998 furthers the objective of providing a forum for interaction on a large scale among researchers, engineers and scientists between Canada and Japan. This second joint workshop is co-organized by Professor Suong V. Hoa of Concordia Centre for Composites, Concordia University and Professor Hiroyuki Hamada from Kyoto Institute of Technology, Kyoto, Japan. Apart from technical presentations, the workshop also includes exhibits from companies and visits to the Industrial Materials Institute and to Bell Helicopter Textron Canada Ltd. The technical presentations cover a range of topics from thermoplastic composites, metal matrix composites, ceramic composites, smart composites, textile composites, manufacturing methods, fatigue and impact failure, analysis and modeling, and design and applications. Sincere appreciation is expressed to many organizations and individuals who have contributed to the success of the workshop. Included among them are the Canadian Department of Foreign Affairs, for the financial support for the workshop; the sponsorship from Canadair Ltd.; the support from the Industrial Materials Institute and Bell Helicopter Textron Canada Ltd.; Concordia Centre for Composites, Department of Mechanical Engineering, Concordia University and Kyoto Institute of Technology for the parametric support; Ms. Sophie Merineau for looking after all important details of the workshop and for the preparation for the publication of the proceedings.

S. V. HOA CONCORDIA CENTRE FOR COMPOSITES DEPARTMENT OF MECHANICAL ENGINEERING CONCORDIA UNIVERSITY, MONTREAL, CANADA H. HAMADA FACULTY OF TEXTILE SCIENCE KYOTO INSTITUTE OF TECHNOLOGY KYOTO, JAPAN page_xi Page 1

THERMOPLASTIC COMPOSITES I page_1 Page 3

Structures and Mechanical Properties of Injection Molded CF/LCP Composites Akihiro FUJITA and Fumiaki BABA Advanced Technology R&D Center, Mitsubishi Electric Corporation 8-1-1, Tsukaguchi-honmachi, Amagasaki, Hyogo 661-8661, JAPAN Keywords: Liquid Crystalline Polymer, Injection Molding, Mechanical Property, Carbon Fiber, Multi Layer Structure 1 Introduction The unique multi layer structure and mechanical properties of injection molded thermotropic liquid crystalline polymer(LCP) are well known. Although there are many studies in injection molding of LCP reinforced composites, only few study have been done about interaction between fiber and polymer in fiber reinforced LCP. In injection molding of thin plate with thickness of more than 1mm, LCP formed multi layer structure, typically consists five layers, namely two skin layers, two highly aligned intermediate layers and core layer. Polymer and fiber are highly aligned along the flow direction in the intermediate layer and are aligned perpendicular to flow direction in the core layer. In the skin layer, polymer are aligned to flow direction and fiber are aligned randomly. This study investigates the effects of gate shape and injection speed on fiber orientation in the injection molded this plates. The structure, fiber orientation and mechanical properties were studied. 2 Experimental Procedures 2.1 Materials The resins used in this study were a non-reinforced liquid crystalline polymer (VECTRA A950, Polyplastics Co.) and carbon fiber reinforced liquid crystalline polymer (VECTRA A230, Polyplastics Co., Vf=30wt%). These two kinds of pellets were mixed under dry condition. Carbon fiber content was 20wt% (A220). 2.2 Injection Molding Two kinds of mold used in this study is illustrated in Fig. 1. These were the plate cavity molds with side (sign: ''S") or film ("F") gates. Dimension of the plates were 55(width) ´ 70(length) × 1.0 (thickness) mm, as illustrated in Fig. 1. A reciprocating screw injection molding machine with a hydraulic accumulator system (V110/75V, Sumitomo Heavy Machinery Co.) were used in this study. Injection molding conditions are summarized in Table 1. Injection speeds were varied in 150mm/sec. (68cc/sec.) and 300mm/sec. (136cc/sec.). Here, 150mm/sec. and 300mm/sec. were regarded as low injection speed (L) and high injection speed (H), respectively. The specimens were distinguished due to the gate shape and injection speed, and each sign, "S" or "F", "L" or "H", were added to the name of the materials, respectively. For example, "A220SH" specimen was fabricated under high injection speed by using the

mold with the side gate. 2.3 Experiments Examinations of individual fiber orientation were performed under an optical microscope using surfaces prepared by the metallographic polishing technique. Specimens for the bending tests were cut out in parallel (machine direction) and perpendicular (transverse direction) to longitudinal direction of the plate, as illustrated in Fig. 1, in order to consider effect of cutting position of the specimen. We called "MD specimen" and "TD specimen" by cutting direction, respectively. The specimens were cut into strips with10mm width, as illustrated in Fig. 1. The bending tests were performed at testing speed 2mm/min and room temperature by using universal testing machine (Autograph, SIMADZU Co.). Span length was 30mm. page_3 Page 4 3 Experimental Results 3.1 Bending Properties Fig.2 and Fig.3 show relation between bending properties and cutting position of the specimen in each cutting direction. The specimens at both side edges exhibited higher bending moduli as compared with the specimens around the center to width direction in the MD specimens. In the specimens around the center to the width direction, the bending moduli were almost same. The bending moduli of the specimens with the film gate were higher than those of the specimens with the side gate in each position. Comparing to effect of the injection speed, A220FL and A220SL specimens under the low injection speed exhibited higher bending modulus as compared with A220FH and A220SH under the high injection speed, respectively. On the other hand, in the TD specimens the bending moduli of the specimens close to the gate were higher than those of the specimens around the edge regardless to the gate shape. Differences of the bending modulus due to the gate shape and injection speed were opposite tendencies to those of the MD specimens. Tendencies of the bending strengths to cutting position were similar to those of the bending moduli in all specimens. 3.2 Observations of Fiber Orientation States Fig.4 and Fig.5 show the optical microphotographs of cross section at the center of the plate along the MD direction in A220FH and A220FL specimens and polished planes at the center of the plate at each layer in A220FL and A220SL specimens. All specimens consisted of five macro layers; two skin layers, two intermediate layers and a core layer. In the film gate, the fibers oriented randomly in the skin and core layer, and to the flow direction in the intermediate layer. Whereas, in the side gate, the random fiber orientation was observed in the skin layer as well as the film gate. However, more random fibers in addition to the flow direction existed in the intermediate layer, as compared with the film gate. In the core layer, the fibers oriented in perpendicular to the flow direction. From these observations, the fiber orientations in the intermediate layer in the film gate and in the core layer with the side gate contributed to the bending properties of the MD and TD specimens, respectively. In all specimens, the fibers locally oriented in the flow direction at the side edges in the intermediate and core layers, and in perpendicular to the flow direction at the center to the width direction in the core layer. These local fiber orientation led to good bending properties at the side edges in the MD specimen and near the gate in the TD specimens. Differences of the fiber orientation due to injection speed were not appeared in each layer. However, thickness of intermediate layer at the low injection speed was larger than that at the high injection speed, as shown in Fig.4. 4 Discussions In this study, the result was that fiber orientation in each layer changed by the gate shape of the plate mold was very interesting. In the skin layer, extent of fiber orientation was small, because resin was immediately cooled by contacting to the mold. Therefore, it is considered that the fibers were aligned randomly in the skin layer regardless to the gate shape. The fibers were aligned to the flow direction by shear force in the intermediate layer, but resin flew in a radial manner from the gate at an initial step in the side gate. The random fiber orientation was observed in the side gate by this flow behavior. In the core layer, it is expected that resin flow showed diverging flow and speed of the resin flow was large. These flow behaviors is remarkable in the small gate as to the side gate, and lead to fiber

orientation in perpendicular to the flow direction. Therefore, the fibers were aligned in perpendicular to the flow direction in the side gate and random in the film gate. 5 Conclusion In this study, effects of injection speed and gate shape of mold on the structure and mechanical page_4 Page 5 properties of the carbon fiber reinforced LCP injection molded plates were investigated. Consequently, all specimens consisted of five macro layers; two skin layers, two intermediate layers and a core layer. The fiber orientation in each layer changed by the gate shape of the plate mold. The differences of the fiber orientation affected the bending properties of the plate. The differences of the mechanical properties and structures due to the injection speed was not remarkable as compared with the effects of the gate shape.

Fig. 1 Plate cavity molds with side and film gates. Table 1 Injection molding conditions. Injection Speed High Low Cylinder Temperature (°C) 320 320 Mould Temperature (°C) 100 100 Injection Speed (mm/sec.) 300 150 Injection Speed (cc/sec.) 136 68 Injection Time (sec.) 0.14 0.3 Holding Pressure (MPa) 25 50 Holding time (sec.) 1 3 Cooling Time (sec.) 30 30

Fig.2 Relation between bending properties and cutting position of the MD specimens.

Fig.3 Relation between bending properties and cutting position of the TD specimens. page_5 Page 6

Fig.4 Optical microphotographs of cross section at the center of the plate along the MD direction in A220FH and A220FL specimens.

Fig.5 Optical microphotographs of polished planes at the center of the plate at each layer in A220FL and A220SL specimens. page_6

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Prediction of Residual Stresses in Continuous Glass Fiber/Polypropylene Composites Younès Youssef and Johanne Denault Industrial Materials Institute, National Research Council 75 De Mortagne, Boucherville (Québec) Canada J4B 6Y4 Keywords: Thermoplastic composites, Polypropylene, Glass fiber, Interfacial strength, Residual stress, Stress relaxation Abstract Thermoplastic composite matrix material undergoes successive thermal phase transformations during the thermoforming process. The processing conditions control these transformations and hence are responsible for the final crystalline morphology, the interface quality, the mechanical properties as well as for the residual stresses developed in the composites. The objective of this work was to develop a model in order to be able to predict the development of residual stresses during the molding process of the continuous glass fiber/polypropylene composites, PP/G. For this purposes, process-induced residual stresses in thermoformed polypropylene/glass composite laminates have been experimentally evaluated by the measurement of cross-ply unsymmetric laminate curvatures. Two composite systems presenting different interfacial properties have been studied. Classical lamination theory model for residual stress prediction has been adapted to take into account the variation of the physical and mechanical properties of the matrix as a function of temperature and cooling rate and to account for stress relaxation either by the viscoelastic behavior of the matrix or by transverse cracking in the composite plies. This model incorporates the effect of the thermal history and processing kinetics (cooling rate) on the composite properties, particularly those controlled by the matrix microstructure and the interface quality. 1 Introduction The processing of thermoplastic composite materials induces inevitably thermal residual stresses due to the mismatch in the fiber and matrix coefficients of thermal expansion. These stresses result in dimensional distortions of the molded parts and may reduce their mechanical performance. The understanding of the development of these stresses, the identification of the variables controlling them may help in controlling these stresses and may also enable to predict their effects on part warpage and mechanical performance. Many investigations of the process-induced residual stresses topic have been carried out [16] but most of the reported work is related to the graphite/polyetheretherketone system. In this case, the most significant portion of the residual stresses builds-up between the glass transition temperature, Tg where relaxation effects are much less important resulting in increasing residual stresses with increasing cooling rates [3]. Polypropylene based composites are very different since their Tg is lower than the ambient temperature. Also, the fiber-matrix interaction in these composites is very particular due to the inert character of PP and brings particular behavior characteristics. The growing interest in this class of thermoplastic composites, particularly for automotive applications, justifies more research efforts in order to understand and characterize their behavior starting with the processing phase. In previous work on thermoforming of continuous glass fiber/polypropylene composites, it has been shown that processing conditions, particularly molding temperature and cooling rate, are responsible for the final matrix morphology and hence control the mechanical performance of the composite [7]. The effect of the cooling rate was found particularly significant in the development of residual stresses [8]. It has also been shown that residual stress predictive models should take into account temperature-dependent parameters and should account for stress relaxation phenomena in order to be more effective [8]. page_7 Page 8 In this work, process-induced residual stresses in thermoformed polypropylene/glass composite laminates have been experimentally evaluated by the measurement of cross-ply unsymmetric laminate curvatures. Two composite systems presenting different interfacial properties have been studied. Classical lamination theory model for residual stress prediction has been adapted to take into account the variation of the physical and mechanical properties of the matrix as a function of temperature and cooling rate and to account for stress relaxation either by the viscoelastic behavior of the matrix or by transverse cracking in the composite plies. This model incorporates the effect of the thermal history and processing kinetics (cooling rate) on the composite properties, particularly those controlled by the matrix microstructure and the interface quality.

The comparison of predicted and measured curvatures showed that the residual stress model should consider temperature-dependent parameters and relaxation phenomena in order to be comparable to experimental measurements. The model predictions, without stress relaxation and variable properties considerations, differ from the measurements not only by the stress amplitude but also by a predicted effect of the cooling rate on these stresses opposite to the observed tendency. The study has also shown that the residual stress build-up depends significantly on the cooling rate since this parameter controls the matrix morphology and the interface properties. The comparison of the two composite systems has emphasized the importance of the fiber-matrix interface quality in the development of the composite stiffness and strength. Weak interface composites could develop significant damage under thermal stresses during the processing. During the processing of thermoplastic composite unsymmetric laminate panel in a hot press, the laminate is kept flat until the end of the process. When the pressure is released (open mold), the laminate is free to shrink and warp due to the thermal forces and moments built-up during the cooling phase. The laminate mid-plane strains and curvatures, and ki, are calculated by solving the governing equation:

where A, B and D are the stiffness and coupling matrices of the laminate, and are the resulting thermal forces and moments respectively. These forces and moments, for a [0n/90n] laminate of thickness h are given by:

and

page_8 Page 9 where and aj are respectively the transformed stiffnesses and the coefficients of thermal expansion (CTE) of the ply. Ti and Tf are the initial and final temperatures of the temperature range over which residual stresses has built-up. For the considered [0n/90n] laminate, the solution of equation (1) for curvatures leads, in dimensionless form, to:

where Qij are the ply stiffnesses at the final temperature, Tf. These equations constitute the basics of the residual thermal stress model based on classical laminate theory (CLT). This model predicts anticlastic saddle shapes for unsymmetric laminates regardless the dimensions of the panel. Hyer [9] investigated experimentally the shapes of unsymmetric laminates and reported that for small thickness to in-plane dimensions ratios, unsymmetric laminates can take right cylinder curvatures. Hyer [10] developed an extended laminate theory (ELT) to account for the non-linear out-of-plane displacements when the thickness is small compared to the in-plane dimensions of the laminate. For a given thickness to in-plane dimensions ratio or when curvatures are measured on narrow strips, the solutions of both CLT and ELT are very close [3]. Since the curvatures studied in the present work are measured on narrow strips, the CLT model will be used. 3 Residual Stress Measurements

The thermoplastic composite considered in the present study is glass fiber reinforced polypropylene. Two different unidirectional prepreg tapes were used. The first one, named PP/G, consisting in a pure PP matrix reinforced with 52 wt% of E-glass fibers, and the second, named mPP/sG, consisting in a 50/50 blend of chemically modified PP with pure PP reinforced with 60 wt% of E-glass fibers coated with a specific thermoplastic sizing. The use of these two different materials aims at highlighting the role of fiber-matrix interface quality in the residual stress build-up. PP/G prepreg has a nominal thickness of 0.4 mm and mPP/sG tape has a nominal thickness of 0.25 mm. These two materials were obtained from BAYCOMP Canada.

Fig. 1. Temperature variation during the cooling phase of the molded panels from 200°C to room temperature Square panels (150 mm´150 mm) of cross-ply unsymmetric laminates were molded in a flat matched mold. Molding temperature was 200°C, molding pressure was 0.7 MPa, residence time at molding temperature was 5 min and three cooling scenarios were applied resulting in the time-temperature curves shown on Fig. 1. The average cooling rates in the temperature range of 140 to 80°C (covering the crystallization range) were evaluated respectively as CR=0.3, 12 and 45°C/min. Six [04/904] panels of page_9 Page 10 mPP/sG and six [03/903] panels of PP/G were molded. All panels took a saddle-like shape when demolded at room temperature. Three to five narrow strips 15 mm´150 mm were cut from each panel and their curvatures were measured. Fig. 2 shows the distribution of the measured curvatures normalized by laminate thickness (k´h) for both materials as a function of the cooling rate. The scatter in the curvature measurements may be attributed to the usual error sources in composites, e.g., precision of measurements, fiber misalignment, disorientation of plies and non-symmetry of ply thickness. These results will be compared to theoretical predictions.

Fig. 2. Normalized curvatures (k´h) as a function of the cooling rate measured on cross-ply unsymmetric laminates of PP/G (a) and mPP/sG (b) composites 4 Temperature-Time Dependent Behavior The mechanical and thermal properties of the composite systems are defined by the well-known rule of mixtures. The stiffness and the coefficient of thermal expansion of glass fibers can be reasonably assumed constant over the temperature range 23°C to 200°C. The temperature-dependent properties are those of the matrix and their variations are discussed here after. The matrix properties are sensitive to temperature and stress/strain history and hence, the composite transverse properties are also temperature-and-time dependent.

4.1 Mechanical Properties of the Matrix DMTA measurements have been carried out on both matrix materials under sinusoidal loading at a frequency of 1 Hz and R=0.1 and for temperature varying from -40°C to the complete softening of the material at around 150°C at a rate of 2°C/min. The results of these tests are shown in Fig. 3. It should be noted that the DMTA does not allow to scan temperature from the melt state back to room temperature. In order to verify the behavior of the material during cooling, complementary measurements of the compression modulus (linear combination of Young's and bulk modulus) using ultrasonic waves in heating and cooling at a rate of 2°C/min have been carried out, details of this method can be found in Ref. [11]. The variation of the compression modulus (which is proportional to Young's modulus) with temperature for PP is also represented in Fig. 3. The ultrasonic measurements show that the material softens while being heated until the melt temperature (Tm»165°C). It could be observed that the cooling path differs from the heating one only between Tm and the crystallization temperature (Tc»120°C). Since the residual stress calculations are limited to temperatures below Tc, the modulus variation during the cooling phase can be assumed identical to that measured during heating by DMTA without significant error. The values of the DMTA measured modulus at T=23°C (1860 MPa for PP and 1980 MPa for mPP) are in the same range than the Young's modulus values of polypropylene (from 1600 to 2000 MPa) obtained by tensile tests. Quasistatic tensile tests have been conducted on unidirectional specimens loaded in the transverse direction and on cross-ply laminates loaded at 45° angle. It was found that the page_10 Page 11 lamina transverse stiffnesses measured experimentally on the composites vary with the cooling rate used when processing the laminates. The slower is the cooling, the stiffer is the lamina and this was explained by the difference in the matrix microstructure [8]. However, these values are generally lower than the nominal moduli (calculated by rule of mixture) especially in the case of PP/G. The difference between the transverse stiffness measured directly on the composite and that calculated by the rule of mixture and nominal PP properties is certainly related to the fiber/matrix interface quality as well as to inter-spherulitic adhesion through the remaining amount of amorphous phase to bind crystalline units. In fact, the rule of mixture does not account for these components and assumes perfect bond between fibers and matrix. To take this parameter into account, the transverse stiffnesses of the lamina in the model evaluation will be those measured experimentally at room temperature. The variation of this properties with temperatures will be scaled to the trend measured by DMTA. Doing so, the interface contribution to the lamina stiffness will be implicitly accounted for.

Fig. 3. Modulus variation with temperature for PP and mPP matrices. Also, from transverse tensile tests, transverse strengths of the composite systems under study have been measured. Regardless the cooling rate, mPP/sG systems showed transverse strengths around 18 MPa. PP/G systems, due to the poor quality of the fiber-matrix interface, failed at stresses as low as 4 MPa. This low transverse strength value can be of prime importance when evaluating residual stresses. 4.2 Coefficients of Thermal Expansion of the Matrix

Fig. 4. Variation of the measured specific volume under constant pressure of 0.7 MPa and the calculated CTE with temperature for mPP Pressure-Volume-Temperature (pvt) tests have been conducted on mPP samples for three pressure levels and with temperature varying from 220°C to 40°C at a cooling rate of 1°C/min, a description of page_11 Page 12 the technique used can be found in Ref. [12]. The variation of the specific volume at the molding pressure of 0.7 MPa and for cooling cycle from 220°C to room-temperature at 1°C/min is shown in Fig. 4. The specific volume shows a marked shrinkage at 131°C corresponding to the crystallization temperature of the mPP cooled down at 1°C/min. After crystallization, the specific volume decreases almost linearly with decreasing temperature. The coefficient of thermal expansion of the tested material is extracted from this data using the following approximation [13]:

where n is the specific volume, T is temperature and To is a reference temperature. The calculated CTE values are reported in Fig. 4 and show mainly a linear variation with temperature. Results of PP are not shown but they are almost identical except for the crystallization temperature which is 123°C instead of 131°C for the mPP. 4.3 Stress-Free Temperature The stress-free temperature is the temperature at which thermal stresses start to build-up during the cooling phase. In the case of semi-crystalline thermoplastics, it is common to take the crystallization temperature as stress-free temperature. The crystallization onset temperature can be considered to account for the crystallization shrinkage, however Unger and Hansen has shown that the contribution of the crystallization shrinkage to the overall thermal stress can be neglected due to the very low stiffness value (Strain precedes modulus hypothesis) [3]. Hence, for the two matrix systems used in the present study and based on results of DSC characterization presented in Ref. [8], the stress-free temperature considered for calculations is the crystallization peak temperature and will be 131, 124 and 109°C for mPP/sG at slow, moderate and fast cooling rates respectively. For the PP/G system, these temperatures will be 123, 111 and 94°C for the same cooling conditions. 4.4 Stress Relaxation Considerations In order to take into account the effect of the viscoelastic nature of the PP matrix in the residual stress relaxation but without dealing with the complexity of this subject, a simple model used in the numerical calculation of residual stresses in injected parts [13] together with relaxation data PP found in literature [14] were included in the evaluation of the residual stress model. The basic principle is to evaluate the thermal stress build-up by incrementing the temperature (and time) from the stress-free temperature to room temperature by a constant step. For each temperature step, it consists in evaluating the increase of stresses in the individual plies and add it to the cumulative thermal stress. If the cumulative transverse stress (same as stress in the matrix) overpasses the equilibrium stress of PP for that temperature, the stress is reduced by a factor that is function of temperature and the time spent at that temperature. The relation between relaxed and unrelaxed stresses is [13]:

where the subscript m refers to matrix, the superscripts r and u refer to relaxed and unrelaxed states, t is the relaxation time and is the reduced time increment as defined in Ref. [15]. The parameters entering in the evaluation of equation (6) are those measured on PP by Ariyama et al. [14]. A lower limit for stress relaxation is defined by the equilibrium stress which is temperature dependent [14]. These calculations are repeated for each temperature/time step. Another stress relaxation mechanism that might take place in the studied systems is the transverse cracking of the single plies. The presence of transverse cracks has been verified by optical microscopy and typical observations are shown in Fig. 5. The PP/G samples molded with cooling rates of 0.3 and 12°C/min show obvious transverse cracks covering the laminate half-thickness and other initiated cracks that did not propagate to the surface of the laminate. Since the lamina transverse strength is very low, it is possible that the thermal stresses developed in the transverse direction of a ply overpass the transverse strength of the ply. In this case, transverse cracks will appear progressively relieving partially the residual stresses. A simple way to consider this effect is to set the transverse strength of the material considered as an upper limit for transverse stresses. Kim et al. studied this phenomena by Finite Element Analysis and showed some correlation between curvature and crack density in unsymmetric graphite/PEEK laminates [2]. page_12 Page 13 At the end of these iterative calculations, the resulting stresses are integrated to evaluate the resulting thermal forces and moments which are used to evaluate the curvatures.

Fig. 5. Optical micrographs showing complete transverse crack (a) and partially initiated crack (b) in PP/G sample 5 Model Validation The model described in the previous section has been applied to evaluate the curvatures of the PP/G and mPP/sG [0n/90n] laminates and the results are compared to the measured curvatures (Fig. 6). The predicted curvatures are marked Model 1, 2 or 3 depending on the way they have been evaluated. Model 1 does not consider any relaxation mechanism but accounts for temperature-dependence of the matrix properties, Model 2 includes viscoelastic relaxation while Model 3 take account of possible transverse failure as described previously.

Fig. 6. Comparison of measured and calculated normalized curvatures for PP/G (a) and mPP/sG (b) as a function of the cooling rate

From Fig. 6, it is clear that without consideration for stress relaxation phenomena, residual stress predictions (Model 1) are ineffective and unrealistic. The predicted curvatures decrease with increasing cooling rate while they show experimentally significant decrease for the slow cooling rate. The fact that these calculations consider temperaturedependent stiffnesses and coefficients of thermal expansion does not make them effective. page_13 Page 14 Calculations considering the stress relaxation due to the viscoelastic nature of the matrix material are closer to the experimental measurements for mPP/sG where the fiber-matrix interface contributes to the stress transfer. In the case of PP/G, characterized by weak fiber-matrix interface, the stresses built-up during the cooling phase are always lower than the equilibrium stress and hence, no viscoelastic relaxation can take place (Model 1 and Model 2 results are identical). The results of Model 3 draw a trend identical to that of experimental measurements. However, the absolute values are generally lower than the experimental ones. The shift in the calculated curvatures can be attributed to the approximate values of some variables used in the calculations and to the simplifications in the development of the model These results demonstrate the contribution of the viscoelastic nature of the matrix and transverse cracking in stress relaxation due to processing thermal cycle. Moreover, these results demonstrate the importance of the fiber-matrix interface quality in the overall performance of the composite. In effect, a weak interface composite can not stand the process-induced residual stresses and can incorporate initially transverse cracks that may be detrimental for its in-service performance. 6 Conclusion In this work, process-induced residual stresses in thermoformed polypropylene/glass composite laminates have been experimentally measured using cross-ply unsymmetric laminates on two composite systems presenting different interfacial properties. Classical lamination theory model for residual stress prediction has been adapted by taking into account the variation of the physical and mechanical properties of the matrix and by accounting for stress relaxation either by the viscoelastic behavior of the matrix or by transverse cracking in the composite plies. The model formulated this way incorporates the effect of the thermal history and processing kinetics (cooling rate) on the composite properties, particularly those controlled by the matrix microstructure and the interface quality. The results showed that the residual stress model should consider temperature-dependent parameters and relaxation phenomena in order to be comparable to experimental measurements. The study has also shown that the residual stress build-up depends on the cooling rate since this parameter controls the matrix morphology and the interface properties. The comparison of the two composite systems has emphasized the importance of the fiber-matrix interface quality in the development of the composite stiffness and strength. Weak interface systems did not stand the process-induced residual stresses. References [1] G. Jeronimidis and A.T. Parkyn, J. Compos. Mater., 22 (1988), p. 401 [2] K.S. Kim, H.T. Hahn and R.B. Croman, J. Compos. Tech. Research, 11 (1989), p. 47 [3] W.J. Unger and J.S. Hansen, J. Compos. Mater., 27 (1993), p. 108 [4] J.A. Barnes and J.E. Byerly, Compos. Sc. Technol., 51 (1994), p. 479 [5] J.T. TZENG, J. Thermoplast. Compos. Mater., 8 (1995), p. 163 [6] C. Wang and C.T. Sun, J. Compos. Mater., 31 (1997), p. 2230 [7] J. Denault and J. Guillemenet, International Plastics Engineering and Technology, 2 (1996), p. 1 [8] Y. Youssef and J. Denault, 42nd International SAMPE Symposium and Exhibition, Anaheim, CA (1997), p. 134 [9] M.W. Hyer, J. Compos. Mater., 15 (1981), p. 175 [10] M.W. Hyer, J. Compos. Mater., 15 (1981), p. 296 [11] A. Sahnoune, F. Massines and L. Piché, J. Polym. Sc., B, Polym. Phys., 34 (1996), p. 341 [12] P. Zoller and D.J. Walsh, Standard Pressure-Volume-Temperature Data for Polymers, Technomic Publishing Co., Lancaster (1995), p. 117

[13] K.K. Kabanemi and M.J. Crochet, International Polymer Processing VII (1992), p. 60 [14] T. Ariyama, Y. Mori and K. Kaneko, Polym. Eng. Sc., 37 (1997), p. 1 [15] S.L. Rosen, Fundamental Principles of Polymeric Materials, 2nd Ed., John Wiley & Sons, Inc., New York (1993), p.298343 page_14 Page 15

Impact Properties of Stampable-Sheet Made of Glass Fiber and Polypropylene Katsuhiro NAGAYAMA *1 and Kazutoshi FUJIWARA *2 *1 Chemical Research Laboratory Technical Research Laboratories, Kawasaki Steel Corporation, 1 Kawasaki-cho, Chuo-ku, Chiba-shi 260 Japan *2 Polymer Mechanics Laboratory, Faculty of Textile Science, Kyoto Institute of Technology, Matsugasaki, Sakyo-ku, Kyoto 606 Japan Keywords: Impact property, Stampable sheet, Glass fiber, Bundle, Polypropylene Abstract The effect of glass fiber bundle on the impact properties of KP-sheets was studied. The impact strength of KP-sheets, which contain 0 and 50 wt% GF bundle of whole GF content, were measured by instrumented charpy test. Up to the maximum load, both samples showed the same energy absorption, however, KP-sheet with 50wt% bundle showed higher total energy absorption. In the case of KP-sheet without bundle, after reaching the maximum load value, the load decreased drastically. On the other hand, in the case of KP-sheet with 50wt% bundle, after reaching the maximum value, the load decreased gradually and showed higher energy absorption than KP-sheet without GF bundle. The difference of energy absorption behavior of these samples can be explained by the fiber pull out model. The influence of GF bundle on energy absorption values of KP-sheets was calculated by using the interfacial properties obtained by fiber pull out test. 1 Introduction KP-sheet is a stampable sheet made of Polypropylene (PP) powder and chopped glass fiber (GF) by a paper making process1). Articles of KP-sheet flow molded under considerable high pressure, show uniform mechanical strength due to uniform flow of the reinforcing glass fibers as well as matrix PP during molding. KP-sheet has excellent mechanical properties and has been used for the molding of Jack-holder, Seat back frame, and so on2). The impact properties and flow mobility of KP-sheet was increased by the addition of GF bundle. In this study the effect of GF bundle on the impact properties of KP-sheet was studied in compare to the quasi-static mechanical property. 2 Experiments 2-1 Single Fiber Pull out Test Glass fiber/polypropylene interfacial properties were determined by the single fiber pull out test. Two kinds of glass fiber (Aminosilane-coupling and non-coupling treatment) and polypropylene were used. The specimen geometry of the single fiber pull out test is shown fig. 1. Test samples were prepared as following; Single glass fibers were set between two-polypropylene films (thickness 150um,

Fig. 1 Specimen geometry of single fiber pull-out test page_15 Page 16 width less than 1mm) with a cover glass on the top, having an aluminum spacer (thickness 120um). A 10g weight was put on the cover glass and heated to 200°C in the oven. The pull out test was carried out at 0.3mm/min. 2-2 Quasi-Static Flexural Test and Instrumented Charpy Test. Materials: Two kinds of KP-sheets (Sample A and B) were used, Sample A was made of 100wt% filament type glass fiber and Sample B was made of a mixture of 50wt% filament and 50wt% bundle type glass fiber. The glass fiber diameter and length of filament and bundle type are 13mm and 11mm respectively. Both KP-sheets are heated to 210°C and the heated substances are compressed in a mold under about 15 MPa of compressive pressure and flow molded articles with 3.8mm thickness was obtained. Quasi-static tensile test: The test was performed using a Shimadzu Autograph Machine (Type AFG 5000) and carried out according to JIS K7054 using A-type test specimens. The crosshead speed during tensile testing was 1mm/min. Instrumented charpy test: Notched charpy tests were carried out using an instrumented test system(Dynatup GRC 8250 type) according to ASTM D256 at an impact velocity 4m/sec. The fractured test specimens were observed using a CCD microscope. 3 Results and Discussion 3-1 Interfacial Properties of GF/PP In Fig. 3-1, the stress/dislocation curves of single fiber pull out test of two kinds of glass fibers (Aminosilanecoupling and non-coupling treatment) are shown. GF with amino-silane treatment had higher maximum value, which is interfacial shear strength, than that of non-coupling treated GF. After reaching the maximum value, the stress decreased and remained almost constant in the case of both samples. This constant stress represents the friction stress during the GF debonding from the PP matrix. The friction stress values, 2.2MPa is 55% of interfacial sheer strength of Amino-silane treated GF/PP.

Fig. 3-1 Interfacial strength of GF/PP measured by single fiber pull-out test This friction stress is a consideration of the radial stress due to thermal shrinkage.

These result suggest that there is an the effect of fiber debonding on the impact properties of GF/PP composite. A simple model calculation based on energy absorption by fiber debonding was carried out using the values obtained from the single fiber pull out test. 3-2 Quasi-Static Tensile Properties of KP-Sheets The tensile strength of KP-sheets A and B as a function of Glass fiber content are shown in Fig. 1. For prediction of the fiber reinforced composite strength, the Kelly and Tyson's model3) has been considered (Eq 3-1).

page_16 Page 17 where sf is the fiber strength and sm is the matrix strength. t is the interfacial strength, nf is the fiber volume fraction, L is the length of the fibers and Lc is the critical length. In the case of KP-sheets, all the fibers are longer than the critical length, so the formula reduces to (Eq 3-2);

Table 1 Glass fiber content and void measurement of used KP-sheets Sample GF Glass fiber content Void aimed (vol.%) GF(wt%) GF(vol.%) 20 19.2 7.7 0.3 A 30 28.8 12.4 0.0 100% 40 37.5 17.2 0.8 filament 50 49.1 24.7 2.1 60 59.2 30.3 10.1 20 19.3 7.8 0.3 B 30 28.7 12.3 0.4 50/50% 40 39.9 18.7 0.8 filament/ 50 48.9 24.9 0.9 bundle 60 59.6 32.5 3.0

Fig.3-2 Tensile strength of Sample A and B in comparison with Kelly and Tyson's prediction

The void content of used KP-sheets samples were calculated as shown in Table 1. hr was a parameter relating to the fiber orientation and the value of 0.5 was obtained by data fitting. For low GF content, sample A(100wt% filament) generate more voids than sample B(50/50wt% filament/bundle). But at high GF content Sample B showed higher void content because of insufficient of impregnation of PP into GF's structure. The comparison between this model prediction and measured strength is shown in Fig. 3-2. The void content has been calculated in these strength values. There is good approximation of the tensile strength up to 40wt%(20vol.%), forever, after that, the measured strength is lower than the predicted values. It is considered that an insufficient impregnation of polypropylene into glass fibers. Sample A and B revealed similar results and it suggests the small effect of glass fiber bundle on quasi- static strength 3-3 Impact Properties of KP-Sheets

Fig. 3-3 Impact properties of KP-sheets with different GF composition

Fig. 3-4 Fracture images of charpy test of sample A and B Fig. 3-3 shows the Load/Deflection curves of instrumented charpy test of two KP-sheets A and B. Up to maximum load, sample A and B showed the similar behavior, however, there is a big difference between two samples after maximum load. In the case of sample A(100wt% filament), the Load decreased drastically and showed smaller energy absorption than sample B. On the other hand, in the case of sample B, after reaching the maximum load, the load decreased slowly and achieved larger energy absorption than sample A. The adsorbed energy before maximum load is defined as E1 and after maximum load is E2, respectively. page_17 Page 18 The fracture images of sample A and B after charpy test is shown in Fig. 3-4. In Sample A, there are few and very short glass fibers are observed in fracture surface and it is considered that most of fibers are broken.In the case of sample B, not a few glass fiber bundles are observed in fracture surface and it appears to be pulled out from matrix PP. As a result, sample B achieved higher energy absorption in charpy test. The effect of glass fiber bundle on the impact properties of KP-sheets is discussed next paragraph. 3-4 Fiber Pull out Model The energy absorption of unidirectional chopped glass fiber dispersed in composite by the glass fiber pull-out is calculated by the following Eq. (3-3);

Figure 3-4 shows the predicted energy absorption for different glass fiber diameter 16, 50 and 100 versus fiber length. The measured values by fiber pull out test were used for calculation and Vf = 18.8(40 wt%). It can be seen the that a maximum energy absorption is predicted to occur when the fiber length is equal to double size of the critical fiber length Lc. With increase in diameter, the maximum energy absorption increased. In the case of KP-sheet with glass fiber bundle, glass fiber bundle behaves same as a larger diameter glass fiber. As a result, the

KP-sheet can achieve the higher energy absorption during notched impact test. It is clear that by the addition of glass fiber bundle, the impact properties of KP-sheet is increased.

