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The Effects of Fibre Moisture Content During Manufacturing on the Properties of Flax/Epoxy Composites

A Thesis submitted to The University of Manchester for the Degree of

Doctor of Philosophy

in The Faculty of Science and Engineering

2020

ABDUL HAKIM ABDULLAH

Schools of Natural Sciences

Department of Materials

Table of Content

Table of Content ...... 2 List of Figure ...... 6 List of Tables ...... 12 List of Symbols ...... 14 List of Abbreviations ...... 16 Abstract ...... 18 Declaration ...... 19 Copyright Statement ...... 20 Acknowledgement ...... 21 Chapter 1 Introduction ...... 22 1.1 Overview ...... 22 1.2 Future of Natural Fibre Based Composites ...... 24 1.3 Research Motivation ...... 25 1.4 Research Objectives ...... 26 1.5 Thesis Outlines ...... 26 Chapter 2 Literature Review ...... 28 2.1 Natural Fibres ...... 28 2.1.1 Flax Fibres and Their Composites ...... 29 2.1.2 Chemical Composition and Internal Structure of Natural Fibre ...... 32 2.1.3 The Origin of Hygroscopic Behaviour in Natural Fibres ...... 33 2.2 Polymeric Matrices ...... 41 2.2.1 Epoxy Resin ...... 41 2.2.2 Moisture Contamination in Uncured Epoxy Resin ...... 43 2.3 Composites Manufacturing ...... 48 2.4 Ideal Humidity Conditions for Composites Processing and Fabrication .... 49 2.4.1 Moisture Presence during Composites Manufacturing ...... 50 2.5 Fibre/Matrix Interface Bonding ...... 51 2.5.1 Interfacial Bond Strength Measurement ...... 51 2.5.2 Relationship between MC and Interfacial Bond Strength ...... 52 2.5.3 Influence of Fibre Swelling and Shrinkage at the Fibre/Matrix Interface ...... 54 2.5.4 Effect of Chemical Treatment to Increase Interfacial Bonding ...... 55 2.6 Surface treatment to Enhance Natural Fibre Composites Performance ..... 56 2.6.1 Chemical Treatments ...... 57

2.6.2 Silane on Alkalized Fibre Treatment ...... 58 2.6.3 Physical Treatments ...... 59 2.7 Effect of Humidity on Composite Mechanical Performance ...... 60 2.8 Fracture Toughness and Impact Properties of Composites ...... 64 2.8.1 Interlaminar Fracture Toughness (ILFT) ...... 64 2.8.2 Low Velocity Impact ...... 66 2.8.3 Compression After Impact (CAI) ...... 68 2.9 Key findings of The Literature ...... 68 Chapter 3 Experimental ...... 70 3.1 Materials ...... 70 3.2 Fabric Surface Treatments ...... 71 3.2.1 Alkaline treatment ...... 71 3.2.2 Silane Treatment ...... 71 3.2.3 Alkaline & Silane Treatment ...... 72 3.3 Conditioning Flax Fibres ...... 72 3.4 Composites Manufacturing ...... 73 3.5 Sample Cutting and Post-Conditioning ...... 74 3.6 Characterization Methods ...... 75 3.6.1 Dynamic Vapour Sorption (DVS) Analysis ...... 75 3.6.2 Simple Weight Gain Method ...... 75 3.6.3 Crimp Measurements ...... 77 3.6.4 FTIR-ATR Spectroscopy ...... 77 3.6.5 Single Fibre Tensile Testing (SFTT) ...... 77 3.6.6 Density Measurements ...... 79 3.6.7 Determination of the Fibre Volume Fraction of Composites ...... 81 3.6.8 Microbond Shear Test Measurements of Fibre-Resin Interfacial Shear Strength (IFSS) ...... 81 3.6.9 Dynamic Mechanical Analysis (DMA) ...... 84 3.6.10 Tensile Testing ...... 85 3.6.11 Flexural Testing ...... 86 3.6.12 Interlaminar Shear Strength (ILSS) ...... 87 3.6.13 Mode I - Interlaminar Fracture Toughness (ILFT) ...... 87 3.6.14 Mode II- Interlaminar Fracture Toughness (ILFT) ...... 91 3.6.15 Low Velocity Impact Testing ...... 92 3.6.16 Ultrasonic C-Scanning Inspection ...... 92 3.6.17 Compression after impact (CAI) Test ...... 93 3.6.18 Scanning Electron Microscopy (SEM) ...... 93 3

Chapter 4 Results and Discussions ...... 94 4.1 Characterization of Untreated Flax Fibre ...... 94 4.1.1 Moisture Content ...... 94 4.1.2 Relative Humidity (RH) Sensitivity ...... 95 4.1.3 Graphical Method for Estimating MC of Flax Fabric ...... 96 4.1.4 Changes of Fibre Density ...... 98 4.2 Interfacial Shear Strength (IFSS) of Flax/Epoxy Composites ...... 99 4.2.1 Microbond Force-Displacement Curves ...... 99 4.2.2 Effect of Embedded Length ...... 101 4.2.3 Interfacial Shear Strength (IFSS) and Weibull Analysis ...... 102 4.3 Mechanical Performance of Untreated Flax/Epoxy Composites ...... 106 4.3.1 Influence of Moisture Content on Fibre Volume Fraction ...... 106 4.3.2 Tensile Properties ...... 107 4.3.3 Flexural Properties ...... 111 4.3.4 Interlaminar Shear Strength (ILSS) ...... 113 4.3.5 Dynamic Mechanical Analysis (DMA) ...... 115 4.3.5 Mode I DCB - Interlaminar Fracture Toughness (ILFT) ...... 120 4.3.6 Mode II ENF - Interlaminar Fracture Toughness (ILFT) ...... 127 4.3.7 Low Velocity Impact ...... 130 4.3.8 Compression After Impact (CAI) ...... 136 4.4 Mechanism of Moisture Interaction at the Interface during Flax/Epoxy Composites Manufacturing ...... 138 4.5 Characterization of Chemically Treated Flax Fibre ...... 140 4.5.1 Moisture Content of Chemically Treated Flax Fibre ...... 140 4.5.2 Fourier Transform Infrared (FTIR) Analysis ...... 144 4.5.3 Changes of Fibre Density ...... 148 4.5.4 Single Fibre Tensile Test ...... 150 4.5.5 Physical Characteristic of Chemically Treated Flax Fabric ...... 157 4.6 Interfacial Shear Strength (IFSS) of Treated Flax/Epoxy Composites ..... 159 4.6.1 Microbond Force-Displacement Curves ...... 159 4.6.2 Interfacial Shear Strength (IFSS) and Weibull Analysis ...... 161 4.7 Mechanical Performance of Chemically Treated Flax/Epoxy Composites ...... 173 4.7.1 Influence of Moisture Content on Fibre Volume Fraction of Composites ...... 173 4.7.2 Dynamic Mechanical Analysis ...... 174 4.7.3 Tensile Properties ...... 180

4.7.4 Flexural Properties ...... 185 4.7.5 Mode I Interlaminar Fracture Toughness ...... 191 Chapter 5 Conclusions ...... 202 5.1 Effect of Moisture Content ...... 202 5.2 Effect of Surface Treatments ...... 203 5.3 Suggestions for Future Work ...... 205 References ...... 208

Final word count: 38,156

List of Figure

Figure 1.1: Plant fibres FRP market share in 2010 out of 2.4 million tonnes [4]. .... 24 Figure 1.2: Examples of natural fibre composites in non-structural application; (a) automotive interior parts [7] and effort for semi-structural/structural applications; (b) a canoe [13] and (c) a bicycle [14]...... 25 Figure 2.1: Classification of natural fibres [23]...... 28 Figure 2.2: Schematic morphology of flax fibre [28]...... 29 Figure 2.3: Comparison between man-made and natural fibres in term of (a) sound absorption coefficient [35] and logarithmic damping decrement [38]...... 31 Figure 2.4: Schematic (a) cross section of a technical fibre and (b) internal structure of an elementary flax fibre [48]...... 33 Figure 2.5: Moisture absorption in flax fibre (Green and Duralin) [64]...... 35 Figure 2.6: Absorption and desorption curves corresponding to RH changes for (a) flax fibre (b) hemp fibre (c) jute fibre and (d) sisal. Curves are redrawn from [58]. 37 Figure 2.7: Messiry’s graphical representation of fibre density versus fibre MC. Curve is redrawn from [73]...... 38 Figure 2.8: Tensile properties of flax fibre (Marylin and Hermes) with respect of RH (a) tensile strength (b) Young’s modulus and (c) tensile strain [77]...... 40 Figure 2.9: Crosslinking reactions between an epoxy group of the resin and an amine group of the hardener [84]. R1, R2 and R3 present the remainders of the molecules. 43 Figure 2.10: Crosslinking reactions between an epoxy group of the resin and an amine group of the hardener in the presence of water [84]. R1, R2 and R3 present the remainders of the molecule...... 44 Figure 2.11: Influence of water addition on the uncured epoxy; (a) cure reaction - data from FTIR [91] (b) glass transition, Tg - data from DSC curve [86]...... 45 Figure 2.12: Influence of nanofiber MC in epoxy during cure, (a) initial reaction rate and (b) final glass transition temperature [68]...... 46 Figure 2.13: Mechanical properties of water contaminated in uncured epoxy; (a) ultimate tensile strength [19] and (c) flexural load strength of FRP-reinforced concrete beam bars bonded with SikaDur 30 epoxy adhesives [93]...... 48 Figure 2.14: Schematic of resin infusion (RI) [97]...... 49 Figure 2.15: Daily plot of humidity in Manchester, United Kingdom on 1st July 2018 [99]...... 50 Figure 2.16: IFSS/load-displacement curves showing strong and weak fibre/matrix interfaces; (a) flax/unsaturated polyester composites [119] and (b) bamboo/vinyl ester composites [18]...... 53 Figure 2.17: Influence of MC/RH on IFSS values; (a) bamboo/vinyl ester composites [123] and (b) flax/unsaturated polyester composites [119]...... 54

