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 << Simulation and Measurement of Reinforcement Deformation During Preform Manufacture >>, Polymers and Polymer Composites, 4(5), 1996. [5] R. Blanlot, J. L. Billoet, << Evolutive Orthotropic Constitutive Equations for the Simulation of the Shaping of Composite Woven Fabrics >>, Journées Nationales des Composites, 1996. [6] C. D. Rudd, A. C. Long, P. McGeehin, P. Smith << In-Plane Permeability Determination for Simulation of Liquid Composite Molding of Complex Shapes >>, 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 (<107 Gy)[2]. Mechanical and thermal properties can also be improved by incorporating various kinds of fibers. These characteristics make epoxies composites excellent candidate materials for nuclear applications. An example of a commonly used epoxy system is the one based on a diglycidylether of bisphenol A (DGEBPA) resin cured with triethylene tetramine (TETA) (Figure 1). page_205 Page 206

Figure 1: Cross-linking reaction between diglycidylether of bisphenol A (DGEBPA) and triethylene tetramine (TETA). The properties of this polymer can be attributed to its chemical structure. The aromatic ring is a dominant component in the DGEBPA composition, and is mostly responsible for the strength, and for the resistance to high temperatures and radiation of this material, resulting from its absorption and dissipation of energy, without bond disruption[2]. Previous investigations on the effects of radiation on Epoxy and Epoxy/carbon composites have dealt mainly with gamma rays and electron beam (eb)[3][4][5]. Research involving the effects of neutrons on similar systems was also conducted, to a lesser extent[6]. Interaction of Radiation with Matter

A wide variety of radiations is produced by nuclear reactors. In the reactor core, the chain reaction of fission liberates photons, neutrinos and fast neutrons. In addition, highly radioactive pairs of fission products are created and undergo decay by one or a series of mechanisms where further types of photons and particles are produced : alpha, beta, gamma, high energy electrons, positrons and neutrinos. Most of these particles such as alpha, beta particles, electrons, and positrons are attenuated in the immediate vicinity of the fuel elements through interactions mechanisms such as electron stripping and collisions with nuclei[7]. In these processes, secondary radiations are emitted, including recoil protons and heavy ions, delta particles (stripped electrons) and X-rays. Photons and fast and thermal neutrons have deeper penetration power and can interact with materials far from the reactor core through mechanisms characteristic of their type of radiation. Photon Interaction. Photons (gammas and X-rays) ionize atoms through three main processes[2]. The photoelectric effect is mainly predominant for low energy photons (<0.05 MeV) and amounts to a total transfer of the photon's energy to one of the electrons from the inner shells (K or L mainly). As a result, this electron is ejected from the atom which is now ionized, and a subsequent rearrangement of the remaining electrons within the electronic shells causes the emission of X-rays and Auger electrons. The second mechanism is the Compton scattering, also predominant for lower photon energies (0.110 MeV). This amounts to a collision between the incident photon and one of the atom's electrons, preferably those located on the valence shells and which is ejected from the atom. As a result, the photon survives the collision, but with a page_206 Page 207 reduced energy and in a direction usually different than the incident direction. Both the scattered photon and the stripped electron may have sufficient energies to ionize several more atoms. Finally, the photon may interact with matter through the pair production mechanism, possible only if the photon's energy is equal to or greater than 1.022 MeV (i.e. twice the electron's rest mass energy). This phenomenon occurs in the vicinity of a heavy nucleus and amounts to the disappearance of the incident photon leading to the creation of an electron-positron pair. Both particles are emitted in opposite directions and interact subsequently by ionizing more atoms. The interplay of all three interaction mechanisms results in a fairly uniform photon energy deposition within the irradiated the irradiated material, such as organic polymers in which the Compton scattering is predominant for photons energies below 2.5 MeV. Neutron Interaction. Fast (up to 12 MeV) and thermal (0.0253 eV) neutrons interact with atoms in a more complicated fashion. Since their electrical charge is neutral, they are not technically ionizing radiation, but in practice they are indeed, albeit indirectly. The ionization is actually produced by the recoil charged particles, mostly protons, from inelastic and elastic collisions between the neutrons and light nuclei through a knock-out process. Since the ionizing power from energetic charged ions such as protons is much larger than that of electrons, higher damage is caused (indirectly) from the incident neutron. In addition, both fast and thermal neutrons can be absorbed into the nuclei of atoms that are transmuted into (usually) radioactive isotopes through the nuclear reaction. These radioisotopes then undergo radioactive decay accompanied by the emission of particles such as beta and gamma. Depending on the atoms present in the irradiated material, the induced radioactivity may remain long after the end of the irradiation. However, for most of the polymers of concern here, the majority of the atoms present, such as C, O, N and H, have either a low neutron absorption cross-section or the activation products are short-lived. Consequently, the activation process is a major concern only when impurities (mostly metallic) are present in significant quantities. The overall effect of the neutrons interacting with matter can be summarized by the high ionization powers of the recoil protons and ions concentrating damages along their specific tracks. However, since the neutrons have a large penetration power, the ionizing recoil particles are created within large volumes and the radiation damage is spread quite evenly throughout the material. Level of Damage. The level of damage caused by radiation depends not only on the quantity of the energy deposited in the material but also on how concentrated the energy deposition is, creating free-radical pairs along the path taken by the energized particle or photon[8]. The destructive potential of radiation is assessed in terms of linear energy transfer (LET) which is the amount of energy deposited per unit length of the ionization track, expressed as dE/dx[9]. Radiation-induced changes. Radiation-induced chemically active species can react in different ways resulting in structural and molecular weight changes. The first change is cross-linking, which increases the molecular weight. This phenomenon can eventually result in the formation of an important network which radically alters transition temperatures, chemical and mechanical properties of the material[8]. The second change is chain scission, where free radicals recombine with low molecular weight molecules, which causes a decrease in the overall molecular weight of the material. The third change is the creation of low molecular weight products resulting from page_207

Page 208 chain scission and followed by the abstraction and recombination reactions. These products are usually gases such as H2, CO2, CO and CH4, but some polymers with particular structural groups may release other gases like SO2, in the case of polysulfones, or HCl for polyvinylchloride[10]. Such small products impact on the overall molecular weight and create small gas pockets in the polymer which reduce mechanical properties. The last induced chemical change is the actual structural modifications by the appearance of new bonds creating new chemical groups or the elimination of existing ones. Such changes can alter both chemical and physical properties of the polymer. Methodology Materials Samples of epoxy were cast in dumbbell shape (ASTM 638) using DGEBPA and TETA at a proportion of 9:1 in weight, at ambient temperature for over 24 hours. This allowed for a reference baseline to study initial radiation effects. Samples of Epoxy/carbon composites had not been subjected to any experimental work at the time of submission of this article. Irradiation Processing Irradiation was performed in the SLOWPOKE-2 nuclear research reactor, producing mostly thermal and epithermal neutrons, and gamma rays. The dose rates for the neutrons and radiation at the irradiation sites in the reactor pool were determined from measurements: 600 ± 50 Gy/h for neutrons, 12 ± 4.5 Gy/h for fast neutrons, and 3600 ± 1250 Gy/h for gamma rays, at full reactor power. Tensile Test The mechanical properties of polymers are driven by their viscoelastic attributes, characterized by the tensile testing. Viscoelasticity is furthermore dependent on the polymers' molecular structure, defined by chemical bonds and chain length, but also by their morphology. The changes in the viscoelastic properties of polymers exposed to aggressive environments such as radiation are often an excellent indicator on the ability of polymeric materials to maintain the mechanical resistance required from them throughout the lifetime of their application. The tensile tests were performed according to ASTM D638 on non-irradiated and irradiated samples with an Instron Tester, model 4206, fully automated. The tensile load cell of 500 N was used. The crosshead speed used was 1 mm/min. Density Measurements It is known that the cross-link degree of a thermoset material such as cured epoxy is directly proportional to its density. The density of samples was measured following the ASTM D792-A procedure in which the samples were weighed in air and in distilled water, successively, at room temperature. page_208 Page 209 Results and Discussion Radiation Effects on Cured Epoxy Results from irradiation of Epoxy are shown on Figure 2. Also displayed are the results for polystyrene (PS), and poly(methyl methacrylate) (PMMA), and poly(ether ether ether ketone) (PEEK) for comparison[11]. Results for Epoxy show a slight increase in mechanical properties low irradiation doses. This may be caused by a continued increase in its cross-linking degree through some radical sites mildly formed by radiation effects. Furthermore, an increase in the density of these weakly irradiated samples (from 2.044 to 2.169 g/cm3) appears to confirm this phenomenon[1].

Figure 2: Tensile Strength Tendencies of Irradiated Polymers[1][11]. Radiation Effects on Epoxy/Carbon Composites As a second step from previous experiments on Epoxies, this work plans to irradiate and test Epoxy/carbon (continuous fibers) composites. The latter will be prepared from prepregs and irradiated in the SLOWPOKE-2 reactor for a combination of different times (or doses) and temperatures. By subsequently evaluating their mechanical properties, a correlation between them, radiation, time and temperature will be investigated. This would lead to a model for ageing in a radiation environment. Other hostile environmental conditions such as the presence of water (moisture) and oxygen can also be incorporated to further investigate resulting effects on material properties. The model could be used to assess the suitability of Epoxy and Epoxy/carbon composites for radioactive waste containers. page_209 Page 210 A literature review reported that fast-neutron irradiation of carbon fibers in air showed an increase in strength followed by a decrease (by as much as 25%) of the control[12]. However, irradiation in a inert environment showed only an increase in the strength[13]. At elevated temperatures, a small increase in both the strength and modulus of carbon was observed[14]. This is due to an increase in the crystallite dimension of carbon. Epoxy/carbon composites irradiated in air at 75°C with a combination of fast and thermal neutrons showed a decrease in flexural strength and modulus[15]. When samples were irradiated in liquid nitrogen, increases in the strength and modulus were observed when tested at liquid nitrogen temperature while a decrease in those parameters occurred when irradiated at liquid nitrogen temperature but tested at room temperature. The results reported for the irradiation by gamma rays and high energy electron beam are consistent with those of earlier works[16]. Design of Container The level of radioactivity is one of the key elements in the design. Since the container must maintain its structural integrity for at least 500 years, a means of assessing the effects of the total dose rate corresponding to this lifetime is to put in place. Based on existing works on Epoxy and Epoxy/carbon composites, these systems could take the total dose generate by Low Level Waste (LLW). As far as HLW is concerned, theoretical dose calculations using the Microshieldä software showed that total dose would exceed dose capacity reported to date[17]. However, the use of radiation absorption substances as filling material in the container would decrease the dose reaching the container wall to acceptable values. Further to investigating a means of simulating the lifetime of the container, a correlation between the combined effects of dose, time and temperature will be investigated. More research is also to be carried out on the effects of a hostile environment, such as thermal cycling, and the exposure to air and water (moisture). References 1. Bonin H.W., et al., Radiation Effects on Aluminium-Epoxy Adhesive Joints, J. Appl. Polym. Sci, 67, 3747 (1998). 2. Encyclopedia of polymer science and engineering, 2nd Ed., John Wiley & Sons, New York, 1986, 13, pp 667703. 3. Seo. K.S., et al., Effects of Ionizing Radiation on Epoxy, Graphite Fiber, and Epoxy/Graphite Fiber Composites. Part II: Radial Types and Radical Decay Behavior, J. Polym. Sci.: Part B: Polym Phy, 26, 533544 (1998). 4. Netravali, A.N., Manji, A., Effects of60Co Gamma Radiation on the Mechanical Properties of Epoxy Blends and Epoxy-Graphite Fiber Interface, Polym. Comp., 12, no.3, 153160 (1991). 5. Fornes, R.E., et al., Effects of 1.33 MeV g Radiation and 0.5 MeV Electrons on the Mechanical Properties of Graphite Fiber Composites, J. Appl. Polym. Sci., 26, 20612066 (1981). page_210 Page 211 6. Egusa, S., Effects of Neutrons and Gamma Rays on Polymer Matrix Composites as Low-Temperature Materials, Radiat. Phys. Chem., 37, no. 1, 147152 (1991). 7. Glasstone S. and Sesonske A., Nuclear Reactor Engineering, Van Nostrand Reinhold, 3rd edition, New York, 1981, pp 438.

8. O'Donnell J.H., ''Chemistry of Radiation Degradation of Polymers", in R.L. Clough and S.W. Shalaby, Eds., Radiation Effects on Polymers. American Chemical Society, Washington, DC, 1991. 9. Sangster D.F., "Early Events in High-Energy Irradiation of Polymers," in E. Reichmanis and J.H. O'Donnell, Eds., The Effects of Radiation on High-Technology Polymers. American Chemical Society, Washington, DC, 1989. 10. O'Donnell J.H., "Radiation Chemistry of Polymers," in E. Reichmanis and J.H. O'Donnell, Eds., The Effects of Radiation on High-Technology Polymers. American Chemical Society, Washington, DC, 1989. 11. Pagé, J.Y.S.D., Master's Thesis: Neutron and Gamma Radiation Effects on the Viscoelasticity Behaviour of Poly(ether ether ether ketone), Department of Chemistry and Chemical Engineering, Royal Military College of Canada, Kingston, Ontario, May 1997. 12. R.E. Bullock, Radiat. Eff., 11, 107 (1970). 13. R.E. Bullock, Fiber Sci. Technol., 7, 157 (1974). 14. B.F. Jones and I.D. Peggs, Nature (London), 239, 95 (1972). 15. R.E. Bullock, Mater. Sci. Eng., 10, 178 (1972). 16. R.E. Fornes, J.D. Memory, and N. Naranong, J. Appl. Polym. Sci., 26, 2061 (1981). 17. Davey, A., et al., Examination of the Feasibility of using Polymeric Composites in the Fabrication of a Container for the Long-Term Disposal of Spent Nuclear Fuel, RMC-CCE-CM417-97-2, Department of Chemistry and Chemical Engineering, Royal Military College of Canada, Kingston, Ontario, February 1996. page_211 Page 212

Glass Fiber from Canada Resist in Acid Condition Y. FUJII Seikow Chemical Engineering & Machinery, Ltd Kukuchi 3-13-33, Amagasaki, 661-0977, Hyogo, JAPAN Key Words: GFRP laminates, Acoustic emission, Environmental-Creep, E glass, ECR glass 1 Introduction Glass fiber reinforced plastics (GFRP) are extensively used in the chemical process industry for applications such as pipe work, reaction vessels, storage tanks, pumps, fans, scrubbers, etc. During service GFRP components are subjected to different chemical environments and load conditions. In most environments GFRP is reasonably inert, especially when it is not subjected to service loads. It is necessary to understand the behavior of GFRP subjected to loads under chemical environments (also known as stress corrosion). References [110] give an account of previous research work on stress corrosion behavior of composite materials. This acid stress corrosion of GFRP is related to the corrosion of reinforcement glass fiber. In this paper, it is shown the GFRP made from anti corrosive glass fiber resist the acid condition. 2 Materials The glass fiber reinforced polyester laminates used in this study were fabricated by a hand lay-up method. Chopped stand mat were used as reinforcements. The matrix resin namely vinylester resin was used. The mat reinforced composites were made by reinforcing with 3 plies 3mm thickness. In this study used 2 types glass chopped strand mat. One is E type glass fiber(ECM 450-193, CG in Japan), another is ECR glass(ECRM 723, OCF, from Canada). All the laminates were post cured at 100°C for 3 hours. Dumbbell shaped specimens were prepared according to the ASTM D638. Polyvinylchloride (PVC) end tabs were glued to the specimen ends using an epoxy adhesive. The authors have found through experimentation that the grip noise can be reduced by using PVC end tabs. 3 Experimental Details

3.1 Test Set-up The details of the creep test set-up are shown schematically in Fig. 1 and 2. The load on the specimen was varied by changing the dead weight. Glass tube with rubber stopper as shown in Fig.2 was used for maintaining several environments. The test environment (5% nitric acid) was introduced into the tube using a hypodermic syringe. The strain gauge was mounted on the non-immersed portion of the specimen for measuring longitudinal tensile strain. page_212 Page 213 3.2 Test Methods For a specified creep stress each specimen was initially creep tested in air for 2hrs for the acoustic emission activity reaches a steady state and then creep tested in environments until the specimen ruptures or testing stopped. Several specimens were tested at different creep stress levels. 3.3 Acoustic Emission Measurement AE signals were detected using a piezoelectric transducer with a resonant frequency of 150KHz. Sensor was mounted on the specimen as shown in Fig.2. During testing only the ring down counts were measured and recorded using a X-T recorder. 4 Results and Discussion 4.1 Step up Loading AE rate on the each step up loading in the air and acid condition are shown in Fig.3. The load vs deflection or strain relation is very similar for both the air and acid conditions. However, because, AE measurement is very sensitive, this environmental effect can be detected easily. This testing is step up loading and measuring the AE count during each 30 min. ECR (show C) type GFRP shows no change between air and acid conditions. But, E (show E) type GFRP is reflect the environmental condition. The effect of environment appear at the large stress levels. In the low stress condition, environmental effect is little, so we could not detect the difference. 4.2 Creep in Acid Condition In acid solution and constant load condition, 2 types glass reinforced mat GFRP creep results are shown in Fig.4. and 5. In Fig. 4, test started in air condition for 2hrs. AE monitoring, and then acid solution input the tube. In air condition, AE occurred a little and similar manner with 2 types GFRP. But in acid condition, very quickly the difference is appeared. In this scale, AE occurrence is no change, but for E-type GFRP AE increased clearly. In Fig. 5 shows long time creep result. In E type GFRP, great number of AE occurred and failured, but C type GFRP occurred little AE signal and did not failure for long times. This AE signal suggested something damage occurred in the GFRP sample before failure. From the measurement of creep deflection, we could not detect the sign of before failure. 4.3 Life Time in Acid Several stress levels creep testing result show in Fig.6. E types mat GFRP failured at the short time with increase of stress. C types mat GFRP failure only 3 samples now, because of stopping test for no AE signal. 5 Conclusions In the acidic condition, Glass fiber types (E and ECR) influence the creep life and AE rate page_213 Page 214

The selection of glass fiber type is very important. This study has clearly indicated that the acoustic emission measurement technique is sensitive to the stress-corrosion damage in GFRP laminates. The creep life of all the composite specimens decreased with increasing creep stress, whereas the rate of AE counts increased with increasing creep stress. AE monitoring shows the something information of damage propagation with corrosion of reinforcement. References 1. F.R.Jones, J.W.Rock and A.R. Wheatley, Composites, 14( 3), 262(1983) 2. W.S.Carswell, ASME Petroleum Division, 24, 105(1988) 3. P.J.Hogg, D.Hull and M.J.Legg, Composite Structures, 106(1981) 4. P.J.Hogg, Composites, 14( 3), 254(1983) 5. P.J.Hogg, Composites Science and Technology, 38, 23(1990) 6. G.P.Marshall and D.Harrison, Plastics and Rubber Processing and Applications, 2( 3), 269(1982) 7. M.Kumosa, D.Hull and J.N.Price, Journal of Materials Science, 22, 331(1987) 8. R.Hill, C.Cowking and W.S.Carswell, Composites, 20(3), 215(1989) 9. Y.Fujii, Z.Maekawa, H.Hamada, T.Kubota, A.Murakami and T.Yoshiki, in A.Miravete,ed., Proceedings of ICCM 9, Woodhead Publishing Limited, Madrid, 562(1993) 10. Y.Fujii, S.Ramakrishna and H.Hamada, Durability Testing of Non-Metallic Materials, 190(1996)

Fig. 1 Schematic illustration of creep testing apparatus

Fig.2 Schematic illustration of creep test with environment page_214

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Fig.3 Effect of environment for 2types GFRP at the step up loading.

Fig.4 Effect of environment for 2types GFRP at the creep condition.

Fig.5 AE monitoring on the creep in acid solution. (Effect of glass fiber type)

Fig.6 Creep life time at the several stress levels in acidic solution. page_215

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Thermoplastic Lined FRP Dual-Laminate Composites for Corrosive Applications - An Overview Paul Habib - C.P.F. Dualam Inc. 11750 J.J. Joubert Montreal, Qc H1E 7E7 Abstract Fabricated Dual-Laminate Equipment such as towers, scrubbers, process vessels and tanks have been used for over 30 years in highly corrosive applications. These are commonly found in the chemical process industry, pulp and paper, and metal refining where chemicals such as chlorine & chlor-alkali products, strong acids, strong bases, organic compounds and others are present. The availability of chemically resistant thermoplastics, ex. polyvinyl chloride, polyolefins and fluoropolymers allows the Design Engineer to tailor materials of construction to a particular chemical service. Other design parameters that must be considered are temperature, pressure (internal & external), specific gravity of liquid, mechanical loads (wind, seismic, and attachments such as platforms), liquid and vapour permeation etc. This paper describes Dual-Laminate Construction in terms of material selection, design and fabrication techniques. Some of the applications where Dual-Laminates are used are also outlined. Introduction Plastics materials have been used in numerous applications where corrosive chemicals are present. Thermoplastics are well known for their excellent chemical resistance and are commonly used in piping, valves, hoods, and tank linings. They unfortunately do not have sufficient mechanical strength to make them suitable for large structures such as storage vessels, towers, and stacks. Fibreglass Reinforced Plastic (FRP), on the other hand, have excellent mechanical properties (flexural, tensile, strength and modulus). They are often used to fabricate large, complex shape equipment. The combination of thermoplastic liners and FRP thermoset composites provides structures, commonly called "DualLaminate", with excellent chemical resistance and structural strength. Dual-Laminates are now used in numerous applications replacing exotic metals & alloys, lined steel (glass, stoneware or rubber). They offer a cost-effective solution in some of the most corrosive services encountered in Industry. Dual-Laminate Construction -Definition Dual-Laminates are defined as a "thermoplastic construction to which is bonded a fibreglass reinforced plastic structure". page_216 Page 217 A typical laminate construction is shown in Fig. 1 below, it consists of:

Fig.1 Typical Dual-Laminate Construction A- Typical thermoplastic liners include polyvinyl chloride (PVC), post-chlorinated polyvinyl chloride (CPVC), polypropylene (PP), polyvinylidene fluoride (PVDF), ethylene chlorotrifluoroethylene (ECTFE), ethylene tetrafuoroethylene (ETFE), fluorinated ethylenepropylene (FEP), Hyflon (MFA), a copolymer of tetrafluoroethylene and perfluoromethylvinylether., and Perfluoroalkoxy (PFA). All liners, with the exception of PVC and CPVC require an embedded fabric backing (fibreglass or polyester) on the back to provide an adequate bond to the FRP. Chemically etched thermoplastic can also be used under certain conditions. B- The thermoset bonding layer provides a mechanical and chemical lock with the FRP.

C- A layer of electrically conductive material (typically a carbon fibre veil) is laminated immediately behind the thermoplastic, allowing for weld spark testing or leak detection in the event of liner damage. D- A secondary corrosion barrier is laminated using thermoset corrosion resistant vinyl ester resins and alternate layers of polyester (Nexus) or glass (C-Veil) fabrics and chopped strand fibreglass mat (CSM). Such laminate must be resin rich to insure adequate secondary chemical corrosion resistance. E- The FRP structure consists of either a hand lay-up, i.e alternate layers of CSM and fibreglass woven roving (WR), or filament winding construction (continuous roving wound at a predetermined wind angle, typically 54.7 degrees). Other winding patterns can also be used providing design and laminate properties have been validated (ex. hoop/unidirectional winding). F- An exterior finish provides for UV resistance, flame retardancy and additional corrosion resistance. FRP Constructions are covered by a number of Industry standards such as : CGSB-41.22-93 Canada (1) PS-1569 - U.S.A. (2) page_217 Page 218 ATM D4097 - U.S.A.(3) ASME RTP-1 - U.S.A.(4) ASME Section X (Pressure Vessels) - U.S.A (5) BS-4994 - United Kingdom (6) Certain practices relating to design and fabrication of Dual-Laminates are specified in BS-4994 (6) as well as other European standards. An active Dual-Laminate subcommittee of the American Society of Mechanical Engineers is currently preparing an RTP-1 Mandatory Appendix M-13 to standardize materials, techniques of design and fabrication. At the present time, however, North American buyers of Dual-Laminate equipment must rely on "Company Specific" standards and on reputable fabricators. Chemical Resistance The main reason Dual-Laminate found success in the Chemical Process Industry is the superior chemical resistance of the thermoplastics compared to FRP only. Suppliers of resin and lining materials produce detailed chemical resistance charts and should be consulted before selecting a particular material. In some cases actual tests need to be undertaken to assess chemical resistance in specific chemical environments. Table 1 compares the chemical resistance of the most widely used thermoplastics in Dual-Laminate construction to FRP. Stress Cracking Some thermoplastics are susceptible to environmental stress cracking in certain liquids. It is important to identify the environments that will cause stress cracking and to take precautions to ensure stress levels in the thermoplastic are within an acceptable range (typically below 450 psi) during the entire service life of the vessel. This is accomplished by heat annealing the thermoplastic structure after moulding and welding. Permeability All plastic materials are susceptible, to various degrees, to permeation of the gases and liquids being handled. Temperature, pressure, liner thickness, etc. will determine the degree of permeation. For that reason, most Dual-Laminate equipment is rated for 110 Deg. C or less even for fluoro polymers which normally would have much higher temperature ratings. At these temperatures (ie under 110 deg C) the level of permeation is low and the FRP secondary corrosion barrier provides sufficient chemical resistance to withstand the chemical. Fabrication of Dual-Laminates Bonding Strong and uniform interfacial adhesion between the thermoplastic and FRP is essential, particularly in vacuum service. To ensure adequate bonding with PVC and CPVC, a bonding

page_218 Page 219 resin (modified isophthalic) is used as a chemical and mechanical lock. For other thermoplastics, the bond is achieved via a fabric backing embedded into the liner under heat and pressure. As such, there is no direct bond between the FRP and the thermoplastic. The adhesion consists of a mechanical lock created by the fabric backing. Adhesion between FRP and thermoplastic is characterized by: shear strength, peel strength and tensile strength. A schematic representation of the three tests is shown in Fig. 2,3 & 4.