Fig. 3-5 Energy absorption estimated by GF-pull out model (Effect of GF-diameter) 4 Conclusions This study has revealed that the quasi-static and impact properties of KP-sheets with 100%-filament glass fiber and 50%-filament/50%-bundle glass fiber; 1) There is good approximation of the quasi-tensile strength up to 40wt%, forever, after that, the measured strength is lower than the predicted values. 2) Glass fiber bundle doesn't effect on the quasi-static properties of KP-sheet. 3)KP-sheet with glass fiber bundle revealed higher energy in the charpy impact test. 4)The effect of glass fiber bundle on the impact property is explained by the fiber pull out model. 5 References 1) T. Takehara and H. Suginobe, Kawasaki-steel, vol.24, No2 (1992) p102104 2) H. Yoshitake, O. Nishimura, K. SE, Y. Araki, T. Sunada and H.Kubo, Plastic age, vol.94, No9 (1996) p124129 3) A Kelly and W.R. Tyson, J. Mech. Phys. Solids, vol.13, No9 (1965) p329350 4) J. L. Thomason and M. A. Vlug, Composites part A vol.28A, (1997) p277288 page_18 Page 19

MANUFACTURING I page_19 Page 21

Special Techniques Used to Manufacture Conductive Composite Laminate Structure for Unique Air-Borne Geophysical System Mike Artus and Vijay Kohli Avcorp Industries Inc., Composites Materials Division 840 Vadnais St., Granby, Quebec, Canada J2J 1A7 Abstract

This presentation will cover various design and fabrication techniques used by AVCORP in the development of the world's largest airborne geophysical system. This system's concept was developed by a company involved in providing airborne systems & conducting surveys for the mining exploration industry. This multi-geometry time domain electro-magnetic system package consisted of a towed instrument platform (structure developed by AVCORP), and an electronics package mounted in the helicopter. Avcorp team was given the challenging task of developing/fabricating the unique structure (two big hexagon shaped structures25 feet in span). These challenges included: selection of right materials, keeping the weight of the structure to a minimum, to provide a balance of stiffness & strength to the structure, minimize aerodynamic drag and to ensure that structure will stand stringent weather conditions. Also bigger challenges were encountered in developing special techniques to embed the transmitter coils in the large composite structure, mechanical inter-connection of segments and provisions for field serviceability. This presentation will cover the manufacturing aspects of this project. page_21 Page 22

Simulation of Resin Injection in Parts with Variable Porosity in Liquid Composite Molding by F. Trochu, X.-T. Pham, Y. Benoit, J. Breard, J.-F. Remacle, R. Gauvin Ecole Polytechnique, Montreal, Quebec, Canada Tel. (514) 340-4711, ext. 4280, Fax (514) 340-5199, Email: [email protected] Keywords: Liquid Composite Molding, finite element, draping, edge effect, simulation Abstract The manufacturing of high performance composite parts by liquid composite molding is becoming increasingly sophisticated because several process variants can be selected to inject preforms with high fiber contents in complex molds. Injection of resin in a fiber reinforcement is no longer only a closed mold type of operation, like in classical Resin Transfer Molding (RTM). Either the mold can be slightly opened at the beginning of the injection to increase the porosity of the preform and reduce the injection pressure, or the mold can be heated to decrease the viscosity of the resin. The use of one of these two types of process variants or a combination of them results in a faster and more efficient filling of the mold. Other ways to improve mold filling are connected with the configuration of the injection gates (central or line gates), the use of runners and the different types of sealing conditions of the mold. The effect of draping must also be accounted for in the numerical simulation for parts with curved geometry, because the shearing of a fabric on a complex surface affects significantly the permeability of the preform. In order to become an effective computer aided design tool, the analysis of liquid composite molding must integrate all these process variants in a comprehensive software simulation tool. This paper presents some of the scientific work necessary to address these questions. More specifically, the usage of runners and the draping of complex shapes will be illustrated by examples of injections performed with LCMFLOT, the new PC version of the injection software developed by the Applied Research Center on Polymers (CRASP) at Ecole Polytechnique of Montreal. Introduction Recent developments in the manufacture of composite parts by injection molding include a series of procedures that aim at producing in a consistent and more reliable way molded components with a higher fiber content (around 60%). These new process variants can be regrouped in two main families of injection techniques: (1) the porosity of the part may change either locally by the usage of runners or everywhere by allowing a motion or deformation of the upper mold wall; (2) the mold may be heated to reduce the viscosity of the resin and hence, facilitate the impregnation of the preform. These new approaches become increasingly popular not only in the aerospace industry for high performance composite parts, but also in the automotive industry for structural and body parts. These new process variants change the whole picture of the now classical ''Resin Transfer Molding" (RTM) process, which becomes only one particular case and in fact, the most simple way of injecting a fibrous preform known as closed mold injection. As presented in the abstract of the IMS - Intelligent Manufacturing Systems - project INTELMAN [1], the various injection techniques page_22

Page 23 now studied under the generic term of Liquid Composite Molding (LCM) are closed mold injection (RTM), non-isothermal LCM (heated mold injection), injection-compression, liquid resin infusion, vacuum assisted resin infusion (VARI), internal pressure resin injection (IPRI) and thermoplastics liquid composite molding. The aim of this paper is to draw a general sketch of LCM technology, discuss its most relevant issues and illustrate by means of examples how numerical simulation can be used effectively as a computer aided design tool in liquid composite molding. We begin by studying the injection of parts with varying porosity. Porosity can be either a function of space or of time, or in the most complicated case of both variables. The most simple change of porosity occurs when runners are connected with the injection gates in order to modify on purpose the shape of the resin flow. A second type of porosity variation is obtained as a result of the curvature of the part. Either the mat is compressed in regions with a small radius of curvature, or a fabric is draped on a complex surface which results in shearing of the fiber network. These two situations change locally the porosity of the part and produce significant variations of permeability which will, in turn, affect the shape of the resin front during mold filling. Porosity can also change in time, for example in vacuum assisted resin infusion when the flexible membrane covering the mold, under which the reinforcement is put under vacuum, is gradually deformed by the gravity driven resin flow. In this latter case, the change in porosity is spreading throughout the preform as it is progressively wetted by the resin. In injection-compression, porosity and hence, permeability change continuously in time everywhere in the reinforcement as the mold upper wall moves during the compression phase. Note that the motion of the mold wall can be controlled either by displacement or pressure, these two possibilities leading to different systems of equations as explained in Pham et al. [2]. In this paper, we concentrate on the two cases of spatial variation of porosity, when runners are used or when a fabric is draped on a complex surface. The resin generally used in liquid composite molding may be considered as a Newtonian fluid, so the resin flow in the fiber bed is governed by Darcy's law

where is the resin interstitial velocity, p the resin pressure, K the permeability tensor, w the porosity of the porous medium and m the dynamic viscosity of the resin. Injection with Runners The purpose of this study is to illustrate how the effect of runners in a RTM mold can be simulated by RTMFLOT software. The part was proposed by Northrop Grumman, a military aircraft company, where experimental injections have been carried out for a carbon fiber preform using a slow reacting resin. The part has a thickness of 0.13 " (3.302 mm). Its geometry is illustrated in Figure 1, and the runners are shown in Figure 2. The first set of simulations were performed at constant flow rate in order to compare simulations using Darcy's law with the results of the presimulation option of the software which is based only on the conservation of the resin mass. It was observed experimentally that the last point to be filled was the upper middle section of the top wall, where the vent was then located. page_23 Page 24

Figure 1 Geometry of the mold The above cavity includes two zones: - the cavity which contains the preform with a volume fiber content of 58%,

- two runners, i.e., free channels with unitary porosity. The illustration of these two zones and the mesh of the cavity are shown in Figures 1 and 2. A vent is located in the upper middle part of the cavity.

Figure 2 Mesh of the cavity and runners The preform is a carbon fiber fabric from Fiberaid (500 satin). The matrix resin is a slow reacting resin 997/3, also from Fiberaid. The resin is injected from two injection points located on each sides of the mold (Figure 1), with a constant flow rate of 1e-5 m3/s at a temperature of 416K. The viscosity of the resin is constant at 1 Pa.s. According to the information transmitted by Northrop Grumman, the permeability of the preform is kx = ky = 1.2 e -10 m2. In order to study the effect of the runners on the filling front, an equivalent permeability will be selected in the runners (between 100 and 10 000 times higher than in the preform). Note that for an infinite crack in a porous medium, the upper limit of the equivalent permeability is given by the formula page_24 Page 25

where d is the thickness of the cavity. Here the thickness of the cavity is 3.3e-03m, so the equivalent permeability is approximately equal to

This indicates that the permeability in the runners should be about 104 times higher than in the cavity. In order to show the critical influence on mold filling of the value of permeability selected in the runners, simulations are now conducted for permeabilities in the runners 100 and 1000 times higher than in the preform. Test 1 The first simulation (test 1) is conducted for a permeability in the runners 100 times higher than in the preform using the presimulation option of the software, i.e., the analysis is based only on the conservation of the resin mass in the cavity. TEST 1: - cavity : kx = ky = 1.2e-10 m2 - runners : kx = ky = 1.2e-8 m2

Figure 3 Flow front at 25 % filling (presimulation, test 1)

Figure 4 Flow front at the end of filling (presimulation, test 1) page_25 Page 26 The results of the presimulation option (filling analysis based only on the conservation of the resin flow rate) are shown in Figures 3 and 4. The total injection time is 134 s. On the contrary to what is expected, the difference between the flow velocity in the cavity and in the runners is low, and the end of filling does not occur at the vent, but exactly at the opposite extremity of the part. The equivalent permeability in the runners is not high enough to model their effect on mold filling. Test 2 In the next simulation (test 2), the permeability in the runners is set 1000 times higher than that of the preform. The analysis is also conducted with the presimulation option of the software. TEST 2: - cavity : kx = ky = 1.2e-10 m2 - runners : kx = ky = 1.2e-7 m2 The results of the presimulation in Figures 5 to 8 show a clear difference between the velocity in the runners and in the cavity. In this case, the flow occurs preferentially in the runners. Once the runners are filled, the resin flows in the cavity. The vent is correctly located where the end of filling occurs.

Figure 5 Flow front at 19% filling (presimulation, test 2)

Figure 6 Flow front at 41% filling (presimulation, test 2)

Figure 7 Flow front at 99% filling (presimulation, test 2)

Figure 8 Flow front at the end of filling (presimulation, test 2) page_26 Page 27 Test 3 In this simulation, Darcy's equation is solved by finite elements for the same 1000 times higher permeability in the runners as in test 2. We obtain in Figures 9 to 12 results very similar as in test 2. Even if the difference in velocity between the cavity and the runners is slightly lower for Darcy simulation than for the presimulation, the last part to be filled is still located in the vicinity of the vent. (Note that the results obtained with Darcy simulation should be more precise because the calculation takes into account Darcy's law.)

Figure 9 Flow front at 19% filling (Darcy simulation, test 3)

Figure 10 Flow front at 41% filling (Darcy simulation, test 3)

Figure 11 Flow front at 99% filling (Darcy simulation, test 3)

Figure 12 Flow front at the end of filling (Darcy simulation, test 3) Simulations performed with a permeability in the runners 104 higher than in the cavity are not shown here because no significant difference was observed from test 3. This shows that the formula of the infinite planar crack gives a fairly reasonable estimation of the upper bound of the equivalent permeability that should be assigned to the runners in order to reproduce the resin race-tracking effect along the mold edges. Moreover, these examples of simulations show that the presimulation analysis provides a good approximation of the resin front progression in time. page_27 Page 28 Draping In liquid composite molding before injection of the resin, the reinforcement is draped or stamped on the mold surface. Fabric layers are compressed between the two rigid parts of the mold. This can cause some displacement and shearing of the fibers that will affect significantly the local porosity of the preform, and hence the resin injection in the cavity and the final mechanical properties of the composite part. Very often only shearing occurs, so a geometrical draping software such as the one developed by Trochu et al. [3] or Long et al. [4] based on the fisherman's net algorithm can be used to determine the angles between the warp and the weft of the fabric. In the examples studied in this paper, the draping software of Nottingham's University is used to predict the fiber orientations in the part after draping [4]. This model is based on the assumption that fiber deformation is restricted to shearing only without any sliding between two intersecting fibers. This is only a first approximation because for some very complex shapes the only way to avoid wrinkling is to provoke inter-fiber sliding. The fiber position and orientation are obtained by a kinematic draping algorithm different from the mechanical approach used by Blanlot and Billoet [5]. The solution of the draping problem is based on the fisherman's algorithm, in which the fibers are allowed to freely rotate around nodes at each intersection between warp and weft. Shearing of the fiber network creates a change of porosity that affects significantly mold filling. The shear angle is defined in Figure 13 below.

Figure 13 Schematics of the shear and flow angles in a sheared fabric The different angles of Figure 13 are defined as follows [6]: a is the fabric shear angle, b the angle of the principal flow axis (main axis of the wetted ellipse in a central injection) and q = 1 - a is the new angle between warp and weft. The fiber volume fraction Vf after shearing can be corrected from the initial fiber volume fraction Vfq by the relation

The permeability is affected by shearing, as well as the principal directions of the flow which are also rotated. This perturbation makes the resin impregnation more difficult in one principal flow page_28 Page 29 direction (KX) than in the other (KY) for example. Note that the value of the angle b is always larger than the shearing angle. The results of draping for a spherical part are presented here as found in the work of Rudd et al. [6]. First the mesh of the part is constructed. Then the fiber orientations obtained with the draping software are illustrated in Figure 14 for a quarter hemisphere.

Figure 14 Fiber orientations and shear angle after draping for a quarter hemisphere (from [6]) The draping algorithm calculates the local fiber orientations required to construct the porosity map on the mesh. Then a new permeability is calculated in each element from the information provided by the curve representing the dependence of permeability in function of porosity. Local variations of porosity either due to runners or to draping of the fabric strongly influence mold filling. So they cannot be neglected in the numerical simulation of Liquid Composite Molding processes. The next step will be to account for time variations of porosity to model injection-compression for example. Acknowledgements The authors wish to thank Sophie Magdeleine for performing the simulations with runners, NSERC (Canada), FCAR (Quebec) for funding this research and C. Husmann, from Northrop Grumman, El Segundo, California for authorizing the publication of the simulation results with runners. References

[1] F. Trochu, G. Ziegmann, M. Hintermann, "Intelligent Manufacturing of Structural Composite Parts by Liquid Composite Molding: Project INTELMAN", abstract, IMS project, January 1998. [2] X.-T. Pham, F. Trochu and R. Gauvin, "Simulation of Compression Resin Transfer Molding with Displacement Control", accepted in Journal of Reinforced Plastics and Composites, 1998. [3] F. Trochu, A. Hammami, Y. Benoit, "Prediction of Fibre Orientation and Net Shape Definition of Complex Composite Parts", Composites: Part A, 27A, 319328, 1996. [4] A. C. Long, C. D. Rudd, M. Blagdon, K. N. Kendall, M. Y. Demeri >, Polymers and Polymer Composites, 4(5), 1996. [5] R. Blanlot, J. L. Billoet, >, Journées Nationales des Composites, 1996. [6] C. D. Rudd, A. C. Long, P. McGeehin, P. Smith >, Polymer Composites, 17(1), 1996. page_29 Page 30

Design Considerations of an All FRP Highway Cargo Tank Alain Chatillon Tankcon FRP Inc. 4250 Marcel Lacasse Boisbriand, Quebec, Canada J7H 1N3 Keywords: FRP, tank trailer, composite, cargo tank, thermoplastics Abstract As early as the 1950's Fiberglass Reinforced Plastic storage tanks were used to contain Hydrochloric acid with great success. Then in 1964, the first FRP cargo tank was manufactured to transport HCl. The first generation was based on the same design principals as steel lined cargo tanks. A self-supporting thin barrel reinforced with external stiffening rings. The running gear components were then welded on a steel subframe that was attached to the FRP barrel. Problems associated with the FRP external stiffening rings were detected as these first cargo tanks were used. In 1974, the stiffening method was changed and a second generation was introduced. This new material offered a better method to stiffen the barrel instead using external stiffening rings a lightweight material was added during construction substantially increasing the stiffness of the barrel. Using a sandwich construction offered other advantages such as increased thermal resistance, and dual containment. Since, the mid 60's in excess of 40 millions road miles have been logged on FRP cargo throughout North America without any incident. To address the recent transportation trends we have combined high performance thermoplastic with the proven performance of the FRP cargo tank for transportation of high purity chemicals. page_30 Page 31 Introduction The use of Fiberglass Reinforced Plastics as a material of construction was documented (1) as early as 1954 in several fields. Even at that time, the characteristics of FRP were recognised in fields such as automobile, aviation, transportation (Figure 1), boats, bathtubs, electrical parts, and pipes. In the early 60's, a Canadian transportation company from Montreal also turned to FRP. This time, its inherent corrosive resistance and structural integrity were required. The intention was to fabricate an all FRP highway cargo tank that would be used to transport corrosive products such as Hydrochloric acid. At that point, rubber lined steel highway cargo tanks were already used in this very corrosive service since the late 40's. Another type of highway cargo was considered one that would not deteriorate if it came in direct contact with the corrosive product and also offered a greater toughness than that of steel rubber lined cargo tanks.

Fiberglass Reinforced Plastics commonly know as FRP, storage was being used very successfully to store corrosives such as Hydrochloric acid. The chemical resistance of FRP in this service had already been proven and was associated with very little maintenance. The next logical step was to manufacture a highway cargo tank from the same material. In the United States, Fiberglass Reinforced Plastic was used to manufactured a transport truck for milk (Figure 1), in the mid 50's. In Canada, the first self-supported FRP highway cargo tank to transport corrosives was manufactured in 1964, as shown in Figure 2, a current picture of one these earlier models. This new cargo tank was designed to integrate the existing working parameters of the transportation industry. Therefore the first design was modelled on existing cargo tanks. The First Generation The Unites States Department of Transportation published a set of regulations that specified minimum requirements based on the type of product transported. Even if not required in Canada, this information was certainly valuable. Under the rules and regulations enforced at that time, when transporting a corrosive, such as hydrochloric acid, an MC-312 class cargo tank would have been specified. These regulations indicated the minimum construction requirements, the accident damage protection required and also the minimum testing parameters. Unfortunately, only ferrous materials were considered, such as steel and aluminium; FRP was not covered under these regulations. Therefore, a detailed analysis was conducted with page_31 Page 32 regards to the design of an all FRP highway cargo tanks. As in the case of ferrous type cargo tanks, the dynamic loads, the applied pressures, and all load distributions were required to determine the barrel construction. Then safety requirements enforced for the MC-312 class of ferrous type cargo tanks such as roll over protection, piping, and emergency venting were evaluated. The design adopted was a 4700 I.G. tandem (2 axles) highway cargo tank. The internal diameter was 60'' and the overall length was 38'-0". The circular shape was reinforced with 8 external stiffening rings with a profile of 2" ´ 2" made of FRP/foam filler. The roll over requirements was met by extending the stiffening rings adjacent to each inlet/outlet. The FRP barrel was designed as a self-supporting structure, but provisions were required to attach components such as axles, upper plate, jacklegs and other items to the barrel. This was achieved by fabricating a steel subframe that was used to transmit the forces from the running gear components to the structural FRP barrel. This steel subframe was built to the exact outside diameter of the FRP barrel, then the running gear components i.e. axles, suspensions, and the bumper were welded to the subframe via a suspension frame; also added were the upper plate, and the landing gear. To compensate for the different coefficient of thermal expansion between the steel and the FRP, a layer of rubber sponge was set at each saddle location between the steel subframe and the FRP barrel. Stainless steel bands were then used to secure the barrel onto the steel subframe. This method of attachment could address maintenance issues, if required, both major components could be unbolted and maintained as needed. This system also ensured that only visual inspection could reveal any deterioration of the steel components. Also, the rubber sponges permitted a uniform distribution of the dynamic loads imposed on the barrel during transportation. This first generation was built for approximately 10 years. As this design was used, it was noticed that the method of stiffening the barrel was creating problems. It was found that hairline cracks were present at the junction area of the external stiffening ring and the barrel wall. Typically, the mid section external rings were the first to have these hairline cracks and the rings towards both heads seldom showed this defect. This was certainly caused by high elongation at the ring/barrel interface due to flexing from the effect of the road during normal operation of the cargo tank. As was previously indicated, the initial design was modelled on existing equipment, consequently using external stiffening rings to solidify a barrel was selected to stiffen the barrel as it was a common practice for steel highway cargo tanks. page_32 Page 33 When using this method, the objective is to increase the stiffness of the barrel without unnecessarily increasing the thickness of the barrel wall.

However, when the analysis is done using steel, stress is typically the governing factor; if only stress is considered when using FRP, a situation can exist where the strain is very high which could lead to microcracks that can grow into hairline cracks. This is reasonable considering that the Young modulus of elasticity of steel is approximately 30 millions and that of FRP is 1.2 millions. Consequently, the strain exerted at the cylinder cross section between the stiffening rings, could not necessarily be a major consideration when fabricating steel barrels but must certainly be looked at when building an FRP barrel. The Second Generation Addressing the hairline cracks issue lead to a second generation of FRP highway cargo tanks which was introduced in the mid 70's (see Figure 3). The cracks were found to be a direct result of the discontinuities created by the use of stiffening rings; consequently, another method to stiffen the barrel was needed. A sandwich construction was adopted. Basically, a pre-determined thickness of lightweight material is incorporated into the construction of the FRP barrel. The total wall stiffness is significantly increased because this layer of material when inserted near the mid point of the barrel construction, moves apart two riding sections of FRP. Using this sandwich construction inherently offered some other very attractive features. The foremost advantage was to avoid hairline cracks associated with the external stiffening rings. The thermal insulation of the barrel was considerably increased. Finally, the accident protection characteristics were increased because one barrel is completely encapsulated into another. This level of accident protection is not typically offered in any other types of barrel construction. The roll over protection was also modified to increase the protection aspect. Even if roll over bars was used on most cargo tanks, it was felt that during a roll over the connections or piping was not completely protected and could be at risk. To reduce the possibility of connection or piping damage, a full square box was incorporated into the design at every inlet/outlet. This second-generation design has been used very successfully since the early 70's and to date we have documented in excess of 30 millions road miles without one incident relating to the material of construction or to the design. page_33 Page 34 Where Is FRP Going from Here? Past experience can certainly attest to the viability of an all FRP cargo tank, but as the concentration and purity transported are continually increased, we have also looked at using thermoplastic for corrosion resistance. It is readily known that simply using thermoplastics to construct the barrel could not withstand the forces imposed on a highway cargo tank. However, if the thermoplastic is used only for its inherent corrosion resistance and the proven FRP tank is used for its structural integrity, this third generation could be of great benefit. An extensive R&D effort (2) was undertaken in 1993 to develop this type of cargo tank. The first step was to understand how an all FRP highway cargo tank trailer reacted to the various types of loading imposed on the cargo tank during its life. The scope of the R & D program was to quantify the strain and acceleration of an FRP tank trailer when loaded and also pulled on the road. This information would then be used to set the parameters for laboratory testing of several possible constructions samples of the new cargo tank. Also, this data would be used to develop a finite element model. Several tests were conducted on new trailers and also on old trailers (see Figure 4). The dynamic forces and the acceleration forces were recorded as the trailer was pulled down a pre-determined circuit. The circuit was selected to represent severe road condition. Several strain gauges (0-90-45) were located inside and outside of the barrel to measure the strain levels, acceleration was measured by accelerometers; both type of devices were plugged to a tape recorder which recorded the complete circuit. The information was viewed and the maximum strains and acceleration were extracted. After laboratory testing of test plates and finite element model; we manufactured the first fluoropolymer lined FRP highway cargo tank trailer in North America, as indicated in Figure 5. Conclusion The first generation of FRP highway cargo tanks were modelled around exiting equipment i.e. rubber lined steel cargo tanks. The use of external stiffening rings proved to be questionable at best. However, the concept of using a steel subframe to page_34

Page 35 distribute the loads from the running gear components to the structural barrel was proven to be best method. Also the possibility of unbolting the barrel from the running gear was considered essential, as maintenance was required on the steel components after several years in a corrosive environment. As the second generation was introduced, it was proven that the use of balsa wood to stiffen the barrel wall in lieu of stiffening rings was much more suited to FRP. It also offered very distinctive characteristics; namely, dual containment (which no other type of cargo tank presently offers), and also increased thermal insulation. This secondgeneration design can easily be combined with thermoplastic such as fluoropolymers to created another type of cargo tank, one that structurally has been proven and also offers practically unlimited chemical performance. Time will dictate if this high performance cargo tank will grow into a third generation. References (1) Ralph H. Sonneborn, Fiberglass Reinforced Plastics, P3, 1954. (2) S.V. Hoa P. Ouillet, NSERC research project, 1993.

Figure #1 FRP milk truck page_35 Page 36

Figure #2 FRP highway cargo tank - built in 1965 - Still in service today

Figure #3 Second generation FRP highway cargo tank page_36 Page 37

Figure #4 Strain gauges and accelerometer locations

Figure #5 Inside the First Fluoropolymer lined FRP highway cargo tank built in North America page_37 Page 39

METAL MATRIX AND SMART COMPOSITES page_39 Page 41

Wear Characteristics of Alumina Particulate Reinforced Aluminum Based Composites J. Lo1, J. Li2, T. Murayama3 and M. Phaneuf2 1CANMET, Dept of Natural Resources Canada, 568 Booth St., Ottawa, Ontario K1A 0G1, Canada. 2Fibics Inc., 568 Booth St., Ottawa, Ontario K1A 0G1, Canada. 3IMRA American Inc., 1044 Woodridge Avenue, Ann Arbor, MI 48105-9774, USA. Abstract Al-1Mg aluminum alloys containing 30 vol% Al2O3 and cast iron (A48 class 30) were evaluated in dry sliding wear against AISI 52100 steel using a block-on-ring apparatus. The wear resistance of the aluminum composite was found to be much superior than the cast iron in the range of high sliding speeds (>1.5 m/s) under an applied load. The cross-section of test samples, were studied using the Focus Ion Beam Microscopy and Electron Microprobe Analysis to determine the wear characteristics of the composites tested at various sliding speeds. Attention was paid to the interfacial characteristics between the particles and the matrix alloys, and an attempt was made to correlate the wear behavior of the materials to their microstructure. Introduction The demand for light weight and high performance in next generation vehicles has led to the search for new materials meeting the stringent property requirements. Metal matrix composites (MMCs) have been considered suitable for automotive components because of their low weights and especially superior wear resistance. The use of discontinuously reinforced aluminum for automotive applications, such as pistons, cylinder liners, and disc brakes, has been demonstrated in the past. With the development of more efficient manufacturing processes and lower material costs, a renew interest in MMCs has developed. Presently, there are several routes for the fabrication of aluminum composite automotive components. The most common approach is the one employed by Duralcan. This process involves the making of composite ingots by stir casting, followed by sand casting of components. In this case, the as-cast components are made of 100% composite material. The advantage of this process is low fabrication cost, but the final machining of composite components could be costly. Other more common fabrication techniques are page_41 Page 42 pressureless infiltration and squeeze casting, both processes require the making of particulate performs and followed with infiltration of molten aluminum into the preforms to form near net-shape composite rotors. Compared to stir casting, both pressureless casting and squeeze casting have the flexibility of using different types of aluminum alloys and incorporating higher volume (>30 vol%) fractions of reinforcement. Further, the unique advantage of reinforcing only the selected region of a component, allows a reduction in the machining costs. In order to fully exploit the advantages of composite materials for wear resistant applications, it is important to understand the wear characteristics of composite materials. Earlier investigations on the tribology of composites covered mostly the areas of abrasive and sliding wear [13], with a little focus on the microstructure/property relationship. In this work, the wear property and mechanisms of squeeze cast aluminum 1% magnesium reinforced with 30 vol% alumina particulates (Al2O3/Al) were evaluated. In addition, the wear properties of Al2O3/Al are compared with those of cast iron (A48 class 30). Experimental Techniques Materials The materials tested in this work are Aluminum 1% Magnesium reinforced with 30 vol% alumina particulates (Al2O3/Al) and cast iron (A48 class 30). The Al2O3/Al composite was made with the squeeze casting technique at CANMET. The Al2O3 particulates were supplied by Alcoa, and the average particle size was 15 mm. The tensile strength of Al2O3/Al in the as-cast condition was 251 MPa. The microhardness of the composite was measured. Rockwell indentations were made on polished surfaces of the composite using a load of 100 kgf. Five indentations

were measured, and the average hardness for Al2O3/Al is RB51.6. As for the counterpart wear surface, AISI52100 steel was used. The Vickers hardness of this steel is Hv 950. Wear Testing Dry sliding wear tests were performed using a block-on-ring apparatus. The slider ring (width 12mm and outer diameter of 38mm) was an AISI 52100 bearing steel. All test specimens were machined to a dimension of 5mm ´ 10mm ´ 10mm, and the narrow rectangular face with the 5mm ´ 10mm dimension was put in contact with the slider. Wear surfaces were polished up to 600 grit SiC paper. Dry sliding tests were performed using sliding speeds of 0.25, 1.5, and 2.5 m/s. In all cases, a sliding distance of 700m was used and a normal load of 2.1 Kgf was applied. The wear damage on the specimens was characterized using optical microscope, Micrion 2500 Focused Ion Beam System and Cameca SX-50 Electron Probe Microanalysis. page_42 Page 43 Results and Discussion Table 1 lists the weight and volume losses of cast iron (A48 class 30) and Al2O3/Al composite tested at various sliding speeds with a counter material of an AISI 52100 bearing steel. Both materials showed that the weight and volume losses are sliding speed dependent. In the case of cast iron, as expected in a typical wear test, the weight/volume loss increased with the increase in sliding speed. As for the composite material, it behaved differently. At both low (0.25m/s) and high (2.5m/s) sliding speeds, the weight/volume losses of the composite are high. And a saddle point occurred at an intermediate sliding speed of 1.5m/s. Comparing the wear behavior of these two materials, it is clear that the weight loss of composite is much lower than the cast iron at high sliding speeds. In order to have a better understanding on the wear characteristics of composite materials, cross sections of wear track surfaces were prepared, using a Microion 2500 high resolution Focus Ion Beam system (FIB). The advantage of using the FIB system is that it provides stress-free cross-sections with a highly focused Ga+ ion beam, thus eliminating some of the common artifacts (such as sub-surface crackings) induced in conventional sample preparation procedures. In addition, it allows in-situ high resolution imaging (~4nm). Figure 1 shows a FIB image of an Al2O3/Al composite sample tested at a sliding speed of 0.25m/s. The dark areas represent the Al2O3 particles; while the top layer with the white phase, is a mechanically mixed layer. Such a layer was observed in all composite samples tested in this work. This layer consists of heavily deformed aluminum alloy matrix interspersed with Al2O3 particles. Iron content from the slider ring was found to have transferred into the mixed layer. The depth of such layer was neither uniform nor continuous along the width of the cross section, and it appears to reach the maximum at the center of the track. The chemical composition of the mechanical mixed layer was identified using the electron microprobe analysis. Table 2 provides the weight percentage of oxygen, iron and aluminum with respect to the distance from the wear surface. It is apparent that a substantial amount of iron is being transferred from the slider ring to the test sample. And the amount of iron is relatively rich in region up to 5 mm beneath the wear surface. In addition, oxygen content was also high even at 6.25 mm beneath the wear surface. The above information indicates that extensive mechanical mixing was happening during wear testing. Iron was being transferred from the slider ring and oxidation of aluminum was resulted from the heat generated during the abrasive wear. On closer examination of this sample, Figure 2 shows that subsurface shear cracks were developed. Such cracks were hundreds of micrometers in length and some propagated along reinforcement-matrix interfaces. One of the unique features of FIB is that the secondary electron images review grain orientation contrast. In this case, the aluminum grain size and the extent of deformation are shown in Figure 3. Comparing the shape of grains near the wear surface (Figure 3) and those away from the surface (Figure 4), a page_43 Page 44 distinct difference is noted. This reflects that substantial plastic deformation was induced to the grains near the surface, but not those away from the surface. Another phenomenon noted was that aluminum grains were heavily deformed in Al2O3 particle free regions, but no deformation was induced on grains surrounded by Al2O3 particles (Figure 3). From the observations made in the worn surfaces and their corresponding weight/volume losses, it is clear that the test conditions used were quite severe. Even at slow sliding speed (0.25 m/s), the frictional force was large enough to abrade the surface of composites to cause large wear debris and islands of mechanical mixed layers. Such severe wear condition had resulted in a high weight loss and substantial plastic deformation. In fact, the plastic deformation was so intense that cracking was observed below the mechanical mixed layers, which would cause sub-surface

delamination. As a result, fresh metal-to-metal contact and transfer would be significantly enhanced, leading to high weight loss. Under this condition, adhesive transfer assisted by subsurface cracking is expected to be the rate-controlling process for wear. The conditions for subsurface shear localization and cracking have been considered by Rosenfield [6], who proposed that under certain sliding conditions plastic deformation will be localized within a region below the contact surface. The localized plastic flow leads to shear instability and fracture of this region. The condition for shear instability is favored by a high coefficient of friction, heavy loading and a steep gradient of shear strength along the depth of deformation. In view of the above results, it is understandable that a high weight loss was experienced by the composite when wear tested at low sliding speed. Firstly, the initial abrasive wear on composite would lead to weight loss as extensive grinding of composite surface was happening. Secondly, the Al-1% Mg matrix is relatively soft, thus weight loss occurred through adhesive wear. Finally, shear cracking on the mixed layers caused layers flake-off, which led to weight loss. As the sliding speed is increased (1.5m/s), the frictional heat has developed to the point that aluminum matrix alloy was softened, thus weakened the frictional force [5], and less abrasion was taken place between the steel and the composite materials. In addition, it is more likely for the wear debris to be embedded into the matrix. As a result, a lower weight loss was noted. For samples tested at higher sliding speed (2.5 m/s), different observations were made on the cross section of the wear surface. Figure 5 shows a FIB image of an Al2O3/Al composite which was wear tested at a sliding speed of 2.5m/s. Similar to the previous sample, a mechanically mixed layer was also observed in this sample. In this case, the mechanical mixed region is extended to a greater depth, and a continuous, compacted (approximately 1.5 mm) and crack free sub-layer was formed on the top surface. page_44 Page 45 At a sliding speed of 2.5m/s, a much higher temperature would have reached at the wear surfaces. The wear resistance becomes worse, and weight loss starts to increase due to melt wear, and seizure of the matrix alloy. In addition, rapid oxidation of is expected. An interesting point to note is that why subsurface shear cracking was only observed in the mixed layer when slow sliding speed (0.25 m/s) was used and not at high sliding speed (2.5m/s). The likely explanation is that at slow sliding speed, only few islands of subsurface layers were developed. Being a mixed layer of iron and aluminum oxide, this layer has higher hardness than the soft aluminum matrix. During wear testing, an applied load of 2.1 Kgf was applied to the surface, such load was automatically transferred to the few islands of harder mixed layer. Consequently, high force per unit area was developed on the mixed layers, and causing unavoidable cracks. At fast sliding speed, substantial abrasion and mechanical mixing was happening at the early stage of wear, and a large continuous mixed layer covering the wear surface was quickly developed due to the high temperature. Therefore, the applied load was distributed on a larger surface of mixed layer, which resulted in a much lower force per unit area on the mixed layer. As a result, no crack was found on the mixed layer. The transition wear phenomenon observed here for the different sliding speeds is believed to be primarily due to the transition of one type of mechanical wear to another. Friction-induced thermal softening is likely to play the role in wear transition at higher velocities. Conclusions 1. The weight loss of Al2O3/Al composite is much lesser than cast iron (A48 class 30) when wear tested at high sliding speeds. 2. At low sliding speed (0.25 m/s), the mechanisms of wear of Al2O3/Al composite are abrasive and adhesive wear, along with surface cracking of mixed layers. 3. At high sliding speed (2.5m/s), the mechanisms of wear of Al2O3/Al composite are abrasive and melt wear. Acknowledgment The authors would like to thank Mr. R.Santos in the assistance of sample preparation and Dr. Alpas for the wear testing work. References 1. I.M. Hutching, Materials Science Technol., 10 (1994), p.513.

2. F.M. Hosking, F. Folgar-Portillo, R. Wunderlin and R. Mehrabian, J. Mater. Sci., 17 (1982), p.477. 3. S.V. Prasad and P.K. Rohagi, J. Met., 39 (1987), p.22. 4. J. Zhang and A..T. Alpas, Mater. Sci Eng., A161, (1993), p.273 5. S.C. Lim and M.F. Ashby, Acta Metall., 35, (1987), p.1. 6. A.R. Rosenfield, Wear, 116 (1987), p.319 page_45 Page 46 Table 1. Weight and volume loss of materials at different sliding speeds Weight Loss Volume Loss Sliding Speed Al2O3/Al Cast Iron Al2O3/Al Cast Iron 0.25 m/s 4.8 3.2 1.56 0.44 1.50 m/s 0.9 7.5 0.29 1.04 2.50 m/s 3.6 16.9 1.17 2.34 Table 2. Chemical analysis by electron microprobe of a sample which was tested at 0.25 m/s sliding speed Distance from wear surface 1.25 mm 2.50 mm 6.25 mm 10.00 mm Weight % Oxygen 28.817 31.587 30.733 1.961 Iron 10.273 9.972 0.129 0.029 Aluminum 60.910 58.441 69.139 98.010 Total 100.000 100.000 100.000 100.000

Figure 1. FIB image of an Al2O3/Al composite wear tested at a sliding speed of 0.25 m/s. page_46 Page 47

Figure 2. FIB image of an Al2O3/Al composite showing the presence of shear cracks in the mixed layer surface.

Figure 3. FIB image of an Al2O3/Al composite showing grains near the wear surface were heavily deformed. page_47 Page 48

Figure 4. FIB image of an Al2O3/Al composite showing no deformation was induced on grains away from the wear surface.

Figure 5. FIB image of an Al2O3/Al composite wear tested at a sliding speed of 2.5 m/s. page_48 Page 49

Design and Fabrication of Smart Composites for Static Shape Control H. Wang 1, M. Giray 2, C. K. Jen 1, S. Kalaycioglu 2 and S. E. Prasad 3 1 Industrial Materials Institute, National Research Council Canada 75 de Mortagne Blvd., Boucherville, Québec, Canada J4B 6Y4 2 Canadian Space Agency 6767 Route de l'Aéroport, Saint-Hubert, Québec, Canada J3Y 8Y9 3 Sensor Technology Ltd. 20 Stewart Road, Collingwood, Ontario, Canada L9Y 3Z4 Keywords: smart structures, composites, sensor, actuator, shape control Abstract A graphite/epoxy test article with embedded optic fiber strain sensors and surface-attached piezoceramic actuators has been developed for the feasibility study of static shape control of composite structures under thermal loading conditions. The study identifies several critical issues in the design and analysis techniques, the sensor and actuator technologies, and the integration of such devices with composite materials. The promise and limitation of practical application of the shape control concept in space structures are also discussed. Introduction Advanced composites such as graphite/epoxy are increasingly used in space structures, because the materials have a high specific stiffness and low coefficient of thermal expansion [1]. Space structures such as antenna reflectors require a high dimensional stability when exposed to the severe thermal environment in space. The temperature fluctuation of typically -100 to 60°C in low earth orbit and -160 to 120°C in geosynchronous orbit [1] may cause excessive thermal deformation and impair the operation of the structures. Using composites can significantly reduce the deformation by designing the laminate with a minimum coefficient of thermal expansion; but small and random temperature-induced shape change may still exist, resulting from uncontrollable factors in material fabrication and in service operations. Therefore, the active structures technology is currently considered to be applied to sense and control the thermal deformation of space structures [1].

Many aspects of the technology need to be evaluated and demonstrated for the potential application. The purpose of this work is to design and fabricate a graphite/epoxy test article with integrated sensors and actuators, which can be used for the feasibility study on shape control of composite structures under thermal loading. It also intends to identify the critical issues in the design and fabrication of such structures, and thus assessing the promise and limitation of the up-to-date technologies of sensor, actuator and composite for practical application of the shape control concept in space structures. page_49 Page 50 System Definition The objective is to develop a 25.4 by 305mm composite beam with embedded fiber optic sensors and surfaceattached piezoelectric actuators, which possesses the following functions. At the temperature loading of DT = 50°C the system can demonstrate a) about 10 mm end deflection under clamp-free configuration; b) a strain response to be detected by the sensors; c) the ability to correct the shape change by the actuators within 400 volt (DC) power supply. Only one surface attachment of the actuators was considered for the system. The composite was AS4/3501-6 graphite/epoxy of Hercules Inc., a material widely used in aerospace industry and extensively studied for its thermomechanical properties [2]. The optic fiber sensor was the extrinsic Fabry-Perot interferometer of FISO Technologies Inc. (Québec, Canada). The sensor provides strain measurement in the range of ± 2000 me with a 0.2 me resolution, while being insensitive to transverse strain and environment temperature fluctuation. The sensor dimension was 250 mm in diameter and about 10 mm in gage length. The actuator selected for design was a modified PZT (lead-zirconate-titanate) piezoceramic, BM532, of Sensor Technology Ltd. (Collingwood, Canada). It was in the form of 0.635 mm-thick rectangular patch with through-the-thickness polarization. When the PZT patch is bonded to the composite and a voltage applied to its electrodes at the opposite surfaces, it can deform by in-plane expansion or contraction, providing actuation to the system. The conception of the system and its functions are schematically shown in Fig. 1. The laminate is a non symmetric type and the sensors are embedded during laying up. After curing at Tcure = 177°C (Fig. 1a) and cooling down to room temperature (RT), the laminate warps into an initial shape (Fig. 1b) due to residual stresses related to the mismatch in thermal expansion between the composite layers with different orientation of the reinforcing graphite fibers. The actuators are subsequently bonded to the laminate surface (Fig. 1b). During shape control testing, the composite is heated above RT and its shape should change with a decreasing curvature (and deflection) because of partial release of the residual stresses. The strain sensor should respond to the shape change (Fig. 1c); and the actuators can be activated accordingly to correct the shape change (Fig. 1d). In the case shown, the voltage applied to the actuators is such that a contracting deformation is achieved thus the reduced beam curvature (or deflection) can be restored to its initial level at room temperature. Similarly the shape change due to a temperature drop below RT can also be studied. This design conception is advantageous for the feasibility study of detecting thermal deformation and shape control at normal laboratory conditions. While it is noted that, contrary to space applications, the thermal deformation is enlarged for the research and demonstrative purposes. The non-symmetric cross-ply laminate family was selected for designing the system. This is because the laminate provides a simple cured shape and, for which, the mechanism of thermal deformation has been well studied (e.g. [2]). Design Analysis In spite of the simple geometry and mild temperature loading conditions, many aspects of the system need to be analyzed in the design stage. Given the limited piezoelectric capacity of the actuators (notably d31 = -274 E-12 m/v for PZT-5H), the essential design loop was to tailor the laminate lay-up so that the specified thermal deformation could be corrected by the actuators at a reasonable PZT coverage over the beam surface and within 400 volt power supply. The lay-up strategy should also enhance the performance of embedded sensors thus the thermal deformation could be reliably detected. It has been suggested in [3] that the sensor should be placed close to the laminate surface (away from the neutral axis) and between two plies of composite prepreg with page_50 Page 51

Fig. 1 Conception of a smart composite beam for the study of static shape control under thermal loading. identical orientation. In addition, the system needs to be verified to be resistant to mechanical failure during fabrication and when the beam experiences thermal deformation. The design analysis was conducted using the commercial finite element code ANSYS. A fully-coupled thermomechanical and piezoelectrical model was established, which consisted of the layered solid elements (SOLID-46) for the composite laminate and the coupled-field solid elements (SOLID-5) for the PZT patches [4]. A user-module was developed to couple the displacement degree-of-freedom at the interface between the two element types and also to enable the parametric design capacity, which allowed the variation of the laminate dimension and lay-up and the number, size and location of the actuators. A model of the cantilever beam with two attached PZT patches is shown in Fig. 2. The PZT patch has the same width as the composite beam. The material properties of AS4/3501-6 composite and the PZT-5H piezoceramics, available in [2] and [5] respectively, have been used in the analysis. page_51 Page 52

Fig. 2 Deflection of the composite beam by thermal loading and by PZT actuation: a) at DT = -154°C (the initial shape after cure), b) at DV = 400 volt actuation. In the first stage, analyses were made to trade off between the laminate lay-up and the PZT-coverage, so that the end deflection of the beam at DT = 50°C could be corrected by applying 400v to the actuators. This was done by considering the following two loading conditions.

where T is the applied temperature, in relative to the reference temperature of 177°C for the composite and 23°C for the PZT, respectively. Vin is the voltage applied to the PZT surface bonded to the composite and Vout to the outer surface. Two predictions were made from this loading condition. That is, when T = 23°C and then 73°C, the model predicted respectively the initial beam shape after curing and bonding of PZT and its shape at 50°C above room temperature. The corresponding shape change was the difference between the two predictions. It is noted that the page_52 Page 53 thermal expansion of PZT patches has been considered in the model which, as discussed in [1], affects the actuator performance under temperature loading.