Figure 2.18: Influence of chemical treatments on (a) moisture absorption and (b) interfacial shear strength of bamboo/vinyl ester composites during composites fabrication [123]...... 56 Figure 2.19: Influence of humidity on flax/Acrodur biocomposites; (a) tensile strength and (b) tensile modulus [149]...... 61 Figure 2.20: Flexural properties of flax/unsaturated polyester composites fabricated at different humidity condition [119]...... 62 Figure 2.21: Influence of humidity during fabrication of unidirectional flax/epoxy composites; (a) flexural and (b) tensile properties [150]...... 63 Figure 2.22: The three modes of fracture toughness [153]...... 64 Figure 2.23: Failure modes of composites due to delamination [153]...... 65 Figure 2.24: Typical impact response of composites (a) rebounding (b) penetration (c) perforation [162]...... 66 Figure 2.25: Schematic of failure mode of laminated composites [166]...... 67 Figure 3.1: Photograph of (a) 2x2 twill flax fabric and (b) 250 tex flax yarn and (c) flax fabric ready for surface treatment and composites fabrication...... 70 Figure 3.2: Chemical structure of APTMS silane agent...... 72 Figure 3.3: Resin infusion of flax/epoxy composites showing (a) lay-up configuration, (b) vacuum check for sign of leakage after sealing (c) resin infusion process without vacuum pump aid and (d) cured composites prior to cutting...... 74 Figure 3.4: The Dynamic Vapour Sorption Analyser (DVS) used in this study...... 75 Figure 3.5: Photograph of (a) produced MC fabric samples and (b) weighting the fabric on the sensitive balance (Sartorius Waage Analytical CP153)...... 76 Figure 3.6: Single fibre tensile test configuration...... 78 Figure 3.7: Excluded samples from SFTT; (a) twisted fibre (b) irregular cross- sectional area...... 78 Figure 3.8: Density measurement configuration set-up on a Metler Toledo analytical balance...... 80 Figure 3.9: Estimating the produced microdroplet size using a size reference...... 82 Figure 3.10: Optical images of (a) accepted microdroplet and (b) rejected microdroplet...... 83 Figure 3.11: Microbond shear test configuration...... 83 Figure 3.12: DMA fixture set-up...... 85 Figure 3.13: Photograph of (a) tensile samples and (b) fractured composites within gauge length after tensile test...... 86 Figure 3.14: Photograph of (a) flexural samples and (b) three point bending test configuration...... 86 Figure 3.15: Specimen geometry and dimension for DCB test...... 89 Figure 3.16: Mode I-ILFT specimen during testing...... 89 Figure 3.17: Example of determination correction factor for delamination length, Δ for DCB test...... 90 7

Figure 3.18: Examples of determination of crack initiation for the DCB test...... 90 Figure 3.19: Specimen and test configuration for 4ENF test...... 91 Figure 3.20: Low velocity impact test configuration [169][163]...... 92 Figure 3.21: (a) Anti-buckling guide CAI fixture and (b) CAI experiment...... 93 Figure 4.1: MC as function of RH of untreated flax fibres at 23°C...... 94 Figure 4.2: Moisture absorption curves of untreated flax fibres subjected to 10% RH and 90% RH at 23°C as a function of time...... 96 Figure 4.3: Graphical presentation for estimating MC of untreated flax fibre. UT1 to UT5 is standardized fabric samples (81 mm x 81 mm)...... 98 Figure 4.4: Comparison between current experimental data and Messiry’s model [73]. The dashed line is a linear-regression plot of the experimental data...... 99 Figure 4.5: Typical microbond force-displacement curves of untreated flax/epoxy composites...... 101 Figure 4.6: Plot of IFSS as function of embedded length for flax/epoxy composites...... 102 Figure 4.7: Average IFSS of flax/epoxy composites as a function of MC...... 103 Figure 4.8: Weibull fitting on the IFSS data distribution as a function of MC ...... 104 Figure 4.9: SEM images showing (a) the rough fibre surface of low MC specimen and (b) smoother surface of high MC specimen...... 105 Figure 4.10: Typical tensile stress-strain curves of untreated flax/epoxy composites with various fibre MC during composites fabrication. The curves were normalised at 30% 푉푓...... 108 Figure 4.11: SEM images showing (a) strong fibre/matrix interface of a 6.3 wt.% MC composite and (b) poor fibre/matrix interface of a 12.9 wt.% MC composite. 110 Figure 4.12: Typical flexural stress-strain curves of untreated flax/epoxy composites with various fibre MC during composites fabrication. The curves were normalised at 30% 푉푓...... 111 Figure 4.13: SEM images showing fracture surface of (a) an untreated 2.1 wt.% MC sample and (b) an untreated 12.6 wt.% MC sample...... 113 Figure 4.14: Typical load versus deflection curves of untreated flax/epoxy composites manufactured at 6.3 wt.% and 12.8wt.% fibre MC...... 114 Figure 4.15: Photograph of ILSS samples (untreated 6.3wt.% MC) showing tension failure, ...... 115 Figure 4.16: Temperature dependence of storage modulus, E´ of untreated flax/epoxy composites at different fibre MC...... 116 Figure 4.17: Temperature dependence of loss modulus, E” of untreated flax/epoxy composites at different fibre MC...... 117 Figure 4.18: Temperature dependence of tan delta of untreated flax/epoxy composites at different fibre MC...... 118 Figure 4.19: Typical load extension curves of untreated 2.3 wt.% MC, 6.6 wt.% MC and 12.7 wt.% MC DCB specimens...... 121

Figure 4.20: Typical resistance curves (R-Curves) of untreated 2.3 wt.% MC, 6.6 wt.% MC and 12.7 wt.% MC DCB specimens...... 121 Figure 4.21: Force-displacement curves of (a) brittle and (b) ductile matrix [228]...... 123

Figure 4.22: Influence of MC on Mode I (a) GIC initiation (NL, VIS, 5% Offset) and (b) GIC propagation as a function of MC of flax/epoxy composites...... 124 Figure 4.23: Photograph of delaminated composites showing colour changes between untreated 2.3 wt.% MC and 12.7 wt.% MC samples...... 124 Figure 4.24: Qualitative observation of fibre bridging (white arrows) on Mode I DCB test showing fibre bridging of untreated (a) 2.3 wt.% MC (b) 6.6 wt.% MC and 12.7 wt.% MC samples...... 126 Figure 4.25: Damage within a fibre of untreated (a) 2.2 wt.% MC and (b)12.7 wt.% MC samples...... 127 Figure 4.26: Photograph of side view of composites after Mode II ENF test of untreated (a) 6.6 wt.% MC and (b) 12.6 wt.% MC samples...... 129 Figure 4.27: Photograph of delaminated of the untreated 12.7 wt.% MC sample surfaces showing the consequences of crack deflection from the mid-plane where the crack path eventually reached the carbon fibre fabric...... 129 Figure 4.28: Typical load-displacement curves and R-curves of Mode II ENF of untreated flax/epoxy composites...... 130 Figure 4.29: Typical load-deflection curves with variation impact loading at different MCs during composites fabrication of 6.5% wt. MC and 12.5% wt. MC...... 132 Figure 4.30: Evolution of (a) peak force, (b) contact time and (c) absorbed energy as a function of impact energy...... 134 Figure 4.31: Ultrasonic C-Scan of 6.5 wt.% MC and 12.5 wt.% MC impacted specimens...... 136 Figure 4.32: The (a) CAI strength and (b) normalised CAI strength as a function of the incident impact energy...... 137 Figure 4.33: Moisture interaction during composites fabrication during (a) wetting stages (b) curing stages. (c) Weak chemical bonding due to excess of MC on the fibre surface...... 139 Figure 4.34: MC of (a) 3.0% NaOH and 4.5% NaOH treated fibre as a function of RH...... 141 Figure 4.35: MC of 1.0% silane treated fibre as function of RH...... 142 Figure 4.36: MC of 3.0% NaOH & 1.0% silane treated fibres as a function of RH...... 143 Figure 4.37: Effect of chemical treatments on the MC of flax fibres...... 144 Figure 4.38: FTIR-ATR spectra of NaOH treated fibres with different NaOH concentration that shows (a) hemicellulose and (b) lignin removal...... 145 Figure 4.39: FTIR-ATR spectra of silane treated fibres with different silane concentrations...... 147

Figure 4.40: FTIR-ATR spectra of combined 3.0% NaOH & 1.0% silane treated fibres...... 148 Figure 4.41: Fibre density as function of MC after chemical treatments...... 149 Figure 4.42: Comparison between current experimental data and Messiry’s model for (a) 3.0% NaOH, (b) 1.0% silane and (c) 3.0% NaOH & 1.0% silane treated fibre. Dashed plot is a linear-regression of experimental data...... 150 Figure 4.43: Typical stress-strain curves of elementary fibres after (a) NaOH, (b) silane and (c) 3.0% NaOH & 1.0% silane treatments. The dash line is purposely added, showing the non-linear behaviour of the curves...... 151 Figure 4.44: Two-parameter Weibull distribution fitting on the tensile strength data of flax fibres after (a) NaOH (b) silane and (c) silane on alkalized treated fibre. ... 153 Figure 4.45: Influence of the chemical treatments to the (a) % crimp and (b) fibre areal density of flax fabric...... 158 Figure 4.46: Photographs of (a) fibre yarn and (b) flax fabric before and after 3.0% NaOH treatment...... 158 Figure 4.47: Typical force-displacement curves of (a) 1.5% NaOH, (b) 3.0% NaOH and (c) 4.5% NaOH, (d) 1.0% silane, (e) 1.5% silane and combined NaOH-silane treated samples: (f) 3.0% NaOH & 1.0% silane...... 160 Figure 4.48: Two-parameter Weibull distribution fitting on the IFSS data; (a-c) NaOH, (d-e) silane and (f) 3.0% NaOH &1.0% silane treated fibres...... 163 Figure 4.49: Influence of MC on the IFSS after (a) NaOH and (b) silane treatments...... 166 Figure 4.50: IFSS comparison between 3.0% NaOH, 1.0% silane and combined treatment as a function of MC...... 167 Figure 4.51: SEM images of flax fibres for different conditions: (a) Untreated, (b) 3.0% NaOH, (c) 4.5% NaOH, (d) 1.0% silane (e) 1.5% silane and (f) 3.0% NaOH & 1.0% silane treatments...... 168 Figure 4.52: SEM images showing surface morphology after microbond test of (a) 3.0% NaOH-2.4% wt. MC and (b) 3.0% NaOH-12.9% wt. MC...... 170 Figure 4.53: SEM images showing surface morphology after microbon test of (a) 1.0% silane-2.5% wt. MC and (b) 1.0% silane-12.7% wt. MC ...... 171 Figure 4.54: SEM images showing surface morphology after microbond test of (a) 3.0% NaOH & 1.0% silane-2.4% wt. MC and (b) 3.0% NaOH & 1.0% silane-13.0% wt. MC...... 172 Figure 4.55: Influence of MC on fibre volume fraction, 푉푓 of composites...... 173 Figure 4.56: Storage modulus of treated composites produced at different MC; (a) 3.0% NaOH, 1.0% silane and 3.0% NaOH & 1.0% silane...... 175 Figure 4.57: Loss modulus of treated composites produced at different MC; (a) 3.0% NaOH, 1.0% silane d and 3.0% NaOH & 1.0% silane...... 176 Figure 4.58: Tan delta of treated composites produced with different MC; (a) 3.0% NaOH treated, 1.0% silane treated and 3.0% NaOH & 1.0% silane...... 178 Figure 4.59: Interfacial strength indicator, B of untreated and chemically treated as a function of MC...... 179 10