Fig. 2 Coupon for Interfacial Shear Adhesion Strength

Fig. 3 Test Determination of Peel Strength of Bond

Fig. 4 Test for Tensile Adhesion Strength Forming & Welding Techniques for forming and welding thermoplastics have been developed to accommodate literally any shape encountered in process vessels : Ex. ledges, baffles, nozzle connections, internal attachments, headers, and etc. A Typical FRP vessel is represented schematically in Fig. 5. The techniques that are specified for FRP vessels must be adhered to for structural integrity with regard to design safety factors. In addition, Dual-Laminate Fabricators must ensure that the stress levels in the thermoplastic are within acceptable levels and all moulded sections and welds are free of pin holes and residual stresses. To ensure welds are pin hole free, a 10,00015,000 volt spark testing is conducted on all welds. page_219 Page 220

Fig.5 Typical FRP Vessel and Fabrication Applications Dual-Laminates have found their niche in numerous applications. The following examples illustrate a few: Chlor-Alkali Plants - chlorine drying towers - piping and headers Pulp Bleach Plants - sodium hypochlorite tanks - chlorine dioxide generators - washer drum suction box Environment Protection - chlorine & chlorine dioxide scrubbers Metal Refining Plants - acid storage tanks Basic Chemicals - bromine process vessel page_220 Page 221 Design All composite equipment used in corrosive service must be designed to meet all mechanical loads such as wind load, seismic load, internal and external pressures (vacuum) internal loads, ex. packing weights, baffles etc.

Fig. 6 Tefzel/FRP Dual-Laminate Vessels; Critical Components to an Acid Regeneration System; 18ft. diameter 40 ft. long; Shipped by barge from Port of Montreal Typical design safety factor of 10:1 is used to ensure equipment would be subjected to low stress and strain levels. Design and fabrication methods must take into account the anisotropic nature of FRP composites. Lamination theory design is often used to evaluate laminate properties. One must also consider the contribution of each layer, the interfacial properties of each layer, fibre orientation, resin content, and mechanical properties of each laminate.

Fig. 7 Teflon FEP/FRP Dual-Laminate Pressure Vessels; 25 psig; built as per ASME Section X Class II; In process of being Acoustic Emission Tested page_221 Page 222 Conclusion Dual-Laminate structures can be designed to withstand, all mechanical loads such as pressure, wind loads, seismic loads, and platform attachments, encountered in most chemical plant environments. In addition, dual-laminate structures offer superior chemical resistance, hence extending the useful life of the equipment and reducing maintenance costs associated with corrosion. It is paramount that the design and fabrication techniques take into account the anisotropic nature of composites, allowing for proper thickness calculation (maintaining a minimum safety factor) and lamination methods, particularly for secondary joints, nozzle attachments, etc. Properly fabricated Dual-Laminate equipment have repeatedly solved numerous corrosion problems in chemical plants, pulp & paper mills, metal refining operations, and wherever corrosive chemicals are being handled. In addition they very often offer a lower cost solution compared to exotic metal alloys. References 1. Canadian General Standards Board CAN/CGSB 41.22-93 Fibreglass-Reinforced Plastic Corrosion-Resistant Equipment, 1993

2. National Bureau of Standards NBS - Voluntary Product Standard PS-1569 Custom Contact Molded Reinforced Polyester Chemical Resistant Process Equipment, Nov 15, 1969 3. ASTM D 4097-88 Standard Specification for Contact-Molded Glass Fiber-Reinforced Thermoset Resin ChemicalResistant Tanks Vol. 8.04, Sept. 1988 4. ASME/ASNI RTP-1 1995 Reinforced Thermoset Plastic Corrosion Reistant Equipment, The American Society of Mechanical Engineers, 1995 5. ASME Boiler and Pressure Vessel Code X, Fiberglass Reinforced Plastic Pressure Vessels, The American Society of Mechanical Engineers, July 1, 1989 6.British Standards Institution, BS 4994 British Standard Specification for Design and Construction of Vessels and Tanks in Reinforced Plastics, 1987 page_222 Page 223 Table 1Chemical Resistance of Thermoplastics & FRP PVCCPVCPPPVDF ECTFE ETFE HALAR TEFZEL Hydrochloric Acid (HCl) @36% A3 A4 A4 A5 A5 A5 Sulfuric Acid (H2SO4) @ 96% A3 A3 B1 A3 A4 A5 Nitric Acid (HNO3) @ 50% A3 A3 B1 A2 A4 A5 Hydrofluoric Acid (HF) @ 40% A2 A2 A3 A4 A5 A5 Sodium Hydroxide (NaOH) @ 50% A3 A4 A3 X A5 A5 Sodium Hypochlorite (NaOCl) with A3 A5 A1 X A5 A5 15% C12 Hydrogen Peroxide (H2O2) @ 30% A2 A2 A2 A5 A3 A5 Chlorine Dioxide Gas (ClO2) A3 A4 X A2 A3 A5 Keytones X X A2 B2 A3 A5 Esters X X A2 B2 A3 A5 Chlorinated Solvents X X B1 A4 A4 A5 Legend:

A- Excellent resistance B- Marginal resistance X- Not Recommended

FEPPFA/FRP MFA A6 A6 B1 A6 A6 X A6 A6 X A6 A6 X A6 A6 A1 A6 A6 A1 A6 A6 A6 A6 A6

A6 A2 A6 A2 A6 X A6 X A6 X

1- up to 20° C ( 68 ° F) 2- up to 40° C ( 104° F) 3- up to 60° C ( 140° F) 4- up to 80° C ( 176° F) 5- up to 100° C ( 212° F) 6- up to 120° C ( 248° F) page_223 Page 224

Prediction of Failure in Unsaturated Polyester Reinforced by Plain Weave Glass Fabric by H. Nguyen-Hoa & T. Vu-Khanh Université de Sherbrooke Faculté des sciences Appliquées Département de genie mécanique 2500 boul. de l'Université Sherbrooke, Québec, Canada, J1K 2R1 Abstract Polyester/glass fiber is probably by far the most commonly used composite. This paper presents the results of an investigation on the failure of this composite made of plain weave glass fabric. A recently proposed sub-plies model was utilized to predict the failure strengths of the fabric composite after shear deformation of the interlaced yarns, induced by the forming process. A procedure for the determination of the equivalent on-axis tensile and compressive strengths of the constituents plies is also proposed to take into account of fiber undulation effect on mechanical properties. The predicted compressive and tensile strengths are compared with experimental measurement as a

function of the deformation angle of the interlaced yarns. Furthermore, the onset of damage was also compared with the initial failure of the sub-plies model. The predicted results are in good with the experimental data. 1 Introduction Woven fabrics are now considered to be one of the most important reinforcing material in the composite technology. Although the structure efficiency of fabric composites is not as high as that of the unidirectional laminates, their versatile properties and low fabrication costs have made fabric composites attractive for structural applications [1], particularly for parts with complex shapes. Prediction of the thermo-mechanical properties of woven fabric composites is a great interested subject and much works have been carried out to model the properties of these materials [37]. The woven fabrics are treated as two-dimensional with two sets of yarns interlaced at a 90° angle, their thermo-mechanical properties are considered as orthotropic. Three main different models, which have been employed to approximate the thermo-elastic behavior of fabric composites, are known as the mosaic model, the fiber undulation model (crimp model) and the bridging model. The failure process in these materials is relatively complex and only a few studies have been reported on the prediction of strength for orthotropic woven fabric composites [9,10]. When the part has a double curvature, the forming process usually results in significant in-plane shear deformation of the interlaced yarns and the fabric composite become a non-orthogonal structure, so these models can no longer be valid. The sub-plies model has been proposed for the characterization of the thermo-elastic properties for any deformed woven fabric composite [1116]. This work aims at applying the sub-plies model in the prediction of tensile and compressive strengths of orthogonal and non-orthogonal woven fabric laminates. To take account of fiber undulation effect, a special method is proposed to determine the equivalent on-axis strength coefficients of the constituent sub-plies. The ultimate strength and the onset of damage as a function of the in-plane shearing deformation angle are discussed. 2 Modeling of Degradation In the sub-plies model, the fabric, consisting of a warp and a weft, is replaced by four fictional unidirectional laminae with orthotropic properties that are laid up in an antisymmetric page_224 Page 225 configuration [0h1/90h2/0h2/90h1] [1116]. Consequently, to predict the failure envelope of composites made of woven fabrics, the strength coefficients such as: longitudinal tensile and compressive strengths X, X', transverse tensile and compressive strengths Y, Y' and shear strength S are needed. A special procedure has been proposed to determine the thermo-elastic coefficients of the fictional constituent plies. In this work, the same approach was used to determine the strength coefficients and the results are compared with that measured on unidirectional composite and on [0/90]s laminate. 3 Experimental The sample were prepared from plain weave fabric WR180Z and AK2100 unsaturated polyester resin, supplied by Armkem Canada, with the fiber volume content of about 50 %. All of specimens were fabricated by hand lay-up method. The fiber volume fraction of the molded samples was always verified after molding by pyrolyzing the resin of the specimen in order to insure a variation of the fiber volume fraction within ±1.0 %. Plaques of unidirectional, [(0/90)n]s, and plain weave fabric composite 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. Tensile tests were carried out according to the ASTM D-3039 standard [18, 20]. The specimens were loaded monotonically to failure at a recommended rate of 1.27mm/min. Strain gages were mounted on the specimens and monitored during the test. Two gages, one in the longitudinal and one in the transverse directions were bonded to the specimen at the center of the test section. The modulus and Poisson's ratios were determined by at least 25 data, measured in the linear region. The compressive strength tests were carried out according to the ASTM D-3410-87 standard (Celanese compression test fixture) [20]. The specimens were loaded at a rate 0.5mm/min. Shear tests were performed using the [(45/-45)2]s coupon to determine the ultimate shear strength and shear modulus. All the tests were carried out on an Instron TTD-3025 machine. All the results were obtained by treating the experimental data according to the Weibull's two-parameter method. 4 Determination of Strength Coefficients

4.1 Measurements Using Woven Fabric Composite Tsai et al. [31], showed theoretically and experimentally that the stress-strain diagram of a cross-ply composite under tension is represented by a bilinear line. The intersection of the two straight segments implies the initial failure, and the intersection was called the first knee-point [32]. In the case of plain weave fabric composite, the tensile stress-strain curve for laminated specimen tested at q = 0° is shown in Figure 1. The result reveals that, as in the case of cross-ply laminated, a bilinear behavior is also observed. The point at which the two linear sections intersect about 23% of the ultimate load is used to determine the transverse strength of the fictional 90° ply in the sub-plies model. The ultimate load and deformation were used to calculate the longitudinal strength of the 0° fictional ply. However, the compression test does not show this bilinear behavior. The transverse compression strength of the fictional constituent ply cannot be directly determined from these samples. The maximum load in the compression test was used to determine the longitudinal compression strength, by considering that the 90° plies had failed. Finally, the shear strength was determined from the coupons cut out from the molded plaque at 45° as discussed later. The measured strength coefficients of the fictional constituent ply in the sub-plies model, using these cross-ply laminates, are shown in Table 1. page_225 Page 226 Table 1: Equivalent on-axis strength coefficients of the fictional constituent sub-ply X (Mpa) X' (Mpa) Y (MPa) S (Mpa) 548.7 306.44 23.74 40.49 4.2 Measurements Using Unidirectional Composite Using unidirectional coupons, the tensile, compressive and shear strength coefficients X, Y, X', Y', S of the unidirectional ply of the composite were also determined and are shown in Table 2. Table 2 : The on-axis strength coefficients of the unidirectional composite X (Mpa) X' (Mpa) Y (MPa) Y' (MPa) S (Mpa) 733.09 386.78 24.74 115.35 36.15 In the above Table, the ultimate shear strength S was obtained from tensile test of [(-45/45)2]s specimens, cut out in 45° from the [0/90]s plaques. It is well known that there are four generally accepted test methods for evaluating the lamina in-plane shear properties. They are the off axis coupon test, the [(45/-45)2]s coupon test, the rail shear test, and the torsion test [18]. The most simple tests to perform are the tensile test using the off axis and the [(45/-45)2]s coupons. The results of the off-axis test indicated that the ultimate shear strength determined by the 15-degree test and the [(45/-45)2]s tensile test compare quite favorably [22]. Chamis and Sinclaire have also recommended the 10-degree off-axis test for determining lamina shear properties [23]. The above results show that the longitudinal tensile strength coefficients X and X' of the fictional constituent ply of the woven fabric, shown in Table 1, are much lower than that of the unidirectional composite. This could be explained by the undulation effect of fiber in the fabric system. In the case of longitudinal tensile strength, since the degree of undulation of the fibers could not all be the same, fibers with less undulation would attain the ultimate strain first. Consequently, the volume fraction of fibers sustaining the load at fracture would be lower than that in a unidirectional composite. It is also surprising to note that the transverse strength Y is the same in both systems. The result suggests that the warp fibers do not affect the transverse properties of the weft fibers and vice versa. 4.3 Measurements Using the [(0/902]s and [(45/-45)2]s Specimens Made From Unidirectional Composite The stress-strain curve of the tensile test on [(0/90)2]s also showed the knee point due to matrix failure in the 90° layers. This point occurs at about 25% of the ultimate load, resulting from fracture of the 0° plies. From these tests, the strength coefficients were determined in the same manner as that in the case of woven fabric composite. The compression test did not show a knee point and only the longitudinal compression strength could be determined. The results are reported in Table 3. Table 3: Equivalent on-axis strength coefficients of the unidirectional composite determined from [(0/90)2]s coupons. X (Mpa) X' (Mpa) Y (MPa) S (Mpa) 693.68 363.90 31.72 36.15

page_226 Page 227 It can be seen that the transverse strength, as measured by cross-ply samples, is higher than that determined from unidirectional coupons. This can be explained by the constrain effect, induced by the adjacent plies, as reported previously for PEEK/Carbon composite [35]. The slightly lower longitudinal tensile strength on Table 5 is probably due to more significant misalignment of fibers, induced by hand lay-up of the cross-ply plaques. The shear strength coefficient S is slightly lower than that of the constituent ply in the woven fabric. This could be due to the weaving effect of the fibers in the fabric system. 5 Determination of Failure Envelop of Deformed Woven Fabric Composites In a laminated composite subjected to loading, the ply with the lowest strength ratio will fail first. This initial failure defines the inner failure envelope of the composite. Once the applied stress exceeds the initial failure, a laminate may or may not be able to sustain additional load. Various approaches to rationalize the lower limit and the ultimate strengths for laminated composites have long been published. In these approaches, matrix degradation models have been proposed to distinguish the plies with cracks due to initial failure from the intact ones [17, 25, 32]. A ply or ply group with cracks will change the internal stress distribution of the laminate. A ply with transverse cracks is not the same as having it completely removed from the laminate. The effective stiffness of the laminate will be reduced, but not as much as if the whole ply is removed. Because conventional stress analysis like laminated plate theory is limited to a continuum or plies without cracks. It has been proposed [17] to replace the cracked plies with a continuum of lower stiffness coefficients so that the conventional stress analysis can be applied. In fact, the observed stiffness of the laminate having partially and totally degraded plies is used to estimate the degree of the ply stiffness reduction. The replacement of cracked plies by quasi-homogeneous plies is done semi-empirically. The onset of first ply failure and the ultimate strength of the laminate were determined using the composite laminate design software Genlam [17]. In the calculation, the strength coefficient determined above were used, and both the maximum strain and the quadratic criteria were applied. The onset of first ply failure of the woven fabric composite was determined from the knee points in the stress-strain curve. The ultimate strength was determined using the procedure proposed in [17] to simulate the gradual degradation of the laminate due to the failure of successive plies with increasing loading. The results showed that the quadratic criterion provides a better prediction at both onset of damage and at ultimate failure of the fabric composite. Figures 2 and 3 present the comparison between prediction using the sub-plies model and experimental measurements. The onset of damage (first cracking) and ultimate tensile strengths of the plain weave laminates deformed to various angle q between the warp and weft threads. The results of experimental measurements are in good agreement with the prediction from the sub-plies model. It is also clear that prediction is much better when the fiber undulation effect is considered, that is, when the strength coefficients determined from the coupons made of the woven composite are used. As discussed above, the transverse compression strength of the fictional constituent ply of the fabric could not be determined from the compression test. The Y' coefficient of the unidirectional composite was used in the preduction. From these figures, it can be seen that the fiber undulation effect is more important when the angle is less than 45°. When the angle is more than 45° the fiber undulation effect becomes less significant. This might be due to the fact that in this region, the fibers are more straightened by the interlacing structure. Figure 4 presents the ultimate compressive strength of deformed laminates as a function of the deformation angle q. In this case, the effect of fiber undulation on the ultimate compression page_227 Page 228 strength of the fabric composite appears to be less significant than that observed for the strength coefficients of the constituent ply in the sub-plies model. The ratio of ultimate tensile strength to the onset of damage determined by the first knee-point (Pu/Po) also varies as a function of the angle q and is shown in Figure 5. When q < 30° the ratio of strength Pu/Po decreases rapidly because the decreasing ultimate tensile strength is accompanied by the increasing strength at first knee-point failure. In this case, the specimens failed mainly under longitudinal stress. When 30° < q < 45° , Pu/Po increase with q, up to a ratio of about one. Since shear failure mechanism is predominant in this region and is maximum at 45°, it can be responsible for the extent of damage in the sample before ultimate failure. When q > 45° the ratio of strength Pu/Po reaches a plateau value (about one) and the failure process is mainly due to the transverse stress. The details of damage progression analysis will be presented in another paper.

6 Conclusion The failure analysis and strength prediction of any deformed woven fabric laminates can be predicted by the sub-plies model. The results of the prediction are in good agreement with the experimental data. The fiber undulation has a strong effect on both transverse cracking and ultimate failure. The effect of fiber undulation can be taken into account by measuring the on-axis strength coefficients of the fictional constituent sub-plies directly on the samples made of woven fabric composites. 7 References 1. S. W. Tsai and H. T. Hahn, '' Inelastic behavior of composite materials, Proc. 1975 ASME Winter Annual Meeting, Houston, TX, Nov.1975, p.73. 2. Nishimura, A., "Proc. 10 th Symp. Mater. Aerospace Tech., Tokyo, Japan, 1980 (in Jap.) 3. Ishikawa, T. and Chou, T.-W, Journal of Material Science, Vol. 17, 1982, p. 32113220. 4. Ishikawa, T. and Chou, T.-W, Proceeding of the 4th International Conference on Composite Materials, Vol.4, 1982, p. 489496. 5. Chou, T.-W. and Ishikawa, T., Textile structural composites, Composites materials series, Vol.3, T.-W. Chou and F.Kko., R.B. Pipes, Series Ed., Elsevier, Newyork, 1989. 6. Chou, T.-W., First conference on Composite Materials, A.S.C., Vol. 1, 1985. 7. Ishikawa, T. and Chou, T.-W., Journal of Composite materials, Vol. 16, 1982, p.219 8. Ishikawa, T., Fib. Sci. Tech., Vol. 15 (1981), p.2. 9. N.K. Naik, P.S. Shembekar, and M.V. Hosur, J. Composites technology and research, Vol.13 p.107116 (1991). 10. K.L Reifsnider and Farshad Mirzadeh, J. Composites technology and Research, Vol. 10 p.156164 (1988). 11. D.Laroche and T. Vu-Khanh, in Development and Design with Advanced Materials, p.223262, S.V. Hoa (Ed.), Elsevier science Publishers B.V., North Holland, 1992. 12. D.Laroche and T. Vu-Khanh, Composite Materials : Testing and Design, ASTM STP 1206, p.386414 (1993). 13. T. Vu-Khanh and B.Liu, Queen's Landing, Niagara on the lake, Ontario, Canada, May 1618, 1994. 14. T. Vu-Khanh and B.Liu, Fifth international Conf. On Marine Application of Composite materials, April 1214, Florida, USA (1994). 15. T. Vu-Khanh and B.Liu, Part I, Pro. 9th Conf. American Society for Composites, Dayton, OH, USA Sept. 2022, 1994. 16. Stephen W. Tsai and Victor D. Azzi, AIAA journal 1996, Vol. 4, No. 2, p. 296301. 17. S.W. Tsai "Composites design-fourth edition, "Thick Composites, 1988. page_228 Page 229 18. James M.Whitoney, Issac M.Daniel and R. Byron Pipes " Experimental mechanics of fibre reinforced composite materials", 1982 19. D. Hull and T. W. Clyne " An introduction to composite materials - Second edition" 1996 20. Adams, Donald F. "Test methods for composite materials", 1992 21. G. P. Sendeckyj, M. D. Richardson and J. E. Pappas, presented at the composite reliability conference, Las Vegas, Nevada, April 1974. 22. Pipes, R.B. and Cole, B.W., Journal of composite materials, 1973, vol. 7, No. 2, p. 246256. 23. Chamis, C.C. and Sinclaire, J.H., Experimental mechanics (1977), Vol. 17, No. 9, p. 339346.

24. Daniel, I. M., Air force technical report AFML-TR-76-244, part 1 (1976). 25. P.H. Petit and M. E. Waddoups, Journal of Composite materials, Vol.3, 1971, p.2 26. L. V. Smith and S.R. Swanson, Proceedings of the American society for composites, ninth technical conference, Sept. 1994. 27. A. K. Roy, Proceedings of the American society for composites, ninth technical conference, Sept. 1994. 28. A. K. Roy, Proceedings of the American society for composites, twelfth technical conference, Oct. 1997. 29. S.W.Tsai and H.T.Halin, Technomic Publishing Co. Lancaster, PA.USA,1994. 30. Kim, R. Y., Test methods and design allowable for fibrous composites, ASTM STP 734, C. C. Chanis, Ed., American society for testing and materials, 1981, p. 91. 31. M. Karayaka & P. Kurath, Journal of Engineering Materials and Technology, vol.116, April 1994, p.222. 32. Ohira, H. and Uda, N., Recent advances in composites in the United states and Japan, ASTM STP 864, J. R. Vinson and M. Taya, Eds., American Society for Testing and Materials, Philadelphia, 1985, p.110. 33. Naik, N. K. and Ganesh,V. K., Journal of composites technology and research, JCTERE, Vol.16, No.1, January 1994, p.3. 34. V. K. Ganesh and N. K. Naik, Journal of composites technology and research, JCTERE, Vol.19, No.2, 1997, p.77. 35. Wang H., Vu-Khanh T., Journal. Compos. Mater., Vol. 28, pp. 684707, 1994.

Figure 1: Stress-strain curve of ((0/90)n)s samples page_229 Page 230

Figure 2: Onset of damage in tensile samples as a function of interlaced yarns angle q

Quadratic criterion with fiber undulation effect; --- Quadratic criterion without fiber undulation effect; · Experimental data.

Figure 3: The ultimate tensile strength as a function of interlaced yarns angle q Quadratic criterion with fiber undulation effect; --- Quadratic criterion without fiber undulation effect; · Experimental data. page_230 Page 231

Figure 4: The ultimate compressive strength as a function of interlaced yarns angle q Quadratic criterion with fiber undulation effect; --- Quadratic criterion without fiber undulation effect; · Experimental data.

Figure 5: The ratio of ultimate tensile strength to the onset of damage determined by the first knee-point (Pu/Po) also varies as a function of the angle q page_231 Page 233

POSTER SESSION page_233 Page 235

Optimum Design of Composites with Functional Properties by Genetic Algorithm Akihiko GOTO*, Atsushi YOKOYAMA** *Osaka Sangyo University, Faculty of Engineering 3-1-1 Nakagaito, Daito, Osaka, JAPAN **Mie University, Faculty of Education Kamihamacho, Tsu, Mie, JAPAN Keywords: Genetic algorithm, Optimum design, Functional properties Abstract Genetic algorithms were applied for the optimum design of composites with functional properties. It is considered that the ply orientation angle of laminated composite is one of the functional properties. The angle is an important factor of the design of the composite because of keeping high mechanical properties. Firstly it is attempted to design laminated composite materials by the algorithm. It was investigated the methodology and the utility of the algorithm for searching the optimum solution in the earlier period on the design of composites. 1 Introduction Genetic algorithm is known to one of the effective method for searching the optimum solution. The algorithm is the engineering model based on processes of biological evolutions. For searching the solution, one of many design parameters is selected and is modified the binary code. Population of the coded parameter of the algorithm is carried out the selection and the crossover and so on. Genetic algorithms have been applied to many fields of optimization problems as techniques and tool that will more effectively search the parameter for the best design (16). The methodology for the structural design is assumed to be included into these application fields. In general, the optimum design of polymeric composites need to be carried out several simulation for trial and error, so that it takes many times step by step for the estimation by human. However, by using the genetic algorithm, the design is possible to be carried out automatically and to be found the optimum solution in the short time.