This condition predicts the corrective shape change by the PZT patches at 400v power supply. It is noted that a flat beam was used as the initial configuration in Case 2; while the corrective shape change should be based on the deformed beam shape as predicted in loading Case 1 at T = 23°C. However, since the model is in the linear range of piezoelectricity and thermoelasticity, the principle of superposition applies. The effect of large deflection was found to be negligible during the analysis. Therefore the flat configuration can be used in both cases and the resulting shape due to the combined loading conditions can be obtained by adding one deflection to the other. Figs. 2 and 3 show the typical results of the design analysis. For each candidate lay-up, the shape at room temperature was predicted from Loading Case 1 at T = 23°C, as shown in Fig. 2a. The shape at T = 73°C was also predicted so the shape change from 23 to 73°C was the difference between the two predictions. The shape correction provided by PZT was determined from Loading Case 2, Fig. 2b. The above calculations were repeated for different PZT coverage ratios: 16%, 33%, 49%, and 65%. The coverage ratio is defined as the PZT-covered area over the beam surface area. Then the beam deflections due to shape change and shape correction are plotted versus the coverage ratio, as shown in Fig. 3 for the typical variations. In Fig. 3a the corrective deflection increases linearly

with the coverage ratio; while the deflection due to shape change remains basically constant, indicating a negligible effect of the PZT patches on the flexural stiffness of the beam. The crosspoint of the two curves corresponds to the minimal PZT coverage required to correct the shape change, which is about 41% for the lay-up shown in Fig. 3a.

Fig. 3 End deflection of composite beams vs. PZT coverage ratio: (a) Type C1A lay-up; (b) Type C3 lay-up. In Fig. 3b the two curves exhibit similar variations as in Fig. 3a; but the lay-up is so stiff that the deflection cannot be corrected even with a PZT coverage of 65%. It is noted that a higher PZT page_53 Page 54 coverage ratio was not considered because distributed PZT patches over the beam length was preferred for shape control by distributed voltage application. A total of seven laminates has been evaluated, as summarized in Table 1. It is seen that the beam deflection can be changed for about 6 to 10 mm by DT = 50°C; but as the beam generally becomes stiffer as its thickness increases, it requires higher PZT coverage for correction. The shape change may eventually become incorrectable for a thicker composite with the given power supply. Table 1 Summary of design analysis results for crossply AS4/3501-6 composite beams. Type Plies/Thickness Lay-up End Deflection by 50°C, Min. PZT (mm) (mm) Coverage for Correction (%) C 10/1.3 [02/902/0/902/0/902] 11 40 C1 12/1.56 [02/90/0/902/0/902 8.5 41 /0/902] C1A 13/1.69 [02/90/0/902/0/903 9.03 44 /0/902] C2 13/1.69 [03/90/0/902/0/902 8.3 45 /0/902] C3 19/2.47 [04/90/0/90/0/903/0 7.5 69 /903/0/903] C4 17/2.21 [03/90/0/90/0/903/0 6.8 58 /903/0/902] C5 16/2.08 [03/90/02/903/0/903 7.6 57 /0/902] As will be discussed later, some lay-ups in Table 1 were fabricated during the molding experiment. Residual stressinduced matrix cracking was observed in the laminates with a thick layer, e.g. [0°]4. The matrix crack disturbs the local strain field thus may affect the response of a optic fiber sensor embedded nearby. Considering the results of both analysis and fabrication, the 13-ply Type-C1A laminate: [0/S/0/90/0/902/0/903/0/902] was selected for building the test article. In the clamped-free configuration, the laminate provides 9.03 mm of end deflection at DT=50°C, which requires more than 44% PZT coverage for correction. The interface between the two 0° surface plies, denoted by ''S", is where the sensor was to be embedded. It has been indicated that, for a better response, the sensor should be embedded close to the laminate surface and in alignment with the two composite plies

between which the sensor is embedded [3]. In the second stage, further analyses were made for the C1A-type beam in order to verify the other design criteria. It was found that, at the temperature loading of DT = 50°C, the axial residual strain of the beam at mid thickness of the [0°]2 surface layer fluctuates for 47 me. This strain response was later proven to be detectable by the sensors embedded at the interface. In addition, at the uniform thermal loading, the strain response is essentially constant over the beam. Thus the sensor performance does not depend on the in-plane location of embeddment. During embedding, the sensors were rather distributed along the beam length to monitor the shape change at different sections of the beam. It is noted that mechanical loading or non-uniform temperature distribution may lead to varying strain filed in the beam. The structural integrity aspect of the system was also examined. Since the composite beam was subjected to a flexural deformation, it was concerned that matrix failure might be induced in the [90°]2 surface layer during thermal loading and actuation. If the actuation was to maintain the room-temperature shape of the beam, the post-cure residual strain page_54 Page 55 in the layer along its transverse direction was the governing parameter. The strain was found to be 22 me, which is a low strain level comparing to the transverse failure strain of the material (about 6000 me); thus matrix cracking should not occur in the [90°]2 surface layer during fabrication and the shape control test. It was later evident that no matrix cracking was found by examining the molded panel. On the other hand, the stress in the PZT patches was also of concern because the material is quite fragile, with a tensile strength of about 63 MPa. Again the critical deformation for the PZT patches corresponds to the initial shape of the beam after cure. It was found that, for the simulated two-patch case, the maximum von Mises stress in the PZT patches, is about 26 MPa. So the patches should not fail during the specified functions. Fabrication and Verification Before embedding the sensors, some molding experiments were conducted to examine the embedding procedure, using bare optic fibers. Five 305 by 305 mm (12" by 12") square panels with typical lay-ups selected from Table 1 were molded using a Baron autoclave. The manufacturer-recommended cure cycle [6] was basically followed, except that the cooling rate of 3°C/minute was lower than the recommended 5°C/minute. It was found that the laminates containing a thick layer (e.g. [0°]3 or [90°]3 for surface layer, or a thicker one for internal layer) exhibited extensive matrix cracking after cure. Such cracks penetrated the layer thickness, ran across the plate dimension, and distributed nearly uniformly over the whole panel. The cracks were caused by excessive thermal residual stress and the low transverse strength of a thick composite ply [7]. They were more likely to occur in a thick surface layer because its neighboring ply provides less constraint (from one side only) to the matrix fracture [7]. It is known that a matrix crack can induce a local disturbance to the strain field. Thus if an embedded sensor happens to locate near or across such a crack, its response will be affected and may provide erroneous readings. In this regard, many lay-ups in Table 1 were prone to matrix cracking and were therefore disqualified for fabrication. It is worthy noting that in [3] it has been suggested that the thickness of the sensor-embedded surface layer should be thicker than twice the sensor diameter in order to guarantee the embedding quality. Given the sensor diameter of 250 mm, more than four plies of AS4/3501-6 prepreg (130 mm in thickness) are needed to contain the sensor; but from the observations discussed above, such a layer can be too thick and prone to matrix cracking. It is also noted that, as the optic fibers were embedded close to the laminate surface, the surface finish and cleanness of the aluminum mold plates required special attention. A small indent or scratch on the mold surface, a dirt or residue from the previous molding cycle, or even a debris of chopped prepreg, could mold into a surface, a dirt or residue from the previous molding cycle, or even a debris of chopped prepreg, could mold into a surface defect of the laminate which can affect the performance of an embedded sensor or even its survival of the molding process.C Based on the molding practice and the design analysis, a type-C1A laminate was molded for producing the composite beam with embedded optic fiber sensors. Two sensors were located colinearly between the 0° surface plies at respectively 115 mm and 216 mm distance from the plate edge. Near the exit of the optic fibers from the composite, for protection purposes, the fibers were jacked by a 800 mm-diameter Teflon tube for a length of 25 mm inside plus 25 mm outside the composite. Careful procedures were followed in handling the optic fibers, in sensor embedding, air-bagging, autoclaving, and the demolding and trimming steps. Post-molding inspection has found that the sensors were close to the surface but well embedded in the material. Though PZT-5H piezoceramic was considered in the design, some 0.19 mm-thick PZT-5A patches were actually used. Four such actuators were mounted on the composite using a commercial glue, covering 62% of the beam surface. The piezoelectric coefficient of PZT-5A (d31 = -171 E-12 m/v) is lower than that of PZT-5H, thus providing less piezoelectric actuation at the given power supply. page_55

Page 56 The system has been subjected to some preliminary tests in order to verify the specified functions. Fig. 4 shows the end deflection initial shape of the beam at room temperature, where the measured data compare reasonably well to the prediction by finite element analysis. The slight discrepancy could be due to the scatter of material properties and the known effect of epoxy cure which was neglected in the analysis. The beam was also tested in a thermal chamber. The end deflection and the strain data were monitored when the beam was heated and cooled over the temperature range from RT to 100°C. Fig. 5 shows the strain variation versus the temperature. The close agreement between the measured and predicted data indicate a successful embeddment of the sensor. Similar performance was also verified for the strain response induced by the actuators.

Fig. 4 Post-cure shape of C1A beam

Fig. 5 Strain response vs. temperature Conclusion The optic fiber sensors can survive the standard cure process of AS4/3501-6 composite if special attention is paid to the handling and protection of the fragile device. Though the embedded sensors can detect the enlarged thermal deformation, further tests are needed to identify their effective resolution and the performance under various loading conditions. The limited actuation capability of the current piezoceramics is one of the major difficulties in using the materials for static shape control of space structures. Finite element method is effective for the design and analysis of the smart composite structures. The fully-coupled thermomechanical piezoelectrical model can simulate the highly interactive response of the structures during shape control under thermal loading. Reference 1. H. L. McManus, in Smart Structures and Intelligent Systems, SPIE v.1917 (1993), p.545. 2. Crasto & Kim, J. of Reinforced Plastics and Composites, v.12 (1993), p.545. 3. L. G. Leka and E. Bayo, J. Composites Technology & Research, v.11 (1989), p.106. 4. ANSYS User's Manual (version 5.4) 5. Material properties provided by Sensor Technology Ltd. 6. Product Data Sheet, No. 843-3, Hercules Inc. 7. Wang, H. and T. Vu-Khanh, J. Composite Materials, v.28 (1994), p.684.

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Fracture Behavior of Adhesively Bonded Composite-to-Metal Lap Joints with Thick Adherends James F. P. Owens1, Pearl Lee-Sullivan2 1Boeing Canada Technology Inc., 99 Murray Park Road Winnipeg, Manitoba, Canada R3J 3M6 2Department of Mechanical Engineering, University of New Brunswick P.O. Box 4400, Fredericton New Brunswick, Canada E3B 5A3 Keywords- adhesive bond, single lap joint, stiffness reduction, composite-to-metal joints, fracture. Abstract This paper presents part of the results of a wider research program on the study of bond fracture and resulting stiffness loss in adhesive bonded composite-to-aluminum single lap joints. Test specimens consisting of long lap joints with thick adherends were subjected to incremental loading until complete failure. The effects of epoxy adhesive rigidity and temperature on joint deformation behavior were investigated. In the temperature studies, tests were conducted at room temperature and at -40°C. In each test, the stiffness reduction due to fracture of the adhesive bond was calculated from load-displacement curves. The results showed that stiffness reduction was higher where delamination joint failures occurred, and that the composite adherend was more prone to delamination when bonded with a flexible epoxy and loaded at -40°C. On the other hand, the joints bonded with a more rigid epoxy failed by adhesive fracture. Introduction There is a growing need to combine metals with polymeric composites in order to optimize the strength, weight and durability of components in aircraft and spacecraft applications. Composites are considered to be more structurally efficient than metals and do not experience galvanic corrosion. Metals on the other hand have better damage tolerance and failure predictability than composites, and are unaffected by the solvents and temperatures which tend to degrade polymers. In order to benefit from the properties offered by both types of materials, hybrid compositeto-metal structures have been developed. An example of these structures is an aircraft engine strut containing lightweight high strength carbon fibre/epoxy fairings joined to damage tolerant aluminum ribs. Although these structures provide an excellent blend of material properties, their success depends upon the integrity of the joints which connect them together. The most efficient method of connecting hybrid structures is either by adhesive bonding or mechanical fastening [1] of shear joints. Mechanical fastening can only achieve a maximum tensile strength of 50% of the weakest adherend in the joint due to the stress concentrations caused by the fastener holes. In comparison, adhesively bonded joints can achieve in excess of 80% of the tensile strength of the weakest adherend even with a simple single shear configuration [2]. Adhesive bonding is therefore particularly efficient for joining composite-to-metal structures which do not require subsequent disassembly for maintenance and inspection. However, the fracture behavior of adhesive bonded joints using dissimilar adherends is not well understood, particularly where the fracture is caused by thermal mismatch in service. Consequently the use of bonded joints has largely been limited to secondary structures such as aerodynamic fairings and wing panels in aircraft. page_57 Page 58 This paper presents part of the results of a research program to study the fracture behavior in shear lap joints with thick adherends when subjected to tensile loading. The joints were bonded with two adhesives with different elastic moduli. In order to investigate temperature effects, tensile testing was performed at room temperature and at -40°C. The onset of failure and crack growth mechanisms were observed. Fracture behavior was characterized by measuring the effects of adhesive crack growth, a, on the overall joint stiffness, K, [e.g. 3,4,5]. Although an analytical model was also developed to predict the stiffness reduction due to crack growth, only the experimental results are presented here. Experimental Details Joint Design and Preparation

There are presently various ASTM Standards [6,7] for testing shear joints. In this work, the composite-to-aluminum joints were designed with a long lap and thick adherends to better represent typical engineering applications (Figure 1). According to Mathews [1], the ratio of the lap length (L) to adherend thickness (t) is one of the most critical factors affecting joint structural efficiency. The joints tested were therefore designed with an L/t ratio of 50/1 as a guideline for obtaining a joint strength which exceeds 70% of the strength of the weakest adherend in the joint [2]. An adherend thickness of 6.35 mm (1/4") was used to simulate load bearing structures in aerospace applications [1]. An adhesive bond line thickness of 0.5 mm (0.010") was used since the thickness is within the optimum range for common structural adhesives [1]. Shimming tabs were bonded at the ends of the adherends to prevent the testing machine grips from twisting the joints during clamping.

Figure 1: Dimensions of single lap joint specimens. The joints were fabricated with 6061-T65 aluminum alloy (E= 72 GPa, sUlt=260 MPa) and Extren 500 pultruded polyester/E-glass composite (EL = 12.4 GPa, sUlt = 137 MPa). Extren consists of a layer of random chopped fibre glass sandwiched between two layers of unidirectional continuous E-glass fibres. The aluminum adherends were machined from 50.8 mm ´ 6.35 mm (2" ´ 1/4") flat bar whereas the composite adherends were machined from 1.2m × 2.44 m × 6.4 mm (4' × 8' × 1/4") sheet. The joints were bonded with either flexible EP21LV (E= 2.6 GPa) or rigid EP45HT (E = 4.0 GPa) Masterbond two-part structural epoxy adhesives. The bond line thickness was controlled by mixing glass beads conforming to ASTM D-1214 (f = 0.2490.297 mm) with the adhesives in amount equal to 0.5% of the total weight. The adherend bonding surfaces were prepared by grit blasting and degreasing with acetone and by applying a solution of K2Cr2O7 + H2SO4 to the aluminum adherends. The joints were laid into a purpose-made jig and cured in accordance with manufacturer's instructions using a hydraulic hot press. Before testing, the joint corners were machined to the radius r. page_58 Page 59 Test Procedure and Data Analysis The test procedure consisted of tension loading the shear lap joints in incremental displacements of 0.1 mm at a rate of 1.3mm/min, Figure 2, using a electronic controlled 1332 servohydraulic Instron machine. In order to simulate actual crack initiation and propagation, this procedure was used instead of employing artificial crack starters at the bond layer. Crack length was measured visually during each incremental cycle after the maximum load had been reached and the crack propagation had arrested. The total joint displacement, d, was measured during each load increment by two transverse mounted pencil LVDT's at the joint ends, Figure 2. Cold temperature testing at -40°C was controlled using an environmental chamber. The load and displacement data were recorded with a data acquisition system interfaced to a personal computer.

Figure 2: Test setup and loading procedure. Four sets of tests were performed for each condition. In each test, the joint stiffness K for each incremental loading cycle was calculated by taking the slope of the load-displacement curves dP/dd. As shown in Figure 3, a 310 kN range was used in all calculations in order to maintain consistency. Each cycle of the incremental load was plotted and the slope was calculated using the sum of the least squares linear regression. As seen in Figure 3, the load-displacement curves were linear except at the upper end of the load cycles where sudden slope changes occurred due to crack growth. As expected, crack growth at the bond due to incremental loading resulted in an overall reduction in stiffness [3,4,5]. page_59 Page 60

Figure 3: Typical load-displacement curves upon incremental loading. Results and Discussion Stiffness Reduction Due to Bond Failure Figures 4 and 5 show the stiffness reduction in joints bonded with flexible epoxy (F) and tested at room temperature (CAFRT) and -40°C (CAFCT), respectively. Figure 6 and 7 show similar results for joints bonded with more rigid epoxy (R) and tested at room temperature (CARRT) and at -40°C (CARCT), respectively. Generally, there is greater scatter in the data sets for the tests performed at room temperature. It also appears that the use of a more rigid adhesive would slightly increase the joint stiffness. In all four conditions the stiffness, K, generally decreased linearly with increasing crack length, a, except at the initial portion of the curves. The initial increase in K was due to aluminum strain hardening which was confirmed by performing individual tests on the adherends and adhesives. Strain hardening is not normally accounted for within analytical models found in the literature, as reviewed in [8].

Nevertheless, its effects dissipated before significant growth occurred (a > 20 mm). According to our analytical model developed for predicting stiffness [8], the relatively small difference in elastic modulus for the two adhesives would not significantly influence the rate of stiffness reduction. Details of the analytical work will be presented elsewhere. In the experimental portion of the work, the rate of stiffness reduction was estimated by the slope of dK/da for the crack length range between 20 250 mm. It was found that there is greater stiffness reduction for the flexible epoxy and testing at cold temperature. At room temperature, the rate was -0.0272 kN/mm for the flexible CAFRT joint as compared to -0.0194 kN/mm for the rigid CARRT equivalent. At -40°C, it was -.0284kN/mm for the CAFCT as compared with -0.0227 kN/mm for the CARCT. The results suggest that the more rigid adhesive had better resistance to crack propagation. The differences in behaviour, although minor, could be attributed to the differences in fracture mechanisms observed in all the specimens tested. Fracture Mechanisms There were two main fracture mechanisms observed during debonding of the lap joints as shown schematically in Figures 8 and 9. The first was adhesive failure at the composite/aluminum interface typically seen in rigid joints irrespective of the temperatures tested, Figure 8. The other failure mechanism was delamination of the unidirectional outer layer of the pultruded bar, Figure 9, which was predominant at -40°C. Delamination with fibres bridging both the adherends was consistently observed in the flexible joints. In all specimens, cracks tend to initiate at the composite fillet end and extend inwards until total debonding. page_60 Page 61

Figure 4: Stiffness as a function of crack length for flexible CAFRT joint at room temperature.

Figure 5: Stiffness as a function of crack length for flexible CAFCT joint at -40°C. page_61

Page 62

Figure 6: Stiffness as a function of crack length for rigid CARRT joint at room temperature.

Figure 7: Stiffness as a function of crack length for rigid CARCT joint at -40°C. page_62 Page 63

Figure 8. Typical adhesive fracture observed in rigid-epoxy joints at room temperature.

Figure 9. Typical delamination and fibre bridging observed in -40°C tests for both epoxies. page_63 Page 64 In comparing the stiffness reduction results, the specimens tested at -40°C showed a higher rate of stiffness loss than those tested at room temperature. This was found to be due to delamination failure of the pultruded composite near the bondline for both epoxies used, Figure 9. It appears then that delamination together with fibre bridging led to a significant reduction in the load-carrying capability of the joint and was triggered by extreme cold. The tendency for delamination would be attributed to both the polyester matrix within the composite and the epoxy bond layer becoming more rigid in the cold, resulting in brittle fracture. Conclusions 1. The fracture behaviour of the composite-to-metal lap joints could be characterized by the rate of stiffness reduction per unit crack growth, which appears to be constant once metal strain hardening is limited. 2. Lap joints bonded with the flexible adhesive were more prone to delamination failure of the pultruded composite. Delamination failure together with fibre bridging between the two adherends led to higher stiffness reduction. 3. Delamination failure near the adhesive bondline was found to occur more frequently in tests at -40°C irrespective of the adhesive used. References [1] Mathews, F.L., ''Bonded and Mechanically Fastened Joints", in "Handbook of Polymer Composites for Engineers", Woodhead Publishing limited, Cambridge, England, 1994. [2] Hart-Smith L.J, "Design of Adhesively Bonded Joints", in "Joining Fibre Reinforced Plastics", Ed., Mathews, F.L., Elsevier Applied Science, London, New York, 1987 [3] O'Brien, T. K., "Stiffness Change as a Non-Destructive Damage Measurement", Structures Laboratory Report, U.S. Army Research and Technology Laboratories (AVRADCOM), NASA Langley Research Center Hampton, VA, 23665, 1979. [4] O'Brien T.K., Reifsnider, K. L.,"Fatigue Damage Evaluation Through Stiffness Measurements in Boron-Epoxy Laminates", Journal of Composites Materials, Vol. 15, January, 1981, pp. 55. [5] Camponeschi E.T., and Stinchcomb W.W., "Stiffness Reduction as an Indicator of Damage in Graphite/Epoxy Laminates Composite Materials:Testing and Design (Sixth Conference), ASTM STP 787, I.M. Daniel, Ed., American Society for Testing Materials, 1982, pp. 226246. [6] "ASTM D5656-95 Standard Test Method for Thick-Adherend Lap-Shear Joints for Determination of the StressStrain Behaviour of Adhesives in Shear by Tension Loading", American Society for Testing Materials, Philadelphia, PA, USA, March, 1995.

[7] "ASTM D3983-81 (Re-approved 1986) Standard Test Method for Measuring Strength and Shear Modulus of Non-Rigid Adhesives by the Thick Adherend Tesile Lap Specimen", American Society of Testing Materials, Philadelphia, PA, USA, 1986. [8] Owens, J.F.P., "Studies on Stiffness and Fracture Behaviour in Adhesively Bonded Composite-to-Metal Shear Joints", MScE thesis, The University of New Brunswick, 1998. Acknowledgements This authors which to thank the Natural Sciences and Engineering Research Council of Canada (NSERC) for funding this work. Great appreciation is also expressed to Strongwell Ltd (Formerly MMFG) of Bristoll, Virginia, U.S.A. for the donation of Extren Series 500 composite. page_64 Page 65

FATIGUE AND DYNAMIC FAILURE page_65 Page 67

Prediction of Tensile Fatigue Life for GFRP/Metal Adhesive Joints Masayuki Nakada1, Sangwook Sihn2, Tomoya Imai1, Yasushi Miyano1, and Stephen W. Tsai2 1Materials System Research Laboratory, Kanazawa Institute of Technology, Yatsukaho, Matto, Ishikawa 924-0838, Japan 2Department of Aeronautics & Astronautics, Stanford University, Stanford, California 94305-4035, U.S.A. Key Words: GFRP, Joints, Fatigue strength, Life prediction, Time-dependent properties Abstract We had proposed a prediction method of fatigue failure load of FRP adhesive joints under an arbitrary frequency, load ratio (minimum load/maximum load), and temperature. This method is based upon the four hypotheses, (A) same failure mechanism for constant elongation-rate (CER), creep, and fatigue failure, (B) same time-temperature superposition principle for all failure loads, (C) the linear cumulative damage law for monotone loading, and (D) linear dependence of fatigue failure load upon load ratio. Tensile tests of GFRP/metal adhesive joints for CER, creep, and fatigue loadings were conducted for various temperatures. As a result, the characteristic time-temperature dependent fatigue behavior of this FRP joints is clarified by using this prediction method. Introduction It is well known that the mechanical behavior of polymer resins exhibits time and temperature dependence, called viscoelastic behavior, not only above the glass transition temperature Tg but also below Tg. Thus, it can be presumed that the mechanical behavior of polymer composites also significantly depends on time and temperature. It has been confirmed that the viscoelastic behavior of polymer resins as matrices is a major influence on the time and temperature dependence of the mechanical behavior of FRP [17]. In previous papers, we proposed a prediction method for the fatigue strength of polymer composites for an arbitrary frequency, stress ratio, and temperature from the data measured by constant strain-rate tests at several strain-rates and various temperatures, and fatigue tests at a single frequency and various temperatures. The validity of this method was proven for the tensile behavior of conical shaped GFRP joint and GFRP/metal adhesive joint as well as several kinds of CFRP [811]. In this paper, the validity of the prediction method is discussed for the case of the tensile behavior of GFRP/metal adhesive joint in which a ductile adhesive resin, PMMA, is used for adhesive resin. Prediction Procedure

A prediction method for fatigue failure load of composite structures for an arbitrary frequency, load ratio(minimum load/maximum load), and temperature rests on the four hypotheses, (A) same failure mechanism for constant elongation-rate (CER), creep, and fatigue failure, (B) same time-temperature superposition principle for all failure loads, (C) linear cumulative damage law for monotone loading, and (D) linear dependence of fatigue failure load upon load ratio. When these hypotheses are met, the fatigue failure load for an arbitrary combination of frequency, load ratio, and temperature can be determined based on the master curves of CER failure load and fatigue failure load for zero load ratio. The master curve of CER failure load can be constructed from page_67 Page 68 the test results at several elongation-rates for various temperatures. On the other hand, the master curve of fatigue failure load for zero load ratio can be constructed from the test results at a single frequency for various temperatures using the time-temperature superposition principle for the CER failure load. The outline of this method is shown schematically in Fig.1 together with definitions of some notations. The detail of the method will be presented with experimental results.

Fig. 1 Prediction procedure of fatigue failure load

Fig. 2 Configuration of FRP joint Table 1 Test conditions Loading Loading rate Frequency Load ratio Temperature type [mm/min] [Hz] [°C] Pmin/Pmax CER 100 25,40,50,60, 1 70,80,90 0.01

Fatigue I

-

5

0.05

Fatigue II Fatigue III

-

0.05 5

0.05 0.5 0.95

25,40,50,60, 70 10,50 25,40,70

page_68 Page 69 Experimental Procedure Preparation of GFRP/Metal Adhesive Joints The GFRP/metal adhesive joints (FRP joint) was made from a GFRP pipe, ductile cast iron rod, and adhesive resin as shown in Fig.2. The adhesive resin is PMMA resin, PLEXUS A0425 (ITW Adhesives). Ductile cast iron rod is made from ductile iron castings Grade 80-55-06 (ASTM A 536-84). The adhesive resin thickness and length of FRP joint are respectively 4mm and 28mm. Test Procedure The tensile tests for CER and fatigue loadings were conducted for various temperatures. The test conditions are shown in Table 1. The tensile CER tests were conducted at 5 testing temperatures between T=25 and 90°C by using an Instron type testing machine. The tensile load was applied at both end screws of the FRP joint. The loading-rates (cross-head speeds) were 0.01, 1 and 100mm/min. The tensile fatigue tests were conducted at 5 testing temperatures between T=25 and 70°C at a frequency f=5Hz, and 10, 50°C at f=0.05Hz, by using an electro-hydraulic servo testing machine. Load ratio R (minimum load/maximum load) was 0.05. Additionally, the fatigue tests were also conducted at T-25, 40, 70°C, f=5Hz and R=0.5, 0.95. Results and Discussion The CER and fatigue failure of FRP joint occurred in the adhesive resin nearby the interface between cast iron rod and adhesive resin. All failed specimens are similar regardless of loading pattern. We consider, therefore, that the failure mechanisms are the same for CER and fatigue loadings. Load-elongation Curves Typical load-elongation curves of FRP joint at various temperatures for CER test are shown in Fig.3. These curves show nonlinear behavior caused by the plastic deformation of adhesive resin. The yield and failure points are defined by the knee and maximum load points on the load-elongation curves. Master Curve of CER Failure Load The left side of Fig.4 shows the CER yield load Py and elongation ly versus time to yield ty at various temperatures of FRP joint, where the ty is defined as the time period from initial loading to Py in constant elongation-rate test. The master curves of Py and ly versus reduced yield time ty1 at a reference temperature T0=40°C as shown in the right side of Fig.4 were constructed by shifting Py and

Fig. 3 Load-elongation curves at various temperatures page_69 Page 70

ly at various temperatures along the log scale of ty untile they overlapped each other. The left side of Fig.5 shows the CER failure load Ps and elongation ls versus time to failure ts at various temperatures of FRP joint, where the ts is defined as the time period from initial loading to Ps. The master curves of Ps and ls versus reduced failure time ts1 at T0=40°C as shown in the right side of Fig.5 can be also constructed. Since Py, ly, Ps, and ls at various temperatures can be superimposed smoothly, the time-temperature superposition principle is applicable for Py, ly, Ps, and ls. Figure 6 shows the time-temperature shift factors aTo(T) for the master curves of Py and Ps of FRP joint. The aTo(T) are quantitatively in good agreement with Arrhenius' equation by using two different activation energies.

Fig. 4 Master curves of CER yield load Py and elongationly

Fig. 5 Master curves of CER failure load Ps and elongationls

Fig. 6 Time-temperature shift factors page_70 Page 71

Where, DH is activation energy [kJ/mol], R is gas constant 8.314´10-3 [kJ/(Kmol)]. The dotted lines in this figure show aTo(T) obtained experimentally for the creep compliance of the adhesive resin. The aTo(T) for Py of FRP joint and Dc of adhesive resin agree well with each other. However, the aTo(T) for Ps of FRP joint are different from that for Dc of adhesive resin. Master Curve of Creep Failure Load We had proposed a prediction method of creep failure load Pc from the master curve of CER failure load using the linear cumulative damage law. Let ts(P) and tc(P) be the CER and creep failure times for the load P. Suppose that the material experiences a monotone load history P(t) for 0£t£* where t* is the failure time under this load history. The linear cumulative damage law states

When P(t) is equal to constant load P0, the above formula implies t* = tc(P0). It is clear from the load-elongation curves shown in Fig.3 that the CER tests employed is approximately equal to constant load tests, that is, creep tests. Therefore, it is not necessary to apply the linear cumulative damage law to the results of CER failure load for predicting the creep failure load. It can be presumed that the CER failure load agrees with the creep failure load. Figure 7 displays the creep failure load Pc versus time to failure tc, where Pc is the fatigue failure load at load ratio R=0.95. The left side shows the experimental data, while right side exhibits the data shifted to T0=40°C using the shift factors for CER failure load. Since Pc at various temperatures can be superimposed smoothly, the time-temperature superposition principle is also applicable for Pc. The right side of this figure also displays the master curve for the CER failure load in the curve of thick line. Since the experimental Pc agrees well with CER failure load, the CER and creep failure loads depend scarcely on loading pattern. Master Curve of Fatigue Failure Load We regard the fatigue failure load Pf either as a function of the number of cycles to failure Nf or of the time to failure tf=Nf/f for a combination of f, R, T and denote them by Pf(Nf; f, R, T) or Pf(tf; f, R, T). Further, we consider that the CER failure load Ps(tf; T) is equal to the fatigue failure load at Nf=1/2 and R=0 by choosing tf=1/(2f). At this point, we introduce special symbols for fatigue failure load at zero and unit load ratios by Pf:0 and Pf:1 where the latter corresponds to creep failure load. To describe the master curve of Pf:0, we need the reduced frequency f1 in addition to the reduced

Fig. 7 Master curve of creep failure load

page_71 Page 72 time tf1, each defined by

Thus, the master curve has the form, An alternative form of the master curve is possible by suppressing the explicit dependence on frequency in favor of Nf as Recall that the master curve of fatigue failure load at Nf= 1/2 reduces to the master curve of CER failure load. The fatigue failure load Pf versus the number of cycles to failure Nf (Pf-Nf curve) for FRP joint at frequency f=5Hz and load ratio R=0.05 are shown in Fig.8. The Pf depends remarkably on temperature as well as Nf. The upper portion of Fig.9 shows Pf versus the reduced time to failure tf1. On the other hand, each point on the master curves of constant reduced frequency represents a number of cycles to failure. Connecting the points of the same Nf with these curves, the master curves of Pf for constant Nf are constructed as shown in the lower side of Fig.9. From this figure, it is found that the fatigue failure load depends scarcely on the number of cycles to failure. The Pf-Nf curves of FRP joint at f=0.05Hz and R=0.05 are shown in Fig.10. The solid lines in this figure indicate the predicted Pf-Nf curves at T=10 and 50°C obtained from the master curves of fatigue failure load as shown in the lower side of Fig.9. The predicted Pf-Nf curves agree with the experimental data. Therefore, the time-temperature superposition principle for CER failure load also holds for the fatigue failure load, and the hypothesis (B) is valid for fatigue failure load. Fatigue Failure Load for Arbitrary Load Ratio from which follows the creep failure load at any We have the master curve for creep failure load temperature T. The creep failure load, in turn, may be regarded the fatigue failure load at unit load ratio R=1 and arbitrary frequency f with tc=tf. Further, from the master curve for fatigue failure load at zero load ratio, we can deduce the fatigue failure load at zero load ratio for any frequency f and temperature T. Invoking the hypothesis (D), we propose a formula to estimate the fatigue failure load Pf(tf; f, R, T) at an arbitrary combination of f, R, T by

Figure 11 shows experimental data of Pf-tf for f=5Hz, R=0.5 and T=25, 40, and 70°C. The curves of R=0.05 and 0.95 respectively represent the least squares fit for experimental data of fatigue test of R=0.05 and 0.95. The curve of R=0.5 is calculated from equation (4) on the basis of the curves for R=0.05 and R=0.95. As can be seen, the predictions correspond well with the experimental data.

Fig. 8 Fatigue failure load versus number of cycles to failure at frequency 5Hz page_72 Page 73

Fig. 9 Master curves of fatigue failure load

Fig. 10 Fatigue failure load versus number of cycles to failure at frequency 0.05Hz Therefore, the hypothesis (D) is valid for fatigue failure load. From this figure, it is found that the fatigue failure load depends scarcely on load ratio at all temperature tested. Conclusion We had proposed a prediction method of fatigue failure load of FRP joint under arbitrary frequency, load ratio, and temperature. The method is based upon the four hypotheses, (A) same failure mechanism for CER, creep, and fatigue failure, (B) same time-temperature superposition principle for all failure loads, (C) the linear cumulative damage law for monotone loading, and (D) linear dependence of fatigue failure load upon load ratio. Tensile tests of FRP joint for CER, creep, and fatigue loadings were conducted for various temperatures. As a result, the characteristic timepage_73 Page 74

Fig. 11 Fatigue failure load for various load ratios temperature dependent fatigue behavior of this FRP joint is clarified by using this proposed method. Acknowledgements Partial support of this work by the National Renewable Energy Laboratory is hereby acknowledged. References 1. Aboudi, J. and G. Cederbaum, Composite Structures, 12 (1989), p.243. 2. Ha, S.K. and G. S. Springer, J. Composite Materials, 23 (1989), p.1159. 3. Sullivan, J.L., Composite Science and Technology, 39 (1990), p.207. 4. Miyano, Y., M. Kanemitsu, T. Kunio, and H. Kuhn, J. Composite Materials, 20 (1986), p.520. 5. Miyano, Y., M. K. McMurray, J. Enyama, and M. Nakada, J. Composite Materials, 28 (1994), p.1250. 6. Miyano, Y., M. K. McMurray, N. Kitade, M. Nakada, and M. Mohri, Advanced Composite Materials, 4 (1994), p.87. 7. Miyano, Y., M. Nakada, and M. K. McMurray, J. Composite Materials, 29 (1995), p.1808. 8. Miyano, Y., M. Nakada, and R. Muki, Mechanics of Time-Dependent Materials, 1 (1997), p.143. 9. Miyano, Y., S. W. Tsai, M. Nakada, S. Sihn, and T. Imai, Proc. ICCM/11,(1997), VI, p.26. 10.Miyano, Y., M. Nakada, M. K. McMurray, and R. Muki, J. Composite Materials, 31 (1997), p.619. 11.Nakada, M., T. Ishiguro, and Y. Miyano, Proc. ICCM/11, (1997), II, p.167. page_74 Page 75

Prediction of the Fatigue Behavior of Graphite-Epoxy Laminates Using Artificial Neural Network Anh Dung NGO1and Yahia OULD ABDESSLAM2 1,2Ecole de technologie supérieure, Université du Québec, Montréal (Québec), CANADA

Keywords: composite materials, graphite-epoxy, static strength, fatigue (materials), life prediction, artificial neural network Abstract The prediction of the fatigue behavior of Graphite-Epoxy laminates from a small database using an artificial neural network is explored. The input vectors of the back-propagation network are the fiber orientation and temperature of the experiments while the output vectors are the static strength and the fatigue strength of the laminate at 106 cycles. The network can regenerate the experimental curves with a precision of equal or less than 9% after the learning process. The predicted curves for the variable of fiber orientation are acceptable while the prediction for the temperature is poor due to the scarcity of the data. A new database is formed by 16 experimental curves and 15 predicted curves generated by the actual neural network in order to improve the size of the database. The new database is used to train the network, which can predict an S-N curve for the [±30]4S laminate with an acceptable precision. Introduction It is known that composite materials are affected by cyclic loading. The fatigue failure criterion is thus, necessary for the design of structures made of composites. The fatigue behavior of composites can be described by strength versus number of cycles to failure (S-N curve). The fatigue function which expresses the degradation in the strength of laminates as influenced by the number of cycles, is affected by many variables, such as fiber orientation, arrangement of layers and fabrication processes. This function is also influenced by external factors: temperature, frequency of cycling, pattern of loading and environments [1]. It is also well known that, the determination of fatigue function of composites is time consuming due to the scatter of the results and the slow frequency of the applied load, that is necessary to avoid the overheating of specimen. For this reason, it is useful to explore some new methodology helping to predict the fatigue performances of these materials from a small available database. Artificial neural networks are adaptive systems containing treatment units, which are mathematical analogies of the biological neurons. It can learn the patterns of an experimental result database and generates new data for any new condition without doing the experimentation. The present paper is concerned with the application of the neural networks to predict the fatigue behavior of Graphite-Epoxy laminates. The data are extracted from the experimental results of the work of Assa Rotem and H.G. Nelson [2]. Back-propagation Network for the Prediction of the Fatigue Behavior In this paper, a back-propagation network was used to predict the fatigue performance of Graphite-Epoxy laminates in different conditions. This technique is known as a suitable method for page_75 Page 76 prediction problem [3]. A neuronal model, an architectural design and training, characterizes an artificial neural network. Fig. 1 illustrates the schematic diagram of a neuronal model. In a neuronal model, synapses are represented by directional links with numerical weights (Wij), which control the importance of the in-coming signal (xj), to the receiving neurons. A neuron is represented by a node (k) carrying a summation function (Sk) and a transfer function (Gk) Weighted input signals (Wkjxj) are added together in the summation function of the node:

Where n is the input vector dimension. The output signal (ok) to the other connected nodes is calculated by the function (Gk)

Figure 1. Schematic diagram of neuronal model A back-propagation network is a collection of nodes and weighted links connected in layers (Fig.2). The network interacts with the outside word through the input and the output layers. The response of a network depends on its internal representation, which is composed of a layer or layers of hidden nodes. All nodes are fully connected to nodes in preceding and succeeding layers. There are no interconnections between nodes within layer. The architectural design of the network depends on the database implying the number of nodes in the input and output layers. These numbers correspond to the respective dimensions of the input and the output vectors. In the past, in order to take into account the complexity of the relations between the input and the output, it was a general practice to use many hidden layers. Recently, it has been demonstrated that one hidden layer was sufficient [4,5]. In this work, the dimensions of the input vectors were the fiber orientation and the page_76 Page 77 temperature of the experiments. Output vectors are the static strength and the fatigue strength of the laminates, at one million cycles.

