Whole

THE UNIVERSITY OF NEW SOUTH WALES FACULTY OF APPLIED SCIENCE SCHOOL OF MATERIALS SCIENCE AND ENGINEERING

IMPROVEMENT OF BEARING STRENGTH OF LAMINATED COMPOSITES BY NANOCLAY AND Z-PIN REINFORCEMENT

by

TRAN PHUONG NAM HUONG

A thesis submitted in fulfillment Of the requirement for the degree of DOCTOR OF PHILOSOPHY

December, 2006

THE UNIVERSITY OF NEW SOUTH WALES Thesis/Dissertation Sheet

Surname or Family name: HUONG First name:

Tran Phuong Nam

Abbreviation for degree as given in the University calendar:

Other name/s:

N/A

Faculty:

Science

Ph.D

School:

Materials Science and Engineering

Title:

Improvement of bearing strength of laminated composites by nanoclay and z-pin reinforcement

Abstract The bearing behavior of bolted composite joints is significantly poorer than that of their metallic counterparts. The objective of the present study was to examine ways of improving the bearing performance of bolted joints in carbon fibre reinforced epoxy laminates. Two strategies were examined, namely stiffening of the matrix using nanofillers and through- thickness reinforcement of the laminates using z- pins. The development of a nanoparticle reinforced matrix resin, and its performance in a composite loaded in bearing, was the focus of the first part of the study. A commercial nanoclay, I30E from Nanocor, was chosen as the reinforcement since the nanoclay particles are modified with long alkyl chain amines which improves dispersion in the epoxy resin. The c onditions for preparing nanocomposites based on the nanoclay were examined for two epoxy resin systems , DGEBA and TGDDM. Significant improvements in the elastic modulus were obtained , with a 20% increase being recorded w ith 8.4 phr nanoclay content in the DGEBA resin and a 50% increase with 20 phr nanoclay content in the TGDDM system. Carbon fibre reinforced laminates w ere prepared from nanoclay reinforced TGGDM matrix resin, using a vacuum assisted prepregging process. The bearing strength and stiffness of the laminated composites was improved by 7% and 15% respectively but the strain to failure was reduced. The addition of nanoclay to the matrix resin was found to change the failure mode. Enhancement of the bearing performance of laminates by through-thickness reinforcement (z-pins) was examined as the second strategy trialed in this thesis. The bearing strength, stiffness and energy to failure of carbon fibre laminates was found to increase progressively with increasing volume content of z-pins , with increases of 10%, 10% and 16%, respectively , being achieved at a z-pin volume content of 4%. No significant change in strength, stiffness or failure energy was observed when the z- pin diameter was changed from 0.28 to 0.51 mm at the same volume content.

Declaration relating to disposition of project thesis/dissertation I hereby grant to the University of New South Wales or its agents the right to archive and to make available my thesis or dissertation in whole or in part in the University libraries in all forms of media, now or here after known, subject to the provisions of the Copyright Act 1968. I retain all property rights, such as patent rights. I also retain the right to use in future works (such as articles or books) all or parts of this thesis or dissertation. I also authorise University Microfilms to use the 350 word abstract of my thesis in Dissertation Abstracts International (this is applicable to doctoral theses only)

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ORIGINALITY STATEMENT

‘I hereby declare that this submission is my own work and to the best of my knowledge it contains no materials previously published or written by another person, or substantial proportions of material which have been accepted for the award of any other degree or diploma at UNSW or any other educational institution, except where due acknowledgement is made in the thesis. Any contribution made to the research by others, with whom I have worked at UNSW or elsewhere, is explicitly acknowledged in the thesis. I also declare that the intellectual content of this thesis is the product of my own work, except to the extent that assistance from others in the project’s design and conception or in style, presentation and linguistic expression is acknowledged.’

Signed....................................................... Date ..........................................................

ABSTRACT

The bearing behavior of bolted composite joints is significantly poorer than that of their metallic counterparts. The objective of the present study was to examine ways of improving the bearing performance of bolted joints in carbon fibre reinforced epoxy laminates. Two strategies were examined, namely stiffening of the matrix using nanofillers and through-thickness reinforcement of the laminates using z-pins. Stiffening of the resin matrix was carried out by reinforcement with nanoclay partic les. It was necessary to establish a technique for uniformly dispers ing the nanoclay particles in the epoxy resin and a systematic study was conducted. The effect of surfactant modification of the nanoclay particles was examined first. The effect of surfactant structure was examined using the surfactants octylamine, diaminoctane and methylbenzylamine, which all have the same chain length (8 carbon atoms in the chain), while the effect of chain length was examined by including hexadecylamine (16 carbon atoms in the chain). The short chain surfactants did not expand the layers in the nanoclay appreciably, but the longer chain surfactant produced a substantial increase in the interlayer spacing (from 11 Å to 18 Å) and was considered a suitable surfactant for modifying the nanoclay for reinforcing epoxy resin. However, similar results were obtained using the commercially available octadecylamine modified nanoclay, N anomer I30E, and this was used in the remainder of the work. DGEBA epoxy resin nanocomposites with up to 8.4 phr Nanomer I30E nanoclay were prepared and the effect of varying the processing parameters explored. It was found that exfoliation of the nanoclay particles was insensitive to the processing conditions within the range examined. Exfoliated nanoc omposites with a d-spacing of 100 Å were obtained for up to 5 phr nanoclay but this was reduced to 63 Å at 8.4 phr nanoclay. The iii

compression modulus of the nanocomposites increased progressively with nanoclay addition, with a 20% increase over that of the neat epoxy resin being achieved at the highest nanoclay loading. Nanocomposites were also prepared using the higher performance aircraft grade TGDDM epoxy resin with loadings of up to 20 phr Nanomer I30E nanoclay. Fully exfoliated nanocomposites with an interlayer spacing greater than 120 Å were obtained up to 5 phr nanoclay, while pre-exfoliated or intercalated nanocomposites, with interlayer spacings from 85 Å (7.5 phr nanoclay) to 60 Å (20 phr nanoclay), were obtained at higher loadings. The compression modulus increased linearly with clay content, with an increase of 50% being achieved over that of the pristine resin at 20 phr clay. The results were in good agreement with the predictions of the Halpin-Tsai model for an aspect ratio of 13. Carbon fibre reinforced laminates with a fibre volume fraction of approximately 55% were prepared using I30E nanoclay modified TGGDM as the matrix resin with clay loadings of 7.5 and 12.5 phr. The addition of nanoclay to the matrix produced a progressive increase in bearing stiffness with improvements of 10% and 23% being obtained at 7.5 and 12.5 phr nanoclay respectively. A more modest improvement in bearing strength was obtained with a 7% increase being achieved at 12.5 phr nanoclay. The strain to failure was reduced by the addition of nanoclay. The relatively poor increase in bearing strength and the reduced strain to failure are attributed to a change in failure mode brought about by the introduction of the nanoclay into the matrix resin. The spacing of the nanoclay layers in the laminates was only half that obtained in the clay-resin nanocomposites with the same clay content. This is attributed to the resin curing prematurely within the clay galleries during the prepreg drying stage. The effect of through thickness reinforcement on bearing strength was examined by reinforcing carbon fibre epoxy laminates locally in the vicinity of the hole with z-pins. Three different volume contents of z-pins were examined (0.5, 2, and 4%) and two different pin diameters (0.28 and 0.51 mm). The insertion of z-pins increased the bearing strength, stiffness and energy to failure. No change in the strain to failure was however observed. The increases are attributed to the bridging effect of the z-pins. The

iv

bearing strength increased linearly with increasing pin content being 7% for 0.5% pin content and reaching 10% at 4% pin content. The bearing stiffness was increased by 8% and 10% at the same pin contents while the energy absorbed to failure was correspondingly increased by 9% and 16% . These increases are attributed to the bridging effect of the z-pins. No significant change in strength, stiffness or failure energy was observed when the pin diameter was changed from 0.28 to 0.51 mm at the same volume content.

v

ACKNOWLEDGEMENTS

Doing a PhD thesis could be compared to navigation of a boat through the ocean till reaching the land. It is a result of interdependence, a cooperation of knowledge and experience of many people. Now I have a pleasant opportunity to express my gratitude for them who have given me uncountable help to complete this project. The first person I would like to gratefully thank with all my heart is Professor Alan Crosky who is not only my direct excellent supervisor, but also as close as my “father”. I owe him a lot of gratitude for giving me great supervision and valuable advice in both fields - research and life. No boundary is drawn between us, a supervisor and a student , only family sense existed instead. He may not realize how much happy I have come to know him in my life. Furthermore, I would like to express my deep gratitude to his wife, Peta, who always shared with me difficult times happening in my life during research period. I would like to thank my co-supervisors, Professor Don Kelly (from UNSW) and Dr. Ben Qi (from CRC-ACS) who kept an eye on the progress of my work as well as gave me great support during project time. I really appreciated valuable comments they made for my thesis. A great thank to Dr. Brian who always sent his great friendship with precious help in improving my writing skill as well as how to make a good structure of thinking. Joining with him, Gavin, Honghua, Singh, Fuhai, all my friends in School of Materials Research and Engineering, UNSW and my flatmates, Hoang, Nadia, Julia, created for

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me a feeling of being home at work. They filled full of the lonely time living far from my family by unforgettable memories originated from this great friendship. I am grateful for Katie Levick, Jenny Norman, Veira and all staff in EM Unit who trained and helped me using varieties of electron microscopes (TEM, SEM, AFM and FIB). Their support straight distributed to my successful work in research. I would like to thank Professor Chris Sorrel who installed a first brick for me to get into UNSW for studying. Special thanks to you, Mrs. Lana, who always took care of me with all her sympathy, to all staffs in the School of MSE for their valuable help and pleasure. The Australian Government and CRC-ACS Ltd. are acknowledged for the supply of IPRS and Nanocomposites scholarships, supporting in finance for my research till fully completed. I would also like to thank Professor Adrian Mouritz, and Dr Paul Chang, RMIT University, Melbourne, Australia for their assistanc e with the z-pinning. Deeply, I am very grateful to my beautiful wife for her love and patience, and to my parents, my brothers and my sisters for their uninterrupted encouragement.

Tran Phuong Nam HUONG

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Table of Contents

TABLE OF CONTENTS

Certificate of Originality.........................................................................................

ii

Abstract.....................................................................................................................

iii

Acknowledgement ....................................................................................................

vi

Table of Contents ..................................................................................................... viii List of Figures........................................................................................................... xvi List of Tables.......................................................................................................... xxviii List of Appendices................................................................................................. xxxi Nomenclature......................................................................................................... xxxii Chapter 1 Introduction..........................................................................................

1

References ..................................................................................................................

5

Chapter 2 Literature Review................................................................................

9

2.1. Introduction.........................................................................................................

9

2.2. Bearing performance ...........................................................................................

9

2.2.1. Introduction ...............................................................................................

9

2.2.2. Failure mechanism .................................................................................... 14 2.2.2.1. Fibre microbuckling ...................................................................... 16 2.2.2.2. Kinking (Kink-banding) ............................................................... 18 2.2.2.3. Shear cracking and delamination .................................................. 20 2.2.3. Factors affecting bearing failure ............................................................... 22 2.2.3.1. Stacking sequence and percentage of 0° plies .............................. 22

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Table of Contents

2.2.3.2. Laminate thickness ....................................................................... 23 2.2.3.3. Hole machine defects.................................................................... 24 2.2.3.4. Lateral clamping ........................................................................... 25 2.2.3.5. Matrix stiffness ............................................................................. 26 2.2.4. Rationale for present study........................................................................ 26 2.3. Epoxy Nanocomposites....................................................................................... 27 2.3.1. Definition and Composition ...................................................................... 27 2.3.1.1. Epoxy resin ................................................................................... 27 2.3.1.2. Nanoclay particles ........................................................................ 33 2.3.2. Structure of nanocomposites ..................................................................... 36 2.3.3. Properties of nanocomposites.................................................................... 38 2.3.3.1. Tensile properties.......................................................................... 38 2.3.3.2. Compressive properties ................................................................ 39 2.3.3.3. Impact properties .......................................................................... 40 2.3.3.4. Flexural properties ........................................................................ 41 2.3.3.5. Toughness ..................................................................................... 42 2.3.3.6. Barrier performance ...................................................................... 42 2.3.3.7. Other properties ............................................................................ 43 2.3.4. Synthesis ................................................................................................... 45 2.3.4.1. Physical methods .......................................................................... 45 2.3.4.2. Chemical methods ........................................................................ 46 2.3.4.3. Curing procedure .......................................................................... 49 2.3. 5. Factors affecting nanocomposite structure................................................ 53 2.3.5.1. Effect of nanoclay......................................................................... 54 2.3.5.2. Effect of surfactants ...................................................................... 54 2.3.5.3. Effect of matrix resin .................................................................... 55 ix

Table of Contents

2.3.5.4. Effect of hardeners........................................................................ 56 2.3.5.5. Effect of temperature .................................................................... 58 2.3.5.6. Effect of time ................................................................................ 59 2.3.5.7. Effect of pressure .......................................................................... 60 2.4. Z -pin reinforcement............................................................................................. 61 2.4.1. Concepts.................................................................................................... 61 2.4.2. Mechanism................................................................................................ 63 2.4. 3. Important factors ....................................................................................... 67 2.4.3.1. Pin density .................................................................................... 67 2.4.3.2. Pin diameter .................................................................................. 68 2.4.3.3. Other factors ................................................................................. 69 2.4. 4. Pinning process ......................................................................................... 69 2.5. Characterization .................................................................................................. 71 2.5.1. Structure Characterization......................................................................... 71 2.5.1.1. Wide Angle X-ray Diffraction ...................................................... 71 2.5.1.2. Transmission Electron Mic roscopy .............................................. 72 2.5.1.3. Scanning Electron Microscopy..................................................... 73 2.5.1.4. Optical Microscopy ...................................................................... 74 2.5.2. Thermal property analysis......................................................................... 75 2.5.2.1. Differentional Scanning Calorimetry ........................................... 75 2.5.3. Mechanical Behaviour ............................................................................... 76 2.5.3.1. Compression Testing .................................................................... 76 2.5.3.2. Bearing Test .................................................................................. 77 2.5.3.3. Pin-contact Bearing Test............................................................... 78 2.6. Aims and Objectives of Thesis ............................................................................ 80 2.7. Scope ................................................................................................................... 81 x

Table of Contents

References .................................................................................................................. 82 Chapter 3 Nanocomposites.................................................................................... 92 3.1. Introduction......................................................................................................... 92 3.2. Experimental procedures..................................................................................... 93 3.2.1. Nanoclay modification .............................................................................. 93 3.2.1.1. Materials ....................................................................................... 93 3.2.1.2. Sample preparation ....................................................................... 94 3.2.1.3. Analysis ....................................................................................... 95 3.2.2. DGEBA nanocomposites .......................................................................... 96 3.2.2.1. Materials ....................................................................................... 96 3.2.2.2. Sample preparation ....................................................................... 97 3.2.2.3. Analysis ....................................................................................... 97 3.2.3. TGDDM nanocomposites ......................................................................... 100 3.2.3.1. Materials ....................................................................................... 100 3.2.3.2. Sample preparation ....................................................................... 100 3.2.3.3. Analysis ....................................................................................... 101 3.2.4. Experimental variables .............................................................................. 101 3.3. Results ................................................................................................................. 102 3.3.1. Nanoclay modification ............................................................................. 102 3.3.1.1. Effect of surfactant types .............................................................. 104 3.3.1.2. Effect of acid/amine ratio ............................................................. 105 3.3.1.3. Effect of surfactant concentration................................................. 106 3.3.1.4. Effect of mixing time.................................................................... 107 3.3.2. DGEBA Nanocomposites ......................................................................... 108 3.3.2.1. Cure temperature .......................................................................... 109 3.3.2.2. Effect of nanoclay type ................................................................. 110 xi

Table of Contents

3.3.2. 3. Effect of mixing temperature........................................................ 116 3.3.2. 4. Effect of mixing speed.................................................................. 118 3.3.2. 5. Effect of mixing time.................................................................... 119 3.3.2.6. Effect of hardener concentration .................................................. 121 3.3.2.7. Effect of curing temperature......................................................... 123 3.3.2. 8. Effect of curing time..................................................................... 124 3.3.2. 9. Effect of nanoclay content ............................................................ 125 3.3.2. 10. Optical properties ........................................................................ 128 3.3.3. TGDDM Nanocomposites......................................................................... 129 3.3.3.1. Effect of surfactant on curing behaviour ...................................... 128 3.3.3.2. Effect of mixing speed.................................................................. 131 3.3.3.3. Effect of mixing temperature........................................................ 134 3.3.3.4. Effect of mixing time.................................................................... 135 3.3.3.5. Effect of degassing time ............................................................... 138 3.3.3.6. Effect of curing temperature......................................................... 140 3.3.3.7. Effect of curing time..................................................................... 143 3.3.3.8. Effect of nanoclay content ............................................................ 146 3.4. Discussion ........................................................................................................... 150 3.4.1. Nanoclay modification .............................................................................. 150 3.4.1.1. Effect of surfactants ...................................................................... 150 3.4.1.2. Effect of acid/amine ratio ............................................................. 151 3.4.1.3. Effect of surfactant concentration................................................. 152 3.4.1. 4. Effect of mixing time.................................................................... 153 3.4.1. 5. Conclusion .................................................................................... 153 3.4.2. DGEBA Nanocomposites ......................................................................... 153 3.4.2.1. Effect of nanoclay type ................................................................. 154 xii

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3.4.2.2. Effect of mixing conditions .......................................................... 156 3.4.2.2.1. Mixing temperature........................................................ 156 3.4.2.2.2. Mixing speed.................................................................. 157 3.4.2. 2.3. Mixing time.................................................................... 157 3.4.2. 3. Effect of curing conditions ........................................................... 157 3.4.2.3.1. Hardener concentration .................................................. 157 3.4.2.3.2. Cure temperature............................................................ 158 3.4.2.3.3. Cure time........................................................................ 159 3.4.2. 4. Effect of nanoclay content ............................................................ 159 3.4.2.5. Optical properties.......................................................................... 160 3.4.2. 6. Conclusion .................................................................................... 161 3.4.3. TGDDM Nanocomposites......................................................................... 161 3.4.3.1. Effect of surfactant on cure ........................................................... 162 3.4.3.2. Effect of mixing conditions .......................................................... 162 3.4.3.2.1. Mixing speed.................................................................. 162 3.4.3.2.2. Mixing temperature........................................................ 162 3.4.3.2.3. Mixing time.................................................................... 163 3.4.3.3. Effect of curing conditions ........................................................... 163 3.4.3.3.1. Degassing time ............................................................... 163 3.4.3.3.2. Effect of curing temperature .......................................... 164 3.4.3.3.3. Effect of curing time ...................................................... 164 3.4.3.4. Effect of nanoclay content ............................................................ 165 3.4.3. 5. Conclusions................................................................................... 170 3.5. Conclusions ......................................................................................................... 172 References .................................................................................................................. 175 Chapter 4 Laminated Nanocomposites................................................................ 178 xiii

Table of Contents

4.1. Introduction......................................................................................................... 178 4.2. Experimental procedure ...................................................................................... 179 4.2.1. Materials .................................................................................................... 179 4.2.2. Sample preparation.................................................................................... 179 4.2.3. Testing and analysis ................................................................................. 182 4.2.3.1. Pin-contact bearing test................................................................. 182 4.2.3.2. Microstructure ............................................................................... 184 4.2.3.3. Fracture surfaces ........................................................................... 184 4.2.3.4. Wide angle X-ray diffraction........................................................ 185 4.3. Results ................................................................................................................. 185 4.3.1. Pin-contact bearing test ............................................................................. 185 4.3.2. Laminate quality........................................................................................ 192 4.3. 3. Microstructural examination ..................................................................... 194 4.3. 4. Fracture surfaces ....................................................................................... 202 4.3. 5. Interlayer spacing ...................................................................................... 204 4.4. Discussion ........................................................................................................... 206 4.5. Conclusions ......................................................................................................... 211 References .................................................................................................................. 213 Chapter 5 Z-pin reinforcement............................................................................. 214 5.1. Introduction......................................................................................................... 214 5.2. Experimental procedure ...................................................................................... 215 5.2.1. Materials .................................................................................................... 215 5.2.2. Sample preparation.................................................................................... 215 5.2.3. Bearing testing........................................................................................... 219 5.2.4. Microstructural examination ..................................................................... 221 5.3. Results ................................................................................................................. 221 xiv

Table of Contents

5.3.1. Bearing testing........................................................................................... 221 5.3.2. Microstructure ........................................................................................... 231 5.4. Discussion ........................................................................................................... 239 5.5. Conclusions ......................................................................................................... 244 References .................................................................................................................. 245 Chapter 6 Conclusions ........................................................................................... 246 Appendix A Compression Tests Appendix B Pin-contact test results for laminated nanocomposites Appendix C Students’ T-Test for laminated nanocomposites Appendix D Bearing test results for Z-pin reinforcement

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List of Figures

LIST OF FIGURES

Figure 2.1

Materials used in F/A 18 fighter aircraft ............................................. 10

Figure 2.2

Schematic of joints used in hybrid metal/composite wing ................... 11

Figure 2.3

Joint efficiency of different materials .................................................. 12

Figure 2.4

Bolted joint failure modes..................................................................... 13

Figure 2.5

Comparison of compression and bearing failure mechanisms ............. 13

Figure 2.6

Bearing damage process detected by acoustic emission (AE).............. 14

Figure 2.7

Schematic illustration of the failure process......................................... 15

Figure 2.8

Initial configuration and buckling modes investigated by Rosen......... 17

Figure 2.9

In-plane buckling of fibres and fibre kink band geometry ................... 18

Figure 2.10 Schematic of fibre failure sequence in shear triggered kinkband formation ...................................................................................... 19 Figure 2.11 Geometry of kink-band formation........................................................ 20 Figure 2.12 Micrograph of shear cracking in bearing failure .................................. 21 Figure 2.13 Shear cracking and delamination .......................................................... 21 Figure 2.14 Open hole average compressive strength as a function of specimen thickness for multidirectional laminates ............................... 24 Figure 2.15 Schematic of curing reaction of epoxy/amine system .......................... 32 Figure 2.16 Schematic of curing reaction of epoxy/anhydride system.................... 32 Figure 2.17 Structure of Montmorillonite................................................................ 34

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List of Figures

Figure 2.18 Schematic illustrations of conventional composite and nanocomposites..................................................................................... 36 Figure 2.19 TEM images of intercalated (left) and exfoliated (right) nanocomposites..................................................................................... 37 Figure 2.20 Representative scattering curves for intercalated and exfoliated morphology for epoxy/SC18 layered silicate nanocomposites at several stages of cure ........................................................................ 37 Figure 2.21 Effect of nanoclay content on the modulus of nanocomposites ........... 39 Figure 2.22 Compressive modulus of nanocomposite and filler composites with clay loading ................................................................................... 40 Figure 2.23 Effect of nanoclays on impact strength of nanocomposite ................... 41 Figure 2.24 Illustration of Neilson’s tortuous path model for barrier enhancement of nanocomposites .......................................................... 43 Figure 2.25 Optical image of 5 mm thick plaques of nanocomposites based on EPON 828 cured by D400 and contains different loadings of nanoparticles..................................................................................... 44 Figure 2.26 Schematic figures for the exfoliation process of clay layers with DGEBA on the mixing process .................................................... 50 Figure 2.27 D-spacing of clay layers with the degree of conversion at some isothermal curing temperatures of 120, 130, 140°C in the DGEBA-DDS-C18 clay (5 phr) ............................................................ 51 Figure 2.28 Schematic illustration of the intercalated state and exfoliation process showing the forces acting on a pair of clay layers: (a) organically modified clay, (b) epoxy-intercalated state, (c) forces acting on a two-particle tactoid .................................................. 52 Figure 2.29 Schematic diagram showing the relationship between the ionic bonding energy and the location of the layers in the tactoid: (a) tactoid, (b) variation of bonding energy along the thickness of the tactoid .............................................................................................. 53

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List of Figures

Figure 2.30 Transmission electron micrographs of 10 phr C18 clay reinforced epoxy nanocomposites cured by (a) MDA (4,4’methylene -dianiline) and (b) DDS (4,4’-diaminodiphenyl sulfone).................................................................................................. 56 Figure 2.31 Small angle X-ray scattering patterns of 1% montmorillonite containing DGEBA epoxy cured by MPDA with (a) 25 phr; (b) 14.5 phr; (c) 5 phr ................................................................................. 57 Figure 2.32 Isothermal time-resolved SXRD patterns of nanocomposites at 120°C with scan time (minute) from bot tom to top as follow: 0, 5, 10, 11, 15, 20, 25, 30, 40, 60 and 89 min.......................................... 59 Figure 2.33 Time-dependent small angle synchrotron X-ray scattering patterns of the epoxy/montmorillonite mixture at 1350C ..................... 60 Figure 2.34 Z-pin performs containing 0.28 mm diameter pins at densities of 0.5, 2 and 4% .................................................................................... 62 Figure 2.35 Micrograph of z-pin reinforcement ....................................................... 63 Figure 2.36 Typical morphology of local area around a z-pin................................. 64 Figure 2.37 Sketch and SEM micrograph showing fibres deflecting around z-pins and weaving through a field of z-pins........................................ 64 Figure 2.38 Crack around z-pin in laminated composite ......................................... 65 Figure 2.39 Scanning electron micrograph showing a crack that initiated near a pin under flexural load ............................................................... 66 Figure 2.40 Delamination of laminate pinned with (a) 0.28 mm and (b) 0.51 mm pins................................................................................................. 68 Figure 2.41 Schematic of the z-fibre reinforcement process ................................... 69 Figure 2.42 Schematic of ultrasonically assisted z-fibre insertion........................... 70 Figure 2.43 Schematic depicting the expected X-ray diffraction patterns for various types of hybrid structures ......................................................... 71 Figure 2.44 TEM images of clay nanocomposite at low (left) and high (right) magnifications ............................................................................ 73

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List of Figures

Figure 2.45 SEM fractograph of 7 wt.% clay nanocomposite ................................. 73 Figure 2.46 Optical micrograph showing large particles the DETDA cured DGEBA system containing 5 wt.% clay............................................... 74 Figure 2.47 DSC of Epon 826/W, 3% SC8/Epon 826/W, and 3% SC18/Epon 862/W at 2°C/min .............................................................. 75 Figure 2.48 Typical stress - strain curve for a compression test.............................. 76 Figure 2.49 Typical stress - strain curve for a bearing test ...................................... 77 Figure 2.50 Fixture assembly for a bearing test ....................................................... 78 Figure 2.51 Schematic diagram of pin-contact test.................................................. 79 Figure 2.52 Typical load - displacement curve for cross-ply laminate .................... 79

Figure 3.1

Schematic of nanoclay modification..................................................... 94

Figure 3.2

XRD patterns of modified nanoclay..................................................... 105

Figure 3.3

XRD patterns of modified nanoclay in different acid/amine ratios...................................................................................................... 106

Figure 3.4

XRD patterns of nanoclay modified with various concentration of hexadecylamine ................................................................................ 107

Figure 3.5

XRD patterns of nanoclay modified with hexadecylamine for various mixing time .............................................................................. 108

Figure 3.6

Exothermal curves of curing reaction for SP Systems amine/epoxy system with different nanoclay types .............................. 110

Figure 3.7

Compression modulus for nanocomposites reinforced with modified nanoclay particles and pure nanoclay (CNa+) as well as for neat epoxy resin (PR) showing (1) effect of surfactants, (2) effect of acid/amine ratio, (3) effect of surfactant concentration and (4) effect of mixing time .......................................... 113

Figure 3.8

WAXD spectra for nanocomposites reinforced with C30B and I30E nanoclay ....................................................................................... 114

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List of Figures

Figure 3.9

TEM micrographs of nanocomposite reinforced with C30B (a) and I30E (b) .......................................................................................... 115

Figure 3.10 Compression modulus of pure resin and nanocomposites reinforced w ith different nanoclay particles ......................................... 115 Figure 3.11 WAXD spectra for nanocomposites made with varying mixing temperatures .......................................................................................... 116 Figure 3.12 Compression modulus of nanocomposites based on I30E nanoclay made with different mixing temperatures.............................. 117 Figure 3.13 WAXD spectra of nanocomposites made with different mixing speeds.................................................................................................... 118 Figure 3.14 Compression modulus of nanocomposites for different mixing speeds.................................................................................................... 119 Figure 3.15 WAXD spectra of nanocomposites made with different mixing times ...................................................................................................... 120 Figure 3.16 Compression modulus of nanocomposites for different mixing times ..................................................................................................... 120 Figure 3.17 Diagrams showing (a) compression modulus of nanocomposites for varying hardener concentrations and (b) % increase over that of neat resin.............................................................. 122 Figure 3.18 WAXD spectra of nanocomposite for varying hardener concentrations ....................................................................................... 122 Figure 3.19 WAXD spectra of nanoc omposites for different curing temperatures .......................................................................................... 123 Figure 3.20 Compression modulus of nanocomposites for different curing temperatures .......................................................................................... 124 Figure 3.21 Compression modulus of nanocomposites for different curing times ...................................................................................................... 125 Figure 3.22 WAXD spectra of nanocomposites for different nanoclay contents .................................................................................................... 126

xx

List of Figures

Figure 3.23 Diagrams showing (a) compression modulus of nanocomposites for different nanoclay contents and (b) % increase over that of neat resin.............................................................. 127 Figure 3.24 Optical appearance of pure resin and nanocomposites based on I30E nanoclay ....................................................................................... 128 Figure 3.25 EDS analysis of nanocomposite: (A) montmorillonite nanoclay particles, (B) aluminum oxide and (C) resin......................................... 129 Figure 3.26 Exothermal curves for cure reaction of TGDDM/DETDA with various nanoclay loadings..................................................................... 130 Figure 3.27 Compress ion modulus of nanocomposites made w ith different mixing speeds ........................................................................................ 132 Figure 3.28 WAXD spectra of nanocomposites with (a) 2.5 phr and (b) 7.5 phr nanoclay.......................................................................................... 133 Figure 3.29 Compress ion modulus of nanocomposites for different mixing temperatures .......................................................................................... 134 Figure 3.30 WAXD spectra of nanocomposites with (a) 2.5 phr and (b) 7.5 phr nanoclay.......................................................................................... 135 Figure 3.31 Diagrams showing (a) compression modulus of nanocomposites for different mixing times and (b) % increase over that of neat resin............................................................................ 136 Figure 3.32 WAXD spectra of nanocomposites with (a) 2.5 phr and(b) 7.5 phr nanoclay............................................................................................ 137 Figure 3.33 Diagrams showing (a) compression modulus of nanocomposite made with 7.5 phr nanoclay for various additional degassing times and (b) % increase over that of neat resin ................................... 139 Figure 3.34 WAXD spectra of nanocomposites containing 7.5 phr nanoclay for various additional degassing times .................................. 139 Figure 3.35 Compression modulus of nanocomposites with varying initial (I), intermediate (II) and final (III) curing temperatures ....................... 141

xxi

List of Figures

Figure 3.36 Percentage increase of compressive modulus over that of neat resin for nanocomposites cured at various initial (I), intermediate (II) and final (III) curing temperatures............................. 141 Figure 3.37 WAXD spectra of nanocomposites for various initial cure temperatures .......................................................................................... 142 Figure 3.38 WAXD spectra of nanocomposites for various intermediate cure temperatures .................................................................................. 142 Figure 3.39 WAXD spectra of nanocomposites for various final cure temperatures .......................................................................................... 143 Figure 3.40 Compression modulus of nanocomposites cured for different initial (I), intermediate (II) and final (III) cure times ............................ 144 Figure 3.41 Percentage increase of compression modulus over that of neat resin for nanocomposites cured at various initial (I), intermediate (II) and final (III) cure times............................................ 144 Figure 3.42 WAXD spectra of nanocomposites for various initial curing times ...................................................................................................... 145 Figure 3.43 WAXD spectra of nanocomposites for various final curing times ...................................................................................................... 145 Figure 3.44 WAXD spectra for nanocomposites for varying I30E nanoclay contents ................................................................................................. 146 Figure 3.45 High magnification TEM images of nanocomposites with various nanoclay contents: (a) 2.5 phr, (b) 7.5 phr, (c) 12.5 phr and (d) 20 phr ........................................................................................ 147 Figure 3.46 Compression modulus of TGDDM/DETDA/I30E nanocomposites .................................................................................... 148 Figure 3.47 Percentage increase in compression modulus over that of neat resin for nanocomposites made with varying nanoclay contents.......... 148 Figure 3.48 Tg of nanocomposites with varying nanoclay contents ........................ 149

xxii

List of Figures

Figure 3.49 SEM images of various nanoclay contents in the resin matrix: (a) 2.5 phr, (b) 7.5 phr, (c) 12.5 phr and (d) 20 phr .............................. 150 Figure 3.50 Light transmission spectra of pure epoxy, Ep/CM nanocomposites with different CM contents ........................................ 160 Figure 3.51 Comparison of experimental results with pre dicted results for aspect ratios of 13, 40 and 100.............................................................. 170

Figure 4.1

Wet prepreg plies drying in laboratory................................................. 180

Figure 4.2

Prepreg lay-up (a) and sealed vacuum bag (b) ..................................... 180

Figure 4.3

Vacuum bagging of lay up (a) and layup being inserted into hot press (b) ........................................................................................... 181

Figure 4.4

Procedure for producing pin-contact bearing test specimens: (a) specimen with drilled hole, (b) testpiece produced by slitting specimen through hole .............................................................. 182

Figure 4.5

Schematic of the pin-contact bearing test employed by Wu and Sun ........................................................................................................ 183

Figure 4.6

Diagram showing positions where thickness was measured on specimens .............................................................................................. 183

Figure 4.7

Production of Mode 1 fracture surface ................................................. 184

Figure 4.8

Typical pin-contact load-displacement curves for laminates made with neat resin (baseline) and nanoclay reinforced resin ............ 185

Figure 4.9

Optical micrograph showing un-reinforced laminate ........................... 192

Figure 4.10 Optical micrograph showing laminate reinforced with 7.5 phr nanoclay (55 vol.% fibres) .................................................................... 193 Figure 4.11 Optical micrograph showing laminate reinforced with 12.5 phr nanoclay................................................................................................ 193 Figure 4.12 Optical micrograph showing un-reinforced 6K fabric laminate ........... 194

xxiii

List of Figures

Figure 4.13 Optical micrograph showing 6K fabric laminate reinforced with 7.5 phr nanoclay............................................................................ 194 Figure 4.14 Optical micrographs showing local failure modes in the laminated composites: kinking (K), matrix cracking (M), shearing cracking (S), delamination (D) ............................................... 195 Figure 4.15 Optical micrograph showing bearing damage in unreinforced laminate containing 56.9 vol.% 3K carbon fibres (a) low magnification view and (b) enlarged view of region A in (a) ............... 196 Figure 4.16 Optical micrograph showing bearing damage in laminate reinforced with 7.5 phr nanoclay containing 57.6 vol.% 3K carbon fibres (a) low magnification view and (b) enlarged view of region A in (a) ................................................................................... 197 Figure 4.17 Optical micrograph showing bearing damage in laminate reinforced with 7.5 phr nanoclay containing 55 vol.% 3K carbon fibres (a) low magnification view and (b) enlarged view of region A in (a) ................................................................................... 198 Figure 4.18 Optical micrograph showing bearing damage in laminate reinforced with 12.5 phr nanoclay containing 51.7 vol.% 3K carbon fibres (a) low magnification view and (b) enlarged view of region A in (a) ................................................................................... 199 Figure 4.19 Optical micrograph showing bearing damage in unreinforced laminate containing 61.3 vol.% of 6k carbon fibres ............................. 200 Figure 4.20 Optical micrograph showing bearing damage in laminate reinforced with 7.5 phr nanoclay and containing 61.1 vol.% 6K carbon fibres (a) low magnification view and (b) enlarged view of region A in (a) ................................................................................... 201 Figure 4.21 Mode 1 fracture surface of baseline (left) and 7.5 phr nanoclay reinforced (right) laminates ................................................................... 202

xxiv

List of Figures

Figure 4.22 SEM images of the fracture surfaces of unreinforced (a, b) and nanoclay reinforced (7.5 phr I30E) (c, d) laminates with 3K tow carbon fibres ................................................................................... 203 Figure 4.23 SEM images of fracture surfaces of unreinforced (a, b) and nanoclay reinforced (7.5 phr I30E) (c, d) laminates with 6K tow carbon fibres ................................................................................... 203 Figure 4.24 WAXD spectra of unreinforce d and reinforced (7.5 phr and 12.5 phr I30E nanoclay) laminates with 3K tow carbon fibres ............ 204 Figure 4.25 WAXD spectra of unreinforced and reinforced (7.5 phr I30E nanoclay) laminates with 6K tow carbon fibres .................................... 205 Figure 4.26 WAXD spectra of uncured 3K tow carbon fibre prepreg containing 7.5 phr I30E nanoclay......................................................... 205 Figure 4.27 Normalised bearing strength from pin-contact test for laminates reinforced with 3K tow fabric (unfilled) and 6K tow fabric (filled) with various nanoclay contents (0, 7.5 and 12.5 phr) ........................................................................................................ 207 Figure 4.28 Slope of load-displacement curves of unreinforced and reinforced laminates with 3k tow fabric (unfilled) and 6k tow carbon fibres (filled) .............................................................................. 208 Figure 4.29 Displacement at failure of unreinforced and reinforced laminates with 3K tow carbon fibres (unfilled) and 6K tow carbon fibres (filled) .............................................................................. 209

Figure 5.1

Area reinforced by z-pins around 10 mm bolthole ............................... 216

Figure 5.2

Photographs showing (a) pins embedded in foam preform and (b) use of ultrasonic horn to insert pins into prepreg stack................... 217

Figure 5.3

Photographs showing (a) completely inserted pins in prepreg and (b) cutting the preform and remaining portion of the pins from the prepreg stack........................................................................... 218

Figure 5.4

Cure cycle used to cure laminates (Hexcel) .......................................... 219 xxv

List of Figures

Figure 5.5

Schematic representation of fixture assembly for pin loaded bearing test ASTM D5961.................................................................... 220

Figure 5.6

Positions at which thickness measurements were made ....................... 221

Figure 5.7

Z-pinned and baseline laminates after bearing test............................... 221

Figure 5.8

Typical load - displacement curves for baseline and z-pinned laminates ............................................................................................... 222

Figure 5.9

Variation in bearing strength with pin density for 0.28 mm pins........................................................................................................ 227

Figure 5.10 Variation in bearing stiffness (slope to failure) with pin density for 0.28 mm pins ................................................................................... 228 Figure 5.11 Variation in energy absorbed to failure with pin density for 0.28 mm pins......................................................................................... 229 Figure 5.12 Comparison of bearing strength for 0.28 and 0.51 mm z-pins at 2% volume density................................................................................ 230 Figure 5.13 Comparison of bearing stiffness for 0.28 and 0.51 mm z-pins at 2% volume density............................................................................ 230 Figure 5.14 Comparison of energy absorbed to failure for 0.28 and 0.51 mm z-pins at 2% volume density.......................................................... 231 Figure 5.15 Optical micrograph of unpinned laminate loaded to 95% of ultimate failure load .............................................................................. 232 Figure 5.16 Optical micrograph of unpinned laminate at final failure..................... 232 Figure 5.17 Optical micrograph of laminate containing 0.5 % 0.28 mm zpins loaded to 95% of ultimate failure load (a) full section view and (b) detailed view of region marked A in (a ) showing fibre kinking .......................................................................................... 233 Figure 5.18 Optical micrograph of laminate containing 0.5 % 0.28 mm zpins loaded to ultimate failure ............................................................... 234 Figure 5.19 Optical micrograph of laminate containing 2 % 0.28 mm zpins loaded to 95% of ultimate failure load (a) full section

xxvi

List of Figures

view and (b) detailed view of region marked A in (a) showing shear bands ............................................................................................ 235 Figure 5.20 Optical micrograph of laminate containing 2 % 0.28 mm zpins loaded to ultimate failure load (a) full section view and (b) - (e) detailed views of regions marked A – D in (a)........................ 236 Figure 5.21 Optical micrograph of laminate containing 4 % 0.28 mm zpins loaded to ultimate failure ............................................................... 237 Figure 5.22 Optical micrograph of laminate containing 2 % 0.51 mm zpins loaded to 95% of ultimate failure load (a) full section view and (b) and (c) detailed views of regions marked A and B in (a) ...................................................................................................... 238 Figure 5.23 Optical micrograph of laminate containing 2 % 0.51 mm zpins loaded to ultimate failure load (a) full section view and (b) detailed views of region marked A in (a).............................................. 239 Figure 5.24 Optical micrograph of curvature of 0° fibres in pinned laminate ......... 241 Figure 5.25 SEM micrograph of a crack beside a z-pin .......................................... 243

xxvii

List of Tables

LIST OF TABLES

Table 2.1

Classification and generalized structural formulae of phyllosilicates.....

33

Table 2.2

Notional structure and chemistry of smectites........................................