In this paper, firstly, we noticed the ply orientation angle of laminated composite which was one of many parameter for designing polymeric composite materials because of modifying each ply to binary code easily. It was attempted to apply the genetic algorithm to the optimum design of the simple plate of laminated composites. It was investigated the methodology and the utility of the algorithm for searching the optimum solution in the earlier period on the design of laminated composites. 2 Analysis Method 2.1 Object Two kinds of carbon fiber reinforced composite laminates as the design object were employed. Table 1 shows material constants of CFRP used in this analysis The number of total ply of laminate was 4 and 8 respectively. However, the composite was consisted of symmetric ply orientation angle, so that half ply was used to calculate the optimum design by the algorithm. The former was called" 2 plies", the latter was called" 4 plies". It was assumed that thickness of each layer was the same. The object model is showed in figure 1. page_235 Page 236 Table 1 Material constants of CFRP used in this analysis. Longitudinal modulus 105.0 GPa Transverse modulus 9.8 GPa Shear modulus 6.0 GPa Poisson's ratio 0.32 Longitudinal tensile strength 1500.0 MPa Transverse tensile strength 40.0 MPa Shear strength 68.0 MPa

Figure 1 The object model in the case of 4 plies. 2.2 Analyzing Procedure Flow chart of the genetic algorithm is showed in figure 2. Parameters of the genetic algorithm include population size (ps), probability of reproduction (pr), probability of crossover (pc), probability of mutation (pm). The ply orientation angle on each layer of laminated composite was replaced three bits parameter as binary code. Table 2 shows binary codes of each ply orientation angle. The population size meant the number for searching a optimum solution. One of population expressed coding information of analysis object. It was selected two population among them, so that new population was generated by the crossover. Mutation was varied gene of population by a constant probability. Figure 3 shows coding information of each process. For example, laminated composite with four plies require the twelve bit parameters to determine the whole ply orientation angle in laminated composite. Standard condition to search the optimum solution the genetic algorithm was determined as follows; ps = 30, pr = 0.0, pc = 0.4, pm = 0.03. Effects of genetic algorithm parameters were estimated by varying the condition. Each varied parameter is shown in table 3. Probability of reproduction was fixed in zero value. Maximum generation was set one hundredth. One parameter was exchanged from standard condition, while others were kept value of standard. Nine kinds of conditions were employed. Moreover, as there were three kinds of primitive random values on each condition, twenty-seven kinds of conditions were carried out. The population was searched the optimum solution by Tsai-Hill theory, so that the population was regenerated for next generation. The reproduction, the crossover and the mutation were carried out till the maximum generation. The criterion of Tsai - Hill theory was given by,

where s||s^ and t were expressed the normal stress of longitudinal axis, the normal stress of transverse axis and shear stress respectively. ' * ' was signified the stress occurring the failure in the laminate. page_236 Page 237 Table 2 Binary codes of each ply orientation. Analysis code Ply orientation angle (deg.) 90 [ 000 ] 60 [ 001 ] 45 [ 010 ] 30 [ 011 ] 0 [ 100 ] -30 [ 101 ] -45 [ 110 ] -60 [ 111 ]

Figure 2 Flow chart of the genetic algorithm.

Figure 3 Coding information of each process. Table 3 Parameter of genetic algorithm. Parameter Value Population size ( ps ) 10, 20, 30, 50

Probability of crossover ( pc ) Probability of mutation ( pm )

0.2, 0.4, 0.6, 0.8 0.01, 0.03, 0.10

3 Analyzing Results and Discussion Variations of minimum objective function value with increasing generation are shown in figure 4. Each curve was calculated the average of results by three kinds of primitive random values. The optimum solution can be searched in the initial period with increasing the population size. The population size of 50 is the best of all. The probability of crossover should be increased for searching the optimum solution in the earlier generation. Moreover the probability of mutation should page_237 Page 238 be increased similar to the above. From these results, it was considered that parameters of the genetic algorithm should be higher value in order to search the optimum solution in earlier period. Therefore the optimum condition for searching the solution in the earlier period was determined as follows; ps=50, pc=0.8, pm=0.1.

Figure 4 Variations of minimum objective function value with increasing generation. Figure 5 shows comparisons of minimum objective function values with the standard condition and the optimum condition. It was tendency that the optimum condition converged the solution in the earlier generation than the standard condition. States of searching the optimum solution both of the standard condition and the optimum condition in 2 plies are shown in figure 6. The population of the standard condition didn't attain to the optimum solution in zero and first generation. However with increasing the generation, the population was tend to converage the solution. In fiftieth and one hundredth generation, most of population found out the optimum solution. Even if not found out the solution, the population searched the solution of 0 deg. On the other hand, in the case of the optimum condition, the population has already found out page_238

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Figure 5 Comparisons of minimum objective function values with the standard condition and the optimum condition. the optimum solution in the first generation. In the fiftieth generation, the more part of population attained the solution. However the number of the population was smaller than that of the standard condition. The other population converged solutions with including 0 deg. In the searching state of the one hundredth generation, the number of the population converged the optimum solution was almost the same as the fiftieth generation. It was considered that values of each parameter of genetic algorithm might be increased for searching the optimum solution as soon as possible in the earlier period.

Figure 6 State of searching the optimum solution both of the standard condition and the optimum condition. page_239 Page 240

Figure 6 Continued. 4 Conclusions The method of genetic algorithm was attempted to design carbon fiber reinforced composite laminates with high mechanical properties against tensile stress. It was determined the ply orientation angle of the laminate as the optimization. We obtained conclusions as follows; For searching the optimum solution as soon as possible in the earlier period, values of each parameter of genetic algorithm might be increased. However it is difficult for the most of population to converge the optimum solution. It is considered that the method of the algorithm is applied the prediction of failure ares and the optimum design took account into several functional properties. The utility of the algorithm and variations of states converging the optimum solution will be investigated on the problem. Reference 1. David L.Calloway, Inter. Conf. Genetic Algorithms, 4, 422 (1991) 2. Donald R.Jones and mark A.Beltramo, Inter. Conf. Genetic Algorithms, 4, 442 (1991) 3. Sushil J.Louis and Gregory J.E.Rawlins, Inter. Conf. Genetic Algorithms, 4, 53 (1991) 4. Prasanna Jog, Jung Y. Suh and Dirk Van Gucht, Inter. Conf. Genetic Algorithms, 3, 110 (1989) 5. J.David Schaffer, Richard A. Caruana, Larry J.Eshelmen and Rajarshi Das, 3, (1989) 6. Terence C. Fogarty, Inter. Conf. Genetic Algorithms,3, 104 (1989) page_240 Page 241

Curing of Thick Angle-Bend Thermoset Composite Part: Curing Cycle Effect on Thickness and Fiber Volume Fraction Variation Malak I. Naji and Suong V. Hoa Concordia Center for Composites, Concordia University Montreal, Quebec, H3G 1M8 CANADA Keywords: Autoclave Cure Cycle, Thick-Angle-Bend Composite Structures. Abstract. In this paper, the effect of cure cycle on the quality of angle-bend thermoset parts is presented. A 2-D finite difference scheme was used to model the processing of arbitrary shaped composite laminates. This model includes heat transfer, viscosity, rate of cure, and resin flow submodels. Experiments were done to investigate the effect of different curing cycles on thickness variation and fiber volume fraction distribution: Results were obtained from both simulation and experiment. 1 Introduction Many laminated composite structures, such as an angle bracket, a co-cured web, or a frame have a curved portion. The final failure in such structures may be a complex progression of ply cracking, delamination, and fiber failure. More recently, efforts at measuring interlaminar tensile strength have been focused on the curved-beam test methods (C-shaped and L-shaped), because of the difficulties in introducing loading onto the flatwise specimens. Because curved thick laminates (especially with small curvature) are very difficult to manufacture with consistent quality, semicircular specimens are of poor quality with a high void content and thickness variations. Furthermore, interlaminar strength is degraded by manufacturing and processing defects such as porosities, resin pockets, and resin-rich interlayers. As composite components increase in complexity, there is a higher possibility of porosity existing in the structure. It is well known that porosities in composites, either artificially induced or processing caused, can result in significant reduction in structural strength and life. Hence, large variations in part thickness and fiber volume fraction, presence of resin-rich layers or voids, and presence of incomplete cure spots would indeed affect the final part quality. There have been many studies on the curing of thick thermosetting composite materials. In Ref's [1, 2, 3], researchers developed cure simulation models that predict temperature and degree of cure distributions within the part as a function of the autoclave temperature history. Researchers who considered the resin flow model include Hojjati [4] who studied curing of thick flat section parts, and Johnston et al [5] who showed, for curved sections, that

resin flow was uneven resulting in non-uniform part thickness. Also, [5] found that springback angle is influenced by many factors as: choice of tool material, bagging condition, laminate thickness and surface friction condition. Many other workers as in [6, 7, 8] also performed studies to come up with optimized curing cycles for their composite parts. Hence, it is desired to gain a fundamental understanding of the curing process unique to thick angle-bend thermosetting composites parts. This is because gradients in temperature page_241 Page 242 and degree of cure are strongly dependent on part geometry, thermal anisotropy, the chemical cure kinetics and the thermal boundary conditions (curing cycle and tooling). Furthermore, dimensional changes which result from the anisotropic properties of the composite part, lead to a phenomenon called 'Spring In' which is present in sharply curved (90°) thick laminates. In this work, a 2-step and 3-step curing cycles were implemented experimentally and theoretically to study the variations in thickness and fiber volume fraction of an angle-bend structure made from graphite/epoxy Hercules AS4/3501-6 prepreg (ply thickness is approximately 0.16 mm and has 36% resin content by weight). The objective is to study the part quality in terms of fiber volume fraction and thickness variation throughout the part length. 2 Problem Formulation A two-dimensional cure simulation code was developed using finite difference method to predict the temperature, degree-of-cure, viscosity, and fiber volume fraction distribution as a function of cure time. Four submodels that are needed to obtain the governing equations describing the fundamental mechanism in curing are described below for a general 2-D problem shown in Fig. 1. I.Heat Transfer Equation

with top surface and sides exposed to a convective boundary condition, while the bottom follows the curing cycle temperature. Here, T is the temperature of the composite, t is the time, the q term represents the internal heat generation (defined below in Eq. 5, and r and cp are the composite density and specific heat, respectively. They are calculated using rule of mixture. kxx, kzz and kxz are the coefficients of effective anisotropic thermal conductivities. The coordinate directions in Eq. 1 are defined in a fixed global (x z) coordinate system, and are related to the composite in its principle (13) material coordinate system via the transformation equations found in Ref [9]. Composite thermal conductivity k11 in the fiber direction is calculated using rule of mixture and that perpendicular to the fiber direction k33 is calculated using Springer and Tsai model [10]. II.Resin Flow Equation according to Dave et al [11]

The top surface and sides having a boundary pressure equal to bag pressure, while the bottom surface has no flow crossing its boundary. P is the hydraulic resin pressure within the pores, Pf is the fiber pressure, vf is the fiber volume fraction, m is the viscosity. Sxx, Szz and Sxz are the coefficients of effective anisotropic specific permeabilities, which depend on the stress level. The off-axis permeabilities are defined in the same way as the off-axis conductivities, while on-axis permeabilities (S11 and S33) are calculated using Van der Westhuizen and Du Plessis [12, 13] expressions:

page_242 Page 243

mv is the coefficient of volume change which describes the stress-strain behavior of a body in confined compression. For a porous medium, it is the ratio of change in porosity to axial stress. To simulate composite processing, Dave et al [11] used a stress-strain relation which was based on experimental data by Gutowski [14]. The same relation will be used in this work and is given below.

III.Thermo-Kinetic Equation

subjected to the initial condition a(x,z) = 0 at t = 0. HR is the total heat of reaction, vr is the resin volume fraction, and a is the degree of cure. A1, A2 and A3 are pre-exponential factors, and E1, E2 and E3 are activation energies. R is the universal gas constant. IV.Viscosity Equation [15]

where U is the activation energy and Km is a constant. Material constants and properties for fiber and resin are shown in Table 1. 3 Experimental Work An experiment was done in order to investigate the manufacturers recommended curing schedules. The mold was machined from a a 25mm thick aluminum plate with 90°C angle and 165mm ´ 153mm sides. The tool surface is covered with non-porous Teflon ply and is shown in Fig. 1. Insulator dams were installed in the y-direction to restrict resin and heat flows. This was done using one-inch thick fire brick stone. A charge consisting of 50 plies (each 254mm long 100mm wide) of Hercules AS4/3501-6 prepreg was hand-layed unidirectionally and was vacuum bagged in a standard manner. The laminate stack was covered from the top with a porous release film, 3 layers of bleeder cloth, and then wrapped with a breather cloth. Finally, a vacuum bag was wrapped all over the mold and sealed with vacuum tape. A vacuum line was fed into this outer layer before inserting the charge into the autoclave. Referring to Fig. 2, the prepreg material was cured according to the recommended standard part 2-step cure schedule with: T1 = 115°C, t1 = 60 min, T2 = 177°C and t2 = 120 min. A pressure of 0.584M Pa (85 psi) was applied initially, and a vacuum was drawn on the bagged charge. A second, modified 3-step curing schedule suggested by the manufacturer for thicker parts, was also applied with: T1 = 115°C, t1 = 120 min, T2 = 150°C, t2 = 90 min, T3 = 177°C and t3 = 240 min. The same pressure and vacuum were drawn on the bag as above. Five samples were processed in the autoclave using both curing cycles. The heating rate was noted to be 4.5°C/min which is more than the recommended value of 3°C/min. The first sample was disregarded due to electrical errors during operation which cause the maximum temperature not to exceed 155°C. The other four samples were prepared as follows: page_243 Page 244 # 2 The 50 layers were layed with debulking after each 10 layers, and the sample was cured using the 2-step curing cycle. 1 layer of porous release film is applied. # 3 The 50 layers were layed with debulking after each 5 layers, and the sample was cured using the 3-step curing cycle. 1 layer of porous release film is applied. # 4 The 50 layers were layed in one shot, and the sample was cured using the 2-step curing cycle. 2 layers of porous release film are applied to restrict resin flow. # 5 The 50 layers were layed with debulking after each 10 layers. The sample was cured using the 3-step curing cycle. 2 layers of porous release film are applied to restrict resin flow.

Fig. 3 shows part thickness variation at the sections indicated with angle f as defined in the figure, and l is a non-dimensional scale (on the thickness direction) that is=0, 1/2 and 1 at the bottom (mold surface), middle and top surfaces, respectively. Table 2 lists the enclosed angle of the final part. The four samples were cut width-wise and across-wise to examine the microscopic structure. Five cuts were made cross wise in samples #4 and #5 to find the fiber volume fraction values. The sections were made at points 2, 4, 6, 4' and 2' and the results at l=1/2 are listed in Table 3. Also, Fig. 4 shows how vf varies across the thickness for sample # 4 (the same trend was observed for sample # 5). From this examination, it was noticed the following: - All samples show a resin-rich layer especially between successive layers at section 6. The frequency of this layer was the least in sample #3, and was the most in sample #4. This is clear by looking to the photos shown in Fig. 5. - Looking to Fig. 4 (for sample # 4), the trend in vf variation from bottom to top is the same for all sections. High vf at the top is due to the presence of bleeder, and low vf to the bottom is due to the presence of mold surface. However, this variation was larger atsection 6 (curved bend) and is demonstrated with photos shown in Fig. 6. - When examining the curved bend, no wrinkling of fibers was noticed. - The maximum difference in the angle (between the mold and the cured sample) was less than 1 degree (0.77 degree). 4 Cure Simulation Running the simulation code for the two curing cycles described in the previous section, the following was obtained: - The difference in part thickness was not large when the part was cured with 2-step and 3-step cycles with 2-step cycle giving higher values than 3-step cycle. This was shown in Fig. 3 along with the experimental results. - Thickness values obtained from model were in good agreement with those obtained from experiment. The difference was only in the straight section which may be due to the pressure boundary conditions implemented in the model. - The same argument holds for the vfm values since thickness and volume fraction are related to each other. 5 Conclusion In this work, thickness and fiber volume fraction variations were investigated for an L-shaped composite part. Both experiment and simulation show reasonable agreement in terms of thickness and fiber volume fraction values. They showed how the quality of the part is not consistent page_244 Page 245 from section to section. Furthermore, they also demonstrated the need for a modified curing process that would produce uniform part thickness and fiber volume fraction across the length. The enclosed angle measurements showed that in this particular case for AS4/3501-6 material, the 'spring in' was less than 1°. This can be reffered to mainly three reasons: (1) The relatively thick material being processed, (2) The type of releasing film used for the mold (Teflon), and (3) The material system used which is used largly in aerospace industry. Johnston et al [5] showed that increasing part thickness and decreasing surface friction tends to decrease the 'spring in' angle. References [1] A. O. Kays, ''Exploratory development on processing science of thick section composites," tech. rep., AFWALTR-85-4090, Air Force Wright Aeronautical Laboratories, Wright Patterson AFB, OH, 1985. [2] T. E. Twardowski, S. E. Lin, and P. H. Geil, "Curing in thick composite laminates: Experiment and simulation," J. Composite Materials, vol. 27, pp. 216250, 1993. [3] T. A. Bogetti and J. W. Gillespie, "Process-induced stress and deformation in thick section thermosetting composite laminates," in 21t SAMPE Technical Conference, New Jersy, Sept. 1989. [4] M. Hojjati, Curing of Thick Thermosetting Composites: Experiment, Simulation, and Scaling. Ph.D dissertation, Mechanical Engineering, Concordia University, Montreal, Canada, 1994.

[5] A. Johnston, P. Hubert, G. Fernlund, R. Vaziri, and A. Poursartip, "Process modeling of composite structures employing a virtual autoclave concept," Science and Engineering of Composite Materials, vol. 5, pp. 235252, 1996. [6] L. N. Hjellming and J. S. Walker, "Thermal curing cycles for composite cylinders with thick walls and thermoset resins," J. Composite Materials, vol. 23, pp. 10481064, 1989. [7] P. R. Ciriscioli, Q. Wang, and G. S. Springer, "Autoclave curing - comparisons of model and test results," J. Composite Materials, vol. 26, no. 1, pp. 90102, 1992. [8] J. S. Kim and D. G. Lee, "Development of an autoclave cure cycle with cooling and reheating steps for thick thermoset composite laminates," J. Composite Materials, vol. 31, no. 22, pp. 22642282, 1997. [9] S. W. Tsai and H. T. Hahn, Introduction to Composite Materials. Technomic, 1980. [10] G. S. Springer and S. W. Tsai, "Thermal conductivity's of unidirectional materials," J. Composite Materials, vol. 1, pp. 166173, 1967. [11] R. Dave, J. L. Kardos, and M. P. Dudukovic, "A model for resin flow during composite processing: Part 1-general mathematical development," Polymer Composites, vol. 8, pp. 2938, 1987. [12] J. V. der Westhuizen and J. P. D. Plessis, "Quantification of unidirectional fiber bed permeability," J. Composite Materials, vol. 28, pp. 619637, 1994. [13] J. V. der Westhuizen and J. P. D. Plessis, "An attempt to quantify fiber bed permeability utilizing the phase average navier-stokes equation" Composites. part a, applied science and manufacturing, vol. 27, pp. 263269, 1996. page_245 Page 246 [14] T. G. Gutowski, T. Morigaki, and Z. Cai, "The consolidation of laminate composites," J. Composite Materials, vol. 21, pp. 172188, 1987. [15] W. I. Lee, A. C. Loos, and G. S. Springer, "Heat of reaction, degree of cure, and viscosity of 35016 resin," J. Composite Materials, vol. 16, pp. 510520, 1982.

Fig. 1. The general curved laminate showing the x z (global) and x - h (local) laminate system of coordinates, and the 13 (principal) ply system of coordinates. Fiber radius Fiber density Specific heat of fiber Thermal conductivity of fiber Resin density Specific heat of resin Thermal conductivity of resin Pre-exponential factor Pre-exponential factor Pre-exponential factor Activation energy Activation energy Activation energy Heat of reaction Activation energy for viscosity

rf 3.5 ´ 10-6m rf 1.79 ´ 103kg/m3 cpf 7.12 ´ 102J/kg.K kf 26 W/m.K rr 1.26 ´ 103kg/m3 cpr 1.26 ´ 103J/kg.K kr 0.167 W/m.K A1 2.101 ´ 109min-1 A2 -2.014 ´ 109min-1 A3 1.96 ´ 105min-1 E1 8.07 ´ 104J/mol E2 7.78 ´ 104J/mol E3 5.66 ´ 104J/mol HR 4.74 ´ 105J/kg U 9.08 ´ 104J/kg

Viscosity constant m¥ 7.93 ´ 10-14Pa.s Viscosity constant Km 14.1 Table 1. Material properties of Hercules AS4/3501-6 Part # 2 3 4 5 Enclosed angle (°) 89.35 89.58 89.23 89.33 Table 2. Enclosed angle results page_246 Page 247

Fig. 2. Processing cure cycles used in 2-step and 3-step cycles, respectively

Fig. 3. Experimental results for thickness variation at different sections. Section #

ufe % ,ufm % 2-step 3-step 2-step 3-step #4 #5 2 64 62 64 66 4 59 59.5 60.5 62 6 57.5 58 58.7 59.4 4 60 62 60.5 62 2 62.5 63.5 64 66 Table 3. Experimental (ufe) and Numerical (ufm) fiber volume fraction values for different sections at l=1/2

Fig.4. vf variation across the thickness obtained from experiment for sample # 4. page_247 Page 248

Fig. 5. Length-wise photomicrographs (50X) at section 6 and l=1/2.