Figure 4.60: Typical flexural stress-strain curves of (a) 3.0% NaOH, (b) 1.0% silane and (c) 3.0% NaOH &1.0% silane samples. The curves were normalised at 30% 푉푓...... 181 Figure 4.61: Influence of MC on the (a) tensile strength and (b) tensile modulus of treated composites. Data points were normalised at 30% Vf...... 183 Figure 4.62: SEM imaging of untreated and treated composites showing internal failure in the fibre due to MC...... 185 Figure 4.63: Typical flexural stress-strain curves of (a) 3.0% NaOH, (b) 1.0% silane and (c) 3.0% NaOH & 1.0% silane samples. The curves were normalised at 30% 푉푓...... 187 Figure 4.64: Influence of MC on (a) flexural strength and (b) flexural modulus of treated composites. Data points are normalised at 30% 푉푓...... 189 Figure 4.65: SEM images of the fracture surfaces of treated composites: (a) 3.0% NaOH, (b) 1.0% silane and (c) combined 3.0% NaOH & 1.0% silane...... 190 Figure 4.66: Typical load-extension curves for (a) 3.0% NaOH (b) 1.0% silane and (c) 3.0% NaOH & 1.0% silane...... 192 Figure 4.67: Influences of MC on the R-Curves behaviour on (a) 3.0% NaOH, 1.0% silane and (c) 3.0% NaOH and 1.0% silane treated samples...... 193

Figure 4.68: Comparison of Mode I GIC initiation-VIS values for untreated and treated flax/epoxy composites as a function of MC...... 195

Figure 4.69: Comparison of Mode I GIC propagation values for untreated and treated flax/epoxy composites as a function of MC...... 196 Figure 4.70: Qualitative observation of fibre bridging (white arrows) on Mode I DCB test between untreated and treated composites when the composites were produced with low MC...... 198 Figure 4.71: Qualitative observation of fibre bridging (white arrows) on Mode I DCB test between of untreated and treated composites when the composites were produced at high MC...... 199 Figure 4.72: SEM images of the delaminated surfaces of treated composites: (a) 3.0% NaOH (b) 1.0% silane and (c) combined 3.0% NaOH & 1.0% silane treatments. White arrows locate the peeling of the cell wall on the fibre or matrix surface...... 200 Figure 4.73: Photograph of delaminated treated composites showing colour changes between (a) NaOH (b) silane and (c) combined 3.0% NaOH & 1.0% silane...... 201

List of Tables

Table 1.1: Advantages and disadvantages of natural fibres [3] ...... 23 Table 2.1: Mechanical properties of selected natural fibres with respect to E-glass fibre [29][22][30][31]...... 30 Table 2.2: Mode I interlaminar fracture toughness of composites [34]...... 30 Table 2.3: Chemical compositions and moisture content of natural fibres [44]...... 32 Table 2.4: Influence of drying (105 °C for 14 hours) on the tensile properties of flax fibre [59]...... 39 Table 2.5: Influence of manufacturing technique with respect of fibre volume fraction and porosity of composites [23]...... 49 Table 2.6: Influence of fibre swelling and shrinkage on IFSS...... 55 Table 2.7: Flexural longitudinal and transversal properties of flax/unsaturated polyester composites produced at 0% and 100% RH [121]...... 60 Table 3.1: Resin properties from manufacturer’s datasheet...... 71 Table 4.1: Fickian diffusion coefficient of flax fibre exposed to 10% RH and 90% RH...... 96 Table 4.2: Example of MC calculation. UT1 to UT5 is standardized fabric samples (81 mm x 81 mm)...... 97 Table 4.3: Influence of MC on flax fibre density...... 98 Table 4.4: IFSS data for specimens manufactured with different fibre MC...... 103 Table 4.5: Theoretical 푉푓 of composites based on the experimental data of fibre areal density and fibre density...... 106 Table 4.6: Thickness, density and theoretical 푉푓 of untreated flax/epoxy composites...... 107 Table 4.7: Tensile properties of untreated flax/epoxy composites with various fibres MC during composites fabrication...... 109 Table 4.8: Flexural properties of untreated flax/epoxy composites with various fibres MC during composites fabrication...... 112 Table 4.9: Average ILSS data of untreated flax/epoxy composites with various fibre MC...... 115

Table 4.10: The 푉푓, tan delta peak height, tan delta FWHM, Tg and interfacial strength indicator of untreated flax/epoxy composites...... 119 Table 4.11: Weibull parameters and experimental tensile properties of untreated and treated flax fibre...... 156 Table 4.12: Weibull parameters and average IFSS of treated flax fibre...... 162

Table 4.13: The average tan delta peak height, Tg (peak max) and tan delta FWHM of untreated and treated composites produced at different fibre MC...... 179 Table 4.14: Tensile properties of untreated and treated composites ...... 182

Table 4.15: Flexural properties of untreated and treated composites ...... 188

Table 4.16: The average of GIC initiation-VIS and GIC propagation of composites...... 195

List of Symbols

α = Weibull modulus

β = epoxy-amine conversion

γ = Level of significance

푎 = Delamination length

푏 = Specimen width

퐵 = Interfacial strength indicator

퐷 = Fickian diffusion coefficient

퐷푓 = Diameter of fibre

E = Young’s modulus

E’ = Storage modulus

E” = Loss modulus

퐹 = Load

퐹푚푎푥 = Maximum load

퐹푑푒푏표푛푑푖푛푔 = Debonding load

GIC = Mode I critical strain energy release rate

GIIC = Mode II critical strain energy release rate

h = Specimen thickness

휌푐 = Density of composites

휌푓 = Fibre density

휌푓푙푎푥 = Density of the flax fibre

휌푙 = Density of rapeseed oil

휌푓푐 = Density of fibre according to the moisture content

휌푓0 = Density of dried fibre

휌푠푎푐 = Fibre areal density

휌푤 = Density of distilled water

MC0.5 = Half of the equilibrium MC

푀푡 = Moisture content at time, 푡

푀푒푞 = Equilibrium of moisture content

Tg = Glass transition

푀푓푎푏푟푖푐 = Weight of fabric

푀푐표푛푑푖푡푖표푛푒푑 = Weight of conditioned fabric

푀푑푟푖푒푑 = Weight of dried fabric

푚푎푓 = Weight of fibre in air

푚푙푓 = Weight of fibre immersed in the rapeseed oil

푚푎푐 = Weight of composite sample immersed in air

푚푤푐 = Weight of composite sample immersed in distilled water

푁 = Number of fabric layer

R1, R2 R3 = Remainders of the molecule

σ0 = Characteristic tensile strength

휎푚 = Tensile strength of matrix

푙푒 = Embedded length

푙푓푎푏푟푖푐 = Length of fabric

푙푠 = Straightened yarn length

훿 = Displacement

휏 = Shear strength

휏푚 = Matrix shear strength

휏0 = Characteristic shear strength

푉푓 = Fibre volume fraction

List of Abbreviations

APTMS Aminopropyltrimethoxysilane

APS Aminopropyltriethoxysilane

ASTM American Society for Testing and Materials

ATR Attenuated Total Reflectance

BVID Barely Visible Impact Damage

CAI Compression After Impact

CTE Coefficient of Thermal Expansion

DCB Double Cantilevered Beam

DGEBA Diglycidyl Ether Of Bisphenol-A

DMA Dynamic Mechanical Analysis

DSC Differential Scanning Calorimetry

DVS Dynamic Vapour Sorption

ENF End Notch Flexural

EU European Union

FRP Fibre Reinforced Polymer

FTIR Fourier Transform Infrared Spectroscopy

FWHM Full Width at Half Maximum

HCL Hydrochloric Acid

HDPE High Density Polyethylene

IFSS Interfacial Shear Strength

ILFT Interlaminar Fracture Toughness

ILSS Interlaminar Shear Strength

IUPAC International Union Of Pure and Applied Chemistry

MBT Modified Beam Theory

MC Moisture Content

MFA Microfibrils Orientation Angle

NaOH Natrium Hydroxide

NDT Non-Destructive Technique

NL Non-linearity

NVH Noise, Vibration and Harshness

PEA Polyesteramide

PHB Polyhydroxybutyrate

PLA Polyactides

RH Relative Humidity

RI Resin Infusion

RTM Resin Transfer Moulding

SEM Scanning Electron Microscope

SFTT Single Fibre Tensile Test

UN United Nation

UT Untreated Fibre

VIS Visual Observation wt.% Weight Percentage

XPS X-Ray Photoelectron Spectroscopy

Abstract The University of Manchester Abdul Hakim Abdullah Doctor of Philosophy The Effects of Fibre Moisture Content During Manufacturing on the Properties of Flax/Epoxy Composites March 28, 2020