34

Table 3.1

Effect of surfactant on nanoclay modification........................................ 101

Table 3.2

E ffect of nanoclay type on properties of DGEBA composites ............... 101

Table 3.3

Effect of processing conditions on properties of I30E/DGEBA composites ....................................................................... 102

Table 3.4

Effect of processing conditions on properties of I30E/TGDDM composites ...................................................................... 102

Table 3. 5 S urfactants and experimental conditions used........................................ 104 Table 3.6 Effect of acid/amine ratio on d-spacing of nanoclay .............................. 105 Table 3.7 Effect of surfactant concentration on d-spacing of nanoclay.................. 107 Table 3.8 Effect of the mixing time on d-spacing of nanoclay............................... 108 Table 3.9 D-spacing in modified CNa+ nanoclay nanocomposites ......................... 111 Table 3.10 Compression modulus for modified CNa+ nanoclay nanocomposites ....................................................................................... 112 Table 3.11 Effect of nanoclay type on d-spacing and compression modulus ........... 114 Table 3.12 Effect of mixing temperature on compression modulus ......................... 117 Table 3.13 Effect of hardener concentration on compression modulus.................... 121

xxviii

List of Tables

Table 3.14 Effect of curing time on d-spacing and compression modulus ............... 124 Table 3.15 Effect of nanoclay content on d-spacing and modulus of nanocomposites ....................................................................................... 125 Table 3.16 Peak and onset temperatures of nanocomposite s having varying nanoclay contents .................................................................................... 131 Table 3.17 Effect of cure temperature on compression modulus of nanocomposites ....................................................................................... 140 Table 3.18 Effect of curing time on compression modulus of nanocomposites ....................................................................................... 143

Table 4.1

Measured values of thickness, failure load, bearing strength, slope to failure and displacement to failure for neat epoxy laminates (baseline) containing 56.9 vol.% carbon fibre........................ 186

Table 4.2

Measured values of thickness, failure load, bearing strength, slope to failure and displacement to failure for laminates with 7.5 phr I30E nanoclay and 57.6 vol.% carbon fibre................................ 187

Table 4.3

Measured values of thickness, failure load, bearing strength, slope to failure and displacement to failure for laminates with 7.5 phr I30E nanoclay and 55 vol.% carbon fibre ................................... 188

Table 4.4

Measured values of thickness, failure load, bearing strength, slope to failure and displacement to failure for laminates with 7.5 phr I30E nanoclay and 51.7 vol.% carbon fibre................................ 189

Table 4.5

Summary of results for baseline and nanocomposite laminates (3K tow fabric) ........................................................................................ 189

Table 4.6

Measured values of thickness, failure load, bearing strength, slope to failure and displacement to failure for baseline (neat resin) laminates containing 61.3 vol.% carbon fibres ............................. 190

Table 4.7

Measured values of thickness, failure load, bearing strength, slope to failure and displacement to failure for laminates with 7.5 phr I30E nanoclay and 61.1 vol.% carbon fibres .............................. 191 xxix

List of Tables

Table 4.8

Summary of results for baseline and nanocomposite laminates (6K tow fabric) ........................................................................................ 191

Table 4.9

Bearing strength of CFRPs with or without nanoclay reinforcement .......................................................................................... 206

Table 4.10 Toughness estimated from area under stress strain curve ....................... 210

Table 5.1

Measured values of thickness, failure load, bearing strength, energy absorbed to failure, slope to failure and displacement to failure for baseline samples..................................................................... 223

Table 5.2

Measured values of thickness, failure load, bearing strength, energy absorbed to failure, slope to failure and displacement to failure for laminates pinned with 0.5% 0.28 mm pins ............................ 224

Table 5.3

Measured values of thickness, failure load, bearing strength, energy absorbed to failure, slope to failure and displacement to failure for laminates pinned with 2% 0.28 mm pins ............................... 224

Table 5.4

Measured values of thickness, failure load, bearing strength, energy absorbed to failure, slope to failure and displacement to failure for laminates pinned with 4% 0.28 mm pins ............................... 225

Table 5.5

Measured values of thickness, failure load, bearing strength, energy absorbed to failure, slope to failure and displacement to failure for laminates pinned with 2% 0.51 mm pins ............................... 225

Table 5.6

Summary of changes in mechanical behaviour produced in zpinning..................................................................................................... 226

xxx

LIST OF APPENDICES

Appendix A

Compression Tests

Appendix B

Pin-contact test results for laminated nanocomposites

Appendix C

Students’ T-Test for laminated nanocomposites

Appendix D

Bearing test results for Z-pin reinforcement

xxxi

NOMENCLATURE

Arabic

Description

d.............................. hole diameter d001 ........................... d-spacing of nanoclay layers E ............................. elastic modulus e .............................. edge distance G ............................. shear modulus F B ............................ bearing strength of laminate F TU ........................... tensile strength J .............................. joint efficiency l............................... length m............................. weight P.............................. bearing load P B............................ maximum bearing load P TU........................... ultimate tensile load t............................... thickness T g ............................ glass Transition Temperature V ............................. volume fraction w ............................. width x .............................. content of nanoclays

xxxii

y.............................. concentration of hardener

Greek

Description

? .............................. shape parameter a .............................. aspect ratio s .............................. stress

Subscripts

Descriptions

c .............................. clay property cr............................. microbuckling property ep............................ elastic-plastic property f............................... fibre property L.............................. longitudinal property m ............................. matrix property n .............................. nanocomposite property p .............................. platelet property T.............................. transverse property xx ............................ compressive direction

Abbreviations

Descriptions

AE........................... Acoustic Emission ASTM..................... American Society for Testing and Materials CEC........................ Cation-Exchange Capacity of clay CFRP ...................... Carbon Fibre Reinforced Polymer CNa+ ....................... Pure sodium nanoclay CRC-ACS............... Cooperative Research Centre for Advanced Composite Structures xxxiii

D ............................. Delamination DDS........................ Diaminodiphenylsufone DETA ..................... Diethylene Triamine DETDA .................. Diethyltoluene Diamine DGEBA.................. Di-glycidyl ether of bisphenol A DSC........................ Differential Scanning Calorimetry hrs........................... hours K ............................. Kinking M ............................ Matrix cracking MDA....................... Methylene Diamine min .......................... minute mPDA ..................... m-Phenylene Diamine PEEK...................... Poly(ether ether ketone) PES......................... Poly(ether sulfone) phr ........................... per hundred of resin rpm ......................... rounds per minute S.............................. Shear cracking SEM........................ Scanning Electron Microscopy STDEV................... Standard Deviation TEM ....................... Transmission Electron Microscopy TETA...................... Triethylene Tetramine TGAP ..................... Triglycidyl p-amino phenol TGDDM................. Tetra-glycidyl diamino diphenyl methane TGMDA ................. Tetraglycidyl methylene dianiline WAXD ................... Wide Angle X-Ray Diffraction XRD ....................... X-ray Diffraction xxxiv

Chapter 1 - Introduction

CHAPTER 1

INTRODUCTION

The development of high performance structur al materials is particularly attractive to the aerospace industry. Fibre polymer composites are one such class of materials that have become increasingly popular in this industry in view of their exceptional strength and stiffness to density ratios. A typical composite consists of strong stiff fibres, such as glass or carbon fibres, embedded in a tough resin matrix, such as epoxy resin. The fibres react the loads while the matrix maintains the fibre orientation, distributes the load, and protects the structure against environmental factors such as moisture and chemical attack. One of the first applications of modern composites in the aerospace industry was in the skins of the empennage of the F14 and F15 fighters, but the structural weight of the composites used was only around 2% [1]. With improved understanding of the material behaviour , the percentage of composites used in military aircraft increased rapidly rising to 19% by weight in the F18 and 24% in the F22. C omposites have also seen growing use in civil transport aircraft where reduced airframe weight can reduce fuel consumption, thereby lowering operating costs. A irbus in particular has made significant use of composite materials for commercial aircraft starting with the A300 and A310. The new A380, one of the largest commercial aircraft ever, contains around 30% composite material by weight. T he Boeing 777 is about 20% composite while the new generation Boeing 787 is 50% composite by weight and boasts high efficiency and performance with reduced weight. Many different parts of the aircraft are now made of

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Chapter 1 - Introduction

composites including wing skins, forward fuselage, flaperons, trailing-edge flaps, spoilers, airlerons, wheel doors, rear pressure bulkhead, and top and bottom skin panels. Effort is now focused on more extensive use of composites in primary airc raft structure such as wing and fuselage structure to further reduce weight and operating costs. This introduces new technical challenges. One of the critical areas is bolted composite joints. Size limitations in fabrication, in addition to economic factors, have necessitated that structur es are manufactured as subcomponents which are subsequently assembled to produce a final vehicle. Due to the inability of adhesive bonds to transfer high mechanical loads between structural components, mechanically fastened joints are often used in design. The joints, however, may be potential weak points in the structure as the bolt hole produces a stress concentration with the potential for crack initiation. The bolts impose a compressive load on the laminate , which can result in bearing failure by a combination of fibre buckling, shear and interlaminar splitting [2]. Since joints have such a critical effect on the safety and efficiency of the aircraft structure, it is vital that the most advanced design methods are used. The need to join composite panels through bolted joints has raised many issues not encountered with metallic joints and has been a major topic of research since the 1970s. For composite laminates, the joint efficiency is substantially lower than for structural metals, being reduced by almost 50% [3]. To avoid catastrophic failure, joints are designed to fail in bearing and the design strength is then limited by the bearing performance of the composite [4, 5]. Bearing results in compression loading of the hole indicating that improved bearing performance should be achieved by increasing the compression strength of the composite. While the stiffness and tensile strength of a composite are influenced principally by the reinforcement , the shear and compressive strength are depende nt more on the resin matrix. Rosen [6] , in a composite analysis, proposed that the compressive strength of a composite is proportional to the shear modulus of the matrix. By taking into account the matrix non-linearity and fibre misalignment, Sun and Jun [7] concluded that the compressive strength of a composite is proportional to the matrix elastic -plastic tangent -2-

Chapter 1 - Introduction

modulus and to the elastic-plastic shear modulus of the composite. This implies that improved bearing performance could be achieved by using a matrix resin with a higher modulus. However the conventional method for improving the modulus is to increase the level of cross linking which reduces toughness. As a result the increased elastic performance is offset by decreased resistance to matrix cracking. In recent years there has been intense interest in the development of polymer nanocomposites. These are a combination of a polymer matrix and nanoparticles having at least one dimension of less than 100 nm. A variety of polymers ha s been used as the matrix, including polyurethane [8, 9], polyimide [10, 11] , polyamide [12, 13] , polyester [14, 15] , and epoxy resin [16, 17]. Many nanometer-sized particles have been introduced into the matrix, including metal oxides, such as T iO 2 [18, 19] , SiO 2 [20, 21] , Al2O3 [22] and ZnO [23] and metallic powders, such as silver [24, 25] and gold [26, 27]. In particular, nanoclay particles have been widely used as the reinforcement in nanocomposites because of their special structure, high aspect ratio, high surface area, high strength and low cost [28-31]. Nanocomposites are promising materials for many industries and are being considered for automotive applications [32, 33], flame retardant materials [34, 35], electronics and electrical engineering [36-38] , packaging and containers for food [39]. When uniformly dispersed at a molecular scale, these nanometer-sized fillers can provide substantial improvements in mechanical properties [40] , thermal properties [35], and barrier performance [17, 41] even at very low volume fraction loadings (15%). These contrast with the high volume fraction loadings (50-70%) of micro-particles used in conventional composites. In particular, the introduction of nanoparticles into a polymer matrix is reported to simultaneously increase both stiffness and toughness providing a potential strategy for improving bearing performance. Additionally, because of the low loadings, there is only minimal change in weight. However, to date no studies have been conducted to examine whether bearing performance can be improved. Bearing fa ilure involves shear displacement beneath the hole resulting in outward lateral movement [7]. This suggests that an alternative strategy for improving bearing performance would be the use of through thickness reinforcement. 3D reinforcement

-3-

Chapter 1 - Introduction

can be produced by techniques such as weaving [42, 43], braiding [44, 45] , stitching [46, 47] , and knitting [48, 49] and its use in composite aircraft components has attracted significant attention since the late 1960s , particularly because of the potential for producing complex shaped structures [50]. However, the cost of manufacture, poor understanding of the in-plane properties and a lack of characterisation of the failure mechanisms have greatly limited the use of 3D fibre composites to date. An effective alternative to through-thickness stitching of composites in the z-direction is through-thickness z-pin reinforcement of traditional two dimension composites [51]. The z-pins are short fibres made of carbon or metal, inserted perpendicular to the stack of prepreg plies during the fabrication pr ocess prior to cur ing the laminate. The presence of z-pins improves the out-of-plane properties of the composite as they effectively link the individual layers together, preventing delamination crack growth. However, the high cost has limited their application. Only limited studies on the use of z-pins have been made and these have been mostly restricted to T-joints [52-54] or areas subject to impact damage [55-57]. However one study has been made very recently on the effect of z-pinning on bearing behaviour [58]. While the z-pins increased the load to failure , the bearing strength was not increased because of local thickening around the hole. None the less the technology looked promising. The present study examines both the use of nanoparticle matrix resin and the use of through thickness reinforcement using z-pinning to improve bearing strength of fibre reinforced composites. While the z-pins were available commercially, nanoreinforced matrix resin was not, and it was therefore necessary to develop the nanoreinforced resin as part of the work.

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Chapter 1 - Introduction

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A. Quilter, Composites in Aerospace Applications. 2006, ESDU International (An IHS White Paper): Canada. M. McCarthy, "BOJCAS: Bolted joints in composite aircraft structures", Air and Space Europe, 3 [3-4] (2001). L.J. Hart-Smith, "Joints", pp. 479-95 in Engineered Materials Handbook. Vol. 1: Composites. AMS International, Metal Park, OH, 1987. M.C.Y. Niu, Composite airframe structures: practical design information and data . Conmilit Press, Hong Kong, 1992. R. Li, D. Kelly, and A. Crosky, "Strength improvement by fibre ste ering around a pin loaded hole", Composite Structures, 57 [1-4] 377-383 (2002). B.W. Rosen, "Mechanics of composite strengthening", pp. in Fibre Composite Materials. ASM International, Metal Park, OH, 1965. C.T. Sun and A. Wanki Jun, "Compressive strength of unidirectional fiber composites with matrix non -linearity", Composites Science and Technology, 52 [4] 577-587 (1994). C.H. Dan, M.H. Lee, Y.D. Kim, B.H. Min, and J.H. Kim, "Effect of clay modifiers on the morphology and physical properties of thermoplastic polyurethane/clay nanocomposites", Polymer, 47 [19] 6718-6730 (2006). M. Song, H.S. Xia, K.J. Yao, and D.J. Hourston, "A study on phase morphology and surface properties of polyurethane/organoclay nanocomposite", European Polymer Journa l, 41 [2] 259-266 (2005). H.-S. Chen, C.-M. Chen, G.-Y. Chang, and S.-Y. Lee, "Study on nanodispersion of PI/clay nanocomposite by temporal analyses", Materials Chemistry and Physics, 96 [2-3] 244-252 (2006). H.-W. Wang, R.-X. Dong, H.-C. Chu, K.-C. Chang, and W.-C. Lee, "Improvements on the synthesis and properties of fluorinated polyimide-clay nanocomposites by using double-swelling agents", Materials Chemistry and Physics, 94 [1] 42-51 (2005). P. Winberg, M. Eldrup, N.J. Pedersen, M.A. van Es, and F.H.J. Maurer, "Free volume sizes in intercalated polyamide 6/clay nanocomposites ", Polymer, 46 [19] 8239-8249 (2005). B. Alexandre, S. Marais, D. Langevin, P. Mederic, and T. Aubry, "Nanocomposite-based polyamide 12/montmorillonite: relationships between structures and transport properties ", Desalination, 199 [1-3] 164-166 (2006). M. Zhang and R.P. Singh, "Mechanical reinforcement of unsaturated polyester by AL2O3 nanoparticles", Materials Letters, 58 [3-4] 408-412 (2004). H.N. Dhakal, Z.Y. Zhang, and M.O.W. Richardson, "Nanoindentation behaviour of layered silicate reinforced unsaturated polyester nanocomposites", Polymer Testing, 25 [6] 846-852 (2006). I. Isik, U. Yilmazer, and G. Bayram, "Impact modified epoxy/montmorillonite nanocomposites: synthesis and characterization", Polymer, 44 [20] 6371-6377 (2003). J.-M. Yeh, H.-Y. Huang, C.-L. Chen, W.-F. Su, and Y.-H. Yu, "Siloxanemodified epoxy resin -clay nanocomposite coatings with advanced anticorrosive properties prepared by a solution dispersion approach", Surface and Coatings Technology, 200 [8] 2753-2763 (2006).

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Chapter 1 - Introduction

[18]

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J. Adebahr, N. Byrne, M. Forsyth, D.R. MacFarlane, and P. Jacobsson, "Enhancement of ion dynamics in PMMA-based gels with addition of TiO2 nano-particles", Electrochimica Acta, 48 [14-16] 2099-2103 (2003). L. Zan, L. Tian, Z. Liu, and Z. Peng, "A new polystyrene-TiO2 nanocomposite film and its photocatalytic degradation", Applied Catalysis A: General, 264 [2] 237-242 (2004). Y. Zheng, Y. Zheng, and R. Ning, "Effects of nanoparticles SiO2 on the performance of nanocomposites", Materials Letters, 57 [19] 2940-2944 (2003). M. Sangermano, G. Malucelli, E. Amerio, A. Priola, E. Billi, and G. Rizza, "Photopolymerization of epoxy coatings containing silica nanoparticles", Progress in Organic Coatings, 54 [2] 134-138 (2005). B. Wetzel, F. Haupert, and M. Qiu Zhang, "Epoxy nanocomposites with high mechanical and tribological performance", Composites Science and Technology, 63 [14] 2055-2067 (2003). Y.-Q. Li, S.-Y. Fu, and Y.-W. Mai, "Preparation and characterization of transparent ZnO/epoxy nanocomposites with high -UV shielding efficiency", Polymer, 47 [6] 2127-2132 (2006). S.N. Abdullin, A.L. Stepanov, Y.N. Osin, and I.B. Khaibullin, "Kinetics of silver nanoparticle formation in a viscous-flow polymer ", Surface Science, 395 [2-3] L242-L245 (1998). M. Rong, M. Zhang, H. Liu, and H. Zeng, "Synthesis of silver nanoparticles and their self-organization behavior in epoxy resin ", Polymer, 40 [22] 6169-6178 (1999). M. Pumera, M. Aldavert, C. Mills, A. Merkoci, and S. Alegret, "Direct voltammetric determination of gold nanoparticles using graphite-epoxy composite electrode", Electrochimica Acta, 50 [18] 3702-3707 (2005). A.J. Vreugde nhil, K.K. Pilatzke, and J.M. Parnis, "Characterization of laser ablated gold nanoparticles encapsulated in epoxy amine crosslinked sol-gel materials", Journal of Non-Crystalline Solids, 352 [36-37] 3879-3886 (2006). J. Park and S.C. Jana, "Adverse effects of thermal dissociation of alkyl ammonium ions on nanoclay exfoliation in epoxy-clay systems", Polymer, 45 [22] 7673-7679 (2004). B. Qi, Q.X. Zhang, M. Bannister, and Y.W. Mai, "Investigation of the mechanical properties of DGEBA-based epoxy resin with nanoclay additives", Composite Structures, 75 [1-4] 514-519 (2006). M.-W. Ho, C.-K. Lam, K.-t. Lau, D.H.L. Ng, and D. Hui, "Mechanical properties of epoxy-based composites using nanoclays", Composite Structures, 75 [1-4] 415-421 (2006). L. Wang, K. Wang, L. Chen, Y. Zhang, and C. He, "Preparation, morphology and thermal/mechanical properties of epoxy/nanoclay composite", Composites Part A: Applied Science and Manufacturing, 37 [11] 1890-1896 (2006). H. Presting and U. Konig, "Future nanotechnology developments for automotive applications", Materials Science and Engineering: C, 23 [6-8] 737-741 (2003). J. Markarian, "Automotive and packaging offer growth opportunities for nanocomposites", Plastics, Additives and Compounding, 7 [6] 18-21 (2005). A. Gu and G. Liang, "Thermal degradation behaviour and kinetic analysis of epoxy/montmorillonite nanocomposites", Polymer Degradation and Stability, 80 [2] 383-391 (2003).

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[35]

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G. Camino, G. Tartaglione, A. Frache, C. Manferti, and G. Costa, "Thermal and combustion behaviour of layered silicate-epoxy nanocomposites", Polymer Degradation and Stability, 90 [2] 354-362 (2005). S.K. Bhattacharya and R.R. Tummala, "Epoxy Nanocomposite Capacitors for Application as MCM -L Compatible Intergral Passives", Journal of Electronic Packaging, 124 [1] 1-6 (2002). H. Koerner, D. Jacobs, D.W. Tomlin, J.D. Busbee, and R.D. Vaia, "Tuning Polymer Composite Morphology: AC Electric Field Manipulation of EpoxyMontmorillonite (Clay) Suspensions", Advanced Materials, 16 [4] 297-302 (2004). R. Gensler, P. Gröppel, V. Muhrer, and N. Müller, "Application of Nanoparticles in Polymers for Electronics and Electrical Engineering ", Particle & Particle Systems Characterization, 19 [5] 293-299 (2002). M. Avella, J.J. De Vlieger, M.E. Errico, S. Fischer, P. Vacca, and M.G. Volpe, "Biodegradable starch/clay nanocomposite films for food packaging applications", Food Chemistry, 93 [3] 467-474 (2005). P.C. LeBaron, Z. Wang, and T.J. Pinnavaia, "Polymer-layered silicate nanocomposites: an overview", Applied Clay Science, 15 [1-2] 11-29 (1999). W. Liu, S.V. Hoa, and M. Pugh, "Fracture toughness and water uptake of high performance epoxy/nanoclay nanocomposites", Composites Science and Technology, 65 [15-16] 2364-2373 (2005). J.A. Soden and B.J. Hill, "Conventional weaving of shaped preforms for engineering composites", Composites: Part A, 29 [7] 757-762 (1998). J. Quinn, R. McIlhagger, and A.T. McIlhagger, "A modified system for design and analysis of 3D woven preforms", Composites: Part A, 34 [6] 503-509 (2003). X. Sun and C. Sun, "Mechanical properties of three-dimensional braided composites", Composite Structures, 65 [3-4] 485-492 (2004). A. Miravete, J.M. Bielsa, A. Chiminelli, J. Cuartero, S. Serrano, N. Tolosana, and R.G. de Villoria, "3D mesomechanical analysis of three-axial braided composite materials ", Composites Science and Technology, 66 [15] 2954-2964 (2006). K. Dransfield, C. Baillie, and Y.-W. Mai, "Improving the delamination resistance of CFRP by stitching --a review", Composites Science and Technology, 50 [3] 305-317 (1994). D.C. Jegley, "Improving strength of postbuckled panels through stitching", Composite Structures, In Press, Corrected Proof (2006). K.H. Leong, S. Ramakrishna, Z.M. Huang, and G.A. Bibo, "The potential of knitting for engineering composites--a review ", Composites: Part A, 31 [3] 197220 (2000). M. Duhovic and D. Bhattacharyya, "Simulating the deformation mechanisms of knitted fabric composites ", Composites: Part A, 37 [11] 1897-1915 (2006). A.P. Mouritz, M.K. Bannister, P.J. Falzon, and K.H. Leong, "Review of applications for advanced three-dimensional fibre textile composites", Composites: Part A, 30 [12] 1445-1461 (1999). I.K. Partridge and D.D.R. Cartie, "Delamination resistant laminates by ZFiber(R) pinning: Part I manufacture and fracture performance", Composites: Part A, 36 [1] 55-64 (2005).

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[52]

[53]

[54] [55]

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[57]

[58]

G. Allegri and X. Zhang, "On the delamination and debond suppression in structural joints by Z-fibre pinning ", Composites: Part A, In Press, Corrected Proof (2006). D.D.R. Cartie, G. Dell'Anno, E. Poulin, and I.K. Partridge, "3D reinforcement of stiffener-to -skin T-joints by Z-pinning and tufting ", Engineering Fracture Mechanics, 73 [16] 2532-2540 (2006). K.L. Rugg, B.N. Cox, and R. Massabo, "Mixed mode delamination of polymer composite laminates reinforced through the thickness by z-fibers", Composites: Part A, 33 [2] 177-190 (2002). X. Zhang, L. Hounslow, and M. Grassi, "Improvement of low-velocity impact and compression -after-impact performance by z-fibre pinning ", Composites Science and Technology, 66 [15] 2785-2794 (2006). A.N. Palazotto, L.N.B. Gummadi, U.K. Vaidya, and E.J. Herup, "Low velocity impact damage characteristics of Z-fiber reinforced sandwich panels -- an experimental study", Composite Structures, 43 [4] 275-288 (1998). C. Scarponi, A.M. Perillo, L. Cutillo, and C. Foglio, "Advanced TTT composite materials for aeronautical purposes: Compression after impact (CAI) behaviour", Composites: Part B, In Press, Corrected Proof (2006). A. Crosky, D. Kelly, R. Li, X. Legrand, N. Huong, and R. Ujjin, "Improvement of bearing strength of laminated composites", Composite Structures, 76 [3] 260271 (2006).

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Chapter 2 - Literature Review

CHAPTER 2

LITERATURE REVIEW

2.1. INTRODUCTION This chapter reviews the literature pertinent to this thesis. It begins with a discussion of bearing behaviour of composites with a review of the mechanisms that have been proposed for bearing failure and an examination of the factors that affect bearing performance. The use of nanoparticle reinforced matrix resin was considered as a potential strategy for improving bearing performance and a detailed review of nanoparticle reinforced composites (nanocomposites) is then presented. A second potential strategy, namely through thickness reinforcement using z-pins, was also trialed in this study and a detailed review of this technology is also included. Finally, the goals of the project are presented.

2.2. BEARING PERFORMANCE 2.2.1. Introduction For over 30 years, there has been a steady increase in the use of composite materials , typically carbon and glass fibre reinforced plastics (CFRP and GFRP respectively), in the aerospace industry. Initially, composites were used only in non-structural parts and

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secondary structures. With further development and improved understanding of the materials, polymer composites have increasingly found use in primary aircraft structures, such as wings, and fuselages, Figure 2.1 [1]. Some of the potential benefits of using composite materials are substantial weight savings, a reduced number of joining operations during assembly, reduced inspection and reduced parts storage and movement, which collectively result in increased reliability and lower operating costs.

Fig. 2.1 Materials used in F/A 18 fighter aircraft (composites are shown in pink) [1]. However, realizing the full value of this potential still presents many technical challenges. One of them is mechanically fastened composite joints. Bolted connections are commonly used in preference to other joining techniques because they allow greater freedom in assembly and repair. An example is shown in Figure 2.2 [2]. The holes in the joints are however potential weak points in the structure. Additionally, introducing a bolted joint into a laminate introduces a new potential failure mode against which the laminate is relatively weak. The bolt imposes a compression load on the laminate and this can result in failure in bearing by a combination of fibre buckling, shear and interlaminar splitting.

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Chapter 2 - Literature Review

Fig. 2.2 Schematic of joints used in hybrid metal/composite wing [2]. The efficiency J of a single row fastened joint in a panel is a the ratio of the maximum bearing load PB that can be transferred by the bolt to the laminate and the ultimate tensile load PTU for the laminate remote from the hole. It is a function of the tensile strength FTU and bearing strength FB of the laminate, the width w and thickness t of the specimen, and the diameter d and local thickness tB of the fastener hole. The joint efficiency J is given by:

J=

FB dt PB PB = = PTU FTU wt FTU wt

(1)

For composite laminates, the joint efficiency is substantially lower than for structural metals [3], as shown in Figure 2.3. The lower efficiency is caused by many factors, such as brittleness, which leads to only minimal stress relief around the highest loaded holes, anisotropy, which leads to higher stress concentration factors, low transverse strength, susceptibility to delamination, and sensitivity to environmental conditions. Combined with these factors, the composite failure modes make the analysis and design of composites joints far more complex than that of metallic joints.

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0.8

Ductile metal

0.7

d

Brittle material

0.6

P B /FTU wt

Composite

PB

0.5 Tension failure

0.4 Bearing failure 0.3 0.2 0.1 0

w

0

0.2

0.4

0.6

0.8

1

d/w

Fig. 2.3 Joint efficiency of different materials (after Hart-Smith [3]).

Indeed, there are three different failure modes that can occur in bolted composite joints, namely net tension, shear-out and bearing failure, as shown in Figure 2. 4 [4]. Net tension and shear-out failures are total structural failures in which the joint components are separated. The net tension failure mode is characterized by fracture of the laminate across its width from the hole to the edges. The load required to fracture the laminate through a cross-section containing holes is less than for a section where there are no geometric irregularities [5, 6]. The shear-out failure mode is characterized by a “pull-out” fracture between the hole and laminate end, as a result of a shear stress concentration that the laminate is unable to adequately support [7]. The bearing failure mode is observed as a form of accumulated compressive damage in the laminate adjacent to the hole [8] and is much less catastrophic.

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Chapter 2 - Literature Review

Fig. 2.4 Bolted joint failure modes [4]. Geometric factors usually render one of the failure modes dominant. These factors include specimen width (w), edge distance (e), hole diameter (d), and thickness (t). For a given laminate, the net tension failure mode is associated with insufficient w/d ratio, where the specimen is not wide enough to prevent a tension failure. The shear-out failure mode is dependant on the e/d ratio; the edge distance can be made larger to prevent this type of failure. For sufficiently large e/d and w/d, the bearing failure mode will result. Various types of compression and bearing failures are illustrated in Figure 2. 5.

Fig. 2.5 Comparison of compression and bearing failure mechanisms [4]. - 13 -

Chapter 2 - Literature Review

An understanding of the mechanisms, as well as the factors affecting bearing failure of composites, is important in order to develop suitable strategies for improving bearing performance. 2.2.2. Failure mechanism The mechanism of bearing failure is very complicated, being a combination of several different fracture mechanisms occurring in the individual component materials and their interfaces, such as microbuckling, kinking, shear-cracking and delamination. Bearing failure in bolted composites is a process of damage accumulation, which is divided into four stages: damage onset, damage growth, local fracture and final structural fracture, as shown in Figure 2. 6 [9].

Fig. 2.6 Bearing damage process as detected by acoustic emission (AE) [9]. In a study by Xiao and Ishikawa [9], it was found that micro-damage, such as fibre microbuckling and matrix cracking started to occur when the applied load reached 50% of the maximum load. Only slight local delamina tion was detected in the epoxy composite laminate. When the load reached 60% of the failure load, the regions of micro-damage - 14 -

Chapter 2 - Literature Review

became more extensive around the hole because of the propagation of kink-bands and the formation of local delaminations. At 75% of the failure load, the extent of the damage became more pronounced as distinct through-thickness shear cracks expanded. In the final fracture, large-scale delaminations and shear cracks dominated. These results were similar to the previous findings of Ireman [10] in his study of the damage process, as illustrated in Figure 2. 7. Matrix cracks occurred in the resin rich surface layer at 25% of the failure load. The resin-rich layer chipped off at the hole edge. More severe damage was observed at approximately 35% of the failure load, where fibre fracture

Fig. 2.7 Schematic illustration of the failure processes [10].

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Chapter 2 - Literature Review

took place some distance from the hole edge, starting in the laminae where the fibres were loaded in an axial direction. At 70% of the failure load, delamination became the main mode of extending the damage area, while a band of accumula ted compression and shearing damage (kinking) was observed in the final failure [10]. In general, during bearing failure, composite materials undergo several different damage phenomena, including fibre microbuckling, kinking, shear cracking and delamination. 2.2.2.1. Fibre microbuckling Microbuckling is a deformation in which the fibres are considered to act individually as columns inside the matrix material [11]. It has been the most widely studied failure mode for composites made of strong fibres and matrices. For high volume fraction composites, microbuckling is expected to be controlled by the matrix stiffness in shear. Defects formed in the composite, such as fibre misalignment or waviness, porosity and residual stresses, can act as sites for microbuckling initiation [12]. Fibre microbuckling in the composite laminate is thought to start from elastic bending of the fibres and manifests itself in the fibres being pushed into a wave pattern causing subsequent stress on the surrounding matrix material [13]. The microbuckling problem was first analyzed in a systematic way by Rosen [14]. He proposed the existence of two possible microbuckling modes in composite laminates – the extension mode (symmetric) and the shear mode (asymmetric), as illustrated in Figure 2. 8. In the extension mode, the matrix material is predominantly in extension and adjacent layers deform out of phase with each other, while for the shear mode they deform in phase and the matrix is predominantly in shear. It was concluded that the nature of microbuckling depended on the distance of fibre separation and was associated with the level of interaction between the fibres in the composite laminate. From the view of microscopic analysis, adjacent fibres in the composite may buckle independently of one another or in a cooperative manner. Wave-like

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Chapter 2 - Literature Review

bending of adjacent fibres in different directions is associated with microbuckling. In this case, the transverse deformation of the fibres is out of phase relative to each other, with the result that deformation of the matrix takes place between adjacent fibres. This form of buckling is referred to as the extension mode and involves strain deformation in the matrix material.

Fig. 2.8 Initial configuration and buckling modes investigated by Rosen [14]. On the other hand, adjacent fibres may buckle cooperatively so as to be in phase with one another. As a result, the deformation of the matrix material between the fibres is primarily in the form of shear, and the mechanism is referred to as the shear mode. It is the more common of the two, as extension mode buckling only takes place when the inter-fibre distance is quite large , as in composites with a low fibre volume fraction. As a result the occurrence of shear mode buckling can be predicted simply from a knowledge of the fibre content in the composite.

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Chapter 2 - Literature Review

2.2.2.2. Kinking (or kink-banding) It is now widely recognized that kink formation and propagation is the dominant compression failure mechanism in unidirectionally fibre-reinforced composites. Kinking is a localized deformation band across the specimen in which the fibres have rotated by a large amount, and the matrix has undergone large shear deformation, Figure 2.9 [11, 15]. It is considered to be a final consequence of microbuckling, and is called a post-buckling event [16-20], although sometimes it is considered to be an independent failure mechanism [13, 21]. Kinking depends on the matrix stiffness and yield behavior in shear, and possibly on the fibre strength as well [14].

Fig. 2.9 In-plane buckling of fibres and fibre kink band geometry [13]. As fibre misalignment occurs in the loaded composite, shear stiffness is lost as a result of matrix yielding. This occurs at a fraction of the stress required to buckle the fibres in a perfectly aligned composite and localized fibre bending then takes place in narrow inclined zones. The final result of local fibre bending is fibre breakage to form distinct kink bands [13, 22, 23]. Garland et al. [24], and Narayanan [25] envisioned a similar failure mechanism in which a small, slant-aligned sequence of fibre breaks develops in shear or crushing failures and

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Chapter 2 - Literature Review

triggers kink-band formation through excessive overloading of adjacent fibres to the point of bending and failure, as illustrated in Figure 2.10.

Fig. 2.10 Schematic of fibre failure sequence in shear triggered kink-band formation [24]. Fibres outside the band above and below it are seen to be straight whereas inside it they are seen to be rotated. The fibres are clearly broken at the boundaries of the band while signs of fibre breaks are also seen inside the kink-band [26]. Once the kink-band forms in the material, the straight fibres adjacent to it tend to kink at a lower stress than was required to initiate the first kink-band. Contact of the straight fibres outside the band with the rotated fibres inside it tends to bend the straight fibres. Excessive bending results in breaking of a narrow strip of these fibres along the boundary. These then rotate further so as to conform with the deformation inside the original band, thereby increasing its width. The process is then repeated as shown in Figure 2. 11.

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Chapter 2 - Literature Review

Fig. 2.11 Geometry of kink-band formation [16]. 2.2.2.3. Shear cracking and delamination According to Wu and Sun [27], microbuckling in the 0° plies caused the formation of kinkbands and was considered as the primary failure mode for damage initiation at the bearing surface. It initiated in the outer-most 0° plies near the laminate surface at around 77.5% of the ultimate failure load. At higher loads, the damage propagated at an angle of between 30° and 40° into the interior of the specimen until it reached a critical length, as seen in Figure 2. 12. This propagation of the damage was referred to as interlaminar shear cracking.

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Chapter 2 - Literature Review

Fig. 2 .12 Micrograph of shear cracking in bearing failure [27]. Wang [28] also stated that local delamination would occur upon the formation of kinkbands. He identified two main mode s of damage to the laminate during bearing failure, as shown in Figure 2. 13.

Fig. 2.13 Shear cracking and delamination [28]. - 21 -

Chapter 2 - Literature Review

Shear cracking often appeared in the area closest to the bearing surface and was formed due to the accumulation of compressive failure in each individual ply of the laminate. Lateral expansion of the laminate was caused by the shear stress under the compressive loading, promoting propagation of the shear cracks. From a microscopic view, shear cracking was comprised of fibre kinking, fibre-matrix shearing and matrix cracking. The propagation of the shear cracks in the laminate ceased once converging cracks met, creating a larger scale unstable delamination [27, 28] , Figure 2.13. The formation of the delamination, or even laminar longitudinal splitting, occurred as a secondary damage mechanism. 2.2.3. Factors affecting bearing failure 2.2.3.1. Stacking sequence and percentage of 0° plies The mechanical properties of composite laminates are strongly dependent on the stacking sequence [29]. The laminates are usually made up of plies in the 0°, 90° and ± 45° directions. Park [30] stated that the stacking sequence had a great effect on delamination and the ultimate bearing strength of laminated composites. For both orthotropic and quasiisotropic laminates, the delamination bearing strength of a lay-up with 90° layers on the surface was stronger than that of a lay-up with 90° layers located at the centre of the laminate. He suggested that 90° plies played an important role in the delamination bearing strength of the laminates. This result is similar to that reported by Quinn and Mathew [31] in a previous study. They found that in a quasi-isotropic laminate, the bearing strength was improved by placing 90° plies at the surface and 0° plies towards the center. The constrainment of the 0° plies by the 90° plies in the outer layers resulted in a higher failure load being sustained by the laminate. According to Colling [5], bearing strength was significantly dependent on the ratio of 0° fibres to fibres in other directions. It increased with the addition of additional 0° plies due to the higher compressive strength in the load direction. However, there was a maximum

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Chapter 2 - Literature Review

percentage of 0° fibres that could be added. Once this percentage was exceeded, the failure mode changed, and a reduction in bearing strength occurred. Longitudinal splitting between the 0° fibres now took place because of the poor in-plane transverse restraint. In contrast, Smith and Pascoe [32] concluded that there was only a minor influence of stacking sequence on the bearing strength of laminates. Only the location of the 0° fibres in the laminate was reported to significantly affect the failure mode. Splitting and breaking away of the laminate occurred with 0° fibres on the surface while more general delamination between layers was observed with 0° fibres towards the interior of the laminate. The bearing strength was higher in the latter case. The elastic stiffness of the laminate in the loading direction and the fracture strain were independent of the stacking sequence of the laminate [33]. However the stacking sequence played a major role in determining the nature of the laminate fracture surface. 2.2.3.2. Laminate thickness An increase in the critical strain has been observed with decreased thickness [34]. More precisely, the transverse strain in the matrix became more important with reduced laminate thickness. Rosen [14] neglected the effect of transverse strain on microbuckling, but its effect should be included [34]. According to Wu and Sun [35], the bearing strength increased significantly with increased laminate thickness. As the thickness increased, kink-bands, which originated from the outmost 0° plies on either side of the laminate, propagated deeper into the interior of the laminate to merge together. Consequently, a higher stress, as well as higher load, was required for a thicker laminate. The effect of laminate thickness on the compression behaviour of composite laminates was investigated by Lee and Soutis [36]. They found that the strength of unidirectional laminates dropped by approximately 2-36 % in going from 2 to 8 mm thick specimens. But for open hole multidirectional specimens, the average strength increased with increasing

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Chapter 2 - Literature Review

specimen thickness, except for the 8 mm thick specimen ([45n /0n /-45n /90n] s) that failed prematurely due to extensive delamination introduced by matrix cracking, as seen in Figure 2.14.

Fig. 2.14 Open hole average compressive strength as a function of specimen thickness for multidirectional laminates [36]. 2.2.3.3. Hole machining defects The mechanical performance of bolted joints can be affected by the quality, as well as the accuracy, of the hole [37]. The quality of the hole is determined by the roughness of the surface and the damage in the area adjacent to the hole. Traditional methods for machining of holes include grinding and drilling. When drilling composites, it is difficult to achieve good hole quality. Hole machining defects which reduce the strength of the laminates, such as delamination, chip-out of fibres or matrix, and degradation of the matrix due to overheating, can easily occur. For example, if the cutting angle is too small and the cutting edge is wedged between two laminae, delamination may occur. However, it usually occurs when the last plies of the panel do not withstand the force exerted by the drill edge [38, 39]. During drilling, fibres and matrix can be torn out of the hole surface, resulting in a rough surface which can generate crack initiation sites. In addition, if the drilling thrust or torque forces are too high, the matrix may be degraded by overheating produced as a result of the friction between the tool and the laminate [37]. - 24 -

Chapter 2 - Literature Review

Increasing attention has been paid to drilling quality recently. For example, a drilling process with a one shot drill bit to get high quality holes has been examined by Fernandes [40]. The process was modeled as five steps, each directly related to the various drilling and reaming processes. 2.2.3.4. Lateral clamping Bearing strength is strongly influenced by lateral clamping which is produced by the pressure of the washer on both sides of the bolted region [41]. There are two mechanisms for explaining the increase in strength in a laterally constrained laminate. Firstly, for clamped bolted joints, failure occurs at the washer edge while it only happens at the pincontact area of an unconstrained laminate. Stress levels at the washer edge are relatively low compared to the hole edge, allowing a greater overall induced stress from bearing in the specimen. Secondly, the frictional force at the interface of the washer and laminate affect the load transfer across the laminate, influencing the distribution of the load over the larger area covered by the washer. The first mechanism often occurs in finger-tightened clamped specimens. The frictional forces in this case are relatively small. However, for a higher clamping torque, the friction between the clamp and the laminate is higher. Consequently, the bearing strength is higher and this is explained by the second mechanism. For clamped bolted joints, the clamping torque presses into the laminate under the washer [30]. Delamination on the loaded side of hole is therefore suppressed. The location of the delamination moves to the outer edge of the washer-constrained region. Clamping pressure of bolted joints suppresses the onset and propagation of laminate delamination, leading to a change in the failure mode. Consequently, the lateral clamping pressure increases both the delamination and ultimate failure strength of bolted joints in the laminate. With increased clamping pressure, a significant increase in ultimate bearing strength occurs while the delamination bearing strength shows a progressive increase.