Fig. 6. Cross-section photomicrographs (600X) for sample # 4. page_248

Page 249

In-Situ Cure Monitoring of Graphite/Epoxy Composites Using Fiber Optics and Ultrasonics J.-Y.CHEN1, S.V.HOA1, C.-K.JEN2 AND H.WANG2 1 Introduction Advanced composites can be fabricated by laminating multiple prepreg plies into the desired shape and then cured in an autoclave with simultaneous application of proper heat and pressure. The knowledge of cure process is very important in order to obtain fully cured and high quality composites at reduced production costs [1]. In-situ sensors capable of monitoring the cure process are therefore desirable. The thermal Differential Scanning Calorimetry (DSC) can monitor the exothermic flow of heat of the cure reaction, and a characterization of the state of cure of composites, thus, is possible. For the AS4/3501-6 prepregs, Figure 1 shows the degree of cure (DOC) as a function of cure time with different cure temperatures. As the cure temperature increased, the DOC increased while the time for completion of the cure reaction decreased. At the cure temperature of 176°C, the cure reaction was almost fully complete after 60 minutes, while at 146°C cure, the DOC only reached around 70% after 120 minutes. However, even though the cure reaction was nearly completed, the mechanical properties of the curing composites may not be well developed because the DOC is a process parameter which gives information on the extent of chemical cure reaction while the extent of modulus is a parameter which gives information on the degree of mechanical property development [2]. The critical objective of the process engineer is to know when the material has been "fully processed", and the cure characterization should be reflective of the ultimate material application, thus, the "end-of-cure" should be determined by the completion of the development of not only chemical properties but also mechanical properties of curing composites. Ultrasonic techniques for cure monitoring have been reported in the literature [28] because it can provide a direct and nondestructive measurement of the viscoelastic properties of a curing composite. However, there are a number of limitations. Specifically, the commonly used piezoelectric ultrasonic transducers (UTs) can only be operated continuously at an upper temperature of about 60°C. This limitation presents difficulties when applying the techniques in actual processing environment. Micromet Instruments, Inc. (Newton Center, MA), has developed an ultrasonic cure monitoring system, which can work with an autoclave by using high-temperature UTs with narrow-band characteristics [9]. The system has been successfully used for some applications on cure monitoring of composite processing, but only in through-transmission mode. For a comparison, a new cure monitoring system being developed by using conventional broadband UTs with clad buffer rods [10] and Extrinsic Fabry-Perot Interferometric (EFPI) sensors will be presented in this paper. The evolution of viscoelastic properties of the curing composite was monitored by continuous recording of the time delay, 1 Dept. of Mech. Eng., Concordia Univ., 1455 de Maisonneuve Blvd. W., Montreal, Quebec H3G 1M8 2 Industrial Materials Institute, NRC, 75 de Mortagne Blvd., Boucherville, Quebec, J4B 6Y4 page_249 Page 250 attenuation and shear reflection coefficient of ultrasonic waves. The EFPI sensors were concurrently used to detect the local strain development in the composite caused by temperature variation, chemical shrinkage and processinduced stresses. The results obtained from the two sensors were compared in order to evaluate their suitability for in-situ monitoring of autoclave processing. Discussions are also made on using the techniques to detect the "end of cure", relative to the traditional "degree-of-cure" method using the DSC technique. 2 Experimental Setup The in-situ cure monitoring system developed in the study consisted of fiber-optic, ultrasonic sensors, and the associated electronics. The system was incorporated into an autoclave via a special feed-through. The material was Hercules AS4/3501-6 composite. 2.1 Fiber-Optic Strain Sensor

A four-channel WFS-100 EFPI fiber-optic strain sensor system, developed by FISO Technologies, Inc. (Quebec City, Canada), was used for the in-situ cure monitoring. The optical fibers were embedded in the composite panels, and aligned with the direction of graphite fibers, so as to minimize the gage disturbance to the host composite and to achieve best load transfer between them. 2.2 Ultrasonic Sensor The ultrasonic sensing system was composed of high performance 5 MHz broadband piezoelectric UTs, couplants, clad buffer rods [10] and air cooling systems which surrounded the UTs, as shown in Figure 2. This system can be operated in either pulse/echo and/or through-transmission mode. The clad buffer rods consisted of aluminum core and a thermal sprayed aluminum cladding. These buffer rods were screwed into the mold, with the end faces flush with the mold inner surface. The principle of the ultrasonic guidance in clad buffer rods is like the clad optical fibers in which the energy is guided in the core, and thus the threading in the cladding region does not disturb the guided ultrasonic energy inside the core. The variation of the amplitude of the reflected ultrasonic signals at the rod/composite interface and those of the amplitude and time delay of signals traversing back and forth through the composite thickness enabled us to monitor the cure. The 100 psi compressed air cooling allowed the cooling of the UTs from 190°C down to 60°C which is the maximum allowable temperature for the UTs used. Ultrasonic waveforms were recorded with a sampling rate of 100 MHz in every 5 or 10 minute period during the entire cure process. 2.3 Panel Processing The commercially recommended cure cycle was used [11], i.e. (1) phase I: the prepreg was debulked under vacuum while the temperature was raised to 116°C and held there for one hour; (2) phase II: the temperature was then raised to the cure temperature of 176°C and held there for 2 hrs. (the length of this dwelling time was called cure time); and (3) cool-down phase: the cured panel was then slowly cooled to room temperature. The autoclave air pressure was kept constant at 85 psi during the entire cure process to secure prepregs page_250 Page 251 compaction and good thermal conduction. The temperature of the panel was monitored by a K-type thermocouple embedded into the panel. For typical cases, effect of cycle conditions on composite cure was evaluated by using the cure temperature of 146°C or 156°C, or by extending cure time from 2 to 3 hrs.. 3 In-Situ Cure Monitoring 3.1 Ultrasonic Pulse/Echo Mode In the pulse/echo mode, the ultrasonic wave traveling through the material is reflected from the back side of the laminate to the originating surface where it is detected by the same UT that generates the pulses. Figure 3(a) shows the ultrasonic waveform as a function of process time for the [016] panel cured at 176°C for 2 hrs., where L1 represents the echo reflected from clad Al buffer rod/top surface of panel interface and L2 the second echo reflected from the back surface of the panel. The recorded variations of process temperature and ultrasonic time delay between L1 and L2 are shown in Figure 3(b). For the 176°C case it is seen that, due to thermal expansion of the buffer rod, echo L1 shifted toward a higher time delay before the constant cure temperature was reached; it came to a standstill during the hold period at 176°C and then gradually returned to its original position at the end of the cool-down stage. The second echo L2, during the initial stage, was too weak to be recorded using the current setup. Until the gelation of the epoxy composite took place, it can be recorded. The time delay between the first echo L1 and the second echo L2 can be determined using a cross correlation algorithm. The sharp decrease in the time delay, in the temperature range after phase I, could be attributed to the gelation of prepregs. The subsequent gradual decrease of the time delay corresponded to the progressive crosslinking reaction of the composite after gelation period. Effects of the state of cure on the time delay can be, therefore, described as a continuous decrease of the time delay as the material changed from a liquid, to a gel and a solid due to increasing stiffness of the curing composite. After cure reaction, the panel started to cool down, and the time delay further decreased. After a certain time, a sudden rise in the time delay was found. This discontinuity resulted from the sudden detachment of the panel from the mold plate at the sensed region. This interpretation is supported by the observation in Figure 3(a) of a sudden phase change of 180° and an increase in amplitude of the echo reflected from the back surface of the panel at that moment. This is due to the change of the interface condition from the "panel-mold plate" to the "panel-air".

Figure 3(b) also shows the time delay recorded from two other runs using the cure temperatures of 146°C and 156°C, respectively. It is noted that increasing cure temperature shifted the curve of the time delay vs. process time to a well distinct lower value, which indicates that ultrasonic in-situ cure monitoring can monitor the change induced by different cure temperatures. 3.2 Ultrasonic Transmission/Reflection Mode In through-transmission mode, the ultrasonic wave passes through the laminate and is detected by another UT located at the opposite side of the mold. Meanwhile, these two UTs can still conduct reflection mode measurements individually. For these measurements of the page_251 Page 252 [016] panel cured at 176°C, the cure time was arbitrarily extended from 2 to 3 hrs., in order to observe the "end-of-cure" in terms of the full development of the stiffness of the curing composite. The variations of the time delay vs. process time obtained from both top UT and bottom UT, in reflection mode, exhibit the same trend, as shown in Figure 4(a). After cured at 176°C for 2 hrs., the time delay reached a plateau, suggesting that the "end-of-cure" has been reached and the extra hour of cure processing did not make significant contribution to the improvement of the stiffness of composites. Here, again, as the panel cooled down, the time delay decreased. After a certain time, because of detaching of panel from the top mold plate at the sensed region, the signal from the top UT completely disappeared, while the time delay obtained from the bottom UT showed a discontinuity. For through-transmission mode, the shape of the time delay variation is the same as those obtained from reflection mode, as indicated in Figure 4(a). After the detachment of the panel from either side of the mold plates, the through-thickness transmission signal was lost. Figure 4(b), obtained from the same cure cycle as in Figure 4(a), presents the variation of the through-thickness ultrasonic attenuation of the prepregs vs. process time during the cure. The variation in Figure 4(b) represents the basic features of the viscoelastic behavior of epoxy resin as it cures. At the beginning of gelation period, the attenuation reached a maximum point: the gel point, and then dropped dramatically during gelation period, indicating the rapid rise of viscosity due to accelerated crosslinking at the early stage of cure reaction. Then, the attenuation gradually decreased because, as crosslinking proceeded, the material behaved more as an elastic medium, with reduced absorption of ultrasonic energy. The plateau was reached when the constant viscoelastic properties of the composite have been developed at the end of cure. Finally, the attenuation further decreased as the material was cooled down. 3.3 Simultaneous Fiber-Optic and Ultrasonic Measurement In order to compare the cure monitoring capabilities of the fiber-optic and ultrasonic sensors, experiment was carried out on a [906/02/SG1, SG2/02/906] panel in which both techniques were operated simultaneously. Two fiber-optic and ultrasonic sensors were different but close to each other. As shown in Figure 5, the two fiber-optic strain gages provide almost identical records of the cure process before detaching of the panel from the mold plates occurred during the cooling stage. The strain gage-detected gelation period, as determined from the strain-temperature variation, agrees with the sharp drop of ultrasonic time delay which corresponds to the rapid increase of viscosity of epoxy at the early stage of cure reaction [12]. However, as observed earlier, the ultrasonic sensor was able to determine the "end-of-cure" where the viscoelastic properties of the composites were fully developed, while the fiber-optic sensor was not. The "end-of-cure" happened as the ultrasonic time delay reached a plateau. During the early cool-down stage, because the panel was attached to the mold plates, the free contraction of the panel was impeded and the thermal strain detected by the EFPI sensors dramatically dropped. After a certain time (D1), the sudden detaching of one side of the panel from the mold plate happened, its thermal strain sharply bounced back. Correspondingly, ultrasonic time delay exhibited a discontinuity (D2). Once more, when the page_252 Page 253 detaching of another side of the panel happened, the readings from the fiber-optic sensor experienced another small bouncing (D3), while the ultrasonic sensor totally lost its signal (D4). 3.4 Ultrasonic Shear Reflection Coefficient Measurement

It has been known for a long time that the shear reflectivity of ultrasonic waves from the interface between the probing end of the buffer rod and monitored curing resin is related to the viscoelastic properties of polymers [13, 14]. This could lead to a direct evaluation of the complex shear modulus and dynamic (high frequency) viscosity. In order to significantly eliminate the adverse effects caused by the UT, couplant and buffer rod on the ultrasonic time delay and amplitude measurements, minor physical discontinuity "a" (e.g. a notch) near the probing end of the clad buffer rod was created to provide a reference signal, as shown in Figure 6. Figure 7(a) shows the reflected 5 MHz shear waves from the curing composite. S1 is the shear wave echo reflected from the probing end of the buffer rod, while Sa is the signal induced by the discontinuity "a". The latter signal is used for normalization. The time delay between S1 and Sa is related to the local average ambient temperature between the discontinuity "a" and the end of the buffer rod. This information may be used to evaluate the temperature of the A1 mold and processed materials during cure processes. Figure 7(b) shows the time difference between the arrival of the echoes from the probe interface and the reference notch, plotted as a function of the process temperature. This time delay reflects the influence of process temperature on the shear modulus, Poisson's ratio, density and length of the buffer rod material that was between the reference notch Sa and the probing surface S1. If this temperature curve was well calibrated, it could be used for in-situ ultrasonic temperature monitoring during the cure processing of the composite. Figure 7(c) shows the variation in shear reflectivity from the probing interface with the process time, indicating that during the gelation period of the composite epoxy, the reflectivity dramatically dropped. Afterwards, the reflectivity decreased quite slowly as the cure reaction proceeded. During the cool-down stage, because of the influence of temperature, the reflectivity slowly increased. Then because of partial detachment of the panel from the A1 mold, a sudden increase in the reflectivity was observed. Obviously, the shear reflectivity of the curing composite is very sensitive to the development of gelation and the early stage of cure reaction of the composite epoxy. 4 Conclusion (1) Both the EFPI fiber-optic and ultrasonic sensors can sense the gelation period of curing composite and the detachment of the composites from A1 mold plates. Ultrasonic monitoring can sense the "end-of-cure", while the fiber-optic monitoring can not. The "end-of-cure" was observed when the ultrasonic time delay or attenuation reached a plateau. Also the timing information of the detachment of the cured panel from A1 mold may provide information concerning the development of residual stress during the cool-down phase, and can be used for tooling design. For the different cure temperatures 146°C, 156°C, and 176°C, increasing cure temperature accelerated the development of the stiffness of composites. page_253 Page 254 (2) The "degree of cure" measured by the DSC technique does not accurately reflect the mechanical property development during cure. Significant changes in the modulus may still exist when the "degree of cure" indicated by the DSC is almost fully developed. Acknowledgment The authors wish to thank FISO Technologies Inc. (Quebec City), for loaning the WFS-100 fiber-optic sensor system. Mr. H. Hebert of IMI, NRC for his designing of cooling system and Mr. S.-S. L. Wen of McGill University, Montreal, for writing LabVIEWâ software. The financial support from a strategic and an operating (STR0192858) grant of Natural Sciences and Engineering Research Council of Canada is also acknowledged. References (1) G.C. Sih. 1994. Advanced technology for design and fabrication of composite materials and structure, Kluwer Academic Publishers, The Netherlands, pp.114 (2) S.R. White and P.T. Mather. 1991. 36th International SAMPE Symposium, pp.1518 (3) G.A. Sofer and E.A. Hauser. 1952. J.Polymer Sci., 8:611620 (4) W.P. Winfree, and F. R.Parker. 1986. Review of progress in QNDE, 5B, pp.10551061 (5) F.R. Parker and W.P. Winfree. 1986. Review of Progress in QNDE, 5B, pp.0631067 (6) T. Saliba, S. Saliba, J. Lanzafame, and L. Gudeman. 1992. 37th International SAMPE Symposium, pp.912 (7) P.J. Biermann, J.H. Cranmer, C.A. Lebowitz, and L.M. Brown. 1996. SPIE, 2948:7283 (8) S.Chiou, P.Kukuchek, D.Echternach, G.Carman, and L.Lai. 1996. SPIE, 2948:338348

(9) D.D. Shepard, K.R. Smith, and D.C. Maurer, Micromet Instruments, Inc., Technical Report No: PR122 (10) C.-K. Jen and J.-G. Legoux. 1996. Proc. IEEE Ultrasonics Symp., pp.771776 (11) Product Data Sheet, No.843-3, Hercules Inc. (12) J.-Y. Chen. 1998. Ph.D. thesis, Concordia University, Montreal, Quebec (13) V. Shah, K. Balasubramaniam, R.-D. Costley, and J.P. Singh. 1997. Review of Progress in QNDE, San Diego, CA (14) F. Cohen-Tenoudji, W.J. Pardee, B.R. Tittmann, L.A. Ahlberg, and R.K. Elsley. 1987. IEEE Transactions on Ultrasonics, Ferroelectrics, and Frequency Control, UFFC-34(2): 263269 page_254 Page 255

Figure 1 Degree of cure vs. cure process time for AS4/3501-6 Gr/Ep prepregs (above)

Figure 2 Illustration of cure monitoring using optic fiber and ultrasonic sensors simultaneously (right)

Figure 3(a) Waveforms obtained from [016] panel cured at 176°C for 2 hrs.

Figure 3(b) Variation of the time delay with cure temperature

Figure 4(a) Time delay vs. process time in reflection/transmission mode

Figure 4(b) Ultrasonic attenuation vs. process time page_255 Page 256

Figure 5 Results of simultaneous fiber-optic and ultrasonic cure monitoring of the panel cured at 176°C for 2hrs.

Figure 6 Reflection coefficient measurements by using a clad buffer rod with one discontinuity ''a", inside the A1 mold

Figure 7(a) Ultrasonic waveforms obtained from the [016] panel cured at 176°C for 3 hrs. by using 5 MHz shear ultrasonic transducer

Figure 7(b) Temperature vs. time delay between Sa and S1

Figure 7(c) Shear reflection coefficient vs. process time (left) page_256 Page 257

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 Chandra Iyer Viswanathan Department of Mechanical and Industrial Engineering University of Manitoba, Winnipeg, MB R3T 5V6, Canada Successful modelling and prediction of creep and creep-rupture of polymer composites depends on how well the influence of various factors, such as fiber volume fraction, moisture, physical again, etc., on the creep and creeprupture of the polymer matrix, is understood and assimilated into the model. In this paper the results of a detailed study, on the influence of various factors, mentioned above, on the visco-elastic properties and fracture of a thermoset polymer matrix, will be presented and discussed. #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_257 Page 258

Impact Fatigue Fracture of Glass Fiber Reinforced Thermoplastics Kazunori Itoh, Takashi Kuriyama, Masaya Kotaki and Ikuo Narisawa Department of Materials Science & Engineering, College of Engineering, Yamagata University, Jonan, Yonezawa city 992-8510, JAPAN Keywords; impact fatigue fracture, fatigue fracture, glass fiber reinforced thermoplastic Abstract The impact fatigue fracture behavior of four kinds of glass fiber or glas bead reinforced thermoplastics has been investigated. The fatigue life of the impact fatigue fracture of the POM and PPE composites was shorter than that of usual the fatigue fracture. On the other hand, for the PPS and SPS composites the fatigue life of impact fatigue fracture was longer than that of the fatigue fracture. Introduction Glass fiber reinforced thermoplastics are recently used as a passage in a hot water supply system. High resistance to water hammer is required for this application, because very high pressure weave occurs in the water passage when the passage is suddenly shut down. In this study, therefore, impact fatigue fracture behavior of four kinds of glass fiber reinforced thermoplastics has been investigated using indentation technique commparing with drop-weight impact behavior.

Experimental Test specimens of 1.6 mm thickness were molded from 30 wt% glass bead reinforced poly(oxymethylene)(POM), 30wt% glass fiber reinforced poly(phenylene ether),(PPE) poly(phenylene sulfide)(PPS) and syndiotactic polystyrene(SPS) sheets by compression molding. Static indentation tests were conducted at a displacement rate of 1 mm/min. The constant energy instrumented drop-weight impact tests were carried out. The weight of the impactor was 6 kg and the height was 1m. Fatigue indentation tests were carried out using pulse weave of 10Hz as shown in Fig. 1(a). Impact fatigue indentation tests were conducted at an indentation speed as shown in Fig. 1(b) in which the pulse weave was repeated with an interval time of 5.9 sec. The fatigue indentation tests were done at e The maximum load for the fatigue and the impact fatigue tests were 50 % of the initial peak load in the static indentation tests. The indentor or impactor of a hemispherical page_258 Page 259 nose of 15 mm in diameter was used in all tests. Acoustic emission (AE) activity was measured during the tests. The gain and the threshold value were s 60 dB and 0.33volt, respectively.

Fig.1 Model of the repeated pulse weave Results In the static indentation tests of POM composites, the load suddenly dropped after the initial maximum load point and then the specimen failed. The AE counts could be detected only just before the specimen fracture. In other specimens, the load dropoed gradually after the initial maximum load point, and the deflection at fracture was larger compared to POM composites. In the instrumented drop-weight impact tests, the load-displacement curves showed quite similar behavior to those of the static indentation tests. The curves were not affected by the indentation rate. The fracture aspects of all specimens were also similar to those of static indentation tests. In the fatigue and impact fatigue tests of the POM composites, the AE counts were also detected just before the specimen breakage. The relation between tan d, AE counts and the number of cycles for PPE composites are shown in Fig. 2 and Fig. 3, respectively, in which N is the number of fatigue cycles and Nf is the number of fatigue cycles when specimens fractured. In the fatigue tests, the AE did not occur in the middle stage of the cyclic loading, but the AE activity continued during the fatigue cycles in the impact fatigue tests. The relation between tan d, AE Counts and the number of fatigue cycles for the SPS composites are shown in Fig. 4 and Fig. 5. The AE was continuously observed during fatigue cycles. However, more active acoustic activity was observed in the fatigue tests than impact fatigue tests. The PPS composites showed the similar AE behavior as the SPS composites. page_259 Page 260

Fig. 5 The relationship between tan d, AE Counts and N/Nf of SPS composites in the impact fatigue tests The number of the repeated cycles (Nf) and the total time (Tf) for specimens fracture are summerized in Table 1. The Tf value of the impact fatigue tests for all specimens was longer than that of fatigue tests. However, there were two characteristics for the fatigue and the impact fatigue tests. For the POM and PPE composites, the Nf value in the impact fatigue tests was shorter than the fatigue tests. In other words, the lifetime decreased by setting the interval time during the cyclic loading. For the PPS and SPS composites, on the other hand, the Nf value of the impact fatigue tests was longer than that of the fatigue tests. Table.1 Relationship of repeated cycles and total time under fatigue and impact fatigue tests Nf(104 cycle) Tf(Min) fatigue tests impact fatigue tests fatigue tests impact fatigue tests POM/GB 2.9 0.8 48 770 PPE/GF 7.5 2.4 124 2350 PPS/GF 3.4 <3.6 57 <3630 SPS/GF 7 <8.1 117 <8100 page_260 Page 261

Fig. 2 The relationship between tan d, AE Counts and N/Nf of PPE composites in the fatigue tests

Fig. 3 The relationship between tan d, AE Counts and N/Nf of PPE composites in the impact fatigue tests

Fig. 4 The relationship between tan d, AE Counts and N/Nf of SPS composites in the fatigue tests page_261 Page 262

Vibration Damping Properties of Adhesive Joints of CFRP Laminates Y. TANIMOTO 1, T. NISHIWAKI 2, A. TANGE 1 and Z. MAEKAWA 1 1Kyoto Institute of Technology, Goshokaido-cho, Matsugasaki, Sakyo-ku, KYOTO 606, Japan 2ASICS Corporation, 6-2-1, Takatsukadai, Nishi-ku, KOBE 651-22, Japan Keywords: Loss factor, Eigenfrequency, Shear deformation, Adhesive joints, Finite element method Abstruct In this study, the vibration properties of adhesive joints of CFRP laminated composites are measured. The specimen used are four kinds of adhesive joints which differ the adherend and adhesion properties. Moreover, in order to investigate the effect of the adherend and adhesion properties, the eigenvibration mode obtained by the eigenvibration analysis are discussed. It is concluded that the 2nd loss factor depends on shear deformability in adhesive layer. 1 Introduction Fiber reinforced Plastic (FRP) can be adapted for many applications in light- weight structures such as spacecraft application, because of its specific stiffness and strength. Recently, the more complicated FRP components are required. In almost structures fabricated by these materials, separate parts or frames are often assembled using adhesively-bonded joints, which furnished bolts and rivets. These joints generally reduce the vibration damping of structures as well as strength properties.1 It is thus very important to improve dynamic properties such as vibration damping property, fatigue resistance and impact property of adhesively-bonded structures. In this study, the vibration properties of adhesive joints of CFRP laminated composites are measured. For this purpose, cantilever beam tests are carried out on 150mm and 16mm wide beam by using B&K T2034 dynamic signal analyzer to measure the eigenfrequency and loss factor. Moreover, in order to investigate the effect of the adherend and adhesion properties, the eigenvibration mode obtained by the new numerical model are investigated. The numerical modeling for the whole single-lap joints are performed by combining the elements constructed independently for adherends and adhesive layers. In the former, the numerical model is proposed and the validity is checked by comparison with experimental results. In the latter, the relationship between the vibration modes simulated the damping factors is discussed. 2 Experiment Method Figure 1 shows a shematic diagram of single-lap joint investigated. The single-lap joint consists of two adherends of composite laminates which are bonded together with adhesive layer.

page_262 Page 263 The overall length is 150mm, the width is 16mm, the lap length is 50mm, thickness of adherends are 1.0mm and thickness of adhesive layer is 0.1mm, as shown in Fig.1. Cantilever beam tests were carried out on 150mm and 16mm wide beam by using B&K T2034 dynamic signal analyzer, as shown in Fig.2. The vibration excitation force was applied by means of a magnetic transducer driving a thin steel disk bonded to the beam tip. The beam response was measured by using a capacitive transducer. Table 1 shows the stacking sequences of composite adherends and two different kinds of adhesion (soft or hard adhesion). Soft adhesion has more lower flexural modulus, as compared with hard adhesion. 3 Modeling The numerical modeling for whole single-lap joints were performed by combining the elements constructed independently for adherends and adhesive layers. The eigenvibration mode is always three-dimensional deformation which includes in-plane and out-plane. This indicates that the simulation requires three-dimensional solid element in the finite element analysis. However, in case that the adherend plane structure is modeled by the conventional solid element, the huge elements and huge solution time are unavoidable. Especially, the adhesive layer thickness is much smaller than other dimensions. It is necessary to shorten solution times and consider adhesive layer. Therefore, we propose new numerical model which constructed independently modeled as adherend with orthtropic-shell element and

Fig.1 Specimen geometry of single lap joints

Fig.2 Schematic illustration of vibration testing system Type A-c A-f B-c B-f

Table 1 Type of single-lap joints Fiber orientation angle of Fiber orientation angle of Adhesive adherend I (clamped part) [deg.] adherend II (free part) [deg.] 0 0 Common 0 0 Flexible 90 0 Common 90 0 Flexible page_263 Page 264

adhesive with beam element respectively.2 Then shell elements are connected with beam elements in the thickness direction to express adhesive layer. The space between shell and beam elements is combined by rigid link to the correspond the beam elements length to the real adhesive thickness. Using the model, eigenfrequencies of single-lap joints of CFRP laminates were analyzed. The boundary condition is clamped in one end and perfectly free for all the other ends. Figure 3 shows the mesh division in the proposed model. The number of shell elements, beam elements and nodes are 128, 50 and 265, respectively. 4 Results and Discussion Table 2 shows comparison between analytical and experimental results in the 1st and 2nd eigenfrequencies in order to check the validity of the proposed model. The good agreements indicate that the eigenvibration analysis of single-lap joints can be successfully performed by the model proposed in this work. The influence of fiber orientation angle on the vibration properties are now discussed. In case of the 1st flexural mode with fixed end, the eigenfrequency of adherend I material with orientation angle of 0 deg.(Type A) exists in the higher frequency side, as compared with 90 deg. orientation angle. This result will be explained from the fact that eigenfrequency depends upon the flexural rigidity of adherend I clamped, because the maximum stress is caused at the clamped area. On the other hand, the difference of eigenfrequency is not significant between adherends of Type B-c and Type B-f. This fact indicates that the sensitivity of flexural rigidity of adherend clamped area to the eigenfrequency is much smaller in the 2nd mode than in the 1st mode. Figure 4 shows the experimental results for 1st and 2nd damping factors of various specimen types. In the 1st mode, no significant difference in loss factor caused by adhesive properties was observed, as can be seen in Fig.4. On the other hand, in case of Type A under the 2nd mode, significant difference in loss factor was observed from the comparison between Type A-c and Type

Fig.3 Finite element division Table 2 Comparison between analytical and experimental results for 1st and 2nd eigenfrequencies 1st flexural frequency (Hz) 2nd flexural frequency (Hz) Type Experiment (Hz) Analysis (Hz) Error(%) Experiment (Hz) Analysis (Hz) Error(%) A-c 52.69 57.62 9.36 454.2 481.7 6.05 A-f 51.00 55.69 9.20 347.2 375.4 8.12 B-c 16.37 17.93 9.53 169.0 186.1 10.12 B-f 16.20 17.80 9.88 159.5 175.2 9.84 Error=(Analysis-Experiment)/Experiment page_264 Page 265

Fig.4 Loss factor of each type A-f. In other words, loss factor of Type A-f were much larger than those of other types. These results are derived from the vibration mode.