Due to increased awareness of environmental issues, considerable attention has recently been given to the development of natural fibre reinforced polymer composites. Flax fibres and other natural fibres have various advantages; such as low cost, abundance, renewability and can offer respectable mechanical properties. However, flax fibre is a hygroscopic material and concern has been raised for composites manufacturing at high fibre moisture content (MC) which could have a negative effect on the mechanical properties of flax fibre composites. In this study, the relationship between weight percentage (wt.%) fibre MC and the mechanical performance of flax/epoxy composites was investigated. Using dynamic vapour sorption (DVS) analysis, an increase in fibre MC with increasing ambient relative humidity (RH) was observed. The flax fibre absorbed as high as 14.8 wt.% MC. A novel method of estimating MC was developed using a standardized flax fabric sample which was useful for monitoring fibre MC during composites manufacturing. The fibre-matrix interfacial shear strength (IFSS) was evaluated using the microbond technique and IFSS was found to reduce consistently with increasing fibre MC for untreated-fibre composites. Tensile, flexural and low velocity impact tests were conducted on the flax/epoxy composites, and in general properties reduced with increasing fibre MC. Scanning electron microscopy (SEM) also indicated weak fibre-matrix interface of composites as shown by fibre pull-out and the formation of gaps between the fibres and the matrix. Dynamic mechanical analysis (DMA) showed high fibre MC plasticised the epoxy matrix, with values of glass transition temperature reducing from 162.0 ºC to 156.6 ºC. The measured propagation Mode-I interlaminar fracture energy, GIC propagation, increased from 1.81 kJ/m2 (at 2.3 wt.% fibre MC) to 2.48 kJ/m2 (at 6.6 wt.% fibre MC) due to toughening of the matrix with increasing MC, resulting in a larger plastic zone at the crack tip. A mechanism of moisture interaction at the fibre-matrix interface during untreated flax/epoxy composites manufacturing is proposed, suggesting weaker fibre-matrix interface resulted from disruption of crosslink formation between reactive groups on the surface of the moisturized fibre and the epoxy matrix. Alkaline and silane fibre surface treatments, and combinations of both, were found to reduce fibre moisture absorption between 10-70% RH. Microbond tests showed that IFSS was generally improved after surface treatments with 3.0% NaOH giving the highest values. In most cases, the tensile and flexural properties of treated composites were lower than untreated composites due to loss of tensile properties of flax fibre after being treated, which the weakened internal structure of fibres as a result of swelling. Mode I GIC propagation of treated composites was significantly improved compared to untreated composites at low MC due to a stronger fibre matrix interface.

Declaration

I declare that no portion of the work referred to in the thesis has been submitted in support of an application for another degree or qualification of this or any other university or other institute of learning

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Acknowledgement

I wish to express my gratitude to my supervisor, Dr. Arthur Wilkinson, for his valuable guidance, encouragement, comments and patience throughout the course of this study. I would like to also express my great to Prof. Constantinos Soutis who have given me lots of advice and useful critics on my research work. I appreciated all the staffs based in National Composites Certification and Evaluation Facility (NCCEF): Dr. Alan Nesbitt, Christopher Cowan, Stephen Cowling, Andrij Zadoroshnyj, Dr. Zijie Wu, Dr. Mayank Gautam, Dr. Kinjal Patel Kumar, Dr. Victor Ramirez Elias and Dr. Shankachur Roy.

Special appreciation goes to my lovely mother, Rusizah Md. Taib and my in-law parents: Hamdan Abdullah and Solehah Ahmad for their support, assistance and encouragement. Special thanks should be given to my siblings; Dr. Nurhidayah, Ahmad Suhaimi, Nurnadiah and Nurdiyanah for their commitment and resources to care our beloved late father, Abdullah Kasim while I was thousand miles away from the country.

Thank you also to my lovely wife, Izzati Hamdan, for her love, support, understanding and patience through thick and thin in finishing this study. Also, my children; Muhammad Haris, Hana Humaira and Muhamad Aqil (Manchester born baby) for giving me the tears and joys all these years.

Thanks also to the Malaysian Community who helped me directly or indirectly during my stay in Manchester, United Kingdom. Finally, I would like to thank to Universiti Teknologi MARA (UiTM) and the Government of Malaysia for sponsoring this PhD study.

Chapter 1 Introduction

1.1 Overview

Applications of modern fibre reinforced polymer (FRP) composites began in 1930s where glass fibre reinforced composites were used as construction materials in boat hulls, radar domes and car body sections [1]. Since then, composites continued to grow to replace conventional metallic structures particularly in the automotive, aerospace and construction industries. Amongst the fibres used for FRP composites, carbon and aramid fibres are superior to glass in strength and stiffness. However, they are much more expensive and glass fibre is normally used in less- demanding general applications.

Environment concerns have become major topics of interest over the last decades. Natural fibres are often considered as candidates for glass fibre replacement because of their renewability, recyclability and sustainability [2]. Fibres such as flax, kenaf, hemp and jute offer acceptable specific properties whilst are lower in cost (presently priced at one-third of the cost of glass fibres or less) and lower in weight (typically half the density of glass fibre)[3]. The advantages and disadvantages of natural fibres are summarized in Table 1.1.

Table 1.1: Advantages and disadvantages of natural fibres [3]

Items Advantages Disadvantages  Low density, thus low weight  Fibres absorb moisture that  Higher specific strength and causes swelling Physical- stiffness than glass fibre  Lower strength properties than mechanical  Good thermal properties glass fibre composites,

particularly on impact strength  Odour generation due to degradation process  Non-abrasive effect aver  The maximum processing screws and other metallic temperatures are limited, Processing parts especially in relation to glass and  No-harmful processing, no fibre Manufacturing tool wear and no skin irritation  It is a renewable resource,  Relatively low durability, due to and is therefore an fungus attack, weathering, etc. inexhaustible supply  Variable quality, depending on  Production energy is only 1/3 unpredictable influences such as of that for glass fibres weather Environmental  The amount of CO2 that the plants absorb during their growth is the same as that given off when they are decomposed It is reported that the European Union (EU) market for FRP composites was 2.4 million tonnes in 2010 [4]. While glass fibres composites constitute the majority of these, the uses of plant fibres are relatively small at 1.9% of the total market as shown in Figure 1.1. However, according to a market size prediction published by Grand View Research Inc. [5] in June 2016, the market for natural fibres is expected to reach USD 10.89 billion by 2024 from USD 3.50 billion in 2015, a rise of 74% in 10 years. Therefore, it is expected soon that more manufacturers will embrace plant fibres in their design considerations and products. The importance of plant fibre industries was voiced by the United Nation (UN) during the UN General Assembly in 2006. The UN declared 2009 as the International Year of Natural Fibres in a hope to raise the profile of natural fibres [6].

Figure 1.1: Plant fibres FRP market share in 2010 out of 2.4 million tonnes [4].

1.2 Future of Natural Fibre Based Composites

Current utilization of natural fibre composites has little or no structural role because they have lower mechanical strength than man-made fibres. Natural fibres have been employed most commonly in automotive interior parts such as door panels, seat backs, headliners, package trays and dashboards as shown in Figure 1.2(a) [7]. The second biggest market is the construction industry. FRP decking, wall panels and composite window frames are some of the examples [8]. Other non- structural applications of natural fibre composites are consumer goods such as toys, food packaging and cases for electronic devices [9].

There have been increased research efforts to employ natural fibres composites in structural or semi-structural components. The general conclusion result of this research is positive, proposing that natural fibres as glass fibre replacements could be possible. Examples of these applications are automotive bumper beams [10], small scale wind turbine blades [4], interiors of helicopters [11], bulletproof vests for military [18], boat hulls [12], canoes [13] and bicycles [14].

Figure 1.2: Examples of natural fibre composites in non-structural application; (a) automotive interior parts [7] and effort for semi-structural/structural applications; (b) a canoe [13] and (c) a bicycle [14].

1.3 Research Motivation

As previously explained, the benefits of natural fibre composites could offer new solutions for environmental issues. The continuing development and improvement of the natural fibre composites could bring these materials on a par with their man-made composite counterparts. However, an issue with plant fibres is that they are known to absorb moisture from the surrounding environment, which leads to poor interfacial bonding with some polymer matrices [15]. Studies of moisture uptake in natural fibre composites is well reported in the literature but typically only deal with post-manufacturing moisture absorption, in which the cured composite is immersed in water [16] or exposed to a weathering condition for a period of time [17]. However, few studies have been undertaken on the effect of moisture content (MC) in natural fibres during composites fabrication. Humidity in a composite manufacturing facility is usually not well controllable and this could lead to high MC in natural fibres used for polymer composites. In

addition, although epoxy is not hygroscopic the amine hardener used for curing is miscible with water [18] and the absorbed water eventually could affect the final mechanical properties of the cured epoxy [19]. Therefore, further study is required of the fibre MC during composite manufacture and its impact on the interfacial bonding and the final mechanical performance of natural fibre composites.

1.4 Research Objectives[AHA1]

In this study, woven flax fabric is used as the reinforcement and epoxy resin as the matrix. The specific objectives are outline below:-

1. To understand the influence of moisture content on flax fibre properties.

2. To investigate the interfacial bonding strength and mechanical properties of flax/epoxy composites manufactured with different fibre moisture content

3. To understand the effect of surface treatments; alkaline, silane and combined

silane-alkalize on both moisture content and flax fibre properties[AHA2].

4. To investigate the effect of surface treatments on the interfacial bonding strength and mechanical properties of flax/epoxy composites manufactured with different fibre moisture content.

1.5 Thesis Outlines

This thesis is divided into five chapters, which are briefly described as follows:-

 Chapter 1: This chapter provides brief information about natural fibre composites and their applications. It is also covers research motivation and the objectives of the current work.

 Chapter 2: This chapter presents a comprehensive review of the relevant literature. This chapter concentrates on the topics of moisture content in

natural fibres, epoxy matrices, fibre/matrix interface properties and fabrication of natural fibre composites.

 Chapter 3: The materials and experimental procedures are explained in this chapter, including fibre conditioning and composites fabrication. This chapter also describes characterization methods and the equipment used for testing.

 Chapter 4: The experimental results are presented and discussed in this chapter. The first part of this chapter covers the characterization of untreated flax fibre, studies of the fibre/matrix interface and the mechanical performance of untreated flax/epoxy composites. A mechanism of moisture interaction of untreated flax/composites during composites fabrication is also proposed. The second part of this chapter deals with the influence of surface treatments on flax fibre properties, the fibre/matrix interface and their composites.

 Chapter 5: The conclusions of this study are presented and suggestions for future work are proposed in this chapter.

Chapter 2 Literature Review

In this chapter, a review of the literature is presented for better understanding and justification of the importance of this work.

2.1 Natural Fibres

Due to environmental concerns, man-made fibrous reinforcements are being replaced increasingly with natural fibres in composite materials. Natural fibre is a fibre taken from an animal, mineral or plant source as shown in Figure 2.1. Animal fibres such as silk, rabbit hair and camel hair consist largely of particular proteins in their construction. In contrast, plant (lignocellulosic) fibres are mainly made up from cellulose in their molecular structure and can be further divided into three categories based on their respective origin or location in a plant [20]. Plant fibres may be extracted from leaves, seeds or bast (stalks). Bast fibres (such as flax, hemp, kenaf, ramie and jute) are taken from the outer stem layer or skin of a plant [21][22].