Figure 2 Back-propagation network for the prediction the fatigue performance of angle ply laminates of Graphite-Epoxy The learning capability of the network enables it to acquire knowledge in response to different input database. Learning is accomplished by adjusting the weights when network errors are back propagated. The network error for an input vector i, according to the Delta learning rule, is computed as follows:

TJ : target output at the output layer oj : computed output at the output layer m : output vector dimension The network is considered to have acquired adequate knowledge when the sum of square errors (SSE) for the database is less than a certain tolerance. The error normally decreases when the number of iterative cycle increases. However, for the generalization, the error passed by a minimum in an optimal area, which is determined by comparing the predicted result with the experimental value. page_77 Page 78 The Neural Network of the MATLAB code, version 1994 was used as simulator for predicting the fatigue behavior of the angle ply Graphite-Epoxy laminates. The learning were accomplished after 250 000 iterations. The precision of the predicted S-N curves were evaluated with a frequency of 30 000 iterations. Fig. 3 shows the experimental S-N curves published by Rotem and Nelson and the predicted curves, in the same conditions, suggested by the neural network at the end of the learning process. The learning of the network was successful since the experimental curves and the predicted curves are very close. The maximum error was 9%.

Figure 3 Experimental and predicted S-N curves of different Graphite-Epoxy laminates at the end of the learning process Fig. 4 shows the predicted S-N curves for [±5]4S and [±20]4S laminates in comparison with the experimental curves of [0]8 and [±15]4S laminates. It can be seen that the results were acceptable. The prediction for the variable of temperature was unfortunately not correct due to the scarcity of the data. In order to improve the situation a new database was formed by 16 experimental curves and 15 predicted curves which were generated by the network. An experimental curve in the new database was not used to train the new network, which had a similar constitution of the previous network. Fig. 5 shows the predicted S-N curve and the experimental curve for the [±30]4S at 74°C. The maximum error is 1.56%. Conclusion The approach using neural network for prediction the fatigue behavior is worth to develop. The model used in this work demonstrated the possibility to predict the fatigue performance of composite materials economically. page_78 Page 79

Figure 4 Experimental and predicted S-N curves of different Graphite-Epoxy laminates with various fiber orientations.

Figure 5 Experimental and predicted S-N curve of Graphite-Epoxy [±30]4S laminates. page_79 Page 80 References 1. B.D. Agarwal, L.J. Broutman, Analysis and performance of fiber composites, John Wiley & Sons (1990). 2. Assa Rotem & H.G. Nelson , Fatigue Behavior of Graphite-Epoxy laminates at elevated temperatures, ASTM STP 723 (1981), p.152. 3. R.C. Eberhart, R.W. Dobbins (editors), Neural Network PC Tools: A Practical Guide, Academic Press Inc, San Diego CA (1990). 4. K. Hornik, M. Stinchcombe and H. White, Multilayer Feedforward Networks are universal Approximators, Neural Networks, Vol.2 (1989), p. 359. 5. K.I. Furnashi, On the approximative realization of continuous mappings by neural Networks, Neural Networks, Vol.2 (1989), p.183. page_80

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Impact Compressive Failure of GFRP Unidirectional Composites Jianming YUAN1,2, Nobuo TAKEDA1, Anthony M. WAAS3 1Center for Collaborative Research (CCR), The University of Tokyo 4-6-1 Komaba, Meguro-ku, Tokyo 153-8904, Japan 2Present Address: Department of Mechanical and Production Engineering, National University of Singapore, 10 Kent Ridge Crescent, Singapore 119260 3Department of Aerospace Engineering, University of Michigan FXB Bldg, Ann Arbor, MI 48109-2118, USA Keywords: GFRP, impact compression, split Hopkinson pressure bar, fiber volume fraction, strain rate effect, temperature effect, kinking, splitting Abstract Compressive impact tests of unidirectional glass fiber reinforced vinyl ester matrix composites (GFRP) were carried out using the split Hopkinson pressure bars. The dynamic stress-strain curves of unidirectional composites of six different fiber volume fractions and pure matrix were obtained at the strain rate of 103 s-1. Impact recovery tests were also performed to study the impact compressive damage evolution in composites. The temperature dependence up to 100°C were examined to study the temperature effect on the compressive strength. Quasi-static compressive tests of the same specimens at the strain rate of 10-3 s-1 were also conducted for comparison. Failed specimens were examined by optical microscopy. Kinking followed by splitting was found to be the main controlling failure mechanism. GFRP exhibited ductile failure for lower fiber volume fractions, but brittle failure for higher fiber volume fractions. As the temperature increased, the failure mode changed from kinking to microbucking. Experiments showed that the strain rate has a strong effect on the compressive strength. Some theoretical prediction of the compressive strength was also made based on the failure mechanism and test data. 1 Introduction For the fiber-reinforced polymer matrix composite, the compressive strength is usually lower than the tensile one. Although numerous studies related to the compressive behaviors have been conducted, the mechanism of composite failure in compression has not been well understood. Much of this understanding of compressive failure mechanism has been obtained on the basis of static loading [1]. Relatively little work has been conducted to study the dynamic compressive behavior of unidirectional composites [2,3]. It is difficult to observe and record the damage evolution process in a static test. Because the recovery impact test [4] is capable of loading a specimen to a certain level followed by unloading it for microscopic analysis, it can be used in studying the compressive failure process. In this paper, the dynamic compressive behavior of glass fiber reinforced unidirectional vinyl ester was studied experimentally. Quasi-static compressive tests of these composites were also conducted for comparison. Research focused on the compressive damage evolution and failure mode in these unidirectional composites as well as the factors influencing the compressive strength. The factors included the strain rate, the fiber volume fraction and the temperature. Prediction of the compressive strength was also analyzed based on the experimental results and some analytical models. 2 Experiments 2.1 Test Specimens Reinforcement and matrix resin used in this study were E-glass fiber and vinyl ester (Derakane 411-C50), respectively. The unidirectional GFRP composites with six different fiber volume fractions ranging from 10% to 60% and pure matrix were fabricated within a long glass tube, then, circular cylindrical specimens with approximately 7 mm in diameter and 5 mm in length were machined. Specimens of all fiber volume fractions were tested under impact and static loadings. GFRP of Vf = 40% was used in impact recovery tests in order to study the compressive damage evolution under impact loading. Both Vf = 50% GFRP and pure matrix were used in impact tests at high temperature up to 100°C in order to study the temperature effect on the compressive failure mechanism. 2.2 Test Procedures Impact compressive tests were performed using the spilt Hopkinson pressure bars (SHPB), in

page_81 Page 82 which the stress waves in the input and output bars were recorded to calculate the dynamic stress-strain curve of specimens. In order to improve the accuracy of the dynamic stress-strain curves of GFRP whose failure strain is small, a method considering wave propagating in specimens was applied in data process of SHPB tests [5]. Recovery impact compressive tests were performed using a compression-type improved SHPB apparatus [2, 4] to recover the specimen unloaded from any specified point in the stress-strain curve. Impact tests under high temperature were conducted by using a heat cable to heat specimens to a desired temperature. Failed or recovered GFRP specimens were observed by the optical microscopy. Steel rings were used to confine specimen ends in some tests. The rings restrained radial expansion of the specimen ends under compression. By comparing the recorded stress-strain curves with and without using rings, it was found that the rings did not alter the mechanical behavior of the specimen. Tests also showed the rings could prevent post-failure in specimens, and therefore, it is helpful in recovering failed specimens for microscopic analysis [6]. 3 Test Results and Analysis 3.1 Effect of Fiber Volume Fraction and Strain Rate The typical stress-strain curves of GFRP of different fiber volume fractions under impact and static loadings are shown in Fig. 1 and Fig. 2, respectively. Figure 3 shows the compressive strength of GFRP versus the fiber volume fraction under static and impact loadings. The failure of GFRP composites changes from ductile to brittle at the fiber volume fraction of 30%40% for both static and impact loadings. Correspondingly, a nonlinear increase of the compressive strength is noted. The ratio between the dynamic compressive strength and the static one ranges from 2.5 for pure matrix to 1.71.8 for GFRP with high fiber volume fractions.

Fig. 1 Stress-strain curves of GFRP and matrix under impact loading (strain rate @103 s-1)

Fig. 2 Stress-strain curves of GFRP and matrix under static loading (strain rate @10-3 s-1)

Fig. 3 Compressive strength of matrix and GFRP with different fiber volume fractions

Fig. 4 Compressive failure mechanism in GFRP (a) interior kinking and (b) surface kinking page_82 Page 83 3.2 Compressive Failure Mechanism and Damage Evolution Process Both kinking and longitudinal splitting were found in failed GFRP specimens. Figure 4 shows kinking band occurring on the surface and in the interior of specimens. When the fiber volume fraction was very small, the interior kinking band could be seen. For a higher fiber volume fraction, by carefully separating GFRP along the crack of longitudinal splitting, two types of kinking bands were found on the separated surface, i.e., the kinking band occurred in plane and out of plane. It can be also found that splitting occurred in the region of kinking band and some fibers were broken by tensile stresses due to bending. Impact recovery tests further revealed the compressive damage evolution in GFRP. Under a low impact loading level, an initial damage (whitening area) occurred first in the region near to the specimen ends. Under a higher impact load, a kinking band was formed, and then, some splitting was also formed. It is concluded that the compressive failure mechanism of GFRP is fiber kinking followed by the longitudinal splitting. 4 Prediction of the Compressive Strength 4.1 Kinking-Controlling Compressive Strength Form the above test results, it is concluded that the same failure mechanism of kinking-splitting controls the dynamic and static compressive failures of the composite with higher fiber volume fraction, but for lower fiber volume fractions, the final controlling factor of compressive strength is the longitudinal compressive property of matrix. It is known that the kinking phenomenon in unidirectional composites is related to the shear property of the composite as well as the initial misalignment of fibers. Misaligned fibers easily undergo microbuckling or kinking under compressive loading. Taking account of matrix plastic deformation and the initial fiber misalignment, Hahn and William [7] provided a predicted compressive strength of unidirectional composites as

where is the ratio of the magnitude to the wavelength of the sinusoidal initial imperfection, GLT is the composite shear modulus, and gLT is the average shear strain in the composite. Moreover, if the shear stress-strain relationship of the composite can be approximated as that of an elastic-perfectly plastic materials, gLT in Eqn. (1) is equal to the yield shear strain [8]. The predicted compressive strength based on Eqn. (1) is plotted in Fig. 5 with the following simplification and approximation: GLT = Gm/(1-Vf), Gm = Em/2(1+ nm), Poisson's ratio of matrix nm = 0.35, and Em = 4.2 GPa. Here Gm and Em are the shear modulus and Young's modulus of matrix, respectively. The values of represents the initial misalignment of fibers, are determined by data fitting.

in which

Fig. 5 Comparison of the predicted compressive strength with test data at room temperature

Fig. 6 Comparison of the predicted compressive strength with test data at high temperatures page_83 Page 84 4.2 Matrix-Controlling Compressive Strength For composites with lower volume fiber fractions, the stress-strain curves show ductile failures. Supposing the final strength is controlled by the yield of matrix, the rule of mixture results in a compressive strength of unidirectional composite as

where e is the yield strain of the matrix under compression, which is a little less than the failure strain and determined from the stress-strain curves of matrix in Figs. 5 and 6. The predicted compressive strength based on Eqn. (2) is also plotted in Fig. 5. The predicted compressive strength is reasonably consistent with the test data. 4.3 Temperature Effects on Compressive Strength

At high temperatures, the compressive failure mechanism is fiber microbuckling. Only considering matrix elastic deformation, Rosen's model gives the compressive strength as [9]

Comparison of the predicted compressive strength based on Eqns. (3) and (1) with test data is shown in Fig. 6. Equation (1) underestimates the compressive strength at higher temperature, it is because of the assumption that in the calculation. In reality, the yield shear strain of composite, gLT, is temperature-dependent, and therefore, Y should also be temperature-dependent. Equation (3) can predict the compressive strength qualitatively. 5 Conclusions The dynamic and static compressive behavior of GFRP was studied experimentally. The following conclusions were drawn. (1) The compressive failure mechanism is fiber kinking followed by longitudinal splitting. The longitudinal splitting is caused due to the required kinematic compatibility of fiber kinking. Splitting causes specimens to lose the total loading capability, but those only with kinking damage still keep some loading capability. Glass fibers often break due to the tensile stresses due to bending in the kinking band. As the fiber volume fraction increases, the compressive failure changes from ductile to brittle, which results in an nonlinear increase in the compressive strength.(2) Within the tested temperature range up to 100°C, the Young's modulus of composites is constant. The compressive failure in GFRP changes from kinking at room temperature to microbuckling at 100°C. The compressive strength decreases significantly at approximately 75°C. The temperature has a strong effect on the mechanical behavior of matrix. The decrease in the compressive strength at high temperatures is due to the temperature-softening effect of matrix. (3) Although the composites have the same compressive failure mechanism under impact and static loadings, the strain rate has a strong effect on the compressive failure strength. Under impact loading, the increase in the compressive strength is due to the increase in the matrix strength. Acknowledgments The authors acknowledge the support of the grant-in-aid for international cooperation from Monbusho throughout the present study. The author also thank Mr. Mikio Hiramatsu for the assitance in experiments. References [1] C.R.Schultheisz and A.M.Waas, Prog. Aerospace Sci., 32 (1996), p. 142. [2] N.Takeda, L.Wan, M.Hiramatsu and J.Yuan, Trans. JSME, Ser. A, 63 (616) (1997), pp. 25982603. [3] J.Lankford, Composites, Part A, 28A (1997), pp. 215222. [4] S.Nemat-Nassar, J.B.Isaacs and J.E.Starrett, Proc. R. Soc. Lond. A435 (1991), pp. 371391. [5] J.Yuan, N.Takeda and A.M.Waas, Accepted for Experimental Techniques (1998). [6] J.Yuan, N.Takeda and A.M.Waas, Proc. 1st Asian-Australasian Conf. Comp. Mater., Osaka (1998). [7] H.T.Hahn and J.G.Williams, Composite Materials, ASTM STP 893 (1986), pp. 115193. [8] D.Hull and T.W.Clyne, Introduction to Composite Materials (2nd Ed.), Cambridge University Press (1996), pp. 177178. [9] B.W.Rosen, Fiber Composite Materials, ASM (1965), pp. 3775. page_84 Page 85

THERMOPLASTIC COMPOSITES II page_85

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Progressive Crushing of Compression-Molded Thermoplastic Composite Tubes Hiroyuki KAWADA*, Takeshi HONDA**, Maiko TAKASHIMA** and Hajime SATOH** *Department of Mechanical Engineering, Waseda University **Graduate School of Waseda University 34-1, Okubo, Shinjuku, Tokyo 1698555, Japan ***Yokohama Rubber Co.LTD, R&D Center Oiwake, Hiratuka, Kanagawa, 254-0047, Japan Key Words: GF/Nylon6 / Specific energy absorption / Progressive crushing / Morphology / Trigger geometry Abstract Compression tests on GF/Nylon6 tubes were performed to clarify the mechanism of the initial failure process and the energy absorption behavior. It was found that all the specimens were crushed in a progressive crushing mode regardless of the trigger geometry, but that the specimens with an asymmetric trigger demonstrated better absorbing characteristics than the symmetric specimens. The formation mechanism of the debris wedge was examined through a step-wise morphology analysis. It was found that changes in the trigger geometries affected the performance of the initial failure process which was characterized by the formation of a debris wedge formation process. It was also shown that at high testing speeds in the range of practical use, GF/Nylon6 tubes exhibited similar energy absorbtion behavior to that of a CF/PEEK system. 1 Introduction Composite materials are actively used for various applications especially as a structural material. Composite materials also have excellent energy absorption characteristics and, therefore, recently, the importance of progressive crushing of FRP tubes1)~4) has been recognized as it is directly applicable to the automobile industry. Some studies on the effect of laminate constitution, testing speed and friction on the energy absorption mechanism of FRP tubes have been reported. It had been shown from static tests that CF/PEEK tubes, which are thermoplastic composites, have high energy absorption behavior. This is because PEEK has a high fracture toughness 5), 6). The crushing of CF/PEEK tubes with different laminate constitutions at different testing speed have been examined 7). Stable progressive crushing, is affected by trigger geometry 8), 9) but the influence of trigger geometry on the energy absorbing characteristics and crushing mechanisms for thermoplastic composites is not well defined. In this paper different kinds of trigger were examined for GF/Nylon6 tubes, in order to investigate the influence of the trigger and testing speed on the crushing behavior and energy absorption characteristics. 2 Experimental Tube specimens which were fabricated using unidirectional prepreg of GF/Nylon6 by an internal pressure method had a fiber volume fraction of 50% and an 18-ply lay-up. The ratio of warp to hoop was 17:3. All of the tubes had a wall thickness of about 3.0mm. To initiate progressive crushing, a chamfer was machined at one end of each tube (Fig.1). Tests were performed using a testing machine at a constant cross-head page_87 Page 88 Typical symmetric triggers, like Type-C, induce buckling failure which makes space between the adjoining laminae. Debris which was produced from buckling accumulated in that space, and built up the amount of debris in the wedge . As crushing progressed, the debris wedge grew, becoming more defined and further producing failure . However, because the trigger was symmetric, after debris wedge formation, shear fracture doesn't occur. And so the debris causing the development of the debris wedge is only the debris caused by friction . At the stage when the trigger section has stop compressing, it can be seen that the size of the debris wedge is smaller than that of the asymmetric triggers. For the symmetric triggers, the central crack is in the center of plate thickness immediately after the start of crushing. After that there is no big shear failure, and the central crack was not observed to move. 3.2 Test Speed

Fig.4 shows the typical load-displacement curves for each test speed. From this figure it can be seen that the medium test speed (7.50 ´ 103mm/min) shows little decline from the maximum load to the stable load.

Fig.4 Compressive load-displacement curves At high speed, the load decline from maximum load is larger and the stable load is lower than the other speeds. Fig.5 shows specific energy absorption Es for each testing speed. These Es values, in Fig.5 were calculated from eq.(1). Fig.5 also shows the results of the CF/PEEK tubes. We can classify Fig.5 into three domains, low speed (1.00 ´ 10~1.00 ´ 102mm/min), medium (1.00 ´ 103~7.50 ´ 103mm/min), and high (4.86 ´ 105~9.18 ´ 105mm/min). From Fig.5, we can see the Es value of GF/Nylon6 tubes increase slightly in the medium speed domain compared with the low speed domain, but considering all testing speed domains the Es values are almost constant. The data implies that the Es value of GF/Nylon6 is not affected by test speed. In the CF/PEEK tube system, which is dependent on the test speed, the Es value of CF/PEEK tube system is higher than the GF/Nylon6 tube system in the low and medium speed domain. In the high speed domain, the Es values of both systems are almost equal, which is very interesting. The tests performed using the impact test machine for the GF/Nylon6 tubes (high speed region), exhibited the same final fracture morphology as the other test speeds. In the case of CF/PEEK tubes this is not the case and the specimens shattered. Since high speeds are of practical use, GF/Nylon6 and CF/PEEK are comparable as both have similar

Fig.5. Dependence of Es value on testing speed page_88 Page 89

Fig.1 Various trigger and dimension of specimen speed of 1.00 ´ 10mm/min to a displacement of 30mm. To study the characteristics at higher speeds, tests were performed using a impact test machine. The energy absorbed, as defined by the area under the load-displacement curve, is a specific energy absorption ''Es" with resect to the specimen's unit mass, and is thus given by

where P is load, u is displacement, V is progressive crushed volume and r is material density. Provided that the crush load becomes stable on the load-displacement curve integral calculus can be performed. In this study the load became stable at about 20mm so it was possible to the integration between 0 and 30mm. 3 Result and Discussion 3.1 Initial Failure Process of Each Trigger Type Typical cross sections of asymmetric trigger Type A and symmetric trigger Type C are shown in Figs. 2 and 3 respectively. Cross sectional photographs of each trigger geometry are at displacements of 2mm, 3mm, 5mm, and are presented alongside their load-displacement curves. Break aspect at the same displacement are different for each trigger geometry. Asymmetric triggers (Type-A,B) have debris which is a factor of the debris wedge between laminae in the early stage . With progressive crushing, debris, which was produced from small and big shear fractures, make up the debris wedge . From these processes, the growth of the debris wedge around the asymmetric trigger can be obsorved. Debris which was produced on the surface of the laminate, as a consequence of friction between the platen and the frond usually became a debris wedge. In particular, in the stable region, debris, which was produced only because of friction, became a debris wedge. The growing process of debris, which was produced due to small and big shear fractures, contribute to the "quick growth of the debris wedge" in the initial failure process, which can be easily observed from the cross-sectional analysis (Fig.2). Asymmetric triggers produce shear fractures, which appear in the area of debris wedge (inner or outer side).

Fig.2. Cross-sectional photographs of initial failure process

Fig.3. Cross-sectional photographs of initial failure process page_89 Page 90 energy absorption characteristics in this domain. 3.3 Morphology Fig.6,7 shows typical cross-sectional photographs for each testing speed. The crushzone morphology, Fig.6 (low speed), shows buckling failure between the outer laminae and the second on both the outer and inner frond. Also no interlaminar cracks were found in that region. Debris accumulating between both the fronds appears as a large wedge shape and results in a frond with a low radius of curvature. At medium speeds (Fig.7), many interlaminar cracks can be observed as well as buckling and shear fracture. These fractures developed continuously in the laminae close to the inner wall. These fractures are most apparent between the outer laminae. The form of the debris wedge is more well defined than at low speeds and makes the curvature of frond higher. From this it can be confirmed that there are many bending fractures at the frond and many more fractures have developed compared with the low testing speeds.

Fig.6. Cross-sectional photograph (1.00 ´ 10mm/min)

Fig.7. Cross-sectional photograph (7.50 ´ 103mm/min) 4 Conclusions Trigger geometries can control behavior of load in the initial failure processes. The formation of the debris wedge is caused from the shape of the trigger cross section. The source of the debris wedge for asymmetric triggers comes from the large shear fractures which are caused by the propagation of a central crack. From the examination of the L.U.F value on the Es value, the Type-A trigger is the most ideal for energy absorption. For GF/Nylon6 tubes, there is no influence of testing speed on the energy absorption. At high speeds GF/Nylon6 tubes have similar energy absorption to CF/PEEK tubes. 5 References 1) D. Hull, Composites Science and Technology, 40, 377 (1991). 2) I. Sigalas, M. Kumosa and D. Hull, Composites Science and Technology, 40, 265 (1991). 3) D. Hull and J. C. Coppola, Materials and Processing-Move into the 90's, Elsevier Science Publishers B. V., Amsterdam, 29 (1989). 4) A. H. Fairful and D. Hull, Wiley interscience, 255 (1989). 5) H. Hamada, J. C, Coppola, D. Hull, Z. Maekawa and H. Satoh, Composites, 23, 245 (1992). 6) H. Hamada, S. Ramakrishna and H. Satoh, Composites, 26, 11, 749 (1995). 7) H. Kawada, T. Okabe, M. Mawatari, H. Satoh, Journal of the Society of Materials Science, Japan, 46, 645 (1997). 8) I. Sigalas, M. Kumosa, and D. Hull, Composite Science and Technology, 40, 265 (1991). 9) M. J. Czaplicki, R. E. Robertson, and P. H. Thornton. Composite Science and Technology, 40, 31, (1991) page_90 Page 91

Development of "Fibro-Composites"Morphology of PBT/polyolefin blend Kazuo Kitagawa*, Hiroyuki Hamada** and Takeshi Semba** * Kyoto Municipal Institute of Industrial Research

Chudoji, Simogyo-ku Kyoto, 600-8813, Japan **Kyoto Institute of Technology Matsugasaki, Sakyo-ku, Kyoto 606-8585, Japan Key word Fibro-Composites, PBT, PE, Morphology Abstract PBT/PE is immiscible polymer blend. We found that PBT component could form a fibro-reinforcement in PE matrix during compounding and injection molding. PBT/PE blend were subjected to drawing at cooling process in compounding. Injection moldings were fabricated with the drawing compound pellets. Processing conditions are shown as follows; cylinder temperature was 190,210,215 and 220°C, other conditions were fixed. The fibroreinforcements of PBT were formed at the skin layer, on the other hand the PBT component was sphere shape at the core layer in both strands. Comparison with drawing and non-drawing strands, mechanical properties of drawing strand was better and the fibro-reinforcements of the PBT were much formed. In the case that cylinder temperature was 190°C, the fibro-reinforcements could not formed and the fracture aspect indicated quite brittle behavior. In 220°C, the PBT component was completely dissolved and formed sphere shape during injection process. In 215°C, the fibro-reinforcements of PBT were formed with optimum condition which the reinforcements were enough long to reinforce. It is cleared that drawing method in compounding encourages for PBT component to form fibroreinforcement and to improves mechanical properties of extruded strand. In injection molding process, the cylinder temperature played a vital role for Fibro-Composites. 1 Introduction In polymer blend, dispersive component can often form reinforcement in matrix. This phenomenon have been studied and developed, furthermore applied to industrial products 1). Takayanagi et al. 2) proposed "molecular composites" whose size of dispersion component accomplish nano-dimension in matrix. Recently, LCP blend system is extensively studied, as the results of these studies, LCP component forms sphere and fibril reinforcement, and it is cleared that dispersive phase played an important role in mechanical properties of composites 3). Blend ratio, viscosity ratio, interaction at of interface and processing conditions have the great influence on the size of dispersion phase in these blend system, the size of LCP component is about from 10-7 to 10-5m. Farikov et al. investigated micro-fibril composites of PET/PA6 blend system. They carried out SEM observation and WAXD pattern, so that it is cleared that PET component forms micro-fibril and orients along flow direction 4) 5). page_91 Page 92 In this study, PBT/PE blend system which is one of the immiscible polymer blend was investigated. The PBT component have been able to form fibro-reinforcement in PE matrix (Fig.1). The size of fibro-reinforcement accomplish nano-dimension, and much longer than micro-fibril. The purpose of this work is development of PBT/PE "Fibro-Composites", and the influence of drawing at compounding and processing temperature in injection process were investigated.

Fig.1 "Fibro-Composite" of PBT/PE blend. 2 Experimental 2.1 Material and Thermal Analysis

Materials used in this study were PE (UP polyethylene LD352, MI=100 :Union Polymer Co., Ltd.) and PBT (Juranex 2002, MI=15 : Polyplastics Co., Ltd.). Melting points(Tm) of these materials were measured by DSC (DSC7: PERKINELMER Co., Ltd.). The temperature of climb was 20°C/min. , and the scanning were measured two times. Tm of second scan decided processing temperatures in this study. 2.2 Compounding PBT/PE blend pellets were compounded by twin screw extruder (LABOTEX300: Japan Steel Works Co., Ltd.). Processing conditions are as follows; cylinder temperature was 220°C, screw revolution was 200rpm. PBT/PE blend ratio was 5/5. Cooling method of the extruded strand was water-cooling. Schematic diagram of cooling methods is shown in Fig.2. Fig.2-(a) shows the cooling method of non-drawing strands, in this method extruded strand was cooled from the nozzle of die to edge of water bath. On the other hand, Fig.2-(b) shows the cooling methods of drawing strands, in this methods extruded strand was slightly cooled at the nozzle of die to frozen only surface, and was subjected to drawing with a roller, and the strand which was completely frozen up in the water bath was cut to the pellets.

Fig.2 Schematic diagram of cooling methods in compound. 2.3 Injection Molding Dumbbell type tensile specimens were fabricated with the compounded pellets subjected to drawing. Injection molding machine was an inline screw type injection page_92 Page 93 molding machine(Plaster Ti30-F6: Toyo machinery Co., Ltd.). Processing conditions are as follows; cavity temperature was 50°C, injection speed was 17.50cm3/sec, cylinder temperature were 190, 210, 215 and 220°C. 2.4 Tensile Test Tensile test for drawing and non-drawing strands and injection molding were performed with a universal testing machine(Instron 4206 type). Test conditions are as follows; gage length was 115mm, tensile speed of extruded strands and injection moldings were 1mm/min.and 2mm/min., respectively. 2.5 SEM Observations SEM observations of internal structure for extruded strand and injection moldings were performed. Cross section along flow direction and fracture surface after tensile test were observed for all of the specimens. In the former, center area for extruded strands and center and surface areas for injection moldings were observed, as shown in Fig.3.

Fig.3 Observation area of extruded strands and injection moldings 3 Results 3.1 Thermal Analysis The Tm of PE was 106.4°C. Double peaks could be measured about the Tm of PBT; the Tm1 was 214.2°C, Tm2 was 223.2°C. 3.2 Tensile Test 3.2.1 Tensile Test of Extruded Strands Fig.4 shows stress-displacement curves of extruded strands. In the case of non-drawing strand the fracture occurred at 0.27mm after maximum load. On the other hand the fracture of drawing strand occurred at 0.73mm after maximum load. Fig.6 and Fig.7 show the tensile moduli and strength of extruded strand, respectively. In both tensile modulus and strength, the drawing strand had higher

Fig.4 Stress-displacement curve for extruded strands.

Fig.5 Tensile modulus for extruded strands. page_93 Page 94 mechanical properties than those non-drawing strand.

Fig.6 Tensile strength for extruded strands.

3.2.2 Tensile Test of Injection Moldings Fig.7 shows load-displacement curves of injection moldings from drawing pellets. The elongate lengths after maximum load about 210, 215 and 220°C were 1.16, 0.98 and 0.90mm. On the other hand, the elongate length after maximum load about process of 210, 215 and 220°C was that the initial fracture arose at internal area and the final fracture occurred at the surface area. On the other hand, the fracture of 190°C was completely brittle fracture.

Fig.7 Load-displacement curve for injection moldings Fig.8 shows tensile modulus, and fig.9 shows tensile strength of injection moldings. In both tensile modulus and strength, the mechanical properties of 190°C was very small, and 215°C had the best mechanical properties.

Fig.8 Tensile modulus for injection moldings.

Fig.9 Tensile strength for injection moldings. 3.3 SEM Observation 3.3.1 Internal Structure of Extruded Strands Fig.10 and Fig.11 show SEM photographs of cross section along flow direction for non-drawing strand, and fracture surface for that. The non-drawing strand forms skin-core structure, the PBT component forms faibro-reinforcement and orient along flow direction at the skin layer, and the PBT component forms sphere shape at the core layer. The diameter of PBT faibro-reinforcement was under 4.4 m m before tensile test, and after tensile test the diameter was under 2.3 m m. Fig.12 and fig.13 show SEM photographs of cross section along flow direction for drawing strand, and fracture surface for that. The drawing strands also forms skin-core structure. However, the PBT component forms a little fibro-reinforcement at the page_94

Page 95 core layer. The diameter of PBT faibro-reinforcement was under 2.4 m m before tensile test, and after tensile test the diameter was under 2.0 m m. Percentage of skin layer was calculated by measuring area of skin and core layer. As the results of that, the percentage of skin layer for drawing strand was 37%, and that about non-drawing strand was 23%.

Fig.10 SEM photographs of cross section along flow direction for non-drawing strand.

Fig.11 Fracture surface for non-drawing strand.

Fig.12 SEM photograph of cross section along flow direction for drawing strand.

Fig.13 Fracture surface for drawing strand. 3.3.2 Internal Structure of Injection Moldings Fig.14 shows SEM photographs for injection moldings along flow direction whose processing temperature was 215°C. In the center area, fibro-reinforcement distributed through whole area, and the diameter of fibroreinforcement at the core layer was larger than that at skin layer. In the surface area, fibro-reinforcement which has a smaller diameter than that in center area distributed through whole area. Fig.15 shows SEM photographs of fracture surface whose processing temperature was 215°C. The fracture surface was composed of three areas as follows; the first was the nearest surface area where the distribution of the fibroreinforcement was very fine, the fracture aspect of the second area was ductile and the diameter of the elongated fibro-reinforcement was under 8.8 m m, the fracture aspect of the third area was brittle and the diameter of the fibro-reinforcement was under 26.0 m m.

Fig.16 shows SEM photographs of the cross section along flow direction for injection moldings whose processing temperature was 190°C. The fibro-reinforcement could be hardly observed and the most of the PBT component was sphere shape at the center area. However, the fibro-reinforcement was widely distributed for injection moldings at the surface area. Fig.17 shows SEM photographs of fracture surface whose processing temperature was 190°C. In this case, the three areas of the cross section along flow direction could not be observed, and a lot of sphere or column shape of the PBT component were observed through the whole area. However a little ductile area was lying on the center page_95 Page 96 of the fracture surface. Fig.18 shows SEM photographs of the cross section along flow direction for injection moldings whose processing temperature was 220 °C. The distributive aspect of the PBT component was like the injection moldings processed at 215°C, but the PBT component was sphere shape at the core layer. Comparison with observation of center and surface area, fine fibro-reinforcement could be much

Fig.14 SEM photographs for injection moldings along flow direction processed at 215°C

Fig.15 SEM photographs of fracture surface processed at 215°C. observed at surface area. Fig.19 shows SEM photographs of fracture surface whose processing temperature was 220°C. In this case, the above described three areas was formed. However, the diameter of the fibro-reinforcement at the ductile fracture area was under 2.5 m m, this value is smaller than 8.8 m m in the 215°C specimen. At the brittle fracture area, very short fibro-reinforcement was generated, so that delamination occurred all over this area.

Fig.16 SEM photographs for injection moldings along flow direction processed at 190°C

Fig.17 SEM photographs of fracture surface processed at 190°C. page_96 Page 97 The internal structure of injection moldings whose processing temperature was 210°C was neutral aspect between 190° and 215°C.

Fig.18 SEM photographs for injection moldings along flow direction processed at 220°C

Fig.19 SEM photographs of fracture surface processed at 220°C. 4 Discussion 4.1 Internal Structure and Tensile Properties of Extruded Strands The drawing and non-drawing strand were fabricated in compound. Comparison with the mechanical properties of these strands, drawing strand had higher mechanical properties than non-drawing strand. A skin-core structure was formed in the internal structure of these strands, and the percentage of skin and core layer is important factor for mechanical properties of extruded strands. The fibro-reinforcement of the skin area showed ductile fracture aspect, from this observation, the fibro-reinforcement resisted to tensile load. In the case of drawing strand, the ratio of the skin layer was large, and a little fibro-reinforcement was distributed at the core layer. From these results fibroreinforcement of the PBT component play a important role in mechanical properties, resulting in good mechanical properties. The compound temperature of 220°C is melting point of PBT, so that PBT and PE were enough melted in extruding cylinder. In die-casting, however, the surface layer of the PBT was formed fibro-reinforcement by shear strain from die wall. On the other hand the core layer of the PBT was not affected by shear strain, so that the PBT component was formed sphere shape. Moreover, the drawing method encourages the fibro-reinforcement at the core layer. 4.2 Internal Structure and Tensile Properties of Injection Moldings

The processing temperature were 190, 210, 215, and 220°C. The mechanical properties of the injection moldings processed at 190°C was the worst, and the injection page_97 Page 98 moldings whose processing temperature was 215°C had the best properties. The difference of the internal structure of these specimens could be observed. In the case of 190°C, the processing temperature is lower than Tm of the PBT, so that the PBT component was not melt. By this reason, the internal structure of the extruded strands directly reflected to that of injection moldings. A part of the fibro-reinforcement whose percentage is 37% in extruded strand directly preserved or disappeared by shear and elongate strain. A part of the PBT component whose shape is sphere and the percentage is 63% in extruded strand became to column shape, but the most of that was preserved sphere shape. Most of the PBT component was sphere shape, and can not reinforce injection moldings, resulting in the low mechanical properties. The injection moldings whose processing temperature was 215°C had the best mechanical properties. The internal structure of this injection moldings quite different from that of 190°C, and the fibroreinforcement distributed through whole area. The processing temperature of 215°C is Tm1 of the PBT, and a crystalline structure which corresponds to the DSC peak of Tm2 is remained. The melting PBT in injection cylinder is oriented on the base of crystalline and drawn along flow direction by shear and elongate strain in injection molding process. The fibro-reinforcement distributed through whole area, so that the injection moldings whose processing temperature was 215°C has good mechanical properties. The internal structure of injection moldings whose processing temperature was 220°C was almost same as that of 215°C, however the PBT component whose shape was sphere distributed at the core layer, resulting in the different mechanical properties from that of 215°C. The crystalline structure is completely melted, the reason why the processing temperature 220°C is nearly to Tm2. In this case, the energy of Tm1 and Tm2 must be radiated, and the time for frozen up is long. Hence, the fibro-reinforcement of core area is easy to relaxed. 5 Conclusion The purpose of this study is development of PBT/PE "Fibro-Composites". In the work the influence for the mechanical properties of drawing at compounding and injection molding temperature were investigated through the morphology observation. Several conclusions from these experimental results are shown as follows; (1) The PBT component in PE matrix could form fibro-reinforcement during compounding and injection molding. (2) The mechanical properties of PBT/PE blend strand can be improved by drawing method at cooling process in compound. (3) It is cleared that the processing temperature in injection molding play a important role in internal structure and mechanical properties. References (1) L. A. Utracti, "Polymer Alloys and Blends" (1989) Hanser Publishers. (2)M. Takayannagi, IUPAC 32nd Inter. Symp. on Macromolecules, Preprint, p.36(1988). (3)P. R. Subramanian and A. I. Isayav, SPE Tech. Pap. 48, 489,(1990). (4)M.Evastatiev, N.Nicolov and S.Fakirov: POLYMER Vol.37 No.20 (1996) (5)J.L. Kardos, and Fakirov and S.Fakirov: J. Polym. Eng. Sci.1975, 15, 183 page_98 Page 99

Mechanism of Fatigue Fracture of Glass Fiber Reinforced Nylon 66 Satoshi Odaka, Takashi Kuriyama, Masaya Kotaki and Ikuo Narisawa Department of Materials Science & Engineering, College of Engineering, Yamagata University, Jonan, Yonezawa city 992-8510, JAPAN

Keywords; Fatigue, Glass fibre reinforced nylon 66, Acoustic velocity, Acoustic emission Abstract The mechanism of the tensile fatigue fracture of glass fiber reinforced nylon 66 composites has been studied using acoustic emission (AE), ultrasonic measurement and optical microscopy (OM) techniques. The acoustic velocity of the edge of specimens decreased by the cyclic loading. This behavior was correlated to matrix cracks and fiber breakages at the edge of the specimens. Introduction The ability to effectively design and manufacture plastics for engineering applications requires a large body of knowledge on long-term performance of the materials. The fatigue life is a typical long-term property which should be characterized in components and structures where cyclic loads are experienced. There has been considerable research into the basic mechanism that cause fatigue of plastics, but still there is not a clear understanding of the mechanism in many plastics. In this study, the mechanism of the fatigue fracture of glass fiber reinforced nylon66 has been identified using acoustic emission (AE), ultrasonic measurement and optical microscopy (OM) techniques. Experimental Test specimens were molded of 33 wt% glass fiber reinforced Nylon 66 (Leona 1402G grade, Asahi Chemical Industry Co., Ltd.) in the shape of the ASTM-D1822 dumbbell using a injection molding machine. The specimens were vacuum dried at 90°C for approximately 24h prior to tests. The fiber orientation in the middle plane of the molded specimens were observed using by SEM. The tensile fatigue tests were using hydraulic controlled servo machine. The fatigue test conditions included a load frequency of 20 Hz as a common condition for all specimens and a maximum stress of smax = 88.2 MPa which was corresponded to 58% of static tensile strength. The minimum to maximum load ratio R is 0.06. The fatigue tests at an ambient temperature is 300 K. AE activities were also measured during fatigue tests. The gain and the threshold value were set at 60 dB and 0.330 V, respectively. The acoustic velocity of the center of the specimens as shown in Fig.1 was measured using ultrasonic measurement techniques. Silicon oil were used as medium of acoustic velocity measurements. The thin section samples, approximately 30 mm thick, for OM observation was page_99 Page 100 prepared using a thin sectioning system (PETRO-THIN, Buehler) and polishing machine (Labopol-5, Struers). Acoustic velocity measurements and OM observation were examined for the damaged specimens during the fatigue tests and the failed specimens.