As already discussed in the previous section, the damping property mainly mainly depends upon the deformed shape of the vibration mode. In addition, the damping property depends upon shear deformability of structure.

Fig.5 Flexural mode of single lap joints Figure 5 shows 1st and 2nd vibration mode obtained by eigenvibration analysis. The deformed shape of 1st vibration mode is similar to the canti-lever beam one, on the other hand, that of 2nd vibration mode is similar to the 3-point bending one. The damping properties in the 1st vibration mode depends upon the flexural rigidity at the clamped area, while 2nd vibration mode depends upon the flexural rigidity in the hoop position. Judging from Fig.5 (b), it is concluded that the large loss factor in Type A-f is caused by the shear deformability of adhesive layer. 5 Conclusion (1) It was concluded that the proposed model which is constructed with shell and beam elements that correspond to adherend and adhesive, respectively is effective for the eigenvalue problems. (2) It was concluded that the 2nd loss factor depends on shear deformability. References 1. Z. Maekawa, H. Hamada, S. Yoshioka, S. Motogi & A. Fukuda, Proceedings of the thirtyfirst Japan Congress on Materials Research, (1988), pp.165171 2. Y. Tanimoto, A. Tange, Z. Maekawa, & T. Nishiwaki, Proceedings of the fifth Japan International SAMPE Symposium, (1997), pp. 17071710 page_265 Page 266

The Effect of the Chemical Metamerism Polyolefin on the Friction and Wear of Bronze Powder Filled High Density Polyethylene Atsushi Saito, Hiromasa Takahashi and Ikjin Um College of Science and Technology, Nihon University 7-24-1 Narashino-dai Funabashi-shi, Chiba 274-8501, Japan KEYWORDS : HDPE, Bronze powder, Chemical metamerism polyolefin, Friction, wear, load-bearing action, fillerholding capacity, filler shape effect. Abstract As previously reported, in order to improve the tribological characteristics of the polymer sliding members, the polyoxymethylene (POM), high density polyethylene (HDPE), and polytetrafluoroethylene (PTFE) were each blended with bronze powder, and their filler effects clarified. As a result, a certain filler effect was achieved from all of those polymer matrices blended with bronze powder, and their wear rates were actually reduced to less than half of their respective unfilled polymers. On the other hand, however, the geometrical shape effect of the bronze powder with respect to wear rate could not be achieved as expected from blended HDPE and PTFE due to the poor adhesiveness between the matrices and the filler. In this paper, the HDPE filled with bronze powder was blended with the chemical metamerism polyolefin (CMP) an adhesive resin to examine the blending effect of the CMP on

the tribological characteristics. As a result, the wear rate was reduced, and the geometrical shape effect of the filled bronze powder appeared. The wear rate was further reduced as the shape index of bronze powder increased. When the bronze powder filled HDPE was blended with CMP at a blending ratio of 5wt%, its wear rate was lowest, and its velocity resistance was best. The coefficient of friction of the HDPE and that of the bronze powder filled HDPE were both increased to some extent by blending them with the CMP. 1 Introduction In the previous paper, in order to enhance the tribological characteristics of the polymer sliding members, the polyoxymethylene (POM), high density polyethylene (HDPE), and polytetrafluoroethylene (PTFE) were each blended with bronze powder, and their filler effects examined [1]. As a result, a certain filler effect was achieved from all of those polymer matrices blended with bronze powder, and their wear rates were actually reduced to less than half of their respective unfilled polymers. On the other hand, however, in the case of the HDPE and PTFE blended with bronze powder, a geometrical shape effect on the wear rate could not be achieved as expected. The reason for this is that both the HDPE and the PTFE are low in mechanical strength, and hence it is difficult for such matrices to hold the filler stably at a constant position all the time. Another reason is considered that the non-polar property of both the HDPE and the PTFE results in poor adhesiveness between the matrices and the filler. To this end, the surface treatment of filler by a coupling agent, and the blending of the matrix with chemical metamerism polyolefin (CMP), etc. are generally known methods of improving the interface strength or adhesiveness between the polymer and the filler of polymer based composite materials [2]. For instance, there is the study of blending glass-fibre-reinforced polypropylene with acid modified page_266 Page 267 polypropylene by Tokaji et al. [3]. It is also reported that another study for the improvement of the mechanical strength of polymer by a polymer blend is under way, in which polymer is blended with the CMP. Moreover, there are a number of studies of tribological characteristics such as the study of adding ionomer to polyethylene/nylon blends by Yelle et al. [4], and the study of the use of polyoxymethylene/stylene-acrylonitrile grafted low density polyethylene blends by Takamatsu et al. [5], etc. The study of using polyamide 6/maleic anhydride grafted high density and low density polyethylene blends by Horiuchi et al. [6] is also well known. Nevertheless, there are not many studies and reports with respect to the improvement of the adhesiveness. In addition, the blending effect of CMP on the tribological characteristics of the polymer sliding members has not been examined in detail to date. In this study, in order to enhance the adhesiveness between matrix and filler, CMP, effective in the enhancement of the adhesiveness between heterogeneous materials, was added to the bronze powder filled HDPE in an attempt to improve the bronze powder-holding capacity of polymer and the geometrical shape effect of filler with respect to the tribological characteristics of the bronze powder filled HDPE. 2 Experimental Details 2.1 Specimens The matrix used in the experiment is HDPE (J-Rex·HD, SS5003B). The bronze powder with different particle shapes (Br, 90wt% Cu and 10wt% Sn, hardness HV115) was used for the filler, whose shape indices (SI, ratio of the major axis to minor axis) were 1.1 and 1.8. The particle size was taken as 38 to 53 mm and the filler content was uniformly taken as 20wt%. CMP (Admer NR106) which is maleic anhydride grafted low density polyethylene (MAH-g-LDPE) was also used to enhance the adhesiveness between the matrix and the filler. The amounts of blended CMP were 5wt%, 15wt%, 25wt%, 35wt% and 80wt%, respectively. Further, 10 types of specimens including unfilled polymers were prepared as shown in Table 1. Specimens were molded to a tubular type by an extruder, and subsequently they were machined to a size of 25 mm outer diameter, 19 mm inner diameter and 30 mm height by a lathe. The frictional surface of the specimens was finished with a 1200 grade emery paper attached to the slider immediately before the experiment. Table 1 Blend compositions of high density polyethylene (HDPE), chemical metamerism polyolefin (CMP) and bronze powder (Br) as percentages of total weight Material code HDPE CMP Br(wt%) (wt%) (wt%) SI 1.1 SI 1.8 Pure CMP 100 CMP+Br(1.1) 80 20 CMP+Br(1.8) 80 20

HDPE+15wt%CMP 85 15 HDPE+15wt%CMP+Br(1.1) 65 15 20 HDPE+15wt%CMP+Br(1.8) 65 15 HDPE+Br(1.8) 80 HDPE+5wt%CMP+Br(1.8) 75 5 HDPE+25wt%CMP+Br(1.8) 55 25 HDPE+35wt%CMP+Br(1.8) 45 35 Note :SI is shape index, ratio of the major axis to minor axis.

20 20 20 20 20

2.2 Apparatus and Experimental Procedure A thrust washer type testing apparatus, in which the end surface of the tubular specimen is rubbed with a slider (counter body), was employed for the friction and wear tests. The slider, made of carbon steel (S45C, hardness HV150), was machined to a disc type to 40 mm diameter and 14 mm height. The frictional surface of the slider was polished with a 1200 grade emery paper to 0.030.05 mmRa, and finally cleaned by scrubbing with tissue paper soaked in ethylalcohol. The experimental conditions were set at 0.05 and 0.1 MPa for the contact pressure (p) and at a range of 0.24 to 2.45 m/s for the sliding velocity (v). All of the tests were conducted under unlubricated condition within a temperature of 23±1°C page_267 Page 268 and a humidity of 50±2%. The coefficient of friction was determined from the torque detected by a dynamometer and frictional force. The weight of the specimen before and after the experiment was measured with an analytical balance (Mettler AE240) to determine the wear volume. 3 Results 3.1 Wear Rate Figure 1 shows the relationship between the wear rate and the sliding velocity for the unfilled CMP. The wear rate reaches a maximum at a sliding velocity of about 1.0m/s, and a minimum at about 1.4m/s. Moreover, when the sliding velocity exceeds 2.0m/s, the wear rate increases sharply. Figure 2 shows the relationship between the wear rate and the sliding velocity for the bronze powder filled CMP. The wear rate increases with the sliding velocity, and the wear rate increases sharply at a sliding velocity of about 2.3m/s. Likewise, the velocity resistance of the CMP is also enhanced by being blended with bronze powder. Further, the wear rate of the bronze powder filled CMP decreases by less than 1/5 when compared with that of the unfilled CMP. The wear rate varies with the shape of bronze powder, and the wear rate with a shape index of 1.8 is less than that with a shape index of 1.1. The above correlation prove the geometrical shape effect of bronze powder on the bronze powder filled CMP. The relationship between the wear rate and the sliding velocity for the unfilled and the bronze powder filled HDPE plus 15wt% CMP is shown in Figure 3. The wear rate of the unfilled HDPE plus 15wt% CMP becomes a maximum at a sliding velocity of 0.6m/s, and a minimum at about 1.0m/s, and afterward the wear rate increases sharply. In contrast, the wear rate of the bronze powder filled HDPE plus 15wt% CMP increases gradually with an increase in sliding velocity. When the HDPE plus 15wt% CMP is blended with bronze powder having a shape index of 1.1, its

Figure 1: Relationship between the wear rate and the sliding velocity for CMP

Figure 2: Relationship between the wear rate and the sliding velocity for bronze powder filled CMP

Figure 3: Relationship between the wear rate and the sliding velocity for unfilled and bronze powder filled HDPE+15wt%CMP page_268 Page 269 wear rate increases sharply at a sliding velocity of about 2.0m/s. Additionally, the velocity resistance of the HDPE plus 15wt% CMP is blended with bronze powder having a shape index of 1.1 is improved by about two times. The wear rate of the bronze powder filled HDPE plus 15wt% CMP is reduced by about 1/3 compared with the unfilled HDPE plus 15wt% CMP. The geometrical shape effect of the bronze powder emerges when the sliding velocity exceeds 1.2m/s, and the wear rate decreases even more at a shape index of 1.8 than at a shape index of 1.1. Additionally, the difference in the wear rate gets greater with the increase in the sliding velocity. The relationship between the wear rate and the sliding velocity for the HDPE plus Br (1.8) blended with CMP is shown in Figure 4. The wear rate of the CMP plus Br (1.8) increases sharply with the increase in the sliding velocity. Although the wear rate of the HDPE plus Br (1.8) blended with CMP increases with the increase in the sliding velocity, its wear rate decreases as the amount of the blended CMP is decreased, resulting in an improvement of the velocity resistance. Additionally, when the amount of blended CMP is 5wt%, its wear rate reaches a minimum. The wear rate of the HDPE plus Br (1.8) blended with CMP is lower than that of the HDPE plus Br (1.8) not blended with CMP. 3.2 Coefficient of Friction Figure 5 shows the variation of the coefficient of friction with the sliding velocity for the unfilled and the bronze powder filled CMP. The coefficient of friction of the unfilled CMP does not depend on the sliding velocity. Although the coefficient of friction of the bronze powder filled CMP varies both upward and downward, it is about constant regardless of the sliding velocity. The coefficient of friction of the unfilled CMP is about 0.38. Both the coefficient of friction of the bronze powder filled CMP and that of the unfilled CMP are nearly the same at about 0.38, and there is also no difference in the

Figure 4: Relationship between the wear rate and the sliding velocity for CMP-blended HDPE+Br (1.8) and CMP+Br (1.8)

Figure 5: Variation of the coefficient of friction with the sliding velocity for unfilled and bronze powder filled CMP

Figure 6: Variation of the coefficient of friction with the sliding velocity for CMP-blended HDPE+Br (1.8) and CMP+Br (1.8) page_269 Page 270 coefficient of friction between the unfilled CMP and the bronze powder filled CMP in terms of the shape index of the bronze powder. The variation of the coefficient of friction with the sliding velocity for the HDPE plus Br (1.8) blended with CMP is shown in Figure 6. The coefficient of friction of the HDPE plus Br (1.8) is not affected by the sliding velocity. The coefficient of friction of the HDPE plus Br(1.8) blended with CMP at a blending ratio ranging from 5 to 35wt% is little affected by the sliding velocity. The coefficient of friction of the HDPE plus Br (1.8) blended with CMP is always at a constant value of about 0.35 regardless of the blending ratio of CMP. This value of 0.35 lies in the middle between the coefficient of friction of the bronze powder filled HDPE, which is 0.33, and that of the bronze powder filled CMP, which is 0.38. Additionally, the coefficient of friction of the unfilled HDPE is 0.26, and that of the HDPE blended with CMP is 0.36. 4 Discussion

In general, particulate filler in a polymer based composite bears mainly the load applied to the frictional surface, and plays a role in reducing the wear rate [7]. This load bearing effect of the particulate filler is affected by the filler holding capability of the matrix and the geometrical shape effect of the filler. The following describes the blending effect of adhesive resin that improves both the particulate filler holding capacity of the matrix, and the geometrical shape effect of the particulate filler. The load bearing capacity can be represented with the maximum stress sl of the following expression [1]. That is,

where tu is the shear strength in the interface between the polymer and the particulate filler, and sc is the compressive stress of the polymer during the strain for the shearing fracture in the interface between the polymer and the particulate filler. Further, it was assumed that the particulate filler has a geometrical shape simplified into a cylinder with a radius of r and height of l and located perpendicularly to the frictional surface. We theoretically analyzed the maximum stress sl with the condition that the particulate filler is in random directions toward the frictional surface. Now, assuming that the load is applied in the direction of Y-axis perpendicular to the frictional surface, the average of the random direction is the direction inclined at an angle of 45° with respect to the X-axis and Z-axis which are perpendicular to the Y-axis. Then, the resolved stress which is the Y-axial component of the maximum stress sl applied in the direction inclined at an angle of 45° in the XY plane becomes and further the resolved stress which is the Y-axial component of the maximum stress sl applied in the direction inclined at an angle of 45° with respect to the Z-axis becomes (cos45° )2´sl, that is, sl/2. Therefore, the equation derived from the multiplication of equation (1) by 1/2 becomes the non-directional maximum stress of the particulate filler. The factors of equation (1) that affect the load bearing capacity are the mechanical properties ( tu, sc ) of the matrix and the geometrical shape of the particulate filler (l /r). Accordingly, the adhesiveness between the matrix and the filler is very important in making the utmost use of the

Figure 7: Relationship between the wear rate and the sliding velocity for unfilled and bronze powder filled LDPE page_270 Page 271 capacity of the matrix to hold the filler at a constant position, i.e. the so-called filler bearing capacity and the geometrical shape effect of the particulate filler. In this study, in order to improve the adhesiveness between the non-polar HDPE (the matrix) and the bronze powder (the particulate filler), we used adhesive resin CMP (MAH-gLDPE). Since HDPE and CMP are the same in PE, their compatibility is desirable, as Imoto [8] stated that two different materials with similar values of solubility parameter (SP value) are easily soluble in each other. Hereinafter, the blending effect of CMP that affects the wear behavior of the bronze powder filled HDPE is examined. In the wear of bronze powder filled polymer, if the mechanical strength of the matrix is weak, the filling effect of the particulate filler, and its satisfactory geometrical shape effect cannot be achieved. The filling effect of bronze powder with a shape index of 1.8 on matrices HDPE, LDPE, and CMP is shown in terms of wear decrement in Table 2. The wear decrement (dw) of each specimen was derived from the following equation (2) by using Figures 1, 2, and 7 and the research paper by Saito et al. [1].

Table 2 Wear decrement(dw) of bronze powder filled polymer for unfilled polymer Matrix of specimen HDPE LDPE CMP dw(%) 68 18 91 Where Wp is the wear rate of the polymer at a sliding velocity of 0.6m/s, Wb is the wear rate of the bronze powder filled polymer. As is clear from Table 2, if the matrix used is the HDPE, its wear decrement is greater than that of the LDPE. We consider that the reason for this is that since the mechanical strength of the HDPE is greater than that of the LDPE as shown in Table 3, the filling effect of bronze powder on the HDPE becomes greater. However, even if the matrices used have nearly the same mechanical strength, for instance when the CMP is used, its wear decrement becomes extremely great as is clear from Table 2. The reason for this is that the load bearing capacity of the filled bronze powder is enhanced by the CMP which is adhesive resin. Further, as is clear from Figure 2, the geometrical shape effect of the bronze powder distinctly appears. Next, an attempt is made to examine the blending effect of the CMP on the wear of the HDPE. Comparing the wear rate between the HDPE and the HDPE plus 15wt% CMP by using Figure 3 and the research paper by Saito et al. [1], their wear rates at a sliding velocity of less than 0.6m/s are little different. When the sliding velocity exceeds 6m/s, the HDPE plus 15wt% CMP has a greater wear rate than the HDPE, and is poorer in velocity resistance. Such influence of the CMP on the wear characteristic of polymer has been studied by Horiuchi et al. [6]. Specifically, the polyamide 6 (PA6) blended with the CMP, that is, maleic anhydride grafted low density polyethylene (MAH-g-LDPE) at a blending ratio of 10 to 20wt% does not cause its specific wear rate to be different from that of the PA6 alone. However, the PA6 blended with MAH-g-LDPE at a blending ratio of 30wt% does enhance its specific wear rate more than that of the PA6 alone. However, when the HDPE plus 15wt% CMP is blended with bronze powder, its wear rate lowers as Figure 3 Table 3 Published properties of HDPE, LDPE and CMP HDPE LDPE CMP Property J-Rex·HD J-Rex·L D ADMER SS5003B F042 NR106 Density (g/cm3) 0.945 0.921 0.928 Melting point 130 112 120 (°C) Melt flow rate 0.35 0.4 9 (g/10min) Tensile yield 25 9.8 9.8 strength (MPa) Elongation at 800 670 500 rupture (%) Note : data are from ref. 9, 10 and 11. page_271 Page 272 shows, and the adhesiveness between the bronze powder and the matrix is improved by the CMP with the geometrical shape effect of the bronze powder also appearing distinctly. Furthermore, when the bronze powder filled HDPE is blended with the CMP at a blending ratio of 5wt%, the wear rate drops lower than that obtained when the matrix is not blended with the CMP. In the HDPE plus Br (1.8), the bronze powder is buried under the frictional surface or inclined in the frictional direction, whereas in the HDPE plus 5wt% CMP plus Br (1.8) such phenomenon does not occur, thereby proving the improvement of the load bearing capacity. Next, an influence of the CMP on the frictional behavior of the bronze powder filled HDPE is examined. As Figure 5 shows, the coefficient of friction of CMP is not varied by filling it with bronze powder. Further, as is clear from Figure 6, when the HDPE plus Br (1.8) is blended with CMP, the coefficient of friction of the matrix increases to some extent regardless of the amount of blended CMP. The reason for this is that the CMP, whose coefficient of friction is greater than that of the HDPE, is transferred to the surface of the metallic counter body to form a film there more easily than the non-polar HDPE. As is clear from Table 3, it is considered that another reason for the above is that CMP has nearly the same melting point as HDPE while the melt flow rate (MFR) of the CMP is greater than 20 times that of the HDPE. Accordingly, blending the HDPE plus Br (1.8) with CMP causes its coefficient of friction to reach that of the CMP. Throughout this study, it has been shown that the wear rate of the polymer blended with particulate filler is affected considerably by the adhesiveness between the matrix and the particulate filler, as well as by the particulate filler holding capacity of the matrix and the geometrical shape effect of the particulate filler. It has been further illustrated that both the filling effect and geometrical shape effect of the particulate fillers vary very considerably with the

mechanical strength of the matrix. To this end, it is considered that the use of a polymer with great mechanical strength as in the polymer matrix blended with particulate filler, makes it possible to improve the particulate filler holding capacity of the polymer and reduce the wear even more by further blending the matrix with the adhesive resin CMP. Moreover, in this case, where the CMP adheres well to the counterface to form a CMP film, the direct contact between the counterface and the metallic filler which mainly bears the load on the frictional surface can be reduced, thereby preventing an increase in friction and wear. We intend to further examine those items described so far as the subject of a future study. 5 Conclusion The following conclusions can be drawn: 1) When filling CMP with bronze powder, the wear rate is reduced to less than about 1/5, the geometrical shape effect of bronze powder appears distinctly, and the bronze powder with its greater shape index reduces the wear rate significantly. 2) When blending bronze powder filled HDPE with CMP at a blending ratio of 15wt%, the wear rate is reduced to about 1/3, and the bronze powder with its greater shape index reduces the wear rate significantly. It is considered that the reason for this is that blending the bronze powder filled HDPE with the CMP causes the adhesiveness between the matrix and the filler to be enhanced, with a resulting improvement in the load bearing capacity and geometrical shape effect of the bronze powder. 3) In the HDPE filled with bronze powder having a shape index of 1.8 and blended with CMP, the wear rate becomes the lowest and the velocity resistance becomes the best when the blending ratio of CMP is 5wt%, indicating that the filling effect of the bronze powder is enhanced by blending the HDPE with the CMP. 4) The coefficient of friction of CMP is not varied by filling the CMP with bronze powder. Further, both the coefficient of friction of HDPE alone and that of bronze powder filled HDPE increase to some extent. page_272 Page 273 Acknowledgments The authors would like to express their appreciation to Japan Polyolefins Co. Ltd. for preparing the samples of HDPE(J-Rex·HD SS5003B) and LDPE(J-Rex·LD F042), and to Mitsui Chemical Co. Ltd. for preparing the sample of CMP(Admer NR106). References 1. A. Saito and H. Takahashi, J. Sci. Eng. Comp. Mater., 6(2), 95109(1997) 2. F. Ide and I. Sasaki, In: Composite Materials and Interfaces, Soc. Mater. Tech. Res.(ed.), Sougo Gijutsu Shuppan, Japan, 217250(1986). 3. K. Tokaji, T. Ogawa, H. Shiota, S. Takaki, T. Shimizu and S. Yumitori, J. Soc. Mat. Sci., Japan, 46(10), 12041209(1997). 4. H. Yelle, H. Benabdallah and H. Richards, Wear, 149, 341352(1991) 5. S. Takamatsu, T. Kobayashi, T. Komoto, M. Sugiura and K. Ohara, Proceedings ofJAST Tribology Conference, Tokyo, 585588(1992). 6. T. Horiuchi, H. Yamane, T. Matsuo and M. Takahashi, Kobunshi Ronbunshu, 53(7), 423433(1996). 7. A. Saito, J. JSPE, 53(3), 432437(1987) 8. M. Imoto, In: Adhesion Handbook(2nd ed.), Japan Adhesion Society(ed.), Nikkan Kogyo Shinbunsha, Japan, Ch.2, 123(1989). 9. J-Rex·HD HDPE resins, Catalog, Japan Polyolefins Co. Ltd., Functional Polymers Division, Japan (1997) 10. J-Rex·LD LDPE resins, Catalog, Japan Polyolefins Co. Ltd., Functional Polymers Division, Japan (1997) 11. ADMER adhesive polyolefins, Admer Technical Catalog AP003, Mitsui Chemical Co. Ltd., Elastomer Division, Japan(1987) page_273

Page 274

Mechanical Behavior of Polyolefin Composites Using Wastes of Fibrous Material As Matrix and Reinforcement T. Kimura*, Y.Kataoka**, Y.Kondo*** and T.Takahashi**** *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 ***Industrial Technology Center of Fukui Prefecture 61 Kawaiwashizuka-cho, Fukui 910-0102, Japan ****Japan Polyolefins Co. Ltd. 3-2, Yakou 2-chome, Kawasaki-ku, Kawasaki 210-8548, Japan Key Words : Waste of fabrics, Injection molding, Thermoplastic composite Abstruct The thermoplastic composites were molded by using the cut waste of non-woven polypropylene fabric as matrix material. The waste cords of cotton, nylon6 and cut waste of rayon were also used as reinforcements. Simple injection molding method was proposed here. Namely, those wastes were fed into the injection molding machine directly together with the matrix material. It is concluded here from the tensile tests that the wastes of cotton and rayon with good dispersion during the molding process are effective for the reinforcements of the thermoplastic composites. Waste cord of nylon6 is not, however, good for reinforcement because of its poor dispersion. Introduction In recent years, the establishment of recycling technique for industrial wastes was required because of environmental protection. Especially, the textile industry takes a growing interest in a recycling system of waste fibers which result from the process of manufacturing products such as textile fabric, non-woven fabric, fishing net, lacy cloth, &c.. In the previous paper, our attention was focused on the material recyclability of such wastes of thermoplastic synthetic fibers as a matrix material of glass fiber reinforced plastics. Usual fibers such as cotton, rayon, etc. have a higher strength than that of plastic resin plates made of polyester, polypropylene, etc.. Therefore, the waste fibers will become a reinforcement of plastic composite. Under these circumstances, the present paper discusses the applicability of such waste fibers as reinforcement for fiber reinforced thermoplastic composites. Recycling System The simple injection molding method was proposed here for the composite molding and the waste fibers were fed into the molding machine directly with the cut waste of non-woven polypropylene as a matrix material. Moreover, spun maleic anhydride grafted polypropylene (g-PP) was used as a binder resin between matrix and reinforcement. The outline of present page_274 Page 275 recycling system is shown in Fig.1. Figures 2(a) and (b) show the injection molding system and its feeding zone respectively. The rolls for the waste fibers are set above the feeding zone. The tensile properties were measured at an ambient temperature using the universal testing machine in accordance with standard procedure JIS K7113. The molded specimen was also shown in Fig.1.