Figure 2.1: Classification of natural fibres [23].

2.1.1 Flax Fibres and Their Composites

Among the bast fibres, flax (Linum usitatissimum L.) has gained popularity for use as reinforcement in polymer composites. Flax is mainly grown in Canada, Europe and New Zealand for its linseed oil and fibre [24]–[26]. Figure 2.2 shows a schematic of the morphology of a flax fibre from the flax stem down to microfibrils. From the flax stem, scutching (beating and scraping of retted stalks) and hackling (combing) processes produce technical fibres. A technical fibre consists from 10 to 40 elementary fibres and is normally with 2-5 cm in length [27]. Microfibrils are the smallest elements in a flax fibre, having a diameter between 1-4 nm.

Figure 2.2: Schematic morphology of flax fibre [28].

A comparison of the mechanical properties of natural fibres is shown in Table 2.1. Strength and stiffness of natural fibres are generally lower than a E-glass fibre. Due to low fibre density, however, the specific strength and modulus for some natural fibres are somewhere higher than E-glass fibre. Flax fibre is amongst the strongest natural fibres as compared to jute, hemp and kenaf. Hence, the uses of flax fibre could provide superior mechanical properties in natural fibre polymer composites.

Table 2.1: Mechanical properties of selected natural fibres with respect to E-glass

fibre [22][29]–[31][AHA3].

Fibre Properties E-glass Flax Hemp Jute Kenaf

3 Fibre density, 휌푓 (g/cm ) 2.5 1.40-1.53 1.14-1.52 1.19-1.52 1.2-1.4

Young modulus, E (GPa) 70-73 27.6-80 30-70 13-60 14–38

Tensile strength (MPa) 200-3500 88-1500 500-920 393-860 295-500

Speciﬁc modulus, (GPa/g/cm3) 28 18.4-38 26.3-52.6 10-39 12–32

Speciﬁc strength, (GPa/g/cm3) 800 - 1400 200 - 1000 370 - 600 300 - 610 -

Elongation at failure [%] 2.5-3.7 1.2-3.27 1.6-4 1.6-2 1.6-6.9

Mechanical properties of flax fibre reinforced polymer composites are widely reported by many authors indicating the potential of flax fibres as reinforcements in polymer composites. Rodríguez et al. [32] studied various natural fibre (flax, jute, sisal) reinforced polyester composites produced via vacuum infusion. They showed that tensile strength and impact strength of flax/polyester composites were the highest of the natural fibres composites. Although flax composites had lower specific strength than glass composites, the study by Oksman [33] found that unidirectional Arctic flax/epoxy composites had greater specific modulus (29 GPa/gcm-3) than the glass/epoxy composites (18 GPa/gcm-3). In term of composites fracture toughness, recent research reported by Bensadoun et al. [34] indicates that the Mode I interlaminar fracture toughness values of flax/epoxy composites are higher than carbon/epoxy composites as shown in Table 2.2.

Table 2.2: Mode I interlaminar fracture toughness of composites [34].

Mode I Interlaminar Fracture Toughness, (J/m2) Composites Initiation Propagation Unidirectional flax/epoxy [90,0]2s 496 663

Unidirectional flax/epoxy [0,90]2s 655 1086

Unidirectional carbon/epoxy 100–200 150–250

Unidirectional glass/epoxy 243–268 N/A

Another interesting characteristic of natural fibres is that they have a good damping and vibration properties which is highly preferable for noise, vibrational and harshness (NVH) application in the automotive sector. Dong and Yan [35] reported that natural fibres has better sound absorption coefficients at any given frequency as compared to glass and carbon fibres as shown in Figure 2.3(a). Several authors also reported that vibration and damping of natural fibres comparable or even superior than glass composites counterpart [36]–[38]. Figure 2.3(b) shows the logarithmic damping decrement of flax and hemp fibre which is higher than glass fibre.[AHA4]

1.0

Ramie

Jute 0.8 Flax Glass Carbon 0.6

0.4

0.2 Sound Sound Absorption Coefficient

0.0 0 500 1000 1500 2000 Frequency Frequency (Hz) (Hz) (a) (b)

[AHA5]

Figure 2.3: Comparison between man-made and natural fibres in term of (a) sound absorption coefficient [35] and logarithmic damping decrement [38].

Higher strength and stiffness of composites can be achieved by increasing fibre volume fraction, 푉푓. Singleton et al. [39] showed that tensile strength and stiffness of flax mat/recycled high density polyethylene (HDPE) composites increased with increasing 푉푓 (0%, 10%, 18%, 20% and 30%). However, there is a maximum 푉푓 that can be attained and beyond this threshold value there is no appreciable strength increase due to high porosity (poor fibre wetting and matrix impregnation), causing inefficient stress transfer at the interface [23]. Shah et al.

[40] proposed that 33.1% was the theoretical maximum 푉푓 for unidirectional flax composites.

2.1.2 Chemical Composition and Internal Structure of Natural Fibre

The main chemical components of flax fibres (Table 2.3) are cellulose, hemicellulose, pectin and lignin but they may also have trace amounts of wax and several water-soluble compounds in their cell wall structure. These essential components provide structural strength/rigidity and provide protection against swelling in the cell wall [41]–[43]. The drawback of natural fibres is that they absorb moisture from the environment and this can be problematical for composites while in-service and also during composites fabrication. For example, flax fibres hold moisture at about 8-12 % at ambient humidity (Table 2.3).

Table 2.3: Chemical compositions and moisture content of natural fibres [44].

Cellulose Hemicellulose Lignin Pectin Moisture Waxes MFA* Fibres (%) (%) (%) (%) Content (%) (%) (º) Flax 71 18.6–20.6 2.2 2.3 8–12 1.7 5–10 Hemp 70–74 17.9–22.4 3.7–5.7 0.9 6.2–12 0.8 2–6.2 Jute 61–71.5 13.6–20.4 12–13 0.2 12.5–13.7 0.5 8 Kenaf 45–57 21.5 8–13 3–5 - - - Ramie 68.6–76.2 13.1–16.7 0.6–0.7 1.9 7.5–17 0.3 7.5 Sisal 66–78 10–14 10–14 10 10–22 2 10–22

* Microfibril orientation angles

The structure of natural fibres is complicated because they consist of several elements that perform different functions. Individual elementary flax fibres are held together by middle lamella which is a pectin-rich region (Figure 2.4(a)). The internal structure of an elementary flax fibre is presented in Figure 2.4(b). A cell wall (elementary fibre) is a collection of two distinct walls; namely, the primary (P) and secondary (S) walls. The secondary wall (S) is then divided with sub-layers: the outer layer (S1); middle layer (S2); and the inner layer (S3). A cell wall has a hollow cavity (lumen) at the centre. The luminal areas of flax and hemp are in the range of 0-5% of total cross-sectional area.

Microfibrils are mainly composed of cellulose and serve as the primary load- carrying elements in the fibres. Microfibrils have the form of a helical structure

embedded in the two-phase (lignin-hemicellulose matrix) amorphous matrix [45]. The mechanical properties of natural fibres are mainly attributed to the middle layer S2 wall where the microfibrils are held [46]. The cellulose content and microfibril orientation angle (MFA) control the mechanical properties of fibres [22][47].

(a) Middle lamella Lumen

(b)

S1 =10º

Secondary S2 cell wall

P Primary cell wall

Figure 2.4: Schematic (a) cross section of a technical fibre and (b) internal structure of an elementary flax fibre [48].

2.1.3 The Origin of Hygroscopic Behaviour in Natural Fibres

The existence of a large number of hydroxyl groups (OH) in natural fibres contributes to their behaviour as hydrophilic and polar materials [41][43] but their water sorption affinity is also related to other oxygen-containing groups that promote hydrogen bonding [50]. As the cell wall absorbs moisture, the water molecules occupy the spaces between the microfibrils and within the amorphous matrix, forcing microfibrils apart and cause swelling in the cell wall. This space is referred to as the transient microcapillary network (also known as pores). Water molecules held by hydroxyl groups is called bound water or monolayer water. Water molecules can be present in the form of multilayer water within transient

microcapillaries that are not closely related to hydroxyl groups. Water present in any inner volumetric structures such as cell cavities (lumen) is called free water. It is reported the lumen contains up to 12% of the total moisture regain for flax and hemp [22][44]–[46].

Although natural fibres are considered highly hygroscopic not all constituents contribute to the moisture affinity. Cellulose is the major constituent of cell walls followed by hemicellulose. It contains linear condensation polymer chains composed of d-glucose units (–CH2OH) and connected with β-1, 4-glycosidic linkages (covalent bonds) [22]. Cellulose has both crystalline and amorphous regions in its molecular structure of which the former contains up to 80% crystalline regions [53][54]. Previous investigations showed that the crystalline regions are inaccessible to water molecules but they may penetrate cellulose through the amorphous regions [2][49].

In contrast, hemicellulose has a random, amorphous structure, making it easily accessible to water molecules [56]. Lignin is considered to be hydrophobic in nature due to the low concentration of hydroxyl (OH) groups in its chemical chains. Lignin in middle lamella layer could establish resistance to the moisture accessible in the cell wall. Lignin partly bonds with hemicellulose-cellulose via covalent bonding, acting as a compatibilizer or coupling agent. This attributes to the rigidity of fibres [22][57]. Pectin can be found in primary and secondary layers composed of carboxyl function where hydrogen bonds may interact with polar water molecules [50][52]. Akram and co-workers [50] point out that hemicellulose is mainly responsible for water absorption, followed by cellulose and others compositions. However, Baley and co-author [59] claimed that moisture absorption for flax fibres is mainly attributed to the pectin at the fibre surface or within the cellulose microfibril matrix.

2.1.3.1 Moisture Behaviour in Natural Fibres

The mechanism of moisture diffusion in natural fibres can be represented by Fickian patterns which can be recognized by the shape of the absorption curve [60]. This behaviour is reported to occurs in both natural fibres and their polymer composites [60][61][38]. Figure 2.5 shows moisture absorption of flax (Green and 34

Duralin) for which the curve increase linearly from the beginning. In Fickian diffusion moisture diffuses from highest to lowest concentrations generated by the concentration gradient [62]. Therefore, the Fickian diffusion coefficient, D is calculated by the following equation [63]:-

1 ⁄2 2.1 푀푡 퐷 = 4 ( 2) 푀푒푞 휋ℎ

Where 푀푒푞 is the equilibrium of moisture content, 푀푡 is the moisture content at time,푡 and h is the sample thickness.