Fig.1 Position of ultrasonic measurement Results The SEM micrograph of the middle plane of the specimen is shown in Fig. 2. The fibers oriented to the parallel direction to the flow direction in the center of the specimen. However, the fibers oriented to the thickness direction of the specimen in the edge of the specimen. Fig.3 shows the compliance and the amount of AE events as a function of number of fatigue cycles. The compliance decreased at the first stage of the cyclic loading and gradually increased to the final stage. The amount of AE events suddenly increased at the first stage and gradually increased for most of fatigue cycles. In final stage, a rapid increase of the amount of AE events was measured. Fig.4 shows the acoustic velocity in each measurement position. In non-damaged specimen, the acoustic velocity of the edge of the specimen was faster than that of the center of the specimen. This tendency was associated with the difference of fiber orientation between the center and the edge of the specimen as shown in Fig. 2. In the damaged

and the failed specimens, the difference of the acoustic velocity between the center and the edge of the specimens became smaller, compared to the non-damaged specimen. Fig.5 shows OM micrographs of the damaged and the failed specimens. In both the damaged and the failed specimens, matrix cracks were observed in the flow orientation region at the edge of the specimen, and fiber breakages were found near the flow orientation region. These might be attributed to the increase of the AE events in the final stage of the cyclic loading and the decrease of the acoustic velocity in the edge of the damaged and the failed specimens. Conclusions In tensile fatigue tests of the glass fiber reinforced nylon 66, the fatigue fracture involving the matrix cracks and the fiber breakages occurred at the edge of the specimen in the final stage of the cyclic loading. The fatigue fracture process in the final stage could be monitored using the ultrasonic measurement. page_100 Page 101

Fig.2 The scanning electron micrograph of the finishing surface of Nylon66/SGF composite.

Fig.3 Variations of storage compliance (J') and amount of AE event with cycle for the Nylon66/SGF composite.

Fig.4 Relationships between acoustic velocity and position. page_101 Page 102

Fig.5 The optical micrograph of Nylon66/SGF composite. page_102

Page 103

MANUFACTURING II page_103 Page 105

Thermal and Electron Beam Curing of Polymer Composites-A Comparison J. Raghavan1,#, Vince J. Lopata2, and Mark R. Baillie1 1Department of Mechanical and Industrial Engineering University of Manitoba, Winnipeg, MB R3T 5V6, Canada 2AECL, Whiteshell Laboratories, Pinawa, MB R0E1L0, Canada Currently, electron beam curing of polymer composites is being pursued as an alternative to the dominant thermal curing. While most of the current research studies in North America are devoted to developing E-beam curable resins, further detailed research is needed to understand the E-beam curing process and its influence on rheological and mechanical properties of the polymer composites. Hence, this study is focussed on studying the E-beam curing process, its kinetics, its influence on rheological and mechanical properties of a polymer composite and its polymer matrix, and on comparing it with conventional thermal curing to ascertain its technical advantages. Results of this comparative analysis will be presented and discussed in this paper. #The presenting and corresponding author. Tel. 204-474-7430; Fax. 204-275-7507; E-mail : [email protected] This abstract is being submitted for presentation during the 2nd Canada-Japan Workshop on Composites to be held at Concordia University in August 1998. page_105 Page 106

An Investigation of Autoclave Convective Heat Transfer Andrew Johnston*, Pascal Hubert, Reza Vaziri, and Anoush Poursartip** Composites Group, Department of Metals and Materials Engineering, The University of British Columbia, Vancouver, B.C., V6T 1Z4, CANADA Keywords: processing, autoclave processing, heat transfer, modelling Abstract Among the most important of the parameters affecting the quality of composite components made using autoclave processing are the thermal ''boundary conditions" to which the parts are subjected during processing. At the temperatures typically encountered in the processing of thermoset matrix composites, the dominant mode of heat transfer between autoclave gases and composite parts and their tooling is generally forced convection. This paper presents an investigation of convective heat transfer rates in three "typical" composites processing autoclaves over a range of temperatures and pressures. This investigation revealed significant differences in the heat transfer characteristics of the autoclaves tested. It was also shown that heat transfer rates are greatly influenced by autoclave pressure and temperature, an effect that could be predicted by a turbulent heat transfer model. Process model simulations using the developed turbulent heat transfer model were used to explore some implications of the differences in autoclave heat transfer characteristics and to suggest some possibilities for taking advantages of the observed pressure effects to reduce process cycle time. Introduction Autoclave processing is one of the most common techniques used for manufacturing advanced composite structures. During this process, composite parts are subjected to a temperature/pressure cycle known as the "process cycle" in order to cure the matrix resin, achieve optimum fibre and resin distribution and to minimize the occurrence of voids. The quality of produced parts is highly dependant on this process cycle; thus it is critical that all aspects of the

process cycle be understood, optimized, and controlled. In recent decades, a great deal of effort has been expended to develop an improved understanding of the behaviour of composite materials during autoclave processing and in creation of models to predict this behaviour (e.g., [17]). Most work done to date has focused on analysis of the composite part itself, with much less emphasis given to analysis of part heat transfer (and other) * Now at National Research Council of Canada, Institute for Aerospace Research, Montreal Road Bldg. M-3, Ottawa, CANADA, K1A 0R6 ** Author to whom correspondence should be addressed. page_106 Page 107 boundary conditions during processing. However, in order to make this work more relevant to industrial practice, a better understanding of these often complex boundary conditions is needed. In recognition of this fact, numerous recent modelling efforts have been undertaken in which tooling thermal effects are considered [812]. Even more complex analysis of thermal boundary conditions by directly modelling autoclave internal gas flow and heat transfer has also been proposed [11], and simple preliminary numerical simulations of these phenomena have been performed [9,10]. At the temperatures typically used in processing thermoset composite materials, the dominant mode of heat transfer between parts and tooling and the surrounding autoclave gases is generally forced convection heat transfer (although as discussed in [13], at the higher temperatures involved in thermoplastics processing, radiation is often an important heat transfer mode). The rate of convection heat transfer (q) is described by the equation:

where A is the surface area, TA is the "bulk" gas temperature, TS is the surface temperature, and h is the "heat transfer coefficient". While empirical correlations exist for calculation of heat transfer coefficients for simple cases (e.g., laminar flow over a flat plate), h is not easy to predict a priori for this case. In an autoclave filled with parts and tooling of varying size and shape, very complex gas flow patterns can develop [14, 15], resulting in variations of as much as 6 8 °C in gas temperature [15] and a factor of 3 in heat transfer rates [14]. Given this complexity, it is perhaps not surprising that quantitative analysis of autoclave heat transfer has been limited. The strategy of composites processors is primarily to attempt to achieve uniform flow and minimize "dead air zones". Process modellers normally assume a uniform heat transfer coefficient which remains constant throughout the process cycle. However, given the importance of the heat transfer boundary conditions seen by a part on processing outcomes, some level of investigation is certainly warranted. Experimental Heat Transfer Coefficient Measurement Heat transfer coefficients were measured over a range of temperature and autoclave gas pressures in the three different autoclaves described in Table 1. In these tests, a simple aluminum plate calorimeter similar to that described in [14] was employed as shown in Fig. 1a. The top and bottom surfaces of the plate were exposed to the autoclave gas and its edges were thermally insulated using several layers of breather cloth. While the large plate thermal mass created a large enough plate-gas temperature difference to allow accurate heat transfer coefficient measurements, the very aluminum thermal conductivity also permitted use of a "lumped" thermal mass assumption (maximum Biot number during any test was about 0.04). Plate temperature was measured at three locations from the plate edge to the centre, midway through the plate thickness. Table 1: Autoclaves examined in heat transfer characterization tests. Autoclave Primary Usage Dimensions (internal) Autoclave A Production 1.8 m ´ 4.5 m Autoclave B Development 0.9 m ´ 1.8 m Autoclave C Development 1.5 m ´ 2.4 m To determine the heat transfer characteristics of the three chosen autoclaves, either one or two calorimeters (depending on autoclave size) were placed in the otherwise empty autoclaves, midway page_107 Page 108

between the autoclave walls, about 30 cm from the autoclave "floor". The calorimeters were then subjected to temperature/pressure cycles similar to that shown in Fig. 2 and autoclave gas temperature and pressure and plate temperatures were recorded. After each test, heat transfer coefficients were calculated using:

where r is the plate density, CP is its specific heat capacity, V is plate volume, A is the exposed plate area (top and bottom surfaces), T¥ is the measured autoclave gas temperature and TC is the plate temperature, taken as the average of the three thermocouple measurements. The temperature of the autoclave gas was measured at a single point between the inner and outer autoclave walls as illustrated in Fig. 1b.

Figure 1: a) Aluminum plate calorimeter; b) Schematic of heat transfer coefficient measurement test in autoclave A showing gas flow and relative plate location The temperature/pressure cycle used to characterize the heat transfer characteristics of autoclave A is illustrated in Fig. 2 along with the average temperatures of the two calorimeters during the test. As shown in this figure, the autoclave pressure was "cycled" in order to more easily isolate temperature and pressure effects on heat transfer rates. Fig. 3 shows the calculated heat transfer coefficients for plates 1 and 2 using Eq. 2. As can be seen in this figure, the heat transfer coefficient at the plate placed near the front of the autoclave (plate 1) was consistently higher than that at the back throughout the test, by about 25%. More significantly, the value of h calculated at both locations was found to change greatly during the test, varying by approximately a factor of four (about 40 160 W/m2K for plate 1). It is clearly apparent from this figure that heat transfer coefficients are highly influenced by pressure, with a lesser but still significant temperature effect. page_108 Page 109

Figure 2: Autoclave A heat transfer coefficient measurement test.

Figure 3: Calculated heat transfer coefficient during autoclave A characterization test. A similar procedure was followed in characterizing the other two autoclaves examined, although for the small autoclave B only a single plate calorimeter was employed. The observed effect on heat transfer coefficient of pressure and temperature changes for both of these autoclaves was similar to that found in autoclave A. For autoclave B, the magnitude of the measured heat transfer coefficients were also about the same as autoclave A, at about 30 W/m2K at room temperature and pressure. Under these same conditions, however, the measured value of h in autoclave C was much lower at about 13 W/m2K. This is believed to be primarily the result of much lower gas flow rates although this was not confirmed through measurements. page_109 Page 110 Analysis and Heat Transfer Rate Model Development Given previous work by researchers such as Ghariban et al. [14], neither the observed internal variations in heat transfer rates nor the differences in heat transfer behaviour observed in different autoclaves is especially surprising. The large effect of pressure (and the lesser effect of temperature) on the heat transfer coefficient, however, is another matter. Although this effect is not entirely unknown in industry [15], it is apparently not generally appreciated by the process modelling community. As mentioned in [15], the source of the observed variations in heat transfer coefficient with changes in pressure and temperature is the attendant change in the density of the autoclave gas. For an ideal gas, density, r, is directly proportional to pressure and inversely proportional to temperature, i.e.,

For a given geometry, the Reynolds number describing the flow of a fluid is directly proportional to fluid density and the "bulk" fluid velocity, V¥, and inversely proportional to the dynamic viscosity, m, i.e.,

As discussed in [16], for fully-developed turbulent flows, the Nusselt number (Nu), is related to the Reynolds number by:

where k is the fluid thermal conductivity. Combining Eqs. 4 6, we obtain:

For air (and nitrogen), the value of is roughly constant over the range of conditions of interest (with about a 10% variation from 20 °C to 180 °C). Thus, assuming constant bulk gas flow velocity, Eq. 7 reduces to:

where C is an empirically-determined constant. For each of the three autoclaves, a "least squares" analysis was used to determine the best fit between Eq. 9 and the measured heat transfer coefficients. A comparison between measured values and turbulent model predictions is shown in Fig. 4. A very good fit was obtained between experiment and predicted trends for autoclave A and autoclave B, with a much poorer fit obtained for autoclave C measurements. The source of the poor fit in the last case is unclear, but is likely due to the fact that a fully-developed turbulent flow did not in fact develop in this case as evidenced by the very low heat transfer coefficients measured. Radiation heat transfer also plays a much more page_110 Page 111 significant role in the case of autoclave C than for the other two autoclaves; a factor which is not included in the model.

Figure 4: Comparison between measured heat transfer coefficients and turbulent heat transfer model. Heat Transfer Rate Model Application Some implications of the measured differences in autoclave heat transfer rates and of the observed effects of temperature and pressure were explored using the autoclave processing model COMPRO [17]. A simple case was examined consisting of a 2.54 cm thick laminate of a typical first-generation CFRP/Epoxy, processed on a 0.64 cm thick invar tool. Convective heat transfer was assumed on the top and bottom surfaces, with the part and tooling edges insulated, resulting in a one-dimensional heat transfer case. Simulation of a conventional "lead/lag" type autoclave control system is also incorporated. As shown in Fig. 5, processing a part in two different autoclaves with very different heat transfer characteristics (in this case autoclave A and autoclave C) can result in a very large difference in the actual process cycle experienced by the part. In this case, the primary difference in the process cycle is the change in its length, about 160 minutes in this case. Fig. 6 illustrates the effect of using a slightly modified process cycle in which we take advantage of the measured effect of pressure on heat transfer rates. By increasing the autoclave pressure early in the process cycle, heat can be more quickly transferred from the autoclave gas into the tool and part. Pressure can then be quickly dropped as the first temperature "hold" to prevent excess resin flow. This effect of using this modified pressure cycle is a reduction in overall process cycle time by approximately 30 minutes for the modelled case, an effect which would be even greater for thicker tools and autoclaves with poorer heat transfer. page_111 Page 112

Figure 5: Effect of autoclave employed on actual process cycle (autoclave A vs. autoclave C).

Figure 6: Comparison of process cycles using conventional and modified pressure cycle illustrating potential cycle time reduction (autoclave A). Discussion and Conclusions This paper describes a simple test for measurement of autoclave heat transfer coefficients employing an aluminum plate calorimeter. Possible applications of this type of test include assessment and improvement of autoclave heat transfer by composites processors (e.g., by improved part placement), and measurement of existing heat transfer rates for inputs into process models. page_112 Page 113 Using the outlined test, heat transfer coefficients were measured in three different autoclaves over a range of temperatures and pressures. It was observed that heat transfer rates varied significantly with pressure with a lesser, but still important, temperature effect. For two of the three autoclaves, a turbulent heat transfer model was found to give very good predictions of observed pressure and temperature effects. It was also found that heat transfer rates varied significantly between autoclaves, with close to a factor of 4 difference from best to worst. Process model simulations using the developed turbulent heat transfer model were used to illustrate some implications of these experimental observations, specifically the potentially large difference in process cycle from one autoclave to the next, and the possibility of taking advantage of observed pressure effects to reduce process cycle time. Acknowledgements The authors would like to acknowledge the financial support of the Science Council of British Columbia, the Natural Sciences and Engineering Research Council of Canada, The Boeing Company, and Integrated Technologies Inc. We would also like to gratefully acknowledge the significant interaction and support of Karl Nelson of The Boeing Company. References

1. A.C. Loos and G.S. Springer, J. of Comp. Mater. 17(2) (1983), pp. 135169. 2. R. Davé, J.L. Kardos and M.P. Dudukovic, Proc. Amer. Soc. Comp., 1st Tech. Conf. (1986), pp. 137153. 3. R. Mallow, F.R. Muncaster and F.C. Campbell, Proc. Amer. Soc. Comp., 3rd Tech. Conf. (1988), pp. 171186. 4. J. Mijovic and J. Wijaya, SAMPE J. 25(2) (1989), pp.3539. 5. T.A. Bogetti and J.W. Gillespie Jr., J. of Comp. Mater. 25(3) (1991), pp. 239273. 6. M. Kenny, Proc. Third Conf. Comp. Aided Des. Comp. Mater. Tech. (1992), pp. 530544. 7. S.R. White and H.T. Hahn, J. of Comp. Mater. 26(16) (1992), pp. 24022422. 8. P.R. Ciriscioli, Q. Wang, and G.S. Springer, J. Comp. Mater. 26(1) (1992), pp. 90102. 9. M.K. Telikicherla, X. Li, M.C. Altan, and F.C. Lai, Proc. 1994 Int. Mech. Engg. Cong. Exp., ASME HTD 289 (1994), pp. 213221. 10. M.K. Telikicherla, M.C. Altan, and F.C. Lai, Int. Comm. Heat Mass Trans. 21(6) (1994), pp. 785797. 11. R.A. Kline and M.C. Altan, United States Patent 5,453,226, Sept. 26, 1995. 12. V. Pillai, A.N. Beris, and P. Dhurjati, Comp. Chem. Engg. 20(3) (1996), pp. 275294. 13. P.F. Monaghan, M.T. Brogan, and P.H. Oosthuizen, Comp. Mnfctg., 2(3/4) (1991), pp. 233242. 14. N. Ghariban, D.Y.S. Lou, and A. Haji-Singh, Heat Trans Eff. Mat. Procg. ASME HTD 233 (1992), pp. 4552. 15. R.W. Roberts, Engineered Materials Handbook, Vol. 1: Composites, (1987), pp. 745760. 16. F.P. Incropera and D.P. DeWitt, Introduction to Heat Transfer, Second Edition (1990). 17. A. Johnston, P. Hubert, G. Fernlund, R. Vaziri and A. Poursartip, Sci. Engg. Comp. Mater. 5(34) (1996), pp. 235252. page_113 Page 114

On the Processing and Testing of FRP Composites Incorporating Fiber Optic Sensors Alexander L. Kalamkarov, Stephen B. Fitzgerald, Douglas O. MacDonald and Anastasis V. Georgiades Department of Mechanical Engineering, Dalhousie University P.O. Box 1000, Halifax, Nova Scotia, B3J 2X4 Canada Keywords: Smart composite reinforcements, fiber optic sensors, pultrusion, process monitoring, residual strains, strain monitoring. Abstract The use of the pultrusion process for the manufacture of fiber reinforced polymer (FRP) composites with embedded fiber optic sensors is discussed. The specific application is the use of smart composite reinforcements for strain monitoring in engineering structures. The Bragg Grating and Fabry Perot fiber optic sensors are embedded during the pultrusion of FRP rods and the process induced residual strains are evaluated using these sensors. The behaviour of optic sensors during pultrusion is assessed, and the effect of the embeddment of optical fibers and their surface coatings on the mechanical properties of the composite is investigated. Monitoring of the output of embedded fiber optic strain sensors during the pultrusion of composite rods gives a unique view of the formation of residual strains within the pultrusion die itself. To verify the operation of the optic sensors embedded in the smart pultruded tendons, mechanical tests were conducted and the output of the fiber optic sensors was compared to that of an extensometer during quasi-static and cyclic tensile tests. Introduction

Composite materials lend themselves as prime candidates for the rapidly expanding field of smart materials [1]. Composites are good candidates for making smart materials because their fabrication techniques inherently allow for the embeddment of sensors, actuators and communication lines. However, in the quest to advance smart composite materials applications it is desirable to reduce their cost through the use of automated production techniques. To date, hand layup in combination with vacuum bagging and autoclaving has been most often used to fabricate smart composites. This can be a labor intensive and time consuming process in which the quality of the final product is significantly affected by the skill and experience of the technician. Pultrusion which is the only continuous process, has received little attention in the area of smart composites, and there are currently only a few publications on this subject [2, 3]. However, by considering the costs and cycle times which are favorable criteria for selection of a manufacturing process, pultrusion is the fastest and most cost effective process. Pultrusion is also well suited to produce prestressing tendons and rebars, because it can provide the structures with a high degree of axial reinforcement. The pultrusion process however, inherently has the potential to generate residual stresses within a composite component for several reasons. The high output or production rate in meters per minute requires a fast cure rate as the raw materials travel through a pultrusion die which is typically just less than one meter in length. The resin matrix is thus subject to a dynamic cure profile created by strip heaters attached to the die surface. Accelerators and promoters are needed to cure the resin in addition to the normal catalysts. All considered, not much is known about the effect of these factors on the development of residual stresses. One must also consider that the infeeding page_114 Page 115 of reinforcing fibers to the pultrusion die is also a dynamic process and problems associated with the balance and symmetry of the fiber distribution may occur. Once again, this effect may generate residual stress within the component. It is therefore useful to investigate the ability of embedded fiber optic sensors to monitor the strains during the processing and to measure the residual strains created by the pultrusion process. In addition to the effects that the embeddment of the fiber optic sensors in the process of pultrusion might have on the mechanical properties of the composite itself, it is necessary to assess the performance of the sensors themselves under conditions of static and dynamic loading. In particular, it would be important to obtain data that would reflect on the repeatability and accuracy of the fiber optic strain measurements and also to compare the output of the pultruded sensors to that from conventional strain gauges such as external extensometers. The objectives of the research reported herein are the following: to evaluate the residual strains induced during the pultrusion of FRP rods, to assess the behaviour of Fabry Perot and Bragg Grating sensors during pultrusion, to determine how the embeddment of optical fibers and their surface coatings affect the mechanical properties of the composite, to confirm the operation of the embedded sensors in the smart pultruded tendons, and to compare the output of the embedded sensors to that from an extensometer. Fiber Optic Sensors In this study, two types of fiber optic sensors were embedded during the pultrusion of carbon and glass fiber reinforced rods: Fabry Perot and Bragg Grating sensors. The Fabry Perot sensor has been developed to use a broadband light source as opposed to laser light. It is highly sensitive and can make precise, linear, and absolute measurements [2]. The other type of fiber optic sensors used was of the Bragg Grating type. Bragg Grating sensors are based on creating a pattern of refractive index differentials directly onto the material of the fiber core. Fiber gratings selectively reflect certain wavelengths and transmit others. Which wavelengths are transmitted and which ones are reflected depend on both the refractive index of the core material as well as the spacing of the pattern. Changes in temperature or pressure will change the refractive index of the core material and hence cause a change in the wavelengths of peak reflection (or transmission). The presence of mechanical strain along the length of the fiber will have a similar effect since it will change the grating spacing. Measurements of these wavelength shifts provide the basis of operation of Bragg Grating sensors. One advantage of Bragg Grating sensors is that the shift in the wavelength of peak reflection and/or transmission is linear with temperature and axial strain. On the other hand, it is not possible to decouple the effects of temperature and strain with just one sensor. In addition, Bragg Grating sensors, unlike Fabry Perot sensors are quite sensitive to transverse strain because of the photoelastic effect. There are several techniques available to determine the wavelength shift, including optical spectrum analyzers and tunable filters. These sensors have a great potential in smart composites and structures. Experimental Materials and Equipment Pultruded carbon and glass FRP rods were produced using a urethane modified bisphenol-A based vinyl ester resin system known for its good mechanical properties and excellent processability. Two types of organic peroxide catalysts were used to cure the resin, di-peroxydicarbonate and tert-butyl peroxybenzoate. Adequate release from the die was achieved using an internal lubricant. The mechanical properties of the resin system and of the carbon fiber rovings are given in [3]. The 9.5 mm diameter carbon rods were pultruded with 22 ends of rovings giving a volume fraction of 62%, while their glass counterparts were pultruded with 26 ends giving a volume fraction of 64%.

page_115 Page 116 The Fabry Perot and Bragg Grating sensors were acquired as prepackaged assemblies. The sensing element was located at the front end of an optical fiber of about 5 m total length. The fiber optic core and cladding were protected by a polyimide coating which acts as the contact surface between the optic fiber and the surrounding composite. Polyimide coating was used to ensue survival of the optical fiber when exposed to the high temperatures in the pultrusion die [3]. In these experiments the actual maximum die temperature was 150°C, whereas the polyimide provides protection up to 350°C. The Fabry Perot sensors used in the experiments were rated for ±5000 or 0-10000 microstrain, while the Bragg Grating ones were rated for ±5000 microstrain. The sensors were not temperature compensated. Pultrusion was carried out on an experimental pultrusion line. To determine how the embeddment of optical fibers and their surface coatings affect the mechanical properties of composite materials a microstructural analysis was carried out on both the pultruded profile's cross section and on fracture surfaces obtained from mechanically fractured samples. Finally, to assess the overall behaviour of the embedded fiber optic sensors, smart tendons were produced with the pultrusion process described above. The samples were tested using an Instron servohydraulic load frame and an appropriate controller. Strain was measured externally on the samples using an extensometer. During the actual tests, four analog signals were read into a data acquisition system, one from the extensometer, one from the load cell, one from an LVDT, and a final strain signal from the Bragg Grating or Fabry Perot sensors. Results and Discussion The microstructural analysis showed that the polyimide coating on optical fiber results in a good interface between optical fiber and host material; whereas acrylate coating cannot withstand the harsh environment (high production temperature up to 150°C) and causes severe debonding of optical fiber and resin. Mechanical tests performed on glass and carbon FRP tendons indicated that the tensile strength and tensile modulus of the tendons were virtually unaffected by the embeddment of a single polyimide coated optical fiber [3]. The suggested explanation is that the fiber reinforcement is the only significant factor directly affecting the tensile properties of a unidirectional composite [4]. The embeddment of a single optical fiber slightly deteriorates the shear strength of glass rods, but no effect is evident in the case of carbon rods. As a first attempt at embedding a fiber optic sensor in a pultruded rod an unmodified Fabry Perot fiber optic sensor was added to the fiber feed side of the pultrusion process. The forward end of the sensor lead was bonded with a 5 minute epoxy to one of the carbon fiber rovings to ensure that it would feed into the die. From the location at which it was bonded, the sensor had to pass through two of the fiber feed cards before entering the die. The sensor was also located towards the outer surface of the carbon fiber rod. After the sensor had passed through the die and had been embedded in the composite rod, the pultrusion process was stopped to enable trimming away of several carbon fiber rovings in order to pass the pigtail and connector through the die. The result of the first trial was a length of carbon fiber rod with an embedded Fabry Perot sensor. The end of the rod which was pulled dry during the shut down process contained the pigtail with connector. When the sensor was tested using the fiber optic readout unit, the readings tended to jump about from low (as expected) microstrains, to very high readings of strain far in excess of the 5000 microstrain limit. Two explanations of the cause of fluctuations in the microstrain readings were suggested. One was that the harsh conditions of temperature, fiber compaction, or resin cure shrinkage in the pultrusion die damaged the sensor. The second explanation was simply that the sensor was handled too roughly before or after processing, or that it may have been damaged by contact with the fiber feed cards or entrance into the die. page_116 Page 117

Fig. 1 Comparison of outputs from not pre-reinforced and pre-reinforced Fabry Perot sensors during normal pultrusion. It was decided to conduct a series of experiments that would expose a Fabry Perot sensor separately to each of the variables in the pultrusion process. These are fiber compaction pressure, elevated temperature, liquid resin, and resin curing shrinkage stresses. Changes were also made to the fiber feed system to allow the sensor to be located more accurately in the center of the rod, and to protect the sensor from damage in the fiber feed system. On account of the low survivability rates of Fabry Perot sensors in the pultrusion process, a novel method was developed to pre-reinforce the sensors before being pultruded in order to offer more protection at the die entrance. The strains observed by this sensor were generally much lower than those of previous trials, see Fig. 1. Unlike Fabry Perot sensors, Bragg Grating sensors showed enhanced survivability during pultrusion and hence it was not considered necessary in this case to perform the various experiments described above whereby the sensors were exposed separately to each of the variables in the pultrusion process. It was not also necessary to pre-reinforce Bragg Grating sensors. Nevertheless, we did subject Bragg Grating sensors to a ''dry" pultrusion run (passing of the sensors and glass fibers through a heated die but with the fiber rovings not soaked in resin), primarily in order to compare the strain output with that from Fabry Perot sensors. In all experiments performed and described in this paper, the sensor was located at the center of the rod using a single fiber roving that traveled a straight line through the feed system. The observed strain readings followed the temperature variation in the die quite closely. These strains which reach a peak of around 1450 microstrain are most likely due to thermal expansion of the sensor. Note that the peak strain is well below the 5000 microstrain capability of the sensor, see Fig. 2. Comparing the strain outputs for dry pultrusion runs from both Bragg Grating and Fabry Perot sensors, it is observed that the plots have a very similar shape as expected since they both closely follow the temperature profile in the die. The strains recorded by the Bragg Grating sensor were about 2.5 times larger than the corresponding strains from the Fabry Perot sensor, and this may be attributed to the difference in the coefficients of thermal expansion of the two sensor types. page_117 Page 118

Fig. 2 Comparison of output from Bragg Grating sensor during normal and dry pultrusion. Subsequently to the dry runs, a number of normal pultrusion experiments with different Bragg Grating sensors were performed, just as it was previously done with Fabry Perot sensors. The overall profile of the plot is very similar to that for the dry run. Fig. 2 shows the strain plots from the "dry" and normal pultrusion runs superimposed. The difference between the two curves is due to the curing of the resin. For example, the peak strain during normal pultrusion is much higher than the equivalent strain during "dry" pultrusion. This is likely to have been caused by an increased thermal expansion of the sensor due to the exothermic reaction accompanying pultrusion. Also to be noted is the fact that as the product exits the die, the difference between the two curves in Fig. 2 represents the residual strains induced by the pultrusion process. As well, Fig. 2 also shows that during normal pultrusion the strains do not go back to zero as they do in "dry' pultrusion, because of the residual strains which are "locked in". To assess and characterize the overall behaviour of the embedded fiber optic sensors, mechanical testing of the pultruded tendons was carried out by applying various loads to the tendons while continuously monitoring strain via the embedded optical sensors and a standard extensometer clipped to the pultruded rod. The smart FRP tendons were subjected to two basic waveforms in order to evaluate their performance and meet the objectives of the research. The first waveform was a trapezoidal waveform whereby the load was ramped from a low value (typically 100 N) to a peak value of about 3000 to 4500 N, at a slow rate of 90 N per sec. The tendons were held at this load for 20 sec and then ramped back down to the initial load at the same rate. The second waveform to which the smart tendons were subjected was a sinusoidal one. The frequency was one cycle per minute (0.0167 Hertz), and a typical range through which the load was cycled was 400 to 5000 N. The glass FRP tendon containing an embedded Bragg Grating was first subjected to a trapezoidal page_118 Page 119

Fig. 3 Strain from the embedded Bragg Grating sensor and the extensometer in a glass FRP tendon subjected to a trapezoidal load. waveform as described above. A graph which illustrates the result of this test is shown in Fig. 3. A carbon FRP tendon containing an embedded Bragg Grating sensor was also subjected to the sinusoidal load. The pertinent data indicates that the Bragg Grating strain output was fairly consistent over the load cycles, consistently reaching approximately 800 microstrain at each peak loading. There was a good agreement between the two measuring devices, see Figs. 4 and 5. The smart FRP tendon containing an embedded Fabry Perot sensor was tested in a similar manner. The sensor was embedded in a carbon fiber tendon and subjected to a sinusoidal tensile waveform, see Fig. 6. Conclusions Bragg Grating and Fabry Perot fiber optic sensors have been successfully embedded into FRP composite parts during pultrusion. It was shown that it was necessary to pre-reinforce Fabry Perot sensors prior to pultrusion. The Bragg Grating sensors show a greater survivability in the pultrusion process than the Fabry Perot sensors and there was no need to pre-reinforce them. Dry pultrusion runs were performed with both sensors and the thermal strain output obtained conformed quite well with the temperature profile within the die. Fabry Perot and Bragg Grating sensors were then incorporated in normal pultrusion runs. The different sensors showed strain outputs combining thermal and residual strains, which had similar basic profile even though the absolute values of the strains varied from sensor to sensor. The information provided by these experiments yields valuable insight to the specifics of the

pultrusion process. Pertinent microscopic analysis indicated that polyimide coating on optical fibers results in a good interface between the optical fiber and the host material. On the other hand, acrylate coating cannot withstand the harsh environment (high production temperature) characterizing the pultrusion process, and causes severe debonding of optical fiber and resin. It was also determined that embedded optical fibers have no significant effect on the tensile properties of the pultruded FRP, but they slightly deteriorate the shear strength of the composites. page_119 Page 120

Fig. 4 Strain from extensometer and embedded Bragg Grating sensor in carbon FRP tendon subjected to a sinusoidal load.

Fig. 5 Strain vs. time plot from extensometer and embedded Bragg Grating sensor in carbon FRP tendon subjected to a sinusoidal load. page_120 Page 121

Fig. 6 Strain from extensometer and embedded Fabry Perot sensor in a carbon FRP tendon subjected to a sinusoidal load. Mechanical testing was carried out in order to assess the overall behaviour of the smart tendons with the embedded Fabry Perot and Bragg Grating sensors. There was a good agreement between the two measuring devices. Acknowledgment This work was supported by ISIS - CANADA, the Intelligent Sensing for Innovative Structures Canadian Network of Centres of Excellence through the Project T3.4 on Smart Reinforcements and Connectors. References 1. A.L. Kalamkarov and A.G.Kolpakov, Analysis, design and optimization of composite structures, Wiley: Chichester, New-York, 1997. 2. A.L. Kalamkarov, S.Fitzgerald, and D.MacDonald, On the processing and evaluation of smart composite reinforcement. Proceedings of the SPIE, 3241 (1997), p.338. 3. A.L. Kalamkarov, D.MacDonald and P.Westhaver, On pultrusion of smart FRP composites. Proceedings of the SPIE, 3042, (1997), p. 400. 4. A.L. Kalamkarov, Composite and reinforced elements of construction, Wiley: Chichester, New-York, 1992. page_121 Page 123

TEXTILE COMPOSITES page_123 Page 125

Predicting Shrinkage in Polyester Reinforced by Glass Fabrics by V. Do-Thanh & T. Vu-Khanh Université de Sherbrooke Faculté des Sciences appliquées 2500 boul, de l'Université Sherbrooke, Québec, Canada, J1K 2R1 Abstract

Polyester is one of the most common resin used in contact lay-up method because of its low cost, room-temperature curing, wide availability, and ease of handing. However, the main disadvantage of this resin is the large volumetric shrinkage after curing (up to 5%7%). This is a major problem because it causes the unexpected deformations of the composite such as warpage, distortion and rippled surface. The effect of resin shrinkage on composite deformations is very complex because of the anisotropic properties of the fibers, especially in woven fabric composites with interlacing yarns. Moreover, in many applications, when the part geometry has a double curvature, the forming process usually results in significant in-plane shear deformation of the interlaced yarns. The angle between the fill and the warp threads is no longer orthogonal because the fabric must follow the shape of the mold. Consequently, the shrinkage of composite parts is very complex and it is essential to reduce the cost of manufacture in predicting the shrinkage instead of an trial and error approach. Introduction In recent years, applications of reinforced polyester has increased constantly in various fields such as automobile, construction, marine, sanitary equipment etc. Polyester reinforced by glass fiber offers major advantages including low cost tooling, ease of fabrication, wide range of available colors, light weigh, high strength, high choke resistance, simple curing cycle, etc. However, several problems related to the quality of parts made of polyester reinforced by glass fiber such as warpage, surface appearance, crack initiation etc., are frequently encountered in practice. This is caused by shrinkage during polymerization of the resin. Much work has been carried to reduce shrinkage of polyester resin. The reported investigations showed that the addition of low-profile additives such as PVAc, PMMA, PU and PS, to unsaturated polyester can lead to a shrinkage compensation, resulting from the phase separation [15]. Some studies have also been carried out to predict the shrinkage and warpage of polyester film and fiber reinforced thermoset composite [69]. However, these reported works are limited to unidirectional or short fiber composites. With the composites made of woven fabrics, the mechanism of resin flow and fiber deformation are more complicated, especially in parts containing double curvature surface [13]. The problem of shrinkage of woven fabric composite can be analyzed in a similar manner as that for thermal expansion in woven fabric composite [15]. Several models have been proposed to analyze the thermo-mechanical behavior of fabric composite such as the mosaic model, the crimp model and the bridging model. However, the major limitation of these is that they cannot be applied to the case of non orthogonal fabric structures, encountered in parts with double curvature. The recently proposed sub-plies model was therefore employed in this work to analyze the shrinkage behavior of deformed woven fabric composites. The purpose of this work is twofold : a) to develop a method to characterize the shrinkage properties in woven fabric composites; and b) to apply the sub-plies model in predicting shrinkage of deformed woven fabric composites. page_125 Page 126 Sub-plies Model In the sub-plies model, the fabric is considered as a laminate consisting of four fictional unwoven unidirectional plies. The lay-up structure of an orthogonal sub-plies model can be considered generally as [(q/2)h1/-(q/2)h2/(q/2)h2 /-(q/2)h1] (Fig.1).

Fig. 1: The sub-plies model. It has been shown in [1011] that the fictional thickness e can be expressed as a function of the fabric thickness t0 and ng. Where ng is a geometrical parameter which presents a wrap thread interlacing with every ng-th fill threads. The stiffness coefficients of this laminate and can be obtained by measuring the stiffness coefficients in 00 and 450 directions of the molded samples of the fabric composite. This form permits the direct use of shell element in any of finite element code to evaluate the shrinkage and residual distortions as a function of the in-plane shear deformation angle of the interlaced yarns. Characterization of Equivalent On-Axis Shrinkage Coefficients. In applying the classical theory of laminate plate in the sub-plies model [14], the in-plane strain e1 and e2 of an angle ply laminate [+q/-q]s can be written as function of the on-axis ply stiffness Qi (i=1..5) and the on-axis shrinkage coefficients gx and gy, in place of the thermal expansions axDT and ayDT [15].

Where

page_126 Page 127 In equation (2), Qxx[eq], Qxy[eq] and Qyy[eq] are the stiffness coefficients of the constituent fictional ply in Fig. 1. The subscript [eq] refers to the equivalent on-axis coefficients of the constituent fictional ply in the sub-plies model, composed of the undulated fibers (See [13] for the definition of these coefficients). The on-axis shrinkage coefficients gx and gy can be solved from Eq.1:

and

where

Experimental Verification Materials The samples were prepared from plain weave fabric WR 180Z and AK2100 unsaturated polyester resin, supply by Armkem Inc. Canada. All specimens were fabricated by hand lay-up method, with a fiber volume content of about 50 %.. The fiber volume fraction of the molded samples was always verified after molding (by burning off the resin and measuring its percentage) in order to issue a variation of the fiber volume fraction within ±1.0 %. Plaques of unidirectional, [0/90]s, and plain weave fabric were molded. The non-orthogonal plain-weave laminates were made by deforming the orthogonal interlaced yarns of the fabric by in-plane shearing to different angles before molding. Elastic Properties Tension tests were carried out according to the standard ASTM (D3039). All measurements were performed at room temperature (23°C) and 50% relative humidity. The equivalent on-axis ply stiffness coefficients Qxx[eq], Qxy[eq], Qyy[eq], and Qss[eq] were determined from the tension tests on samples cut out in 0° and 45° directions of the molded plaques of the plain weave composite. With the assumption of Ey[eq] = Ey[unidireccction] measured on unidirectional sample [18]. It was found that, at room temperature, Qxx[eq]= 37.826 GPa, Qxy[eq]= 2.896 GPa, Qyy[eq]= 7.418 Gpa and Qss[eq]= 5.105 Gpa. From these values, the equivalent unidirectional engineering properties can be determined [13,15,17] and are presented in Table 1. Table 1: Properties of equivalent unidirectional composite from plain weave and unidirectional. Ex (Gpa) Ey (Gpa) Es(Gpa) Vxy Equiv. Properties (Plain weave) 36.695 7.196 5.1049 0.39 Unidir. Composite 39.064 7.196 4.729 0.25

page_127 Page 128 To verify the accuracy of the above properties, the plain weave fabric was deformed to different angles before molding. The off-axis modulus of the deformed woven fabric composite is shown in Fig. 2 as a function of the angle, q, between the interlaced yarns. Predictions based on sub-plies model with the equivalent unidirectional properties derived from the in-plane stiffness Qxx[eq], Qxy[eq], and Qyy[eq] measured on the plain weave samples and on unidirectional samples [18] are also shown. The results suggest that there is a relatively good agreement between experimental data and the sub-plies model with properties measured by the proposed procedure [13,18].