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Chapter 2 - Literature Review

2.2.3.5. Matrix stiffness Wu and Sun [27] used Sun and Jun’s [42] microbuckling model to show that the critical stress in the direction of compression (σ xx ) cr is given by

(σ xx ) cr = Gmep /(1 − V f )

(2)

where G mep is the elastic-plastic shear modulus of the matrix and V f is the fibre volume fraction. This shows that the bearing strength is dependent on the shear modulus of the matrix indicating that the bearing strength should be increased by increasing the modulus. 2.2.4. Rationale for present study As noted in Section 2.2.1 and shown in Figure 2.3 the joint efficiency for bolted composite laminates is substantially lower than for metals. Reinforcement of composite laminates in the vicinity of the hole using layers of steered fibres has been used successfully as a strategy for improving bearing performance. Several different methods have been used to define the trajectories for the steered fibres, as described in References 43-47. However although substantial improvements were achieved, the incorporation of steered fibres into the laminate clearly complicates the production process. A review of the literature revealed two alternate strategies that might be effective in improving bearing performance and these were the focus of the work undertaken here. The findings of Wu and Sun [27] indicate that bearing strength should be improved by increasing the modulus of the matrix, as discussed in Section 2.2.3.5. However, increasing the modulus generally leads to a reduction in toughness for matrix resins but recent work has shown that simultaneous increases in both stiffness and toughness can be achieved in epoxy resin by incorporating nanoparticles into the resin [48]. Development of a nano particle reinforced matrix resin and its performance in a composite loaded in bearing was therefore the focus of the first part of the study. Accordingly, a detailed review of nanocomposites was undertaken and this is presented as Section 2.3. - 26 -

Chapter 2 - Literature Review

Additionally, it has been reported that bearing strength is also strongly affected by lateral constraint (clamping force) at the loaded hole [9, 30, 41] as discussed in Section 2.2.3.4. One benefit of lateral constraint is that it should oppose generation of the kink-bands since this involves outward lateral displacement of the material below the kink bands. This suggests that bearing performance may be improved by through thickness reinforcement, such as z-pins, and their use was examined as the second strategy trialed in this thesis. Zpin reinforcement is a relatively new technology and is therefore reviewed in detail in Section 2.4.

2.3. EPOXY NANOCOMPOSITES 2.3.1. Definition and Composition An epoxy nanocomposite is a composite material consisting of an epoxy resin matrix as the continuous phase, and nanometer-sized fillers, such as nanoparticles [49-55] and nanofibres [56-59], as the dispersed phase. The fillers must have at least one dimension in the order of nanometers (smaller than 100 nm) and can range from essentially isotropic elements to highly anisotropic needle-like or sheet-like elements [60]. 2.3.1.1. Epoxy resins Epoxy resins are used extensively in composite materials for a variety of demanding structural applications. The properties and structure of epoxy resins vary and strongly depend on the chemical structure of the resin and the curing agent, the presence of modifying agents and the curing conditions. All epoxy resins are characterized by the presence of a three-membered ring containing two carbon atoms and an oxygen atom. This is called an epoxy group or, less commonly, an epoxide or oxirane or ethoxylene ring: O C

R C

- 27 -

Chapter 2 - Literature Review

where R represents the point of attachment to the remainder of the resin molecule. The epoxide function is usually a 1,2- or a-epoxide that appears in the form: O H2C

CH

CH 2

called the glycidyl group, which is attached to the remainder of the molecule by an oxygen, nitrogen, or carboxyl linkage, hence, the terms glycidyl ether, glycidyl amine, or glycidyl ester [61]. In spite of the diversity and complexity of the available epoxies, the majority of resins are based on only three compounds: TGMDA (tetraglycidyl methylene dianiline), also called TGDDM (tetra-glycidyl 4,4’-diamino-diphenylmethane resin), DGEBA (diglycidyl ether of bisphenol A), and phenol formaldehyde novolac epoxy. The major difference between the molecules is that the cross-link density of cured TGDDM and novolacs is higher than that of DGEBA, resulting in higher values of Young’s modulus and glass transition temperature (Tg) but lower values of failure strain [62, 63]. TGDDM is the major component of the high performance resin formulations. These are amongst the stiffest available, and as a result, laminates based on TGDDM are amongst the highest performing. These resins are often used as matrices for carbon and aromatic fiber reinforced composites in the aerospace industry because they fulfill the requirements of high modulus and high temperature performance [64, 65]. Well-known trade names of these resins are Alradite MY 0510 (TGAP), Alradite MY 720 (TGDDM), Alradite XVMY 0505 (all supplied by Ciba). The structure of TGDDM is: O H2C

O CH

H2C

CH

H CH 2 CH 2

N

C H

O

CH 2

HC

CH 2

CH 2

HC

CH 2

N

O

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Chapter 2 - Literature Review

There are drawbacks to the high cross-link density of TGDDM resins. The most disadvantageous feature is that the resin failure strain is low, approximately 1.5%. This leads to the development of large delaminations upon impact, and consequently, low compression strength after impact. A second problem is water absorption, as every epoxy amine reaction results in a hydroxyl group. This leads to a reduction in the glass transition temperature (Tg). DGEBA is the most widely used resin, bearing common trade names such as EPON 828, EPON 826, EPON 825 (Shell), DER 332, DER 331, DER 323 (Dow), Alradite CY 225, CYD-128 (Ciba). Diglycidyl ether of bisphenol A made by reacting bisphenol A with epichlorohydrin is widely used in industry due to its fluidity, processing ease, and the desirable physical properties of the cured resin. The structure of the resin is as follows. O H2C

O

CH 3 CH

CH 2

O

O

C

CH 2

HC

CH 2

CH 3

Compared with TGDDM, bisphenol A epoxy cures to a lower cross-link density. Thus the resin modulus and Tg are lower, and consequently, mechanical properties and high temperature performance are also reduced. However, the cured bis A epoxy has a higher failure strain and a lower water absorption. Novolac epoxies are one of the most important classes of epoxy resins. They are products of the reaction of epichlorohydrin with various phenolic novolacs, cresol novolacs or bisphenol A novolacs. The structure of novolac epoxy is shown below. O CH 2

HC

O CH 2

CH 2

HC

O

O

O CH 2

CH 2 O

CH 2

CH 2

- 29 -

HC

CH 2

Chapter 2 - Literature Review

Novolac epoxies have a higher functionality, and thus cure to a higher cross-link density than DGEBA. The addition of novolac to a formulation increases the resin Tg, but decreases the failure strain. These resins are extensively used in prepreg formulations. Epoxy resin curing agents are divided into three technologically important classes: (a) active hydrogen compounds, which cure by polyaddition reactions; (b) ionic initiators, which are subdivided into anionic and cationic; and (c) crosslinkers, which couple through the hydroxyl functionality of higher-molecular-weight bisphenol A-type epoxy resins [66]. The most widely used of these are amine compounds, such as aliphatic amines, aromatic amines, polymercaptans , polyphenols, polybasic acids and acid anhydrides. Among them, the most common curing agents are the amines, in which each of the amino hydrogens reacts with an epoxide group. Depending on the number of amino hydrogens or structure of the curing agents, as well as their reactivity, a variety of properties of the material can be obtained. The aliphatic amines and their derivatives, such as diethylenetriamine (DETA), triethylenetetramine

(TETA),

are

recommended

for

ambient-temperature

curing.

Consequently, they would be expected to have a limited pot life or shelf life. They are often used in the form of an epoxy resin adduct, at a sacrifice of lower viscosity. The moderate temperature cured cycloaliphatic and tertiary/primary aliphatic amines offer a compromise between the room-temperature curing agents and the high-temperature curing aromatic amines. Aromatic

amines,

for

example,

m-phenylenediamine

(mPDA),

4,4’-

diaminodiphenylsulfone (DDS), are almost exclusively used in composite fabrication as these provide moderate viscosities at room temperature to the epoxy resin, long pot life, excellent chemical resistance, and good elevated-temperature performance [66]. Although less widely used than polyamines, acid anhydrides, such as nadic methyl anhydride, methylendic anhydride, hexahydrophthalic anhydride, and dodecenyl succinic anhydride, are the second most commonly used curing agents. They are preferred in some applications of polyamines due to long useful pot lives, low peak exotherms, and good

- 30 -

Chapter 2 - Literature Review

electrical properties. However, being hygroscopic, they cannot be exposed to the atmosphere for extended periods, and they require cures involving an added accelerator (for example, bezyldi-methylamine BDMA) at elevated temperature for optimum properties of the final material. Anionic initiators, such as tertiary amines (pyridine, benzyldimethylamine), secondary amines (piperidine, diethanolamine, imidazole derivatives), and metal alkoxides, are used as curing agents for epoxy resins, but only to a limited extent. Amongst the cationic initiators, boron trifluoride in the form of its monoethylamine complex is the most important member. While boron trifluoride gives an excessively rapid gelation, high hygroscopicity and light instability, its complex with amine overcomes these disadvantages [66]. A distinguishing feature of the chemistry of epoxy-curing is that the reaction involves the opening of the epoxy ring and the addition of molecules and reagents, symbolized as follows:

X: + Y

X:Y

where X: and Y are the electron-donating and the electron-accepting species, respectively. During curing, epoxy resins can undergo three basic reactions: (i)

Epoxy groups are rearranged and form direct lingkages between themselves.

(ii)

Aromatic and aliphatic hydroxyl group react with epoxy groups.

(iii)

Cross-linking occurs with the curing agent via various radical groups.

Under basic or neutral conditions, all ring-opening reactions are essentially similar and involve the attack of a nucleophile on one of the epoxide carbon atoms, as in the case of an amine curing agent. The proportion of epoxy resin to amine hardener is modified to achieve optimum cured properties based on the reactions given in Figure 2.15.

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Chapter 2 - Literature Review

OH

O CH

R

CH2

H ~C

+ H2N

R’

OH

N

C

R

CH

H

C

+ C

R’

OH

O

C

N

CH2

C’

~C

N

C

C

C~

C C

OH

C Schematic of curing reaction of epoxy/amine system.

Fig. 2.15

Since both the epoxy resin and the amine are polyfunctional, a highly cross-linked structure is created. In the curing reaction with anhydrides, there is a combination between the epoxy groups and the hydroxyl groups present in most epoxy resins, Figure 2.16. The reaction is initiated by the formation of the half ester through the anhydride joining the hydroxyl group. Selfpolymerization of the epoxy group through the epoxy-hydroxyl reaction may also occur [66, 67]. O ~~C

C

~~C

C

O

OH +

O

~~C~~

~~C

C

~~C

C

O

C

OH

O

O

O ~~C

C

O O OH

~~C

O

C~~ + ~~C

C O Fig. 2.16

~~C

C

O

C~~

~~C

C

O

C

C

O Schematic of curing reaction of epoxy/anhydride system.

- 32 -

C~~ OH

Chapter 2 - Literature Review

2.3.1.2. Nanoclay particles All nanoclay particles, also called phyllosilicates, contain articulated silicate or aluminosilicate layers in which sheets of tetrahedrally coordinated cations are linked through shared oxygens to sheets of cations which are octahedrally coordinated to oxygens and hydroxyls. When one octahedral sheet is linked to one tetrahedral sheet a 1:1 layer is formed; when one octahedral sheet is linked to two tetrahedral sheets, one on each side, a 2:1 layer is produced. There are six main groups of mineral clays distinguished according to structural type (Table 2. 1) Table 2.1 Classification and generalized structural formulae of phyllosilicates [68].

Mineral group

1. Kaolinite Serpentine

Nature of octahedral sheet(s)

Negative Within silicate layer charge per silicate Octahedral Tetrahedral cations cations layer

Between silicate layers Anions

Cations Hydroxide sheet cations

OH or H2O

Dioctahedral Trioctahedral

0 0

Y4 Y6

Z4 Z4

O 10 (OH)8 O 10 (OH)8

2. Pyrophyllite Dioctahedral Talc Trioctahedral

0 0

Y4 Y6

Z8 Z8

O 20 (OH)4 O 20 (OH)4

3. Micas

2 2 4 4

Y4 Y6 Y4 Y6

Z8 Z8 Z8 Z8

O 20 (OH)4 O 20 (OH)4 O 20 (OH)4 O 20 (OH)4

Y4 Y4 Y6

Z8 Z8 Z8

O 20 (OH)4 O 20 (OH)4 O 20 (OH)4

0.5-1.2 0.5-1.2 1.2-1.9 1.2-1.9

Y4 Y6 Y4 Y6

Z8 Z8 Z8 Z8

O 20 (OH)4 O 20 (OH)4 O 20 (OH)4 O 20 (OH)4

X 0.5 -1.2 X 0.5 -1.2 X 1.2 -1.9 X 1.2 -1.9

nH 2 O nH 2 O nH 2 O nH 2 O

-

Y4 Y8

Z8 Z12

O 20 (OH)2 (OH2 )4 O 30 (OH)4 (OH2 )4

X? X?

4H 2 O 8H 2 O

Brittle micas 4. Chlorite

5. Smectite

Dioctahedral Trioctahedral Dioctahedral Trioctahedral

Dioctahedral Variable Di,Trioctahedral Variable Trioctahedral Variable

Dioctahedral Trioctahedral Vermiculite Dioctahedral Trioctahedral

6. Palygorskite Sepiolite

- 33 -

nH 2 O

X2 X2 X2 X2 A4 A6 A6

(OH)12 (OH)12 (OH)12

Chapter 2 - Literature Review

Smectites have a crystal structure characterized by a 0.96 nm thick silicate layer consisting of two silica tetrahedral sheets fused to an edge -shared octahedral sheet of nominally alumina or magnesia. When three-valent aluminum ions are partially substituted by divalent magnesium ions (montmorillonite), or divalent magnesium ions are partially replaced by lithium ions (hectorite), the individual layers are anionically charged. The presence of gallery ions is responsible for the swelling of clay minerals in water. (The gallery is a term used to describe the stacked array of clay sheets separated by a regular spacing). Some subclasses of smectites used as fillers in nanocomposites are shown in Table 2.2. Table 2.2 Notional structure and chemistry of smectites [68].

Montmorillonite Hectorite Saponite Fluorohectorite Laponite

Type

Substitution

Dioctahedral Trioctahedral Trioctahedral Trioctahedral Trioctahedral

Octahedral Octahedral Tetrahedral Octahedral Octahedral

Unit Cell Formula (x ~ 0.67) M x[Al4-xM gx ]Si8 O 20 (OH)4 M x[Mg6-xLix]Si8 O 20 (OH)4 M y-x[Mg6-xFex3+][Si8-y Aly ]O 20 (OH)4 Li1.12[Mg4.88Li1.12][Si8 O 20 ]F4 Na0.7 [Mg5.4 Li0.4 ]Si8 O 20 (OH)4

Amongst them, montmorillonite, a crystalline, 2:1 layered clay mineral in which a central aluminium octahedral sheet is sandwiched between two silica tetrahedral sheets, is especially important in making polymer nanocomposites because a particular feature of the montmorillonite structure is that water and other polar molecules can enter between the unit layers due to the relatively weak forces between the layers. Its structure is given in Figure 2.17.

Fig. 2.17 Structure of montmorillonite. - 34 -

Chapter 2 - Literature Review

The ideal structural formula for montmorillonite is (Al3.15Mg0.85)(Si8.00)O 20 (OH)4 X0.85 nH2O, in which, X is a monovalent interlayer cation. The charge arises from divalent cations, usually Mg, in octahedral sites. The main attractive force is the electrostatic interaction of the negatively charged layers and positively charged interlayer material, frequently an alkali or alkaline earth cation. The size of this force depends on the surface charge density and also on the separation of the positive and negative charges, principally in the direction normal to the plane of the layers (dependent on layer separation) but also within the plane of the layers (dependent on distribution of substitutions in the layers). When the layers are in contact, or nearly so, there are Van der Waals attractive forces between atoms in the surfaces of adjacent layers and also residual electrostatic forces between charged, or partially charged, atoms in adjacent layers. The latter, like the overall electrostatic attraction between the layers and interlayer material, are dependent on layer separation and structure. The main repulsive force arises from the solvation of interlayer cations supplemented by that from interaction of the solvate (intercalate) with the surface oxygens of the 2:1 layers. In addition, there are repulsive forces between the hydrogens of the hydroxyl groups linked to the octahedrally coordinated cations in the 2:1 layers, and the interlayer cations. The magnitude of this repulsion depends on the proximity of the hydroxyl hydrogen to the seat of charge in the interlayer and is greatest when the O-H dipole is normal, or nearly so, to the plane of the 2:1 layer [68]. In general, the individual sheets in the silicates are stacked together and hydrophilic; thus they are not compatible with the hydrophobic organic polymer matrix.

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Chapter 2 - Literature Review

2.3.2. Structure of nanocomposites There are two possible types of epoxy/clay nanocomposites: intercalated and exfoliated. Intercalated nanocomposites are formed when one , or a few, molecular layers of epoxy resin are inserted into the clay galleries with fixed interlayer spacings (producing a corresponding small increase in the interlayer spacing of a few nanometers) and the expanded silicate layers are still in order. Exfoliated na nocomposites are formed when the silicate nanolayers are well separated from one another and individually dispersed in the continuous polymer matrix, with the average distance between the segregated layers being dependent on the clay loading. The separation between the exfoliated nanolayers may be uniform (regular) or variable (disordered) [69-74] as shown in Figures 2.18 and 2.19.

Fig. 2.18 Schematic illustrations of conventional composite and nanocomposites [48].

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Chapter 2 - Literature Review

Fig. 2.19 TEM images of intercalated (left) and exfoliated (right) nanocomposites [75]. The ideal nanocomposite structure is characterized by homogeneous dispersion of the inorganic platelets in a polymer matrix. Exfoliated nanocomposites have higher phase homogeneity than that of the intercalated counterpart, thus the exfoliated structure is the most desirable for improving mechanical strength as well as the other properties of the nanocomposites.

Fig. 2.20

Representative scattering curves for intercalated and exfoliated morphology for

epoxy/SC18 layered silicate nanocomposites at several stages of cure [76].

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Chapter 2 - Literature Review

The formation of intercalated or exfoliated nanocomposites can be observed using x-ray scattering, as shown in Figure 2.20, where the extent of separation of the clay layers is given by the d spacing. The structure produced in the nanocomposites leads to a significant enhancement in their properties compared to those of conventional composites. 2.3.3. Properties of nanocomposites Nanocomposites usually exhibit substantial improvement s in strength and modulus. The dispersion of individual nanosheets of the layered silicates in the epoxy resin and the interfacial coupling between the individual sheets and the polymer matrix facilitate stress transfer to the reinforcement phase, increasing the effective volume fraction of the reinforcement entities leading to improve d mechanical properties [71]. There is a strong correlation between the degree of dispersion of the silicate aggregates on a microscale and the degree of separation of the silicate layers on a nanoscale. In the context of fracture properties, the degree of microscale dispersion is considered an important factor contributing to the nanoscale structured properties [64]. This dependence of the properties of composites on such a wide range of variables means that although they have the potential for improved properties, sometimes these cannot be achieved as desired. 2.3.3.1. Tensile properties The tensile modulus of epoxy nanocomposites is substantially increased with increased nanoclay content [77-79], as seen in Figure 2.21. This significant improvement is achieved due to contributions from several different factors, such as the high modulus of the layered silicate, the orientation of the silicate layers with respect to the loading direction, the good dispersion of the nanosize clay particles and the good interfacial adhesion between the nanoclay and the epoxy matrix, creating a higher effective volume fraction of nanoclay reinforcement.

- 38 -

Chapter 2 - Literature Review

Fig. 2.21 Effect of nanoclay content on the modulus of nanocomposites [79]. Improvements in tensile strength have also been observed, with a 20 % increase in the ultimate tensile strength having been achieved with 6 wt.% nanoclay content [54, 78]. However, in some cases the tensile strength has been found to be reduced with increased nanoclay content [80, 81]. This has been attributed to the stress concentration effect of agglomerated clay nanoparticles [82]. With increased nanoclay concentration, the agglomerated particle size increases and inter-particle distance decreases as the nanoclay particles fill-in the available volume [83]. 2.3.3.2. Compressive properties The presence of nanoclay particles in the epoxy matrix improves both compressive strength and modulus of the epoxy resin nanocomposites in comparison to the pure resin, Figure 2.22 [48, 84]. The stiffness of the nanoparticles and the good interfacial interactions between the nanoclay layers and the epoxy matrix are responsible for this improvement.

- 39 -

Chapter 2 - Literature Review

Fig. 2.22

Compressive modulus of nanocomposites and filler composites with clay

loading [84]. 2.3.3.3. Impact properties Incorporation of nanoclays into epoxy resin improves the impact properties of the nanocomposites, as shown in Figure 2.23 [81, 85]. After exposure to mechanical stress, the nanoparticles are likely to orientate perpendicular to the applied load, creating nanovoids and initiating shear yielding of the epoxy interlayers at the tip of the propagating crack and also throughout the entire volume [85]. Thus the fracture surfaces of epoxy/clay nanocomposites show massive shear deformation, meaning that energy is absorbed for such shear yielding, leading to the increase in impact strength.

- 40 -

Chapter 2 - Literature Review

Fig. 2.23 Effect of nanoclay addition on impact strength of nanocomposites [85]. Inhomogeneous dispersion of the nanoclay at higher clay content can however reduc e the impact strength. Although the nanoparticles serve as binding agents, modifying the morphological structure of epoxy network and promoting cavitation at the particle-polymer boundaries, stress concentration sources that lead to a reduction in impact strength are produced when significant aggregates occur in the nanocomposite [86]. Therefore, to enhance the impact properties of nanocomposites, a uniform dispersion of the nanoparticles in the epoxy matrix resin is necessary. 2.3.3.4. Flexural properties The flexural properties have been found to vary depending on the type of epoxy resin, the curing agent and the nanoclay concentration. The flexural strength of nanocomposites shows an improvement in both the glassy and rubbery regions of the epoxy resin in comparison to the pure epoxy resin [87]. Generally, the rubbery state shows a greater improvement in flexural strength than the glassy state. This is expected to be caused by the extra reinforcement provided by nanosheet alignment in the rubbery region [71]. The

- 41 -

Chapter 2 - Literature Review

nanoparticles tend to occupy small hole defects formed from local shrinkage during curing of the epoxy resin and act as a bridge interconnecting more molecules. This result s in a reduced total free volume as well as an increase in the cross-link density. Consequently, the flexural strength of nanocomposites is higher than that of conventional composites as well as that of neat epoxy resin [88]. 2.3.3.5. Toughness Some workers have reported improved toughness with the addition of clay nanoparticles [84, 89] while others have reported the reverse [71, 83]. The improvement in toughness has been attrubuted to the presence of agglomerates of nanoclays in the network. The tactoids act as stress concentrators, causing matrix yielding thus improving toughness [80]. Wellbonded agglomerates impede the propagation of cracks existing in the network, temporarily pinning them [84]. The propagating cracks are assumed to be impeded by rigid, impenetrable and well-bonded particles. When a propagating crack encounters these obstacles, it becomes temporarily pinned and tends to bow out between the particles, forming secondary cracks. Consequently, the energy absorption is increased because of the creation of the new fracture surface and the formation of the new non-linear crack front. Crack tip blunting can also occur due to localized shear yielding and damage zones, such as debonding of the particle/matrix interface and fracture of particles. The decrease in toughness found by other workers [71, 83] was attributed to the size of the nanoparticles being too small to provide toughening through a crack-bridging mechanism and thus not effectively enhancing crack-trajectory tortuosity [90]. 2.3.3.6. Barrier performance The hindered diffusion pathways through the nanocomposite caused by the dispersion of the individual sheets of the layered silicate result in enhanced barrier properties of nanocomposites, such as resistance to the uptake of water and solvents [91, 92] , reduced gas permeability [93], reduction of thermal expansion and increased flame retardance [94-

- 42 -

Chapter 2 - Literature Review

97]. Platelet particles enhance the barrier performance of polymers according to a tortuous zigzag path model developed by Neilson, in which the platelets obstruct the passage of gases and other permeates through the matrix polymer (Figure 2. 24) [70, 98]. In fully exfoliated nanocomposites, individual clay platelets will have the highest aspect ratio and thus the highest barrier improvement is expected.

Pristine polymer Fig. 2.24

Nanocomposite

Illustration of Neilson’s tortuous path model for barrier enhancement of

nanocomposites [98]. 2.3.3.7. Other properties Due to the unique structure and good dispersion of nanoparticles, epoxy resin nanocomposites have some other noteworthy features. Exfoliated nanocomposites have good optical transparency since the nanometer-thickness sheets of the silicate are fully separated and dispersed at a molecular level in the epoxy matrix resin to a size smaller than the wavelength of visible light [71, 99]. Figure 2.25 shows the level of optical transparency of exfoliated and intercalated nanocomposites based on epoxy resin filled with 10-20 wt.% nanoparticles [75].

- 43 -

Chapter 2 - Literature Review

Fig. 2.25

Optical images of 5 mm thick plaques of nanocomposites based on EPON 828

cured by D400 containing different loadings of nanoparticles [75]. Additionally, it has been reported that the glass transition temperature can be increased by increasing the amount of organoclays [82, 100]. This may be due to the layered silicates, when dispersed and wetted relatively well in the epoxy resin, hindering the motion of molecules in the epoxy network at least in the vicinity of the silicate surface. However a reduction in Tg has also been observed. This has been attributed to effects, such as a decrease in the cross-link density of the epoxy network [64, 79], and the plasticizing effects of small free molecules present in the network [78, 88]. Nanocomposites are also lower in density than equivalent conventional composites [88]. Loadings are much smaller than the amount of micro-size particles used in equivalent conventional composites, thus the change in weight of the nanocomposite is minimal. For instance, the density of pure epoxy is 1.2 g/cc and is found to increase by only 3% with an addition of 10 wt.% of clay [101].

- 44 -

Chapter 2 - Literature Review

2.3.4. Synthesis There are essentially two different approaches to synthesizing epoxy resin nanocomposites: solution and in-situ polymerization. The Solution method: is where both the organo nanoparticles and the epoxy resin are dissolved in a polar organic solvent. With clay filler, the entropy gained by desorption of solvent molecules allows the polymer chains to diffuse between the clay layers, compensating for their decrease in conformational entropy. After evaporation of the solvent, an intercalated nanocomposite is obtained. The major disadvantage of this method is that a large amount of solvent is required [71, 75, 102-105]. The in-situ polymerization method: is similar to the solution method except that the role of the solvent is replaced by a polar monomer solution. First, the layered silicate is rendered organophilic by exchanging the inorganic cations placed between the layers, with organic surface modifiers. Then, the organosilicate is swollen in the monomer before adding the curing agent [65, 72, 79, 91, 98, 106]. Dispersion of the nanoclay particles throughout the epoxy resin plays a key role in the successful preparation of nanocomposites. Effective methods for the uniform dispersion have attracted much attention and can be classified into two main techniques, namely physical and chemical methods. 2.3.4.1. Physical methods Mechanical mixing This method is based on generating high-shear forces by using high-speed mixing equipment, such as the Speedmixer DAC 150FV (an orbital mixer) [107] and the Molteni Planimax high shear mixer [108] , to isolate and disperse the nanoparticles homogenously in the epoxy resin. Nanoclay particles must be dried at a temperature higher than room

- 45 -

Chapter 2 - Literature Review

temperature to aid movement and reduce electrostatic interaction between the particles. The epoxy resin should also be preheated to lower its viscosity as well as to enable better wetting of the particles. The particles are then mechanically mixed into the epoxy resin to achieve a uniform dispersion of nanoparticles in the matrix resin. This is done under vacuum to remove entrapped air that may significantly lower the mechanical properties [109]. Ultrasound sonication According to ultrasound theories, the densities and sound velocities of a particle and a medium determine the particle behavior in that medium. A particle undergoes a gravity (sedimentation) force as well as the acoustic force in a vertically formed ultrasound standing wave field. According to classical mechanics, the particles in an ultrasonic radiation - gravity coupled field move to an equilibrium position determined by the balance between these two forces [110]. Cavitation, formation and implosion of microbubbles in a high-intensity ultrasonic field propagate shock waves through the liquid. This intense energy accelerates both physical and chemical reactions, enhancing surface chemistry and causing violent particle motion, thereby generating high-velocity interparticle collisions. Ultrasound substantially reduces the particle size of ultrafine suspensions in one tenth the time of traditional ball milling methods. In addition, disaggregation and deagglomeration of clumps occur (particle size reduction), as does degassing of the carrier liquid, increasing slurry flow properties, and producing a higher homogeneity. Nanoparticles can be homogenously dispersed in epoxy resin under ultrasonic waves in an ultrasonic bath, such as Bransonic Model B220 [69, 106]. This can be followed by a high-speed homogenizer with a rotational speed of 24,000 rpm [111]. 2.3.4.2. Chemical methods The purpose of these methods is to improve the wetability of hydrophilic nanoclay particles in organophilic resin so that a homogenous mixture formed by uniform dispersion of nanoparticles in the epoxy matrix can be obtained.

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Chapter 2 - Literature Review

Surfactants To prevent agglomeration and stabilize the surface of nanoclay particles, repulsive interparticle forces are essential, and surfactants are used for this purpose [112]. Organic cation intercalation plays an important role in epoxy - clay nanocomposite formation by providing a hydrophobic environment for intergallery adsorption of the polymer precursor. Organic modifiers make the hydrophilic surfaces of inorganic nanoparticles hydrophobic, improving the interfacial properties between the organic and inorganic phases when the epoxy resin is added. Furthermore, the surface modifier reduces the surface energy of these inorganic particles in order to decrease the physical or electrostatic interactions between the silicate layers, favouring the diffusion of epoxy and curing agent molecules into the layer gallery. If there are no organic surfactants, the unmodified clay particles quickly precipitate out of solution and refuse to form a homogeneously dispersed nanocomposite [83]. Alkylammonium ions and related ions are the most commonly used for the preparation of clay-epoxy nanocomposites and provide two functions. The first is to make the intergallery regions organophilic so as to enable the intercalation of the matrix resin between silicate layers. The second function of the surfactant is to catalyze the epoxy polymerization reaction via the onium ions [107]. The Van der Waals galleries in clay layers contain charge-compensating cations such as alkali and/ or alkali earth cations and the interaction between the layers is relatively weak. Thus, the exchange of interlayer cations is relatively easy, making cation-exchange possible. The ammonium ions replace the original cations in the gallery of the clay. The basal spacing of a homoionic organoclay depends on the chain length of the alkylammonium ion and the charge density of the layered silicate. Exfoliated nanocomposites show greater phase homogeneity than intercalated ones. This structural distinction is the primary reason for the exfoliated state improving the performance properties of polymer-clay composites [73]. To improve results, dispersion of organic modified nanoclay particles (montmorrilonite) in the matrix resin is often aided by using stirring during mixing under pressure, followed by sonication and finally degassing in vacuum.

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Chapter 2 - Literature Review

Monomers Monomers with a low molecular weight can penetrate into agglomerated nanoparticles easily, reacting with activated sites inside and outside the agglomerates. The interstitial volume inside the nanoparticle agglomerates will be partly filled with grafted macromolecular chains, and the agglomerated nanoparticles will be separated further. Additionally, the surface of the nanoparticles will also become organophilic due to an increased hydrophobicity resulting from the grafted polymer. When the pre-grafted nanoparticles are mechanically mixed with epoxy resin, the former will keep their more stationary suspension state due to interaction between the grafted polymer and the matrix. After the mixture is cured, adhesion between the epoxy matrix and the nanometer-sized fillers will be substantially enhanced by chain entanglement and/or chemical bonding between the grafted polymer and matrix polymer [104, 113]. Reactive solvents Melt mixing of nanoclays with high performance epoxy resin is not feasible due to severe shear heating and formation of particle aggregates. One effective method has been to use low molecular weight reactive solvents as dispersing agents, such as acetone, chloroform, tetrahydrofuran, toluene, N,N-dimethylformamide [71, 75]. The role of the solvent is to reduce the viscosity of epoxy resin in mixing as well as to separate the nanoparticles. When using solvents to aid dispersion of nanoparticles, there was no observed change in curing reaction, morphology or final mechanical properties of nanocomposites [75]. The epoxy resin was fully dispersed in acetone by mechanical stirring and sonicating. Then, the desired amount of organoclay was added to the mixture under mechanical stirring and sonication before adding the curing agent. The system was held under vacuum to remove the solvent that had been used to help disperse the organoclay particles, while the mechanical stirring and sonication were maintained [71].

- 48 -

Chapter 2 - Literature Review

Modified matrix resin To obtain nanoscale dispersed epoxy composites, it is possible to modify the epoxy resin to be an interpenetrated network with another engineering plastic, such as aromatic polyester, which is used as a dispersing medium for the nanoparticles [49]. This modification, can also significantly affect the matrix properties. For example, cyanate esters can be used to improve the thermo-mechanical properties as well as the water absorption resistance of the epoxy nanocomposite [114], while the epoxy matrix toughness can be improved by using a hyperbranched epoxy resin [115]. 2.3.4.3. Curing procedures A number of phenomena occur during curing of the epoxy-clay system. Theoretical and experimental investigations have been carried out to determine the thermodynamics and kinetics of polymer melt intercalation. Most have focused on cross-linked polymer nanocomposites derived by polymerizing the intercalated monomer. The optimum preparation conditions are determined to promote the exfoliated nanostructure by investigating the initial and the final state of the clay layers [116]. A critical balance must be established between the relative rates of reagent intercalation, chain formation, and network cross-linking in order to achieve exfoliation before network formation and matrix gelation [73]. When the interfacial interaction between the polymer and layered organosilicate is strong enough, intercalated or exfoliated nanocomposites will be formed. As polymerization and network formation develop, the unit layers begin to expand slightly. It is possible that during the network formation, the layers are not uniform introducing slight losses in registry, contributing to general disorder [76]. Different curing speeds in bulk epoxy and at the silicate interfaces can account for internal stresses and substantial loss of mechanical properties and even formation of powdery materials [117]. Therefore, the balance of polymerization rate between the intergallery and the extragallery regions plays an important role in the exfoliation of clay layers in the

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epoxy resin. As long as the intergallery polymerization occurs at a rate comparable to the extragallery polymerization, an exfoliated nanocomposite structure will be formed. However if the extragallery polymerization is quicker than the intergallery diffusion and polymerization, or if intergallery polymerization is retarded, an intercalated structure will result [69, 107, 118]. According to Park [119], faster intergallery polymerization expedites exfoliation of the clay galleries but is not a requirement for exfoliation. A complete state of exfoliation of the clay is only achieved under specific conditions of curing times and curing temperatures. Kong and Park [116] investigated the intercalation behavior of DGEBA epoxy resin in the intergallery of clay modified by C18 amine was investigated. They showed that at 50°C epoxy resin has sufficient thermal energy and mobility to diffuse into the intergallery of the clay. This may be due to a sudden increase in the interlayer distance resulting from rapid reorientation of the C18 amine molecule causing displacement of the silicate layers, as shown schematically in Figure 2.26.

Fig. 2.26 Schematic diagram showing the exfoliation of the clay layers in DGEBA resin during mixing [116]. At the first stage of curing (ie the mixing state), the epoxy resin ha s a linear structure rather than a 3-D network structure since the self-polymerization of epoxy resin by protonated amine is faster than the curing reaction with unprotonated amines. This is, however, limited due to the confined geometry of the intergallery region.

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Figure 2. 27 indicates that the clay layers are exfoliated stepwise rather than gradually. The unique S-shaped plot of the interlayer distance against degree of conversion is due to the distinct three part behavior. The first (0 - 30% conversion) is where the epoxy resin is cured by the catalytic effect of protonated amine to form a linear polymer in the intergallery area rather than a cross-linked network. During this period, the interlayer space is expanded by the epoxy resin growing parallel to the clay layer.

Fig. 2.27 D-spacing of clay layers as function of degree of conversion at isothermal curing temperatures of 120, 130, 140°C for DGEBA-DDS-C18 clay (5 phr) [116]. The second is associated with the cross-linking reaction between the epoxy resin and the curing agent continuing in the intergallery region. The protonated alkylammonium cations are already consumed, and the cross-linking reaction of the epoxy resin increases the interlayer distance further. Temperature strongly affects the degree of curing in epoxy resin, for example at 130 and 140°C, cross-linking from the epoxy-amine reaction is enhanced by the catalytic effect of the hydroxyl groups produced from self-polymerization and from the amine, whereas at 120°C cross-linking of the epoxy-amine is too slow, producing an exfoliated clay layer with an interlayer distance no greater than 20 nm.

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In the third part, when conversion is 50% or more completed, the rate of the epoxy resin reactions in the intergallery decreases and the reactions in the extragallery regions increase since the viscosity of the mixture increases. The interlayer distance increases slowly until it is saturated. Therefore, to obtain a well-exfoliated nanostructure, cross-linking in the intergallery region must be promoted during the curing period. Additionally, in a study by Park et al [119] , it was found that the elastic forces developed in the clay galleries during epoxy curing were responsible for the exfoliation of the clay structure (Figure 2.28).

Fig. 2.28

Schematic illustration of the intercalated state and exfoliation process showing

the forces acting on a pair of clay layers: (a) organically modified clay, (b) epoxyintercalated state, (c) forces acting on a two-particle tactoid [119]. When the elastic force overcame the attractive force between the quaternary ammonium ions and the negatively charged clay particles and the viscous force increased with curing of the epoxy resin, exfoliation of the clay occurred; otherwise, intercalated clay structures, with cross-linked epoxy molecules in them, were retained in the cured system.

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Furthermore, the ionic bonding energy between the clay layers gradually increased from the surface to the center of the clay particles, meaning that the inner layers had a higher ionic bonding energy than the surface layers (Figure 2.29). Thus, the edge layers could be separated more easily than the inner layers and the exfoliation process began with separation of the surface layers away from the tactoids. Fully exfoliated nanocomposites were obtained when all the layers were individually separated.

Fig. 2.29

Schematic diagram showing the relationship between the ionic bonding energy

and the location of the layers in the tactoid: (a) tactoid, (b) variation of bonding energy along the thickness of the tactoid [119]. 2.3.5. Factors affecting nanocomposite structure Depending on the layered silicate charge density, the nature of the interlayer ion exchange , the curing conditions, the curing agent, and the epoxy resin, the resultant epoxy nanocomoposite will show either an intercalated or exfoliated morphology [65].

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2.3.5.1. Effect of nanoclay The amount of nanoclay particles introduced into the matrix resin has a large influe nce on the structure, as well as the properties, of the final nanocomposites. When the loading of clay is lower, the intergallery spacing of the nanosheets is increased [69], most likely due to more epoxy resin entering between the nanosheets [71]. The low content of the organoclay in the epoxy matrix also makes the distance between the neighboring organoclay clusters generally greater. When the nanoclay gallery expands, interaction between these clusters should be much smaller, and the expansion of the gallery can be continued more freely [120]. In almost all reported studies, the concentration of nanoparticles in the epoxy resin varied over a relatively small range, usually from 1 to 20 wt.% , with the loadings being predominatly less than 10 wt.%, although loadings up to 70 wt.% have been examined [102]. In the synthesis of epoxy-clay nanocomposites, the cation-exchange capacity (CEC) controls the amount of surfactant ions, which can be intercalated between the layers [121]. Lan et al. [122] reported that the CEC of the clay determine d the content of alkylammonium ions present between the clay layers and thus controlled the space available for diffusion of epoxy molecules during mixing of the organoclays with the epoxy resin. When the charge density was increased, the interlayer spacing decreased and achieved its minimum value at the highest CEC [91]. 2.3.5.2. Effect of surfactants The structure of organic surfactants has a significant influence on the nanoclay layer expansion behavior as well as the intercalation capability of epoxy resins and curing agents in the intergallery layer of clays. Increasing the alkyl chain length of alkylammonium cations increases the interlayer distances of the organoclays and the corresponding epoxy composites [117, 120].

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Lan and coworkers [122, 123] found that clays with primary and secondary onium ions formed exfoliated nanocomposites, whereas those with tertiary and quaternary onium ions retained an intercalated structure. Generally, varying the chain length of alkylammonium ions facilitates diffusion of the epoxy resin and curing agent molecules between the silicate layers because it reduces the electrostatic interactions present between them. According to Kornmann et al [64], the acidity of the alkylammonium ions used for catalytic epoxy homopolymerization may have a negative effect on the properties of nanocomposites. The homopolymerization of epoxy can substantially change the structure of the polymer network. The major problem with the use of alkylammonium ions for synthesis of epoxy-layered silicate nanocomposites is their relatively poor compatibility with the polar epoxy matrix because of the non-polar nature of their alkyl chain. Therefore, all amines are only partially protonated before being used to modify the surface of clay layers, so as to retain some unprotonated amino reactive groups which can then react with the epoxy network during nanocomposite synthesis [77]. It may be of interest to find alternative ways of catalyzing the intergallery polymerization. 2.3.5.3. Effect of matrix resin Different epoxy resins have different effects on the curing behavior, as well as the morphology, of the nanocomposites due to their mobility and reactivity, their structure and their functionality. In Becker et al’s study [65], three different epoxy resins were used to investigate the influence of the matrix on the structure of the nanocomposites, and the extent of exfoliation of the clay layers in the epoxy resins was examined. The resins were conventional epoxy (DEGBA) and high performance epoxies (TGDDM and TGAP). The DEGBA produced a well-intercalated or exfoliated nanocomposite structure easily at a variety of temperatures (80-160°C), whereas the other two resins only for med intercalated nanocomposites. When these two higher functionality epoxy resin systems were compared, TGDDM, in spite of its significantly higher viscosity, experienced better exfoliation than the less viscous TGAP system. Both systems showed improved exfoliation at higher

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temperature. This was attributed to increased catalytic effect between the epoxy resins and organic modifiers, as well as the increased mobility of the macromolecular chains. In preparing adducts, epoxy resin is often treated with prot onated amines in a stoichiometric ratio. The epoxy containing amine molecules is more compatibile with the surface modifier and consequently improved separation of silicate layers results [77]. 2.3.5.4. Effect of hardeners Various hardeners are used to cure epoxy resin for nanocomposites, including aliphatic polyamines, aromatic diamines, and anhydrides, such as 3,3’-dimethylmethylenedi (cyclohexylamine) [124] , methyltetrahydrophenol anhydride [72], diethylene triamine [125], m-phenylenediamine (mPDA) [123], 4,4’ diaminodiphenyl sulfone (DDS) [64], and imidazole [126]. Depending on the reactive level of the curing agent, a nanocomposite will have either a simple intercalated, partially exfoliated or sufficiently exfoliated structure.

20 nm

Fig. 2.30

20 nm

Transmission electron micrographs of 10 phr C18 clay reinforced epoxy

nanocomposites cured by (a) MDA (4,4’ -methylene -dianiline) and (b) DDS (4,4’diaminodiphenyl sulfone) [116].

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With low electronegativity amines, polymerization occurs slowly enough to provide a low viscosity medium to the nanoclay layers, therefore, layered silicates can be exfoliated before gelation of the epoxy resin [116]. From TEM studies, Kong and Park [116] found that low reactivity and highly flexible curing agents, such as DDS, promote an exfoliated nanostructure, Figure 2.30. Slow curing of the epoxy resin delays extragallery gelation and reduces the extragallery polymerization rate. This provides enough time for intergallery polymerization. These findings are similar to results obtained by Kornmann et al. [124]. They investigated the cure kinetics of epoxy systems cured by three hardeners and concluded that the lower the reactivity of the curing agent, the higher the degree of exfoliation. Therefore, the flexibility (i.e. molecular mobility) and the reactivity of the curing agent are important parameters, which influence the balance between the extragallery and the intergallery polymerization rates.