Green

Duralin MC(%)

0.5 Time (min )

Figure 2.5: Moisture absorption in flax fibre (Green and Duralin) [64].

2.1.3.2 Moisture-Induced Swelling

It is known the when a natural fibre absorbs moisture; it swells larger in the radial direction than the axial direction [65]. Swelling of bamboo fibre in the width direction was reported to be around 6% when the samples were exposed to 90% RH. Pucci et al. [66] studied the swelling of elementary flax fibre under dry (as received) and wet conditions. Flax fibres were wetted with water and the diameter expansion was observed using a microscope. Swelling ratio (wet fibre diameter/dry fibre diameter) was calculated and was found to be 1.27 ± 0.13. Moisture absorption in cured natural fibre composites has been reported to degrade the fibre/matrix interface and, ultimately, lower the mechanical properties of the composites [67][57]. As explained by Azwa et al. [62], a swelled fibre develops internal stress at the interface, causing microcracking in the matrix around the swollen fibres. In the

case of moisture presence during composite fabrication, the swelling and shrinkage could also affect the interface properties of composites and this is presented in Section 2.5.3.

2.1.3.3 Evaluating Moisture Content in Natural Fibre

The conventional methods of using desiccators, saturated salt solutions or simple weight gain measurements for evaluating moisture content (MC) are time- consuming, expensive, labour extensive and could provide inaccurate data. Recently, more studies have been using the Dynamic Vapour Sorption (DVS) analyser for materials such as polymers, ceramics, food industries, pharmaceuticals and natural fibres [59][62]–[66]. DVS uses gravimetric measurements to monitor the mass changes as a function of time when a specimen is exposed to a series of steps changes in the RH. The MC will correspond to the RH but this takes a definite time to achieve. DVS performs real-time measurements, yields reproducibility data and gives accurate data analysis.

The study of moisture behaviour in natural fibres using a DVS analyser was reported by several authors [52][59][63]. The interactions of atmospheric water vapour and water molecules located within the cell wall pores can be described into 6 distinct shapes with reference of pore size developed by The International Union of Pure and Applied Chemistry (IUPAC) [62][65][66]. The moisture sorption isotherms of natural fibres are always associated with sigmoid IUPAC Type II with strong moisture uptake in the beginning and towards the end of the curves [70]. Figure 2.6 showed typical absorption and desorption curves for several natural fibres i.e flax, hemp, jute and sisal. It can be seen at higher humidity; hemp, sisal and jute have higher MC compared to flax fibre. Therefore, the RH values are not interchangeable with MC. It is observed in Figure 2.6 that absorption-desorption processes are found to be imbalanced for many cases and open hysteresis loops are observed [52]. During absorption, the incoming water molecules create nanopores with the resulting creation of new internal surfaces. In contrast, desorption is delayed and affecting nanopores to collapse and this would be explanation for the hysteresis loops [58].

25 25 (a) (b) 20 20

15 Flax 15 Hemp

10 10 MC (%) MC 5 5 Absorption 0 0 Desorption

0 20 40 60 80 100 0 20 40 60 80 100 25 25 (c) (d) 20 20 15 Jute 15 Sisal 10 10

MC (%) MC 5 5 0 0

0 20 40 60 80 100 0 20 40 60 80 100 RH (%) RH (%)

Figure 2.6: Absorption and desorption curves corresponding to RH changes for (a) flax fibre (b) hemp fibre (c) jute fibre and (d) sisal. Curves are redrawn from [58].

2.1.3.4 Changes in Fibre Density Due to Absorbed Moisture

Fibre density is one of the parameters used to determine the fibre volume fraction, Vf of composites materials. In a review paper, Summerscales et al. [21] explained the determination of Vf for natural fibre composites can be challenging because of the MC and they recommended a graph of fibre density against MC could provide useful data. Messiry [73] proposed a simple mathematical expression to include the weight of MC in the fibre density calculation. The new density of fibre according to the moisture content 휌푓푐 is given by;

(1 + 푀퐶) 휌 = 푓푐 1 2.2 [( ) + 푀퐶] 휌푓0

Where 휌푓0 is the density of dried fibre and MC is moisture content. Using equation 2.2, Messiry [73] showed graphical relations between the density of fibre

as a function of MC as shown in Figure 2.7. The dried fibre has the highest density and the inclusion of moisture causes the density of fibre to reduce. Although this estimation of fibre density seems to be encouraging, the mathematical expression was not supported by experimental studies and this could present bias in Vf and void contents.

1 1.1 1.8 1.2 1.3 1.4

) 1.5 3 1.6 1.6 1.7 1.8 1.4

Fibre Density (g/cm Density Fibre 1.2

1.0

0 5 10 15 20 25 30 35 40 45 MC (%) [AHA6]

Figure 2.7: Messiry’s graphical representation of fibre density versus fibre MC. Curve is redrawn from [73].

2.1.3.5 Fibre Density Measurement

There are two method of density assessment typically reported in the literature, which are the pycnometer and buoyancy methods. The former is either a liquid or helium pycnometer. The accuracy and effectiveness of this density measurement of natural fibres was discussed by Truong [74]. Helium pycnometry was found to be effective to provide reproducible density of flax fibre with small deviations. Shah [45] explained that the helium pycnometer could measure the true density of fibres because the instrument measures the amount of gas displacement of porous material such as natural fibres. The porosity of natural fibres originates from the lumens or air spacing between the elementary fibres if the fibres are twisted to form fibre yarn. The density of natural fibres however depends on the MC and the use of helium pycnometry is not applicable for this purpose.

Buoyancy methods have been favoured by many researchers due to the inexpensive equipment and low cost of immersion liquids. Using the gravimetric

technique by measuring the mass of a sample in air and in an immersion liquid, the density of natural fibres across the MC could be measured. The accuracy of this method is dependent on the immersion fluid. Truong et al. [74] reviewed several immersion fluids such as methanol, acetone, ethanol, water and canola oil and compared the results with pycnometry data. They found that canola oil had comparable results to pycnometry. Recently, Amri [75] demonstrated the effect of immersion time between canola oil and distilled water. Although the density of fibres increased as the immersion time increased, the differences in canola oil was not as significant compared to distilled water.

2.1.3.6 Influence of Humidity on the Mechanical Performance of Natural Fibres

As natural fibres are being used increasingly as reinforcements in polymer composites, their mechanical performance with respect to MC is highly important. It is evidently shown that removal of water in natural fibres could lower their tensile strength and Young’s modulus compared with the as-received fibre (Table 2.4) [59][18]. This means the natural fibre are strongest with presence of MC and removing MC prior to composites fabrication might not be ideal to retain optimum performance.

Table 2.4: Influence of drying (105 °C for 14 hours) on the tensile properties of flax fibre [59].

Designation Failure Stress[W7] Young’s Modulus (GPa) Failure Strain (%) (MPa)

As-received 1499 ± 346 64.10 ± 13.65 2.93 ± 0.74

Dried 870 ± 266 59.24 ± 19.36 2.07 ± 0.30

For variations of humidity in natural fibres, Davies and Bruce [76] showed that the tensile moduli of flax and nettles fibres are reduced with increasing % RH while the tensile strengths of the fibres were not affected. However, recently, the majority of studies have shown that the tensile strength and Young’s modulus were reduced when natural fibres were conditioned at lower humidity while the greatest tensile strength was achieved on exposure to medium humidity (close to ambient

RH) [18][64][77][78]. Above this certain threshold value, the literature data indicates that the tensile strength and Young’s modulus levelled-off or dropped, depending on the type of fibre. To illustrate this, experimental data from Thuault and co-workers [77] for flax fibre (Marylin and Hermes) are presented in Figure 2.8(a-b) which clearly show that the tensile strength and Young’s modulus increased considerably up to a threshold value (68% RH). The rapid fall of tensile strength at higher humidity (85% RH) was due to disorganisation of the microfibrils network when a high amount of water penetrates the fibre. The threshold value at which the greatest tensile properties were achieved was reported at 50% RH for bamboo [18] and 50% RH for hemp fibre [78]. In contrast, due to plasticization in natural fibres, the tensile strain was reported to increase over the full RH range as shown in Figure 2.8(c) [77].

(a) (b)

Young’s Young’s Modulus (GPa) Tensile Strength (MPa) RH (%) RH (%)

(c)

Tensile (%) Strain

RH (%)

Figure 2.8: Tensile properties of flax fibre (Marylin and Hermes) with respect of RH (a) tensile strength (b) Young’s modulus and (c) tensile strain [77].

2.2 Polymeric Matrices

The polymeric matrices used in natural fibre composites can be either thermoset or thermoplastic [25]. More recently, research of natural fibres composites with biopolymer matrices derived from renewable resources have been reported [79][80]. Examples of these biopolymers are polyesteramide (PEA), polyhydroxybutyrate (PHB), polyactides (PLA)[8]. However, most of these biopolymers are more expensive than common non-renewable matrix polymers [25]. Thermoset resins, such as polyester, vinyl ester, and epoxy resins are often favourable because thermoset-based composites generally perform better at higher temperatures than thermoplastic-based composites [23]. Once a thermoset resin is cured, it cannot be melted or reformed due to the formation of a three-dimensional cross-linked molecular network [8]. Advantages of thermoset resins include low processing temperatures which is suited to natural fibres, lower viscosity for better impregnation and fibre wettability, leading to better interface formation with hydrophilic natural fibres [45].

2.2.1 Epoxy Resin

Among the thermoset polymers, epoxy resins are the most common for high performance composites and have been used as composite matrices, adhesives and laminates for more than 40 years [81]. Epoxy resins are characterized by the epoxide functional group, a three-membered ring with two carbons and one oxygen atom (Figure 2.9). Depending on the desired properties, the ring can be added at the resin chain ends and epoxy resins for composite applications are highly aromatic in order to benefit from both increased strength and stiffness [82]. The most widely used epoxy resin are based on the diglycidyl ether of bisphenol A (DGEBA) [83].

In order to establish the crosslinking reaction, curing agents (hardeners) are used. Hardeners can be amine-, acid-, isocyanides- or phenol-based [83]. Amine- based curing agents are most commonly used for the epoxy reaction and these are characterized by NH2 groups, typically at the ends of the molecule (Figure 2.9). Curing reactions can be done either at room temperature (catalysed) or assisted by heating (heat-activated). There are three type of organic compounds for amine-based

curing agents; aliphatic, cycloaliphatic and aromatic. Based on their nucleophilic character, aliphatic amines are less reactive and are used when the curing process is required at low temperature. Adhesives and coating applications are applications for which curing is desired at room temperature. In contrast, aromatic amines are more reactive than cycloaliphatic amines and are generally used in matrix systems for composite applications [83].