Fig.1 Outline of recycling system.

Fig.2 Injection molding system. page_275 Page 276 Results and Discussion The results of tensile strength for rayon, cotton and nylon6 fiber reinforced polypropylene composites are plotted in Figs.3(a)-(c), respectively. The figure shows that the strength increases with increasing contents of rayon and cotton by using g-PP. It should be noted, however, that the fraction of reinforcement fibers exhibits only a small effect on tensile strength for nylon6 composites. It is cleared from SEM observation, that the lack of strength is caused by the fairly poor dispersion of nylon6 filaments constructing waste cords. To obtain a good dispersion of the reinforcements, the molded specimens were remolded by using the cut specimen into small pieces as shown in Fig.4. Figures 5(a)-(c) show the tensile strength of the re-molded specimens. Recycling time shown in the figures was indicated the time of re-molded. As shown in these figures, the tensile strength of the re-molded specimen increases with increasing recycling time. It is noted from these results that

Fig.3 Tensile strength. page_276 Page 277 the span maleic anhydride grafted polypropylene is good for the binder material between matrix and reinforcement of polyolefin composite. Conclusion Although minor problems were encountered for the composite with nylon6 waste cord as a reinforcement, the simple injection molding method described in the present paper shows promise as a contribution towards the recycling of various kinds of waste fibers as a reinforcement of thermoplastic composites.

Fig.4 Cut specimen for re-molding.

Fig.5 Tensile strength of re-molded composites. page_277 Page 278

Identifying Delamination in Composite Beams Using Response Surface Methodology Yoshinobu SHIMAMURA 1, Akira TODOROKI 1, Hideo KOBAYASHI 1, Haruo NAKAMURA 1 and Ken-ichi IWASAKI 1 1Tokyo Institute of Technology, O-okayama 2-12-1, Meguro-ku, Tokyo 152-8552, Japan Keywords: Delamination, Smart Composite, Response Surface Methodology, Identification Abstract FRP composite structures have been attempted to many structures of vehicles. Damage causes degradation of the stiffness and strength of the structures. Since the damage is interior and invisible from outside, it is important to inspect the damage by non-destructive testing. Smart structures using piezoelectric materials have been investigated to identify the damage. For the smart structures, piezoelectric patches are adopted as sensors and actuators for modal analysis. Since the damage changes modal frequencies of the structures, we can detect the presence of the damage by measuring the change. In the present paper, we propose an identification method using response surface methodology (RSM) for the smart structures. To calculate the change of modal frequencies, FRP cantilever beam with a delamination is used as an analytical model. There are three stages for the identification. First, modal frequencies for the delaminated cantilever beams are calculated from the analytical model. Second, by using RSM, data sets of modal frequencies for the cantilever beams with various delamination lengths and positions are complemented. Third, response surfaces to identify the delamination lengths and position from modal frequencies are regressed from the calculated data sets. As a result, it is shown that delamination length can be predicted easily and inexpensively by using the identification method. Introduction FRP composite structures have been attempted to many structures of vehicles. Damage, such as delamination, by an accidental impact causes degradation of the stiffness and strength of the structures. Since the damage is interior and invisible from outside, it is important to inspect the damage by non-destructive testing. In general, ultrasonic

inspection or X-ray inspection is used to inspect the damage. The inspection techniques are so expensive that health monitoring techniques have been studied by many researchers. Since the damage changes modal frequencies of the structures, we can detect the presence of the damage by measuring the change. The modal frequencies can be measured easily by modal analysis [1,2]. Smart structures using piezoelectric materials have been investigated to identify the damage [36]. In the smart structures, piezoelectric patches are adopted as sensors and actuators for a modal analysis. It is known that the change of the modal frequencies clearly shows the presence of the damage. On the contrary, it is shown that the length and position of the damage is difficult to identify. To identify the length and position, artificial neural network (ANN) is often adopted [3,4]. ANN is, however, time-consuming for teaching, and can not validate the regression results obtained from the ANN by using a statistical test. In the present paper, we propose an identification method using response surface methodology page_278 Page 279 (RSM) for the smart structures. To calculate the change of modal frequencies, FRP cantilever beam with a delamination is used as an analytical model. As a result, it is shown that the length of delamination can be predicted easily and inexpensively by the identification method. Analytical Model The modal analysis is carried out with piezoelectric patches. The instrument set-up of the smart composite cantilever beam is shown in Fig.1. A PVDF patch is adopted as a sensor, and a PZT patch is adopted as an actuator. Piezoelectric patches are bonded to the composite beam, and sine sweep method is adopted for excitation. Longitudinal strain of the composite beam is measured from PVDF response. PZT input and PVDF response at 27.3 Hz are shown in Fig.2. Table 1 shows modal frequencies of a perfect CFRP beam of stacking sequence [0/902/0]s measured from impulse hammer method and the sine sweep method. There are good agreements with the modal frequencies. Since the damage changes modal frequencies of the composite beam, we can detect the presence of the damage of the composite beam by measuring the change. Table 1 Modal frequencies of a composite beam. Modal order 1st [Hz] 2nd [Hz] 3rd [Hz] 27.3 159.4 410.8 Sine sweep Impulse 27.0 159.0 410.0 In the present paper, we investigate an identification method of delamination in a composite beam. FRP cantilever beams with a delamination are used to calculate the change of modal frequencies. Analytical model is shown in Fig.3. The cantilever beam is divided into four Euler beams to calculate modal frequencies [1]. The stacking sequence is [0/902/0]s, and there is a delamination between 0° ply. The analytical model does not include bending and extensional coupling term. Identification Method To identify the length and position, artificial neural network (ANN) is often adopted to identify the damage. ANN is,

Figure 1 Instrument set-up of smart composite cantilever beam with piezoelectric patches.

Figure 2 The time history of PZT input and PVDF response at 27.3 Hz.

Figure 3 Analytical model of a composite beam. page_279 Page 280 however, time-consuming for teaching and can not validate the regression results obtained from the ANN by using a statistical test. In the present study, we propose an identification method using response surface methodology. Using RSM, we can identify the position and length of the delamination easily and inexpensively. There are three stages for the identification method. Flow chart of the identification is shown in Fig.4. First, data sets of modal frequencies for the cantilever beams with several delamination lengths and positions are calculated from the analytical models. For the calculation, material properties and thickness of the beam with errors of ±5% are used. Second, data sets of modal frequencies for the cantilever beams with various delamination lengths and positions are complemented by using RSM. Taking account of the data sets with errors of material properties and thickness of the beam in the first stage, robust response surfaces can be obtained. We can obtain the data sets inexpensively by using RSM. Third, response surfaces to identify the delamination lengths and position from modal frequencies are regressed from the complemented data sets. For basic problems like this, the second stage is not always necessary. In case of complicated structures, more numerous finite element analyses are required for calculating modal frequencies, and they are expensive. Identification Results By using the identification method, we predicted delamination lengths and positions of the FRP cantilever beams. Figure 5 shows the relationship between actual and predicted delamination length. Figure 6 shows the relationship between actual and predicted delamination position. In each figure, horizontal axis means actual value, and vertical axis means predicted value. A diagonal line of each figure means agreement between actual and predicted value. Figure 5 shows that there are good agreement between actual and predicted delamination lengths. Figure 6 shows, however, that positions of delamination can not be predicted by the identification method. To predict more accurate delamination position, the third stage of the identification method was improved as follows.

Figure 4 Flow chart of proposed identification method.

Figure 5 Prediction of delamination lengths. page_280 Page 281

Figure 6 Prediction of delamination positions.

Figure 7 Prediction of delamination positions by an improved identification method. 1) A response surface to identify the delamination length from modal frequencies is regressed from the complemented data sets. 2) The delamination length is predicted from the response surface. 3) A new response surface for the position is calculated from data sets corresponding to the predicted delamination length 4) The delamination position is predicted from the new response surface. Figure 7 shows the relationship between actual and predicted delamination positions calculated form the improved identification method. Though accuracy of the predicted positions is improved in some degree, the predicted positions are not always close to the actual positions. It is necessary to investigate better identification method to predict the more accurate position. Conclusions An identification method of delamination in a smart composite beam with piezoelectric patches was proposed. The identification method is simple and inexpensive to carry out because of using response surface methodology. It is shown analytically that the identification method is effective to predict the length of delamination. References 1. P.M.Mujumdar and S.Suryanarayan, Journal of Sound and Vibration, 125-3 (1988), p.441. 2. J.J.Tracy and G.C.Pardoen, Journal of Composite Materials, 23 (1989), p.1200. 3. A.S.Islam and K.C.Crag, Smart Materials and Structures, 3 (1994), p.318. 4. A.C.Okafor, K.Chandrashekhara and Y.P.Jiang, Smart Materials and Structures, 5 (1996), p.338. 5. H.Luo and S.Hanagud, AIAA, 97-1218 (1997), p.720. 6. C.H.Keilers,Jr., AIAA, 97-1343 (1997), p.1707. page_281 Page 282

Effects of Flexible Interphase on Mechanical Properties of Unidirectional Carbon Fibre Reinforced Composites K.Kitagawa1, S.Hayasaki2, H.Hamada2 1 Kyoto Municipal Institute of Industrial Research, Chudoji, Shimogyou-ku, Kyoto 600, JAPAN 2Kyoto Institute of Technology, Matsugasaki, Sakyo-ku, Kyoto 606, JAPAN Keywords: Unidirectional CFRP, Flexible interphase, Tensile test, Scatter Abstract

Basic mechanical properties of unidirectional carbon fiber reinforced plastics and their scatter depend on fiber, matrix and interfacial properties. We progressed the interphase concept further with a new concept, the flexible interphase. This concept is that the interphase is not a surface; it is a region which has some amount of volume. Flexible epoxy resin was coated on carbon fiber strand. The effectiveness of the interphase on mechanical properties of unidirectional CFRP were investigated. 1 Introduction Carbon fiber reinforced plastics are still now attractive advanced material in application of high modulus, strength and light weight construction. Actual application fields are aerospace, civil engineering fields and so on. Carbon fiber itself possesses high modulus and strength and their properties should be effectively used, however, performance of composite structure is not determined by only matrix and fiber properties. Interfacial properties between fiber and matrix greatly affect on whole composite performances[12]. Although importance of interface in composite materials are well recognized, evaluation system does not appear. The reason of the difficulties of interfacial research work is that interface is not surface between fiber and matrix, it is a new material generated during fabrication. Recently, instead of ''interface", technical word of "interphase" is using. Design of composite structure is greatly performed by using macroscopic database such as tensile modulus, strength, bending, impact properties and so on. In order to design high reliable structures, reliability analyses have been introduced. According to their analytical method, scatters of materials is very important. Even if the material possesses high strength, design allowable strength should be lower in the case that the scatter is large. Allowable concept has been common for designing composite structure. Composite material was inherently large scatter materials due to gathering of two or three components. Therefore, it leads to general image, that is composite material is unreliable material. Further this image make a hesitation of usage of composite material for structural designer. page_282 Page 283 In order to solve this problem concrete data base is constructed by large number of specimen, because the data base should include the scatter of properties. However, this approach is not essential solution, instead,. We have to suppress the scatter of composites. As mentioned before interphase is a new material which have own material constants. Therefore it should be considered that fiber reinforced composite consists of fiber, matrix and interphase; namely three consistence material. If we understand interfacial properties accurately, we can control the performance of composite. This means that the scatter can be also controlled. According to this concept an importance of interphase is emphasized greatly. More actively, if some functions are provided to interphase this leads to creation of new composite materials, which involve low scatter materials. In this paper an attempt to fabricate the composite materials which have low scatter mechanical properties was performed. Flexible interphase layer was formed on the reinforcing fibers and this layer absorb fracture energy which generate at fiber fracture. Therefore it can be considered that fracture behaviour would be uniformed. Finally the scatter of mechanical properties can be reduced. By using flexible rein or sizing agent tensile strength might be reduced, however, the scatter is also reduced. This would contribute high allowable strength structure in reliable design of composite and image of "reliable" material can be arisen in composite materials. 2 Background Table 1 lists data base for unidirectional carbon fiber reinforced plastics using several kinds of carbon fiber and matrix systems. Not only mean value in tensile strength but also C.V. value are important for designing actual composite structure. Here, we focused two material systems, T300/#3601 and T800/#3631. T300 is intermediate strength carbon fiber, while T800 is high strength carbon fiber. #3601 is brittle epoxy resin and #3631 is toughened epoxy resin. Their ultimate strains are 2.0% and 4.7% respectively. Apparently T800 composites show higher strength than T300 composites moreover the important result is C.V. value of both composites. C.V. of T300 composites is higher than T800 composites. According to observation of fractured specimens there TABLE 1 TENSILE PROPERTIES OF UNIDIRECTIONAL CFRP. Tensile strength Ultimate strain Elastic modulus Reinforcing fibre / Fibre alignment N mean (MPa) C.V.(%) N mean (%) C.V.(%) N mean (GPa) C.V.(%) Matrix resin M50J / Epoxy 40 2085.5 10.4 37 0.66 7.5 35 313.8 7.3 M40J / Epoxy 40 2126.1 11.5 35 0.98 12.0 40 226.9 2.9 MR50 / Epoxy 38 2241.8 13.7 25 1.36 9.8 38 163.1 4.2 IM400 / Epoxy 39 1964.2 12.5 36 1.23 9.2 39 165.3 9.4 T300 / #3601 278 1515.3 10.7 276 1.06 9.5 275 138.0 2.1 T300 / #3631 0° 74 1740.5 6.3 74 1.22 5.2 74 135.5 1.3 T800 / #3601 79 2518.4 3.0 56 1.45 2.6 77 162.2 1.9

T800 / #3631 AS4 / 1908 AS4 / PEEK AT-400 / Nylon 6 AT-400 / PPS

279 85 83 100 15

2572.7 2132.6 2253.1 1511.3 1560.1

4.8 5.6 2.4 6.2 4.7

244 32 65 91 15

1.45 1.39 1.55 1.28 1.31

3.9 4.5 3.2 5.2 2.0

276 34 87 93 15

165.7 138.8 139.2 113.6 114.5

2.8 2.4 1.5 3.3 2.2

page_283 Page 284 are two kinds of fracture aspects in T300 composites. They are straight failure and brush failure, on the other hand in T800 composites show only brush failure. Several kinds of fracture aspects increase scatter of strength because the strength distribution exists in each fracture aspects and the strength in whole composites have combined distribution. Therefore the scatter increases. The differences between straight failure and brush failure were caused by energy released from broken fiber, matrix toughness and interfacial properties. SEM observation on the fractured fibers suggested weak interface in T800 composites. Crack generated at fiber propagates interface region due to weak interface. 3 Fabrication Materials used in this study were PAN-based carbon fiber, T300 (TORAYCA, Toray Co., Ltd), and bisphenol-A type epoxy resin, EPIKOTE828 (Shell Chem. Co., Ltd) as flexible interphase. Fabrication process of prepreg is shown in Fig.1, which is normal wetting method. Carbon fibers which extracted from bobbin are dipped in mixture of epoxy resin and hardner and they are wound onto the stell dram. In the case of flexible interphase fibers firstly are dipped in flexible epoxy resin and after they are dipped again in epoxy resin. Concentration of flexible epoxy was 0.5wt%, 1.0wt%, and 5.0wt%. The treated carbon fibers were impregnated with MEK solution of a matrix epoxy resin. Prepreg fabricated were cut into 30cm´30cm and stacked. Unidirectional composites were fabricated and the number of layers was four. (0/90/90/0) stacking sequence was also used for 1.0wt% flexible resin. In this stacking sequence four kinds of specimen can be fabricated. The first one consists of normal prepreg and the second one consists odf prepreg with flexible interphase. The first one consists of normal prepreg and second one consists of prepreg with flexible interphase. The last two was hybrid interfacial laminates. Flexible prepregs were stacked on the outer layer or normal prepregs were in the outer

Figure 1 Schematics of Prereg Fabricating Machine

Figure 2 Pressure, Tempertature and Time Profile for Press Cure. page_284

Page 285 layer. Laminated plate was fabricated by airbag system and curing cycle was indicated in Fig.2 Tensile specimens were cut out from the fabricated plate and their dimensions are based on ASTM D3039. 4 Results 4-1 Unidirectional Composites Table 2 lists results of tensile tests for unidirectional composites. The important result was that 1.0wt% material showed higher tensile strength than 0wt% material. 0wt% material means normal epoxy composites without flexible interphase. Figure3 shows fracture aspects in each concentration of flexible interphase. Apparently 0wt% showed two different fracture aspects; straight mode and brush mode. This tendency was similar to T300/#3601 composite described at section2. On the other hand composites with flexible interphase showed only brush failure mode. Accordingly, C.V. values of than that of 0wt% material. 4-2 (0/90/90/0) Material Table 3 shows tensile test results of (0/90/90/0). For the tensile strength 828 only material showed the higher vale than 550 material. However, hybrid interphase material showed the almost 8% higher value than the 828 only material, in which flexible 0° layers are set at the outer layer and 90° rigid layer are set at the middle. TABLE 2 TENSILE PROPERTIES OF CFRP LAMINATES. Tensile strength Ultimate strain Elastic modulus Concentration of flexible epoxy resin Fiber N mean C.V. N mean C.V. N mean C.V. (wt%) alignment (MPa) (%) (%) (%) (GPa) (%) 0 29 1680.9 5.8 24 1.12 5.0 28 148.8 3.2 0.5 0° 24 1682.3 4.1 23 1.15 4.9 23 147.0 3.8 1.0 26 1727.0 3.7 26 1.24 5.4 25 146.3 4.5 5.0 25 1563.8 3.9 24 1.10 4.3 25 141.7 3.6

Figure 3 Typical fracture aspects. page_285 Page 286 TABLE 3 TENSILE PROPERTIES OF (0/90/90/0) LAMINATES. Elastic modulus Tensile strength Material Fiber alignment N mean (GPa) C.V. N mean (MPa) C.V. 828 only 5 79.6 3.3 5 842.7 9.2 550 only 0/90/90/0 4 81.9 2.8 5 765.5 7.1 550/828/828/550 5 77.2 3.1 5 912.7 5.1 828/550/550/828 4 75.2 4.0 5 828.9 4.8 5 Discussion

Interphase generated through our fabrication process is shown in Figure 4 as schematic drawings. Carbon fibers supplied from manufacturing company originally were treated, so that "sizing agent" might be unknown. In the normal specimens a matrix resin was applied on the sizing agent layer, and they makes a new interface. In the case of flexible interphase the flexible resin was added on the sizing agent layer, and normal resin, next a matrix resin was applied. Therefore, thick interphase was created, however new interface between flexible resin and matrix resin was not

Fig. 4 Design of flexible interphase.

Fig. 5 Axial tensile fracture mechanisms of each unidirectional CFRP laminate page_286 Page 287 cleared and it would be a kind of gradient layers mixture of rigid and flexible resin, because resins with same main molecular chain were used. Although a gradient interphase was created, the resin near sizing agent layer has flexibility. Fracture mechanism of unidirectional composites is summarized in Figure 5. The first fiber fracture occurs due to different fracture strain. The normal interphase as shown in Figure 4-a or sizing layer can not absorb the fracture energy of fibers Further brittle matrix can not suppress the crack propagation and finally fracture mode would be straight mode which usually appeared in low strength specimens. Sometimes the crack propagation is stopped by the adjacent reinforcing fibers, and it leads to brush mode and high strength.. In the case of flexible interphase the fracture can be absorbed in high toughness interphase. This mechanism leads to interfacial fracture, and finally bruch mode appears, but straight mode is hardly observed these are preliminary explanation. However,

very thick flexible interphase would reduce stress transmission to reinforcing fibers so that tensile strength of 5 wt% material was smaller than others. For flexible interphase an optimum contents existed; !.0 wt% in this study. 6 Conclusion The concept of flexible interphase was proposed. Flexible epoxy resin was applied on the carbon fibers with different concentration firstly, and next normal rigid epoxy resin was impregnated to fabricate unidirectional prepreg. Tensile strength value was increased and C. V. value was decreased in unidirectional composites. The usefulness of flexible interphase was confirmed in (0/90/90/0). References 1) K. Kitagawa, S. Hayasaki N. Ikuta and Z. Maekawa, Comp. Interfaces, 4, 417(1997). 2) N. Ikuta, Y. Suzuki, Z. Maekawa and H. Hamada, in Proc. 4th Japan Int. SAMPE Symp.933, Tokyo(1995). page_287 Page 288

Impact Properties of Braided Composites with Flexible Interphase Asami Nakai1, Hiroyuki Hamada2, Kazuo Kitagawa3 and Nobuo Takeda4 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 Kyoto Municipal Institute of Industrical Research, Chudoji, Shimogyo-ku, Kyoto 606, Japan 4 Center for Collaborative Research (CCR), The University of Tokyo Keywords: Flexible Interphase, Braided Composite, Charpy Impact Test, Absorbed Energy Abstract In this study, flexible interphase was introduced to fabricate braided composite materials which have improved impact properties without sacrificing static properties. Carbon fiber reinforced braided composites with flexible interphase were fabricated, and Charpy impact tests were performed for flat braided composites with different concentrations of flexible interphase. The impact bending strength was improved by sizing treatment with flexible epoxy resin, and showed the maximum value at the optimum concentration of flexible epoxy resin. The effects of the sizing treatment on increase in absorbed energy depended on the properties of matrix resin. The improvement of the impact propagation energy could be explained by the difference in the failure process. 1 Introduction Textile composites reinforced with woven, knitted and braided fabric preforms using well-developed textile technologies, are being considered for potential structural applications in various fields. The composites are highly suitable for automated process, so that this technique has much 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. Woven, knitted and braided fabrics, in spite of differences in fabrication process and mechanical behavior, are fabricated by intertwining the fiber bundles, and the whole structure consists of repeating units of crossing fiber bundles. Therefore, the architecture of a textile fabric in textile composites is complex, and the parameters controlling the behavior at fiber crossing parts become numerous. Moreover, the properties of the interface between fibers and matrix are essential to determine the behavior at fiber crossing parts. In fiber reinforced plastics (FRP), properties of interfaces between fibers and matrix are so important that many studies have been conducted on surface treatments to improve the adhesive properties of fiber/resin matrix interface13). A silane coupling agent is widely used as an adhesive between glass fibers and resin4,5). Recently, extensive studies have suggested that the resin region around the fiber/resin interface was affected by the surface treatment on fibers, and was defined as an interphase6). To

page_288 Page 289 extend the design allowances of FRP as a structure, material design in consideration of interphase become extremely important. More actively, providing new performance or functions to interphase leads to creation of new composite materials which have improved mechanical properties as a whole. In this study, flexible interphase was introduced to increase the energy absorption ability to the interphase and an attempt to fabricate the composite materials which have improved impact properties without sacrificing static properties was performed. Carbon fiber reinforced braided composites with flexible interphase were fabricated and Charpy impact tests were performed for flat braided composites with different concentrations of flexible interphase. 2 Fabrication of Specimens Materials used in this study were PAN-based carbon fiber bundles, HTA (Toho Rayon) and two types of bisphenol-A type epoxy resin, EPOMIK R-140 (Mitui Petro-Chemical Industry Co.) and EPIKOTE 828 (Shell Chem. Co., Ltd.) as matrix resin. As flexible interphase, modified epoxy resin KR-550 (Ko-ei Chem. Co., Ltd.) was used with methyl ethyl ketone (MEK) as a solvent. Figure 1 shows the tensile properties of the two types of epoxy resin and flexible epoxy resin for flexible interphase. Flexible epoxy resin showed quite a ductile property with the ultimate failure strain of 60%. R-140 have a ductile property compared with a common epoxy resin, EPIKOTE 828.

Fig.1 Tensile properties of epoxy resin. Sizing treatment with flexible epoxy resin was performed using a prepreg fabricating machine (Fig.2) as follows: Flat braided fabric with 2/2 intersection repeat was fabricated by using a flat braiding machine with 25 spindles. Fabricated braided fabric was dipped in MEK solution of flexible epoxy resin and a hardner, and they were wounded onto a steel dram and dried on an electric oven at 135°C for 25 minutes. The concentration of flexible epoxy resin was 0.0, 1.0, 3.0, and 5.0wt%. The treated fabric was impregnated with resin, degassed in a vacuum chamber and subsequently a flat bar was manufactured by hand lay-up. The number of layers was three, and the thickness of the flat specimen was approximately 2.5mm. page_289 Page 290

Fig.2 Prepreg fabricating machine.