Fig. 2.31

Small angle X-ray scattering patterns of 1% montmorillonite containing

DGEBA epoxy cured by MPDA with (a) 25 phr; (b) 14.5 phr; (c) 5 phr [69]. montmorillonites, indicating the absence of exfoliation. For the stoichiometrically equivalent mixture, the interlayer spacing was 40.5Å (b), which showed that exfoliation was suppressed by the reaction involving the curing agent. When the curing agent concentration was below the stoichiometric amount, the layer spacing increased

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dramatically to approximately 180Å (c). This meant that the individual layers of montmorillonite were now well separated in a regular manner and were well exfoliated. From the above it is seen that the lower the curing agent concentration, the higher the level of exfoliation of the clay. An excess amount of curing agent promotes extragallery crosslinking of DGEBA and MPDA (curing agent) with insufficient diffusion of both epoxy and hardener molecules into the intergallery regions. The rapid rate of extragallery crosslinking, in comparison to the slower intergallery diffusion, limits the exfoliation [69]. Alternatively, the higher concentration of curing agent in the intergalleries speeds up the curing reaction of the hardener/epoxy system occurring in these intergalleries. Therefore, good penetrating ability of the curing agent into the clay is also an important factor for exfoliation of nanoclays [127]. 2.3.5.5. Effect of temperature Curing epoxy resin is strongly temperature dependent. Determination of the optimum temperature for curing is crucial as it directly affects the extent of cross-linking in the matrix network, and the level of exfoliation of the nanoclay. If the temperature is not high enough, the movement of the epoxy and curing molecules is very slow, and there will be only a few organic molecules of matrix resin and curing agent in the gallery of the clay. The epoxy network has too low an extent of conversion to exfoliate the clay layers and only intercalated or partial exfoliated nanocomposites are formed. On the other hand, if the temperature is too high, the curing reaction occurs too quickly. There is not enough time for molecules of epoxy resin and curing agent to penetrate into the galleries of the nanoclay. The final composite is either a conventional composite or an intercalated nanocomposite. The best temperature for curing provides sufficient thermal energy to the matrix resin and curing agent to intercalate easily into the clay layers while giving a the balance between polymerization between the intergallery and extragallery of the clay layers [116]. Lan et al. [123] demonstrated that exfoliation of the clay was not only dependent on the reactivity of the epoxy system but also on the rate of intercalation of epoxy and curing

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agent. The term “rate of intercalation” corresponds to the rate of diffusion of the organic molecules between the clay layers. If this rate is too slow, the extragallery polymerization rate will be faster than the intergallery polymerization rate and only intercalated nanocomposites will form. Higher curing temperatures increase the reactivity of the epoxy systems and the diffusion rate of the epoxy and the curing agent between the layers, favouring the intergallery cure kinetics. This leads to exfoliation of the nanoclay. This trend was confirmed in Tolle and Anderson’s study [76]. The exfoliation of the nanocomposite was completed in shorter times at higher isothermal cure temperatures. 2.3.5.6. Effect of time The curing time has a strong effect on the extent of conversion, as well as the extent of exfoliation of the clay layers. In Figure 2.32, for isothermal curing at 120°C, the initial diffraction peak, which represents the intercalation of epoxy molecules into the clay layers, broadens until it disappears with time, whereas the peak representing exfoliation of the clay layers develops with time [116].

Fig. 2.32 Isothermal time-resolved SXRD patterns of nanocomposites at 120°C with scan time (minute) from bottom to top as follow: 0, 5, 10, 11, 15, 20, 25, 30, 40, 60 and 89 min [116]. - 59 -

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In another study, the extent of exfoliation of clay layers was found to increase with time at constant temperature, from 35Å to 120Å after 195 minutes (Figure 2.33).

Fig. 2.33

Time-dependent small angle synchrotron X-ray scattering patterns of an

epoxy/montmorillonite mixture at 1350C [69]. 2.3.5.7. Effect of pressure Pressure can strongly affect the dispersion of nanoparticles as well as the extent of exfoliation of the layered clay when preparing nanocomposites. Under different pressures the interdistance between the nanoparticles has been reported to vary, while the time for curing was reduced at higher temperatures since the interaction between the epoxy and curing agent molecules was increased. Pressure also influenced the diffusion of macromolecules between layers. However, there is little literature available on this matter, the limited work being at 13 mbar [108, 117] , 100 psi [104, 128] and 15,000 psi [102]. The effect of pressure may warrant more detailed examination.

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2.4. Z-PIN REINFORCEMENT 2.4.1. Concepts Laminated fibre reinforced polymer composites have be come extremely popular as structural materials because of their high in-plane strength to weight ratios. One drawback, however, to the use of laminated composites is that they have poor through-thickness strength, making interlaminar delamination a major concern in the application of composites. Several methods have been examined for improving the through-thickness properties of laminated composites, such as the use of stitching, toughened resin matrices and through-thickness pin reinforcement. Stitching involves the sew ing of fibres through laminates to improve interlarminar strength and delamination resistance, since the bridging action of the through-thickness threads suppresses delamination crack growth by exerting a high crack closure at the interface between the delaminated layers [129-131]. However, the in-plane tensile, compressive and flexural strengths are usually reduced due to damage of the fibres during the stitching process while the stitches also act as stress concentration sites under loading [132, 133]. In toughened resin systems, the formulation of the resin is changed in order to increase the fracture toughness so as to improve delamination resistance. The most common way to increase the fracture toughness is to incorporate a rubber into the epoxy resin, such as carboxyl terminated butadiene acrylonitrile [134, 135], amine terminated butadiene acrylonitrile and epoxy terminated butadiene acrylonitrile [136, 137]. However, the rubbertoughened highly crosslinked epoxy systems have low moduli and decreased thermal stability. In order to get improved toughness without losing thermal and mechanical properties, engineering thermoplastics, such as poly(ether sulfone) PES [138], hydroxyl terminated poly(ether ether ketone) PEEK [139], polypropylene and nylon [140] have been used as an alternative toughening agent to rubber. However, this method is costly.

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The most recently developed method of improving through thickness strength is the use of through-thickness pin reinforcements, known as “z-fibres” or “z-pins”. The name signifies that the reinforcing pins are parallel to the ‘z-axis’ of the analytical plate models [141]. Zpins are usually made from structural materials such as pultruded carbon fibre/epoxy or glass/epoxy, titanium alloy, and stainless steel. The range of pin diameters is from 0.15 to 1 mm. The z-pins are supplied in blocks, known as “preforms”, consisting of structural foam containing the reinforcing fibers. The preform is characterized by the material of the z-pins and by the areal density of the pins, i.e., the ratio between the combined surface area of cross sections of the z-pins and the total cross-sectional area of the preform [142-144]. The aerial densities range from of 0.5-10%. Examples are shown in Figure 2.34.

Fig. 2.34 Z-pin performs containing 0.28 mm diameter pins at densities of 0.5, 2 and 4% [144]. The pins are inserted into the plane of the prepreg plies during the manufacturing process, effectively pinning the individual layers together prior to curing. With a pin content of 0.5 5 vol.% the z-pinned composite laminates have been reported to exhibit high delamination resistance [145, 146], increased compression-after-impact strength (up to 45%) [147] , a significant improvement in the shear strength (up to 40%), and increased fatigue strength (up to 40%) [148].

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Other studies, however, have reported that the presence of through-thickness pins decreased the tensile and compression strengths of the laminates [143, 149] , and caused a steady decline in the flexural strength and fatigue life [150]. 2.4.2. Mechanism In z-pinned composite laminates, the z-fibres bridge delamination cracks, reducing the driving force for crack growth. The bridging tractions arises from the debonding and frictional forces at the interface of the z-pin with the matrix and from the work necessary to bend the z-pins under shear load and force them through the surrounding material [141, 145]. The pins suppress the formation of delamination cracks through the specimen, even at the lowest pin content (0.5%) [148]. The effect of z-pin reinforcement in crack bridging is illustrated in Figure 2.35.

Fig. 2.35 Micrographs showing z-pin reinforcement [145]. In some cases, the presence of z-pins results in a change in the failure mechanism from delamination to microbuckling of the laminate. Note that the failure mode also depends on the material and geometrical properties of the laminate plies, especially their thickness and waviness as well as the resin flow stress. Competing failure mechanisms include delamination initiation, the arrest of delamination cracks after growth past bridging z-

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fibres, the conversion of the crack tip field to mode II, and compressive failure (microbuckling) under bending loads. In through-thickness pin reinforcement, resin rich pockets are created around the sides of the z-pins, filling the space where the laminate fibres are pushed apart. The extent of these pockets depends upon the z-pin diameter and pin-to-pin spacing. During fracture under mode I loading, the pins pull out of the resin envelope and the major energy absoprtion mechanism is the friction associated with this pull-out. In the mode II loading case, the zpins undergo significant bending deformation prior to their failure by internal shear [142].

Fig. 2.36 Typical morphology of local area around a z-pin [143, 151].

Fig. 2.37

Sketch and SEM micrograph showing fibres deflecting around z-pins and

weaving through a field of z-pins [143].

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Steeves and Fleck [143] concluded that the compressive strength of z - pinned laminates was reduced by the misalignment of the fibres adjacent to the z-pins (Figure 2.36). The worst fibre misalignments are caused when fibres weave through a field of z-pins. As a consequence, the compressive strength of the z-pinned composites is inversely related to the orientation angle. The weaving effect dominates the compressive strength of z-pinned laminates Figure 2.37. Out-of-plane crimp of the tows is observed in the vicinity of the z-pins [148]. Crimping is caused by the pin tips pressing into the tows as they penetrate the prepreg stack and by friction dragging fibres along the sides of the z-pins. Lateral distortion of the fibres also occurs around the pins, causing the formation of resin-rich regions. Sweeting and Thomson [151] determined that the thermal residual stresses predicted by modeling were higher than the failure stress of the resin. As a result failure should occur in the region around the z-pin (Figure 2.38). Evidence of cracking around the z-pins was obtained experimentally. The cracking occurred both inside and outside the z-pin, but rarely at the interface, indicating that the resin in, or surrounding the z-pin, failed rather than the interface between the z-pin and the laminate. The good interfacial strength is most likely due to mechanical interlocking and friction, rather than a physical bond.

Fig. 2.38 Crack around z-pin in laminated composite [151].

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Cracks initiated in the damaged regions surrounding the pins were also observed by Chang and his colleagues [150] , as seen in Figure 2.39. They proposed that the cracks could have initiated in the clusters of broken fibres near the pins, and then spread radially towards similar cracks propagating from neighbouring pins before finally coalescing into a single through-width crack that caused final failure.

Fig. 2.39 Scanning electron micrograph showing a crack that initiated near a pin under flexural loading [150]. They also expected that the loss in flexural strength could be minimized by staggering the location of the pins in a zig-zag pattern, rather than having the pins located in straight, parallel rows because it would increase the spacing between pins across the laminate. Damage can also be diminished by reduction of the pin diameter, and it has been proposed that reinforcement with ultra-fine pins with diameters in the nanometer to micron-size range would cause less damage than the existing sub-millimeter size pins [150]. Under mixed-mode loading, the pins exhibit several different mechanisms: (i) debonding from the laminate and axial sliding against an interfacial friction traction; (ii) lateral deflection into the laminate against a nearly uniform resisting force per unit length of deflected segment of the pin, which arises from irreversible deformation of the laminate; - 66 -

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(iii) enhancement of interfacial friction (snubbing effect); (iv) stable pull-out up to a critical load, which, as a result of snubbing, can be substantially higher in mixed-mode conditions than the load required for pull-out in mode I; (v) ultimate failure by either unstable pull-out or shear failure of the pin, the latter possibly occurring prior to the end of the stable pull-out phase, but not after it, since the load is then decreasing [148]. 2.4.3. Important factors 2.4.3.1. Pin density Increasing pin density will increase crack bridges, strongly suppressing crack propagation [144, 145]. For a higher density of pins, delamination cracks grow in a stable manner and the laminate can withstand a higher failure load before the crack runs into a pin and is artificially stopped. Consequently, the load carrying capacity of the laminate is considerably increased with a higher pin density, up to four fold by a 2 % increase of z-pin density [142]. In an investigation of T-joints, Allegri and Zhang [152] found that the efficiency of z-pinning increased with higher density by delaying the flange-to-base debond growth, with up to 10% increase in ultimate strength and failure displacement. In Chang et al’s study [148] , the ultimate shear strength and elongation limit were greatly improved with increased pin loading, with improvements of 41% and 56% being achieved respectively at 2% pin density. The effective stiffness of the bridging mechanism was expected to rise proportionately with pin density. Stiffer bridging tractions tend to favour stable crack propagation and the possibility of arrest after some growth. However, a decrease of flexural strength has been reported with increased pin content, falling approximately 6% for every 1% additional density of pins [150]. It was suggested that swelling and micro-structural damage caused by the pinning process could be significant factors in the strength degradation.

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2.4.3.2. Pin diameter The maximum load per pin which the specimen can carry is higher for larger diameter pins [144]. Increasing the pin diameter is essentially equivalent to raising the frictional sliding shear between the pin and the matrix, as insertion of a thicker pin provides a larger displacement for the surrounding laminate, which will react by applying a larger residual stress on the pin [152]. However a significant loss of in-plane mechanical properties may occur due to the larger local misalignment of the in-plane laminate. Thus larger Z-pin diameters can have a detrimental effect on the in-plane mechanical properties. At the same density of pins, the static tensile strength of joints was considerably higher with laminates pinned by smaller pins (0.28 mm) than by larger pins (0.51 mm) [148]. Varying the pin diameter induced a transition in the failure mechanism of the pinned joints. At a density of 2%, the thinner pins arrested bond-line cracks and the joints failed in the laminate adherend. In contrast, for thicker pins with the same aerial density, the delamination crack propagated along the entire bond-line, as seen in Figure 2.40, and the mechanism of ultimate failure was a combination of complete pin pull-out and pin pull-out interrupted by pin shear failure.

(a)

(b)

Fig. 2.40 Delamination of laminate pinned with (a) 0.28 mm and (b) 0.51 mm pins [148]. Increasing the pin diameter also probably results in a greater number of broken fibres at each pin location [150]. The force needed to drive the pins through the laminate increases

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with thicker pins, creating a higher stress exerted by the pin tip onto the fibres during the pinning process. 2.4.3.3. Other factors Regardless of the pinning effect, there remain other factors affecting the bearing properties of pinned laminates, such as stacking sequence, laminate thickness, hole machining, lateral clamping, and matrix stiffening as discussed in Section 2.2.3. 2.4.4. Pinning process The original typical z-fibre reinforcing process relied on a combination of heat and pressure to compact the preform and push in the z-pins during the autoclave cure cycle [141] , as illustrated in Figure 2.41.

Fig. 2.41 Schematic of the z-fibre reinforcement process [141]. In this method, a release film, a z-fibre preform and a rigid tool are placed onto a laid-up prepreg laminate. During the cure in a standard autoclave, the heat softens the preform which collapses under the applied pressure, driving the z-pins into the laminate. On removal from the autoclave, the residual foam is removed and discarded. The process is completed with the removal of any pin material that projects above the laminate. Large areas of composite laminates can be reinforced in one autoclave cycle using this method but

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it is limited to using very low pinning density performs (below 0.5%) because of the high pressure required to push the pins in. At the present time, the insertion of z-pins into laminates is a two stage process, and involves the use of a specialized ultrasonic insertion gun and sequential removal of the collapsible foam sandwich in which the z-pins are held [142]. The preform is located on the top of an uncured laminate, in the exact area to be reinforced. A layer of Teflon coated glass fabric is placed between the laminate and the preform. An ultrasonic horn is then used to insert the z-pins into the laminate by vibrating at 20 kHz. After insertion, the compacted foam is removed and the excess pin length is cut flush with the surface of the laminate and the laminate is then autoclave cured. A schematic of the process is shown in Figure 2.42.

Fig. 2.42 Schematic of ultrasonically assisted z-fibre insertion [148].

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2.5. CHARACTERIZATION 2.5.1. Structure characterization 2.5.1.1. Wide Angle X-ray Diffraction (WAXD) WAXD analysis is commonly used to probe the structure of the clay nanocomposites due to its ease and availability. By monitoring the position, shape and intensity of the basal reflections from the distributed silicate layers, the nanocomposite structure may be identified (intercalated or exfoliated morphology).

Fig. 2.43 Schematic depicting the expected X-ray diffraction patterns for various types of hybrid structures [153].

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The d-spacing can be determined using Bragg’s Law, as given by equation (3) [81]:

nλ = 2d sin θ

(3)

where n is an integer, ? is the wavelength, ? is the glancing angle of incidence and d is the interplanar spacing of the crystal. The diffractograms produced by different types of the hybrid composite are illustrated in Figure 2.43. For the intercalated morphology, there is a visible peak around 2 degrees as the d-spacing of the clay is 35 - 40 Å. When the layers are more disordered, a smaller and broader peak is seen instead. There is no peak observed for exfoliated nanocomposites. Peak broadening and intensity decreases are very difficult to study systematically. Thus, conclusions concerning the mechanism of the formation and structure of the nanocomposites based on WAXD patterns are only tentative [154]. 2.5.1.2. Transmission Electron Microscopy (TEM) TEM is a very powerful technique, allowing a qualitative understanding of the internal structure, spatial distribution of the constituent phases and the defect structure, through direct visualization [154]. However, special care must be exercised to guarantee a representative cross-section of the specimen. The basic principle is the same as light transmission microscopy, in which the beam passes through the section and is affected by interacting with it. Typical TEM images of clay nanocomposites at low and high magnification are shown in Figure 2.44. The distribution of the nanoclay particles at microscale can be seen at low magnification while visualization of the nanoscale layer separation is clearly shown at high magnification.

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Fig. 2.44

TEM images of clay nanocomposite at low (left) and high (right) magnifications

[155]. 2.5.1.3. Scanning Electron Microscopy (SEM) A finely focused electron beam scanned across the surface of the sample generates secondary electrons, backscattered electrons, and characteristic X-rays. These signals are collected by detectors to form images of the sample displayed on a cathode ray tube screen. Features seen in the SEM image may then be immediately analyzed for elemental composition using energy dispersive spectroscopy EDS. Secondary electron imaging shows the topography of surface features in nanoscale or microscale. Materials are viewed at useful magnifications up to 100,000 times without the need for extensive sample preparation and without damaging the sample.

Fig. 2.45 SEM fractograph of 7 wt.% clay nanocomposite [101]. - 73 -

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The use of SEM in characterization of nanocomposites is concentrated on visualization of the distribution of nanoclay particles in the matrix as well as examination of the fracture topography of the materials. Figure 2.45 shows the fracture surface of an epoxy-clay nanocomposite. The bright spots correspond to clay aggregates finely dispersed in the epoxy matrix. Additionally, the fibres and the matrix of polymer composite laminates can be readily seen using SEM [156, 157]. Thus it is a useful tool for investigation of both nanocomposites and fibre reinforced composites. 2.5.1.4. Optical microscopy This method can be used for imaging nanocomposites at the microscale and for examining fracture surfaces for features such as separation of the nanoparticles from the matrix [84, 88]. It can also be used for examining crack propagation in composite laminates [158]. An optical micrograph is shown in Figure 2.46 of a DGEBA system containing 5 wt.% organoclay. Large particulates can be seen indicating that even for the DGEBA-based system not all the clay particles are fully dispersed in the polymer phase [88].

Fig. 2.46 Optical micrograph showing large particles in DETDA cured DGEBA system containing 5 wt.% clay [88].

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2.5.2. Thermal property anal ysis 2.5.2.1. Differentional Scanning Calorimetry (DSC) DSC measures the temperatures and heat flow associated with transitions in materials as a function of time or temperature, under controlled temperature conditions. DSC is an effective method for investigating the curing characteristics of epoxy/hardner/nanoparticle systems as well as for determining physical parameters of nanocomposites, such as glass transition temperature (Tg) [159]. Figure 2.47 shows the onset temperature of curing and the temperature of the exothermal heat peak for two nanocomposites and also for the neat resin Epon 862/W. Both the onset temperature for curing and the temperature of the exothermal heat peak (121 and 159°C, respectively in the neat resin) can be seen to be shifted to lower temperatures in the nanocomposites (80 and 142°C for 3% SC8/Epon 862/W, and 100 and 152°C for 3% SC18/Epon 862/W). This is caused by the catalytic effect from the acidic RNH3+ group in the gallery of the organoclay [120].

Fig. 2.47 DSC of Epon 826/W, 3% SC8/Epon 826/W, and 3% SC18/Epon 862/W at 2°C/min [120].

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2.5.3. Mechanical Behaviour 2.5.3.1. Compression Testing A method for the determination of the mecha nical properties of un-reinforced plastics and fibre reinforced high-modulus plastic composites is given in ASTM D695 [160]. The standard test specimen can be in the form of either a right cylinder or prism whose length is twice its principle width or diameter. A typical stress-strain curve is illustrated in Figure 2.48. The compressive modulus is calculated as the tangent to the initial linear portion of the load deformation curve. The compressive strength is determined by dividing the maximum compressive load carried by specimen during the test by the original minimum crosssectional area of the specimen.

Fig. 2.48 Typical stress - strain curve for a compression test.

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2.5.3.2. Bearing Test Bearing tests are used to determine the bearing response of polymer composite laminates. A test method is described in ASTM D5961/D5961M [161]. A typical load - deformation curve obtained from a bearing test is shown in Figure 2. 49. The bearing strength FB of the laminate is calculated by: FB = P/td Where FB = bearing strength (Pa) P

= bearing load (N)

d

=

t

= specimen thickness (m)

bearing hole diameter (m)

Fig. 2.49 Typical stress - strain curve for a bearing test [162].

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(4)

Chapter 2 - Literature Review

On the curve, the yield strength corresponds to a 2% hole elongation in the composite laminate, and the ultimate load represents the maximum load sustained by the joint prior to failure. The fixture assembly for the procedure is shown in Figure 2. 50.

Fig. 2.50 Fixture assembly for a bearing test [157]. In the test, the load is applied to the specimen by means of a double-shear clevis, using a pin. The pin can be torqued to allow a transverse compressive load to be applied to the coupon around the periphery of the hole. 2.5.3.3. Pin-contact Bearing Test As an alternative to the standard bearing test described above , Wu and Sun [27] have developed a simple bearing test, in which the load is applied through a hardened steel pin to

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a vertically clamped specimen with just the lower half of the hole at is upper edge , as illustrated in Figure 2.51

Fig. 2.51 Schematic diagram of pin-contact test [27]. This simple method allows the location and progression of contact damage to be observed more easily than in the conventiona l bearing test for bolted joints. A typical load displacement curve for a cross-ply laminate is shown in Figure 2.52.

Fig. 2.52 Typical load - displacement curve for cross-ply laminate [27].

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2.6. AIMS AND OBJECTIVES OF THESIS The aim of this investigation was to examine potential methods for improving the bearing strength of laminated fibre reinforced plastic composites. Two different strategies were evaluated, nanoparticle reinforcement of the matrix resin and through thickness reinforcement using z-pins. Objective 1: Nanocomposite preparation i)

Nanoclay modification with a series of surfactants was carried out to provide an understanding of the effect of surfactants on the nanoclay particles. The results were used to determine the most suitable modified nanoclay for reinforcing epoxy resin.

ii)

Nanocom posites based on low performance DGEBA epoxy resin reinforced with nanoclay were prepared. The work was undertaken to provide background knowledge which was subsequently used to prepare nanoclay composites from high performance TGDDM epoxy resin

iii)

TGDDM/nanoclay composites were synthesised and the key factors controlling the quality of the epoxy nanocomposites were identified.

Objective 2: Laminated nanocomposite investigation i)

Carbon fibre laminates were prepared using the nanoclay reinforced TGGDM epoxy resin as the matrix.

ii)

The nature of the nanocomposite matrix in the carbon fibre laminates was examined.

iii)

The effect of nanoclay reinforcement of the matrix on bearing performance was investigated. - 80 -

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Objective 3: Z-pin reinforcement investigation i)

An improved method of inserting z-pins into carbon fibre laminates was developed which minimized local thickening of the laminates and minimised the level of defects, such as voids, fibre waviness and resin rich areas.

ii)

The effect of z-pins on the bearing behaviour of laminated composites was evaluated.

2.7. SCOPE The experimental work is presented as three separate chapters. The preparation and analysis of epoxy nanocomposited reinforced with nanoclay is described in Chapter 3. The use of the epoxy nanocomposites as the matrix for carbon fibre laminates is examined in Chapter 4. Chapter 5 presents the study of through thickness reinforcement of carbon fibre laminates using z-pins. The conclusions are then summarized in Chapter 6.

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Nanocomposites with a Reduced Organic Modifier Content", Chemistry of Materials, 14 [10] 4088-4095 (2002). S.C. Jana and S. Jain, "Dispersion of nanofillers in high performance polymers using reactive solvents as processing aids", Polymer, 42 [16] 6897-6905 (2001). M.-M. Shen, M.-G. Lu, Y.-L. Chen, and C.-Y. Ha, "Nanocomposites based on liquid -crystalline epoxy-clay: synthesis and morphology", Polymer International, 54 [8] 1163-1168 (2005). P.B. Messersmith and E.P. Giannelis, "Synthesis and Characterization of Layered Silicate-Epoxy Nanocomposites", Chemistry of Materials, 6 [10] 1719-1725 (1994). J.S. Chen, C.K. Ober, M.D. Poliks, Y. Zhang, U. Weisner, and E. Giannelis, "Study of the interlayer expansion mechanism and thermal-mechanical properties of surface-initiated epoxy nanocomposites", Polymer, 43 [18] 4895-4904 (2002). C. Zilg, R. Mühlhaupt, and J. Finter, "Morphology and toughness/stiffness balance of nanocomposites based upon anhydride-cured epoxy resins and layered silicates", Macromolecular Chemstry and Physics., 200 [3] 661-670 (1999). B. Wetzel, F. Haupert, K. Friedrich, M.Q. Zhang, and M.Z. Rong, "Impact and Wear Resistance of Polymer Nanocomposites at Low Filler Content", Polymer Engineering and Science, 42 [9] 1919-1927 (2002). T. Masudo and T. Okada, "Ultrasonic Radiation - Novel Principle for Microparticle Separation", Analytical Sciences, 17 1341-1344 (2001). Y. Zheng, Y. Zheng, and R. Ning, "Effects of nanoparticles SiO2 on the performance of nanocomposites", Materials Letters, 4268 1-5 (2002). H. Ishida, S. Campbell, and J. Blackwell, "General Approach to Nanocomposite Preparation", Chemistry of Materials, 12 [5] 1260-1267 (2000). M.Q. Zhang, M.Z. Rong, S.L.Y.B. Wetzel, and K. Friedrich, "Effect of particle surface treatment on the tribological performance of epoxy based nanocomposites", Wear, 253 [9-10] 1086-1093 (2002). K. Dinakaran and M. Alagar, "Preparation and characterisation epoxy-cyanate ester interpenetrating network matrices/organoclay nanocomposites", Polymer for Advanced Technologies, 14 [8] 574-585 (2003). D. Ratna, O. Becker, R. Krishnamurthy, G.P. Simon, and R.J. Varley, "Nanocomposites based on a combination of epoxy resin, hyperbranched epoxy and a layered silicate", Polymer, 44 [24] 7449-7457 (2003). D. Kong and C.E. Park, "Real Time Exfoliation Behavior of Clay Layers in EpoxyClay Nanocomposites", Chemistry of Materials, 15 [2] 419-424 (2003). C. Zilg, R. Thomann, J. Finter, and R. Mülhaupt, "The influence of silicate modification and compatibilizers on mechanical properties and morphology of anhydride-cured epoxy nanocomposites", Macromolecular Materials and Engineering, 280-281 [1] 41-46 (2000). Q. Wang, C. Song, and W. Lin, "Study of the exfoliation process of epoxy-clay nanocomposites by different curing agents", Journal of Applied Polymer Science, 90 [2] 511-517 (2003). J.H. Park and S.C. Jana, "Mechanism of Exfoliation of Nanoclay Particles in EpoxyClay Nanocomposites", Macromolecules, 36 [8] 2758-2768 (2003).

- 88 -

Chapter 2 - Literature Review

[120] C. Chen and D. Curliss, "Preparation, Characterisation, and Nanostructural Evolution of Epoxy Nanocomposites", Journal of Applied Polymer Science, 90 [8] 2276-2287 (2003). [121] X. Kornmann, H. Lindberg, and L.A. Berglund, "Synthesis of epoxy-clay nanocomposites: influence of the nature of the clay on structure", Polymer, 42 [4] 1303-1310 (2001). [122] T. Lan, P.D. Kaviratna, and T.J. Pinnavaia, "Epoxy self-polymerization in smectite clays", Journal of Physics and Chemistry of Solids, 57 [6-8] 1005-1010 (1996). [123] T. Lan, P.D. Kaviratna, and T.J. Pinnavaia, "Mechanism of Clay Tactoid Exfoliation in Epoxy-Clay Nanocomposites", Chemistry of Materials, 7 [11] 2144-2150 (1995). [124] X. Kornmann, H. Lindberg, and L.A. Berglund, "Synthesis of epoxy-clay nanocomposites. Influence of the nature of the curing agent on structure", Polymer, 42 [10] 4493-4499 (2001). [125] W. Xu, S. Bao, S. Shen, W. Wang, G. Hang, and P. He, "Differential scanning calorimetric study on the curing behavior of epoxy resin/diethylenetriamine/organic montmorillonite nanocomposite", Journal of Polymer Science: Part B, 41 [4] 378386 (2003). [126] W.-B. Xu, S.-P. Bao, S.-J. Shen, G.-P. Hang, and P.-S. He, "Curing Kinetics of Epoxy Resin-Imidazole-Organic Motmorillonite Nanocomposites Determined by Differential Scanning Calorimetry", Journal of Applied Polymer Science, 88 [13] 2932-2941 (2003). [127] L. Jiankun, K. Yucai, Q. Zongneng, and Y. Xiao-Su, "Study on intercalation and exfoliation behavior of organoclays in epoxy resin", Journal of Polymer Science: Part B, 39 [1] 115-120 (2001). [128] B.P. Rice, C. Chen, L. Cloos, and D. Curliss, "Carbon Fiber Composites: Organoclay-Aerospace Epoxy Nanocomposites, Part I", SAMPLE Journal, 37 [5] 7-9 (2001). [129] K. Dransfield, C. Baillie, and Y.-W. Mai, "Improving the delamination resistance of CFRP by stitching--a review", Composites Science and Technology, 50 [3] 305-317 (1994). [130] R. Velmurugan and S. Solaimurugan, "Improvements in Mode I interlaminar fracture toughness and in-plane mechanical properties of stitched glass/polyester composites", Composites Science and Technology, 67 [1] 61-69 (2007). [131] H.-J. Chun, H.-W. Kim, and J.-H. Byun, "Effects of through-the-thickness stitches on the elastic behavior of multi-axial warp knit fabric composites", Composite Structures, 74 [4] 484-494 (2006). [132] M.Z.S. Khan and A.P. Mouritz, "Fatigue behaviour of stitched GRP laminates", Composites Science and Technology, 56 [6] 695-701 (1996). [133] A.P. Mouritz, K.H. Leong, and I. Herszberg, "A review of the effect of stitching on the in-plane mechanical properties of fibre-reinforced polymer composites", Composites: Part A, 28 [12] 979-991 (1997). [134] M.L. Arias, P.M. Frontini, and R.J.J. Williams, "Analysis of the damage zone around the crack tip for two rubber-modified epoxy matrices exhibiting different toughenability", Polymer, 44 [5] 1537-1546 (2003). [135] R. Thomas, J. Abraham, S.T. P, and S. Thomas, "Influence of carboxyl-terminated (butadiene-co-acrylonitrile) loading on the mechanical and thermal properties of - 89 -

Chapter 2 - Literature Review

[136]

[137]

[138]

[139]

[140] [141] [142]

[143]

[144]

[145]

[146] [147]

[148]

[149] [150] [151]

cured epoxy blends", Journal of Polymer Science: Part B, 42 [13] 2531-2544 (2004). J.-F. Hwang, J.A.M.R.W. Hertzberg, G.A. Miller, and L.H. Sperling, "Structureproperty relationships in rubber-toughened epoxies", Polymer Engineering and Science, 29 [20] 1466-1476 (1989). E. Butta, G. Levita, A. Marchetti, and A. Lazzeri, "Morphology and mechanical properties of amine-terminated butadiene-acrylonitrile/epoxy blends", Polymer Engineering & Science, 26 [1] 63-73 (1986). I. Blanco, G. Cicala, M. Costa, and A. Recca, "Development of an epoxy system characterized by low water absorption and high thermomechanical performances", Journal of Applied Polymer Science, 100 [6] 4880-4887 (2006). B. Francis, G.V. Poel, F. Posada, G. Groeninckx, V. Lakshmana Rao, R. Ramaswamy, and S. Thomas, "Cure kinetics and morphology of blends of epoxy resin with poly (ether ether ketone) containing pendant tertiary butyl groups", Polymer, 44 [13] 3687-3699 (2003). P.J. Hogg, "Toughening of thermosetting composites with thermoplastic fibres", Materials Science and Engineering: A, 412 [1-2] 97-103 (2005). D.J. Barrett, "The mechanics of z-fiber reinforcement", Composite Structures, 36 [12] 23-32 (1996). I.K. Partridge and D.D.R. Cartie, "Delamination resistant laminates by Z-Fiber(R) pinning: Part I manufacture and fracture performance", Composites: Part A, 36 [1] 55-64 (2005). C.A. Steeves and N.A. Fleck, "In-plane properties of composite laminates with through-thickness pin reinforcement", International Journal of Solids and Structures, 43 [10] 3197-3212 (2006). D.D.R. Cartie, M. Troulis, and I.K. Partridge, "Delamination of Z-pinned carbon fibre reinforced laminates", Composites Science and Technology, 66 [6] 855-861 (2006). K.L. Rugg, B.N. Cox, and R. Massabo, "Mixed mode delamination of polymer composite laminates reinforced through the thickness by z-fibers", Composites: Part A, 33 [2] 177-190 (2002). W. Yan, H.-Y. Liu, and Y.-W. Mai, "Mode II delamination toughness of z-pinned laminates", Composites Science and Technology, 64 [13-14] 1937-1945 (2004). X. Zhang, L. Hounslow, and M. Grassi, "Improvement of low-velocity impact and compression-after-impact performance by z-fibre pinning", Composites Science and Technology, 66 [15] 2785-2794 (2006). P. Chang, A.P. Mouritz, and B.N. Cox, "Properties and failure mechanisms of pinned composite lap joints in monotonic and cyclic tension", Composites Science and Technology, 66 [13] 2163-2176 (2006). M. Grassi, X. Zhang, and M. Meo, "Prediction of stiffness and stresses in z-fibre reinforced co mposite laminates", Composites: Part A, 33 [12] 1653-1664 (2002). P. Chang, A.P. Mouritz, and B.N. Cox, "Flexural properties of z-pinned laminates", Composites: Part A, In Press, Corrected Proof (2006). R.D. Sweeting and R.S. Thomson, "The effect of thermal mismatch on Z-pinned laminated composite structures", Composite Structures, 66 [1-4] 189-195 (2004).

- 90 -

Chapter 2 - Literature Review

[152] G. Allegri and X. Zhang, "On the delamination and debond suppression in structural joints by Z-fibre pinning", Composites: Part A, In Press, Corrected Proof (2006). [153] S.L. Marco Zanetti, Giovanni Camino, "Polymer layered silicate nanocomposites", Macromolecular Materials and Engineering, 279 [1] 1-9 (2000). [154] S. Sinha Ray and M. Okamoto, "Polymer/layered silicate nanocomposites: a review from preparation to processing", Progress in Polymer Science, 28 [11] 1539-1641 (2003). [155] H. Miyagawa, K.H. Foo, I.M. Daniel, and L.T. Drzal, "Mechanical properties and failure surface morphology of amine-cured epoxy/clay nanocomposites", Journal of Applied Polymer Science, 96 [2] 281-287 (2005). [156] N.A. Siddiqui, R.S.C. Woo, J.-K. Kim, C.C.K. Leung, and A. Munir, "Mode I interlaminar fracture behavior and mechanical properties of CFRPs with nanoclayfilled epoxy matrix", Composites: Part A, In Press, Corrected Proof (2006). [157] G. Kelly and S. Hallstrom, "Bearing strength of carbon fibre/epoxy laminates: effects of bolt-hole clearance", Composites: Part B, 35 [4] 331-343 (2004). [158] F. Edgren, L.E. Asp, and R. Joffe, "Failure of NCF composites subjected to combined compression and shear loading", Composites Science and Technology, 66 [15] 2865-2877 (2006). [159] F.X.Q. T. Hatakeyama, Thermal Analysis: Fundamentals and applications to polymer science. John Wiley and Sons, 1999. [160] ASTM D695-02a Standard Test Method for Compressive Properties of Rigid Plastics. ASTM Annual Book of Standards, 2003. [161] ASTM D5961/D5961M-05 Standard Test Method for Bearing Response of Polymer Matrix Composite Laminates. ASTM Annual Book of Standards, 2005. [162] M.S. Wang and T.J. Pinnavaia, "Clay-Polymer Nanocomposites Formed from Acidic Derivatives of Montmorillonite and an Epoxy Resin", Chemistry of Materials, 6 [4] 468-474 (1994).

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Chapter 3 - Epoxy nanocomposites

CHAPTER 3

EPOXY NANOCOMPOSITES

3.1. INTRODUCTION One of the aims of the present study was to examine the possibility of improving the bearing strength of fibre reinforced composite laminates using nanocomposite s based on high performance epoxy resin as the matrix resin. This part of the study is presented in Chapter 4. It was not possible , however, to obtain nanoreinforced matrix resins at the time this work was commenced. A program of work was therefore undertaken to develop a suitable material and this work is presented in this chapter. From a review of the literature it was decided to use nanoclay as the reinforcement. The initial work examined the effect of nanoclay modification to obtain a suitable nanoclay for the work. Based on these findings a commercial nanoclay was chosen. To explore the effect of various variables, nanocomposite s based on conventional epoxy resin (DGEBA) were first examined. Using the experience gained from this part of the work, nanocomposites based on high performance epoxy resin (TGDDM) were made and then used as the matrix resin for carbon fibre reinforced laminates. The work in this chapter is presented separately for the nanoclay modification, the DGEBA nanocomposites and the TGDDM nanocomposites.

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Chapter 3 - Epoxy nanocomposites

3 .2. EXPERIMENTAL PROCEDURES 3 .2.1. Nanoclay modification Nanoclay modification was examined using four different surfactants. 3.2.1.1. Materials Nanoclay The nanoclay used was Cloisite Na+ (CNa+) nanoclay supplied by Southern Clay Products. This is a natural montmorillonite clay w ith an off-white colour and a density of 2.86 g/cm3. The cation exchange capacity is 92.6 meq/100g and the distance between the clay layers (d001) is 11.7 Å [1]. Surfactants The following four surfactants were used to modify the nanoclay: Octylamine (99%) with molecular formula: CH3 - (CH2)6 - CH2 - NH2 supplied by Sigma-Aldrich. This is a liquid with a density of 0.782 g/ml. Octamethylene diamine or 1,8-d iamino octane (98%) with molecular formula: H2N CH 2 - (CH2)6 - CH 2 - NH2 from Sigma-Aldrich. This is a powder with a melting point of 50-52°C. P-xylylamine or 4-methylbenzylamine (97%) with molecular formula: CH 3 - C 6H 4 - CH2 - NH2 from Sigma -Aldrich. This is a liquid with a density of 0.952 g/ml. Cetylamine or hexadecylamine (98%) w ith molecular formula: CH 3 - (CH2)14 - CH 2 NH 2 supplied by Sigma -Aldrich. This is a powder with a melting point of 45-48°C.

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Chapter 3 - Epoxy nanocomposites

3.2.1.2. Sample preparation Approximately 2g of pure nanoclay (CNa+) was oven dried at 80°C. A volume V1 (ml) of 1M hydrochloric acid (HCl) was mixed with 50ml of 50 volume % ethanol in distilled water at 50°C for 15 minutes. A stoichiometric mass m2 (g) of surfactant was then dissolved in 50ml of 50 volume % distilled water in ethanol, also at 50°C for 15 minutes, then mixed with the first solution for 30 minutes until the mixture became transparent. The dried nanoclay was taken from the oven then immediately added to the mixture and dispersed using a magnetic stirrer for 1 – 4 hrs, at a constant temperature of 50°C. The mixture was then vigorously mixed in an ultrasonic bath for 30 minutes at room temperature. The mixture was subsequently centrifuge d a t a speed of 3000 rpm for 5 minutes to separate out the treated nanoclay. This was then placed on a filter paper and washed in 50 volume % ethanol/distilled water until no trace of chloride was detected using 0.1N AgNO3. After that, the treated nanoclay was dried at 80°C for 24 hours and ground in a clean mortar to obtain small particles with a size of less than 53 μm, as collected by sieves. Finally, the modified nanoclay particles were stored in desiccators. The preparation technique is shown schematically in Figure 3.1. Pure clay 80°C HCl 1M

Mixing 50°C, 15 min

Water/Ethanol Amines

Mixing 50°C, 15 min

Stirring 50°C, 1-4 h

Stirring Sonication Centrifuge (3000 rpm)

Particles (100.00

CHE2

Hexadecylamine

18.25

94.00

CHE3

Hexadecylamine

17.82

>100.00

CHSC2

Hexadecylamine

CHSC4 CHMt2

Surfactant

Acid/amine ratio

18.25

94.00

Hexadecylamine

Surfactant concentration

18.14

>100.00

Hexadecylamine

Mixing time

18.25

94.00

CHMt8 Hexadecylamine 17.93 >100.00 The na nocomposites have been designated by the ID used for the modified CNa+

nanoparticles in Tables 3. 5-3.8. It can be seen from Table 3. 9 that the d-spacing was not appreciably changed in the composites made from the CNa+ clay modified with the short chain molecule surfactants

octylamine, diaminoctane and methylbenzylamine. However good

exfoliation was observed for the composites made with the long chain hexadecylamine modified nanoclay, with the d-spacing increasing from 18 Å in the modified nanoclay to

- 111 -

Chapter 3 - Epoxy nanocomposites

94 Å in the corresponding nanocomposite. Consistent with these results [15] , the nanocomposite made from the hexadecylamine modified nanoclay was transparent while the composites made with the other nanoclays were opaque. Varying the acid/amine ratio, the surfactant concentration and the nanoclay/surfactant mixing time had no appreciable effect on the d-spacing of the composites, Table 3. 9. Table 3.10 Compression modulus for modified CNa+ nanoclay nanocomposites.