Figure 2.9 shows the three major reactions during curing between the epoxy group of a resin and an amine group of the hardener. The curing reaction begins with the epoxide group reacting with reactive hydrogen atom on the primary amine group, forming a secondary amine and generating a hydroxyl group. The secondary amine can then react with an epoxide group to form a tertiary amine and also generating a hydroxyl group. As the formation of hydroxyl groups proceeds, an etherification reaction with epoxy groups can occur and large numbers hydroxyl groups can catalyst the epoxy-amine cure reaction [84][85].

Epoxide group reaction with a primary amine

H O H OH

R1—N—H + H2C—CH—R3 R1—N—CH2—CH—R3

Primary Epoxide Secondary Amine Amine

Epoxide group reaction with a secondary amine

H O R2 OH

R1—N—R2 + H2C—CH—R3 R1—N—CH2—CH—R3 Secondary Epoxide Tertiary Amine Amine

Epoxy-hydroxy reaction (also called etherification)

R2 OH O OH

R1—N—CH2—CH—R3 + H2C—CH—R3 R3—CH—CH2

R2 O

R1—N—CH2—CH

Figure 2.9: Crosslinking reactions between an epoxy group of the resin and an amine

group of the hardener [84]. R1, R2 and R3 present the remainders of the molecules.

2.2.2 Moisture Contamination in Uncured Epoxy Resin

Epoxy resin in general is not miscible with water but the amine-based hardener is well known to attract water molecules [86]-[87]. A solubility study of 19 amines with water was reported by Stephenson [88]. He showed that there was a great decrease in solubility with increasing molecular weight of the amines. The source of moisture contamination in the uncured epoxy is from ambient atmosphere exposure [89]-[90]. However, in some studies discussed in this section to study the effect of moisture on polymerisation, water is introduced purposely into the 43

epoxy/hardener mixture. The detrimental effect of water contamination in amine- cured epoxy resin is most commonly seen in coatings and flooring application [87]. The coating surface will develop oily spots, a defect known as ‘amine blush’, if epoxy polymerization is initiated in a humid environment. Figure 2.10 shows a schematic of the cure reaction between epoxy and amine hardener in the presence of water. Water is seen to provide additional hydroxyl group during the cure reaction [84].

H—O—H

O - O + R NH + H C—CH—R + H O R N — CH — CH — 2 2 3 2 2 2 R

OH OH

+ R2N — CH2—CH—R3 + H2O R2NH — CH2 — CH —R3 + - HO Figure 2.10: Crosslinking reactions between an epoxy group of the resin and an

amine group of the hardener in the presence of water [84]. R1, R2 and

R3 present the remainders of the molecule.

2.2.2.1 Effect of Moisture Content on Cure Reaction and Degree of Cure of Epoxy

Moisture and others hydroxyl containing compounds such as alcohols and phenols induce a catalytic effect to the polymerization of epoxy by contributing to the ring opening of epoxides [19][89][91][92]. The studies of water contamination in an epoxy resin or epoxy adhesive during cure were reported by several authors [19][80][85]–[87]. The water was purposely added into the hardener and mixed with the epoxy resin for these studies. Choi and co-workers [91] employed fourier transform infrared spectroscopy (FTIR) to characterize the cure reaction of Epon 836 epoxy resin with poly(oxpropylene)diamine hardener whilst Nathan [86] studied the cure reaction of epoxy resin (EPON 862) with diethyltoluenediamine hardener via Differential Scanning Calorimetry (DSC). Their results were similar and agreed that the cure reaction was accelerated in the presence of water. The epoxy-amine

conversion, β, versus time with different amounts of added of water is shown in Figure 2.11(a), showing evidence of cure reaction acceleration [91].

The presence of water during epoxy polymerization was also studied in term of degree of cure. The uses of the glass transition temperature (Tg) is widely accepted as a measurement of degree of cure of thermosetting polymer [92][94][95]. Wu and co-workers [93] studied two adhesive epoxy systems (SikaDur 30 and Epon 828) with moisture addition (2% wt., 4% wt., 6% wt.) during cure. They found that the glass Tg of epoxy increased with water addition (2% wt. and 4% wt.), indicating higher degree of cure (crosslink density) of these samples which is due to the increased mobility of polymer chains. However, they observed that the Tg was reduced with 6 wt% water addition. They explained excess of water remaining in the epoxy could cause plasticization effect and hence, lowering the degree of cure. In another research, Nathan [86] reported that the Tg of epoxy (EPON 862/ diethyltoluenediamine hardener) was consistently decreasing for every increment of weight percentage water addition (0.3% wt., 0.4% wt., 0.9% wt. and 1.7% wt.) as shown in Figure 2.11(b). The same effect was also reported by Chian et al. [92] who studied underfills for flip-chip industry. The epoxy resin used was a cycloaliphatic type and anhydride hardener. They found that the Tg decreased with higher amount of water concentration and longer storage times at the ambient temperature.

(a) (b)

)

Heat Flow Flow (W/g) Heat Conversion ( Conversion Tg

Time (Min) Temperature (K)

Figure 2.11: Influence of water addition on the uncured epoxy; (a) cure reaction -

data from FTIR [91] (b) glass transition, Tg - data from DSC curve [86].

The study of fibre MC in composites during fabrication is rarely reported. Heijden and co-workers [68] studied the cure of epoxy resin in the presence of humidified polyamide 6 nanofibres using DSC. Samples having a fibre content of 26 ± 2 wt% were produced using EPIKOTE Resin 828 LVEL with a tetrafunctional methylenedianiline hardener. In general, they observed that the initial reaction rate is increased with increasing nanofiber MC (wt.%) as shown in Figure 2.12(a). The final glass transition temperature (Tg) of the samples, however, showed no influence of MC as shown in Figure 2.12(b).

(a) 0.06

0.05

0.04

0.03

0.02 2

Initial reaction Initial rate (W/g) R = 0.9891 0.01 0 2 4 6 8 Nanofibre MC (wt.%) [AHA8] (b) 180

170

175

165

(ºC)

T 160

155

150 0 1 2 3 4 5 6 7 Nanofibre MC (wt.%)

Figure 2.12: Influence of nanofiber MC in epoxy during cure, (a) initial reaction rate and (b) final glass transition temperature [68].

2.2.2.2 Effect of MC During Cure on the Mechanical Properties of Epoxy

The mechanical performance of an epoxy system after being cured with water was reported by some authors [19][90][93]. The earliest attempt to characterize the mechanical properties of epoxy adhesives contaminated with water during cure was conducted by Chen et al. [19] in 2001. They found that ultimate tensile strength (Figure 2.13(a)), flexural strength, compression strength, adhesive shear strength reduced with increasing water content. Using only exposure to ambient atmosphere (RH) to facilitate the moisture absorption in uncured epoxy resin, Ian and Polak [90] also observed reduction of tensile strength, compression strength and shear strength compared to uncontaminated epoxy samples. In contrast, Wu et al. [93] found that a small amount of water (2% wt.) in an epoxy adhesive (SikaDur 30) increased the three point bending flexural load of FRP-reinforced concrete beam bars bonded with SikaDur 30 adhesive as shown in Figure 2.13(b). However, the flexural load was subsequently decreased at higher MC (4% wt. and 6% wt.) and the trend was consistent with dynamical flexural modulus result from DMA. In summary, previous studies show that mechanical properties of epoxy resins are strongly affected by the presence of water during cure and this can be a drawback for natural fibres where the occurrence of moisture is inevitable.

(a) (b) 25

)

2 - 20 50 15

Maximum FlexuralMaximum Load (KN) 0 No Neat Ultimate Tensile Strength (Nmm + 2% + 4% + 6% Reinforcement Sikadur water water water 0 0.2 30 Water Content (%) Adhesive used for bonding composite sheet to concrete [AHA9]

Figure 2.13: Mechanical properties of water contaminated in uncured epoxy; (a) ultimate tensile strength [19] and (c) flexural load strength of FRP- reinforced concrete beam bars bonded with SikaDur 30 epoxy adhesives [93].

2.3 Composites Manufacturing

The manufacturing of natural fibre polymer composites uses many methods such as hand lay-up, compression moulding, prepreg moulding, resin transfer moulding (RTM) and resin infusion (RI). Among these manufacturing methods, RI or vacuum infusion is a cost-effective way to produce composite laminates for aerospace, marine and wind power applications using liquid resins [45]. Table 2.5 compares different manufacturing technique for natural fibres polymer composites. The advantage of RI is that it offers low porosity in cured composites which is very important as porosity reduces mechanical properties such as interlaminar shear strength (ILSS) and compression test [96].

Table 2.5: Influence of manufacturing technique with respect of fibre volume fraction and porosity of composites [23].

Manufacturing technique Fibre volume fraction (%) Porosity volume fraction (%)

Compression moulding Up to 85 % (typically 25–50 %) Up to 25 % (typically 2–8 %)

Prepreg (autoclave) Up to 60 % (typically 35–50 %) Up to 10 % (typically 0–4 %)

RI/RTM Up to 60 % (typically 25–50 %) Up to 10 % (typically 1–4 %)

A schematic of RI is shown in Figure 2.14. RI is an open moulding system where only one side of a solid mould is used. A vacuum bag is applied on top of laminates to form the mould. RI utilizes vacuum-only (negative) pressure to compact the dried laminates initially, allowing the resin to flow inside and wet the laminates [97]. Once infusion is completed, the wet laminate is then cured (typically in an oven).

To vacuum pump Cup containing mixed resin Outlet tube Inlet tube Mesh flow media Vacuum bag

Sealant tape

Peel Ply Laminate mould Figure 2.14: Schematic of resin infusion (RI) [97].

2.4 Ideal Humidity Conditions for Composites Processing and Fabrication

The atmospheric moisture absorption of natural fibre is due to humidity exposure. To illustrate this, a plot of relative humidity on a single day (1st July 2018) in Manchester, UK is presented in Figure 2.15. It can be seen that the RH fluctuates significantly within hours and consequently, MC in natural fibres could vary accordingly. Davis [98] suggested that between 65% RH at 18°C and 45% RH at 23°C is ideal to restrict moisture uptake in epoxy adhesive but that lower RH is desirable. Unfortunately, composite fabrication facilities for both storage and

composite fabrication are often not well controlled. This would mean a huge impact on the final mechanical properties of natural fibre polymer composites.