3 Experimental Method Charpy impact tests were performed with unnotched specimens using an instrumented Chapy impact machine. The distance between the specimen anvils, namely the span length, was fixed at 45mm as shown in Fig.3. The impact load was given in flatwise direction of the specimen. The specimens were 15mm wide, 70mm long and the other dimension was equivalent to the plate thickness. The impact machine is instrumanted so that the dynamic load vs time can be recorded on a digital oscilloscope, and this signal is also integrated to record the energy vs time trace.

Fig.3 Specimen arrangement for Charpy impact test. page_290 Page 291 The load-time history can be divided into two distinct regions - a region of fracture initiation and a region of fracture propagation7). In the initiation region, as the load increases, elastic strain energy is accumulated in the specimen during the contact with the striking head of the pendulum. In the region, no gross failure takes place but failure mechanisms on a microscale are possible. When a critical load is reached at the end of the initiation phase, the composite specimen may fail. At this point, the fracture propagates either in a catastrophic brittle manner or in a progressive manner continuing to absorb energy at smaller loads. The total impact energy is the sum of the initiation energy and propagation energy. The energy can be divided by the cross sectional area of each specimen in order to obtain their normalized values. 4 Results and Discussion 4.1 Dynamic Load-displacement Curves Figures 4 and 5 show dynamic load-displacement curves in Charpy impact tests for composites with different matrix resins. In the case of specimens in which R-140 was used as matrix resin (Fig.4), the impact load increased to the maximum and then decreased gradually without sudden reduction. This indicates an ability for the material to carry certain amount of load even after the initial failure has occurred. The load-displacement history up to the maximum load was independent on the concentration of flexible epoxy resin. After the maximum load it depended on the concentration. The impact time until the ultimate failure increased with increase in the concentration. In the case of specimens in which EPIKOTE 828 was used as matrix resin (Fig.5), it can be seen that the impact load increases to the maximum and then decreases rather rapidly, but another local peak also exists. This also indicates an ability for the material to carry some load even after the initial failure has occurred. The load-displacement history both before and after the maximum load depended on the concentration of flexible epoxy resin. The first maximum load have a peak at the concentration of 1.0 wt%, whereas, the impact time after the maximum load is shortest in the case of 1.0 wt%. Obviously, the load-displacement history with EPIKOTE 828 was different from that with R-140.

Fig.4 Load-displacement curves for Charpy impact (R-140). page_291 Page 292

Fig.5 Load-displacement curves for Charpy impact (EPIKOTE 828). 4.2 Microscopic Fracture Observations Figure 6 shows the SEM photographs of the failed specimens with R-140 with the concentration of 0 wt% and 5 wt%. The pulled-out fiber bundle at the tensile side of the specimen was observed. In the case of 0 wt%, or specimen without sizing treatment with flexible epoxy resin, the laceration by shear fracture was clearly observed. This indicates the fracture progressed in a rather brittle manner. On the other hand, in the case of 5 wt%, or specimen with sizing treatment with flexible epoxy resin, resin was elongated near a fiber for plastic deformation. This means the fracture progressed in a comparatively ductile manner. As a consequence, the fracture progress around or inside fiber bundles of braided composites was changed from brittle to ductile fracture with flexible epoxy resin. 4.3 Impact Strength and Absorbed Energy Tables 1 and 2 indicate that the normalized initiation, propagation, total energies and impact bending strength of the specimen with R-140 and EPIKOTE 828 as matrix resin, respectively. The impact bending strength of the specimen both with R-140 and EPIKOTE 828 was improved by sizing treatment with flexible epoxy resin, and showed the maximum value at 3 wt% in the case of R-140 and at 1 wt% in the case of EPIKOTE 828. For absorbed energies, in the case of R-140, the normalized initiation energies Ei showed almost the same value regardless of the concentration of flexible epoxy resin, but, the normalized propagation energies Ep and the normalized total energies Et increased with increase in the concentration. On the other hand, in the case of EPIKOTE 828, Ei was also independent on the concentration of flexible epoxy resin. However, Ep showed the maximum value at 0 wt%, and decreased by sizing treatment with flexible epoxy resin and then increased with increase in the concentration from 3 wt% to 5 wt%. Thus, the effects of sizing treatment with flexible epoxy resin on increase in absorbed energy depended on the properties of matrix resin. page_292 Page 293

Fig.6 Typical fracture surfaces after Charpy impact. Table 1 Impact absorbed energy (R140). Flexible resin content (wt%)Ei (KJ/m2)Ep(KJ/m2Et (KJ/m2)Strength (MPa) 0 57.1 34.4 91.6 1295 1 56.0 42.6 98.7 1304 3 56.0 42.7 98.7 1349 5 55.6 56.6 112 1327 Table 2 Impact absorbed energy (EPIKOTE828). Flexible resin content Ei Ep(KJ/m2) Et Strength (wt%) (KJ/m2) (KJ/m2) (MPa) 0 60.2 67.6 128 1229 1 61.4 49.4 111 1396 3 61.0 48.2 109 1386 5 54.8 62.0 117 1293 page_293 Page 294 A closer examination of the failed specimens by an optical microscope showed that four different modes of failure were dominant in impact fracture surfaces of braided composites, namely delamination, fiber breakage, cracks inside fiber bundles and cracks around fiber bundles (including both between fiber bundles and between layers). Therefore, the total energy absorption after the maximum load Etotal was expressed by the following equations as the sum of each contributed energy absorption;

where Edel is for delamination, Efiber for fiber breakage, Einside for crack propagation inside fiber bundles, and Earound for crack propagation around fiber bundles. Ek is the kinetic energy stored in the specimen at fracture. From our results, it can be concluded that the surface treatment does not have a significant effect on the amount of bending deflection and fiber breakage, so the energy Ek and Efiber are assumed to be constant regardless of the concentration of flexible epoxy resin. As a result, only Edel, Einside and Earound depend on the concentration of flexible epoxy resin. In the case of R-140, delamination was not observed at the failed cross section, and the crack propagated both inside fiber bundles and around fiber bundles. Figure 7 shows optical microscopic photographs at the cross section of failed specimens with R-140 with the concentration of 0 wt% and 5 wt%.

Fig.7 Typical fracture aspects at cross section after Charpy impact (composites with R-140 as matrix). page_294 Page 295 The total crack length increased with increase in the concentration of flexible epoxy resin for cracks both inside and around fiber bundles. The increase in the total crack length means the increase in Einside and Earound. It was considered that, therefore, Ep increased with increase in the concentration. In the case of EPIKOTE 828, the delamination dominated the fracture at the 0 wt% specimen. The delamination led to the pull-out of fiber bundles and allowed the 0 wt% specimen to absorb more energy Edel. The delamination was diminished by sizing treatment with flexible epoxy resin, whereas cracks inside fiber bundles were dominant in impact fracture of braided composites. The total crack length inside fiber bundles, namely Einside, increased with increase in the concentration. This is the reason why Ep increased with increase in the concentration from 3 wt% to 5 wt%. 5 Conclusions In the present study, the following conclusions were obtained; 1) The impact bending strength of the specimen was improved by sizing treatment with flexible epoxy resin, and showed the maximum value at the optimum concentration of flexible epoxy resin. 2) The fracture progress around or inside fiber bundles of braided composites changed from brittle to ductile fracture by the sizing treatment. 3) The effects of sizing treatment with flexible epoxy resin on increase in absorbed energy depended on the properties of matrix resin. 4) The total propagation energy absorption after the maximum load was expressed as the sum of each contributed energy absorption. Each term depended on the properties of matrix resin and flexible interphase. In consequence, it was concluded that the material design which meets the demands for various usage can be achieved by selecting not only fiber and matrix but also interphase. Reference 1. Y.Nakanishi, K.Hana and H.Hamada, Composites Sci. and Tech., 57, 1139(1997). 2. J.Ivens, M.Wevers and I.Verpoest, Composites, 25, 722(1994). 3. N.H.Ladizesky and I.M.Ward, Journal of Material Science, 18, 533(1983). 4. H.Ichihashi, H.Hamada, N.Ikuta and Z.Maekawa, SEN-I-GAKKAISHI, 49, 169(1993). 5. T.Morii, H.Hamada, N.Ikuta, M.Desaeger and I.Verpoest, Proc.Japan-U.S.CCM-VII, 207(1995).

6. Y.Suzuki, Z.Maekawa, H.Hamada, M.Kibune, M.Hojo and N.Ikuta, J. of Material Science, 27,6782(1992). 7. P.Yeung and L.J.Broutman, Polymer Engineering and Science, 18, 62(1978). page_295 Page 296

Evaluation of Delamination Energy Release Rates by Layerwise Higher-Order Finite 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 Abstract As quite well-known, composite laminates exhibit a variety of failure and damage phenomena which conventional homogeneous, isotropic materials do not experience, such as fiber breakage, matrix cracking, delamination and their combinations. In particular, delamination is one of the most feared failures for structural reliability of composite laminates. Therefore, a precise and reliable numerical tool for evaluating the delamination properties of composite laminates is becoming of more and more importance. Although numerous finite element models have been applied to delamination modelings and the strain energy release rate calculations, almost all of them are limited to some two-dimensional (2-D) or quasi-three-dimensional (quasi-3-D), local regions in the laminates [13]. Consequently, in the context with respect to delamination modelings and the strain energy release rate calculations, multi-layered (layerwise) plate/shell finite elements [46] will serve as one of the most attracting models for this purpose because they are capable of evaluating delamination properties of fully 3-D, arbitrary shaped composite laminates with various kinds of delamination failures. In this study, a delamination modelings and the energy release rate calculations by the layerwise higher-order plate element of the authors [5] is presented. In the present finite element, instead of detaching nodal points like in the conventional schemes, delaminations and crack extensions are modeled by setting the penalty numbers [7,8] at the delaminated layer interface equal to zero or considerably small. The advantages of the present scheme are many, for instance, (a) One can easily model adhesion and delamination by merely changing the magnitude of the penalty numbers and any iteration or any non-linear consideration is not required ; (b) One can easily control delamination modes, which have been difficult in the conventional schemes. For instance, when one wishes to model mode II delamination alone, then one can set only the corresponding penalty number equal to zero and leave the others unchanged ; (c) In the penalty method, there is a possibility for modeling the interfacial properties for the ''interphases" in laminates by setting the penalty numbers some finite magnitude, neither very large nor nearly equal to zero. In addition to this flexible and rational delamination and crack extension modelings by the penalty numbers, the virtual crack closure integral (VCCI) scheme [1] is also sophisticated in the present study. In a customary finite element implementation of VCCI, the nodal relative displacements and the nodal forces are used for the energy release rate calculations. This approximation requires substantially fine meshes in the vicinity of the delamination front because the conventional *JSPS Research Fellow page_296 Page 297 scheme omits the area integrals. On the other hand, in the present formulation, the energy release rates are calculated by the exact integral using the penalty numbers and the gap vectors, and hence fine meshes are not required. For validating the present layerwise higher-order finite element model for the strain energy release rate calculations, a simple edge-delamination problems are investigated in this study. Consider a symmetric laminated rectangular strip with its both ends clamped and extended as illustrated in Fig.1. By virtue of Saint Venant's principle, quasi-3-D strain/stress state can be assumed in the area sufficiently far from the loaded ends. The laminate has straight edge delaminations along the free edges and delamination length (or delamination depth) are denoted by a. Three different types of the stacking sequence are considered, i.e.,

cross-ply [90/0]s, angle-ply [45/45]s and quasi-isotropic [45/-45/0/90]s. Each of the dissimilar material laminae was precisely modeled by a specific layer and thus the total number of layers is 4 for the cross-ply and the angle-ply laminates, and 8 for the quasi-isotropic laminate. Symmetric edge delaminations are introduced by setting the penalty numbers equal to zero at the interfaces between 90 and 0 layers for the cross-ply, 45 and -45 layers for the angle-ply, and -45 and 0 layers for the quasi-isotropic laminate, respectively. For the rest of perfectly adhered interfaces, the penalty numbers are set equal to 5´105 GPa/mm. The strain energy release rate G was calculated by using the virtual crack closure integral (VCCI) scheme (see Fig.2) expressed as ;

where gx', gy' and gz' are the gaps after crack extension, while lx', ly' and lz' are the interlaminar stresses prior to crack extension. Note that the present VCCI scheme does not employ any kind of nodal values of relative displacements and forces, but interfacial gaps after crack extension and reaction forces (i.e. interlaminar stresses) prior to crack extension. The integrals for VCCI are exactly carried out by Gauss numerical integration. Aoki and Kondo [3] examined this problem by using quasi-3-D isoparametric finite element analysis. In accordance with Aoki and Kondo, material data and laminate thickness are as follows ; 2h = 1.12 mm, EL = 130.9 GPa, ET = 8.934 GPa, vLT = 0.3213, GLT = GTT = 4.648 GPa Finite element meshes are produced for each of the three cases, so that the virtual crack length Da is set equal to 0.025 mm when the delamination length a is less than or equal to 0.1 mm and 0.05 mm when a is greater than 0.1 mm. It should be noted that, in order to achieve the quasi-3-D state of the problem by the present fully 3-D layerwise plate element, only one quadrant of the laminates can be considered and only one finite element division in the axial direction are used in the case of the cross-ply laminate, while in the cases of the angle-ply and the quasiisotropic laminate, the entire laminate are considered and fine meshings in the axial direction are required because of the dominance of mode III of fracture. Fig.3 through 5 present the strain energy release rate G against edge-delamination length a for cross-ply, angle-ply and quasi-isotropic laminates, respectively. These results consistently indicate that the results by the present FE agree fairly well to the results by Aoki and Kondo. Consequently, it is ensured that the present finite element can be successfully applied to the strain energy release rate calculations. The present FE model is not a particular model for evaluating local stresses, but a general and versatile model for fully three-dimensional structural analysis. Therefore, there are many practical cases considered to which the developed FE can be applied, such as calculations of stress fields and strain energy release rates for fully 3-D arbitrary shaped composite laminates. page_297 Page 298 References 1. E.F. Rybicki, D.W. Schmueser and J. Fox, An energy release rate approach for stable crack growth in the free-edge delamination problem, J. Compos. Mater. 11 (1977), pp.470487. 2. A.S.D. Wang and F.W. Crossman, Initiation and growth of transverse cracks and edge delamination in composite laminates : Part 1. An energy method, J. Compos. Mater. Suppl. 14 (1980), pp.7187. 3. T. Aoki and K. Kondo, Free-edge delamination of anisotropic composite laminates (I) Theoretical approach, J. of the Japan Society for Aeronautical Space Sciences 37(420) (1989), pp.2938. 4. J.N. Reddy, An evaluation of equivalent-single-layer and layerwise theories of composite laminates, Compos. Struct. 25 (1993), pp.2135. 5. K. Suzuki, Layerwise Higher-Order Finite Elements for Laminated Composite Material Structures, Dissertation of Doctor of Engineering at the University of Tokyo (1997). 6. I. Kimpara, K. Kageyama and K. Suzuki, Hierarchical layerwise higher-order finite elements for laminated composite, Proc. Second Canada-Japan Workshop, Montreal, (1998). 7. F. Fraternali and J.N. Reddy, A penalty model for the analysis of laminated composite shells, Int. J. Solid Struct. 30 (1993), pp.33373355.

8. J.P. Fuehne and J.J. Engblom, Finite element/penalty function method for computing stress near debonds, AIAA J. 30 (1992), pp.16251631.

Figure 1 Concept of virtual crack closure integral (VCCI)

Figure 2 Problem definition of a infinitely long laminated strip with symmetric edge delamination page_298 Page 299

Figure 3 Strain energy release rate G versus delamination length a for cross-ply laminate [90/0]s

Figure 4 Strain energy release rate G versus delamination length a for angle-ply laminate [45/-45]s

Figure 5 Strain energy release rate G versus delamination length a for quasi-isotropic laminate [45/-45/0/90]s page_299 Page 300

CAE for SMC Molding H.Hamada1, T.Hasegawa1, E.Tanigaki1 and H.Naito2 1Kyoto Institute of Technology, Matsugasaki, Sakyo-ku, Kyoto 606-8585, JAPAN 2Sekisui Chemical Co., Ltd., Minami-ku, Kyoto 601, JAPAN Keywords: CAE / SMC / slippage flow Abstract Deformation and flow behavior of SMC during compression molding is very important to design of SMC products as well as curing behavior. The progression of cure also greatly affects on the slippage phenomena. This paper describes numerical analysis method for flow behavior of SMC in order to establish total CAE system for SMC compression molding. Finite element analysis was performed to understand fundamental aspects of slippage flow. Three spring elements were employed for representing the initiation of slippage. Validity of this concept was confirmed by both experimental and numerical results. This proposed analysis was included into total CAE system for SMC compression molding. Introduction SMC (Sheet Molding Compound) is still attractive for high volume fabrication composites. The fabrication system is rather simple, namely easy to fabricate complex shape products. Furthermore SMC products have the feature such as good surface without any particular technique. Therefore it is used for relatively large composite structures such as exterior panel, water tank, bathtub and so on. However, heterogeneity of SMC due to short fiber and large amount of fillers causes difficulty of structural design because scatter of basic mechanical properties is large. Not only material heterogeneity itself but also complex flow pattern would be source of unexpected properties, which are sometimes undesirable. In another difficulty of SMC compression molding, a set of mold would be expensive and change of mold design takes from a large loss of time. Therefore CAE system is required for mold and products design. SMC material is usually set in a die as laminated materials. Several SMC laminates are stacked together. Accordingly material flow is not uniform such as metal forming. We have already found slippage flow both interlamina and intralamina, which is typical flow behavior of SMC at initial stage in compression molding. The initial flow behavior would have great influences upon following flow and filling process. In this paper we proposed total CAE system as shown flow chart in the Fig.1. In CAE method proposed by our previous paper, the initial flow analysis was set in important part as well as head conduction and curing process. page_300 Page 301

In this paper, the original analysis method of initial flow behavior, which characterized by slippage, was described. Particularly double nodes were set between both interlamina and intralamina, and separation of the double nodes expressed slippage flow.

Fig.1 Total CAE system in SMC compression molding

Fig.2 Deformation state of SMC.

Fig.2 Finite element mesh. Experimental Results At the initial stage of compression molding the most important deformation state is slippage flow between interlamina and/or intralamina. Figure 2 shows deformation state of 3 layers SMC. The materials were set in the heated lower die at 140°C and the upper die downward to touch the material. It took about 30 seconds. After touching the upper die to the material, compression process immediately started. "0 sec" indicated in this figure means this process, whereas "150 sec" was time between touching material and starting compression process. In the case of 0 sec, top and bottom layers near both dies were soft due to decreases of viscosity by thermal diffusion. Therefore both layers flow exceed the middle layer. On the other hand, in the cases of 150 and 250 sec curing was progressed in top and bottom layers, so that middle layer flows exceed. Apparently in these two cases interlamina slippage flow could be observed. And in the case of 0 sec intralamina slippage flow might be occurred. 150 sec and 250 sec cases would be unreal condition, however, an importance of slippage flow due to curing progression could be recognized by these cases. page_301

Page 302 Analytical Method Three dimensional solid elements were used in large deformation analysis. Materials properties were assumed as elastic-plastic behavior. Figure 3 shows finite element division used in this study. Boundary conditions are follows; Plane-ABCD

:x-,y-displacement are fixed z-displacement are applied -0.1mm by the step

plane-ABRQ, -DCS

:x-displacement are fixed

plane-ADQ

:y-displacement are fixed

plane-QRS

:x-,y-,z-displacement are fixed

In order to express slippage flow between layers double nodes were set on plane -EFG, -HIJ, -KLM, and -NOR. Elastic modulus of top and bottom layers had higher value than that of middle layer, because extreme condition, which was similar to 150 or 250 sec case in above section was selected. According to experimental results as shown in above section, slippage flow layers were significant in these cases. Three Spring Elements Method The approach method to express the slippage flow was that three spring elements were set between double nodes in x, y and z direction respectively. For slippage flow in xy plane, spring constants in x and y direction were smaller that in z direction. If spring constant in z direction would be small, upper material tended to invade into the lower material. These spring elements are imaginary elements, however, from physical meaning of initiation of slippage flow is regarded as a kind of frictional problem. Therefore this method was an application of numerical modeling in frictional problem. In this paper spring constants in x and y direction were 0.1, whereas that in z direction was 100. Schematic drawings of numerical method around double nodes are shown in Fig.4. Results and Discussion Figure5 shows the deformed mesh of proposed analytical model in 25 incremental step.

Fig.4 Schematic drawing of finite element modeling. page_302

Page 303 Clearly the middle layer flow exceeds to top and bottom layers. Furthermore not only middle layer but also both top and bottom layers deformed. The outer layer was not deformed due to high elastic modulus. The gaps could be obtained between interlamina and intralamina, so that there would be regarded as slippage flows. From these results, our proposed model was appropriate model for initial deformation state in SMC molding. Figure6 shows total of reaction force in the top surface, the slope coefficient and the length of spring element models at p1 and p2 shown in Fig.3. The reaction force decreased at second incremental steps and after this point, the reaction forces gradually increased. The length of spring element of p2, which means intralamina, increased at second incremental steps. On the other hand that of p1, which means interlamina, increased at incremental step of third. From these results it is clear that reaction force in the top layer decreases when length of spring element increases. In other words slippage behaviour cause the decrease of reaction force in the top layer. The force applied into z direction was transferred into x and y directions by means of the spring element, because of slippage flow. Therefore it is considered that the reaction force in z direction was decreased. This decrease of reaction force in z direction can be regard as the decrease of molding pressure at actual molding process. In SMC molding the decrease of molding pressure causes bad surface smoothness.

Fig.5 Deformed mesh of hybrid laminate model

Fig.6 Reaction force, the slope coefficient and spring length of p1 andp2. Conclusion In this paper two finite element modelings for initial deformation state were proposed. The aim of modeling was to express the slippage flow. It is concluded that our proposed analytical method will be included into CAE method for SMC compression molding. page_303

Page 304

Energy Absorption Capability of Braided Composites H.Hamada1, A.Nakai2, K.Kameo1 and N.Takeda2 1 Kyoto Institute of Technology, Matsugasaki, Sakyo-ku, Kyoto 606 Japan 2 The University of Tokyo, 4-6-1 Komaba, Meguro-ku, Tokyo 153 Japan Key Words: GF/Epoxy Braided Rod, Braided Pultrusion Process, Resin Curtain, Energy Absorption Capability, H/A Ratio Abstract Braiding has a potential corresponded to a continuous manufacturing process such as Braided Pultrusion Process (BPP). For crushing elements in cars, braided composites show high energy absorption characteristics. During the braiding process unidirectional fibers can be easily inserted to the braid. Under compressive load, unidirectional fibers aligned in the loading direction can sustain high load and braiding bundles prevent crack propagation, which is generated along the unidirectional fibers, by the high circumferential strength. These two facts create progressive crushing behavior in which the high load is kept constant approximately. In this paper, the curshing performance of braided rod was presented. 1 Introduction Safety is one of the emergency issues among various requirements in cars, because to protect passengers from accidents or collision is a daily demand for car-users. The car body, which has high performance in protecting the passengers, have been arisen and proposed in many ways. Polymeric composite tubes have been recognized that they have high energy absorption performance. Table 1 lists specific energy absorption value (Es) in many kinds of materials including metallic materials [13]. Particularly, carbon/PEEK composite tube has very high energy absorption capability, which exceeds 200kJ/kg. This is good example for adding unique performance and using in small parts. Regarding to high volume and low cost fabrication for polymeric composites are pultrusion process in which reinforcing fibers are aligned in longitudinal direction is one of possible process, so that pultruded composites possess high strength. However, transverse strength is very low due to unidirectional composites. In order to solve this problem chopped mat sheets or fiber bundles are wound around the unidirectional fibers. However, this technique can not provide sufficient transverse strength. Our group has developed Braided Pultrusion Process (BPP) to provide high strength in both directions, the longitudinal and the transverse directions [4,5]. In this composite, unidirectional fibers are surrounded by braided fabric. And BPP can produce high strength composites at low-cost. In this paper, we introduce fabrication method of braided pultruded composites. It will contribute to fabricate crushing elements by low-cost fabrication process. The crushing performance was measured and the cross-section of the composite was observed microscopically to understand the crushing mechanism. page_304 Page 305 Table 1 Energy absorption capabilities. Reinforcement Matrix Es value (kJ/kg) Glass Fiber Epoxy 53.7 Kevlar Fiber Epoxy 57.9 Carbon Fiber Epoxy 82.1 PEEK 225.0 PEI 155.4 PI 131.4 PAS 128.1 Steel 33.7 Aluminum 66.9 2 Braided Rod by BPP

2-1 Composition of Braided Rod Figure 1 schematically illustrates a braided rod. Unidirectional fibers (axial fibers) are surrounded by tubular braided fabric. The braiding fibers follow the + s and - s directions. The angle is called braiding angle. The longitudinal fibers (middle-end-fibers) also can be simultaneously inserted in the braided preform during the braiding process. It is possible to change both of braiding angle, and the number of the middle-end fibers and the axial fibers.