Fig. 2 Variation of Young 's modulus as a function of q/2 for the plain weave composite. Shrinkage Coefficients To measure the shrinkage coefficient of constituent fictional ply, the plain weave fabric was deformed by in-plane shearing to an angle of 30°. The measurement was performed on the molded fabric samples at room temperature, after removing the specimens from the mold for 24 hours. These conditions permit the resin to complete the shrinkage process. The relative displacements e1 and e2 (see Fig.1) were measured by optical microscopy. The shrinkage coefficients in the fiber direction and the transverse direction gx[eq] and gy[eq] were evaluated using equations (3) and (4). The shrinkage measurement is rather difficult since the shrinkage of the woven fabric composite is small with respect to the scatter of experimental measurements. The results are shown in Table 2 that are average values determined from experimental measurements on 6 specimens. It is surprising to find that the value of gx[eq] is positive. This implies that the resin shrinkage leads to an expansion in the fiber direction of the constituent unidirectional ply of the sub-plies model. In fact, micromechanics models have long been developed to predict the hygrothermal properties of a lamina. The longitudinal and transverse thermal expansion coefficients can be expressed by [19]:

page_128 Page 129 Using the above equation, it can be seen that if shrinkage occurs only in the matrix, both the longitudinal and transverse shrinkage coefficients should be negative. The results shown in Table 2 are therefore quite surprising and cannot be explained by the increase in temperature of the curing process. Table 2 : On-axis shrinkage coefficient. e1 Sb e2 -3.7879´10-4 2.13698´10-5 12.197´10-4 Sb : Standard Deviation

Sb 8.4632´10-5

gx[eq] 4.4296´10-4

gy[eq] -26.512´10-4

As mentioned above, the deformations were only measured 24 hours after removing the sample from the mold so that its temperature is that of room temperature. From the laminate theory, it can be shown that a positive shrinkage coefficient gx[eq] should result in an expansion of angle-ply laminates for low values of j. Fig. 3 shows the deformation e1 (see Fig. 1) after curing of two plain weave composites with the angles between interlaced yarns of 60° and 70°. The results confirm that curing of the matrix results in an expansion of the laminates in the direction 1, validating therefore the positive value of gx[eq]. A possible explanation for the expansion in the fiber direction of the constituent ply is that matrix shrinkage could straighten the undulated fibers in woven fabric composites. For large values of q, it could be expected that e1 are negative since the coefficient gy[eq] becomes predominant. This effect is confirmed by Figure 4 for the values q of 90°, 105°, and 110°. The values of gx[eq] and gy[eq] in Table 2 were subsequently used in the sub-plies model to predict the deformation e1 of the plain weave fabric, deformed to different angles q (see Fig.1). The prediction was then compared with experimental measurements. The measured data are listed in Table 3. Fig. 5 presents the variation of the relative displacement e1 as a function of the yarns angle q. It is seen that prediction based on the sub-plies model is in good agreement with experimental measurements. For a comparison purpose, prediction using the theoretical shrinkage coefficients calculated from Equations (6), (7) for different percentages of shrinkage of the matrix is also shown. It can be seen that the calculated shrinkage coefficients using micromechanics models for aligned fibers without undulation result in a very strong discrepancy between prediction and experimental measurements of shrinkage of the fabric composite after curing.

Fig. 3 Deformation e1 with the angles between interlaced yarns of 60°, 70°. page_129 Page 130

Fig. 4 Deformation e1 with the angles between interlaced yarns of 90°, 105°, and 110°. Table 3: The shrinkage of deformed woven fabric laminates

Yarn 's angle No of sample Relative deformation on axis of woven fabric 1 (e1) q (°) Mean Std. [+35/-35] 6 2.614´10-4 5.056´10-5 [+45/-45] 6 -2.462´10-4 5.065´10-5 [+52.5/-52.5] 6 -5.076´10-4 7.829´10-5 [+55/-55] 6 -7.652´10-4 9.714´10-5 Std. : Standard deviation In order to further verify the accuracy of this approach for predicting shrinkage due to matrix curing in woven fabric composites, measurements and calculation were carried out on orthogonal fabric samples. Eight laminates of undeformed plain weave fabric were fabricated in order to measure the relative displacement e in the principal directions. The prediction was made by using the same shrinkage coefficients gx[eq] and gy[eq] in Table 2. The experimental results are compared to the predicted values in Table 4. Again, there is a good page_130 Page 131 agreement between experimental measurement of shrinkage in the plan weave composite and the proposed approach of prediction (with an error of only about 3%). Table 4: On-axis deformation due to matrix shrinkage in orthogonal plain weave composite Experimental measurement Prediction by sub-plies model e = 0.00151 e = 0.00147

Fig. 5: Variation of deformation due to shrinkage,e1, as a function of angle q/2. ( ·) : Test data; ( ) : Prediction based on sub-plies model using equivalent unidirectional properties derived from experimental measurement with sub-plies model; ( . . .),(-.),(--) : Prediction based on the micro mechanic model neglecting the effect of undulation correspond to different value of g. Conclusion In this work, an approach to measure the shrinkage coefficients of the interlaced yarns of fabric structure has been developed. The method consists of deforming the woven fabric by shear to before laminating the composite plaque. Experimental measurements of relative deformations due to shrinkage were carried out on two perpendicular directions. The proposed sub-plies model has been used to predict deformations du to resin shrinkage in woven fabric composite. It has been found that there can be an expansion in the fiber direction of a lamina with undulated fibers in polyester resin. This could be due to a straightening effect of resin shrinkage on the fibers. The expansion due to matrix shrinkage has been verified on several woven laminates. The prediction of deformations due to matrix shrinkage in woven fabric composites by the sub-plies model is in good agreement with experimental measurements. page_131

Page 132 References [1] M. Ruffier, G.Merle, J.P.Pascault, '' The Shrinkage Compensation Of Unsaturated Polyester Resins-Polyvinyl Acetate Blends Polumerization Proceeds Through Fractal Morphologies: Characterization And Simulation ", Journal of materials sciences 31, 1996, p.46794687. [2] Huang, Yan-Jyi, Liang, Chiou-Ming, " Volume Shrinkage Characteristics In The Cure Of Low-Shrink Unsaturated Polyester Resins", Polymer v37, 1996, p401412. [3] Saito, R.; Kan, W.-M.J.; Lee, L.J " Thickening Behaviour and Shrinkage Control of Low Profile Unsaturated Polyester Resins", Polymer v 37,1996, p 35673576. [4] Huang, Yan-Jyi, Liang, Chiou-Ming, " Effect of Low-Profile Additives on Volume Shrinkage Characteristics in The Cure of Unsaturated Polyester Resin.", Annual Technical Conference - ANTEC, Boston, MA, USA, 1995. [5] Piggott, M.R., Zhou, W. " Shrinkage Control In Fiber Reinforced Polymers III : Carbon Fibre Reinforced Polyesters With Expanding Monomers And Low Profile Additives", Polymer & Polymer Composites v3 n 6 1995, p.395402. [6] Ahang, L.Ernst, L.J, Brouwer, H.R., " Transverse Behavior of a Unidirectinal Composite ( Glass Fibre Reinforced Unsaturated Polyester). Part II. Influence of Shrinkage Strains", Mechanics of Materials v27 n 1 Jan 1998, p3761. [7] Shih, Wayne K., "Shrinkage Modeling of Polyester Shrink Film ", Polymer Engineering and Science v43 Jun 1994, p11211128. [8] Shi-Chang Tseng and Tim A. Osswald, " Prediction of Shrinkage and Warpage of Fiber Reinforced Thermoset Composite Parts", Journal of Reinforced Plastics And Composite v 13 Aug. 1994, p.698721. [9] S.F. Walsh,"Shrinkage and Warpage Prediction for Injection Molded Components", Journal of Reinforced Plastics And Composite v 12 July. 1993, p.769777. [10] Ishikawa, T. and Chou, T,-W., " Stiffness and Strength Behavious of Woven Fabric Composites", Journal of Material Science, Vol 17, 1982, pp.32113220. [11] Ishikawa, T. and Chou, T,-W., " Stiffness and Strength Properties of Woven Composites", Proceeding of the 4th International Conference on Composite Materials, Vol. 4, 1982, pp. 489496. [12] Ishikawa, T. and Chou, T,-W., " In-Plane Thermal Expansion and Thermal Bending Coefficients of Fabric Composites", Journal of Composite Materials, Vol 17, 1983, pp.92104. [13] Laroche, D. and Vu-Khanh, T., " Modeling of the Themo-Elastic Properties of Wowen Fabric Composites in Complex Shapes ", Composite Materials : Testing and Design (Eleventh Volume), ASTM STP 1206, E. T. Camponeschi, Jr., Ed,, American Society for Testing and Materials, 1993, pp.263275. [14] S.W.Tsai and H.T.Haln " Introduction to composite materials", Technomic Publishing Co., Lancaster. PA. USA(1980). [15] T. Vu -Khanh and B.Liu, " Prediction of Fibre Rearrangement and Thermal Expansion Behaviour of Deformed Woven-Fabric Laminates." Composite Science and Technology, Vol. 53, 1995, pp.183191. [16] T. Ishikawa, T.W.Chou, " Stiffness and Strength behaviour of Woven Fabric Composites", Journal of materials science, vol. 17, 1982, pp.32113220. [17] T. Vu -Khanh and B.Liu, " Characterization of thermo-elastic behaviour of woven fabric composites at the elevated temperatures.", Science and Engineering of Composite Materials, Vol.6,1997, pp. 5162. [18] H. Nguyen-Hoa and T. Vu-Khanh, " Prediction of Failure in Polyester Reinforced by Plain Weave Glass Fabric.", Université de Sherbrooke, Faculté de Science Appliquées, Département de genie mécanique. [19] Kuno K.U. Stellbrink, " Micromechanics of Composites ",Hanser/Gardner Publication, Inc., Cincinnati,1996. page_132 Page 133

Experimental and Numerical Analysis of Micro-fracture Behavior in Textile Composites Asami Nakai1), Hiroyuki Hamada2) and Nobuo Takeda3) 1) Graduate School of Interdisciplinary Engineering Science, The University of Tokyo, 4-6-1 Komaba, Meguro-ku, Tokyo 153-8904, Japan 2) Faculty of Textile Science, Kyoto Institute of Technology, Matsugasaki, Sakyo-ku, Kyoto 606, Japan 3) Center for Collaborative Research (CCR), The University of Tokyo Keywords: Braided Composites, Fiber Crossing Part, In-Situ Observation, Replica Observation, Micro-grid Abstract Woven, knitted and braided fabrics are fabricated by intertwining fiber bundles, and the whole structure consists of repeating units of crossing fiber bundles. Clearly the deformation and fracture behavior at the fiber crossing part determine the mechanical behavior of textile composites, so that it is important to understand the behavior at the fiber crossing part quantitatively. In this paper, for the purpose of investigating the microscopic fracture behavior at the fiber crossing part in textile composites, microscopic in-situ observation was carried out during tensile tests in flat braided composites. The micro-fracture process of flat braided composites was clarified by combining the results at the specimen edge by the replica method and those at the specimen surface by the micro-grid method. Moreover, a numerical analysis model was proposed for simulating the micro-fracture behavior of braided composites. 1 Introduction Using well developed textile technologies, such as weaving, knitting and braiding, structural preforms may be produced with an architecture such that the continuous orientation of fibers at any point is not restricted to a plane. The composite with such preforms as reinforcements is highly suitable for automated process, so that this technique has the potential for low cost production. Also, reliable composite structural components of complex shapes can be achieved. Because of the three dimensional nature of the fiber architecture, such structures are less prone to delamination and their impact resistance is increased significantly. As a result, the design allowances increase together with the cost reductions, and remove barriers to use of composite structure for many applications. The architecture of a fiber preform in the composite is complex, so that the parameters controlling its mechanical properties are numerous, e.g. type of fiber bundles, fiber orientation angle, weaving structure, crimp ratio of fiber bundles, fiber volume fraction in the composite, and so on. Since these factors are related to each other, it is difficult to control them independently. In other words, the best possible combination for desirable dimension and mechanical properties can be obtained by controlling these factors properly. Woven, knitted and braided fabrics, in spite of differences on fabrication process and mechanical behavior, are fabricated by intertwining fiber bundles and the whole structure consists of repeating units of crossing fiber bundles. Because of the fiber crossing, the undulation of fiber bundles is page_133 Page 134 generated and the particular transmission mechanism of force between fiber bundles is produced. Clearly the deformation and fracture behavior at the fiber crossing part determines the mechanical behavior of textile composite, so that it is important to understand the behavior at the fiber crossing part quantitatively. These circumstances require the approach in terms of experimental micro-fracture mechanics, which visualize, quantitate and model the actual micro-fracture. In this paper, for the purpose of investigating the microscopic fracture behavior in textile composites, microscopic in-situ observation was carried out during tensile tests in flat braided composites. The replica method was introduced to examine the fracture process at the edge, and the micro-grid method was used for the fracture process at the surface. The micro-fracture process of flat braided composites was clarified by combining the results by the replica method and the micro-grid method. Moreover, an attempt was made to propose a numerical analysis model for simulating the micro-fracture behavior of braided composites. 2 Materials and Experiments 2.1 Flat Braided Composites A flat braided fabric used in this study is illustrated in Fig.1(a). For the flat braided fabric, fiber bundles initially move to the left (or right), reverse themselves at the side edge and move to the right (or left). Accordingly, fiber bundles are continuous in the fabric. Materials used in this study were glass fiber bundles (ER520-F165: Nippon

Electric Glass Co., Ltd.) and epoxy resin (EPOMIK R-140: Mitui Petro-Chemical Industry Co.). A flat braided fabric with 2/2 intersection repeat was fabricated by using a flat braiding machine with 25 spindles. The braiding angle was about 20 degree. The fabric was impregnated with resin, degassed in a vacuum chamber, and subsequently a flat bar was manufactured by hand lay-up. To investigate the effects of fiber continuity at the side edge on deformation and fracture behavior of composites, two types of specimen were prepared as shown in Fig.1: the first one was a normal flat braided composite (Noncut specimen, (a)), the second type was a flat braided composite whose fiber bundles at the side edges were cut (Cut specimen, (b)) for comparison. Dimensions of the specimen are 1.0mm thick, 100mm long, 13mm wide for Cut specimens and 15mm wide for Noncut specimens. FRP end-tabs were placed on the specimen, which left the gage length of 50mm.

Fig. 1 Flat braided composite ((a) Noncut specimen) and flat braided composite whose fiber bundles at side edges were cut ((b) Cut specimen). page_134 Page 135 2.2 Preparation for In-Situ Observation The fracture process was measured by the replica observation at the edge surface. For the replica observation, the surface at the one side of edge of the specimen was polished. In the case of Cut specimens, the other side of edge which was not polished, was cut by diamond cutter for symmetry. For the Noncut specimens, a resin rich region was provided to polish the edge surface without cutting the continuous fiber bundle. The deformation and fracture process at the surface of specimen was observed by the micro-grid method1). The micro-grids were printed on specimen surfaces by using the photolithography technique. First, the surface of the specimen was polished, then coated by the photo-resist or photo-chemical reactive resin. The specimens was heated to cure the resist, and the surface was exposed to the light through a photomark, or glass plate with micro-grids. After that, the exposed part of the resist was removed in the developer, and vacuum-evaporated metal was deposited on the surface. Finally, the remaining resist was removed in the solvent to prepare the micro-grids on the specimen surface. The size of the micro-grids is 15mm ´ 15mm, and the interval is 5mm. 2.3 Experimental Procedure Tensile tests were performed using Tensilon UTM-1 500-W (Toyo Measuring Instruments Co., Ltd.) with a crosshead speed of 0.5mm/min. During the tensile tests, the testing machine was periodically stopped and the polished edge surface of a specimen was replicated on a replica film (acetyl cellulose film) with methyl acetate as resolvent. The replica film was observed by optical microscopy. In the same way, the testing machine was periodically stopped and the micro-grids at the surface of specimen were observed by a video microscope (Keyence, VH-6300) and the image was recorded in VCR. 3 Experimental Results 3.1 Tensile Properties First, to obtain the tensile properties of each specimen, the tensile tests without replica observation was performed. Figure 2 shows relation between tensile stress and strain in Cut and Noncut specimens. A knee-point is observed at small strain for both Cut and Noncut specimens and the decrease in the slope of stress-strain curve for Noncut specimens is smaller than Cut specimens. Noncut specimens possess tensile strength of approximately 1.6 times as high as Cut specimens.

From the visual observation during the testing, in the case of Cut specimens, the debonding of the fiber crossing part at the cut out side edges was found at the stress level where the slope of the stress/strain curves began to decrease. Also, the whitening of the fiber crossing part at the center of the specimen was observed. As the tensile stress increased, the debonding region propagated from the specimen edge to the inside of the specimen, and final fracture of the specimen occurred from the edge. On the failed specimen surface, pull-outs of fiber bundles were observed at the edge, and fractures of fiber bundles were observed at the center of the specimen. In the case of Noncut specimens, after the slope of the stress/strain curves began to decrease, the whitening at the fiber crossing part was observed. As the tensile stress increased, number of whitening increased and the final fracture of the specimen occurred from the edge due to the fracture of fiber bundles. As mentioned above, the tensile properties were greatly dependent on whether the fiber bundles were continuous at the edge or not. page_135 Page 136

Fig.2 Relation between tensile stress and strain in Noncut and Cut specimens. 3.2 Fracture Process by In-Situ Observation From the obtained stress/strain curves, the strain at the knee-point was 0.2% for Cut specimens and 0.3% for Noncut specimens. Accordingly, the testing machine was stopped at every 0.2% for Cut specimens and every 0.3% for Noncut specimens to conduct replica observation of polished edge surfaces or observation of the micro-grids at the surface in flat braided composites. In the case of Cut specimens, at the polished edge surface, the crossing part of cut fiber bundles could be observed. Figure 3 shows delamination onset and growth between fiber bundles at the edge surface of Cut specimens obtained by replica observation ( (a)strain=0.8%, (b)1.2%, (c)1.4%).

Fig.3 Delamination onset and growth between fiber bundles in the edge surface of Cut specimen page_136

Page 137 In these Cut specimens, the first microscopic damage is delamination of fiber bundles at the center of fiber crossing part (see (a)). The delamination progressed at the fiber crossing part and then along the fiber bundles (see(b)). As the tensile stress increased further, the delamination reached the surface of the specimen (see(c)), and finally, the fiber bundles were pulled out of matrix. In the case of Noncut specimens, at the polished edge surface, a part of fiber bundles turning at the edge could be observed. In Noncut specimens, the first microscopic damage is fracture of a filament in fiber bundles at 0.6% of strain. Figure 4 shows the filament fracture multiplication at the edge surface of Noncut specimen obtained by replica observation ((a) strain=0.6%, (b)1.2%, (c)1.8%). As the tensile stress increased, the number of filament fracture increased (see(b)) and the number was saturated at a regular interval of filament fractures.

Fig.4 Filament fracture multiplication at the edge surface of Noncut specimens. For the observation of the micro-grids at the surface of both Cut and Noncut specimens, the deformation of the micro-grids was observed at fiber crossing part near the edge of the specimen. In the case of Cut specimens, the micro-grids, which were arranged in a straight line at strain of 0.0%, began to lean to the left at strain of 1.4% at fiber crossing part. As the tensile stress increased, the deformation of the micro-grids increased in quantity and a gap between adjacent micro grids was page_137 Page 138 observed along fiber bundle at strain of 1.8% as shown in Figure 5(a). The strain of 1.4% corresponded to the strain in which delamination reached the surface of the specimen in the replica observation. As a result, it is considered that the binding force between fiber bundles was lost by delamination at the edge and the pulled out of fiber bundles was found as a deformation of the micro-grids as shown in Figure 5(b). On the other hand, in the case of Noncut specimens, a gap between adjacent micro grids was observed locally at strain of 1.5%, and clearly opened at strain of 1.8% as shown in Figure 6. The strain of 1.5% corresponded to the strain in which the number of filament fracture was saturated in the replica observation. As a result, it is considered that filaments in fiber bundles turning at the specimen edge were fractured at the edge surface, and then the fracture of filaments in same fiber bundle propagated from the edge to the inside.

Fig.5 Deformation of micro-grids of Cut specimens.

Fig.6 Deformation of micro-grids of Noncut specimens. page_138 Page 139 4 Numerical Analysis Model So far, the micro-fracture process of flat braided composites was clarified by combining the results by the replica method and the micro-grid method. Next, the above experimental results were related to numerical analysis for predicting micro-fracture properties of braided composites which have been developed2,3). Figure 7 shows the weaving structural model for simulating the mechanical behavior of braided composites, in which the weaving structure of the unit-cell was expressed by connecting beam elements. The thick lines and the fine lines express the fiber bundle and resin in the braided composites, respectively. This model consists of fiber bundle elements, surface resin elements, and cross resin elements which express the resin existing between crossing fiber bundles. The micro-fracture process of flat braided composites obtained by the experiments were expressed by the weaving structural models for Cut and Noncut specimens as shown in Figures 7 and 8, respectively. Figure 7(a) shows the fracture of a cross resin element corresponded to the delamination

Fig.7 The micro-fracture process on the weaving structural model for Cut specimen.

Fig.8 The micro-fracture process on the weaving structural model for Noncut specimen. page_139 Page 140 between fiber bundles at the edge of Cut specimens. Fiber bundles without binding force between intertwined fiber bundles were pulled out (Figure 7(b)). Figure 8(a) shows the fracture of a fiber bundle element corresponded to the filament fracture in the fiber bundles turning at the edge of Noncut specimens. After that, the fiber bundle element was fractured at the surface as shown in Figure 8(b). Moreover, on the basis of experimental results, a more microscopic numerical analysis model is proposed to simulate the mechanical behavior of the braided composites in consideration of the micro fractures. In the weaving structural model, a fiber crossing part is expressed by one beam element. In order to treat more microscopic deformation and fracture behavior at the fiber crossing part such as delamination onset and growth between fiber bundles of Cut specimens, an analytical model which represents the fiber crossing part is necessary, for example, as shown in Figure 9. Also, in order to consider the microscopic fracture of a filament inside fiber bundles for Noncut specimens, an analytical model which represents the filaments impregnated with resin is required.

Fig.9 Analytical model which represents the fiber crossing part. 5 Conclusion In this study, for the purpose of investigating the microscopic fracture behavior at the fiber crossing part in textile composites, microscopic in-situ observation in flat braided composites was carried out. The micro-fracture process of flat braided composite was clarified by combining the results at the edge by the replica method and at the surface by the micro-grid method. In the case of flat braided composites, whose edge were cut, or Cut specimens, the delamination of fiber bundles at the fiber crossing part of the specimen edge occurred first, and then the fiber bundle was pulled out. In the case of flat braided composite in which fiber bundles were continuous, or Noncut specimens, filaments turning at the edge of fiber bundles were fractured at the edge and the filament fractures led to the fracture of fiber bundles. Experimental results in consideration of micro fractures are very useful in the construction of the model for prediction of the macroscopic properties of the composites from the microscopic characteristics, that is, design of textile composites in consideration of micro-fracture. Reference 1) N.Takeda, H.Niizuma, S.Ogihara and A.Kobayasi, Experimental Mechanics, 37, 182(1997). 2) A.Nakai, A.Fujita, A.Yokoyama and H.Hamada, Composite Structure, 32, 501(1995). 3) H.Hamada, A.Fujita, Z.Maekawa, A.Nakai and A.Yokoyama, J. of Science and Engineering of Composite Materials, 4,109(1995).

page_140 Page 141

In Situ Observation of Micro Damage under Tensile Load Single Fiber, Fiber Bundle & Woven Fabric Kazuaki NISHIY ABU 1 and Masaru ZAKO 1 1 Osaka University, Graduate School of Eng., 2-1 Yamada-oka, Suita, Osaka 565-0871, JAPAN Keywords; In situ SEM observation, Single fiber, Fiber bundle, Woven cloth, Glass fiber, Vinyl ester resin, Surface treatment, Interfacial debonding, Crack propagation Abstract This study describes about how the fiber surface treatments are affected on both mesoscopic damages, such as interfacial debonding and matrix cracking in the embedded single fiber or strand, and macroscopic ones, such as inter-bundle delamination of woven fabric composites. The damage observations have been carried out under two types of tensile testing: (1) tensile test into scanning electron microscope (in situ SEM observation) and (2) universal tensile test (in situ macro observation). Test specimens are vinyl ester resin matrix embedded by single filament or strand, and reinforced by laminated woven cloth. In order to investigate the effects of interfacial properties, each specimen was fabricated using three types of surface treated fiber. It was obvious that the interfacial properties were able to be evaluated by in situ observation of each failure process, such as interfacial debonding, crack linking, and crack propagation. The effect of the fiber surface treatment on damages has been revealed in each scale level, i.e., filament, strand, and woven fabric. Introduction It is important for an improvement of reliability on strength to correlate the macroscopic fracture mechanism with microscopic damages such as transverse crack into the fiber bundle, and the interfacial debonding around the single fiber under the applied load. Microscopic damage behaviors are remarkably affected by the interfacial properties, i.e. surface treatment conditions of the fiber [1]. As shown in Fig.1, there are three viewpoints for composite interfaces. There are manufacturing level, mechanical level and molecular level. In this study, we focus on the effect of interfacial property on the damage behaviors in mechanical level. Macroscopic damage of woven fabric laminated composite is supposed to initiate from a mesoscopic damage into the fiber bundle and a microscopic damage around the single fiber. These damage behaviors might be caused by the interfacial properties in molecular level. We have proposed the embedded single fiber transverse tensile (ESFTT) test in order to investigate the effect of fiber surface treatments on the microscopic damage [2]. The numerical simulation of microscopic damage for ESFTT test has been carried out [3]. The purpose of this study is to establish the evaluation technique by the visualization of the interfacial properties in fiber reinforced composites in order to understand the effect of the interfacial properties on the damage behavior and the connectivity of the interfacial property in mechanical level. page_141 Page 142

Fig. 1 Connectivity of interfacial properties. Materials

Fiber and Matrix The embedded fibers or the reinforcement are single filament, single strand and plain woven cloth. Vinyl ester resin supplied from Showa polymer Co.LTD (R-806) is used for all test specimens. Single filament is E-glass fiber of about 11 m m in diameter, which is picked up the fiber strand. Woven cloth is glass fabric supplied from Asahi fiberglass Co.LTD (MS250), which is commercially surface-treated in woven cloth. Surface Treatment As shown in Fig.2, we treat three types of surface treatment. There are two types of silane coupling agents, which are acrylic silane agents and epoxy silane agents and heat cleaning, respectively. Acrylic silane treatment is to make chemical reacted interface between the glass fiber and vinyl ester matrix. Epoxy silane treatment is to make chemical unsuitable reaction with glass fiber - vinyl ester matrix interface. Heat cleaning is to make no-chemical reaction (i.e. mechanical contact) with glass fiber - vinyl ester matrix interface.

Fig.2 Interfacial properties and silane coupling agents. page_142 Page 143 Test Specimens The configurations of test specimens for in situ SEM observation and in situ macro observation are shown in Fig.3 and Fig.4 respectively. A single filament or/and a single strand is embedded into the matrix, and is aligned transversely to tensile direction. Woven cloth specimen is made from GFRP laminate of 5 plies is fabricated by hand-lay up method. The warp and weft directions of the fabric are aligned to longitudinal and transverse direction of laminates respectively. All the specimens are post-cured at 373K for 2 hours. Polyvinyl chloride end tabs are glued to the specimen ends using an epoxy adhesive. In the test specimen for in situ SEM observation (Fig.3), the observation surface is grounded by buffing (Al2O3, 0.05 m m) and is sputtered by platinum.

Fig.3 Geometry of test specimen for in situ SEM observation.

Fig.4 Geometry of test specimen for in situ macro observation. Experimental Details In Situ SEM Observation Experimental apparatus used for in situ SEM observation is shown in Fig.5. The tensile stage is mounted on scanning electron microscope produced by Hitachi Co.LTD (type No. S-2460N). Tensile displacement for test specimen is applied at constant speed. For the test specimen embedded by single filament or strand, the transverse tensile load is applied to the longitudinal direction to fiber. Semiconductor-loading cell and AE (Acoustic Emission) sensor produced by NF Co.LTD (type No. AE-901S, 140KHz) are installed in the tensile stage as shown in Fig.5. An electron beam is radiated at the edge surface of specimen, and SEM images are recorded in video as animation. Tensile load and AE events also are recorded simultaneously. The influence of surface treatment conditions on the microscopic damage behavior, e.g. the initiation and propagation of micro cracks occurred at intra/inter-bundle of the fiber is investigated by correlating the observation of VTR images to the changes of tensile load or AE event counts. page_143 Page 144 In Situ Macro Observation As in situ SEM observation restricts to the damages which occurred at only the edge surface of specimen, the optical experimental method is applied for the internal observation of micro damages occurred at the embedded strand or laminate composites. The scheme of experimental apparatus used is shown in Fig.6. The influence of fiber surface treatments on the behavior of damages occurred into the fiber strand, or fiber bundle is investigated by the tensile test.

Fig.5 Overview of tensile stage into SEM.

Fig.6 Experimental apparatus for in situ macro observation. Results and Discussions Embedded Single Filament Transverse Tensile Test VTR images of the damage process in each time are shown in Fig.7, 8 and 9. In case of acrylic silane treated fiber (Fig.7), no damage is seen even until a considerably long time (i.e. high load level) as found clearly from Fig.7(a). It is obvious that the interfacial bonding strength is improved by the surface treatment. The initial debonding occurs at outer edge of the filament, and it progresses quickly along the interface around the filament as appeared in Fig.7(b), and cracking is opened (Fig.7(c)). The interfacial debonding stops at certain angle, the progression of crack transit to the matrix as shown in Fig.7(d), and the macro fracture occurs. In contrast, for epoxy silane treated fiber (Fig.8), the interfacial damage initiates at very low load level (Fig8(a)), because there is chemical mismatch of reactive radical between the treated fiber and matrix. The obvious interfacial debonding appears gradually, and it progresses at the both side to the circumference direction of filament as shown in Fig.8(b), (c) and (d). Moreover, for non-treated fiber(Fig.9), the interfacial damage initiates at lower load level(Fig9(a)), because there is no chemical reactive radical between the fiber and matrix. Fig.9(b) shows the interfacial debonding progresses gradually at the both side as well as the fiber treated by epoxy silane(Fig.8). As shown in Fig.9(c) and (d), the crack progresses in matrix at vertical to tensile direction. From the experimental results, it is found clearly that the micro damage behaviors are quiet different according to the surface treatment conditions. In other words, we can state that the experiment for single filament is useful for an evaluation of the interfacial properties. It can examine for only the effect of chemical parameter on damage behaviors, because there is few effects of fiber orientation parameter and wetting parameter (i.e. molding conditions) etc. page_144 Page 145

Fig.7 VTR images of the embedded filament under transverse tensile load. (Acrylic silane)

Fig.8 VTR images of the embedded filament under transverse tensile load. (Epoxy silane)

Fig.9 VTR images of the embedded filament under transverse tensile load. (Heat cleaning) Embedded Single Strand Transverse Tensile Test VTR images of the embedded single strand specimen are shown in Fig.10. In case of acrylic silane treated fiber (Fig.10(a)), a large cracking occurs into the strand, and the final rapture occurs quickly. For epoxy silane treated fiber (Fig.10(b)), the interfacial debonding is not easy to connect each other, because the filaments disperse into the matrix. On the other hand, in case of epoxy silane treated fiber (Fig.10(c)), the interfacial debonding occurred around the filament is easy to link each other, because the filaments are closely arranged. It does not only progress to vertical tensile direction, but also to any directions into the strand. It is thought that the test results remarkably depend on the effect of filament arrangement. The effect of the fiber surface treatments appears in damage process of the edge surface into the strand under transverse tensile test. The weaker adhesive interface is easier to initiate the debonding, and the crack links to closer interfacial debonding. Thus it is guessed that the interfacial properties appeared in the embedded single filament transverse tensile test (Fig.7-8), have much influences to tendency of damage progression in the embedded single strand transverse tensile test (Fig.10). page_145 Page 146

Fig.10 VTR images of the embedded strand under transverse tensile load. Fig.11 shows the internal damages occurred in the embedded strand just after rupturing of the specimen. Whitening occurs into the strand with an increasing the tensile load. The change of whitening is quantified by means of image processing technique. The pixel percentage for whitening and tensile stress-displacement curve are shown in Fig.12. This percentage for acrylic silane treated fiber (Fig.12(a)) increases in higher stress level than epoxy silane treated fiber (Fig.12(b)), but it gradually increases in case of heat cleaning (Fig.12(c)). This damage behavior agrees with the results obtained by in situ SEM.

Fig.11 Damages into strand after rupture.

Fig.12 Stress-displacement curve and damage fraction. page_146 Page 147 Woven Cloth Laminate Tensile Test Fig.13 shows the relations between tensile stress, AE events and testing time for acrylic silane, epoxy silane and heat cleaning, respectively. Fig.14 shows VTR images just before final breaking. It can be suggested that the interfacial bonding strength for acrylic silane is higher than any other treatments, because the cracks into the bundle are few. It is also suggested from AE test results. In case of epoxy silane and heat cleaning, however, many cracks occur into the bundle as the displacement increase. It is obvious that the interfacial adhesive strength for both epoxy silane and heat cleaning is weaker than that for acrylic silane. Though similar damages occur for epoxy silane and heat cleaning, it behaves that AE events for epoxy silane are more than that for heat cleaning.

Fig.13 Tensile load, AE vs. testing time.

Fig.14 SEM images just before fracture. page_147 Page 148 Moreover it is guessed that the crack occurs easily, because the surface treatment agents of epoxy silane does not have a reactive radical with the resin, and the energy of crack initiation for heat cleaning is lower, because of only a mechanical contacted interface. Fig.15 shows the damage images occurred into woven cloth fabric lamina after rupturing. The internal cracks into the transverse fiber bundle are few for acrylic silane (Fig.15(a)), and there are many cracks for epoxy silane(Fig.15(b)) and heat cleaning (Fig.15(c)). It is verified that the interfacial adhesive strength for both epoxy silane and heat cleaning is weaker than that for acrylic silane. It can be recognized that many transverse intra-bundle cracks for heat cleaning are much finer than that for epoxy silane. Moreover, there are no damages in fiber bundle along tensile direction for coarse woven density cloths used in this experiment.