Nanocomposite ID1

Surafactant used

Parameter varied

Compression modulus (MPa)

STDEV

PR

Pure resin

--

2309

56

Pure clay

--

2512

89

COE2

Octylamine

Surfactant

2462

181

CDE2

Diaminoctan

2393

231

CME2

Methylbenzylamine

2485

112

CHE2

Hexadecylamine

2570

55

CHE1

Hexadecylamine

2565

69

CHE2

Hexadecylamine

2570

55

CHE3

Hexadecylamine

2558

210

CHSC2

Hexadecylamine

2570

55

CHSC4

Hexadecylamine

Surfactant concentration

2564

129

CHMt2

Hexadecylamine

Mixing time

2570

55

CHMt8

Hexadecylamine

2517

103

CNa

1

+

Acid/amine ratio

The na nocomposites have been designated by the ID used for the modified CNa+

nanoparticles in Tables 3. 5-3.8. The addition of unmodified CNa+ nanoclay to the Ampreg 22 epoxy resin increased the compression modulus by 9% (from 2309 MPa to 2512 MPa, Table 3. 10), Figure 3.7. The increase was however more modest in the modified nanoclay nanocomposites prepared with octylamine, diaminoctane and methylbenzylamine (6%, 3% and 7% respectively). However, the nanocomposites prepared with the hexadecylamine modified nanoclay had a higher modulus (11%) than the unmodified nanoclay nanocomposites, Figure 3.7.

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Chapter 3 - Epoxy nanocomposites

Varying the acid/amine ratio, the surfactant concentration and the nanoclay/surfactant mixing time had no significant effect on the modulus of the composites, Figure 3.7.

2900

1

2

3

4

Compressive modulus (MPa)

2800 2700 2600 2500 2400 2300 2200 2100

CO E2 CD E2 CM E2 CH E2 CH E1 CH E2 CH E3 CH SC 2 CH CS 4 CH M t2 CH M t8

PR CN a

2000

Samples

Fig . 3.7 Compression modulus for nanocomposites reinforced with modified nanoclay particles and pure nanoc lay (CNa+) as well as for neat epoxy resin (PR) showing (1) effect of surfactants, (2) effect of acid/amine ratio, (3) effect of surfactant concentration and (4) effect of mixing time (error bars indicate one standard deviation). The d-spacings and compressive moduli for the composites made with the commerciall y available nanoclay particles C30B and I30E are shown in Table 3.11. An exfoliated nanocomposite with a d-spacing greater than 80 ? was obtained using the I30E nanoclay. This is attributed to the catalytic effect of the octadecylamine surfactant [3]. An intercalated nanocomposite with a d-spacing of 39? was obtained with the C30B nanoclay, Figure 3.8. C30B is modified by methyl tallow bis-2 hydroxylethyl quaternary ammonium [1]. Although this makes the particles very hydrophobic , the results indicate that the catalytic effect was weaker than in the 130E.

- 113 -

Chapter 3 - Epoxy nanocomposites

Table 3.11 Effect of nanoclay type on d-spacing and compression modulus.

Nanocomposite Nature ID Pure resin -+ CNa Pure C30B Commercial I30E Commercial

x

y

-2.5 2.5 2.5

33 33 33 33

d-spacing (Å) -13 39 >80

Modulus (MPa) 2309 2512 2447 2586

STDEV 56 89 89 82

450

Intensity (counts)

400 350

C30B

300

I30E

250 200 150 100 50 0 0

1

2

3

4

5

6

7

2Theta (degree)

Fig . 3.8 WAXD spectra for nanocomposites reinforced with C30B and I30E nanoclay. The results of the d-spacing measurements obtained from the WAXD patterns shown in Figure 3.8 were confirmed for the two commercial nanoclays by TEM, as shown by the micrographs in Figure 3.9 in which the dark lines delineate the nanoclay layers. This confirmed that WAXD, which was less time consuming than TEM, provided a reliable method for determining the distance between the layers of nanoclay in the nanocomposites.

- 114 -

Chapter 3 - Epoxy nanocomposites

(a)

(b)

Fig. 3.9 TEM micrographs of nanocomposites reinforced with C30B (a) and I30E (b).

2700

Modulus (MPa)

.

2600 2500 2400 2300 2200 2100 2000 PureResin

CNa

C30B

I30E

Samples

Fig. 3.10

Compression modulus of pure resin and nanocomposites reinforced w ith

different nanoclay particles (error bars indicate one standard deviation). The increase in compression modulus over that for the neat resin was lower for the nanocomposite made with the C30B nanoclay than for the nanocomposite made with the unmodified CNa+ nanocla y (6% compared with 9%), Table 3. 11 and Figure 3.10. However an increase of 12% was obtained for the nanocomposite reinforced with I30E nanoclay, similar to the level achieved for the CNa+ nanoclay modified with

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Chapter 3 - Epoxy nanocomposites

hexadecylamine. Accordingly, I30E nanoclay was selected as the reinforcement for the nanocomposites for the remainder of the work using the Ampreg 22 DGEBA resin. 3.3.2.3. Effect of mixing temperature The viscosity of the epoxy resin affects the level of dispersion of the nanoclay particles as well as the separation of the nanoclay layers in the matrix resin. An optimal viscosity needs to be found such that the nanoclay particles become wet by the resin sufficiently quickly to completely detach the clay layers and disperse them uniformly throughout the resin medium. The viscosity was varied by varying the mixing temperature, with the viscosity decreasing with increasing mixing temperature. An upper temperature limit for mixing of 60°C was experienced for mixing for the Ampreg 22 resin since excessive diluent vaporisation occurred above this temperature. Accordingly, the mixing temperature was varied within the range 30°C to 60°C. A mixing time of 1 hour was used with a mixing speed of 10,000 rpm. Otherwise the nanocomposites were prepared using the same parameters as described previously.

Relative Intensity

60°C

50°C

40°C

30°C

0

1

2

3

4

5

6

7

2Theta (degree)

Fig . 3.11 WAXD spectra for nanocomposites made with varying mixing temperatures.

- 116 -

Chapter 3 - Epoxy nanocomposites

No peak was observed by WAXD , indicating that all the composites obtained were exfoliated with a d-spacing larger than 80? , Figure 3.11. The compression modulus of the nanocomoposites also showed no significant variation with mixing temperature, Table 3.12 and Figure 3.12. Table 3.12 Effect of mixing temperature on compression modulus.

Sample ID Pure Resin MT30 MT40 MT50 MT60

Tmix (°C) -30 40 50 60

x -2.5 2.5 2.5 2.5

y 33 33 33 33 33

Modulus (MPa) 2309 2560 2510 2560 2506

STDEV 56 79 55 116 79

2800

Modulus (MPa)

.

2700 2600 2500 2400 2300 2200 2100 2000 PR

MT30

MT40

MT50

MT60

Samples

Fig . 3.12 Compression modulus of nanocomposites based on I30E nanoclay made w ith different mixing temperatures (error bars indicate one standard deviation). It was found that at the lower mixing temperatures it was necessary to continually cool the system to prevent the temperature increasing as a result of the frictional heating. No cooling was necessary at 60°C and, accordingly, this temperature was used in subsequent examination of the mixing parameters.

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Chapter 3 - Epoxy nanocomposites

3.3.2.4. Effect of mixing speed In mechanical mixing, the mixing speed plays an important role in the dispersion of the nanoclay particles at the microsca le, and even down to several hundred nanometers [14]. The shear force, however, is principally responsible for separating the nanoparticles at the nanoscale and for detaching the individual layers of the nanoclay. To examine the effect of mixing speed, the mixer speed was varied from 10,000 rpm to 26,000 rpm. Mixing was carried out for 1 hour at 60°C. All other experimental parameters were as described previously. The distance between the nanoclay layers in the nanocomposites was measured by WAXD and the results are shown in Figure 3.13. Again, no peak was detected, indicating that all the composites were exfoliated with a d-spacing larger than 80? . 450 400

26K rpm

.

350

10K rpm

Intensity (counts)

300 250 200 150 100 50 0 0

1

2

3

4

5

6

7

2Theta (degree)

Fig. 3.13 WAXD spectra of nanocomposites made with different mixing speeds. The compression modulus was higher in the nanocomposite made using a mixing speed of 26,000 rpm (2586 MPa) than in the one made using a speed of 10,000 rpm (2508 MPa), however , in view of the magnitude of the error the difference is not considered

- 118 -

Chapter 3 - Epoxy nanocomposites

significant, Figure 3.14. None the less it was decided to use a speed of 26,000 rpm in the remainder of the work. 2700 2600

Modulus (MPa)

2500 2400 2300 2200 2100 2000 PureResin

10K rpm

26K rpm

Samples

Fig . 3.14 Compression modulus of nanocomposites for different mixing speeds (error bars indicate one standard deviation). 3.3.2.5. Effect of mixing time The effect of mixing time was examined by mixing for times from 1 to 4 hours at 60°C at 26,000 rpm. The other experimental parameters were as before. No peak was observed in the WAXD patterns indicating that the nanocomposites were exfoliated with a d-spacing larger than 80? , Figure 3.15. It was also apparent that the modulus of the composites was not increased with increased mixing time, Figure 3.16. A mixing time of 1 hour was therefore used in subsequent work.

- 119 -

Relative Intensity

Chapter 3 - Epoxy nanocomposites

4 hrs

2 hrs

1 hr

0

1

2

3

4

5

6

7

2Theta (degree)

Fig. 3.15 WAXD spectra of nanocomposites made with different mixing times.

2800 2700

Modulus (MPa)

2600 2500 2400 2300 2200 2100 2000 PureResin

1hr

2hrs

4hrs

Sample

Fig . 3.16

Compression modulus of nanocomposites for different mixing times (error

bars indicate one standard deviation).

- 120 -

Chapter 3 - Epoxy nanocomposites

3.3.2.6. Effect of hardener concentration To examine the effect of hardener concentration the ratio of hardener/resin (y) was varied from 28 to 38 vol.%. Mixing was carried out for 1 hour at 60°C at a speed of 26,000 rpm. Otherwise, the experimental conditions were as before. Table 3.13 Effect of hardener concentration on compression modulus.

Sample ID Conc. (vol.%) Pure resin 33 HC28 28 HC33 33 HC38 38

x -2.5 2.5 2.5

y 33 28 33 38

Modulus (MPa) 2309 2685 2586 2543

STDEV 56 50 82 54

The compression modulus was found to decrease progressively with increased hardener content, Table 3.13, with a 5% decrease being observed at 38% , Figure 3.17. No peak was observed in the WAXD patterns, Figure 3.18, indicating that exfoliation had occurred at all hardener concentrations. Although a 4% decrease in modulus was observed with the recommended hardener concentration of 33% it was decided to continue with that concentration for the remainder of the work with Ampreg 22.

2800 2700

Modulus (MPa)

2600 2500 2400

(a)

2300 2200 2100 2000 PureResin

HC28

HC33

Samples

- 121 -

HC38

Chapter 3 - Epoxy nanocomposites

18 16.24

16

Percentage (%)

14 12

12 10.13

10

(b)

8 6 4 2 0 23

28

33

38

43

Hardener Concentration (phr)

Fig . 3.17 Diagrams showing (a) compression modulus of nanocomposites for varying hardener concentrations (error bars indicate one standard deviation) and (b) % increase

Relative Intensity

over that of neat resin.

38 vol.%

33 vol.%

28 vol.%

0

1

2

3

4

5

6

7

2Theta (degree)

Fig . 3.18 WAXD spectra of nanocomposite s for varying hardener concentrations.

- 122 -

Chapter 3 - Epoxy nanocomposites

3.3.2.7. Effect of curing temperature The effect of curing temperature was examined by curing at 60, 80 and 100°C for 2 hours followed by postcuring for 2 hours at 120°C. The other experimental conditions were as before. No peaks were observed in the WAXD patterns, Figure 3.19, indicating that for all three cure temperatures the nanoclay particles were well separated in the

Relative Intensity

resin medium at the nanoscale.

100°C

80°C

60°C

0

1

2

3

4

5

6

7

2Theta (degree)

Fig . 3.19 WAXD spectra of nanocomposites for different curing temperatures. The compression modulus of the nanocomposite s is shown in Figure 3. 20. The modulus was significantly lower in the sample cured at 60°C (2278 ± 121 MPa) than for the other two curing temperatures, but did not differ significantly between the sample cured at 80°C (2586 ± 82 MPa) and the one cured at 100°C (2552 ± 50 MPa).

- 123 -

Chapter 3 - Epoxy nanocomposites

2700

Modulus (MPa)

2600 2500 2400 2300 2200 2100 2000 PR

60°C

80°C

100°C

Samples

Fig . 3.20 Compression modulus of nanocomposites for different curing temperatures (error bars indicate one standard deviation). 3.3.2.8. Effect of curing time Curing time may also affect the degree of separation of the nanoclay particles, however no examination of this effect was available in the literature. Accordingly, this effect was examined by curing samples for 1, 2 and 4 hours (followed by postcuring for 2 hours at 120°C). Otherwise, the experimental conditions were as before. Table 3.14 Effect of curing time on d-spacing and compression modulus .

Sample ID Pure resin Ct1 Ct2 Ct4

Time (hrs) -1 2 4

X -2.5 2.5 2.5

y 33 33 33 33

WAXD -No peak No peak No peak

Modulus (MPa) 2309 2664 2586 2655

STDEV 56 92 82 42

No peaks were observed in the WAXD patterns, Table 3.14, indicating that the nanocomposites were exfoliated, while there was no significant difference in the modulus values, Figure 3.21.

- 124 -

Chapter 3 - Epoxy nanocomposites

2800 2700

Modulus (MPa)

2600 2500 2400 2300 2200 2100 2000 PR

1hr

2hrs

4hrs

Samples

Fig . 3.21

Compression modulus of nanocomposites for different curing times (error

bars indicate one standard deviation). 3.3.2.9. Effect of nanoclay content Nanocomposites with nanoclay/resin (x) contents of 0.5, 2.5, 5, 8.4 phr were prepared using the optimum conditions identified in the previous sections. Accordingly, the nanocomposites were prepared from I30E nanoclay mixed into the resin for 1 hour at a 60°C at a speed was 26,000 rpm. The ratio of hardener/resin (y) was 33 vol.% and the two stage curing procedure described in Section 3.3.2.1 was used to cure the nanocomposites. The results are given in Table 3. 15. Table 3.15 Effect of nanoclay content on d-spacing and modulus of nanocomposites.

Sample ID Pure resin NC1 NC5 NC10 NC15

Cont. (phr) -0.5 2.5 5.0 8.4

x -0.5 2.5 5.0 8.4

y 33 33 33 33 33

d-spacing (Å) Modulus (MPa) -2309 >100 2443 >100 2586 >100 2604 63 2790

- 125 -

STDEV 56 161 82 74 91

Relative Intensity

Chapter 3 - Epoxy nanocomposites

8.4 phr

5 phr

2.5 phr 0.5 phr

0

1

2

3

4

5

6

7

2Theta (degree)

Fig. 3.22 WAXD spectra of nanocomposites for different nanoclay contents. WAXD analysis showed that all the nanocomposites had d-spacings greater than 80 ? , except for the composite with the highest nanoclay loading of 8.4 phr for which a peak was detected at around 1.2 degrees, Figure 3. 22, corresponding to a d-spacing of 63 ? Table 3.15. The compression modulus increased progressively with nanoclay content, Table 3. 15. The addition of 0.5 phr nanoclay produced only a small increase of just over 1% (from 2309 MPa to 2443 MPa) over that of the neat resin but this increased progressively to 21% at 8.4 phr nanoclay, Figure 3.23. It was not possible to introduce nanoclay contents higher than 8.4 phr into the resin due to the high viscosity of the resulting mixture.

- 126 -

Chapter 3 - Epoxy nanocomposites

3000 2900

Modulus (MPa)

2800 2700 2600 2500

(a)

2400 2300 2200 2100 2000 0

0.5

2.5

5

8.4

Nanoclay content (phr)

25

20.83

Percentage (%)

20

15 12.78

12

(b)

10

5.8

5

0 0

1

2

3

4

5

6

7

8

9

10

Nanoclay content (phr)

Fig . 3.23 Diagrams showing (a) compression modulus of nanocomposites for different nanoclay contents (error bars indicate one standard deviation) and (b) % increase over that of neat resin.

- 127 -

Chapter 3 - Epoxy nanocomposites

3.3.2.10. Optical propert ies The addition of filler particles to transparent resins normally produces opaqueness. However when the particles are of nanodimensions (ie, smaller than the wavelength of light) they no longer impede the passage of light and the materials remains transparent [15]. Indeed, retention of transparency is a simple indicator of successful dispersion of nanoparticles in a host material. Figure 3.24, shows cylindrical samples approximately 20 mm in diameter and approximately 50 mm high made with and without additions of I30E nanoclay. The pure Ampreg 22 resin sample was translucent so that the word “UNSW” underneath the sample could not be seen. However the samples became transparent on the addition of the nanoclay.

Pure resin

0.5 phr nanoclay

2.5 phr nanoclay

8.4 phr nanoclay

Fig . 3.24

Optical appearance of pure resin and nanocomposite s based on I30E

nanoclay. While successful addition of nanoparticles should not reduce transparency, the fact that their addition lead to increased transparency was surprising. SEM and EDS analysis of the Ampreg 22 resin matrix in the nanocomposites indicated that it contained submicron sized particles of an aluminium rich material, presumably alumina , Figure 3.25. It would appear that these particles must agglomerate in the neat resin, but become more dispersed when the nanoclay is added. However further examination of this phenomenon was not considered within the scope of this thesis.

- 128 -

Chapter 3 - Epoxy nanocomposites

Si

(A)

Al

(B)

(C)

Fig . 3.25 EDS analysis of nanocomposite: (A) montmorillonite nanoclay particles, (B) aluminum oxide and (C) resin. 3 .3.3. TGDDM Nanocomposites High technology applications of fibre reinforced composites, such as in the aerospace industry, require a high performance resin matrix. TGDDM resins are amongst the stiffest available and laminates based on TGDDM are therefore amongst the highest performers. These resins have a high modulus (and also high temperature performance) because of their high cross-link density. However they have a low failure strain, leading to low compression strength after impact, which can be disadvantageous . In this section the introduction of nanoclay into the TGDDM epoxy resin is examined. As discussed earlier, Section 2.2.4, the introduction of nanoclay to improve the compressive modulus of the nanocomposite could be of significant benefit to improving the bearing performance of high performance composites reinforced with carbon fibres.

- 129 -

Chapter 3 - Epoxy nanocomposites

TGDDM epoxy resin has a high viscosity which needs to be considered when making nanocomposites using this resin. As with the DGEBA resin the effect of various processing parameters was systematically examined in order to optimize the conditions for producing a nanoclay reinforced matrix resin based on TGDDM The TGDDM resin used was Araldite® LY 568 containing 40 vol.% diluents which act as swelling agents during the intercalation process of nanocomposite preparation. Either intercalation or exfoliation can occur depending on the molecular weight. High molecular weight resin tends to exfoliate to a lesser extent than their lower molecular weight counterparts [16]. Based on the results obtained with the DGEBA resin, I30E nanoclay was used as the reinforcement. DETDA at a concentration of 35 phr, as recommended by the supplier, was used as the hardener for the LY 568 TGDDM resin. 3.3.3.1. Effect of surfactant on curing behaviour

3.5 0phr

3

2.5phr 2.5

.

7.5phr 10phr

Heat Flow (W/g)

2 1.5 1 0.5 0 -0.5 -1 -1.5 0

50

100

150

200

Temperature (oC)

Fig . 3.26

Exothermal curves for cure reaction of TGDDM/DETDA with various

nanoclay loadings. - 130 -

Chapter 3 - Epoxy nanocomposites

The effect of the surfactant (octadecylamine) in the I30E nanoclay on the cure behaviour of the resin was determined for 0-10% nanoclay additions using DSC. The results are shown in Figure 3.26. With increased nanoclay content, the exothermal peak temperature shifted to the left, indicating that the energy required for curing the system decreased progressively with increasing nanoclay content, as is also shown in Table 3.16. Table 3.16

Peak and onset temperatur es of nanocomposites having varying nanoclay

contents.

Nanoclay content in nanocomposites (phr) 0 1 2.5 7.5 10

Onset temperature (°C) 123.0 119. 5 111. 9 94.0 91. 6

Peak temperature (°C) 166. 1 165. 5 165. 2 160. 6 159.0

Based on this data a cure cycle of 1 hr at 110°C, followed by 2hrs at 155°C, with a final cure of 1 hr at 195°C , was chosen for the subsequent work. 3.3.3.2. Effect of mixing speed To examine the effect of mixing speed, nanoclay/TGDDM mixtures were stirred at 70°C for 1 hour at speeds of 10,000 and 26,000 rpm. The hardener/resin ratio was as given in Section 3.3.3 while the cure schedule was that described in Section 3.3.3.1. Nanoclay contents of 2.5 phr and 7.5 phr were examined. No significant change in the modulus with stirring speed was detected for either of the nanoclay contents, Figure 3. 27.

- 131 -

Chapter 3 - Epoxy nanocomposites

3100

Modulus (MPa)

.

2900

2700

2.5 phr 7.5 phr 2500

2300

2100 PureResin

10K

26K

Speed (rpm)

Fig . 3.27 Compression modulus of nanocomposites made with different mixing spe eds (error bars indicate one standard deviation). The interlayer spacing of the nanoclays was checked by WAXD, Figure 3. 28. No peak was detected for the composites with 2.5 phr nanoclay content but a peak was visible at around 1 degree for the 7.5 phr nanocomposites indicating a d-spacing of 80-85Å. No significant difference was however evident for the two different stirring speeds.

- 132 -

Chapter 3 - Epoxy nanocomposites

900 800 700

10K rpm 26K rpm

Intensity (counts)

600 500

(a)

400 300 200 100 0 0

1

2

3

4

5

6

7

2Theta (degree) 1000 900 800 10K rpm

Intensity (counts)

700

26K rpm

600

(b)

500 400 300 200 100 0 0

1

2

3

4

5

6

7

2Theta (degree)

Fig. 3.28 WAXD spectra of nanocomposite s w ith (a) 2.5 phr and (b) 7.5 phr nanoclay.

- 133 -

Chapter 3 - Epoxy nanocomposites

3.3.3.3. Effect of mixing temperature To examine the effect of mixing temperature, nanoclay/TGDDM mixtures were stirred at temperatures from 60-80°C for 1 hour at a speed of 26,000 rpm. The other experimental parameters were as in Section 3.3.3. 2. As noted earlier, higher mixing temperatures reduce the viscosity of the resin, improving the ability of the epoxy molecules to penetrate into the galleries of the nanoclay. Within the range of temperatures examined there was, however, no significant change in the modulus, Figure 3. 29. WAXD patterns also showed no effect of mixing temperature on the level of exfoliation, Figure 3. 30. In view of these findings a temperature of 70°C was chosen as the mixing temperature for the remainder of this part of the study.

3100

Mudulus (MPa)

.

2900

2700

2.5phr 7.5phr 2500

2300

2100 PureResin

60

70

80

Temperature (oC)

Fig. 3.29 Compression modulus of nanocomposites for different mixing temperatures (error bars indicate one standard deviation).

- 134 -

Chapter 3 - Epoxy nanocomposites

900 800 700

60C 70C

Intensity (counts)

600

80C 500

(a)

400 300 200 100 0 0

1

2

3

4

5

6

7

2Theta (degree) 1000 900 800

60C

Intensity (counts)

700

70C

600

80C

(b)

500 400 300 200 100 0 0

1

2

3

4

5

6

7

2Theta (degree) Fig. 3.30 WAXD spectra of nanocomposites with (a) 2.5 phr and (b) 7.5 phr nanoclay. 3.3.3.4. Effect of mixing time Nanoclay/TGDDM mixtures were stirred at 70°C for 1 - 4 hours using the same experimental conditions as before. There was no significant change in the compression - 135 -

Chapter 3 - Epoxy nanocomposites

modulus with stirring time for the nanocomposites containing 2.5 phr nanoclay, Figure 3.31. 3300

.

3100

Modulus (MPa)

2900

(a) 2.5 phr

2700

7.5 phr 2500

2300

2100 PR

0.5

1

2

4

Time (hours) 35 29.04

Percentage (%)

.

30

25

(b)

22.11 18.62

20

2.5 phr 7.5 phr

15 9.48 8.93

10

8.67

8.33

5

0 0

1

2

3

4

5

Mixing Time (hrs)

Fig. 3.31 Diagrams showing (a) compression modulus of nanocomposites for different mixing times (error bar s indicate one standard deviation) and (b) % increase over that of neat resin.

- 136 -

Chapter 3 - Epoxy nanocomposites

However the modulus increased progressively with stirring time (from 2790 to 3035 MPa) for the composites with the higher clay loading. There was, however, no noticeable change in the WAXD patterns with stirring time for either of the nanoclay loadings , Figure 3.32. 1000

Intensity (counts)

900 800

0.5 hr

700

1 hr 2 hrs

600

(a)

4 hrs

500 400 300 200 100 0 0

1

2

3

4

5

6

7

2Theta (degree)

1000 900 800

1 hr

Intensity (counts)

700

2 hrs 4 hrs

600

(b)

500 400 300 200 100 0 0

1

2

3

4

5

6

7

2Theta (degree) Fig. 3.32 WAXD spectra of nanocomposites with (a) 2.5 phr and(b) 7.5 phr nanoclay.

- 137 -

Chapter 3 - Epoxy nanocomposites

3.3.3.5. Effect of degassing time The effect of degassing time after stirring was examined for 7.5 phr nanoclay/TGDDM mixtures stirred for 1 hour using the parameters described previously . The mixtures were then degassed under vacuum until visible bubble removal ceased. They were then degassed for a further period of 15, 30 or 60 minutes. The compressive modulus of the nanocomposites is shown in Figure 3. 33. There appeared to be an improvement from the additional 15-30 minutes degassing although, in view of the magnitude of the error bars, this is considered to be of marginal significanc e. No significant effect of additional degassing time was seen in the WAXD patterns. However, it was decided to use an extra 15-30 minutes in the subsequent work.

3100

Modulus (MPa)

2900

2700

(a)

2500

2300

2100 PR

0 min

15 min

30 min

Additional degassing time

- 138 -

60 min

Chapter 3 - Epoxy nanocomposites

30

25.17

25.04

Percentage (%)

25

20

18.62 16.45

(b)

15

10

5

0 0

15

30

60

Additional degassing time (min)

Fig . 3.33

Diagrams showing (a) compression modulus of nanocomposite made w ith

7.5 phr nanoclay for various additional degassing times (error bars indicate one standard

Relative Intensity

deviation) and (b) % increase over that of neat resin.

60 min 30 min 15 min 0 min

0

1

2

3

4

5

6

7

2Theta (degree)

Fig . 3.34 WAXD spectra of nanocomposites containing 7.5 phr nanoclay for various additional de gassing times.

- 139 -

Chapter 3 - Epoxy nanocomposites

3.3.3.6. Effect of curing temperature The effect of curing temperature was examined for 7.5 phr nanoclay/TGDDM mixtures prepared using the conditions described previously. The mixtures were then cured using a three stage curing temperature but with the initial temperature T1 (1hr), the intermediate temperature T2 (2hrs) and the final temperature T3 (1hr) being varied according to the schedules I (initial temperature varied), II (intermediate temperature varied) and III (final temperature varied) , as shown in Table 3. 17. Table 3.17 Effect of cure temperature on compression modulus of nanocomposites.

Group

Standard I II III

T1 (1hr) - T2 (2hrs) - T3 (1hr)

Modulus (MPa)

Percentage (%)

T1

T2

T3

110

155

195

2941 ± 79

25.04

90

155

195

2909 ± 74

23.68

130

155

195

3012 ± 77

28.06

110

140

195

2986 ± 49

26.96

110

170

195

3014 ± 32

28.15

110

155

180

2940 ± 48

25.00

110

155

210

3068 ± 77

30.44

While there was a general upward trend in the modulus with increased temperature in stages I and III, Table 3.17 and Figures 3.35-3.36, the differences were reasonably small and were considered to be generally within the experimental error. No significant differences were also apparent from the WAXD spectra, Figure 3. 37-3.39. In view of these findings the temperatures of the standard cure schedule (as given in Section 3.3.3.1 and shown in Table 3.17) were used in the remainder of the work.

- 140 -

Chapter 3 - Epoxy nanocomposites

3300

I

II

III

Modulus (MPa)

3100

2900

2700

2500

2300

2100 PR

90

110

130

140

155

170

180

195

210

Temperature (oC)

Fig . 3.35

Compression modulus of nanocomposites with varying initial (I),

intermediate (II) and final (III) curing temperatures (error ba rs indicate one standard deviation). 35 30.44 28.06

30

26.96

25.04

25.04

I

II

23.68

25

Percentage (%)

28.15 25

25.04

20 15 10

III

5 0 80

100

120

140

160

180

200

220

Temperature (oC)

Fig . 3.36

Percentage increase of compressive modulus over that of neat resin for

nanocomposites cured at various initial (I), intermediate (II) and final (III) curing temperatures.

- 141 -

Relative Intensity

Chapter 3 - Epoxy nanocomposites

130°C

110°C 90°C

0

1

2

3

4

5

6

7

2Theta (degree)

Relative Intensity

Fig . 3.37 WAXD spectra of nanocomposites for various initial cur e temperatures.

170°C 170°C

155°C 140°C

0

1

2

3

4

5

6

7

2Theta (degree)

Fig . 3.38

WAXD spectra of nanocomposites for various intermediate cure

temperatures.

- 142 -

Relative Intensity

Chapter 3 - Epoxy nanocomposites

210°C

195°C 180°C

0

1

2

3

4

5

6

7

2Theta (degree)

Fig. 3.39 WAXD spectra of nanocomposites for various final cure temperatures. 3.3.3.7. Effect of curing time The effect of cure time was examined by varying the times of the initial stage t1, the intermediate stage t2, and the final stage t3, of the standard cure schedule between 0 and 2 hours, according to the schedules I (time at initial temperature varied), II (time at intermeditae temperature varied) and III (time at final temperature varied), as shown in Table 3. 18. All other experimental parameters were as used previously. Table 3.18 Effect of curing time on compression modulus of nanocomposites. 110 (t1) - 155 (t2) - 195 (t3)

Group

Standard I II III

Modulus (MPa)

Percentage (%)

t1 (hrs)

t2 (hrs)

t3 (hrs)

1

2

1

2941 ± 79

25.04

0

2

1

2996 ± 39

27.38

2

2

1

2898 ± 75

23.21

1

0

1

2926 ± 77

24.40

1

1

1

2944 ± 65

25.17

1

2

0

2926 ± 38

24.40

1

2

2

3026 ± 29

28.66

- 143 -

Chapter 3 - Epoxy nanocomposites

The results showed a general decrease in the modulus with increased time at the initial temperature, no change with increased time at the intermediate tempe rature and a general increase in the modulus with increased time at the final temperature, Table 3. 18 and Figures 3.40-3.41. 3300

I

II

III

Modulus (MPa)

3100

2900

2700

2500

2300

2100 PR

0

1

2

0

1

2

0

1

2

Time (hours)

Fig . 3.40

Compression modulus of nanocomposite s cured for different initial (I),

intermediate (II) and final (III) cure time s (error bars indicate one standard deviation). 35 30

I

II

25.04

25

Percentage (%)

III

28.66

27.38 23.21

24.4

25.17

25.04

24.4

25.04

1

2

0

1

20 15 10 5 0 0

1

2

0

2

Time (hours)

Fig . 3.41

Percentage increase of compression modulus over that of neat resin for

nanocomposites cured at various initial (I), intermediate (II) and final (III) cur e times. - 144 -

Chapter 3 - Epoxy nanocomposites

However, with the exception of the increase at the longest time for the final stage, the changes are considered to be within the experimental error. No significant differences

Relative Intensity

were apparent from the WAXD spectra, Figure 3.42-3.43.

2 hrs

1 hr 0 hr

0

1

2

3

4

5

6

7

2Theta (degree)

Relative Intensity

Fig. 3.42 WAXD spectra of nanocomposites for various initial curing time s.

2 hrs

1 hr

0

1

2

3

4

5

6

7

2Theta (degree)

Fig. 3.43 WAXD spectra of nanocomposites for various final curing time s.

- 145 -

Chapter 3 - Epoxy nanocomposites

3.3.3.8. Effect of nanoclay content The effect of varying the content of nanoclay particles from 1 phr to 20 phr was examined for composites prepared using the optimum conditions identified in the previous sections except that final cure time was kept at 1 hour as in the standard cure schedule given in Section 3.3.3.1.

20 phr

Relative Intensity

17.5 phr 15 phr 12.5 phr 10 phr 7.5 phr 5 phr 2.5 phr 1 phr 0

1

2

3

4

5

6

7

2Theta (degree)

Fig . 3.44 WAXD spectra for nanocomposites for varying I30E nanoclay contents. The distance between the layers of nanoclay, as measured by WAXD, varied with clay loading, as indicated in Figure 3.44. For a diffraction angle larger than 0.75°, no peaks were observed with nanoclay contents from 1 to 5 phr, indicating that the d-spacing of the nanoclay layers was greater than 80 Å, consistent with the nanoclay particles being fully exfoliated. However , with further increase in the clay loading the distance between the nanoclay layers decreased, being only around 70-80 Å for the composites with 7.520 phr of nanoclay. These nanocomposites were not considered to be fully exfoliated and are therefore referred to as pre-exfoliated nanocomposites.

- 146 -

Chapter 3 - Epoxy nanocomposites

(a)

(b)

(c)

(d)

Fig. 3.45 High magnification TEM images of nanocomposites with various nanoclay contents: (a) 2.5 phr, (b) 7.5 phr, (c) 12.5 phr and (d) 20 phr. These results were confirmed by TEM observations as shown in Figure 3.45. The dark lines define the nanoclay layers. For the 2.5 phr nanoclay composite the spacing of the nanoclay platelets was about 120-125Å, Figure 3.45(a). This then decreased to 85Å for 7.5 phr of nanoclay, Figure 3.45(b), 75-80Å for 12.5 phr of nanoclay, Figure 3.45(c), and 60Å for 20 phr of nanoclay, Figure 3.45(d). The nanoclay particles were present in groups, with almost all the platelets of nanoclay within a group being parallel and the individual groups being well dispersed, indicating that partially exfoliated and full intercalated nanocomposites had been obtained. It was also noticed that the number of platelets in a group (stack) increased from approximately 10 to approximately 15, with increased nanoclay content, while the thickness of the stacks increased from approximately 60 to approximately 100 nm.

- 147 -

Chapter 3 - Epoxy nanocomposites

3700

Modulus (MPa)

.

3500 3300 3100 2900 2700 2500 2300 2100 0

1

2.5

5

7.5

10

12.5 15

17.5

20

Nanoclay Content (phr)

Fig . 3.46 Compression modulus of TGDDM/DETDA/I30E nanocomposites (error bars indicate one standard deviation).

60 49.63

50

.

44.86

Percentage (%)

40

34.86

33.29

28.83

30 18.62

20 12.29

10

8.218.93

0 0

2.5

5

7.5

10

12.5

15

17.5

20

22.5

Nanoclay content (phr)

Fig . 3.47

Percentage increase in compression modulus over that of neat resin for

nanocomposites made with varying nanoclay contents.

- 148 -

Chapter 3 - Epoxy nanocomposites

The compressive modulus of the nanocomposites increased progressively with increased nanoclay content, Figure 3.46. The modulus increased from a value 2352 MPa for the pristine resin to 3030 MPa at a loading of 10 phr of nanoclay, eventually attaining a value of 3517 MPa at a loading of 20 phr of nanoclay. These values correspond to improvements of almost 30% for 10 phr nanoclay and 50% for 20 phr nanoclay additions, Figure 3.47. An increase in modulus in polymers is usually also accompanied by an increase in the glass transition temperature of the material. However, to the contrary the Tg of the nanocomposites decreased progressively with clay content, as shown in Figure 3.48.

200 R2 = 0.9767

.

160

Tg (oC)

120

80

40

0 0

5

10

15

20

25

Nanoclay content (phr) Fig. 3.48 Tg of nanocomposites with varying nanoclay contents. The dispersion of the nanoclay in the resin at a microscale was examined on metallographically polished surfaces using SEM, Figure 3.49. The particles observed were spherical, or nearly spherical, in shape, with diameters of 100-150 nm. It was noticed that smoother surfaces were obtained for the nanocomposites containing lower nanoclay contents. With the higher nanoclay loadings, the nanoparticles tended to agglomerate into larger particles, typically 400 nm in diameter, but the se were still

- 149 -

Chapter 3 - Epoxy nanocomposites

uniformly dispersed in the resin matrix. Holes obs erved on the polished surfaces appeared to have resulted from particle pull-out during polishing.

(a)

(b)

(c)

(d)

Fig. 3.49 SEM images of various nanoclay contents in the resin matrix: (a) 2.5 phr, (b) 7.5 phr, (c) 12.5 phr and (d) 20 phr. 3.4. DISCUSSION 3 .4.1. Nanoclay modification 3.4.1.1. Effect of surfactants The alkylammonium amines octadecylamine, diaminoctane, methylbenzylamine all had the same number of carbon atoms (8 carbons) but differe nt structures and their use was examined to investigate the effect of alkyl chain structure on nanoclay modification. These short-chain amines are considered as very strong catalysts for the epoxy curing reaction [17]. However, they produced only a minimal increase in the interlayer spacing of the nanoclay (from 12.77 to 13.58 Å), Table 3. 5 and Figure 3. 6. This suggested that only a monolayer of the surfactant had occupied the gallery of the nanoclay [6].

- 150 -

Chapter 3 - Epoxy nanocomposites

Hexadecylamine however has a substantially longer alkyl chain (16 carbons), and produced a subs tantial increase in the gallery height of the nanoclay from 12.77 to 18.25 Å (ie, 43%). The results of the present study agree with the findings of previous studies conducted by Lan et al [6] , Zilg et al [7], Chen and Yang [13] , Chen and Curliss [18]. By treating nanoclay with long-chain alkyl ammonium ions, the inorganic cations are replaced by organic ones, and convert the nanoclay to being hydrophobic. As the chain length increases, the intercalated chains adopt a more ordered, liquid-like structure in the gallery. The ordered chain form should have a better ability to penetrate into the nanoclay gallery than the disordered form produced by the shorter alkyl chains since the latter is likely to produce obstacles which would hinder ingress of other chains into the gallery. As a result, it is considered that the arrangement of onium ions in the gallery of the nanoclay changes from a monolayer to bilayer. 3.4.1.2. Effect of acid/amine ratio The acid/amine ratio affects the protonation process of the amine to obtain alkylammonium cations since the active hydrogen cations of the acid protonate the -NH2 amine groups. The density of exchanged sodium cations in the gallery depends on the available alkylammonium cations. When the ratio of acid/amine is less than 1:1 (mol/mol) , there would be un-protonated amines remaining which would have a low affinity for the inorganic cations. Intercalation of the nanoclay by the amine surfactant would be expected to decrease since there would still be sodium cations in the gallery. These could restrict expansion of the gallery as the un-exchanged sodium cations may act as pinning points, preventing swelling of the nanoclay in the surfactant environme nt. This appeared to have minimal effect in the present study. No reduction in d-spacing was evident for octadecylamine and diaminoctane, Table 3.6, while only minimal reduction (from 18.3 to 17.5 Å) was seen for the third surfactant, hexadecylamine, for which this effect was examined, Figure3.3. While the results above did not indicate any appreciable effect of acid amine ratio (within the range of 0.5-2 examined) it was however found to be difficult to centrifuge the modified nanoclay samples with an acid/amine ratio of less than 1.0. - 151 -

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The presence of acid also affects the characteristics of nanoclay. According to Park’s [9] research, acid treatment affects the pH and acid-base values of the clay, increasing the Boehm’s acid value or decreasing the base value. The acid value of the modified organoclay is higher than that of as -received clay due to an increase in cation groups on the clay surface or reaction of the OH gr oups at the surface. 3.4.1.3. Effect of surfactant concentration It was expected that increasing surfactant concentration could make the interlayer spacing of the nanoclay more expanded. However, the experimenal results, Table 3. 7 and Figure 3.4, showed that there was no significant change with increasing surfactant concentration. This implies that the gallery of the nanoclay was saturated with surfactant when the ratio of surfactant/nanoclay was 1:1 (mol/mol) indicating that all the sodium cations in the gallery of the nanoclay had been exchanged by the alkylammonium cations. If the stoichiometric concentration of these organic cations is less than that of the sodium cations in the galleries of the nanoclay, it would restrict the expansion of the nanoclay layers by the pinning effect described above. Again this did not appear to be the case here. Only one of the surfactants (octylamime) was examined at a lower than stiochiomentric concentration (0.5:1). While the d-spacing was marginally lower than at the stoichiometric concentration (13.5 Å compared with 13.6 Å) this difference is considered to be within the experimental error. It is noted that the effect of surfactant concentration could become more important when the nanoparticles are incorporated into epoxy resin since there is a correlation between the onium ion concentration of the bulk solution and the amount of intercalated gallery cations within the clay layers [13].

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3.4.1.4. Effect of mixing time Mixing time was examined to see whether increased penetration of the amine molecules into the gallery of the nanoclay particles could be achieved. However this was not observed, with no improvement in d-spacing being seen when the mixing time was extended from 1 to 4 hours, Table 3. 8 and Figure 3.5. It appears that 1 hour of mechanical stirring, together with 30 minutes subsequent sonication, was sufficient to produce effective breakdown of the nanoclay particles and to allow the surfactant molecules to diffuse into the galleries to produce full saturation. Thus no increase in interlayer spacing of the nanoclay was produced with further mixing time. 3.4.1.5. Conclusions •

Alkylammonium surfactants with short alkyl chains (8 carbon atoms) amines did not expand the layers in the nanoclay appreciably.

•

The longer alkyl chain (16 carbon atoms) surfactant produced a substantial increase in the interlayer spacing and is considered suitable for modification of the nanoc lay for reinforcing epoxy resin

•

Within the ranges examined, the acid/amine concentration, surfactant concentration and mixing time had no appreciable influence on the interlayer spacing.