Figure 1 Scheme of braided rod. 2-2 Braided Pultrusion Process A new pultrusion system has been developed. This system is the result of applying "Integrated Braiding Technique" to the conventional pultrusion system [6]. With this system, high-performance composites can be fabricated at low-cost. In addition, it is possible to not only manufacture pultrusion products of various cross-sectional shapes, but also obtain the product with various mechanical characteristics by changing the braiding angle, number of the middle-end-fibers and the axial fibers. The construction of this fabrication system is shown in Figure 2. From the left, in order, that consists of braider, resin impregnator and heated mold station, puller and cutter.

Figure 2 Braided pultrusion process (BPP) page_305 Page 306 3 Results and Discussion Figure 3 shows the fracture aspects of the specimens after compressive test and the load-displacement curves. Unidirectional rod (UD) caused catastrophical crack propagation along the fibers through the specimen. On the other hand, the fractured material of the braided rods were splayed out in a radial pattern beautifully, and obvious fiber fracture in axial fibers was not observed except a debris wedge and all braided rods showed progressive crushing. Namely, the energy absorption of a braided rod became high.

Figure 3 Crushed aspects and load-displacement curves of UD rod and braided rod. The cross-section of the crushed braided rod was shown in Figure 4. A longitudinal crack at the center of the material (central crack) is approximately 12 mm long. In this material axial fibers aligned to load direction can sustain high load to progressive crushing and tubular braid plays an important role in a restrain effect for the axial fibers to prevent a rapid propagation of the central crack. Material properties of the pultruded composite rods are shown in Table 2. Pultruded braided rod exhibited high bending strength and modulus because of its high fiber volume fraction. Additionally, pultruded braided rod has good properties for crushing test.

Figure 4 Cross-section of the crushed braided rod. page_306 Page 307 Specimen Braided Rod UD Rod

Table 2 Material properties of UD rod and braided rod by BPP. Fiber Composition Diameter Density Vf 3pt Bending 3pt Bending Es (mm) (tex. ´ bundles) (g/cm3 (%) Strength (MPa) Modulus (GPa) (kJ/kg) Core 2310 ´ 55 11.9 2.20 74 1006 47.73 49.5 Braid 1150 ´ 48 Skin 2310 ´ 12 Core 2310 ´ 55 11.8 2.05 62 1003 41.81 --Skin 2310 ´ 20

4 Conclusions Braided pultrusion process, which satisfy both cost and strength requirements, was proposed and developed. With this system (BPP), high-performance braided composite rod could be produced continuously and automatically. Since braided pultruded composites possess inherently high transverse strength, high-impregnated composites lead to excellent performance. In addition, not only manufacturing pultrusion products of various cross-sectional shapes, but also creating various mechanical characteristics to the product can be realized by changing the braiding angle or the number of axial fibers and braiding fibers. And an application field of braided composites would be extended. 5 References

[1] H. Hamada, S. Ramakrishna, T. Nishiwaki and Z. Maekawa, ''Energy absorption characteristics of composite tubes with different cross-sectional shapes", Proceedings of 10th ASM/ESD Conference, (1994) pp.523534. [2] H. Hamada, S. Ramakrishna and H. Sato, "Effect of fiber orientation on the energy absorption capability of carbon fiber/PEEK composite tubes", Journal of Composite Materials, 30-8 (1996) pp.947963. [3] H. Hamada, J. C. Coppola, D. Hull, Z. Maekawa and H. Sato, "Comparison of energy absorption of carbon/epoxy and carbon/PEEK composite tubes", Composites, 23-4 (1992) pp.245252. [4] Y. Hisa, T. Uozumi, A. Fujita, H. Hamada, A. Nakai and A. Yokoyama, "Braiding pultrusion process (BPP)", 27th International SAMPE Technical Conference, (1995) pp.371379. [5] T. Uozumi, Y. Hisa, A. Nakai, H. Hamada, K. Kameo and N. Takeda, "Development of Braided Pultrusion Process", 54th Annual Conference, Composites Institute, The Society of the Plastics Industry, Inc. (1998) Session 1-B. [6] H. Hamada, A. Nakai, A. Fujita, Z. Maekawa, A. Yokoyama and T. Uozumi, "CAE in integrated braided composite", Journal of Science and Engineering of Composite Materials, 4-2 (1995) pp.109120. page_307 Page 308

Stochastic Characteristics of Interlaminar Shear Strengths of Laminated Composites Chao Zhang, R. Ganesan and Suong V. Hoa Concordia Centre for Composites Department of Mechanical Engineering Concordia University Montreal, Quebec, Canada H3G 1M8 Abstract Experimental observations clearly show that delamination failure of laminated composites exhibits two distinct characteristics; (1) the corresponding failure strengths display considerable variability, and (2) the morphologies of the fracture surfaces and thus the interlaminar strengths are strongly dependent on the magnitude of the fiber orientation difference (abbreviated as f.o.d.) between the two adjacent layers of the interface. These characteristics need to be thoroughly investigated in order to have accurate information on the delamination strength. The stochastic analysis technique, in which the interlaminar strength is represented as a stochastic process with respect to the f.o.d. angle, can be an efficient method for this investigation. In the present work, a stochastic model for the interlaminar shear strengths is developed based on a number of discrete experimental sample data sets. The associated experimental scheme is also described, which is performed to measure the randomness in the interlaminar shear strengths using the double notch shear specimens with multidirectional lay-up. 1 Introduction Delamination is one of the fundamental failure modes in multidirectional fiber-reinforced laminated composites. A large amount of experimental research work has shown that the failure strength governed by delamination failure mode exhibits considerable variability from specimen to specimen. However, due to the lack of efficient experimental and modeling techniques, the understanding of the stochastic behavior of delamination failure is still lacking. As the in-plane transverse tensile and shear strength components of a single layer in laminates strongly depend on its thickness and the fiber orientations of its surrounding layers [1,2], the interlaminar strength components of a interface in laminates are also expected to be dependent on the f.o.d. between the two adjacent layers. Some investigations have been performed [3,4] on the critical strain energy release rate using angle-ply specimens to observe the effects of f.o.d. on the corresponding values. The experimental results show a strong dependence of the critical values on the f.o.d. angles. Despite the valuable work on delamination energy release rates, no research work has been performed on the aspect of interlaminar strength components. With the above considerations, a stochastic modeling of the delamination failure is proposed, in which the interlaminar strength components are represented to be a stochastic process with respect to the f.o.d. angle, i.e., s(b). An elaborate experimental scheme is designed for the CYTECâ G40-800/5276-1 carbon/epoxy material system so as to yield the individual and joint probability distributions of the interlaminar shear strength components at 10 discrete values of f.o.d. angle b. The angle values are set to be 0°, 10°, 20°, 30°, 40°, 50°, 60°, 70°, 80°, 90°, and the sample

size is specified to be 36 at each angle. Based upon the experimental sample data sets at 10 discrete f.o.d. angles, the stochastic process for interlaminar shear strength components is established, thus the individual and joint probability distribution functions at all values of f.o.d. angles can be page_308 Page 309 determined. The autocorrelation functions are then obtained and the important stochastic characteristics of interlaminar shear strengths are demonstrated. 2 Stochastic Modeling of Interlaminar Shear Strengths The objective of the present stochastic modeling is to obtain the probability distribution functions of the interlaminar shear strengths at any given f.o.d. angles through the stochastic formulation. In most of the probabilistic analyses, the probability distribution function of a random variable is assumed to be one of the standard distribution forms, such as the Gaussian and Weibull distributions, and then the parameters involved in these standard distribution functions are obtained by the curve-fitting techniques. In the present model, the true individual probability density functions of the interlaminar shear strengths are determined without such an assumption using the maximum entropy technique [5]. The individual probability density function is represented in a general form as

where li's, i = 0, 1, . . ., p, are the parameters to be determined from the given sample data sets. One of the salient features of the maximum entropy technique is that the form of the probability density functions are given a priori, and the order p of the polynomial in the power is determined by the stochastic characteristics of the random variables. Hence, the parameters li's can be evaluated by the moments of the random variables up to the p-th order, which can be estimated from the sample data using the following equation:

In the above, N is the sample size and i denotes the order of a moment. The estimation accuracy of the probability density functions can be improved by using the higher orders of moments. For most kinds of probability functions, however, the moments higher than the 3rd or 4th orders do not make much contribution to the estimation accuracy, that is, the parameters li's with i greater than 3 or 4 are very close to zero [6]. Two algorithms are proposed to constitute the stochastic process of the interlaminar shear strengths, namely the parametric algorithm and the ARMA (autoregressive moving average) algorithm, respectively. 1) Parametric Algorithm Suppose that the nonstationary stochastic process s(b) can be approximately represented by a 10th order of Fourier cosine series in the interval [0, p/2] (here the order is determined by the number of discrete f.o.d. angles at which experiments are performed),

where Ai's are correlated random coefficients [7]. The analytical expressions for various orders of joint moments of the stochastic process s(b) can be derived in terms of the joint moments of the random variables Ai's based on Eq. (3). Therefore, if the joint moments of Ai's can be determined, then those of s(b) are also known. The joint moments between the interlaminar strengths at 10 discrete f.o.d. angles at which experiments are performed can be estimated by

page_309 Page 310 which are then substituted into the analytical expressions of joint moments of s(b). A set of linear homogeneous equations are established regarding the joint moments of Ai's, which are thus solved using a very simple algorithm. At this point, the joint moments of s(b) in the interval [0, p/2] are determined, and then the individual probability density functions at any given f.o.d. angles are obtained through the maximum entropy technique.

2) ARMA Algorithm This algorithm employs the time series analysis methodology, i.e., a sequence of random data is modeled as one realization of the response of a stochastic system to uncorrelated or independent discrete white noise inputs [8]. Therefore, the variations of the experimental data from sample to sample at each discrete f.o.d. angle is viewed as a discrete time series, for which the associated stochastic process is strictly stationary. This type of time series can be simulated by the ARMA(n,m) model, that is,

where m is the mean value. All the ARMA coefficients, fi's and qi's, and the parameters representing the response of the system to discrete white noise inputs can be determined using an iterative procedure of identification, estimation, and diagnostic checking. Once again, the orders of n and m of the discrete time series are determined by the stochastic characteristics of the random variables. In practice, it is usually true that adequate representation of the actual stationary time series is simulated by an ARMA model in which n and m are not greater than 2. All the ARMA coefficients and white noise parameters in the interval [0, p/2] can then be obtained by the interpolation technique. A discrete time series can be written as a linear combination of a number of Green's functions (which are only functions of ARMA coefficients) and independent discrete white noise inputs [8]. Therefore, the sample data of interlaminar shear strengths at any given f.o.d. angle can be obtained by substituting the discrete white noise inputs that are numerically generated by the Monte Carlo simulation technique [9] into the above mentioned linear combination. Finally, the individual probability density functions of the shear strengths at any given f.o.d. angles are evaluated by the maximum entropy technique. 3 Experimental Scheme Experiments are conducted to measure the stochastic behavior of interlaminar shear strengths at interfaces corresponding to the 10 f.o.d. angles that were mentioned before. 1) Specimen Design and Test Setup The double notch shear test, recommended in ASTM D 3846, is chosen in the present work, because a pure shear failure always takes place near the central planes between two notches of unidirectional specimens. The interlaminar shear strengths have to be measured from multidirectional lay-up specimens, thus the specimens and test setup are modified. Because of the two asymmetrically located notches, an undesirable bending moment appears in the specimens. The bending moment can cause a severe stress concentration near the notches, thereby decreasing the measured strengths. Moreover, since the matrix tensile strength is lower than the interlaminar shear strength, a tensile load may cause transverse matrix cracks prior to interlaminar failure. Hence, a compressive load is applied to the double notch shear specimen that is clipped between a clamp by evenly tightened bolts, as shown in Fig. 1. A series of [(0/b)3(b/0)3(b/0)3(0/b)3] (b =0°, 10°, 20°, 30°, 40°, 50°, 60°, 70°, 80°, and 90°) laminates are designed that constitute a number of 0/b interfaces in the middle thickness of the specimen, thereby ensuring that the shear failure takes place on a 0/b interface even though the notches are not cut precisely to half the specimen thickness. The results from the Classical Laminated Plate Theory show that neither residual page_310 Page 311 bending and twisting deformations nor coupling effects occur in this type of laminates although they are asymmetric with respect to the central planes. The geometry and dimensions for the double notch shear specimens are also given in Fig. 1. The ratio of the notch space to the laminate thickness, l/t, is critical to the distribution of the interlaminar shear stress at the central planes; the smaller the value is, the more uniform the shear stress distribution will be. The nominal thickness of the manufactured panels is 4.5 mm, and the notch space is selected to be 7 mm. Furthermore, the notch bottoms are chosen to be of round shape with a fillet radius of 0.9 mm for the convenience of notch-cutting using a diamond blade. A 2D stress analysis is conducted on the unidirectional specimens using the ANSYS software [10], and the notches with and without fillets are considered in order to observe the influence on the stress distribution at the central planes. The stress analysis results are shown in Fig. 2, where the shear stresses are demonstrated with respect to the specimen middle point (i.e., at the point of half length L/2). It can be seen that the fillets do affect the stress distribution; the shear stress is maximum and immediately drops down at the tips of the notches without fillets, while it reaches a maximum value a little away from the notches and then slowly drops down inside the fillet regions. Thus the nominal uniform shear stress at the central planes is modified as

where s0 is the applied compressive load and d is a correction parameter; for non-fillet specimen, d = 0, and for fillet specimens, the best selection is found to be d = 0.5 mm for the geometry specified in the present work. This means that the materials inside the fillet regions are capable of carrying a part of loadings. The interlaminar shear strengths

are thus calculated using Eq. (6). 2) Specimen Preparation and Testing Procedures All the laminates are fabricated from the CYTECâ G40-800/5276-1 carbon/epoxy pre-impregnated tapes with a nominal thickness of 0.185 mm and cured in an autoclave following a specified curing cycle. The C-Scan non-destructive evaluation is performed on the manufactured composite panels and the half-thickness notches are cut using a diamond blade with a higher grit order. Each cured panel has a geometry of 7 in. ´ 14 in.; from each plate 36 specimens are cut using a diamond blade. The edges of the specimens are ground using two types of sand papers with grit orders of 120 and 320 respectively to remove the damage due to cutting. Each specimen for a separate data set is numbered, which is essential for the calculation of joint probability and autocorrelation functions. The specimens are then clipped into the clamp and mounted on an MTS testing machine. A compressive load is applied to the specimens by moving the steel platen upward at a constant speed of 1 mm/minute and terminated when the load suddenly drops down close to zero. During the loading process, relative displacements and the corresponding loads are directly recorded from the MTS machine. 3) Experimental Results A summary of the stochastic characteristics of the interlaminar shear strengths is given in the next section. As expected, all specimens fail in a shear mode roughly along their central planes. The load-displacement curves recorded from the MTS machine for all specimens show a linear relationship until the final failure. The strength values display considerable variations and the coefficients of variation (as well as mean values) are listed in Table 1. The failure loads and the morphologies of fracture surfaces are found to be strongly dependent on the f.o.d. angles; rough surfaces correspond to higher strengths and smooth surfaces to lower strengths. For example, all the specimens with 0/90 interfaces have fracture surfaces with either fiber bridging or crack jumping through individual plies, while all the specimens with 0/80 interfaces have page_311 Page 312 smooth fracture surfaces. The phenomenon of crack jumping has been observed in a few of the specimens with 0/0 and 0/10 interfaces, but most of them have relatively smooth fracture surfaces. Since two different fracture morphologies simultaneously happen to these two types of specimens, their strengths present larger variations, as shown in Table 1. Table 1. Probabilistic moments of the experimental data of the interlaminar shear strengths. 0° 10° 80° 90° Mean Value, MPa Coefficient 50.23 53.87 40.36 67.01 of Variation, % 22.36 19.40 16.72 9.09 4 Stochastic Characteristics Based on the experimental sample data sets, some stochastic characteristics of the interlaminar shear strengths, such as individual probability distributions, joint probability distributions and autocorrelation functions, are presented, as follows. 1) Individual Probability Distributions Histogram is a simplest and most straightforward technique to describe the probability distributions of random variables. Basically, the range of the sample data of a random variable is divided into a number of intervals. The number of sample data within each interval is counted so as to calculate the relative frequency. The determination of the interval number is important; the probability distribution of the random variable can not be properly presented if the number is too small. On the other hand, if the interval number is too large, then zero frequency may occur within some intervals. In the present work, the interval numbers are selected to be 7 based on a careful judgment. Fig. 3 displays the histograms of the interlaminar strengths at b = 0°, 10°, 80°, and 90°. 2) Joint Probability Distributions 3-D histogram is a standard and classical technique to describe the joint probability distributions of two random variables, that are calculated directly from the two sample data sets. A reasonable selection of interval numbers is important to properly present the joint distributions. Based on a judgment, it is found that 25 is a good interval number for the present sample data sets. In the present work, the 3-D histograms are further modified as follows; define smooth surfaces using the cubic interpolation technique through the midpoints of all the top areas of the intervals corresponding to the frequencies, and make the volumes under the surfaces equal to one. Fig. 4 presents the modified histograms for the joint probability distributions of interlaminar shear strengths between various combinations of sample data sets.

3) Autocorrelation Functions The correlation characteristics of the interlaminar shear strengths are quantified by the autocorrelation functions of the stochastic process s(b). As shown in Table 1, the mean value of s(b) is not a constant but changes with the f.o.d. angle b, therefore the stochastic process s(b) is nonstationary; the corresponding autocorrelation function is a function of both the initial f.o.d. angle and the lag distance (i.e., the difference between two f.o.d. angles can be estimated from the experimental sample data sets using the following equation:

page_312 Page 313 where s(bi) denotes the estimated standard deviation of each experimental sample data set. The estimated values of the coefficients of correlation are tabulated in Table 2. Table 2. The coefficients of correlation of the interlaminar shear strengths 0° 10° 80° 90° 0° 1 -0.1025 0.0551 0.0834 10° 1 -0.0335 -0.1038 80° 1 0.0285 90° 1 5 Conclusions and Discussions In the present stochastic model, the interlaminar shear strength is considered as a stochastic process, which is represented by a paramatric algorithm and a ARMA algorithm developed respectively. Experiments for the interlaminar shear strength are performed to yield the sample data set on which the development of the present stochastic model is based. The morphologies of fracture surfaces and thus the shear strengths are found to strongly depend on the fiber orientations of the neighboring plies of the fracture interface. Based on the experimental sample data sets, some important stochastic characteristics for the interlaminar shear strengths are presented. As can be seen from Fig. 3, not only the mean values but also the forms of individual probability distributions change with respect to the f.o.d. angles, thereby indicating the nonstationarity in the shear strengths. Also the randomness of the strength values does not follow the Gaussian or Weibull distributions, thus the maximum entropy technique is essential to adequately quantity the corresponding probability density functions. The values of the coefficient of correlation are smaller, and thus the correlation between the shear strengths corresponding to different f.o.d. angles is not strong. Therefore, the stochastic mechanism behind the shear strengths is different from that behind the cumulative damage process of laminated composites under fatigue loading [6], wherein the former damage states determine the impending damage in the process of damage evolution. References [1] Flaggs, D. L. and Kural, M. H., "Experimental Determination of the In Situ Transverse Lamina Strength in Graphite/Epoxy Laminates," Journal of Composite Materials, Vol. 16, 1982, pp. 103116. [2] Chang F. K. and Chen M. H., "The In Situ Ply Shear Strength Distributions in Graphite/Epoxy Laminated Composites," Journal of Composite Materials, Vol. 21, 1987, pp. 708733. [3] Robinson, P. and Song, D. Q., "A Modified DCB Specimen for Mode I Testing of Multidirectional Laminates," Journal of Composite Materials, Vol. 26, 1992, pp. 15541577. [4] Nicholls, D. J. and Gallagher, J. P., "Determination of GIC in Angle Ply Composites Using a Cantilever Beam Test Method," Journal of Reinforced Plasticsand Composites, Vol. 2, 1983, pp. 217. [5] Siddall, J. M., "Probabilistic Engineering Design: Principles and Applications," Marcel Dekker, Inc., New York. 1983. [6] Ganesan, R., Hoa, S. V., Zhang, S. and El-Karmalawy, M., "A Stochastic Cumulative Damage Model for the Fatigue Response of Laminated Composites," Proceedings of the 11th International Conference on Composite Materials, Australia, 1997.

[7] Shinozuka, M., "Stochastic Fields and Their Digital Simulation," Stochastic Methods in Structural Dynamics, Eds. Schueller, G. I. and Shinozuka, M., Martinus Nijhoff Publishers, Boston, 1987, pp. 93133. [8] Pandit S. M. and Wu, S. M., "Time Series and System Analysis with Applications," John Wiley and Sons, New York, 1983. [9] Ang, A. H. S. and Tang, W. K., "Probability Concepts in Engineering Planning and Design: Basic Principles," John Wiley and Sons, New York, 1984. [10] Swason Analysis Systems, INC., "ANSYS User's Manual," Version 5, 1992. page_313 Page 314

Fig. 1. The double notch shear test; (a) test setup, (b) specimen geometry.

Fig. 2. The distribution of interlaminar shear stress at the central planes, in which the specimens are subjected to a compressive load of -100 MPa; (a) notches without fillets, (b) notches with fillets. The straight lines denote the nominal stress with a uniform distribution at the central planes.

page_314 Page 315

Fig. 3. The histograms of the interlaminar shear strengths; (a) b = 0°, (b) b = 10°, (c) b = 80°, (d) b =90°.

Fig. 4. The joint probability distributions of the interlaminar shear strengths: (a) b = 0°and b = 10°, (b) b = 80° and b =90°, (c) b = 0° and b =80°, (d) b = 10° and b= 90°. In each plot, x and y axes (of the right hand coordinate system) correspond to angles b in the same order. page_315 Page 317

AUTHOR INDEX A Abdesslam, Y. O., 75 Akamizu, H., 184 Artus, M., 21 B Baba, F., 3 Baillie, M. R., 105

Bates, P. J., 205 Benoit, Y., 22 Bonin, H. W., 205 Breard, J., 22 Bui, V. T., 205 C Chatillon, A., 30 Chen, J.-Y., 249 D Denault, J., 7 Do-Thanh, V., 125 E Evangelista, E., 183 F Fitzgerald, S. B., 114 Fujii, Y., 212 Fujita, A., 3 Fujiwara, K., 15 G Ganesan, R., 308 Gauvin, R., 22 Georgiades, A. V., 114 Giray, M., 49 Goto, A., 235 H Habib, P., 216 Hamada, H., 91, 133, 149, 165, 282, 288, 300, 304 Hasegawa, T., 300 Hayasaki, S., 165, 282 Hirai, T., 199 Hoa, S. V., 149, 241, 249, 308 Honda, T., 87 Hosokawa, K., 153 Hubert, P., 106 I

Imai, T., 67 Itoh, K., 258 Iwasaki, K.-I., 278 J Jen, C. K., 49, 249 Johnston, A., 106 K Kageyama, K., 191, 296 Kalamkarov, A. L., 114 Kalaycioglu, S., 49 Kameo, K., 304 Kataoka, Y., 161, 274 Kawada, H., 87 Kijima, K., 184 Kim, P., 175 Kimpara, I., 191, 296 Kimura, T., 161, 274 Kitagawa, K., 91, 165, 282, 288 Kobayashi, H., 278 Kohli, V., 21 Kondo, Y., 274 Kotaki, M., 99, 258 Kuriyama, T., 99, 258 Kwan, E., 149 L Lee-Sullivan, P., 57 Legault, J.-F., 205 Li, J., 41 Lo, J., 41 Lopata, V. J., 105 M MacDonald, D. O., 114 Maekawa, Z., 262 McQueen, H. J., 183 Miyano, Y., 67 Murayama, T., 41

N Nagayama, K., 15 Naito, H., 300 Naji, M. I., 241 Nakada, M., 67 Nakahira, A., 184 Nakai, A., 133, 288, 304 Nakamura, H., 278 Narisawa, I., 99, 258 Ngo, A. D., 75 Nguyen-Hoa, H., 224 Nishijima, S., 184 Nishiwaki, T., 165, 262 Nishiyabu, K., 141 O Odaka, S., 99 Owens, J. F. P., 57 P Page, J. Y. S. D., 205 Pham, X.-T., 22 Phaneuf, M., 41 Poursartip, A., 106 Prasad, S. E., 49 R Raghavan, J., 105, 257 Remacle, J.-F., 22 S Saito, A., 266 Sakata, T., 153 Satoh, H., 87 Semba, T., 91 Shimamura, Y., 278 Sihn, S., 67 Suzuki, K., 191, 296 Suzuki, Y., 184 T

Takahashi, H., 266 Takahashi, T., 274 Takashima, M., 87 Takeda, N., 81, 133, 288, 304 Tange, A., 262 Tanigaki, E., 300 Tanimoto, Y., 262 Todoroki, A., 278 Trochu, F., 22 Tsai, S. W., 67 U Ueno, S., 184 Um, I., 266 V Vaziri, R., 106 Viswanathan, C. I., 257 Vu-Khanh, T., 125, 224 W Waas, A. M., 81 Wang, H., 49, 149, 249 Watanabe, Y., 199 page_317