Fig.15 Damages occurred into woven cloth fabric lamina after rupturing. Conclusion The effect of fiber surface treatment conditions on micro damage behaviors in each scale level, such as single filament, single strand and woven cloth laminate, have been investigated by both in situ SEM observation and in situ internal observation. From the results of the embedded single filament transverse tensile test, it is found that the fiber surface treatment affects strongly on microscopic damage behaviors around a single fiber. For the embedded single strand specimen, the damage progression remarkably depends on the filament arrangement. Such a tendency agrees with the results of edge surface damage observation by in situ SEM. With woven cloth laminate, the initiation and propagation of interfacial debonding into the fiber bundle were remarkably affected by kinds of surface treatment. From the experimental results, we have made a clear that microscopic behaviors of the interfacial debonding occurred around the single filament with different surface treatments caused to affect on the both mesoscopic damage mechanisms for the embedded strand specimen and macroscopic ones for woven cloth laminates. It is thought that the effect of interfacial properties in practical composites should be clarified by comparing the microscopic experimental results with macroscopic ones. In addition, in situ observation is useful for understanding the damage behaviors in composites with different interfacial properties surface treated fiber. Acknowledgment We would like to express our thanks to Dr.Yoshimichi Fujii of Seikow Chemical & Machinery Co.LTD, Japan for his valuable suggestions and cooperation. References [1] K.Nishiyabu and M.Zako, Proceedings of 7th International Conference on Composite Interfaces (ICCI-VII) (1998) [2] K.Nishiyabu, A. Yokoyama and M.Zako, Proceedings of 5th Japan International SAMPLE Symposium, 847852(1997) [3] K.Nishiyabu, A. Yokoyama and H.Hamada, Jour. of Compos. Sci. Technol., 57, 11031111(1997) page_148 Page 149

Impact Resistance of Multi-Reciprocal Braided Composites E.Kwan1, X.R.Xiao1, H.Wang2, S.V.Hoa1, H.Hamada3 Concordia Center For Composites, Concordia University, Montreal, Canada1 Industrial Materials Institute, National Research Council, Montreal, Canada2 Kyoto Institute of Technology, Kyoto,Japan3 It is well known that laminated composite structures are susceptible to impact loading. Among the damage modes caused by low-velocity impact, delamination is found to be most detrimental to the structural performance such as stiffness, strength and fatigue resistance. Extensive research has shown that although impact behaviour of composite materials may be affected to certain extent by a number of factors such as stacking sequence, fiber, matrix and interface, the most significant improvement in impact resistance is achieved by introducing fiber reinforcement through the thickness by methods such as stitching or 3D-fabrics. During the 1980's two distinctly different methods for 3D braiding were developed: two-step and four-step. Presently these methods are still in the laboratory stage, i.e. only small size samples have been produced. To improve the impact resistance of braided composites using the available technology, Hamada et al together with Murata Machinery have developed techniques to produce multi-reciprocal (MR) braids using maypole-type braiders. Composites made from these preforms are expected to have superior impact resistance than that of conventional composite materials. This paper presents the experimental results of the low energy impact properties of composite panels made from multi-reciprocal in comparison with those made from conventional 2D braids. The mechanism of enhanced resistance of MR composites is also discussed. page_149 Page 151

ANALYSIS AND MODELING page_151 Page 153

Free Vibration Analysis of Cantilevered Laminated Trapezoidal Plates Kenji HOSOKAWA, Jimin XIE and Toshiyuki SAKATA Department of Mechanical Engineering Chubu University 1200 Matsumotocho, Kasugai, Aichi, 487-8501 Japan Keywords: free vibration, vibration of continuous system, composite material, laminated plate, trapezoidal plate, cantilevered plate Abstract A numerical approach for analyzing the free vibration problem of a laminated FRP (fiber reinforced plastic) composite plate has already been proposed by the authors. In the present study, this approach is applied to a cantilevered laminated trapezoidal plate. First, it is attempted to estimate numerically the natural frequency of the plate, and the convergence and accuracy of the results are discussed. Next, the natural frequency and mode shape of the plate are calculated. Furthermore, the natural frequency and mode shape of the plate are obtained experimentally. These experimental results are found to agree well with the numerical results computed using the measured material properties of the lamina. 1 Introduction Since a laminated plate is an important structural member, many studies on the free vibrations of a symmetrically laminated plate have been reported in the literature. Also, the vibrations of cantilevered laminated trapezoidal plates are interesting in the field of a structural engineering and an aerospace engineering. However, one can find only a few reports on the vibrations of symmetrically and unsymmetrically laminated trapezoidal plates. For example, Qatu calculated, using the Ritz method with algebraic polynominals, the natural frequencies of laminated composite

trapezoidal plates with completely free or cantilevered boundary conditions[1,2]. The free vibrations of symmetric and unsymmetric laminated plates with trapezoidal planform and arbitrary boundary conditions were analyzed by using the pb-2 Ritz method[3,4]. As above mentioned, one can find the numerical researches. However, one can find few reports on an experimental analysis of the free vibration of the cantilevered laminated trapezoidal plate. The authors have already proposed a numerical approach for analyzing the free vibration problem of a laminated FRP composite plate[5,6]. In the present paper, this approach is modified for applying to the cantilevered laminated trapezoidal plate. First, it is attempted to estimate numerically the natural frequency of the plate, and the convergence and accuracy of the results are discussed. Secondly, the natural frequency and mode shape of the cantilevered laminated trapezoidal plate are computed. Finally, the natural frequency and mode shape of the plate are page_153 Page 154 obtained by using the experimental modal analysis technique. And then, the experimental results are compared with numerical ones estimated using the measured material properties of the lamina. 2 Numerical Approach 2.1 Frequency Equation Consider a cantilevered laminated trapezoidal plate as shown in Fig.1. By using the classical laminate theory, the free transverse vibration of the laminated plate is governed by

where h is the thickness, u(x,y,t) and v(x,y,t) are displacements in the x and y directions, respectively, w(x,y,t) is the transverse deflection, ¨w = ¶2w / ¶t2, t is the time variable, and r is the density. The symbols L1[ ], L2[ ], and L3[ ] are the differential operators for the static bending problem of the laminated plate. The solution of equation (1) is written as

Fig.1 Cantilevered laminated trapezoidal plate. page_154 Page 155 In equation (2), the functions G1(x,y,x,h), G2(x,y,x,h), and G3(x,y,x,h) satisfy the boundary condition and the differential equation represented by

where d( ) is Dirac's delta function. Because it is very difficult to solve analytically integral equation (2), the integration is carried out approximately under the assumption that the integrands in equation (2) are constants in the small regions hatched in Fig.1(b). The plate transverse deflection w(x,y,t) is assumed as

From the third equation in equation (2) and equation (4), the frequency equation is given by

where DSn and (xn, yn) are the area and typical point of the nth small region of the plate shown in Fig.1(b), respectively, w is a radian frequency, and dm,n is Kronecker's delta. Also, in the case in which the concentrated mass is attached to the Ith dividing point (xI,yI) of the plate, from d'Alembert's principle, the frequency equation is expressed as

2.2 Estimation of Functions G[1],G[2], and G[3] The functions G1,G2, and G3 are assumed in a power series form as the following equations,

page_155 Page 156 where Ai(x,h), Bi(x,h), and Ci(x,h) are constants determined by the position (x,h) where unit load acts, and k and l are positive integers determined according to the positive integer i. The functions y1(x,y), y2(x,y), and y3(x,y) are determined such that the functions G1,G2, and G3 satisfy the boundary conditions of the plate, respectively. The strain energy U due to bending and the work T due to unit load are estimated. According to the Ritz minimizing process, a set of simultaneous equations with respect to Ai(x,h), Bi(x,h), and Ci(x,h) is obtained as

By substituting Ai(x,h), Bi(x,h), and Ci(x,h) estimated from the above equation into equations (7), (8), and (9), respectively, one can obtain the functions G1,G2, and G3. For the case of the cantilevered plate shown in Fig. 1, the functions y1(x,y), y2(x,y), and y3(x,y) may be expressed as

3 Convergence and Accuracy of Numerical Approach

To discuss the convergence and accuracy of the proposed approach, natural radian frequencies of the cantilevered antisymmetrically laminated trapezoidal plate were estimated. The stacking sequence is [30°/ - 30°/30°/ - 30°] and the plate material is assumed to be E1 / E2 = 40, G12 / E2 = 0.5, and n12 = 0.25. The dividing pattern of a trapezoidal plate is shown in Fig. 1(b). A natural radian frequency is non-dimensionalized by using the bending stiffness parameter D0 = E1h3 / 12(1 - n12n21). For the case of the cantilevered antisymmetrically laminated trapezoidal plate, the functions G1,G2, and G3 are obtained as follows. The strain energy U and the work T are estimated by equations (13) and (14), respectively.

page_156 Page 157

where Ai,j(i,j = 1,2,6) are the extensional stiffnesses, Bi,j(i,j = 1,2,6) are the coupling stiffnesses, and Di,j(i,j = 1,2,,6) are the bending stiffnesses. These stiffnesses are evaluated from the moduli of elasticity E1 and E2 in the direction of parallel and normal to the fiber, Table 1 Convergence of natural radian frequencies of cantilevered four layered trapezoidal plate; [30° / - 30°/30° / - 30°], N = 272 = 16 ´ 17, b /a=1.0, c /a=0.5, E1 / E2 = 40, G12 / E2 = 0.5, n12 = 0.25 Number of terms I 3 6 10 15 21 28 36 45 55 Ref.[3]

s=1 2.984 2.905 2.803 2.768 2.753 2.749 2.747 2.745 2.747 2.751

s=2 13.25 10.35 9.818 9.646 9.501 9.467 9.445 9.440 9.439 9.450

s=3 23.02 14.91 14.45 14.36 14.26 14.13 14.07 14.04 14.05 14.14

s=4

s=5

34.48 27.51 24.99 24.03 23.62 23.61 23.52 23.57 23.70

39.75 30.18 27.36 26.79 26.57 26.38 26.33 26.33 26.43

Table 2 Convergence of natural radian frequencies of cantilevered four layered trapezoidal plate ; [30°/ - 30°/30°/ - 30°], I=55, b /a=1.0, c /a=0.5, E1 / E2 = 40,G12 / E2 = 0.5, n12 = 0.25 Division number N 20 42 72 110 156 210 272 Ref.[3]

s=1 2.672 2.716 2.732 2.739 2.743 2.746 2.747 2.751

s=2 9.101 9.302 9.372 9.404 9.421 9.432 9.439 9.450

s=3 12.44 13.33 13.69 13.86 13.96 14.02 14.05 14.14

s=4 21.53 22.70 23.13 23.34 23.46 23.52 23.57 23.70

page_157

s=5 23.24 25.06 25.71 26.00 26.17 26.26 26.33 26.43

Page 158 respectively, shear modulus G12, poisson's ratio n12, n21, and fiber orientation angle qr of each layer (See Fig.1(a)). As described in the previous chapter, by substituting equations (13) and (14) into equation (10), one can obtain the functions G1,G2, and G3. Table 1 shows the effect of the number of terms I of the functions Gj(j = 1,2,3) on the convergence of the natural radian frequencies of the plate. The effect of the number of divisions N on the convergence of the natural radian frequencies is presented in Table 2. From Tables 1 and 2, it follows that one can obtain sufficiently converged values for lower modes by using I=55 to estimate the functions Gj(j = 1,2,3) and by dividing the plate into N = 272 = 16 ´ 17 small regions. Also, one can see the similar convergence for the cantilevered symmetrically laminated trapezoidal plate. The numerical results obtained by the other investigator[3] are also tabulated in these tables. Judging from the comparison it follows that the values obtained by the present approach are accurate. 4 Justification of Numerical Results To justify the numerical results, experimental studies were carried out for the cantilevered symmetrically and antisymmetrically laminated trapezoidal plates. The plates were clamped by using a rigid clamping fixture. The plate configuration of the trapezoidal plates is as follows: a=0.2[m], b=0.2[m], c=0.1[m]. The stacking sequence of the symmetrically laminated plate is [30°2/ - 30°2/ - 30°2/30°2] and the plate thickness h is 1.61 ´ 10-3[m]. The stacking sequence of the antisymmetrically laminated plate is [30°2/ - 30°2/30°2/ - 30°2] and the plate thickness h is 1.60 ´ 10-3[m]. Each layer material is a carbon fiber reinforced plastic (CFRP). Table 3 Material properties of layer ; CFRP E1 (GPa) E2 (GPa) G12 (GPa) n12 r (kg/m3) 97.6 6.26 5.18 0.33 1535 Table 4 Natural frequencies of cantilevered eight layered trapezoidal plates ; CFRP, [30°2/ - 30°2/ - 30°2/30°2], [30°2/ - 30°2/30°2/ - 30°2] Natural frequency [Hz] Model Symmetrically laminated plate Antisymmetrically laminated plate order Experiment Computation Experiment Computation 1 43.75 43.89 43.75 43.96 2 138.8 135.5 166.3 148.7 3 245.0 244.0 232.5 232.0 4 359.7 343.6 408.3 403.7 5 446.3 436.0 442.5 423.6 page_158 Page 159 The measured material properties of the lamina are listed in Table 3. Natural frequencies and mode shapes of the plates were obtained by using the experimental modal analysis technique. Table 4 presents the natural frequencies determined numerically and experimentally. As the vibration pick-up, a strain gage was affixed at the location where nodal lines disappear. Figures 2 and 3 show the mode shapes and natural frequencies of the plates obtained numerically and experimentally. In these figures, the dotted lines present the nodal lines obtained by the experimental modal analysis technique and the solid lines show those calculated by the numerical approach. The notation · represents the location of the attached accelerometer. For the numerical results, the mass of an accelerometer (4.8[g]) was considered as a concentrated mass. Therefore, to compute the natural frequency, frequency equation (6) was used. From Table 4 and Figures 2 and 3, one can see that the difference between experimental and

Fig.2 Nodal patterns and natural frequencies of symmetrically laminated trapezoidal plate ;[30°2/ - 30°2/30°2/30°2], CFRP.

Fig.3 Nodal patterns and natural frequencies of antisymmetrically laminated trapezoidal plate ; [30°2/ -30°2/30°2/ - 30°2], CFRP. page_159 Page 160 numerical natural frequencies is about 6% at the most except for the second natural frequency. Also, one can observe it in Ref.[7] that the second experimental natural frequency is bigger than the computational one. On the other hand, it follows that one can find the good agreements between all experimental and numerical mode shapes. 5 Conclusions The numerical approach using the functions for the static bending problem of the plate was applied to the cantilevered laminated trapezoidal plate. The natural frequencies of the cantilevered symmetrically and antisymmetrically laminated trapezoidal plates were calculated. From the numerical results, it follows that one can estimate sufficiently converged values for lower modes by using this approach. The natural frequencies and mode shapes of the plates were obtained by the experimental modal analysis technique. From the comparison of experimental and numerical results, one can see the good agreements between these results. Accordingly, it follows that one can accurately estimate natural frequencies and mode shapes by using the numerical approach proposed by the authors. Acknowledgment The authors are grateful to Mr. Kawahara for helping in the computations. References 1. M. S. Qatu, International Journal of Mechanical Science, 36(9), (1994), p.797. 2. M. S. Qatu, Composite Science and Technology, 51, (1994), p.441. 3. K. M. Liew and C. W. Lim, Journal of Sound and Vibration, 183(4), (1995), p.615. 4. C. W. Lim and K. M. Liew, Journal of Acoustical Society of America, 100(6), (1996), p.3674. 5. K. Hosokawa, T. Yada, and T. Sakata, JSME International Journal. 36(3), (1993), p.296. 6. K. Hosokawa, Y. Terada, and T. Sakata, Journal of Sound and Vibration, 189(4), (1996), p.525. 7. K. Hosokawa and T. Sakata, Proceedings of 16th Canadian Congress of Applied Mechanics 97, 1, (1997), p.219. page_160 Page 161

Mechanical Behavior of Sandwich-type Composites with Waste of Fibrous Material As Core Layer

T. Kimura* and Y. Kataoka** *Advanced Fibro-Science, Kyoto Institute of Technology Matsugasaki, Sakyo-ku, Kyoto 606-8585, Japan **Department of Mechanical Engineering, Fukui University 9-1 Bunkyo 3-chome, Fukui 910-8507, Japan Key Words : Waste of fabrics, Compression molding, Sandwich-type composite Abstruct The sandwich-type composites with fiber assemblage as a core layer was molded by applying a compression molding method to the waste cord assemblage and the mechanical behavior of molded composites was measured by the three points loading test. The waste cord of synthetic fabrics, cut waste of non-woven fabrics and pellet-type resin with/without glass fiber reinforcements were used as a surface layer respectively. A fairly poor adhesion at the interface between resin/reinforcements can be seen in the case of pellet-type resin with reinforcements used as a surface layer. Apart from some problem, the compression molding method described in this paper must contribute to material recycle of various wastes of fibrous materials as a sandwich-type composite with good insulating property. Introduction In recent years, increased emphasis has been placed on developing techniques for industrial waste products, with the goals of protecting the environment. Especially, the textile industry has taken a growing interest in developing a system for recycling waste fiber which result from the process of manufacturing products such as textile fabrics, non-woven fabrics, fishing net, lacy cloth etc.. Most of these waste are, however, now destroyed by fire or buried underground. Our recent interest is the material recyclability of waste cords of synthetic fabrics as a heat insulating board, because the fiber assemblage has a small thermal conductivity in itself. The consideration of this work focuses on the mechanical behavior of sandwich-type heat insulating board with waste of fibrous material as a core layer. Compression Molding Method The molding test was performed in the special closed furnace having separated 18 infrared heaters attached at both upper and lower walls as shown in Fig.1. The die was made of aluminum plate of 3mm thick. The waste cord assemblage was stuffed into the die with dimensions of 100´100´100mm for molding test. The waste cord of synthetic fabrics, cut waste of non-woven fabrics and pellet-type resin with/without glass fiber reinforcement were page_161 Page 162

Fig.1 Heating furnace with infrared heaters.

Fig.2 Used wastes of fabrics.

Fig.3 Molding type of insulating board. page_162 Page 163 used as a surface layer respectively and are shown in Figs. 2(a)-(c) In the case of Type I shown in Fig.3, the upper surface layer of the waste cord assemblage of polyester set in the die is melted by the infrared heating system in the first place. Aspects of the waste cord assemblage during the process of our molding system were photographed at the outside of the die as shown in Figs.4 (A)(E). Figures A and B show the aspects before and after melting process, respectively. After the melting process the compression molding is performed with cooling at the outside of the furnace. As a result, the waste cord assemblage is solidified at the upper surface layer. The aspect of this state becomes C. The reverse side of the waste cord assemblage is also melted in the furnace as shown by D and the compression molding is performed again. The finished sandwich-type insulating board with non-melted fiber assemblage in the core layer can be obtained as shown by E. In the cases of Type II and TypeIII shown in Fig.3, the pre-molded plate-type waste of non-woven fabrics of polypropylene and the pellet-type resins of polypropylene and polyester were set on the surface of the waste cord assemblage of polyester and were melted by the infrared heating system in the same way of Type I. In the cases of Type IV and V, the pellet-type resin with glass fiber were used as a surface layer, where the glass fiber was set in the pellet in advance for Type IV and was separated from the pellet for type V. Three points loading test with 80mm span length was performed for the molded specimens with dimensions 100mm´100mm´20mm(thickness) to discuss the mechanical behavior of molded board. Aspect of loading test is shown in Fig.5. Thickness of surface layer is 3.5 mm for all specimens. Results and Discussion The results of fracture and maximum stresses are shown in Figs.6 (a)(e) for Types I ~ V, respectively. The mass of waste cord assemblage as a core layer was varied in the experiments. Therefore, resultant density, namely the void ratio of core fibrous layer was varied for molded specimens. X-axis in these figures means the void ratio of core fibrous layer of waste cord assemblage of polyester. Needless to say, the larger void ratio is good for the heat insulating board.

Fig.4 Aspects of waste cord assemblage during molding.

Fig.5 Aspect of bending test. page_163 Page 164 It is noted from these figures that the fracture stress is decreases with increasing void ratio for all Types. The higher maximum stress can be obtained for TypeIII using pellet-type resin as a surface layer. The lower maximum stress can be, however, seen in Type I. The observation results of fracture surface suggest that the higher strength of such insulating board molded in this paper can be achieved for the board with uniform thickness of surface layer. Namely, the lower maximum stress for Type I may be caused by the non-uniform thickness of surface plastic layer. Though the pellet-type resin with glass fiber was used as a surface layer for Types IV and V, the maximum stress is smaller than that of Type III. This fact may be caused by the fairly poor adhesion at the interface between resin/reinforcement in our compression molding method. Conclusion The higher strength can be achieved for the board with uniform thickness of surface layer. The pellet-type resin with glass fiber is not good for the molding material of surface layer because of the fairly poor adhesion between resin and glass fiber. Apart from some problem, the compression molding method described in this paper must contribute to material recycle of various wastes of fibrous materials.

Fig.6 Fracture and maximum stresses. page_164 Page 165

Numerical Modeling Method of GFRP Laminate with Flexural Interphase

Tsuyoshi NISHIWAKI 1, Satoru HAYASAKI2, Kazuo KITAGAWA 3 and Hiroyuki HAMADA 2 1 ASICS Corp., 6-2-1 Takatsukadai, Nishi-ku, Kobe 651-2271 JAPAN 2 Kyoto Inst. of Tech., Matsugasaki, Sakyo-ku, Kyoto 3 Kyoto Municipal Inst. of Indus. Res, Chudoji, Shimogyo-ku, Kyoto Key words : Flexible Interphase, Finite Element Method, GFRP Laminate, Quasi-three-dimensional Model, Flexural Strength Abstract The interphase has an important role which supports the stress transmission between fiber and resin. Authors introduced the concept of flexural interphase, established the fabrication machine of prepregs. In this study a new numerical model of considered the fiber/resin interphase is proposed, in order to design the flexural interphase. Using the proposed model, the mechanical behaviors of unidirectional GFRP laminates with various flexural interphase thickness were simulated. Therefore the validity of the proposed model was confirmed. It was also concluded that the model was effective for the fiber/resin interphase designing in the GFRP composite structures. Introduction An interphase between fiber and resin in a laminated composite has the 3 dimensional region with a thickness value. The interphase has an important role, which supports the stress transmission between fiber and resin. The various mechanical properties of the laminated composites depends upon not only properties of fiber and matrix resin but also the interphase property. However an influence of this interphase upon the mechanical properties of the whole laminated composite has not been clarified. This indicates that the establishment of the designing guide of the fiber / matrix interphase leads to the widely development of the laminated composites. Authors have actively fabricated the new CFRP and ArFRP materials involving the interphase with new properties, investigated the mechanical properties of these new laminated composites. One of the most representative laminated composites is the GFRP laminate with the flexural interphase which has the higher toughness and less modulus than the matrix resin.[1][2] For the designing of the flexural interphase, an important key is the interphase thickness. It is clear that the contribution of the flexural interphase to the mechanical properties of the whole page_165 Page 166 laminated composites depends upon the flexural interphase thickness. In this study a new numerical modeling method considerd the interphase property is proposed. The model called as 'Quasi-three-dimensional model' is a construction of shell and 2 kinds of beam elements, which represent the fiber plate, interphase layer and interlaminar layer, respectively. Using the proposed model, the mechanical behaviors of unidirectional GFRP laminates with various flexible interphase thickness subjected to the 3-point flexural load are simulated in order to check the influence of the flexural interphase thickness upon the flexural moduli and strength. Also the damage propagation is simulated. Quasi-Three-Dimensional Model[3][5] In the application of finite element method to the damage propagation analysis of the multi-component composite materials, the consideration of the heterogeneity is an important factor. For the prediction of damage propagation of the laminated composite, the various local failure modes such as interlaminar delamination, transverse crack, fiber breakage and interphase fracture must be independently considered, because responses caused by the these failures are quite different. Fig. 1 shows the quasi-three-dimensional modeling method. Fig. 1(a) shows a cross-section of an each layer of the laminated composite. The cross-section has been usually regarded as a homogeneous object in the conventional numerical model. According to our concept, the cross-section is considered to be heterogeneous, as shown in Fig. 1(b). Namely the each layer is subdivided into 3 regions, fiber concentrated region, interphase region and resin matrix region, as shown in Fig. 1(c). In the fiber concentrated region, the fiber coated flexible interphase is assumed to be a hexagonal closet packed arrangement with Vf = 90.7 %.[6] In Fig. 1(c), t2 is given from Eq. 1.

Here, r and ti indicate the reinforcement fiber diameter and interphase thickness, respectively. Also, t1 is calculated from Eq.2.

Here, t0 and Vf0 is the thickness and fiber volume fraction of the prepreg, page_166 Page 167 respectively

Fig. 1 Basic concept for new numerical modeling The proposed model is constructed by using orthotropic shell and 2-type isotropic beam elements. The shell elements correspond to the fiber-plate with t1 thickness in Fig. 1(c). The 2 kinds of isotropic beam elements correspond to the flexural interphase and remaining resin layers with Vf = 0. Then shell elements are connected with 2 kinds of beam elements in the thickness direction in order to express the interphase and interlamina, as shown in Fig.2. The beam elements have various cross-sections depended upon their positions. The sum of beam cross-sectional areas in the same interlamina is set to be same as the whole laminate area.

Fig.2 Numerical modeling example Analytical Procedure The GFRP plates analyzed are set up to from an 6-layered unidirectional prepreg with the various flexural interphase thickness. The overall length of the model is 40mm, width is 15mm and the nominal thickness is 2mm. The span length is 32mm. Here x-, y- and z-directions indicate the longitudinal, width page_167 Page 168 and thickness directions, respectively. The quasi-three-dimensional model used is shown in Fig.3. The load are applied by a load incremental method. Since the 3-point flexural damage propagation analysis is a nonlinear problem involving the stiffness reduction due to the damage propagation, the stress components in all elements are calculated for every load increments.

Fig. 3 Numerical model used in this study In the model, three kinds of local failure modes, transverse crack, interlaminar delamination, and interphase fracture appear. The transverse crack is represented by the fracture in shell elements. Tsai-Hill criteria is applied. In this analytical procedure, failure of shell elements by compressive stress is not considered. After shell elements failed, the elastic modulus in the transverse direction reduces the modulus of the matrix resin in order to consider the stress redistribution. The interlaminar delamination and interphase fracture are represented by the yielding of 2-type beam elements, respectively. In both the yielding, Von-Mises criteria are applied. After these yielding, the properties of 2 kinds of beam elements are changed as shown in Fig.4. This yielding is mainly caused by the out-plane shear stress. Then the propagation of yielded beam elements is continuously investigated. The 3-point flexural experiments of these GFRP plates were carried out in order to check the validity of the above analytical procedure Fig. 5 shows an illustration of the fabrication machine of

Fig.4 Beam element properties used in this analytical procedure page_168 Page 169 prepregs. Using this machine, the thickness and of the interphase and surface treatments can be changed. Glass fibers were covered with flexible interphase

Fig.5 Fablication Method of Unidirectional Prepreg Sheet with Flexible Interphase

by immersing strand in MEK solvent. These glass strands were impregnated with convensional epoxy resin ( #828, Koei Chemical Co.Ltd.) mixed and were prepregnated by heat treatment. The thickness of flexible interphase are controlled by the concentration of flexible epoxy resin( #550, Koei Chemical Co.Ltd.) in MEK solvent. The fiber volume fraction Vf of the specimen was constant 61%. The concentrations of #550 tested are 0, 2, 5 and 7 wt%. Here 0 wt% indicates the normal GFRP prepregs. With increasing the concentration of #550, the thickness of interphase formed between fiber and resin matrix is also increasing. Results & Discussions Fig.6 shows both the analytical and experimental results in case that #550wt% = 2 and 7. As already mentioned, the flexural interphase thickness represented by 2 wt% is smaller than 7 wt%. In this figure, reasonable agreement is evidenced between both the results for stress-deflection relationships. In other cases, reasonable agreements can be also confirmed Fig.7 shows the damage propagation maps of 0 wt%. In this figure, bold lines page_169 Page 170 indicate yield beam elements. In all quasi-three-dimensional analyses, the fracture of shell elements cannot be predicted. In the standard unidirectional GFRP laminate, 0wt% any local failures cannot be appeared till 64kgf loading, as shown in Fig.7(a). Under the 65 kgf loading, beam element yielding corresponded to interlaminar delamination are occured in 3rd/4th interlamina. This position is the neutral surface of this model. In the all yielded beam elements, the xz shear stress component is much larger than any other components.

Fig.6 Load-deflection curves for GFRP laminates with 2 and 7 wt% flexible interphases. As already reported,[7] this indicates that the interlaminar delamination is mainly

Fig. 7 Damage propagation maps of quasi-three-dimensional model with no interphase beam page_170 Page 171 caused by the out-plane interlaminar shear stress. With increasing flexural load, the yielding spreads in the neutral surface. On the other hand, in the model with 7 wt% flexural interphase, almost interphase beam elements yielded under 10 kgf flexural loading, as shown in Fig.8(a). With increasing the load, yielding of interlaminar beam elements appeared. With further increasing load, yielding of interlaminar beam elements spread 3rd/4th interlamina. The evaluational method of the critical 3-point flexural strength was also investigated for the quasi-threedimensional model. The proposed flexural strength law is based on the occurrence of interlaminar delamination. When the delamination occurred in laminates, its flexural rigidity declined. Particularly, in case that the delamination goes through the width direction, the reduction in the flexural stiffness is large. Based on the above analytical results, flexural strength was defined by the load level corresponding to initial yielding of all interlaminar beam elements along the width direction. Table 1 shows the comparison between the analytical and experimental critical strengths. From this table, it is confirmed that the control of the flexible interphase thickness is effective for the application of the GFRP laminate to the structural components.

Fig.8 Damage propagation maps of quasi-three-dimensional model with 7 wt% interphase beam elements page_171 Page 172 Particularly, it must be noted that the flexural strength of GFRP laminate with 2 wt% is much larger than that with 0wt%. Table 1 Comparison of flexural modulus and strength in analytical and experimental results Concentration of Flex Flex. Modulus Flex. Strength Interphase [GPa] [MPa] [wt%] FEM Exp. FEM Exp. 0 40.3 32.0 658 637 2 36.0 37.5 875 851 5 21.2 24.3 576 514 7 14.9 17.2 482 510 Conclusions The mechanical behaviors of GFRP laminates with various flexible interphase under 3-point flexural load were descrived. The numerical simulation method based on the quasi-3-dimensional model was established. It was found that the stiffness reduction due to the damage propagation could be simulated. The analytical strength defined as the 3-point flexural stress level where interlaminar delamination goes through width direction entirely and experimental

strength were compared, the good agreements could be confirmed. Using the proposed model, the influences of the flexible interphase thickness upon the modulus and strength of GFRP laminate subjected to the 3-point flexural load were predicted. Therefore it was concluded that the interposition of 2wt% interphase between fiber and matrix resin stiffened GFRP laminate. References 1. K.Kitagawa, S.Hayasaki, et.al, Proc.5th Int. SAMPE, (1997), pp.765770 2. S.Hayasaki, K.Kitagawa, H.Hamada, J.Mater.Sci. Lett., (submitted) 3. T.Nishiwaki, et.al, JSME Int.J., 38,1, (1995), pp.1620 4. T.Nishiwaki, et.al, Compos.Struct., 32, (1995), pp.635640 5. T.Nishiwaki, et.al, Compos.Struct., 32, (1995), pp.293298 6. D.Hull, An Intro. to Compos. Mater., Cambridge Univ.Press, (1992) 7. T.Nishiwaki, et.al, Compos.Struct., 25, (1993), pp.6167 page_172 Page 173

CERAMIC/METALS/POLYMER HYBRID COMPOSITES page_173 Page 175

Design and Applications of Metal/FRP Hybrid Structures Patrick Kim Shonan Institute of Technology 1-1-25 Tsujido Nishikaigan, Fujisawa 251, Japan Key words: Hybrid structure, lightweight design, optimization, failure modes, cost analysis Abstract This paper discusses the design, production, performance, cost, and applications of metal/FRP hybrid structures. Hybrid structures combine the relatively high shear stiffness and ductility of metal alloys with the high specific stiffness, strength, and fatigue properties of FRPs. Their performance is comparable to all-CFRP structures at a cost that is competitive with all-metal structures. Representative design examples of flexural and torsional structural elements are given. In both cases, a significant weight saving over equivalent aluminum structures is achieved. The relatively simple methods for producing hybrid components circumvent the need for the complex and expensive equipment that is used for advanced composites processing. The various advantages and outstanding problems of hybrid structures as well as existing and potential applications are discussed. Introduction Fiber-reinforced polymers (FRPs) are popular in many fields of engineering for the design and production of stiff, lightweight structures. However, the materials and facilities required to process structural FRPs makes their use too expensive for many applications. Furthermore, the multiple fiber orientations in FRP structures translate into an ineffective use of the outstanding fiber properties. As an alternative we propose a metal/FRP hybrid concept that makes appropriate use of the respective advantages of both classes of materials [1]. In this concept, the aluminum contributes shear stiffness, ductility, and ease of forming, while the FRP contributes axial stiffness and fatigue resistance (Fig. 1). Hybrid structures can be produced as an adhesively bonded assembly of semi-finished sections, such as extruded aluminum and pultruded FRP sheets, thus reducing the difficulty and cost often associated with manufacturing FRP parts with complex shapes.

Fig. 1. Advantages and drawbacks of metal, FRP, and hybrid structures. page_175 Page 176 Design of Hybrid Structures Materials Selection For hybrid structures to be competitive on a mechanical basis when low weight is desired, the composite must have a specific stiffness (stiffness/mass) at least equal to that of the metal it is supposed to replace. One can use an FRP with a lower modulus than the metal for the hybrids, as long as this deficit is compensated by a sufficient mass advantage. Combinations of steel or aluminum with GFRP are thus ruled out unless other functional considerations such as electrical insulation or corrosion resistance have a high priority. Carbon or graphite and aramid fibers are considered the most appropriate types as they offer a large range of stiffnesses and strengths. Care must be taken in design, as high-modulus fibers (e.g. HM carbon) do not necessarily achieve a higher weight saving than high-strength (HT) fibers. HT fibers can be a better solution in particular when strength as well as stiffness requirements are active. One central idea in hybrid design is keeping fiber orientations to a minimum and the fibers aligned with strongly oriented dominant loads. Parts of a structure that are subject to complex stress states or that undergo strong directional load variations should be isotropic and therefore made of metal. The most effective use of FRPs is achieved by using unidirectional (UD) materials, and although two fiber orientations may be necessary in some cases, a single orientation is preferred. The Metal/FRP Interface Hybrid structures derive their principal advantages from combining two or more classes of materials with greatly differing properties and degrees of isotropy. This combination however also raises problems of stress transfer between the metal and the FRP. Stiff bonds often fail in a brittle manner, as they concentrate stresses in a small volume [2]. A compliant and toughened adhesive should therefore be used for the bonded interface. The durability of the bond is increased by using an adhesive with minimal aging and moisture-related property degradation. Stresses are generated at the interface by external mechanical loads and processing or service-related thermal loads. In order to keep the residual processing stresses to a minimum, the hybrid structure should be processed at or near the service temperature. This may require using a low-temperature curing thermoset adhesive. Stiffness and Ultimate Behavior The overall stiffness of hybrid structures can be calculated as the sum of the contributions from the metal and the FRP. The contribution of the adhesive layer can be neglected as long as this layer is comparatively thin and end-effects are small. For example the flexural stiffness of hybrid, as well as composite beams, is approximated by simple beam theory:

where the subscript i denotes the contribution of parts with different Young's moduli E and corresponding bending moment of inertia I. For the case of torsional elements, E is replaced by the shear modulus, G, and I by the torsional moment of inertia J. This simplified approach slightly overestimates the stiffness of hybrid tubes under torsion [3], but it is very accurate for hybrid beams [4] under flexure. When Ei is expressed as a function of strain, Eq. 1 predicts anelastic behavior such as the flexural stiffness of hybrid beams after yielding of the aluminum with fair accuracy. The occurrence of first failure is determined from the strains calculated from the beam-bending or other appropriate load-strain equation using Eq. 1 for the overall stiffness and the constitutive equations of the individual materials [5]. Additional equations are required for structural failure related to local and global elastic instability. Bondline failure is mainly of concern in the case of

torsional hybrid tubes, as discussed page_176 Page 177 below, and when flaws are present in the bondline. The analysis of bondline flaws requires a special attention, as it can lead to local buckling or fracture propagation that undermine the overall structural performance of the hybrid structure [6]. Connections and Load Induction Connections and load induction are a big problem with all-composite structures, as the materials exhibit little plasticity and thus accommodate high local stresses poorly. This problem is mitigated in hybrid structures by an appropriate design in which the connections and concentrated load areas are assigned to the metal parts of the structure. Connections can then be made using conventional mechanical or thermal joining techniques. Areas subject to end effects, such as the transition from all-metal to hybrid sections should be located in moderately stressed parts of the structure. Stress concentrations can be avoided by an adapted design of the geometry of the interface. Structural Optimization Hybrid structures offer a larger design space than structures made of a single isotropic material and thus some interesting possibilities for structural optimization, in particular when the objective is mass reduction. Optimization of hybrid structures begins with an appropriate choice of structural layout. The layout must match the advantages of metals and FRPs to the various stress conditions within the structure. A representative example is the optimization of an aluminum/CFRP hybrid box beam using a non-linear constrained optimization approach [5]. The objective function of mass is minimized, subject to constraints given by the stiffness and strength requirements, failure mechanisms, and arbitrary limiting values on certain dimensions. In some cases, a compromise between mechanical optimization and manufacturability is necessary. In summary, the effectiveness of hybrid structures depends on the use of unidirectional high-modulus or high-strength FRP lamina, assembly at or near service temperature, a compliant, durable adhesive bond between the metal and the FRP, and structural optimization. Production, Performance, and Cost: Illustrative Examples Two examples are given to illustrate the points made above: hybrid flexural beams and torsional tubes. The representative properties of the materials used for these cases are given in Table 1. The moduli of angle-ply FRP layers are calculated by classical lamination theory from the UD properties. Table 1 also gives representative current unit cost for the processed materials. These include processing costs for the methods typically used for these materials. Table 1. Aluminum and FRP properties used for the calculations Property Aluminum GFRP HT CFRP HM CFRP Fiber content [vol%] -60 50 50 Elastic modulus, UD FRP [Gpa]* 70 / 70 / 70 45 / 16 / 5.0 120 / 15 / 4.3 200 / 15 / 4.4 Poisson's ratio, UD 0.34 .27 0.22 0.22 Yield stress or strength [Mpa]* 300 /300 /200 1100 / - /110 1500 / - /110 1000 / - /110 Density [g/cm3] 2.70 1.90 1.55 1.65 Approx. unit cost [US$/kg] 6 15 40 60 * longitudinal / transverse / shear page_177 Page 178

Fig. 2. Representative cross-section of metal, composite, and hybrid box beams.

Hybrid Flexural Beam Description and Fabrication Aluminum/CFRP hybrid beams give an interesting tradeoff between mechanical performance, ease of fabrication, and cost [7]. These hybrid beams consist of an aluminum box profile with pultruded unidirectional (UD) FRP adhesively bonded to the flanges. The processing consists of pultrusion of the composite, extrusion of the aluminum, and surface preparation such as abrasion and degreasing. All are standard, low-cost methods. Profiled boards and clamps are the only necessary equipment for assembly. A comparative all-FRP beams typically consists of two layers of filament-wound, angle-ply FRP at about 45°, with an intercalated UD layer in each flange. Fig. 2 shows typical cross sections of metal, all-FRP, and metal/FRP hybrid box beams. These beams were optimized to achieve minimum mass for a given stiffness or strength requirement. Behavior A typical force-deflection diagram for aluminum and hybrid beams is shown in Fig. 3. Bonding a thin layer of CFRP to an aluminum beam results in a large increase in stiffness and strength with very little increase in mass. After initial yielding and subsequent failure of the CRFP, the hybrid beam is still capable of carrying a significant residual load. The residual load-carrying capacity is determined by the design. This ductility is essential in particular for civil engineering structures. Table 2 gives a comparative overview of the performance of the different beams. In a design governed by strength requirements, a GFRP beam is 21% lighter, while CFRP beams using HM and HT fibers are 6669% lighter, respectively, than the minimum-weight aluminum beam with the same Table 2. Comparison between aluminum, FRP, and hybrid beams giving a flexural moment resistance M=60kNm and shear-to-moment ratio V/M=.001/mm. Beam type Beam depth EIeq [x1011N/mm] Mass [g/m] Cost [US$/m] [mm] Aluminum 200 7.02 4902 29.41 GFRP 167 2.63 3857 70.03 HT CFRP 126 1.93 1534 65.72 HM CFRP 135 4.01 1650 103.33 Al / HT-CFRP hybrid 141 4.92 2743 49.30 Al / HM-CFRP hybrid 122 4.25 2163 53.96 page_178 Page 179

Fig. 3. Force-deflection diagram of aluminum and hybrid beams

Fig. 4. Moment-twist diagram of hybrid tubes with different bond strengths ultimate flexural moment resistance. The hybrid beams using HT-CFRP and HM-CFRP are 44% and 56% lighter, respectively, than the aluminum beam. The weight saving is slightly less for a design controlled by the stiffness requirement only. However, as the percentage of CFRP in the beam is lower, the total cost is lower as well. In this example the cost saving is 25 to nearly 50%. The optimized hybrid beams have a 3040% smaller depth than their aluminum counterpart. This fact is important for applications such as transportation, where a gain in usable space can be valuable. Partly as a result of this, the CFRP and hybrid beams have a lower flexural stiffness than the aluminum beam for a given strength requirement. However, due to their lower mass, their specific stiffness (EI/mass) is significantly higher than for the aluminum beam. Finally, the critical failure modes of the hybrid beams remain the same over a wide range of aluminum and CFRP moduli and limit stresses [7]. Failure of the bondline is not among the initial failure mechanism under three or four-point bending. Bondline failure may however occur when the ends of FRP lamellas are in highly stressed zones or as a result of other mechanisms such as local buckling of the flange. Hybrid Tube under Torsion Description and Fabrication The joint between a CFRP tube and metallic end fittings is a weak link in which the stress concentrations can lead to a premature failure [8, 9]. A solution to this problem, as well as a way to reduce the mass of torsional tubes, is an aluminum/CFRP hybrid tube design consisting of an aluminum core tube with a reduced center section overlaid with CFRP (Fig. 5). Transition stresses are kept low by an adapted geometry of the transition between the all-metal ends and the hybrid section. The overall properties of the tubes can be tailored to meet combined shear and flexural requirements such as stability against dynamic buckling during rotation. For example a combination of torsional and flexural properties practically identical to that of an aluminum shaft is obtained using angle-ply CFRP layers at ±45° and an outer layer at 0°. page_179 Page 180

Fig. 5. Geometry and layup of the hybrid tube Adapted processing methods for the hybrid tubes are table rolling or tape-laying. We fabricated hybrid tubes by laying prepregs up and curing them on the aluminum core. The splice in the CFRP was made parallel to the fibers, and did not constitute a source of weakness in the tube. This approach avoids fiber cross-overs that reduce the stiffness and strength of the tube, as in filament-wound tubes [10]. A thin adhesive film with a glass carrier fabric

provides a strong, tough bond between the aluminum and the CFRP [11] as well as an insulating layer that reduces corrosion between the aluminum and the carbon fibers. This hybrid tube design in particular avoids the problem of FRP damage due to mandrel extraction after filament winding. Behavior An exploratory study has compared a hybrid tube with an aluminum tube having the same torsional stiffness [3]. The simplified design equation (Eq.1) overestimated the actual torsional stiffnesses of the hybrid tubes by 1725%. However, as the hybrid tubes weighed 37% less than their aluminum counterpart, the normalized stiffness JG/M of the hybrids was 1527% higher than that of the aluminum tubes. The moment-twist relations for the hybrid tubes are shown in Fig. 4. Two types of failure were observed. Sample 1 underwent a sudden drop in load after a rise to a high peak, but continued to carry a residual load. Sample 2 failed progressively, reaching a constant load value after a sharp inflection of the loading curve. This second type of failure is likely due to residual thermal stresses [12], which lead to debonding before the CFRP could fail in shear. In both cases, the CFRP delaminated from the aluminum starting at the transition between the all-metal and the hybrid section. This shows a need to optimize the geometry for a smoother stress transfer in the transition zone. The load-bearing capacity at large deformation corresponded simply to that of the aluminum part of the tube. However, it should be noted that the tubes underwent anelastic deformation before the CFRP delaminated, and were thus loaded higher than would be the case in service. Table 3. Properties of the aluminum and hybrid tubes (calculated values in parentheses) Sample Stiffness J G Mass M [g/m] JG/M Tdam(a) [x103Nm2 [Nm2/g·m] [Nm] All-aluminum 2.44 (2.52) 774 3.15 16.5 Hybrid, sample 1 1.79 (2.40) 495 3.62 10.7 Hybrid, sample 2 1.94 (2.35) 486 3.99 9.1 (a) Yield for aluminum, other departure from linearity for hybrid. page_180 Page 181

Fig. 6. Break-even price of CFRP as a function of aluminum price for CFRP and Al/CFRP hybrid beams, for a strength-based design, (price of angle-ply) - (price of UD = $5/kg. Cost Analysis Factors such as the choice of material preform, processing method, added value in the form of structural details, and production volume affect the relative cost of all-metal, all-composite, and hybrid structures. For example, if a large production volume brings down the cost of structural composites as a result of increased use of CFRP in civil engineering [13], all-CFRP and hybrid structures would become more cost-competitive. Fig. 6 shows the price of processed CFRP for which CFRP or hybrid beams would have the same price as an equivalent aluminum beam. At a unit price of aluminum of $6/kg, an all-CFRP would be competitive at a CFRP cost under $12.50/kg. The hybrid beam is cost-competitive at a CFRP cost of $27.50/kg, a much more realistic figure that could be achieved in the near future. For the time being, aluminum/CFRP hybrid beams provide a significant weight savings over aluminum at a much lower cost than when using only CFRP [7]. Applications of Aluminum/CFRP Hybrid Structures

Existing and Potential Applications We consider hybrid structures to have a good potential as an alternative to all-composite structures or as replacement for all-metal designs for simple flexural and/or torsional structural elements. Extruded aluminum and pultruded FRP sections are readily available at a relatively low cost from a large number of manufacturers. The simplicity and low cost of the equipment necessary to process them makes their manufacture easily accessible with a relatively small investment, and thus could be appealing to companies with little experience with composite materials. Some of the principal current applications of aluminum/FRP hybrids are: beams for lightweight civil engineering structures [4] flatbed trailer for heavy road transports compressed gas tanks (aluminum with GFRP [14], AFRP, and CFRP [15]) safety enhancement of tanker trucks (aluminum/AFRP [16]) Applications that could be realizable in the near future are: high-speed 2-D cutting installations and 2-D or 3-D robot arms bus frames and train wagons floor beams in airplanes, trains, or busses lightweight upper structures for high-speed ships page_181 Page 182 Limitations and Outstanding Problems The principal limitations and problems with hybrid structures arise from the durability of the interface and the thermal property mismatch. The resistance of the polymeric adhesive bondline to moisture and corrosion is a critical issue that needs to be investigated more extensively. Although processing stresses can be limited by using low-temperature curing resin materials, the thermal expansion mismatch between the aluminum and the CFRP is a source of potentially severe internal stresses or warpage during service [17]. This limits the use of hybrids to applications with relatively small temperature fluctuations. More data is needed on the static and impact fracture resistance, damage tolerance, environmental resistance, fatigue, and vibration of aluminum/CFRP hybrids before these can be used with confidence in a broader range of applications. Conclusions This paper has discussed the design and applications of metal/FRP hybrid structures. Hybrid structures combine the shear stiffness and ductility of metals with the high specific stiffness and fatigue properties of FRPS in a mechanically appropriate way. The combination makes hybrids lightweight, fatigue resistant [18, 19], and ductile [5]. The weight reduction afforded by hybrid structures is not as great as with CFRP, but the manufacturing and joining are simpler and significantly cheaper. Hybrid structures can in many cases be made of inexpensive semi-finished products, such as pultruded composites, that are simply bonded together using adhesives. When needed, reinforcement can be retrofitted cost-effectively with modest equipment [20]. Hybrid structures are a potentially competitive alternative to metal and FRP structures in a range of applications in mechanical, transportation, and civil engineering. References 1. P. Kim, EMPA report No.134'325, Swiss Federal Laboratories for Materials Research and Testing, (1991). 2. R.D. Adams et al., Structural Adhesive Joints in Engineering, 2nd ed, London, Chapman & Hall, (1997). 3. P. Kim and T. Tanimoto, 22th JSCM Composite Materials Symp, Fukuoka, Japan, Nov. 67, (1997), p. 68. 4. P. Kim and H. Meier, EMPA report No.126'537/2, Swiss Federal Lab. for Materials Research and Testing (1991). 5. T.C. Triantafillou et al., International Journal of Mechanical Science, 33 (1991), p. 729. 6. P. Kim, P., 22th JSCM Composite Materials Symp, Fukuoka, Japan, Nov. 67, (1997), p. 70. 7. P. Kim, to appear in Applied Composite Materials (1998).