•

Surfactant modified nanoparticles with acid/amine ratios less than 1:1 proved difficult to centrifuge.

3 .4.2. DGEBA Nanocomposites The addition of 2.5 phr I30E nanoclay to the DGEBA resin produced a decrease in the curing temperature of around 4°C , Figure 3.6. This is attributed to the catalytic effect of the octadecylamine cations occupying the galleries of the I30E nanoclay on the curing reaction of the epoxy/amine system. A similar phenomenon was observed by Tolle and

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Anderson [19], in which the addition of I30E resulted in a significantly earlier onset temperature of polymerization, as well as a decrease of T max , as compared with that for unreinforced epoxy-hardener system. 3.4.2.1. Effect of nanoclay type During the curing reaction of the epoxy resin, the gradual diffusion of epoxy molecules into the nanoclay galleries pushes the individual nanoclay layers out of the tactoids. Intragallery polymerization takes place providing the energy to overcome the attractive electrostatic force between the layers of the nanoclay, leading to their separation. As a result, exfoliation of the nanoclay can be achieved. However no appreciable increase was observed in the distance between the layers of the nanoclay modified by the short-chain amines, octylamine, diaminoctane, and methylbenzylamine, after introduction of the clays into the DGEBA resin (Table 3. 9), even though primary short alkyl-chain amines are normally more versatile than longer ones. The lack of improvement obtained in the present study is attributed to the small initial d-spacing of these modified nanoclay particles. It is considered that the space inside the galleries of the nanoclay would have become essentially fully filled by these protonated short-chain amines, limiting the available space for the organic epoxy resin molecules. This is consistent with the views of Lan et al. [6] who considered that short alkylammonium ions, exchanged in the gallery of the nanoclay, restrict the intragallery diffusion of the epoxy molecules. It thus becomes more difficult to achieve the curing conditions required to get a balance between the intra- and extragallery polymerization rates. Consequently, un-intercalated nanocomposites were formed, in which the nanoclay particles often aggregated together, forming clumps in the resin matrix, resulting in serious defects in the network. Thus, the compression modulus of the composites was even lower than that of the composites reinforced with the pure nanoclay, as seen in Figure 3.7. In contrast to the findings of the present study, Zilg et al [7] and Chen and Curliss [18] , obtained exfoliation of the nanoclay in epoxy resin with nanoclay modified with amine surfactants having six carbon atoms in the chain. This difference may be due to a difference in the DGEBA resin system used. - 154 -

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In general, the final spacing of nanoclays in the resin is controlled by the chain length of the alkylammonium cation, regardless of the initial orientation of the ion in the gallery [20]. With hexadecylamine, exfoliated nanocomposites were obtained with an interlayer spacing of the nanoclay layers of around 100Å (Table 3. 9). It is clear that with the larger initial d-spacing in the modified nanoclay produced by the longer alkyl chain amine , it was easier for the organic epoxy polymer molecules to penetrate into the galleries of the nanoclay and occupy the available space. Salahuddin [15] also concluded that a greater distance between the layers of nanoclay allowed new organic species to diffuse into the gallery more easily and substantially exfoliate the nanoclay. To get clay exfoliation, the attractive electrostatic force between the negatively charged silicate layers and the gallery cations must be overcome by the polymerization reaction of the epoxy resin. In addition, Lan et al. [20] found that the primary amine cations modifying the nanoclay acted as acid catalysts rather than as curing agents for the epoxy resin, and considered this to be very important in improving the nanoclay layer expansion in the resin environment. Although exfoliated nanocomposites were obtained for the composites made with hexadecylamine modified nanoclay, the increase in the modulus of was only about 10% higher than for the composites made with unmodified nanoclay, Figure 3.7. This is not , however, surprising in the view of the relatively low content (2.5 phr ) of nanclay used. The two commercial nanoclays, C30B and I30E, behaved quite differently to one another. Only limited separation of the clay layers (around 39 Å) was achieved for the C30B nanoclay when added to the DGEBA resin, Figure 3. 8. This is similar to the results obtained by Basara et al. [21] and can be explained by the weak catalytic effect of the quaternary amine which is used to modify the C30B nanoclay [20, 22] . The intercalated nanocomposites obtained with the C30B nanoclay produced only a modest increase in the modulus (6%) compared to that of the neat resin, which was again slightly lower than that obtained with the unmodified nanoclay (9%),Figure 3.10. Substantialexfoliation (around 100 Å) was achieved when I30E nanoclay was added to the epoxy resin , as determined from the WAXD pattern (Figure 3. 8) and confirmed by TEM (Figure 3.9). This is attributed to dissociation of the octadecylamine cations, - 155 -

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CH 3(CH 2)17NH3+, occupying the clay galleries generating protons, causing acid catalyzed ring opening homopolymerization of the epoxide . According to Lan et al.’s [23] study, intragallery polymerization was accelerated by the presence of CH 3(CH 2)17NH3+ in I30E and became comparable to that of the extragallery polymerization. As a result, exfoliation of the nanoclay could be achieved. Consequently, the modulus of the exfoliated nanocomposites based on I30E nanoclay was higher than that of the intercalated composites (12% cf 6%), Table 3.11. This increase is attributed to the high aspect ratio of the layered silicate when separated at the nanoscale [24, 25]. 3.4.2.2. Effect of mixing conditions Nanocomposite s reinforced with uniformly dispersed nanoclay can be achieved through high-shear force mixing techniques [26]. The purpose of high-shear mixing is to break the primary clay particles down at the nanoscale. Decreasing the resin viscosity could assist exfoliation of the clay layers by allowing better diffusion of the epoxy molecules into the clay galleries. Similarly, increasing the mixing speed and mixing time could also increase penetration of the epoxy molecules into the galleries. 3.4.2. 2.1. Mixing temperature The viscosity of the resin can be decreased by increasing the mixing temperature. T his makes the epoxy molecules more mobile and would allow them to penetrate more easily into the gallery of the nanoclay. The more epoxy molecules in the intragallery of the clay, the greater the expansion of the nanoclay layers. The mixing temperature was varied within the range 30-60°C. It was noted that the viscosity was reduced substantially at 60°C but higher temperatures could not be used because of excessive diluent evaporation. There was however no significant change with mixing temperature with the composites being exfoliated at all temperatures, Figure 3.11 and the modulus remaining constant, Figure 3.12. It would appear therefore that the viscosity was sufficiently low at all temperatures examined.

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3.4.2. 2.2. Mixing speed High shear force stirring also reduces viscosity due to the high friction and could also enhance exfoliation. However, no significant difference was again seen in the level of exfoliation nor the compressive modulus for the two speeds examined (10,000 and 26,000 rpm), Figure 3.13 and 3.14. This would indicate that the viscosity was already sufficiently low at the lower speed to produce effective exfoliation. 3.4.2. 2.3. Mixing time The use of a longer mixing time would increase the time available for the epoxy molecules to diffuse into the nanoclay layers and times from 1 to 4 hours were examined. However again, there was no significant change in the level of exfoliation nor any increase in the modulus with increased mixing time , Figures 3.15 and 3.16. To the contrary, the modulus of the nanocomposites was , in fact, significantly reduced with mixing times of longer than 1 hour (Figure 3.16). This was probably caused by evaporation of the diluents in the DGEBA at 60°C during mixing. As result, the viscosity would increase so as to restrict the uniform separation of the nanoclay in the epoxy matrix. Jiankun et al. [27] also examined mixing time and found that prolonging the mixing time beyond a certain point did not further improve intercalation. In their work, a time of 20 minutes mixing at 70 - 80°C for a DGEBA resin/nanoclay was enough to create full intercalation. It is noted that the present work was carried out at a nanoclay content of 2.5 phr. At higher nanoclay contents the effects of mixing time, speed and temperature could become more significant. 3.4.2.3. Effect of curing conditions 3.4.2. 3.1. Hardener concentration As seen in Table 3.13 and Figure 3.17, the compressive modulus of the nanocomposites decreased progressively with increased hardener concentration. This is consiste nt with the work of Chin et al. [28] and Deng et al. [29], both of whom found that the degree of

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exfoliation decreased with higher amounts of curing agent. They considered that higher concentrations of hardener caused the extra gallery reactions to become faster. As a result, the viscosity of the epoxy resin in extragallery increases so that less resin and hardener can enter the intergallery of organoclay. During the curing reaction, there is competition between the intra- and extragallery polymerization speeds [10, 27, 30]. If intragallery polymerization is faster, more epoxy molecules can migrate into the clay galleries as the viscosity of the remaining resin is still not high. Additionally, expansion of the clay layers will only meet minimal resistance from the resin in the extragallery. Consequently, exfoliated nanocomposites will readily form. When the extragallery cures faster, the high viscosity resulting from the curing process will prevent the epoxy molecules penetrating into the nanoclay galleries. Moreover, the high viscosity of the resin will oppose expansion of the clay restricting the degree of layer separation. 3.4.2. 3.2. Cure temperature The temperature of curing is an important factor controlling the properties of epoxy nanocomposites since it not only affects the rate of the cure reaction but also the rate of diffusion of the epoxy and hardener molecules into the clay layers [6]. If the rate is too slow, the polymerization speed outside the layers (extragallery polymerization) will be faster than the polymerization speed between the layers (intragallery polymerization) and only intercalated nanocomposites will form. While for the temperature range examined (60-100ºC) all composites appeared to be fully exfoliated, Figure 3. 19, the modulus was significantly lower in the sample cured at 60ºC than for those cured at higher temperatures (80 and 100ºC), Figure 3.20. At low nanoclay contents, as used here, the absence of any peak in the WAXD spectra is not always reliable for establishing full exfoliation [31] and it is considered that the decreased modulus observed at 60ºC is indicative of a reduced level of exfoliation. The cur ing temperature of 60°C is presumably too low to allow sufficient mass diffusion of resin and curing agents into the clay layers before gelation and vitrication locks in the morphology. This may also be coupled with bridging of the silicate layers caused by reactions at their edges. - 158 -

Chapter 3 - Epoxy nanocomposites

The compressive modulus was increased at the higher temperatures indicating better exfoliation of the nanoclay layers in the resin. According to Kornmann [32] a higher curing temperature favours the the intragallery cure kinetics as a result of the increased reactivity of the epoxy system and the increased diffusion rate of the epoxy and the curing agent between the layers. This leads to exfoliation of the clay. This was also found by Tolle and Anderson [33] who showed that a higher temperature not only resulted in earlier initiation of the exfoliation process, but also resulted in greater silicate spacing. Dean et al. [34] also concluded that increasing the curing temperature increased the interlayer spacing. At the higher temperatures, the combination of a lower prepolymer viscosity and faster intragallery polymerization are sufficient to yield more significant expansion of the layers. 3.4.2. 3.3. Cure time Increasing the cure time from 1to 4 hours did not increase the compressive modulus of the nanocomposites, Figure 3.21. This indicates that with the postcure of 2 hours at 120ºC a satisfactory network was formed using an initial cure of one hour curing at 80°C. It is noted that at higher nanoclay contents even shorter cure times may be suitable since the presence of increase d amounts of organoclays has been shown to reduce the gel time and also completed cure time [35]. Increased temperature also reduces cure time. 3.4.2.4. Effect of nanoclay content The effect of nanoclay content was examined by adding amounts of up to 8.4 phr nanoclay (0.5, 2.5, 5 and 8.4 phr). Above 8.4 phr the viscosity of the resin/nanoclay mixture became excessive and higher loadings could not be used. With the content of nanoclay less than 5 phr, exfoliated nanocomposites with a nanoclay interlayer spacing larger than 100 Å, were obtained except at the highest clay loading, for which an intercalated composite with an interlayer spacing of 63 Å was obtained, Table 3.15 and Figure 3.22. As the clay loading is increased, the volume of intergallery requiring filling increases while the volume of resin available decreases. This would slow the rate of filling of the galleries and account for the lower level of separation of the nanoclay layers. Thus intercalated nanocomposites would be formed [36]. - 159 -

Chapter 3 - Epoxy nanocomposites

With increased nanoclay content, the compressive modulus of the nanocomposites increased progressively over that of the neat resin, Figure 3. 23, with a 21 percent increase in the modulus being achieved with 8.4 phr nanoclay content. The present findings are similar to those of Becker et al. [37], in which a monotonic increase in the modulus with increasing organoclay loadings was observed. A 20 percent increase for an organoclay content of 10 wt% was achieved. Yasmin et al. [38] also found that the modulus of DGEBA nanoclay composites increased linearly w ith increasing clay content. They considered that the improvement in modulus was due to the restricted mobility of the polymer chains brought about by good dispersion of the nanosize clay particles and good interfacial adhesion between the particles and the epoxy matrix . The orientation of the silicate layers and polymer chains with respect to the loading direction could also contribute to the reinforcement effect. 3.4.2.5. Optical properties The optical transparency of the Ampreg 22 resin was increased with the addition of nanoparticles, Figure 3.24. Transparency is to be expected when the level of dispersion of the nanoclay is on a scale smaller than the wavelength of visible light [15]. None the less, the level of light transmission is generally found to decrease, rather than increase, when nanoclay is added as shown in Figure 3.50 [29].

Fig . 3.50

Light transmission spectra of pure epoxy, Ep/CM nanocomposites w ith

different CM contents [29].

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The Ampreg 22 resin was found to contain submicron sized particles of an aluminium rich material, presumable alumina. As noted earlier it would appear that these particles must agglomerate somewhat in the neat resin, thereby reducing transparency, but become better dispersed when the nanoclay is added. 3.4.2.6. Conclusions •

Selection of a suitable surfactant for modification of the nanoclay particles is important for epoxy resin nanocomposites. The longer chain surfactant hexadecylamine and octadecylamine were found to give the best results.

•

For a 2.5 phr addition of nanoclay the level of exfoliation and the compression modulus of the nanocomposites was insensitive to the mixing conditions over the range examined.

•

The modulus of the nanocomposites was reduced when the cure temperature was lower than 80ºC. This is attributed to insufficient diffusion of the resin into the nanoclay gallery.

•

The modulus of the nanocomposites increased progressively with clay content. This is attributed to the disperse d clay layers restricting the mobility of the polymer chains.

•

Exfoliated nanocomposites with an interlayer spacing of 100 Å were obtained up to a loading of 5 phr nanoclay, but the composites obtained at 8.4 phr were intercalated only with a layer spacing of only 63 Å.

3.4.3. TGDDM Nanocomposites Based on the findings for the DGEBA resin, the commercially available octadecylamine modified nanoclay Nanomer® I.30E

was selected for the TGDDM epoxy

nanocomposites.

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3.4.3.1. Effect of surfactant on cure Both the onset and peak temperature of the curing reaction were gradually reduced with increased nanoclay content, Figure 3.26 and Table 3.16. A similar effect has been observed in previous studies [5, 23, 39]. The reduction in the cure temperature was attributed to the catalytic effect of octadecylamine cations on the epoxy-amine polymerization [23]. Increasing the nanoclay content provided more of these organic cations to the nanocomposite network, el ading to a gradual decrease in the energy required for epoxide ring opening polymerization during the curing process. 3.4.3.2. Effect of mixing conditions While the mixing conditions did not affect the composites over the range examined for the DGEBA resin, the TGDDM resin had a higher viscosity and so the mixing conditions were again examined. The work was carried out for both 2.5 and 7.5 phr clay loadings. 3.4. 3.2.1. Mixing speed No significant effect of varying the mixing speed from 10,000 rpm to 26,000 rpm was seen in either the interlayer spacing or the compression modulus for either of the clay loadings examined, Figures 3.27 and 3. 28. This indicates that both speeds produced a sufficiently strong shear force to separate the nanoclay particles at both the micro and nanoscale. 3.4. 3.2.2. Mixing temperature Even though the viscosity of the TGDDM resin was higher than that of the DGEBA resin, no change in either the interlayer spacing or the compression modulus was again observed when the mixing temperature was varied between 60 and 80ºC, Figure 3. 29 and 3. 30. Evidently the viscosity was sufficiently low even at the lowest temperature of 60°C to allow good diffusion of the resin into the clay galleries.

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3.4. 3.2.3. Mixing time For the 2.5 phr nanoclay content, both the interlayer spacing and the compression modulus were insensitive to the mixing time for the range examined (0.5-4 h), Figures 3.31 and 3.32. For the 7.5 phr nanoclay loading there was again no change in the interlayer spacing but the compression modulus increased progressively with time , with the modulus increasing by 8% when the time was increased from 1 to 4 hours. This indicates that a longer time of mixing is beneficial at higher clay loadings. Longer mixing times are expected to be required at higher clay loadings since there are more clay particles to be broken down to the nanoscale and this would therefore require longer time. 3.4.3.3. Effect of curing conditions 3.4. 3.3.1. Degassing time Increasing the degassing time above that required for visible bubbling to cease was examined for the 7.5 phr composites. An initial increase of 6% in the modulus was observed with 15-30 minutes additional degassing time, but this was then lost when the additional degassing time was further increased to 60 minutes, Figure 3.33. A lthough visible bubbling had ceased at the start of the extra degassing period it appears that some entrapped air was still being removed. Entrapped air produces voids in the composite. Thus reduction of the amount of entrapped air would reduce the level of voiding, thereby increasing the modulus [40]. The subsequent reduction in the modulus with longer degassing time is surprising. T his may, however, be due to excessive evaporation of volatiles from the resin. As noted earlier, this would raise the viscosity of the resin , making it more difficult to separate the nanoclay layers during curing. However no significant difference was seen in the interlayer spacing. Alternatively, loss of volatiles from the resin could affect the level of cure, reducing the cross link density, thereby reducing the modulus of the resin. The reduction in modulus arising from this source would oppose the increase in modulus produced by additional degassing.

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3.4. 3.3.2. Effect of curing temperature There was a general upward trend in the modulus with increases in both the initial curing tempe rature (Stage I) and the final curing temperature (Stage III), Table 3.17 and Figures 3. 35-3.36, however the differences were considered to be generally within the experimental error. Nonetheless an increase in the modulus might be expected to accompany an increase in the temperature of initial cure. Chen and Curliss [18, 41] , found that intragallery polymerization was favoured at the onset temperature of curing (the initial curing temperature used here). Thus increasing the temperature of initial cure could produce greater separation of the clay layers and a corresponding increase in the modulus. This follows since the nanoclay layers would be expanded more by the extra exothermal heat inside the gallery allowing them to better overcome the attractive electrostatic forces between the nanoclay sheets. As a result, the epoxy resin molecules could penetrate more readily into the gallery, pushing the nanosheets further apart. Increasing the temperature would also reduce the viscosity and increase the diffusion rate of the resin and curing agent into the nanoclay layers, both of which would again be beneficial for separation of the clay layers [32]. However no significant change in the interlayer spacing was observed in the present study from the WAXD spectra , Figure 3.37-3.39, indicating that the effect was minimal, if any. An improvement in modulus might also be expected from increasing the final cure temperature since this could produce a greater level of cross-linking in the cured network, and the matrix surrounding the layered silicate [18]. The magnitude of this effect w ould depend on the level of cross linking produced in the earlier stages (ie, Stages I and II). Consideration of the error bars in Figure 3.35 indicates however that this effect is of marginal, if any, statistical significance in the present study. 3.4. 3.3.3. Effect of curing time The results showed a trend of decreasing modulus with increased time at the initial curing temperature, no change with time at the intermediate curing temperature, and an increase with time at the final curing temperature. However, with the exception of the difference observed between the shortest (0 hr) and longest (2 hr) times at the final curing temperature, the changes were considered to be within the experimental error. - 164 -

Chapter 3 - Epoxy nanocomposites

Since the shortest time at each cure temperature was 0 hr, the results also provide information on the effect of omitting each of the three cure stages (ie, reducing the cure schedule from 3 stages to 2 stages). Omission of either the first or second stage had no significant effect on the modulus. However a significant increase in the modulus was obtained when the third stage was included. This confirms that additional cross linking takes place during this stage. This postcuring stage would not, however, affect the layer spacing, consistent with experimental findings. 3.4.3.4. Effect of nanoclay content Unlike the DGEBA resin where no more than 8.4 phr nanoclay could be added to the resin becoming excessively viscous, it was possible to add up to 20 phr nanoclay to the TGDDM resin. Fully exfoliated nanocomposites with an interlayer spacing of greater than 120 Å were obtained up to 5 phr nanoclay, while pre-exfolitaed or intercalated nanocomposites , with interlayer spacings of from 85 Å (7.5 phr nanoclay) to 60 Å (20 phr nanoclay), were obtained at higher loadings, Figure 3.44 and 3.45. The nanoclay platelets were present within spherical, or nearly spherical, particles approximately 100150 nm in diameter. These tended to agglomerate into larger particles, typically 400 nm in diameter, at the higher clay loadings. In all cases the particles (or agglomerates) were uniformly dispersed with almost all the platelets of nanoclay within a particle being parallel. According to Chen and Curliss [18] , low organoclay loadings generally show much larger interlayer spacings due to the comparatively larger volume of epoxy resin molecules available to diffuse into the clay gallery. In addition, the distance between neighbouring nanoclay particles is greater at low er nanoclay concentrations. Thus, when the gallery of the organoclay is expanded, the interaction between the adjacent organoclay clusters is less and expansion of the gallery can continue more freely, leading to full exfoliation of the nanoclay in the resin matrix. The size of the nanoclay clusters may also affect the ease of pre -intercalation and, to a lesser extent, that of exfoliation [26]. At low nanoclay contents, it is easier to separate the nanoclay particles in the resin so that smaller clusters are obtained. This allows the

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resin and curing agent to diffuse more readily to the centre of the clusters than is possible for the larger clusters produced at higher nanoclay contents. The compressive modulus increased in an essentially linear manner with increased nanoclay loading, Figure 3.46 and 3.47, with an increase of 50 % being achieved over that of the pristine resin at a nanoclay content of 20 phr. Becker [37] also reported a monotonic increase in the modulus with increased organoclay concentration. A 20 % increase in the modulus of TGDDM-DDS at 9 phr nanoclay was obtained in Liu et al.’s [42] research, although the increase was only half this (ie, 10 %) for untreated clay at the same content. While, overall, the modulus increased linearly with clay content the values for the two lowest clay loadings (1 and 2.5 phr) were anomalously high, Figure 3.47. It appears that the better exfoliation achieved at the low clay loadings provided more effective reinforcement for the volume fraction of clay used. This in turn would suggest that the level of cross linking within the resin-filled galleries between the clay layers was increased in the better exfoliated composites. It is noted that the wider the galleries between the clay layers, the easier would be the access of the hardener to the resin in the galleries. It may be that in the composites with thinner interlayer spacings (ie, those with high clay loadings) insufficient hardener was able to penetrate into the galleries to provide full cross-linking of the resin. Although only fully intercalated or partially exfoliated nanocomposites were obtained with nanoclay contents of more than 7.5 phr, the nanoclay layers still provided effective reinforcement of the resin network. An increase in stiffness is to be expected because of the higher modulus of clay compared to that of the matrix (170 GPa [43] compared with 2.35 GPa for the cured neat resin (Section 3.3.3.8). However the stiffness improvements brought about by the nanocomposite structure are known to be substantially higher than that of conventional micro-composites [44]. The mechanical property improvement, however, is not just related to the state of dispersion of the nanoclay at the nanometer scale , but also to the level of dispersion at the micro-scale [39]. Bassara et al. [21] found that while the modulus increase d initially with clay content in their nanocomposites, no further improvement was obtained at clay loadings greater - 166 -

Chapter 3 - Epoxy nanocomposites

than 8% . This was attributed to agglomeration decreas ing the polymer-clay surface interactions. The continual improvement in modulus with increased clay content obtained in the present study would appear to indicate that agglomeration had not occurred to the level obtained by Bassara et al. , w ith the clay being dispersed more uniformly throughout the resin matrix. It was noted that while the interlayer spacing decreased with increased clay loading the number of clay platelets in a stack increased. During the curing process, epoxy molecules migrate into the spaces between the nanoparticles detaching them at a microscale prior to migrating into the nanoclay gallery. An adequate amount of epoxy resin is required to make the nanosheets totally separated. With increasing nanoclay content the volume of resin in the mixture decreases correspondingly. Thus there is less resin available to diffuse into the clay galleries during the swelling period. Consequently, there would be more layers of nanoclay closely stacked together, forming more clumps or nanoparticle clusters in the resin matrix [45]. A gradual linear decrease in Tg was observed with increasing nanoclay content, as illustrated in Figure 3.48. The decrease in T g of the nanocomposites is attributed to the plasticizing effect of the octadecylamine surfactant, an aliphatic amine with long flexible linear chains [44]. Since the concentration of surfactant in the resin would increase with increasing nanoclay content, T g would be expected to decrease proportionately, as observed here. Additionally, a reduction in Tg can arise in nanocomposites from a combination of several other factors, such as unreacted resin plasticization [37, 46], a decrease in the cross-link density of the epoxy network [44] , and a lack of surrounding molecular entanglements [37, 44, 46]. The variation in the compression modulus with nanoclay content was compared with the predictions of the Halpin-Tsai model. According to this model, the longitudinal and transverse moduli EL and ET are given by [47, 48].

EL = Em

1 + η L ζVc 1 − η LVc

(1)

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Chapter 3 - Epoxy nanocomposites

ET = E m

1 + ηT Vc 1 − η T Vc

(2)

where Ec −1 Em ηL = Ec +ζ Em

(3)

Ec −1 Em ηT = Ec +2 Em

(4)

and V is the volume fraction and ζ is the shape parameter. The subscripts of L, T, m, and c refer to the longitudinal direction, the transverse direction, the matrix and the nanoclay, respectively. For the resin and clay used here E m = 2.35 GPa (Section 3.3.3.8) and E c = 170 GPa [43]. According to Van-Es t he shape parameter ζ is given by [48] :

2 2l ζ = αc = c 3 3 tc

(5)

where a c, lc and tc refer to the aspect ratio, length and thickness of the exfoliated clay nanoplatelets. For a randomly -oriented nanocomposite the elastic modulus En can be obtained from the longitudinal and transverse moduli using [48] :

3 5 En = EL + ET 8 8

(6)

- 168 -

Chapter 3 - Epoxy nanocomposites

While a fibre reinforces only in the longitudinal direction (ie, along the fibre) a platelet can reinforce in all directions parallel to the platelet surface [15, 49] . Nanoclay fillers, which may be in the form of platelets or platelet stacks (particles), are often simulated by squares or circular discs with varying thickness. The shape parameter ζ can be calculated from the average length lc and thickness t c of the platelet stacks. From the SEM micrographs shown in Figure 3.49 the nanoclay particles appeared to be circular discs with diameters of 100 - 400 nm, depending on the nanoclay content. The thickness of the stacks also changed from 60 to 100 nm. Using these values gives lc/tc= 2 – 4. While for intercalated nanocomposites, the aspect ratio of the stacks is the length versus the thickness of the stacks, for fully exfoliated nanocomposites t c = tp = 1 nm (tp = thickness of a platelet), so the aspect ratio would then be 100-400 depending on the length of the nanoclay platelets. In the present study full exfoliation was only achieved at the very low clay loadings for which lc = ~100 giving an aspect ratio of ~100. Considering for the moment only the intercalated composites (ie, those with clay loadings of 5 phr or higher), the d-spacing of the nanoclay particles was 6-8 nm depending on the nanoclay content, Figure 3. 45. It was also noticed that the number of platelets in the stack, and also the thickness of the stack, increased with reduced dspacing (ie, with increased nanoclay content). This relationship is illustrated by the following equation: t c = (n − 1)d 001 + t p

where tc = thickness of the stack n = number of platelets (= 10-15)

d 001 = 6 − 8 nm - 169 -

(7)

Chapter 3 - Epoxy nanocomposites

t p = 1 nm

The predicted modulus values calculated using Equation 6 are compared with the experimentally obtained values for various values of the aspect ratio a c = l c/tc in Figure 3.51. The experimental results agree well with the predicted values for an lc/tc ratio of 13 which is in reasonable agreement with the experimental value of 2-4. As noted earlier the modulus values for the 1 and 2.5 phr nanocomposites were anomalously high. The value for 1 phr fits reasonably well to an l c/tc ratio of 100, Figure 3.51, consistent with the value of lc/tc determined above for a fully exfoliated composite, but the result for the 2.5 phr corresponds to a much lower value of l c/t c. However, examination of Figure 3.51 indicated that the clay platelets did have a tendency to stack which would then result in a reduction in the aspect ratio, l c/tc. 6

Young's Modulus E (GPa)

5

4

3

lc/tc=100

2

lc/tc=40 lc/tc=13 1

Experimental

0 0

2

4

6

8

10

Nanoclay content (vol.%)

Fig. 3.5 1 Comparison of experimental results with predicted results for aspect ratios of 13, 40 and 100. 3.4.3.5. Conclusions •

For the clay loadings of 2.5 and 7.5 phr examined the compression modulus and interlayer spacing was insensitive to the mixing speed and mixing temperature over

- 170 -

Chapter 3 - Epoxy nanocomposites

the ranges examined. However longer mixing times were found to beneficial at higher clay loadings. This is attributed to the increased number of clay particles that need to be broken down during mixing. •

Degassing the resin-nanoclay mixture after stirring for an extra 30 minutes beyond the time required for visible gas removal to cease was found to be beneficial. Longer additional degassing times were found to be detrimental. This is attributed to excessive loss of volatiles from the resin.

•

There was no significant effect of curing temperature within the range examined for the three stage cure cycle.

•

Varying the cure time of the first two stages of the cure cycle produced no significant effect. Indeed, either of these steps could have been omitted without causing any significant change. This indicates that both the temperatures used were effective for separation of the nanoclay layers. The modulus was however increased by the inclusion of the final postcuring stage. This confirms that additional cross linking takes place during this stage.

•

Up to 20 phr nanoclay could be added to the TGDDM resin without the mixture becoming excessively viscous. Fully exfoliated nanocomposites with an interlayer spacing of greater than 120 Å were obtained up to 5 phr nanoclay, while preexfolitaed or intercalated nanocomposites, with interlayer spacings of from 85 Å (7.5 phr nanoclay) to 60 Å (20 phr nanoclay), were obtained at higher loadings.

•

For clay loadings of 5 phr and above, the compression modulus increased linearly with clay content, with an increase of 50% being achieved over that of the pristine clay at 20 phr clay. The results were in good agreement with the predictions of the Halpin -Tsai model for an aspect ratio of 13.

•

The modulus was anomalously high for the 1 and 2.5 phr nanocomposites. This is attributed to the better exfoliation achieved at these clay loadings.

- 171 -

Chapter 3 - Epoxy nanocomposites

3.5. CONCLUSIONS The conclusions are summarised below. Nanoclay modification •

Alkylammonium surfactants with short alkyl chains (8 carbon atoms) amines did not expand the layers in the nanoclay appreciably.

•

The longer alkyl chain (16 carbon atoms) surfactant produced a substantial increase in the interlayer spacing and is considered suitable for modification of the nanoc lay for reinforcing epoxy resin

•

Within the ranges examined, the acid/amine concentration, surfactant concentration and mixing time had no appreciable influence on the interlayer spacing.

•

Surfactant modified nanoparticles with acid/amine ratios less than 1:1 proved difficult to centrifuge.

DGEBA nanocomposites •

Selection of a suitable surfactant for modification of the nanoclay particles is important for epoxy resin nanocomposites. The longer chain surfactant hexadecylamine and octadecylamine were found to give the best results.

•

For a 2.5 phr addition of nanoclay the level of exfoliation and the compression modulus of the nanocomposites was insensitive to the mixing conditions over the range examined.

•

The modulus of the nanocomposites was reduced when the cure temperature was lower than 80ºC. This is attributed to insufficient diffusion of the resin into the nanoclay gallery.

- 172 -

Chapter 3 - Epoxy nanocomposites

•

The modulus of the nanocomposites increased progressively with clay content. This is attributed to the dispersed clay layers restricting the mobility of the polymer chains.

•

Exfoliated nanocomposites with an interlayer spacing of 100 Å were obtained up to a loading of 5 phr nanoclay, but the composites obtained at 8.4 phr were intercalated only with a layer spacing of only 63 Å.

TGDDM nanocomposites •

For the clay loadings of 2.5 and 75 phr examined the compression modulus and interlayer spacing was insensitive to the mixing speed and mixing temperature over the ranges examined. However longer mixing times were found to beneficial at higher clay loadings. This is attributed to the increased number of clay particles that need to be broken down during mixing.

•

Degassing the resin-nanoclay mixture after stirring for an extra 30 minutes beyond

•

The time required for visible gas removal to cease was found to be beneficial. Longer additional degassing times were found to be detrimental. This is attributed to excessive loss of volatiles from the resin.

•

There was no significant effect of curing temperature within the range examined for the three stage cure cycle.

•

Varying the cure time of the first two stages of the cure cycle produced no significant effect. Indeed, either of these steps could have been omitted without causing any significant change. This indicates that both the temperatures used were effective for separation of the nanoclay layers. The modulus was however increased by the inclusion of the final postcuring stage. This confirms that additional cross linking takes place during this stage.

•

Up to 20 phr nanoclay could be added to the TGDDM resin without the mixture becoming excessively viscous. Fully exfoliated nanocomposites with an interlayer - 173 -

Chapter 3 - Epoxy nanocomposites

spacing of greater than 120 Å were obtained up to 5 phr nanoclay, while preexfolitaed or intercalated nanocomposites, with interlayer spacings of from 85 Å (7.5 phr nanoclay) to 60 Å (20 phr nanoclay), were obtained at higher loadings. •

For clay loadings of 5 phr and above, the compression modulus increased linearly with clay content, with an increase of 50% being achieved over that of the pristine clay at 20 phr clay. The results were in good agreement with the predictions of the Halpin -Tsai model for an aspect ratio of 13.

•

The modulus was anomalously high for the 1 and 2.5 phr nanocomposites. This is attributed to the better exfoliation achieved at these clay loadings.

- 174 -

Chapter 3 - Epoxy nanocomposites

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Chapter 3 - Epoxy nanocomposites

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Chapter 3 - Epoxy nanocomposites

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D. Chen and P. He, "Monitoring the curing process of epoxy resin nanocomposites based on organo-montmorillonite - a new application of resin curemeter ", Composites Science and Technology, 64 [16] 2501-2507 (2004). M.-M. Shen, M.-G. Lu, Y.-L. Chen, and C.-Y. Ha, "Nanocomposites based on liquid -crystalline epoxy-clay: synthesis and morphology", Polymer International, 54 [8] 1163-1168 (2005). O. Becker, R. Varley, and G. Simon, "Morphology, thermal relaxations and mechanical properties of layered silicate nanocomposites based upon high functionality epoxy resins", Polymer, 43 [16] 4365-4373 (2002). A. Yasmin, J.J. Luo, J.L. Abot, and I.M. Daniel, "Mechanical and thermal behavior of clay/epoxy nanocomposites ", Composites Science and Technology, 66 [14] 2415-2422 (2006). L. Le Pluart, J. Duchet, and H. Sautereau, "Epoxy/montmorillonite nanocomposites: influence of organophilic treatment on reactivity, morphology and fracture properties ", Polymer, 46 [26] 12267-12278 (2005). M.-W. Ho, C.-K. Lam, K.-t. Lau, D.H.L. Ng, and D. Hui, "Mechanical properties of epoxy-based composites using nanoclays", Composite Structures, 75 [1-4] 415-421 (2006). D.C. Chenggang Chen, "Processing and morphological development of montmorillonite epoxy nanocomposites", Nanotechnology, 14 [6] 643-648 (2003). W. Liu, S.V. Hoa, and M. Pugh, "Organoclay-modified high performance epoxy nanocomposites", Composites Science and Technology, 65 [2] 307-316 (2005). A.C.D. Newman, Chemistry of clays and clay minerals. Longman Scientific and Technical, 1987. X. Kornmann, R. Thomann, R. Mülhaupt, J. Finter, and L.A. Berglund, "High Performance Epoxy-Layered Silicate Nanocomposites", Polymer Engineering and Science, 42 [9] 1815-1826 (2002). X. Kornmann, H. Lindberg, and L.A. Berglund, "Synthesis of epoxy-clay nanocomposites: influence of the nature of the clay on structure", Polymer, 42 [4] 1303-1310 (2001). V. Nigam, D.K. Setua, G.N. Mathur, and K.K. Kar, "Epoxy-Montmorillonite Clay Nanocomposites: Synthesis and Characterization ", Journal of Applied Polymer Science, 93 [5] 2201-2210 (2004). H. Miyagawa, M.J. Rich, and L.T. Drzal, "Amine-cured epoxy/clay nanocomposites. II. The effect of the nanoclay aspect ratio", Journal of Polymer Science: PartB, 42 [23] 4391-4400 (2004). H. Miyagawa, K.H. Foo, I.M. Daniel, and L.T. Drzal, "Mechanical properties and failure surface morphology of amine-cured epoxy/clay nanocomposites", Journal of Applied Polymer Science, 96 [2] 281-287 (2005). X. Li, H. Gao, W.A. Scrivens, D. Fei, V. Thakur, M.A. Sutton, A.P. Reynolds, and M.L. Myrick, "Structural and mechanical characterization of nanoclayreinfo rced agarose nanocomposites ", Nanotechnology, 16 [10] 2020-2029 (2005). C.-K. Lam, H.-y. Cheung, K.-t. Lau, L.-m. Zhou, M.-w. Ho, and D. Hui, "Cluster size effect in hardness of nanoclay/epoxy composites", Composites: Part B, 36 [3] 263-269 (2005).

- 177 -

Chapter 4 - Fibre laminated nanocomposites

CHAPTER 4

FIBRE LAMINATED NANOCOMPOSITES

4 .1. INTRODUCTION Compressive failure of composites occurs by fibre microbuckling which leads to kink formation in the 0° plies [1]. Numerical analysis by Drapier and Wisnom [2] and Paris et al [3] showed that this takes place as a result of shear instability of the fibre tows. As a result, the compressive strength is proportional to the matrix elastic-plastic shear modulus (i.e, the stiff ness of the matrix) [4, 5]. Since there is a strong correlation between compression and bearing behaviour [6] , it is expected that increasing the shear modulus of the matrix should also improve the bearing strength. In the previous chapter it was shown that nanoclay particles could be successfully introduced into epoxy resin and that this increased the compressive modulus of the resin by up to 50%. Since the shear and compressive moduli are interrelated it would be expected that the use of nanoclay reinforced resin as the matrix could improve the bearing performance of fibre reinforced composites. Laminated carbon fibre composites were fabricated using nanoclay/TGDDM epoxy nanocomposite as the matrix resin. The bearing strength of the laminates was determined using the pin-contact bearing test developed by Wu and Sun [1]. The failure mechanism was examined using optical and scanning electron microscopy. The dspacing of the nanoclay was measured in the uncured prepreg plies, as well as in the cured composites, using wide angle x-ray diffraction (WAXD).

- 178 -

Chapter 4 - Fibre laminated nanocomposites

4.2. EXPERIMENTAL PROCEDURE 4.2.1. Materials Matrix Araldite® LY568 TGGDM epoxy resin reinforced with Nanomer® I.30E nanoclay and cured with Ethacure 100 DETDA hardener was used as the matrix resin. Details of these materials have already been given in Section 3.2.3.1, while the nanocomposites produced from these materials have been discussed in Section 3.3.3. Carbon fibre Two carbon fibre fabrics were used in this study. Both fabrics were made from T300 carbon fibres. The first fabric was a 5-harness satin weave fabric made from 6K tows and having an aerial weight of 370 grams per square metre. The second was a plain weave fabric made from 3K tows with an aerial weight of 200 grams per square metre. 4 .2.2. Sample Preparation The composites were fabricated using a vacuum assisted prepreg method. The carbon fibre fabric was cut into pieces 150 x 150 mm. The TGDDM / DETDA resin, with and without I30E nanoclay particles, was prepared as described in Section 3.3.3 and preheated to 70°C to reduce the viscosity. The resin was then applied to the fabric using a brush and a roller and the impregnated fabric dried at room temperature in the laboratory, with the fabric hanging vertically, for a period of one week, as shown in Figure 4.1.

- 179 -

Chapter 4 - Fibre laminated nanocomposites

Wet prepreg plies

Fig. 4.1 Wet prepreg plies drying in laboratory. After drying, the prepreg plies were trimmed and then laid up on an aluminium plate, over a layer of FEP release agent, to produce the required thickness of the panel.

(a) Layed-up prepreg

Aluminium plate

(b)

Fig. 4.2 Prepreg lay-up (a) and sealed vacuum bag (b). - 180 -

Chapter 4 - Fibre laminated nanocomposites

20 plies were applied for the 3k tow fabric and 12 plies for the 6k tow fabric. A layer of perforated FEP followed by a layer of breather cloth was placed over the lay-up which was then vacuum bagged at a pressure of 90 kPa, Figures 4.2 and 4.3. After vacuum bagging the lay up was placed in a Carver 10 tonne hot press, with the vacuum still applied, Figure 4.3 (b), and cured at a pressure of 700 kPa using a cure schedule of 2 hours at 110°C, then 1 hour at 155°C, followed by 1 hour at 200°C .

(a)

(b)

Fig. 4.3

Vacuum bagging of lay up (a) and layup being inserted into hot press (b).

- 181 -

Chapter 4 - Fibre laminated nanocomposites

4.2.3. Testing and analysis 4.2.3.1. Pin-contact bearing test Samples 40 x 80 mm were cut from the cured panels using a water-lubricated diamond saw. A 10mm diameter hole was precision drilled in the centre of each of the samples. A schematic of the samples is shown in Figure 4.4 (a). Each drilled sample was sectioned through the centre line of the hole using a water -lubricated slitting wheel so as to divide it into two test pieces, each containing a semi-circular hole, as illustrated in Figure 4.4 (b). The holes were then polished using emery paper (1200 then 4000 grit SiC ) on a rotating mandrel inserted into a drill press. This removed approximately 0.1 mm from the hole surface.

(a)

(b)

Fig. 4.4 Procedure for producing pin -contact bearing test specimens: (a) specimen with drilled hole, (b) testpiece produced by slitting specimen through hole. Testing was conducted using the pin-contact bearing test devised by Wu and Sun [1]. The test pieces were mounted vertically in a fixture with two steel supporting plates clamped either side of the laminate to prevent global buckling, as shown in Figure 4.5. - 182 -

Chapter 4 - Fibre laminated nanocomposites

The support plates extended to within 10 mm of the top of the specimen (the free span), ie, to within 5 mm of the bottom of the hole. A vertical load was applied through a 10 mm diameter horizontal steel pin using an Instron 1185 electromechanical testing machine with an MTS upgrade at a cross-head speed of 1 mm/min. The loading was applied until ultimate failure occurred. Six replicate samples were tested for each condition examined. Prior to testing the thickness of each specimen was measured at the positions marked B, C and D in Figure 4.6. The fibre volume fraction of the laminates was determined by matrix digestion.

Fig. 4.5 Schematic of the pin-contact bearing test employed by Wu and Sun [1].

C

B

D

Fig. 4 .6 Diagram showing positions where thickness was measured on specimens.

- 183 -

Chapter 4 - Fibre laminated nanocomposites

4.2.3.2. Microstructure The damage produced during bearing failure was examined, using optical microscopy, in longitudinal sections perpendicular to the laminate cut through the base of the hole. The specimens were prepared by first reinforcing the surface of the hole and the surrounding region with a layer of room temperature curing epoxy resin. The test piece was then cut vertically through the base of the hole using a water-lubricated slitting wheel. The cut was made about 0.5 mm off the centre of the hole so that, after grinding and polishing of the cut face, the section would be essentially through the base of the hole. A sample 10 mm long along the cut face of the section and 5 mm wide was removed from the sectioned test piece and mounted on the cut face using room temperature curing epoxy resin. The mounted specimens were gr ound using 1200 then 4000 grit SiC papers and then polished using 3µm then 1 µm diamond paste. The samples were examined using a Nikon EPIPHOT 200 light microscope. Images were recorded at 10-100x magnification using a Nikon digital camera fitted with Nikon ACT-1 software. 4.2.3.3. Fracture surfaces

Fig. 4.7 Production of Mode 1 fracture surface. - 184 -

Chapter 4 - Fibre laminated nanocomposites

Mode 1 fracture surfaces were produced by driving a metal wedge into the end of samples 40 x 12 mm as shown in Figure 4.7. A slot was first cut into the end of the sample to receive the wedge. The wedge was driven into the sample using an Instron 1185 testing machine at a cross-head speed of 1 mm/min. The fracture surfaces obtained were coated with evaporated chromium and examined using a Hitachi S4500 SEM at an excitation voltage of 20 kV. 4.2.3.4. Wide angle X-ray diffraction Rectangular samples 5 x 12 mm were cut from the cured laminates using a waterlubricated slitting wheel and then polished using 1200 grade SiC paper to produce a smooth surface. Samples were also cut from the uncured prepreg after drying at room temperature for week. The samples were examined by wide angle X-ray diffraction using the procedure described in Section 3.2.2.3. 4 .3 . RESULTS 4.3.1. Pin-contact bearing test 25 Baseline Nanoclay

Load (kN)

20

15

10

5

0 0

0.5

1

1.5

Displacement (mm)

Fig . 4.8

Typical pin-contact load-displacement curves for laminates made with neat

resin (baseline) and nanoclay reinforced resin. - 185 -

Chapter 4 - Fibre laminated nanocomposites

Typical load displacement curves for test pieces made with and without (baseline) nanoclay reinforcement of the matrix are shown in Figure 4.8 while the individual curves are given in Appendix B. Similar curves were obtained for both the baseline and nanoclay reinforced laminates with the curve rising linearly, after an initial settling in period, to the point of ultimate failure. A similar result was obtained by Wu and Sun [1]. a) 3K tow carbon fibres (3K CFRP) The measured values of thickness, failure load, bearing strength, slope to failure (slope in linear region of load displacement curve) and crosshead displacement at failure for the test pieces made with neat epoxy are shown in Table 4.1. The fibre volume fraction was 56.9 vol.%. The bearing strength of these baseline samples was 485.3 ± 18.2 MPa with the slope to failure being 21.4 ± 2.1 kN/mm. The displacement at failure was 0.95 ± 0.10 mm. Table 4.1 Measured values of thickness, failure load, bearing strength, slope to failure and displacement to failure for neat epoxy laminates (baseline) containing 56.9 vol.% carbon fibre.

Specimen

1

2

3

4

5

6

B

3.95

3.87

3.77

3.85

3.83

3.96

Thickness C (mm) D

3.92

3.87

3.79

3.81

3.81

3.93

3.95

3.91

3.83

3.81

3.9

3.99

3.94

3.88

3.80

3.82

3.85

Bearing area (mm2)

39.40

38.83

37.97

38.23

Failure load (kN)

19.25

19.33

18.82

Bearing strength (MPa)

488.58

497.77

Slope to failure (kN/mm)

19.41

Displacement a t failure (mm)

1.04

Ave

Ave

STD

3.96

3.88

0.07

38.47

39.60

38.75

0.65

17.2

19.08

19.15

18.805

0.81

495.70

449.87

496.01

483.59

485.25

18.15

22.25

23.03

21.17

23.87

18.59

21.39

2.07

0.95

0.89

0.86

0.85

1.08

0.95

0.10

- 186 -

Chapter 4 - Fibre laminated nanocomposites

Table 4.2 shows the results obtained from the pin-contact bearing test for the composite s reinforced with 7.5 phr nanoclay. The specimen thickness was slightly lower (4%) than for the baseline samples while the fibre volume fraction was slightly higher (57.6% compared with 56.9%). The bearing strength was similar to that of the baseline samples being 472.2 ± 35.8 MPa. However the slope to failure was higher being 24.0 ± 1.0 kN/mm while the displacement to failure was smaller at 0.80 ± 0.07 mm. Table 4.2 Measured values of thickness, failure load, bearing strength, slope to failure and displacement to failure for laminates with 7.5 phr I30E nanoclay and 57.6 vol.% carbon fibre.

Specimen

1

2

3

4

5

6

3.73

3.72

3.70

3.76

3.72

3.74

3.69

3.69

3.69

3.76

3.69

3.73

3.70

3.67

3.68

3.78

3.67

3.73

3.71

3.69

3.69

3.77

3.69

Bearing area 2 (mm )

37.07

36.93

36.90

37.67

Failure load (kN)

15.49

18.24

16.65

Bearing strength (MPa)

417.90

493.86

Slope to failure (kN/mm)

24.25

Displacement a t failure (mm)

0.7

B Thickness C (mm) D Ave

Ave

STD

3.73

3.71

0.03

36.93

37.33

37.14

0.30

19.48

18.16

17.21

17.54

1.40

451.22

517.17

491.70

460.98

472.14

35.74

24.19

22.29

23.61

25.14

24.20

23.95

0.95

0.82

0.82

0.91

0.78

0.77

0.80

0.07

In view of the absence of any improvement in bearing strength produced by the nanoclay reinforced matrix, de spite the increased bearing stiffness (as evidenced by the increase in the slope to failure), a second set of samples was prepared. It was however difficult to obtain exactly the same fibre volume fraction since the thickness of the laminates (which determines the volume fraction) was found to be very sensitive to the applied pressure in the hot press and this was difficult to control to the required accuracy with the equipment available. The second set of specimens had a fibre volume fraction of 55%. The results from the pin-contact test are given in Table 4.3.

- 187 -

Chapter 4 - Fibre laminated nanocomposites

Table 4.3 Measured values of thickness, failure load, bearing strength, slope to failure and displacement to failure for laminates with 7.5 phr I30E nanoclay and 55 vol.% carbon fibre.

Specimen

1

2

3

4

5

6

4.16

4.09

4.12

4.12

4.07

4.10

4.16

4.08

4.16

4.13

4.08

4.10

4.17

4.12

4.15

4.12

4.12

4.12

4.16

4.10

4.14

4.12

4.09

Bearing area 2 (mm )

41.63

40.97

41.43

41.23

Failure load (kN)

19.53

19.55

20.17

Bearing strength (MPa)

469.10

477.22

Slope to failure (kN/mm)

20.15

Displacement a t failure (mm)

1.03

B Thickness C (mm) D Ave

Ave

STD

4.11

4.12

0.03

40.90

41.07

41.21

0.28

19.85

19.45

18.88

19.57

0.43

486.81

481.41

475.55

459.74

474.97

9.52

18.59

26.47

19.56

25.55

25.18

22.58

3.51

1.1

0.83

1.08

0.84

0.83

0.95

0.13

The measured bearing strength was a very similar to that obtained with the first set of specimens (475.0 ± 9.5 MPa compared with 472.1 ± 35.8 MPa) indicating the validity of the earlier result bearing strength of the reinforced laminate was 474.97 ± 9.52 MPa. Samples with an increased nanoclay content of 12.5 phr nanoclay were then prepared and tested. In this case the fibre volume fraction was 51.7%. The test results are shown in Table 4.4. Again no improvement in bearing strength was obtained, with a value of 472.9 ± 11.9 MPa being recorded. However, the slope to failure was increased substantially (23%) over that of the baseline samples, while the displacement to failure was again reduced (0.83 ± 0.08 compared with 0.95 ± 0.10). The results for all the composites fabricated from the 3k carbon fibre fabric are summarized in Table 4.5. While no improvement in bearing strength was obtained, the results showed a progressive increase in bearing stiffness (slope to failure) with nanoclay addition (10% at 7.5 phr and 23% at 12.5 phr nanoclay) with a general reduction in the strain to failure (displacement to failure).

- 188 -

Chapter 4 - Fibre laminated nanocomposites

Table 4.4 Measured values of thickness, failure load, bearing strength, slope to failure and displacement to failure for laminates with 7.5 phr I30E nanoclay and 51.7 vol.% carbon fibre.

Specimen

1

2

3

4.16

4.14

4.18

4.2

4.20

4.17

4.20

4.12

4.15

4.18

4.17

4.17

4.17

4.16

4.21

4.2

4.17

4.15

4.18

4.14

4.18

4.19

4.18

Bearing area 2 (mm )

41.77

41.40

41.80

41.93

Failure load (kN)

19.86

19.53

19.91

Bearing strength (MPa)

475.50

471.74

Slope to failure (kN/mm)

28.35

Displacement a t failure (mm)

0.77

B Thickness C (mm) D Ave

Table 4.5

4

5

6

Ave

STD

4.16

4.17

0.02

41.80

41.63

41.72

0.18

19.36

20.59

19.13

19.73

0.51

476.32

461.69

492.58

459.49

472.89

11.94

29.05

23.12

26.56

25.79

24.53

26.23

2.25

0.73

0.95

0.80

0.88

0.86

0.83

0.08

Summary of results for baseline and nanocomposite laminates (3K tow

fabric ).

Nanoclay

Fibre

Thickness

Bearing

Slope to failure

Displacement

content

volume

(mm)

strength (MPa)

(kN/mm)

at failure (mm)

(phr)

(vol.%)

0

56.9

3.88 ± 0.02

485.25 ± 18.15

21.39 ± 2.07

0.95 ± 0.10

7.5

57.6

3.71 ± 0.03

472.1 4 ± 35.7 4

23.95 ± 0.95

0.80 ± 0.07

7.5

55.0

4.12 ± 0.03

474.97 ± 09.52

22.58 ± 3.51

0.95 ± 0.13

12.5

51.7

4.17 ± 0.02

472.89 ± 11.94

26 .23 ± 2.25

0.83 ± 0.08

- 189 -

Chapter 4 - Fibre laminated nanocomposites

b) 6K tow carbon fibre (6K tow CFRP) A set of samples made from 6K satin weave carbon fibre fabric was also examined. The baseline results obtained using neat resin as the matrix resin are presented in Table 4.6 while those for the 7.5 phr nanoc lay reinforced resin are given in Table 4.7. Both sets of results are summarised in Table 4.8. The fibre volume fraction for the baseline specimens was 61.3 % while that for the nanocomposite reinforced laminates was 61. 1%. The measured bearing strength for the baseline samples was 452.6 ± 20.6 MPa. No significant improvement was obtained with the addition of 7.5 phr nanoclay for which a value of 462.2 ± 17.8 MPa was achieved. The bearing stiffness (slope to failure) was again increased by 11% (25.2 ± 2.4 kN/mm compared with 22.7 ± 3.3 kN/mm) by the addition of the nanoclay while the strain to failure (displacement at failure) was reduced (0.81 ± 0.14 mm compared with 0.87 ± 0.10 mm). Table 4.6 Measured values of thickness, failure load, bearing strength, slope to failure and displacement to failure for baseline (neat resin) laminates containing 61.3 vol.% carbon fibres.

Specimen

1

2

3

4

5

6

3.99

3.89

4.08

4.07

4

3.91

4.00

3.91

4.09

4.05

3.95

3.88

4.01

3.95

4.05

4.09

4

3.94

4.00

3.92

4.07

4.07

3.98

Bearing area 2 (mm )

40.00

39.17

40.73

40.70

Failure load (kN)

18.58

18.15

16.83

Bearing strength (MPa)

464.50

463.40

Slope to failure (kN/mm)

26.11

Displacement a t failure (mm)

0.77

B Thickness C (mm) D Ave

Ave

STD

3.91

3.99

0.07

39.83

39.10

39.92

0.71

18.31

18.12

18.37

18.06

0.63

413.18

449.88

454.90

469.82

452.61

20.60

22.72

26.35

20.26

22.82

17.72

22.66

3.34

0.85

0.69

0.98

0.85

1.07

0.87

0.14

- 190 -

Chapter 4 - Fibre laminated nanocomposites

Table 4.7 Measured values of thickness , failure load, bearing strength, slope to failure and displacement to failure for laminates with 7.5 phr I30E nanoclay and 61.1 vol.% carbon fibres.

Specimen

1

2

3

4

5

6

4.06

4.07

4.08

4.05

4.04

3.97

4.04

4.05

4.05

4.03

3.97

4.04

4.06

4.08

4.04

4.07

3.97

4.04

4.05

4.07

4.06

4.05

3.99

Bearing area 2 (mm )

40.53

40.67

40.57

40.50

Failure load (kN)

18.33

19.38

17.98

Bearing strength (MPa)

452.22

476.56

Slope to failure (kN/mm)

23.46

Displacement a t failure (mm)

0.85

B Thickness C (mm) D Ave

Table 4.8

Ave

STD

4.02

4.04

0.03

39.93

40.17

40.39

0.28

19.80

18.55

17.99

18.67

0.75

443.22

488.89

464.52

447.88

462.22

17.82

21.08

26.61

27.37

26.64

25.83

25.17

2.42

0.99

0.73

0.78

0.76

0.74

0.81

0.10

Summary of results for baseline and nanocomposite laminates (6K tow

fabric ).

Nanoclay

Fibre

Thickness

Bearing

Slope to failure

Displacement

content

volume

(mm)

strength (MPa)

(kN/mm)

at failure (mm)

(phr)

(vol.%)

0

61.3

3.99 ± 0.0 7

452.6 1 ± 20.6 0

22 .66 ± 3.34

0.87 ± 0.14

7.5

61.1

4.04 ± 0.0 3

462.2 2 ± 1 7.8 2

25.17 ± 2.42

0.81 ± 0.10

- 191 -

Chapter 4 - Fibre laminated nanocomposites

4.3.2. Laminate quality Post-mortem examination of failed specimens was conducted on metallographically polished sections cut through the base of the loaded hole perpendicular to the plane of the laminate in the direction of loading. First, however , the quality of the laminates was evaluated by examination of sections away from the failed region. a) 3K tow fibre Sections through the unreinforced laminates and those reinforced with 7.5 and 12.5 phr nanoclay are shown in Figures 4.9, 4.10 and 4.11 respectively. The laminates can be seen to be of high quality with the void level being generally very low. The highest void level (calculated from the ratio of the area of the voids to the surface area of the laminate in the field of view) of around 2% was observed in the 7.5 phr nanoclay laminate with a fibre volume fraction of 57.6%.

100µm

Fig. 4.9 Optical micrograph showing un-reinforced laminate.

- 192 -

Chapter 4 - Fibre laminated nanocomposites

100µm

Fig. 4.10 Optical micrograph showing laminate reinforced with 7.5 phr nanoclay (55 vol.% fibres).

100µm

Fig. 4.11 Optical micrograph showing laminate reinforced with 12.5 phr nanoclay. b) 6K tow carbon fibres Figure 4.12 shows a micrograph of the un-reinforced laminate made using the 6K tow fabric while a micrograph of the 7.5 phr nanoclay reinforced laminate is shown in Figure 4.13. Localised voids can be seen in both laminates but overall the void content was low , being less than 1.5% in both cases.

- 193 -

Chapter 4 - Fibre laminated nanocomposites

100µm

Fig. 4.12 Optical micrograph showing un-reinforced 6K fabric laminate.

100µm

Fig. 4.13

Optical micrograph showing 6K fabric laminate reinforced with 7.5 phr

nanoclay. 4.3. 3. Microstructural examination Microstructural examination of cross sections through the damage zone beneath the surface of the hole did not reveal any notable differences between the laminates made with and without nanoclay reinforcement nor between those made with the different tow fabrics. Several failure modes were observed, these being kink-band formation (K)

- 194 -

Chapter 4 - Fibre laminated nanocomposites

matrix cracking (M), shear cracking (S), and delamination (D), examples of which are shown in Figure 4.14. As can be seen in Figure 4.14 shear cracking is essentially a cooperative propagation of the kink bands through several adjoining plies.

Kinking (K)

Matrix cracking (M)

Shear cracking (S )

Fig. 4.14

Delamination (D)

Optical micrographs showing local failure modes in the laminated

composites: kinking (K), matrix cracking (M), shearing cracking (S), delamination (D). Representative micrographs of the bearing damage for all sample sets examined are shown in Figures 4.15 – 4.20. While there is considerable variability in the detail this was not outside the variability seen amongst samples of the same kind. The damage was characterised by the development of shear bands (shear cracking) emanating from the bearing surface at an angle of approximately 45°. These often (though not always) initiated at one or both of the free surfaces and propagated deep into the laminate to link - 195 -

Chapter 4 - Fibre laminated nanocomposites

up and then produce a longitudinal delamination. A typical example is seen in Figure 4.17. In some cases several shear bands formed and linked up to form delaminations as seen in Figure 4.20. Localised regions of matrix cracking and isolated kink bands were often seen. Waviness of the fibre bundles was also apparent, as is typical of woven fabric laminates.

(a)

S

M

D

(A)

M

1 mm

(b)

K

S

Fig. 4.15

Optical micrograph showing bearing damage in unreinforced laminate

containing 56.9 vol.% 3K carbon fibres (a) low magnification view and (b) enlarged view of region A in (a).

- 196 -

Chapter 4 - Fibre laminated nanocomposites

(a)

D

S

D

Void

(A) 1 mm M

K

(b)

M

Fig. 4.16 Optical micrograph showing bearing damage in laminate reinforced with 7.5 phr nanoclay containing 57.6 vol.% 3K carbon fibres (a) low magnification view and (b) enlarged view of region A in (a).

- 197 -

Chapter 4 - Fibre laminated nanocomposites

D

(a)

(A) 1 mm

(b) K

S

Fig. 4.17 Optical micrograph showing bearing damage in laminate reinforced with 7.5 phr nanoclay containing 55 vol.% 3K carbon fibres (a) low magnification view and (b) enlarged view of region A in (a).

- 198 -

Chapter 4 - Fibre laminated nanocomposites

(a) (A)

S 1 mm

(b) D

K

Fig. 4.18

Optical micrograph showing bearing damage in laminate reinforced with

12.5 phr nanoclay containing 51.7 vol.% 3K carbon fibres (a) low magnification view and (b) enlarged view of region A in (a).

- 199 -

Chapter 4 - Fibre laminated nanocomposites

b) 6K tow carbon fibre The 6k tow fabric composites showed similar types of bearing damage to those seen in the 3k tow material. Again there were no significant differences between the composites with and without nanoclay reinforcement of the matrix.

D

K

Void

S 1 mm

Fig. 4.19 Optical micrograph showing bearing damage in unreinforced laminate containing 61.3 vol.% of 6k carbon fibres.

- 200 -

Chapter 4 - Fibre laminated nanocomposites

(a) (A)

D

S

1 mm M

(b)

K

Fig. 4.20 Optical micrograph showing bearing damage in laminate reinforced with 7.5 phr nanoclay and containing 61.1 vol.% 6K carbon fibres (a) low magnification view and (b) enlarged view of region A in (a).

- 201 -

Chapter 4 - Fibre laminated nanocomposites

4.3.4. Fracture surfaces The Mode 1 (interlaminar) fracture surfaces of the 3K fabric laminates without and with the incorporation of nanoclay (7.5 phr) are shown in Figure 4.21. The fracture surfaces were produced using the procedure described in Section 4.2.3.3. The fracture surfaces can be seen to be quite different, with the baseline laminate being quite reflective while the 7.5 phr nanoclay laminate is matte in appearance.

Fig . 4.21

Mode 1 fracture surface of baseline (left) and 7.5 phr nanoclay reinforced

(right) laminate s. Scanning electron micrographs of the fracture surfaces are shown at two different magnifications in Figure 4. 22. The laminate made with neat resin as the matrix shows essentially clean separation of the fibres from the matrix with little or no matrix resin adhering. In contrast, substantial amounts of matrix are seen adhering to the fibres in the nanoclay reinforced laminate. Additionally , the matrix adhering to the fibres has a rough nodular appearance suggesting that failure has occurred around the cla y particles.

- 202 -

Chapter 4 - Fibre laminated nanocomposites

(a)

(b)

(c)

(d)

Fig. 4.22 SEM images of the fracture surfaces of unreinforced (a, b) and nanoclay reinforced (7.5 phr I30E) (c, d) laminates with 3K tow carbon fibres.

(a)

(b)

(c)

(d)

Fig. 4.23

SEM images of fracture surfaces of unreinforced (a, b) and nanoclay

reinforced (7.5 phr I30E) (c, d) laminates with 6K tow carbon fibres.

- 203 -

Chapter 4 - Fibre laminated nanocomposites

Similar features to those observed in the 3k fabric laminates were also observed in the 6k fabric laminates, as shown in Figure 4.23. Again the laminate with the neat resin matrix produced a Mode 1 fracture surface , in which the fibres were essentially devoid of any adhering matrix, while the nanoclay (again 7.5 phr) reinforced laminate had substantial quantities of rough nodular matrix adhering to the fibres. 4.3.5. Interlayer spacing Figure 4.24 gives the WAXD spectra for the 3K tow carbon fibre laminates with and without nanoclay reinforcement. No peaks are evident in the neat resin matrix baseline sample as would be expected in view of the absence of nanoclay. Sharp peaks are present for both the 7.5 phr and 12.5 phr nanoclay reinforced laminates corresponding to interlayer spacings of 42.8 Å and 41 Å, respectively, indicating that the clay had an intercalated morphology. 1400 1200 Baseline-3K

Intensity (counts)

1000

7.5 phr I30E 12.5phr I30E

800 600 400 200 0 1

2

3

4

5

6

7

2Theta (degree)

Fig. 4.24

WAXD spectra of unre inforced and reinforced (7.5 phr and 12.5 phr I30E

nanoclay) laminates with 3K tow carbon fibres. The WAXD spectrum of the 6K tow carbon fibre laminate reinforced with 7.5 phr nanoclay is shown in Figure 4.25. A clear peak is again evident corresponding to an

- 204 -

Chapter 4 - Fibre laminated nanocomposites

interlayer spacing of 38.5 Å in the nanoclay, again indicative of an intercalated nanocomposite matrix in the carbon fibre laminate.

1400 1200

Intensity (counts)

1000 Baseline

800

7.5 phr I30E 600 400 200 0 1

2

3

4

5

6

7

2Theta (degree)

Fig. 4.25

WAXD spectra of unreinforced and reinforced (7.5 phr I30E nanoclay)

laminates with 6K tow carbon fibres 250

Uncured prepreg - 1 week

Intensity(counts)

200

150

100

50

0 0

1

2

3

4

5

6

7

2Theta(degree)

Figure 4.26

WAXD spectra of uncured 3K tow carbon fibre prepreg containing 7.5

phr I30E nanoclay. - 205 -

Chapter 4 - Fibre laminated nanocomposites

The interlayer spacing was also measured in the prepreg before curing. The WAXD spectrum of the uncured 3k prepreg 7.5 phr nanoclay is presente d in Figure 4.26. A broad peak is present corresponding to an interlayer spacing of 41 Å which is only marginally less than that obtained in the cured sample. 4.4. DISCUSSION As noted earlier it was difficult to accurately control the laminate thickness us ing the prepreg method and this resulted in different fibre volume fractions. However, using the method of Siddiqui et al. [7], it was possible to normalize the results for this effect and, accordingly, the bearing strength was normalized to a fibre volume fraction of 57% for the 3k tow fabric laminates and 61% for 6k tow fabric laminates. The results are presented in Table 4. 9 and shown in Figure 4. 27. Table 4.9 Bearing strength of CFRPs with or without nanoclay reinforcement.

Nanoclay content in matrix (phr)

Carbon fibre Bearing strength content (vol.%) (MPa)

3K tow CFRP (normalized to 57 vol.%) 0 56.9 485.25 ± 18.15 7.5 57.6 472 .14 ± 35 .74 7.5 55.0 474.97 ± 09.52 12.5 51.7 472.89 ± 11.94 6K tow CFRP (normalized to 61 vol.%) 0 61.3 452 .61 ± 20 .60 7.5 61.1 462.22 ± 17.82

Normalized strength (MPa)

Percentage improvement (%)

486.10 ± 18.20 467 .22 ± 35 .37 492.24 ± 09.87 521.34 ± 13.16

---7.25

450 .39 ± 20 .50 461.46 ± 17.79

---

A slight improvement can be seen in the bearing strength at 7.5 phr for both fabrics examined, but thus not considered to be statistically significant according to the student t-test [8] Appendix C. However a more substantial improvement of 7% was obtained at 12.5 phr nanocla y content and this difference was found to be highly significant according to the student-t test, Appendix C.

- 206 -

Chapter 4 - Fibre laminated nanocomposites

600

Bearing strength (MPa)

500

400

300

200

100 0

7.5

7.5

12.5

0

7.5

Nanoclay content (phr)

Fig. 4.27 Normalised bearing strength from pin-contact test for laminates reinforced with 3K tow fabric (unfilled) and 6K tow fabric (filled) with various nanoclay contents (0, 7.5 and 12.5 phr). Error bars indicate one standard deviation. In view of the finding for the 12.5 phr nanoclay addition, and the increase observed at 7.5 phr (despite it being within the experimental scatter ) it is concluded that the stiffening of the matrix by nanoclay reinforcement does translate into increased bearing strength as hypothesized in this thesis on the basis of the findings of Wu and Sun [1]. It is however noted that the level of clay layer separation achieved in the laminates was substantially less than that achieved in the resin nanoclay composites described in Chapter 3 at both nanoclay loadings (42.8 Å at 7.5 phr and 41 Å at 12.5 phr compared to 85 Å and 75 Å respectively in the epoxy nanocomposites) indicating that the method used to produce laminates from the reinforced resin had in some way hindered exfoliation of the clay layers. According to Wu and Sun [1] the improvement in be aring strength is brought about by a stiffening of the matrix and this would also stiffen the composite during bearing loading resulting in a steepening of the stress-strain curve, and, correspondingly, in the loaddisplacement curve. To examine this, the slope of the linear region of the load displacement curve was determined and this has been given in the results in the previous

- 207 -

Chapter 4 - Fibre laminated nanocomposites

sections as the slope to failure. These results are shown again (without normalization) in Figure 4.28. Increases of 11 % for 7.5 phr nanoclay and ~22 % for 12.5 phr nanoclay were obtained, both of which were confirmed as being significant by the student t-test Appendix C. Concurrent with the increases in bearing strength and bearing stiffness was a decrease in the strain to failure (as measured and listed in the previous sections as displacement to failure), Figure 4.29. This indicates that the addition of nanoclay has caused the laminates to fail prematurely. This is also evident from the fracture surfaces since the mode of failure is dramatically different in the nanoclay reinforced composites to that in the baseline ones, Figures 4.21 and 4.22. 30

Slope to failure (kN/mm)

25

20

15

10

5

0 0

7.5

7.5

12.5

0

7.5

Nanoclay content (phr)

Fig. 4.28 Slope of load-displacement curves of unreinforced and reinforced laminates with 3k tow fabric (unfilled) and 6k tow carbon fibres (filled). Error bars indicate one standard deviation. The fracture surfaces in the nanoclay reinforced composites showed substantial amounts of nodular -appearing matrix adhering to the fibres while in the baseline unreinforced composites the fibres were essentially devoid of adhering matrix. This indicates that the fracture mechanism has changed from interfacial failure to matrix failure. In turn, this

- 208 -

Chapter 4 - Fibre laminated nanocomposites

indicates that the addition of the nanoclay has weakened the matrix relative to the fibre matrix interface

1.2

Displacement at failure (mm)

1

0.8

0.6

0.4

0.2

0 0

7.5

7.5

12.5

0

7.5

Nanoclay content (phr)

Fig. 4.29

Displacement at failure of unreinforced and reinforced laminates with 3K

tow carbon fibres (unfilled) and 6K tow carbon fibres (filled). Error bars indicate one standard deviation. The nodular appearance of the resin adhering to the fibres suggests that nanoclay agglomerates have been exploited during the fracture process. These would provide an easy crack path thus accountin g for the reduced strength of the matrix. The above argument implies that the presence of the nanoclay has led to a reduction in the toughness of the matrix. An estimate of the toughness can be obtained from the area under the stress strain curve which in turn is proportional to half the area obtained by multiplying the normalized bearing strength by the displacement at failure. These values are listed in Table 4.10 and show a general trend of decreasing toughness with addition of nanoclay. - 209 -

Chapter 4 - Fibre laminated nanocomposites

Table 4.10 Toughness estimated from area under stress strain curve. Nanoclay content (phr)

Normalised strength (MPa)

Displacement at failure (mm)

Toughness*

0

486

0.95

231

7.5

467

0.80

187

7.5

492

0.95

234

12.5

521

0.83

216

0

450

0.87

196

7.5

461

0.81

187

(MPa.mm)

3K tow fabric composites

6K tow fabric composites

* 0.5 x normalised strength x displacement at failure As noted in Chapter 2 the addition of nanoparticles to a polymer has the potential to produce simultaneous improvements in both stiffness and strength. However, the improvement in toughness is often not realized due to the nanoparticles not being appropriately dispersed at either (or both) the nano or microscale and this appears to have been the case here. If the inherent toughening due to nanopa rticle addition could be realized, premature matrix failure could then be avoided and an increase in bearing strength could then result. For example, if, by improving the toughness of the matrix, the same strain to failure could be achieved in the 12.5 phr nanoclay 3k tow composite as in the baseline 3k tow composite, the bearing strength would be increased to 596 MPa providing a 23% improvement. It is noted that the d-spacing of the nanoclay particles in the carbon fibre laminates was only about half that in the nanoclay-epoxy nanocomposites discussed in Chapter 3 (~40 ? compared with 70-80 ? ). However it was noted that the d-spacing in the prepreg after drying room temperature for seven days was also about 40 ? with no further increase taking place during curing. This indicates that the prepregging process had affected the ability of the nanoclay to exfoliate to the expected level in the carbon fibre laminates. It is considered that while the prepreg was held at room temperature there was sufficient time for the octadecylamine surfactant occupying the gallery of the nanoclay particles to catalyze the curing reaction within and around the clay layers before the laminates were cured. As a result , the clay layers were accidentally fixed (while almost all the - 210 -

Chapter 4 - Fibre laminated nanocomposites

remaining resin was still uncured), restricting subsequent expansion of the nanoclay layers when the laminates were cured. A similar effect was observed by Timmerman et al [9] when trying to fabricate nanoclay reinforced matrix laminate s using the prepreg method. The distance between the layers of nanoclay in the laminates was found to be 32-35 ? , similar to the result obtained here. With interlayer spacings of 40 ? the nanocomposites are intercalated rather than exfoliated. Zilg et al. [10] found that intercalated nanoclay promoted toughness of their epoxy resin matrix whereas exfoliated clay platelets principally improved stiffness. They attributed the improvement in toughness to energy-absorbing shearing of the intercalated nanoclay layers. The intercalated platelets were believed to promote crack pinning at the microscale. Becker et al. [11] reported that both toughness and stiffness had been improved through organoclay incorporation and attributed this to a blend of intercalated and exfoliated morphologies. The layered silicate with its high modulus contributed to the higher stiffness of the nanocomposite while the tactoids acted as stress concentrators and caused matrix yielding, thus improving toughness [12]. Miyagawa et al. [13] also reported that intercalated clay nanoplatelets were more effective than exfoliated nanoplatelets in preventing crack propagation in the matrix of carbon fibre reinforced laminates. In view of these findings it is concluded that the reduced toughness obtained in the present study is due to the microscale distribution of the clay rather than the nanoscale distribution. 4.5. CONCLUSIONS •

The addition of nanoclay to the matrix produced a progressive increase in bearing stiffness with improvements of 10% and 23% being obtained at 7.5 and 12.5 phr nanoclay respectively.

•

A more modest improvement in bearing strength was obtained with a 7% increase a being achieved at 12.5 phr nanoclay indicating that stiffening of the matrix does translate into improved bearing strength.

- 211 -

Chapter 4 - Fibre laminated nanocomposites

•

The strain to failure was reduced by the addition of nanoclay. This is attributed to a change in failure mode brought about by the introduction of the nanoclay into the matrix resin.

•

The spacing of the nanoclay layers in the laminates was only half that obtained in the clay-resin nanocomposites with the same clay content. This is attributed to the resin curing prematurely within the clay galleries during the prepreg drying stage.

- 212 -

Chapter 4 - Fibre laminated nanocomposites

REFERENCES [1] [2]

[3]

[4] [5] [6]

[7]

[8] [9]

[10]

[11]

[12]

[13]

P.S. Wu and C.T. Sun, "Modeling bearing failure initiation in pin-contact of composite laminates ", Mechanics of Materials, 29 [3-4] 325-335 (1998). S. Drapier and M.R. Wisnom, "Finite -element investigation of the compressive strength of non-crimp-fabric-based composites ", Composites Science and Technology, 59 [8] 1287-1297 (1999). F. Paris, A. Gonzalez, E. Graciani, M. Flores, and E.d. Castillo, "A 3D FEM study of compressive behaviour of non-crimp fabrics ". Presented at 11th European conference on composite materials , Rhodes, Greece, 2004. B. Budiansky and N.A. Fleck, "Compressive failure of fibre composites", Journal of the Mechanics and Physics of Solids, 41 [1] 183-211 (1993). C.T. Sun and A. Wanki Jun, "Compressive strength of unidirectional fiber composites with matrix non -linearity", Composites Science and Technology, 52 [4] 577-587 (1994). A. Crosky, D. Kelly, R. Li, X. Legrand, N. Huong, and R. Ujjin, "Improvement of bearing strength of laminated composites", Composite Structures, 76 [3] 260271 (2006). N.A. Siddiqui, R.S.C. Woo, J.-K. Kim, C.C.K. Leung, and A. Munir, "Mode I interlaminar fracture behavior and mechanical properties of CFRPs with nanoclay-filled epoxy matrix", Composites: Part A, In Press, Corrected Proof (2006). http://www.biology.ed.ac.uk/research/groups/jdeacon/statistics/tress4a.html. School of Biological Science, The University of ENDINBURGH. [cited 2006]. J.F. Timmerman, B.S. Hayes, and J.C. Seferis, "Nanoclay reinforcement effects on the cryogenic microcracking of carbon fiber/epoxy composites", Compos ites Science and Technology, 62 [9] 1249-1258 (2002). C. Zilg, R. Mühlhaupt, and J. Finter, "Morphology and toughness/stiffness balance of nanocomposites based upon anhydride-cured epoxy resins and layered silicates", Macromolecular Chemstry and Physics., 200 [3] 661-670 (1999). O. Becker, R. Varley, and G. Simon, "Morphology, thermal relaxations and mechanical properties of layered silicate nanocomposites based upon high functionality epoxy resins", Polymer, 43 [16] 4365-4373 (2002). D.G. Jo¨rg Fro¨hlich, Ralf Thomann, Rolf Mu¨lhaupt, "Synthesis and Characterisation of Anhydride-Cured Epoxy Nanocomposites Containing Layered Silicates Modified with Phenolic Alkylimidazolineamide Cations", Macromolecular Materials and Engineering, 289 [1] 13-19 (2004). H. Miyagawa, R.J. Jurek, A.K. Mohanty, M. Misra, and L.T. Drzal, "Biobased epoxy/clay nanocomposites as a new matrix for CFRP", Composites: Part A, 37 [1] 54-62 (2006).

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Chapter 5 - Z -pin reinforcement

CHAPTER 5

Z-PIN REINFORCEMENT

5.1. INTRODUCTION As seen in the previous chapter, bearing failure occurs by the formation of shear bands at approximately 45° to the bearing surface. This requires lateral displacement of the laminate in an outward direction. In view of this it was considered that through thickness reinforcement could retard initiation of failure and this was investigated using z-pins inserted in the though thickness direction. Z-pin reinforcement is considered to retard crack growth by providing a bridging action [1]. The use of z-pins in laminated composites has yielded significant improvements in strength, w ith a 45% increase in the compression-after-impact strength being reported in one case [2] and a 40% increase in the shear strength of a bonded joint in another [3]. High delamination resistance was also produced [4]. However, insertion of the z-pins can lead to defects such as voids, fibre waviness (causing localised fibre misalignment), in-plane distortion, and resin-rich pockets, and these can produce a marked reduction in the in-plane properties of z-pinned laminates [5-7]. The present study involved inserting carbon fibre/BMI z-pins over a localised region in the vicinity of the hole in bearing loaded test coupons. Z-pins of three different volume densities (0.5%, 2.0% and 4%) and two different diameters (0.28 and 0.51) were examined. Particular care was taken to minimize defects due to insertion of the z-pins and to minimize misalignment of the pins. The bearing response was measured in the zpinned samples during pin loaded bearing using a standard ASTM bearing test. Post - 214 -

Chapter 5 - Z -pin reinforcement

mortem failure analysis was conducted to compare the damage produced in the pinned and unpinned laminates. 5.2. EXPERIMENTAL PROCEDURE 5.2.1. Materials 5.2.1.1. Prepreg Hexcel W3G 282 bi-directional plain weave carbon fibre fabric prepreg impregnated with F593 epoxy resin was used to fabricate the laminates. This material is a standard aircraft grade sytem. 5.2.1.2. Z-pins Carbon fibre T300/BMI Z-FibersTM supplied by Aztex Inc., Waltham, Massachusetts, USA were used for the work. To investigate the effect of z-pin content , 0.28 mm diameter z-pins were inserted in three different volume densities of 0.5%, 2.0% and 4%. The pins were arranged in a square-element grid with the spacing between the pins being 3.5, 1.75 and 1.2 mm, respectively. A set of samples was also prepared using 0.51 mm carbon fibre T650/BMI Z-FibersTM (also supplied by Aztex Inc., Waltham, Massachusetts, USA) with a volume density of 2.0% to examine the effect of pin diameter. The spacing between these pins was 3.2 mm. 5.2.2. Sample preparation Sixteen ply quasi-isotropic laminates 90 mm by 300 mm were prepared with a [±45°/0°90°/±45°/0°90°/0°90°/±45°/0°90°/±45°]s stacking sequence with the 0° fibres aligned along the longitudinal axis of the specimen. The thickness of cured unpinned (baseline) laminates was measured prior to fabrication of the z-pinned laminates to determine the required z-pin length. The z-pinned panels were prepared by initially laying up all but the outermost ply on each side (ie, only 14 of the 16 plies). A 25 x 40 mm foam preform containing the z-

- 215 -

Chapter 5 - Z -pin reinforcement

pins was then placed over the prepreg stack in the vicinity of where the 10 mm diamete r hole was to be drilled. The z-pins were inserted so as to be symmetric on both sides of the hole, and extended back 4 mm from the hole centre to resist backward delamination, as shown in Figure 5.1. Z-pin embedded in an area 4 mm over the centreline of φ10 mm bolt hole 70 mm 40 mm 25 mm

290 mm

Fig . 5.1 Area reinforced by z-pins around 10 mm bolthole . The z-pins were inserted into the prepreg stack using an ultrasonic horn. The pins were chamfered at an angle of about 45° to ease their insertion into the prepreg stack. The pins were progressively inserted by moving the ultrasonic horn over the foam pr eform until all the pins had fully penetrated through the prepreg stack, as shown in Figure 5.2.

- 216 -

Chapter 5 - Z -pin reinforcement

(a)

(b)

Fig . 5.2

Photographs showing (a) pins embedded in foam preform and (b) use of

ultrasonic horn to insert pins into prepreg stack. The excess length from both sides of the fully inserted pins was then cut off using a fine-tooth hacksaw blade , as illustrated in Figure 5.3. Two additional plies of scrap prepreg, with the plastic backing st ill attached, were then placed on each side of the 14 ply prepreg stack and the stack then compressed in a 10 tonne press until the thickness of the 14 ply laminate was equal to that of the cured 16 ply baseline laminate. The two plies of scrap prepreg were then removed from each side and the new protruding portion of the pins sanded back using 600 grit emery paper so as to be flush with the surface of

- 217 -

Chapter 5 - Z -pin reinforcement

the 14 ply prepreg stack. This ensured that the pins were of the exact length required in the cured laminate. (a)

(b)

Fig . 5.3 Photographs showing (a) completely inserted pins in prepreg and (b) cutting the preform and remaining portion of the pins from the prepreg stack. Finally, the two remaining plies were placed on either side of the 14 ply stack to produce the 16 ply lay-up. The lay-up was then debulked under a vacuum pressure of 100 kPa for 30 minutes. The debulked lay-up was placed on an aluminium plate over a layer of FEP release film. A layer of FEP release film was then placed on top of the stack followed by an aluminium caul plate 90 mm x300 mm. This was followed by a layer of pierced FEP, then breather cloth. The laminates were then vacuum bagged and cured in an industrial autoclave at 630 kPa using the cure cycle shown in Figure 5.4. The same procedure was used to fabricate the unpinned baseline laminates.