8. K.S. Kim et al., Composite Structures, 21 (1992), p. 163. 9. J.H. Choi and D.G. Lee, Journal of adhesion, 44 (1994), p. 197. 10. D.C. Jegley. and O.F. Lopez, AIAA Journal, 30 (1992), p. 205. 11. P. Kim and T. Tanimoto, 6th SIMS Symposium on Interfacial Materials Science, Osaka, Japan (1997), p. 527. 12. H. Orsini and F. Schmit, Journal of Adhesion, 43 (1993), p. 55. 13. S. Ashley, Mechanical Engineering, 118 (1996), p. 76. 14. L. Varga et al., Composites, 26 (1995), p. 457. 15. J.M. Lifshitz and H. Dayan, Composite Structures, 32 (1995), p. 313. 16. Erhöhte Unfallsicherheit für Tankfahrzeuge, No. 27, Ciba-Geigy Kunststoff-Aspekte, 1991. 17. C.T. Lin et al., Composites, 25 (1994), p. 303. 18. R.J. Bucci et al., Aluminum Alloys-Contemporary Research and Applications, Academic Press (1989), p. 295. 19. C.T. Lin et al., Composites, 22 (1991), p. 135. 20. U. Meier et al., in Alternative Materials for the Reinforcement and Prestressing of Concrete, Blackie Academic & Professional (1993), p. 153. page_182 Page 183

Mechanical Forming of Aluminum Matrix Composites H.J. McQueen and E. Evangelista Mech. Eng., Concordia Univ., Montreal, H3G 1M8, Canada Mechanics, Univ. of Ancona, I-60131, Italy Abstract Traditional mechanical forming processes are an economical route for secondary shaping of fully dense billets continuously cast from aluminum matrix particulate composites initially mixed in the molten alloy. Their increased strength, along with augmented modulus and wear resistance, is associated with diminished ductility at 20°C. In consequence, hot fabrication is the most suitable route to produce components that are free of particle decohesion and cracking. The hot workability of A356, A359, 6061, 2618 and 7075 matrix alloys with 10, 15 and 20% of 1015 mm particles, either of Al2O3 or of SiC, were determined over the range 300 to 540°C at rates of 0.1 to 10 s-1. Ductilities in torsion equivalent to 90% reduction were obtained between 400 and 500°C. Constitutive equations providing the temperature and strain rate dependence of the flow stress were derived. Both extrusion and forging have been modeled and the forces for forming were predicted and compared to those of the base alloys. page_183 Page 184

Effect of Plasma Treatment on Surface of Glass Fiber for Plastic Based Composites A. Nakahira1, Y. Suzuki2, S. Ueno3, H. Akamizu1, K. Kijima1, S. Nishijima3 1Dept. of Chem. and Materials Tech., Kyoto Institute of Technology, Gosho Kaido-cho, Matsugasaki, Sakyo-ku, Kyoto 606-8585, Japan 2National Industrial Research Institute of Nagoya, 1-1 Hirate-cho, Kita-ku, Nagoya 462, Japan 3ISIR, Osaka University, 8-1 Mihogaoka, Ibaraki, Osaka 567, Japan Keywords: surface / glass fiber / plasma treatment / microstructure / interface

Abstract Glass fiber and sheets were treated by the low temperature argon plasma process in order to improve the surface properties for glass fiber reinforced plastic (GFRP). Especially some surface properties, for example, roughness and wetting behaviors, of glass fiber and sheet were studied before and after plasma treatment. Interfacial strength between the plasma treated fiber and plastic was measured. The effect of plasma treatment on surface of glass fiber and sheet was mainly evaluated and discussed. Introduction Better glass fiber and ceramics fiber as reinforcement are necessary to develop the high performance fiberreinforced ceramic matrix composites (FR-CMC) and fiber-reinforced plastic matrix composites (FR-PMC). However, in order to apply FR-CMC and FR-PMC into structural component under sever condition the further improvement of the mechanical properties, for example, interfacial strength, is also demanded1. Actually, surface treatment techniques for materials, especially polymer, are widely studied to improve interfacial properties. Various surface treatment techniques were attempted with the aim of increasing the interfacial bonding between the fiber and matrix. Especially, glow discharges (cold plasma) have been studied for surface modification of materials, because these dry process have the advantage from the view point of circumstance pollution, compared to wet process, such as, coupling agent treatment2. The purpose in this paper is to examine the effect of argon plasma treatment on glass surfaces by microscopic observation by SEM and to evaluate the interfacial shear strength between the glass fiber and epoxy with the pull-out test of fiber embedded with epoxy resin. The effect of plasma treatment on surface modification and mechanical properties of glass fibers and sheet was mainly discussed. Experimental Procedure E-glass sheet and E-glass fiber were used in the present study. Glass sheets were cut by diamond blade with the dimension of 10´10´5mm and polished mirror-like with diamond paste below 0.1mm. A strand of E-glass fiber was consited of 200 fibers with the diameter of 9mm. The sizing materials of surface of glass fiber was removed by thermal treatment. Fig. 1 illustrates the plasma equipment system used in the present work, in which the system consists of Pyrex reaction chamber, gas inlets system and vacuum system, operating at the R.F. frequency of 4MHz. The chamber was evacuated to 2Pa by using mechanical pump and then purged page_184 Page 185 well with argon before plasma initiation. Plasma power was varied from 100W and 200W in this experiment. The argon plasma was initiated at pressure of 60Pa, 90Pa, and 120Pa at required power level. Typical plasma treatments time was 120sec under conditions above-mentioned. Glass samples were usually supported on BN jig and mounted in the center of chamber. SEM observations of glass sheets and fiber were done before and after plasma treatments. Wettability measurement of surface of E-glass sheet before and after plasma treatments was evaluated by measuring the contact angle with using the equipment (Face contact-angle meter: Kyowa-Kagaku Co.). Contact angle of sample was measured only for E-glass sheet samples before and after the various plasma treatment of different power levels with distilled water at ambient temperature. E-glass fiber was mounted on aluminium-plate (10´10´1mm) using epoxy resin (Araldite, Ciba-Geigy). Pull-out test, which was shown in Fig.2, were performed for measuring the interfacial shear strength of E-glass fiber/epoxy sample using Instron-type at a cross-head speed of 1mm/min for interface shear strength evaluation between glass fiber and epoxy resin 3. Interfacial shear strength, t, was calculated by the following equation:

where P : load, d : fiber diameter, 1 : fiber length embedded with epoxy resin, and N : number of fiber. The embedded fiber length, 1, was measured by optical scope. Fracture surfaces of samples were observed by SEM.

Fig. 1 Illustration of plasma equipment system

Fig. 2 Schematic drawing of pull-out test. page_185 Page 186 Results and Discussion Fig. 3 shows the relationship between the contact angle and gas pressure during argon plasma treatment. As the performance of plasma treatment was well-known to be dependent on the power, gas pressure and flow rate, in this experiment the contact angle values were also varied with gas pressure and power of plasma treatment4. The as-polished E-glass sheet showed the contact angle of approximately 66.2°. Furthermore the surface of E-glass ground with 400# diamond wheel indicated very high contact angle of aproximately 80°, though approximately 30mm in roughness. However the contact angle of E-glass sheet treated by argon plasma decreased with the gas pressure during plasma treatment with 100W. The high power plasma treatment of 200W showed also lower contact angle than un-treated, though the contact angle increased with high gas pressure. Wettability of surface of E-glass sheet was enhanced by argon plasma treatment, compared with that of as-polished E-glass sheet. Therefore the plasma treatment makes the surface of E-glass more hydrophilic, as pointed out by V. Krishnamurthy et al 5. It is thought that this decrease of contact angle was caused not only by the remove of contamination, but also the modification of surface chemistry of E-glass sheet by plasma treatment.

From the results of SEM observation for E-glass fiber after plasma treatment, which were not shown here, plasma treatment in the present work induced the damage of surface of E-glass fiber though the change was subtle in comparison to polymer fiber. The surface of E-glass fiber was a little damaged by the plasma treatment of 100W power with high gas pressure. When gas pressure was relatively high, 200W plasma treatment induced the severer damage of fiber surface. The microscopic change of roughness and damage was confirmed on surface of E-glass fiber by SEM observation. The effect of gas pressure on the interfacial shear strength with the plasma treatment of 100W and 200W power was shown in Fig.4. The interfacial shear strength of sample with 200W plasma treatment showed the maximum value at gas pressure of 60Pa and decreased with gas pressure over 90Pa. On the other hand, samples plasma-treated with plasma power of 100W tended to show the increase in interfacial shear strength with gas pressure. The changes in interfacial shear strength were thought to be attributable to wettability and optimum increase in roughness on the surface of E-glass fiber by plasma treatment.

Fig. 3 Relation between the contact angle and gas pressure during argon plasma treatment. page_186 Page 187

Fig.4 Effect of gas pressure on the interfacial shear strength with the plasma treatment. Summary The effect of argon plasma treatment on glass surfaces by microscopic observation was examined by SEM. The roughness and damage on of E-glass fiber surface was dependent on the plasma power and gas perssure. Wettability of surface of E-glass sheet was decreased by argon plasma treatment. Although the rough surface on micron-meter scale tends to have higher contact angle, the increase in surface roughness of E-glass on nano-meter level were found to be not reflected in contact angle but be associated with modification of surface chemistry of E-glass sheet on the control of contact angle by plasma treatment.

The plasma treatment in the persent work increased the interfacial shear strength between glass fiber and epoxy resin. From these results, it was found that the argon plasma treatment of glass fiber surface produced the improvement of wettability and optimum increase in roughness on the surface of E-glass fiber and resulted in the improvement of the interfacial shear strength between E-glass fiber and epoxy. References 1. A. Rose and J. T. A. Pollock, J. Mater. Sci., 23, 1752(1988). 2. R. Foerch, N. S. Mcintyre, and R. N. S. Sodhi, J. App. Poly. Sci., 40, 1903(1990). 3. A. Nakahira, Y. Suzuki, S. Ueno, H. Akamizu, K. Kijima, S. Nishijima, contributed to J. of Sci. Eng.of Composites. 4. H. Yasuda, in ''Thin Film Processes", Edited by J. L. Vossen and W. Kern (Academic, New York, 1978), p.361. 5. V. Krishnamurthy and I. L. Kamel, J. Mater. Sci., 24, 3345(1989). page_187 Page 189

DESIGN AND APPLICATIONS page_189 Page 191

Hierarchical Layerwise Higher-Order Finite Elements for Laminated Composite I. Kimpara, K. Kageyama and K. Suzuki* Department of Naval Architecture and Ocean Engineering, the University of Tokyo 7-3-1 Hongo, Bunkyo-ku, Tokyo 113-8656, Japan Abstract As practical applications of advanced composite materials (ACM) are increasing in number and scale, accurate and proper numerical analysis models for the structures composed of such advanced materials are becoming more and more important. Ironically, the unique and superior engineering properties of ACM compel their numerical modelings more complicated and skillful than those of the conventional materials. Especially, laminated composite material structures (composite laminates) possess various kinds of mechanical, geometrical and material inhomogeneities in every phase of the analysis, ranging from fiber/matrix interfaces to adhesive joints of structural components, as presented in Fig.1. Furthermore, in practical problems, these inhomogeneities should be treated at a time as shown in Table 1. In particular, in structural analysis and design, the inhomogeneity in "Laminate-like Phase" must not be overlooked. Therefore, an accurate and proper numerical model for composite laminates should take their inhomogeneities into consideration efficiently. In this study, a flexible and versatile finite element (FE) model is proposed, which answers these modeling demand correctly.

Figure 1 Comparison of composite materials and homogeneous materials with respect to inhomogeneity *JSPS Research Fellow page_191 Page 192 Table 1 Possible combinations of different inhomogeneities in composite laminates combination of inhomogeneity levels examples matrix crackings meso-mechanics homogenization method Structural Member + Laminate-like delaminations CAI failure of sandwiches interfacial stresses interlayer properties

free-edge effects overlayed joints So far, as applicable numerical models for composite laminates, two different kinds of models have been proposed in the literature, that is, Equivalent-single-layer (ESL), higher-order theories (HOTs) and their finite elements (FEs); Multi-layer (ML) theories and their FEs. The family of ESL theories and FEs is probably the most dominant mainly because they have been well tested, robust and have been proven to produce satisfactory predictions of global responses of thin to moderately thick composite laminates. Another reason for their popularity is that degrees of freedom (DOF) in ESL-based theories and FEs, whether classical or refined, are independent of the number of laminae, which means drastically decreased DOF in number, and at the time of inferior computational circumstances, only such a small scale calculation was acceptable. One of the most successful ESL-based HOTs is perhaps Reddy's third-order theory (HOT of Reddy) [1]. This theory assumes cubic variations of in-plane displacement components over the thickness while suppressing excessive DOFs via the traction-free conditions on the exterior surfaces of laminates. However, ESL-based HOTs do not precisely model multi-layered kinematical configurations of the actual laminates. Besides they don't provide meaningful predictions of transverse stress (and interlaminar stresses) unless pointwise re-calculations by means of the 3-D equilibrium equations are carried out. In order to improve these drawbacks of ESL, the family of multi-layer (ML) theories and their FEs have been proposed in the literature. Roughly three different types of ML models have been proposed in accordance with variational considerations they adopt, that is, (a) assumed stress hybrid elements [24]. This type of element assumes stress distributions for each layer by introducing stress parameters so that equilibriums of stresses at the interlayers and the surfaces are satisfied a priori. Especially, a series of recent works by Hoa and his colleagues [4] is prominent. (b) Other mixed variational method [5]. (c) Displacement-based layerwise theories [68]. In this study, the displacement-based ML models will be called "Layerwise" models for distinguishing it from the other ML models. These ML theories and their FEs will make a mainstream in the numerical analysis and design for composite laminates, because they can provide sufficiently good results for both global values (e.g. page_192

Page 193 deflections and flexural stresses) and local values (e.g. transverse stresses and interlaminar stresses) of thin-to-thick laminates. Stress-based and mixed ML models are, if anything, appropriate for local stress analysis, not for structural analysis and design, of composite laminates. Displacement-based layerwise models on the other hand match the aim of the present study very well because they model laminate-like inhomogeneity through the thickness in the most natural manner. Besides, they are easy to integrate into the existing systems and extend to dynamic problems, non-linear problems and modeling of delaminations. Increasing DOFs proportional to the numbers of discrete layers will not be so serious in near future due to the recent accelerating advancements of computational facilities symbolized by massive parallel computers equipped with vary large extensional core storages. Furthermore, one will witness less amount of total DOFs in the layerwise models than those discretized by 3-D continuum solid elements which is necessary to be meshed obeying a proper geometrical aspect ratio. Hence, as a proper numerical model for analysis and design of composite laminates, a "layerwise" and "higherorder" model will be appropriate, because this model is equipped with both simple two-dimensional geometry of ESL and three-dimensional/multi-layer state of displacements and stresses. It should be emphasized that such a numerical model that accounts for both of the layerwise and the higher-order concepts is still rare. In the present layerwise higher-order theory, a laminate is divided into a certain number of layers in the laminate thickness direction. There are several possible schemes and their combinations considered for dividing a laminate into layers as depicted in Fig.2. It should be noted that layers can be defined independently of physical laminae. The most natural manner of layer division is probably to divide a laminate according to the individual dissimilar material laminae (natural ML). However, in the present modeling, a laminate can be divided into sub-laminates which include several dissimilar material laminae in themselves (coarse ML). On the contrary, in some cases, an individual lamina should be divided into several layers so that complicated displacement fields can be modeled (fine ML). For clarifying this ML discretization scheme, the number of layers is denoted by NK, while the actual number of laminae is denoted by N. This flexible way of layer descretization enables designers to select an appropriate model corresponding to their analysis demands. By using the present model, a hierarchical numerical modeling scheme can be constructed.

Figure 2 Schematics of various multi-layer discretization schemes page_193 Page 194 The present model assumes the three displacement components in each discrete layer in the form of the power series expansions in terms of the thickness coordinates as follows:

where the superscript k denotes each one of layers of a laminate, numbered from the bottom layer to the top, and U(k)i P are the unknown coefficients of the series expansion. The geometric configuration of a laminate and the deformation of the typical layer k are schematically illustrated in Fig.3.

Figure 3 Coordinate system and geometry of a laminate, and higher-order deformation of the kth layer Geometrically, and are the translational components along x1, x2 and x3 axes, respectively, and and denote small rotations about x2 and -x1 axes, respectively. The rest of coefficients of power series expansions are higher-order influences such as parabolic stretching, cubic warping of the cross-section, the elongation of a transverse normal and so forth. It will be necessary to discuss whether or not these polynomial series assumptions can be sufficiently correct representations for displacement variations through the layer thickness. However at least it can be said that this polynomial approximation will be adequate except certain "unusual" local regions such as the vicinity of concentrated loads or the free edges. The series in Eq.1 should be terminated at a proper order of expansion, however, at this moment, the orders of termination is not determined yet for preserving and

the flexibility of the present theory. For conceptual sake, when the orders of the series for kth layer are for

and

respectively, the notation

will be used to express the order of displacement assumptions, and the notation

will also be used for the entire laminate made up of NK layers. Note that the parentheses denote layers and the brackets denote an entire laminate. page_194 Page 195 In the present general theory, there are possible choices undetermined, that is, layer discretization scheme : coarse ML, natural ML or fine ML; orders of terminations of higher-order displacement assumptions : up to lower-order like (110) or up to higher-order such as (332). In Fig.4, possible numerical models are charted with respect to the order of deformation assumptions and the order of layer discretization. In the conventional numerical models, all of these have been determined a priori. However, in the present modeling, these are individually determined corresponding to the actual analysis demand. In some cases, ESL-based first-order terminations will be enough, while in other cases such as evaluations of interlaminar stress singularity, extremely fine layer discretizations and higher-order terminations might be required. In order to establish an efficient numerical analysis and design for composite laminates, the most important task is to arrange an appropriate specific model to each individual problem. Hence in the present study, the numerical modeling methodology are proposed in Table 2. The spirit of the proposed methodology is that, since the laminate-like, through-the-thickness inhomogeneity of composite laminates plays a dominant role in numerical modeling accuracy, the numerical model is supposed to be determined in accordance with the inhomogeneity levels of the problem. For instance, when global responses of the entire

Figure 4 A schematic chart of possible combinations of deformation assumptions and layer discretizations Table 2 Possible numerical models for various levels of inhomogeneity of composite laminates inhomegeneity examples applicable number of layers and orders levels examples of supplementary conditions laminated members [(1,1,0)] + shear corrections, [(3,3,0)] + traction free Structural Member sandwich members [(110) (111) (110)], [(331) (332) (331)] + displ. continuities hybrid members [(110) (110) . . . (110)], [(330) (221) . . . (330)] + displ. continuities behavior of delamiantion [(110) (110)] + sparation and slipping failures of sandwiches [(110) (111) (110)], [(331) (332) (331)] + sparation and slipping Structural Member + Laminate-like CAI [(110) (110) (110)] + sparation and slipping interlaminar shear [(110) (111) (110)] + very thin middle layer stresses free-edge effects [(331) (332) (331)] + very thin middle layer Structural Member + Single Material toughened interlayer [(332) (332) . . . (332)] + Lagrange multipliers Laminate-like + interlaminar stresses Single Material Joints with overlays [(331) (332) (331)] + penalty numbers + [(332)] Strctural Member + Laminate-like + Single Material page_195 Page 196 laminated structure are of primary interest, ESL models like [(110)] or [(330)] will serve as the most appropriate model. In the case that "Laminate-like Phase" can not be neglected, such as in sandwich structures and delaminated plates, coarse ML models like [(110)(110)], [(110)(111)(110)] or [(331)(332)(331)] will be required. On the other hand, when one wishes to evaluate localized quantities such as interlaminar stresses, fine ML and higher-order models like [(332)NK] will be available. Finally, to illustrate advantage of the present model and methodology, an numerical example is shown. Structural joints are one of the most challenging aspects for engineers and designers in the analysis and design of composite laminates. No composite products are free from joints, even though the number of portions to be joined will be substantially decreased compared to that of metallic counterparts. Since the overall structural strength of composite products is dominantly governed by strength of joint regions, nowadays the context of experimental and numerical investigations of joints are very active. Hence, numerous attempts have been made. For instance, Shenoi [9,10] have been extensively performing detailed studies on structural composite T-Joints in marine structures by using both experimental and numerical approaches.

Figure 5 T-Joint construction (source : Dulieu-Smith el al.)

Figure 6 Design details of T-Joint construction The layerwise higher-order model of the author is modified and applied to stress analysis of a structural joint in laminates. The structural T-joint picked up herein is schematically shown in Fig.5. The T-Joint comprises a 560 mm flange and a 260 mm web which are 15 mm thick and made of 17 lay-ups of glass fiber woven roving (WR) cloth set. The flange and the web are connected by a resin fillet with a radius of 50 mm. The fillet is overlayed to form boundary angles. The boundary angles form 12 layup laminates, the detail of which is given in Fig.6. Delieu-Smith et al.[11] have quantitatively obtained stress distributions across the thickness of the boundary angles by thermoelastic stress analysis. As the experimental equipment for the thermoelastic stress analysis, they used the SPATE (Stress Pattern Analysis by the measurement of Thermal Emissions) system. page_196 Page 197 Further details of the thermoelastic stress analysis of this T-joint construction can be found in the original work of Delieu-Smith et al.[11]. In the present finite element analysis, [(332)5] and [(332)3] elements are used for modeling the T-joint. The web, the boundary layers and the fillet layers are individually represented by a specific layer. Elastic properties used for the present finite element analysis are, if any, taken from the original works of Delieu-Smith et al. [11] and given in Fig.6. Fig.7 shows the gridwork of finite element meshes and boundary conditions in the present finite element analysis. In order to compare the present finite element results to the experimental results in the original work, stress values are transformed into sp + ast. The directions of sp and st are schematically shown in Fig.6 as well. a is a experimentally determined constant and according to Delieu-Smith et al., a is set equal to 3.1 in this case. In Fig.8, thermoelastic stress distributions along Line 1 are compared. In addition to the present finite element results and the experimental results by SPATE system, Delieu-Smith et al. also obtained numerical results by 2-D continuum plain-strain element in ANSYS package, and therefore, they will also be shown for comparison. It is observed that the present FE results as well as the results by the 2-D continuum elements indicate fairly satisfactory predictions of the stress distributions across the thickness when compared to the experimental results. Therefore, the applicability of the present FE to detailed stress analysis of structural joints in composite laminates are, if partly, ensured. Applications of the layerwise higher-order models to joint problems are quite beneficial to the efficient analysis and design of

Figure 7 Gridwork of finite element meshes and boundary conditions

Figure 8 Thermoelastic stress distribution across the thickness along Line 1 page_197 Page 198 composite laminates. The example shown here implies that, by using the layerwise higher-order models alone, designer can evaluate almost all design quantities precisely, and conduct a satisfactory design of composite laminates. References 1. J.N. Reddy, A simple higher-order theory for laminated composite plates, ASME J. Appl. Mech. 51 (1984), pp.745752. 2. S.T. Mau, P. Tong and T.H.H. Pian, Finite element solutions for laminated thick plates, J. Compos. Mater. 6 (1972), pp.304311. 3. R.L. Spilker, Hybrid-stress eight-node elements for thin and thick multilayer laminated plates, Int. J. Numer. Meth. Engng. 18 (1982), pp.801828. 4. J. Han and S.V. Hoa, A three-dimensional multilayer composite finite element for stress analysis of composite laminates, Int. J. Numer. Meth. Engng. 36 (1993), pp.39033914. 5. H. Murakami, A laminated composite plate theory with improved in-plane responses, ASME J. Appl. Mech. 53 (1986), pp.661666. 6. J.N. Reddy, A generalization of two-dimensional theories of laminated composite plates, Commun. Appl. Numer. Meth. 3 (1987), pp.173180. 7. J.N. Reddy, An evaluation of equivalent-single-layer and layerwise theories of composite laminates, Compos. Struct. 25 (1993), pp.2135.

8. K. Suzuki, Layerwise Higher-Order Finite Elements for Laminated Composite Material Structures, Dissertation of Doctor of Engineering at the University of Tokyo (1997). 9. R.A. Shenoi and F.L.M. Violette, A study of structural composite tee joints in small boats, J. Compos. Mater. 24 (1990), p.644666. 10. R.A. Shenoi and G.L. Hawkins, Influence of material and geometry variations on the behaviour of bonded tee connections in FRP ships, Composites 23 (1992), pp.335345. 11. J.M. Smith-Dulieu, S. Quinn, R.A. Shenoi, P.J.C.L. Read and S.S.J. Moy, Thermoelastic stress analysis of a GRP tee joint, Applied Composite Materials 4 (1997), pp.283303. page_198 Page 199

Design of an Implant and External Fixation for the Treatment of Bone Fracture in Consideration of Mechanical Properties of Cortical Bone Tsuneo HIRAI INTEC-HIRAI Ltd. Miyamaki-nanasegawa, Kyotaname, Kyoto, 610-0313, JAPAN Yoshinobu WATANABE Kyoto Prefectural University of Medicine Kawaramachi-Hirokoji, Kamigyo-ku, Kyoto, 620, JAPAN Atsushi YOKOYAMA Kyoto Institute of Technology Matsugasaki, Sakyo-ku, Kyoto, 606-8585, JAPAN Keywords: Cortical Bone, Composites molded Rigid and Liquid System on Living Bone Abstract Treatment of damaged bone cannot proceed properly without the preservation of the bio-activity of the living bone, that is cells. It is necessary to preserve the active cells, which fulfill many functions in the damage region. In the treatment of bone fractures using implants, it is necessary to achieve mechano-compatibility as well as bio-compatibility. So it is desirable to investigate the mechanical characteristics on cortical bone. Fundamental elements of cortical bone consist of hydroxyappatite and symmetric ring sandwich laminate of collagen with different directions of reinforcement, called osteon. A solution for the behavior of such composites using quasi three dimensional computational analysis of the coupling moment caused by the osteon help to clarify the physical characteristics of cortical bone. It can be understood how the high load capacity depends on the coupling moment and the mechanical impedance determine the good well shock absorbing ability of osteon caused by the combination of various materials. The coupling moment caused by osteon should be analyzed by three dimensional numerical methods. 1 Introduction Cortical bone consists of inorganic materials, in the form of mineral salts, and on organic component forming a composite structure. At the microscopic level, the fundamental structure unit of bone is the osteon which consists of concentric of layers, called lamellae, in which the bone mineral is embedded in helically oriented fibers of the protein collagen. The orientation of the collagen is changed alternately in each lamella, forming laminate materials. According to this structural formulation, tension and compression to the cortical bone produce coupling moments. In this study, the relative values of the coupling moments in cortical bone are page_199 Page 200 nemerically found by three dimensional finite element analyses to explain the mechanical properties of the cortical bone. It was not possible to achieve competent mechano-compatible, whether it is secured theoretically mechanical in either case.

Fig.1 Model components at cortical bone

Fig.2 Soft X-image of a slice of knee joint 2 Toughness Model for Cortical Bone in Relation to the Loading System Fundamental elements of cortical bone consist of lamella molded into a laminated form of osteon with hydroxyappatite as interfaced matrix and a random combination of lamella in a matrix between many elements of osteon. The physical characteristics obtained for various arrangements of the elements very widely. Externally applied loads are transmitted through cartilage and cancellous bone to the objective cortical bone being the main bearing element dependent on the surrounding system as shown in Fig.2. The stress transmission through cancellous bone is dependent on the complex arrangement of solid and liquid phases, including marrow fluid in the bone. The synthesized structure is determined by bio-dynamic requirements. A honeycomb-shape element with liquid in all spaces1) and an optimized axial arrangement of the bone elements for stress dispersion could be used as a simplified model for theoretical calculation in cancellous bone. Results using the Newmark b system shown as in Fig.3 give good bio-mechanical compatibility. For a specified loading pattern on the cancellous bone. The progressive transmission of stress is shown into the cortical bone. The result of principal stress distribution using numerical analysis is shown in Fig.3, for eccentric loading. Initially the load transmitted extends over the upper domain of the cancellous bone and the stress developed in the cortical bone as a uniform distribution. The figure shows increasing incremental development without reference to the strength of cancellous bone. The increasing load develops an incremental increase of stress in the cortical bone. It could be assumed from the load history of the cortical bone that uniaxial loading could be used to investigate the physical characteristics. But load history should be page_200 Page 201 osteon in the cortical bones. Load history with on an AE count for an individual case has 8 steps of behavior determined by the points of inflection of the diagram shown in Fig.4. Through the average slope over the full range scams to have a useful physical meaning, the various features over at shorter stages of the diagram could be thought of as caused by realistic physical changes in the straining behavior. It should be requested on the experimental procedure to investigate symmetry behavior as living condition as remained relabolism. Initial conditions at above steps have a liquid phase showing iso-tensor condition, but the subsequent process might show with the current of a

bone characteristics. Then the behavior show as following: the first linear, the second become strain hardening and shows knee point, and then behave the sequent of solid bone structure, that is, cause a fracture as debonding or delamination, then load-strain history transmit uneven as shown in previous paper2). As described above, coupling moments will be generated between lamellae of each osteon pure axial loading. In long bones, the osteons usually run longitudinally, but the remodeling in living bone creates frequent branches and anastomose with one another, which is likely to produce the debonding of osteons rather than the delamination of lamellae under axial loading. The debonding of osteons will align the orientation of collagen more parallel, ultimately they will be broken. 3 Quasi-Three Dimensional Analysis of Osteon From the experimental results, it is considered that the deformation and fracture behavior of each component such as the osteon and lamella affect the complex behavior of cortical bone, for example, variations in slope of the stressstrain curve.

Fig.3 Stress dispersion in cancellous Bone on cortical bone

Fig.4 Stress-strain for load history and acoustic event in tensile test page_201 Page 202 It is believed that deformation behavior of the osteon largely contributes to the slope in the stress-strain curve. Coupling moments of cortical bone under loading caused by the osteon formation system should be obtained by three dimensional analysis to investigate the behavior. Finite element analysis may be useful to clarify factors of the behavior. Here, an important purpose is to establish a numerical model to simulate the behavior of the cortical bone that consists of osteon, lamella and collagen. Only the osteon part with collagen is included in the FEM analysis as the first step in this report.

3.1 Numerical Model Our proposed model is constructed of orthotropic shell elements and isotropic beam elements, which represent fiber and matrix, as illustrated in Fig.5. A cylindrical shape is built up by these elements. Collagen fiber and osteon matrixes are expressed by shell elements with orthotropic material constant. The osteon layer is regarded as a unidirectional composite, which consists of collagen fiber and osteon matrix. When the fiber orientation angle of the inner layer is +a°, that of the outer layer is -a°. It is assumed that the collagen has a helical structure. Interlamina between osteon layers, that is, the osteon matrix is expressed by beam elements. Inner and middle layer of shell elements as shown in Fig.5 express the structures of this osteon. Moreover, to express interface between the osteon and outer lamella, the outer cylinder of shell elements is set and is connected to middle layer as the osteon by beam element as interface region. Material constants of each element use the value in Table 1.

Fig. 5 Numerical model for osteon under outer layer loading Table 1 Material constant Young's Shearing Poisson's ratio modulus modulus Element for osteon 8160 3138 0.3 Element for lamella 1020 408 0 page_202 Page 203 In this material constant, Poison's ratio of lamella is 0 to express that lamella is surrounding to the osteon and have large volume as compared with that of the osteon. Tensile load is applied on the outside layer of the osteon as illustrated in Fig.5. Nodal forces on the cross section of the cylinder are treated as the applied load. The sum of these forces becomes the tensile load in the cylinder. Tensile stress is obtained, by dividing the tensile load by the section area of the cylinder. Also, strain is equal to displacement divided by the length of the cylinder.

Fig.6 Coupling deformation state of osteon as a sample of twin-layer of osteon

Fig.7 Normal stress of beam element at interface 3.2 Numerical Results Fig.6 illustrates deformation states at the end edges of the inner and outer layers obtained by calculation. In both the inner and outer layers, torsional deformation occurs due to the helical structure of the collagen fibers. Fig.7 shows the tensile stresses of interface between the osteon and the lamella and that of inner interface in the osteon. These values are calculated by beam element in the numerical model. From this result, the damage at the interlamina between the osteon layers might be large as compare with inner damage in the osteon. So the numerical results suggest that the delamination at interface between the osteon and the lamella is specific fracture mode of microstructure in cortical bone. 4 Features of Dynamical Behavior Caused by Physical Characteristics Osteon might produce the coupling moment due to the multiple quasi-cylindrical helical inclined elements of collagen. Though it might be expected that there will be delamination evident on some interfaces at the ends of lamellae under loading, shown by microscopic line markings, the figure of fracto-graphy(SEM) showing in Fig.8 indicates precious little such trace of osteon and the idea has been discounted excepting shown only a debonding at outer interface page_203 Page 204 with matrix in the figure. The theoretical results confirm that bi-plied coupling behavior shown a similar moment vector to that given in the previous figure. The stress distribution in the osteon of numerical results suggests the delamination of the interface between the osteon and the lamella in microstructure of cortical bone from actual fracture mode of the bone.

Fig.8 SEM Graph(on Cortical bone for younger generation) 5 Conclusions Cortical bone should be represented as composites consisting of complex elements with solid and liquid phases. In order to examine the behavior on osteon, it is necessary to use an analytical computation method for such a complex composite structure. References 1) T.Hirai, Proceedings of FRP Symposium Plenary Lecture II(1992) 2) T.Hirai and Y.Watanabe, JCOM:JSMS COMPOSITES-26,121 and 122(1997) page_204

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Design of a Polymer-Based Composite Container for the Long-Term Storage of Radioactive Materials H.W. Bonin, V.T. Bui, P.J. Bates, J.-F. Legault and J.Y.S.D. Pagé Department of Chemistry & Chemical Engineering Royal Military College of Canada - Collège militaire royal du Canada P.O. Box 17000, Station Forces Kingston, Ontario K7K 7B4 Keywords: Epoxy, carbon, radiation effects, cross-linking, chain scission Abstract This work aims at studying the feasibility of using polymeric-based composite systems in the design of a long-term storage container for radioactive materials. More specifically, the combined effects of radiation, heat transfer, and time (> 500 years) on these systems are to be assessed, in addition to structural stability and resistance to the environment of the container. Although the primary objective of the container is for surface storage, the design could eventually be extended for deep underground disposal. The conditions considered in the evaluation of this system are based on the long-term disposal concept proposed by Atomic Energy of Canada Limited (AECL), using containers made of copper or titanium. Polymer-based composites could be used as an alternative to these metals with the potential net advantages of less material degradation, lighter weight and cost reduction. Epoxy and Epoxy matrix reinforced with carbon fibres have been considered for this work. Preliminary results on Epoxy based on diglycidylether of bisphenol A (DGEBPA) cured with triethylene tetramine (TETA) showed an increased in cross-linking density after receiving a total dose of 2´104 Gy[1]. Provided that the container wall could be sufficiently shielded from the radiation emitted by the radioactive material (e.g. HLW), these types of materials represent an interesting option for this application. Introduction The aerospace, communications and nuclear industries rely increasingly on new polymeric materials for advanced engineering applications demanding high resistance to radiation. Among the most reliable and cost effective polymeric materials are phenoxy epoxies. The latter are resistant to high temperatures in addition to displaying good stability in both chemically active and low-dose radiation environments (