- 218 -

Chapter 5 - Z -pin reinforcement

Fig . 5.4 Cure cycle used to cure laminates (Hexcel). 5.2.3. Bearing Testing The cured laminates were trimmed to a size of 70 mm x 290 mm using a water lubricated diamond saw. A 10 mm diameter hole was drilled at a position on the central axis of the specimen 60 mm from the end, as shown in Figure 5.1. Particular care was taken to avoid damaging the laminate during drilling. Double shear bearing tests were carried out accor ding to a modified version of ASTM D5961 Standard test method for bearing response of polymer matrix composite laminates using an Instron 8504 servohydraulic universal testing machine. The specimens were mounted in the loading fixture shown in Figure 5.5. Aluminum tabs were attached by adhesive tape to the end of the laminate which was inserted into the lower grips. The modification to the ASTM procedure was that no clamping force was applied. However, a pair of washers bonded to the loading plates was used to align the laminate with the load direction. The washers were chosen with a larger internal diameter (25 mm) than the bolt hole (10 mm) to avoid transverse clamping forces on the bearing edge or the adjacent area. The hole was loaded at a crosshead rate of 1 mm/min using a high strength steel bolt with a diameter of 10 mm. Six specimens of the baseline

- 219 -

Chapter 5 - Z -pin reinforcement

and each of the pinned laminates were tested until ultimate failure. Additionally, one sample from each set was loaded only to the point at which audible cracking could be heard (generally, around 90-95% of the failure load) and then removed from the test fixture to allow examination of the early stages of failure.

Steel washers Test specimen Lower grip

Al tabs

Loading direction Fig . 5.5

Schematic representation of fixture assembly for pin loaded bearing test

ASTM D5961. Prior to carrying out the bearing tests the thickness of the specimens was measured at the locations shown in Figure 5.6. Bearing strength was calculated as the ultimate load divided by the pin diameter x average thickness (average of the measurements made at the three locations B, C and D).

- 220 -

Chapter 5 - Z -pin reinforcement

D C

B

Fig . 5.6 Positions at which thickness measurements were made. 5.2.4. Microstructural examination The microstructure of the tested laminates was examined using optical microscopy to reveal the damage produced during the bearing test. The procedure used was the same as that described previously in Section 4.2.3.2. 5.3. RESULTS 5.3.1. Bearing testing

4%

2%

0.5%

0.28 mm pins

Fig. 5.7

2%

0.51 mm pins Baseline

Z-pinned and baseline laminates after bearing test.

- 221 -

Chapter 5 - Z -pin reinforcement

Representative example s of the unpinned (baseline) and pinned specimens after testing are shown in Figure 5.7. It can be seen that the presence of the z-pins has not made the pinned region of the laminate any thicker. Typical load versus displacement curve s for the z-pinned and baseline laminates are shown in Figure 5.8 while the individual curves are given in Appendix D. It can be seen that in both cases the load rises to ultimate failure without any subsequent recovery. This is contrary to what is observed when the hole is clamped, in which case the load subsequently recovers after the initial failure [8].

20 Baseline

Load (kN)

16

Pinned

12

8

4

0 0

0.5

1

1.5

2

2.5

Displacement (mm) Fig. 5.8 Typical load - displacement curves for baseline and z-pinned laminates. The results obtained from the bearing tests for the baseline laminates are given in Table 5.1, while those from the 0.28 mm z-pinned laminates with 0.5, 2 and 4 % z-pins are given in Tables 5.2, 5.3 and 5.4, respectively, and those for the 2% 0.51 mm z-pinned laminates are given in Table 5.5. For the baseline laminates the average bearing strength was 451 ± 35 MPa, the bearing stiffness (measured as the slope of the load displacement curve) was 22.2 ± 2.5 kN/mm, the strain to failure (measured as the displacement to

- 222 -

Chapter 5 - Z -pin reinforcement

failure) was 0.82 ± 0.11 mm and the energy absorbed to failure (area under load displacement curve) was 6111 ± 970 J. The addition of z-pins increased the bearing strength by 7.3 - 9.8%, the bearing stiffness by 7.5- 9.6% and the energy absorbed to failure by 8.5 - 16.3%. However the strain to failure remained unchanged having a value of 0.8 mm in all cases. The changes are summarised in Table 5.6. The bearing strength, bearing stiffness, and energy absorbed to failure all increased in an essentially linear fashion with pin density, Figures 5.9 to 5.11. However there was no significant difference between the 0.28 mm and 0.51 mm pins, Figures 5.12-5.14. Table 5.1

Measured values of thickness, failure load, bearing strength, energy

absorbed to failure, slope to failure and displacement to failure for baseline samples.

Specimen

1

2

3

4

5

6

B

3.52

3.52

3.53

3.54

3.52

3.56

C

3.53

3.52

3.52

3.54

3.51

3.56

D

3.52

3.52

3.53

3.53

3.52

3.57

Ave

3.52

3.52

3.53

3.54

3.52

Bearing area (mm2)

35.25

35.20

35.25

35.35

Failure load (kN)

16.29

15.85

16.26

Bearing strength (MPa)

461.99

450.26

461.37

Thickness (mm)

Ave

STD

3.56

3.53

0.02

35.15

35.65

35.31

0.18

17.91

14.55

14.75

15.93

1.22

506.67

413.83

413.71

451.30

34.93

Energy absorbed 7579.07 5798.76 5813.49 7010.61 5142.63 5326.18 6111.79 to failure (J)

969.97

Slope to failure (kN/mm)

17.48

22.50

23.97

24.35

22.37

22.50

22.19

2.46

Displacement at failure (mm)

1.00

0.80

0.79

0.70

0.88

0.76

0.82

0.11

- 223 -

Chapter 5 - Z -pin reinforcement

Table 5.2

Measured value s of thickness, failure load, bearing strength, energy

absorbed to failure, slope to failure and displacement to failure for laminates pinned with 0.5% 0.28 mm pins.

Specimen

1

2

3

4

5

6

Ave

STD

B

3.51

3.53

3.51

3.52

3.51

3.52

C

3.53

3.55

3.55

3.51

3.52

3.56

D

3.55

3.56

3.55

3.53

3.54

3.55

Ave

3.53

3.55

3.54

3.52

3.52

3.54

3.53

0.01

Bearing area 2 (mm )

35.40

35.55

35.50

35.20

35.30

35.55

35.42

0.14

Failure load (kN)

17.17

17.51

17.74

16.82

15.92

17.80

17.16

0.71

Bearing strength (MPa)

485.07

492.57

499.83

477.77

450.91

500.84

484.50

18.66

Thickness (mm)

Energy absorbed 6533.77 6703.65 7292.43 6406.93 6154.19 6694.36 6630.89 to failure (J)

383.09

Slope to failure (kN/mm)

24.50

23.64

23.68

24.03

22.61

24.71

23.86

0.75

Displacement at failure (mm)

0.82

0.72

0.98

0.89

0.77

0.88

0.84

0.09

Table 5.3

Measured values of thickness, failure load, bearing strength, energy

absorbed to failure, slope to failure and displacement to failure for laminates pinned with 2% 0.28 mm pins .

Specimen

1

2

3

4

5

6

B

3.55

3.54

3.54

3.54

3.54

3.54

C

3.60

3.60

3.59

3.54

3.59

3.57

D

3.58

3.59

3.59

3.6

3.59

3.57

3.58

3.58

3.57

3.56

3.57

35.90

35.95

35.90

35.70

Failure load (kN)

17.17

17.23

17.58

Bearing strength (MPa)

478.29

479.31

489.81

Thickness (mm)

Ave Bearing area (mm

2

Ave

STD

3.56

3.57

0.01

35.90

35.70

35.84

0.11

17.36

17.30

18.98

17.61

0.69

486.37

481.96

531.77

491.25

20.32

Energy absorbed 6521.13 6561.51 6572.18 6692.55 6655.90 7394.58 6732.98 to failure (J)

330.28

Slope to failure (kN/mm)

23.99

23.54

24.78

23.92

23.72

24.77

24.12

0.53

Displacement at failure (mm)

0.77

0.82

0.77

0.80

0.76

0.82

0.79

0.03

- 224 -

Chapter 5 - Z -pin reinforcement

Table 5.4

Measured values of thickness, failure load, bearing strength, energy

absorbed to failure, slope to failure and displacement to failure for laminates pinned with 4% 0.28 mm pins .

Specimen

1

2

3

4

5

6

B

3.57

3.56

3.54

3.59

3.57

3.58

C

3.65

3.62

3.60

3.61

3.62

3.65

D

3.65

3.64

3.59

3.62

3.62

3.65

Ave

3.62

3.61

3.58

3.61

3.60

Bearing area 2 (mm )

36.50

36.30

35.95

36.15

Failure load (kN)

16.91

18.43

17.77

Bearing strength (MPa)

463.38

507.67

494.32

Thickness (mm)

Ave

STD

3.63

3.61

0.02

36.20

36.50

36.27

0.21

18.86

18.04

17.84

17.97

0.66

521.65

498.31

488.79

495.69

19.58

Energy absorbed 6465.90 7449.99 6889.57 7872.80 7156.71 6798.82 7105.63 to failure (J)

501.96

Slope to failure (kN/mm)

23.24

23.88

24.63

24.62

24.51

25.10

24.33

0.66

Displacement at failure (mm)

0.79

0.96

0.82

0.92

0.80

0.73

0.84

0.09

Table 5.5

Measured values of thickness, failure load, bearing stre ngth, energy

absorbed to failure, slope to failure and displacement to failure for laminates pinned with 2% 0.51 mm pins .

Specimen

1

2

3

4

5

6

B

3.55

3.53

3.55

3.54

3.53

3.53

C

3.62

3.64

3.59

3.59

3.57

3.59

D

3.61

3.62

3.60

3.58

3.6

3.59

3.59

3.60

3.58

3.57

3.57

36.15

36.30

35.95

35.85

Failure load (kN)

18.47

17.52

17.33

Bearing strength (MPa)

510.79

482.78

482.07

Thickness (mm)

Ave Bearing area (mm

2

Ave

STD

3.57

3.58

0.01

35.85

35.90

36.00

0.18

15.64

17.67

19.47

17.68

1.27

436.34

492.91

542.37

491.21

35.12

Energy absorbed 7190.55 6725.37 6460.66 6033.01 6723.20 7960.14 6848.82 to failure (J)

663.18

Slope to failure (kN/mm)

24.09

23.73

24.63

22.25

24.35

25.03

24.01

0.97

Displacement at failure (mm)

0.77

0.78

0.78

0.83

0.86

0.80

0.80

0.04

- 225 -

Chapter 5 - Z -pin reinforcement

Table 5.6 Summary of changes in mechanical behaviour produced in z-pinning.

Specimen

Thickness Failure load Bearing strength Energy absorbed (mm) (kN) (MPa) (J)

Baseline

3.53 ± 0.02 15.93 ± 1.22 451.30 ± 34.93

6111.79 ± 969.97 22.19 ± 2.46

3.53 ± 0.01 17.16 ± 0.71 484.50 ± 18.66

6630.89 ± 383.09 23.86 ± 0.75

Z-pin 0.5% (0.28 mm)

Value

Z-pin 2% (0.28 mm)

Value

Z-pin 4% (0.28 mm)

Value

Z-pin 2% (0.51 mm)

Value

Change

Change

Change

Change

0%

7.70%

7.36%

3.57 ± 0.01 17.61 ± 0.69 491.25 ± 20.32 1.13%

10.49%

8.85%

3.61 ± 0.02 17.97 ± 0.66 495.69 ± 19.58 2.27%

12.81%

9.83%

3.58 ± 0.01 17.68 ± 1.27 491.21 ± 35.12 1.42%

10.99%

8.84%

- 226 -

8.49%

Slope to failure

7.52%

6732.98 ± 330.28 24.12 ± 0.53 10.16%

8.67%

7105.63 ± 501.96 24.33 ± 0.66 16.26%

9.63%

6848.82 ± 663.18 24.01 ± 0.97 12.06%

8.20%

Chapter 5 - Z -pin reinforcement

550 500

Bearing strength (MPa)

450 400 350 300 250 200 150 100 0

0.5

2

4

Pin density (%)

Change of bearing strength (%)

12

10

9.83 8.85

R2 = 0.9602

8 7.36 6

4

2

0 0

1

2

3

4

5

Pin density (%) Fig . 5.9 Variation in bearing strength with pin density for 0.28 mm pins. Error bars indicate one standard deviation.

- 227 -

Chapter 5 - Z -pin reinforcement

26

Slope to failure (kN/mm)

24 22 20 18 16 14 12 10 0

0.5

2

4

Pin density (%)

12

Change in stiffness (%)

10

9.63 8.67

R2 = 0.9821

8 7.52 6

4

2

0 0

1

2

3

4

5

Pin density (%)

Fig . 5 .10 Variation in bearing stiffness (slope to failure) with pin density for 0.28 mm pins. Error bars indicate one standard deviation.

- 228 -

Chapter 5 - Z -pin reinforcement

.

8000

Absorbed energy (J)

6000

4000

2000

0 0

0.5

2

4

Pin density (%)

18 16.26

16

Change in energy (%)

14 R2 = 0.9455 12 10.16

10 8.49

8 6 4 2 0 0

1

2

3

4

5

Pin density (%)

Fig . 5.11 Variation in energy absorbed to failure with pin density for 0.28 mm pins. Error bars indicate one standard deviation.

- 229 -

Chapter 5 - Z -pin reinforcement

550 500

Bearing strength (MPa)

450 400 350 300 250 200 150 100 0

0.28

0.51

Pin diameter (mm)

Fig . 5.1 2 Comparison of bearing strength for 0.28 and 0.51 mm z-pins at 2% volume density. Error bars indicate one standard deviation.

26

Slope to failure (kN/mm)

24 22 20 18 16 14 12 10 0

0.28

0.51

Pin diameter (mm) Fig . 5.1 3 Comparison of bearing stiffness for 0.28 and 0.51 mm z-pins at 2% volume density. Error bars indicate one standard deviation.

- 230 -

Chapter 5 - Z -pin reinforcement

8000

Absorbed energy (J)

6000

4000

2000

0 0

0.28

0.51

Pin diameter (mm)

Fig . 5.1 4 Comparison of energy absorbed to failure for 0.28 and 0.51 mm z-pins at 2% volume density. Error bars indicate one standard deviation. 5.3.2. Microstructure Representative micrographs of the baseline and the z-pinned laminates are shown at approximately 95% of the failure load (representing the early stages of failure) , and after ultimate failure, in Figures 5.15-5-16 (baseline) 5.17-5.21 (0.28 mm pins) and 5.22-5.23 (0.51 mm pins). The laminates all showed similar features to those described in Chapter 4 for the pin-contact loaded specimens, these being fibre kinking, shear band formation, shear cracking and delamination. After ultimate failure, the damage had resulted in the development of large cavities within the samples. There were no significant differences in the nature of the damage produced in the z-pinned samples for either the different pin densities or the different pin sizes. However, the damage was substantially different in the baseline and the z-pinned samples.

- 231 -

Chapter 5 - Z -pin reinforcement

Microcraking

1 mm

Fig . 5.15 Optical micrograph of unpinned laminate loaded to 95% of ultimate failure load.

1 mm

Fig. 5.16 Optical micrograph of unpinned laminate at final failure.

- 232 -

Chapter 5 - Z -pin reinforcement

Pin

Pin

(A)

500 µm

(a)

50µ m

(b) Fig . 5.17 Optical micrograph of laminate containing 0.5 % 0.28 mm z-pins loaded to 95% of ultimate failure load (a) full section view and (b) detailed view of region marked A in (a) showing fibre kinking. The z-pins were not present in the plane of section but their position beneath the surface is marked in (a).

- 233 -

Chapter 5 - Z -pin reinforcement

Pin

Pin

500 µm

Fig . 5.18 Optical micrograph of laminate containing 0.5 % 0.28 mm z-pins loaded to ultimate failure. The z-pins were not present in the plane of section but their position beneath the surface is marked.

Pin

(A)

500 µm

(a)

- 234 -

Chapter 5 - Z -pin reinforcement

(b) Fig . 5.19

Optical micrograph of laminate containing 2 % 0.28 mm z-pins loaded to

95% of ultimate failure load (a) full section view and (b) detailed view of region marked A in (a) showing shear bands.

(D)

(B )

500 µm (A) (C)

(a)

- 235 -

Chapter 5 - Z -pin reinforcement

200µm

Fig . 5.20

(b)

(c)

(d)

(e)

Optical micrograph of laminate containing 2 % 0.28 mm z-pins loaded to

ultimate failure load (a) full section view and (b) - (e) detailed views of regions marked A – D in (a). The regions circled in (d) and (e) show crack arrest by the pins.

- 236 -

Chapter 5 - Z -pin reinforcement

500 µm

Fig . 5.2 1

Optical micrograph of laminate containing 4 % 0.28 mm z-pins loaded to

ultimate failure.

(B)

(A) 500 µm

(a)

- 237 -

Chapter 5 - Z -pin reinforcement

200µm

(b)

(c)

Fig . 5.2 2 Optical micrograph of laminate containing 2 % 0.51 mm z-pins loaded to 95% of ultimate failure load (a) full section view and (b) and (c) detailed views of regions marked A and B in (a).

(A)

(a)

- 238 -

500 µm

Chapter 5 - Z -pin reinforcement

(b) Fig . 5.2 3

Optical micrograph of laminate containing 2 % 0.51 mm z-pins loaded to

ultimate failure load (a) full section view and (b) detailed views of region marked A in (a). The most striking difference was seen at 95% failure load where the lateral spreading produced by bearing could be seen to become progressively more restricted with increasing pinning density. A decreased level of damage was also evident. Ultimate failure produced either fracture (Figures 520, 521 and 5.23) or pullout (Figure 5.22) of the pin nearest the bearing surface but the damage was then generally contained between this pin and the next pin. The high quality of the laminates in the vicinity of the z-pins is also evident in the micrographs, for example Figure 5.20 (b) (0.28 mm pins) and Figure 5.22 (b) (0.51 mm pins). 5.4. DISCUSSION The incorporation of local through thickness reinforcement in the form of z-pins increased the bearing stiffness by 8-10%, the bearing strength by 7-10% and the energy absorbed during bearing failure by 9-16%. No change in the strain to failure was

- 239 -

Chapter 5 - Z -pin reinforcement

however observed. The stiffness, toughness and strength all increase d linearly with increasing pin content but there was no difference between the effect produced by 0.28 mm pins and 0.51 mm pins at the same volume density of pins. An improvement in bearing load with the introduction of z-pins has been reported previously, but in this case a caul plate was not used, and no improvement in bearing strength was obtained due to the laminates bulking out in the vicinity of the z-pins [8]. The z-pins were the same 4% 0.28 mm pins as used here. As a result of the bulking out , the increase in bearing load was offset by an increase in thickness at the hole , so that no improvement in bearing strength was achieved. In subsequent work, in which a caul plate was used, it was found that the pins rotated during curing since they were slightly longer than the cured laminate thickness and again no significant improvement in bearing strength was achieved [9]. The method used in the present work was develo ped specifically to avoid bulking out and pin rotation. For the 0.28 mm pins, the cured laminates did increase slightly in thickness with increased pin volume density but the maximum increase was only 2.3%, Table 5.6. In contrast, a 10% increase in thickness was produced in the previous study where no caul plate was used [8]. Similar increases in thickness were seen with both pin diameters (1.1% and 1.4 % for the 0.28 and 0.51 mm pins respectively at 2 volume %). These results indicate that local thickening of the laminates in the vicinity of the pins is a function of the volume fraction of the laminate replaced by the pins and is essentially independent of the pin diameter for the sizes examined. With increased volume fraction more of the matrix resin is displaced by the pins and this becomes progressively harder to redisperse throughout the laminate. Examination of the micrographs in Figures 5.15-5.23 indicates that some rotation of the pins still occurred despite the care taken to ensure that the pin length was the same as the cured laminate thickness. This appears to be due to the pins not fully penetrating the outermost ply on each side of the laminate. These plies were added subsequent to inserting the pins and it was hoped that they would be penetrated by the pins during curing of the laminates. However, it is clear that this did not occur to the level desired. Evidently, it was easier for the pins to rotate than to fully penetrate the outermost ply.

- 240 -

Chapter 5 - Z -pin reinforcement

Curvature

Fig . 5.24 Optical micrograph of curvature of 0° fibres in pinned laminate [3]. As noted in Section 5.1, insertion of the z-pins can lead to defects such as voids, fibre waviness (causing localised fibre misalignment), in-plane distortion, and resin-rich pockets. These defects were largely eliminated using the procedure employed in the present study, as is evident from Figures 5.15-5.23. In particular, fibre waviness was kept to a minimal level as can be seen from Figure 5.20 (b) (0.28 mm pins) and Figure 5.22 (b) (0.51 mm pins). These micrographs contrast with that shown in Figure 5.24 which shows the level of localized fibre curvature that can occur adjacent to the pins. Steeves and Fleck [6] concluded that misalignment of the fibres adjacent to the pins leads to a reduction in the compressive strength of the laminate. Thus, restriction of the waviness (or curvature) of the fibres adjacent to the z-pins is important in bearing applications. The results showed that bearing strength increased linearly with z-pin content. The increase in bearing strength brought about by z-pinning is due to the increased bearing load that the z-pinned laminates can withstand. This is consistent with the findings of the earlier study where z-pinning was found to increase the ultimate bearing load even though no improvement in bearing strength was achieved (due to thickening of the laminates) [8].

- 241 -

Chapter 5 - Z -pin reinforcement

The results obtained correspond well with other studies in which z-pins have been used to improve strength. Chang et al [3] found that the ultimate shear strength of composite lap joints was greatly improved by z-pinning, with up to a 40% increase being achieved with only a 2% density of 0.28 mm pins. However , contrary to what was observed in the present study, the improvement did not increase continually with pin density but instead decreased to 23% at a 4% pin density. The reduction in strength was attributed to a change in the failure mechanism [3]. The improvement in bearing strength achieved by z-pinning is attributed to the bridging action of the z-pins reported by Barret [1] and Rugg [4]. This would become more effective with increasing pin content, consistent with the findings here. The bridging tractions tend to favour stable crack propagation, and thus the possibility of crack arrest after some growth, providing an increase in load carrying capacity [10]. Examples of crack arrest by the z-pins are shown in Figures 5.20 (d) and (e). The bearing stiffness also increased linearly with pin density, Figure 5.10. The improvement in bearing stiffness obtained by z-pinning is in contrast to previous work where the bearing stiffness (again determined as the slope to the failure) was reduced 17% by the insertion of 4% 0.28 mm pins [8]. Farley [11] also reported a reduction in the in-plane stiffness and attributed this to fibre curvature and voids adjacent to the pins. The reduced bearing stiffness obtained in the previous study is probably also similar in origin. The minimisation of these defects in the present study is considered to be responsible for the absence of any loss in stiffness. Greenhalgh et al [10] and Partridge et al [12] reported improved stiffening as a result of z-pinning. This was attributed to the through-the-thickness bridging actions of the z-pins [12]. The effective stiffening from the bridging mechanism is expected to rise proportionately with pin density [3] and this would account for the stiffening observed here. The energy absorbed by the laminates was also increased by z-pinning and showed a linear increase with increasing pin density. This is attributed to the z-pins producing a large scale bridging zone [2, 13]. Once a delamination crack has propagated to this zone, the pin will bridge the crack. The bridging can slow down, or temporally arrest, - 242 -

Chapter 5 - Z -pin reinforcement

the delamination crack. The mechanism for bridging is the z-pin pull-out process during stable crack growth, which applies a traction force to the crack surface forcing it shut. This process absorbs a considerable amount of ene rgy [2, 14]. Crack deflection at the zpins would also increase the energy absorption. Note that cracks can also occur along the edge of z-pins as a result thermally induced stressing during curing. An example is shown from the work of Sw eeting and Thomson in Figure 5.25.

Fig . 5.2 5 SEM micrograph of a crack beside a z-pin [15]. There was no significant change in bearing strength, stiffness or energy absorption when the pin diameter was doubled at the same volume density. As noted by Cartie [16] et al, the maximum load per pin is, however, greater for larger diameter pins. Chang et a l [3] also found that the ultimate strength and elongation limit were improved for both 0.28 and 0.51 mm pins at the same density (2%), but a transition in the failure mechanism was induced by varying the pin size. Partridge and Cartie [12] reported that resin rich pockets formed around the sides of z-pins, filling the space where the laminate fibres were pushed apart. The extent of these pockets depended on both the zpin diameter and pin-to-pin spacing, i.e., the pin density. Significant resin rich pockets were not however observed here.

- 243 -

Chapter 5 - Z -pin reinforcement

5.5. CONCLUSIONS •

The insertion of 0.28 mm and 0.51 mm z-pins locally in the vicinity of the hole increased the bearing strength, stiffness and energy to failure. No change in the strain to failure was however observed. The increases are attributed to the bridging effect of the z-pins.

•

The bearing strength increased linearly with increasing pin content being 7% for 0.5% pin content and reaching 10% at 4% pin content. The bearing stiffness was increased by 8% and 10% at the same pin contents while the energy absorbed to failure was correspondingly increased by 9% and 16%. The increases are attributed to the increasing bridging effect provided by increasing the pin content.

•

No significant change in strength, stiffness or failure energy was observed when the pin size was changed from 0.28 mm to 0.51 m at the same volume content.

•

The method develo ped to prepare the z-pinned laminates resulted in only a minimal (2% max) increase in laminate thickness and also a minimal level of defects such as voids, fibre waviness (causing localised fibre misalignment), in-plane distortion, and resin-rich pockets, which are common in z-pinned laminates.

- 244 -

Chapter 5 - Z -pin reinforcement

REFERENCES [1] [2]

[3]

[4]

[5]

[6]

[7] [8]

[9] [10] [11]

[12] [13]

[14]

[15]

[16]

D.J. Barrett, "The mechanics of z-fiber reinforcement", Composite Structures, 36 [1-2] 23-32 (1996). X. Zhang, L. Hounslow, and M. Grassi, "Improvement of low-velocity impact and compression -after-impact performance by z-fibre pinning ", Composites Science and Technology, 66 [15] 2785-2794 (2006). P. Chang, A.P. Mouritz, and B.N. Cox, "Properties and failure mechanisms of pinned composite lap joints in monotonic and cyclic tension ", Composites Science and Technology, 66 [13] 2163-2176 (2006). K.L. Rugg, B.N. Cox, and R. Massabo, "Mixed mode delamination of polymer composite laminates reinforced through the thickness by z-fibers", Composites: Part A, 33 [2] 177-190 (2002). M. Grassi, X. Zhang, and M. Meo, "Prediction of stiffness and stresses in z-fibre reinforced composite laminates ", Composites: Part A, 33 [12] 1653-1664 (2002). C.A. Steeves and N.A. Fleck, "In-plane properties of composite laminates with through-thickness pin reinforcement", International Journal of Solids and Structures, 43 [10] 3197-3212 (2006). P. Chang, A.P. Mouritz, and B.N. Cox, "Flexural properties of z-pinned laminates", Composites: Part A, In Press, Corrected Proof (2006). A. Crosky, D. Kelly, R. Li, X. Legrand, N. Huong, and R. Ujjin, "Improvement of bearing strength of laminated composites", Composite Structures, 76 [3] 260271 (2006). R. Li, A. Crosky, and D. Kelly, School of Materials Science and Engineering, University of New South Wales, personal communication. 2005. E. Greenhalgh, A. Lewis, R. Bowen, and M. Grassi, "Evaluation of toughening concepts at structural features in CFRP --Part I: Stiffener pull-off", Composites Part A: Applied Science and Manufacturing, 3 7 [10] 1521-1535 (2006). G.L. Farley, "A mechanism responsible for reducing compression strength of through -the-thickness reinforced composite material", Journal of Composite Materials, 26 1784-1795 (1992). I.K. Partridge and D.D.R. Cartie, "Delamination resistant laminates by ZFiber(R) pinning: Part I manufacture and fracture performance", Composites: Part A, 36 [1] 55-64 (2005). M. Grassi and X. Zhang, "Finite element analyses of mode I interlaminar delamination in z-fibre reinforced composite laminates", Composites Science and Technology, 63 [12] 1815-1832 (2003). M. Meo, F. Achard, and M. Grassi, "Finite element modelling of bridging micro mechanics in through -thickness reinforced composite laminates", Composite Structures, 71 [3-4] 383-387 (2005). R.D. Sweeting and R.S. Thomson, "The effect of thermal mismatch on Z-pinned laminated composite structures", Composite Structures, 66 [1-4] 189-195 (2004). D.D.R. Cartie, M. Troulis, and I.K. Partridge, "Delamination of Z-pinned carbon fibre reinforced laminates", Composites Science and Technology, 66 [6] 855861 (2006).

- 245 -

Chapter 6 - Conclusions

CHAPTER 6

CONCLUSIONS

The work was carried out in three stages. The first involved development of nanoclay reinforced epoxy resin nanocomposites. The effect of surfactant on separation of the clay layers prior to addition to the epoxy resin was examine d first. Preliminary work was then carried out by adding nanoclay at loadings of 0-8phr to DGEBA resin.The knowledge gained was then applied to the preparation of TGGDM nanoclay nanocomposites with clay loadings of 0-20phr. In the second stage of the work the nanoclay reinforced TGGDM resin was used to fabricate carbon fibre reinforced laminates and the bearing response of the laminates evaluated using the pin -contact bearing test. In the final stage of eth work the effect of through thickness reinforcement on bearing behaviour was examined by z-pinning conventional carbon fibre laminates using three different volume contents of z-pins and two different pin diameters. The conclusions are summarized below. Nanoclay modification •

Alkylammonium surfactants with short alkyl chains (8 carbon atoms) amines did not expand the layers in the nanoclay appreciably.

•

The longer alkyl chain (16 carbon atoms) surfactant produced a substantial increase in the interlayer spacing and is considered suitable for modification of the nanoclay for reinforcing epoxy resin

- 246 -

Chapter 6 - Conclusions

•

Within the ranges examined, the acid/amine concentration, surfactant concentration and mixing time had no appreciable influence on the interlayer spacing.

•

Surfactant modified nanoparticles with acid/amine ratios less than 1:1 proved difficult to centrifuge.

DGEBA nanocomposites •

For a 2.5 phr addition of nanoclay the level of exfoliation and the compression modulus of the nanocomposites was insensitive to the mixing conditions over the range examined.

•

The modulus of the nanocomposites was reduced when the cure temperature was lower than 80ºC. This is attributed to insufficient diffusion of the resin into the nanoclay gallery.

•

The modulus of the nanocomposites increased progressively with clay content. This is attributed to the dispersed clay layers restricting the mobility of the polymer chains.

•

Exfoliated nanocomposites with an interlayer spacing of 100 Å were obtained up to a loading of 5 phr nanoclay, but the composites obtained at 8.4 phr were intercalated only with a layer spacing of only 63 Å.

TGDDM nanocomposites •

For the clay loadings of 2.5 and 7.5 phr examined the compression modulus and interlayer spacing was insensitive to the mixing speed and mixing temperature over the ranges examined. However longer mixing times were found to beneficial at higher clay loadings. This is attributed to the increased number of clay particles that need to be broken down during mixing.

•

Degassing the resin-nanoclay mixture after stirring for an extra 30 minutes beyond

- 247 -

Chapter 6 - Conclusions

•

The time required for visible gas removal to cease was found to be beneficial. Longer additional degassing times were found to be detrimental. This is attributed to excessive loss of volatiles from the resin.

•

There was no significant effect of curing temperature within the range examined for the three stage cure cycle.

•

Varying the cure time of the first two stages of the cure cycle produced no significant effect. Indeed, either of these steps could have been omitted without causing any significant change. This indicates that both the temperatures used were effective for separation of the nanoclay layers. The modulus was however increased by the inclusion of the final postcuring stage. This confirms that additional cross linking takes place during this stage.

•

Up to 20 phr nanoclay could be added to the TGDDM resin without the mixture becoming excessively viscous. Fully exfoliated nanocomposites with an interlayer spacing of greater than 120 Å were obtained up to 5 phr nanoclay, while preexfolia ted or intercalate d nanocomposites, with interlayer spacings of from 85 Å (7.5 phr nanoclay) to 60 Å (20 phr nanoclay), were obtained at higher loadings.

•

For clay loadings of 5 phr and above, the compression modulus increased linearly with clay content, with an increase of 50% being achieved over that of the pristine clay at 20 phr clay. The results were in good agreement with the predictions of the Halpin -Tsai model for an aspect ratio of 13.

•

The modulus was anomalously high for the 1 and 2.5 phr nanocomposites. This is attributed to the better exfoliation achieved at these clay loadings.

Nanoclay reinforced epoxy matrix carbon fibre laminates •

The addition of nanoclay to the matrix produced a progressive increase in bearing stiffness with improvements of 10% and 23% being obtained at 7.5 and 12.5 phr nanoclay respectively.

- 248 -

Chapter 6 - Conclusions

•

A more modest improvement in bearing strength was obtained with a 7% increase a being achieved at 12.5 phr nanoclay indicating that stiffening of the matrix does translate into improved bearing strength.

•

The strain to failure was reduced by the addition of nanoclay. This is attributed to a change in failure mode brought about by the introduction of the nanoclay into the matrix resin.

•

The spacing of the nanoclay layers in the laminates was only half tha t obtained in the clay-resin nanocomposites with the same clay content. This is attributed to the resin curing prematurely within the clay galleries during the prepreg drying stage.

Z-pinned carbon fibre laminates •

The insertion of 0.28 mm and 0.51 mm z-pins locally in the vicinity of the hole increased the bearing strength, stiffness and energy to failure. No change in the strain to failure was however observed. The increases are attributed to the bridging effect of the z-pins.

•

The bearing strength increas ed linearly with increasing pin content being 7% for 0.5% pin content and reaching 10% at 4% pin content. The bearing stiffness was increased by 8% and 10% at the same pin contents while the energy absorbed to failure was correspondingly increased by 9% and 16%. The increases are attributed to the increasing bridging effect provided by increasing the pin content.

•

No significant change in strength, stiffness or failure energy was observed when the pin size was changed from 0.28 mm to 0.51 m at the same volume content.

•

The method developed to prepare the z-pinned laminates resulted in only a minimal (2% max) increase in laminate thickness and also a minimal level of defects such as voids, fibre waviness (causing localised fibre misalignment), in-plane distortion, and resin-rich pockets, which are common in z-pinned laminates.

- 249 -

Appendix A

APPENDIX A

COMPRESSION TESTS

1. DGEBA Nanocomposites

160 140

No1 120

No2

Stress (MPa)

No3 100

No4 No5

80 60 40 20 0 0

0.1

0.2

0.3

0.4

0.5

Strain (mm/mm)

Fig. A.1 Stress versus strain curves of neat resin DGEBA resin.

Appendix A

160 140

No1 120

No2

Stress (MPa)

No3 100

No4 No5

80 60 40 20 0 0

0.1

0.2

0.3

0.4

0.5

Strain (mm/mm)

Fig. A.2 Stress versus strain curves of 0.5 phr nanoclay nanocomposites.

160 140

No1 120

No2

Stress (MPa)

No3 100

No4 No5

80 60 40 20 0 0

0.1

0.2

0.3

0.4

0.5

Strain (mm/mm)

Fig. A.3 Stress versus strain curves of 2.5 phr nanoclay nanocomposites.

Appendix A

180 160

No1 140

No2 No3

Stress (MPa)

120

No4 100

No5

80 60 40 20 0 0

0.1

0.2

0.3

0.4

0.5

Strain (mm/mm)

Fig. A.4 Stress versus strain curves of 5 phr nanoclay nanocomposites.

180 160

No1 140

No2 No3

Stress (MPa)

120

No4 100

No5

80 60 40 20 0 0

0.1

0.2

0.3

0.4

0.5

Strain (mm/mm)

Fig. A.5 Stress versus strain curves of 8.4 phr nanoclay nanocomposites.

Appendix A

2. TGDDM Nanocomposites 300

250

No1 No2 No3

Stress (MPa)

200

No4 No5

150

100

50

0 0

0.1

0.2

0.3

0.4

0.5

Strain (mm/mm)

Fig. A.6 Stress versus strain curves of neat TGDDM resin. 250

No1

200

No2

Stress (MPa)

No3 No4

150

No5 100

50

0 0

0.05

0.1

0.15

0.2

0.25

0.3

0.35

Strain (mm/mm)

Fig. A.7 Stress versus strain curves of 2.5 phr nanoclay nanocomposites.

Appendix A

250

No1

200

No2

Stress (MPa)

No3 No4

150

No5 100

50

0 0

0.05

0.1

0.15

0.2

0.25

0.3

0.35

Strain (mm/mm)

Fig. A.8 Stress versus strain curves of 5 phr nanoclay nanocomposites.

180 160

No1 140

No2 No3

Stress (MPa)

120

No4 100

No5

80 60 40 20 0 0

0.05

0.1

0.15

0.2

0.25

Strain (mm/mm)

Fig. A.9 Stress versus strain curves of 7.5 phr nanoclay nanocomposites.

Appendix A

250

No1

200

No2

Stress (MPa)

No3 No4

150

No5 100

50

0 0

0.05

0.1

0.15

0.2

0.25

0.3

0.35

Strain (mm/mm)

Fig. A.10 Stress versus strain curves of 10 phr nanoclay nanocomposites.

250

No1

200

No2

Stress (MPa)

No3 No4

150

No5 100

50

0 0

0.05

0.1

0.15

0.2

0.25

0.3

0.35

Strain (mm/mm)

Fig. A.11 Stress versus strain curves of 12.5 phr nanoclay nanocomposites.

Appendix A

180 160

No1 140

No2 No3

Stress (MPa)

120

No4 100

No5

80 60 40 20 0 0

0.05

0.1

0.15

0.2

0.25

Strain (mm/mm)

Fig. A.12 Stress versus strain curves of 15 phr nanoclay nanocomposites.

180 160

No1 140

No2 No3

Stress (MPa)

120

No4 100

No5

80 60 40 20 0 0

0.05

0.1

0.15

0.2

0.25

Strain (mm/mm)

Fig. A.13 Stress versus strain curves of 17.5 phr nanoclay nanocomposites.

Appendix A

180 160

No1 140

No2 No3

Stress (MPa)

120

No4 100

No5

80 60 40 20 0 0

0.05

0.1

0.15

0.2

0.25

Strain (mm/mm)

Fig. A.14 Stress versus strain curves of 20 phr nanoclay nanocomposites.

Appendix B

APPENDIX B

PIN-CONTACT TEST RESULTS FOR LAMINATED NANOCOMPOSITES

25

No1 20

No2 No3

Load (kN)

No4 15

No5 No6

10

5

0 0

0.2

0.4

0.6

0.8

1

1.2

Displacement (mm)

Fig. B.1 Load versus displacement curves of baseline laminates (3K CFRP - 56.9 vol.%).

Appendix B

25

No1 20

No2 No3

Load (kN)

No4 15

No5 No6

10

5

0 0

0.2

0.4

0.6

0.8

1

Displacement (mm)

Fig. B.2

Load versus displacement curves of 7.5 phr nanoclay laminates (3K CFRP -

57.6 vol.%).

25

No1 20

No2

Load (kN)

No3 No4 15

No5 No6

10

5

0 0

0.2

0.4

0.6

0.8

1

1.2

Displacement (mm)

Fig. B.3 55 vol.%).

Load versus displacement curves of 7.5 phr nanoclay laminates (3K CFRP -

Appendix B

25

No1 20

No2 No3

Load (kN)

No4 15

No5 No6

10

5

0 0

0.2

0.4

0.6

0.8

1

1.2

Displacement (mm)

Fig. B.4 Load versus displacement curves of 12.5 phr nanoclay laminates (3K CFRP 51.7 vol.%).

20 18

Load (kN)

No1 16

No2

14

No3 No4

12

No5 No6

10 8 6 4 2 0 0

0.2

0.4

0.6

0.8

1

1.2

Displacement (mm)

Fig. B.5 Load versus displacement curves of baseline laminates (6K CFRP - 61.3 vol.%).

Appendix B

25

No1 20

No2

Load (kN)

No3 No4 15

No5 No6

10

5

0 0

0.2

0.4

0.6

0.8

1

1.2

Displacement (mm)

Fig. B.6 Load versus displacement curves of 7.5 phr nanoclay laminates (6K CFRP 61.1 vol.%).

Appendix C

APPENDIX C

STUDENTS’ T-TEST FOR LAMINTED NANOCOMPOSITES

a/ Effect of nanoclay on bearing strength of the laminated nanocomposites

7.5 phr nanoclay laminate (1) with 3K CF – Nomalised 57 vol.% 7.5 phr (1) 0 phr Mean 467.22 486.1 ± SE 35.37 18.2 n 6 6 df 10 SE² 1251.037 331.24 s² 1582.277 s 39.77784 t stat -0.47464 t critical (p=0.05) 2.23 t stat < t Critical => Difference is not significant 7.5 phr nanoclay laminate (2) with 3K CF – Nomalised 57 vol.% 7.5 phr (2) 0 phr Mean 492.24 486.1 ± SE 9.87 18.2 n 6 6 df 10 SE² 97.4169 331.24 s² 428.6569 s 20.70403 t stat 0.296561 t critical (p=0.05) 2.23 t stat < t Critical => Difference is not significant 12 .5 phr nanoclay laminate with 3K CF – Nomalised 57 vol.%

Appendix C

Mean ± SE n df SE² s² s t stat t critical (p=0.01)

12.5 phr 521.34 13.16 6 10 173.1856 252.2128 9.16902 3.843377

0 phr 486.1 18.2 6 331.24

3.17

t stat > t Critical => Difference is highly significant 7.5 phr nanoclay laminate with 6K CF – Nomalised 61 vol.% 7.5 phr 0 phr Mean 461.46 450.39 ± SE 17.79 20.5 n 6 6 df 10 SE² 316.4841 420.25 s² 736.7341 s 27.14285 t stat 0.407842 t critical (p=0.05) 2.23 t stat < t Critical => Difference is not significant b/ Effect of nanoclay on bearing stiffness of the laminated nanocomposites

i. Laminate reinforced by 7.5 phr (1) nanoclay – 3K CF Replicates 7.5 phr (1) 0 phr 1 24.25 19.41 2 24.19 22.25 3 22.29 23.03 4 23.61 21.17 5 25.14 23.87 6 24.2 18.59 t-Test: Two-Sample Assuming Equal Variances

Mean Variance Observations Pooled Variance Hypothesized Mean Difference df t Stat P(TThe difference is significant ii. Laminate reinforced by 7.5 phr ( 2) nanoclay – 3K CF Replicates 7.5 phr (2) 0 phr 1 20.15 19.41 2 18.59 22.25 3 26.47 23.03 4 19.56 21.17 5 25.55 23.87 6 25.18 18.59 t-Test: Two-Sample Assuming Equal Variances Variable 1 Variable 2 Mean 22.58333 21.38667 Variance 12.33167 4.277667 Observations 6 6 Pooled Variance 8.304667 Hypothesized Mean Difference 0 df 10 0.719238 t Stat P(T