100

80 70

40 Relative Humidity (%) Humidity Relative 30

20 23:58 02:00 04:00 06:00 08:00 10:00 12:00 14:00 16:00 18:00 20:00 22:00 23:58 Time

Figure 2.15: Daily plot of humidity in Manchester, United Kingdom on 1st July 2018 [99].

2.4.1 Moisture Presence during Composites Manufacturing

For natural fibres composites, often the main source of moisture during manufacturing is from the moist fibres themselves. Moisture can possibly diffuse into the epoxy during resin infusion process (fibre wetting) and ultimately, the presence of moisture will interfere with the curing kinetics of the epoxy [68][100]. A number of researchers have investigated pre-cautionary efforts to reduce the presence of moisture in natural fibres during composites manufacturing [53][95][96]. Typically, the process began with drying the natural fibres in a heating oven at between 80 ºC and 105ºC for several hours. However, natural fibres are extremely hygroscopic materials and the moisture uptake could immediately initiate once it has been removed from the oven and exposed to the atmosphere humidity. This method requires further processing cost.

2.5 Fibre/Matrix Interface Bonding

Bonding at the interface in a composite is a consequence of the adhesion between the fibres and matrix. Stronger interfaces allow more effective transfer between fibre and matrix. The adhesion mechanism is not straightforward and adhesion mechanisms common in polymer-fibre composites are described below [103]–[105]:-

Interdiffusion. Occurs when molecules of the fibre and matrix may interact or diffuse at the fibre-matrix interface, which result from hydrogen bonding or van der Waals forces. The strength of this bonding is defined by the degree of molecular entanglement and the numbers of molecules involved.

Adsorption and wetting. When two surfaces (one of which is in liquid form) are brought together and come into intimate (atomic scale) contact, ‘wetting’ is said to occur for which adhesion is primary caused by the van der Waals forces.

Chemical bonding. Occurs when a functional group at the fibre surface reacts with a compatible functional group in the matrix. The strength of bonding depends on the number of bond per unit area and type of chemical groups.

Mechanical Adhesion. This bonding implies that the two surfaces provide mechanical interlocking such as when a liquid resin wets entirely the rough fibre surfaces.

It is worth noting that one of more adhesion mechanisms can possibly occur at the same interface at the same time [104].

2.5.1 Interfacial Bond Strength Measurement

The measurement of interfacial bond strength can be done qualitatively and quantitatively for purposely designing composites. The methods can either directly or indirectly measure stress at the interface. Examples of indirect measurements are interlaminar shear strength (ILSS) and transverse properties. Recently, the interface bonding of natural fibre polymer composites has been characterized using Dynamic Mechanical Analysis (DMA)[106][107].

2.5.1.1 Microbond Test

Several micromechanical methods have been developed in order to determine the interfacial shear strength (IFSS) including the fibre pull-out test [108], single fibre fragmentation test [109], microbond test [110] and microindentation [111]. Each technique has its own unique advantages and limitations. Among these micromechanical methods, the microbond test has been widely used to characterize the interfacial bonding of composites. Microbond testing was pioneered by Miller [112] as an upgraded version of the fibre pull-out test. The development of the microbond test was aimed at overcoming the common problem from the pull-out test, which requires a small embedded length or else the fibre will break before it pulls free. Hence, a small amount of resin in form of a microdroplet is applied on the fibre. The test specimens have a simple configuration and various types of fibres and matrices can be successfully employed [113]. However, to this date, there is no standard procedure available on how this test should be carried out and most laboratories develop their own specimen preparation and testing procedure.

Numerous studies have been conducted using microbond methods to determine variation in IFSS properties due to procedures such as fibre chemical treatments [114], hygrothermal and/or other aging conditions [115][116]. The microbond method also has been used successfully for natural fibre composites, in which either the fibre is in the form of a technical fibre (bundled elementary fibres) or a single elementary fibre [109][111]–[114]. The choices between technical fibres and single elementary fibres for microbond depends on the personal preferences or type of fibre, but the IFSS of single elementary fibres is usually higher than those of technical fibres. According to Fuentes et al. [121], the strength of composites is probably controlled by the elementary fibres rather than the technical fibres.

2.5.2 Relationship between MC and Interfacial Bond Strength

Only a few studies have reported the interfacial shear strength (IFSS) behaviour of natural fibre composites produced at different humidity/MC. Using the microbond technique, Zhang et al. [119] showed there was a different behaviour of the IFSS-displacement curves on flax/unsaturated polyester composites when the

samples were manufactured at either low or very high RH as shown in Figure 2.16(a). High IFSS values were characterized by a sudden failure of the interface where the IFSS drop abruptly. However, there was no sudden drop of IFSS when the samples were subjected to 90% RH. Using the fibre pull-out technique, a similar observation was reported by Chen et al [18] for bamboo/vinyl-ester composites (Figure 2.16(b)). They observed a broad peak in the force-displacement curves when the samples were produced at >80% RH. According to Desarmort and Favre [122], this behaviour indicates very poorly bonded interfaces.

(a)

Strong interface

IFSS (MPa) IFSS Weak interface

Displacement (mm) (b)

(Low RH) (>80% RH)

Strong interface

Weak interface

Load (N) Load Load (N) Load

Displacement (mm) Displacement (mm)

Figure 2.16: IFSS/load-displacement curves showing strong and weak fibre/matrix interfaces; (a) flax/unsaturated polyester composites [119] and (b) bamboo/vinyl ester composites [18].

In regards to IFSS values with humidity variation during fabrication, Chen et al. [123] observed the highest IFSS values for bamboo/vinyl ester composites with dried reinforcement (~11 MPa) compared to the humidified conditions (8.5 MPa and lower) as shown in Figure 2.17(a). They also observed that IFSS was reduced with increasing MC until at highest MC of ~28 wt.%, the IFSS almost reached zero.

Microstructure studies of the cross-section of bamboo/vinyl ester composites, using optical microscopy, indicated that at 80% RH multiple wide voids were present around the fibre interface.

For flax/unsaturated polyester composites, Fuentes et al. [115] observed that the IFSS was highest when manufactured using dried fibres (13.6 MPa) and IFSS reduced to 9.6 MPa for 50% RH fabrication as shown in Table 2.6. Zhang et al. [113] studied the influence of humidity between 40%, 50%, 70%, 80% and 90% RH for flax/unsaturated polyester composites fabrication as shown in Figure 2.17(b). They observed that the IFSS were consistent between 40% and 70% RH (~7 MPa) but the IFSS reduced significantly (by >40%) at 80% RH and by ≈85% for samples were produced at 90% RH. SEM images of test specimens indicated strong interfaces for samples ≤70%RH, characterized by resin adhesion and fractured resin segments on the surfaces, while weak interfaces exhibiting smooth fibre surfaces were found for the higher RH specimens

(a) (b)

IFSS (MPa) IFSS IFSS (MPa) IFSS

Bamboo MC (wt.%) RH (%)

Figure 2.17: Influence of MC/RH on IFSS values; (a) bamboo/vinyl ester composites [123] and (b) flax/unsaturated polyester composites [119].

2.5.3 Influence of Fibre Swelling and Shrinkage at the Fibre/Matrix Interface

Swelling and shrinkage in natural fibres due to changes in MC can have significant effects on IFSS and this depends on the humidity during composites manufacturing, post-conditioning and test conditions [18][121]. A summary of two 54

studies is shown in Table 2.6. Chen et al. [18] produced bamboo/vinyl ester composites at dried fibre conditions and the samples were post-conditioned at 80% RH. The IFSS was slightly reduced from 10.82 MPa to 9.58 MPa despite bamboo experienced swelling, which was proposed to increase mechanical interlocking at the fibre/matrix interface. In contrast, Fuantes et al. [121] observed a slight increase from 10.3 MPa to 11.7 MPa for flax/unsatured polyester composites as RH during fabrication was increased from 50% to 70%, as shown in Table 2.6. Fibre swelling was proposed to exert extra pressure against the solidified matrix during the pull-out test.

In case of fibre shrinkage after post-conditioning, both Chen et al. [18] and Fuantes et al. [121] agreed that the shrinkage has a negative effect on IFSS as shown in Table 2.6. The former reported that samples manufactured at 80% RH and post conditioned in dried conditions had lower IFSS (0.25 MPa) as compared to the samples without post-conditioning (0.80 MPa). Fuantes et al. [121] showed that IFSS was reduced slightly from 10.3 MPa to 9.6 MPa when the samples at 50% RH in the fabrication were post-conditioned at 0% RH.

Table 2.6: Influence of fibre swelling and shrinkage on IFSS.

Fibre Conditioning, Post-Conditioning, IFSS Samples Ref RH (%) RH (%) (MPa) 0 (dried) None 10.82 Bamboo/vinyl 0 (dried) 80% 9.58 [18] ester composites 80 None 0.80 80 0 (dried) 0.25 0 0 13.6 Flax/unsaturated 50 0 9.6 polyester [121] 50 50 10.3 composites 70 70 11.7

2.5.4 Effect of Chemical Treatment to Increase Interfacial Bonding

To improve the interfacial adhesion of untreated bamboo/vinyl ester composites in humid environments during composites fabrication, Chen et al. [123] performed various fibre surface treatments (alkaline, acetylation, oxidation –

potassium permanganate (KMnO4) and silane). Except for the alkaline treatment, they found that the other chemical treatments gave lower moisture absorption for a given humidity (40%, 60% and 80% RH) in the bamboo strips as shown in Figure 2.18(a). These chemical treatments also improved the interfacial shear strength as shown in Figure 2.18(b), with silane giving the highest improvement (36.5%) as compared to the untreated sample. The enhanced of IFSS after chemical treatment was attributed to increase mechanical interlocking and chemical bonding at the interface.

(a) (b)

Figure 2.18: Influence of chemical treatments on (a) moisture absorption and (b) interfacial shear strength of bamboo/vinyl ester composites during composites fabrication [123].

2.6 Surface treatment to Enhance Natural Fibre Composites Performance

The hygroscopic/polar characteristics of natural fibres induce lower compatibility with non-polar/hydrophobic matrices [124]. Hence, fibre surface treatments are a common method to reduce moisture regain in natural fibres while at the same time improving the fibre/matrix interface, thermal resistance and mechanical properties of natural fibre reinforced polymer composites [123][125]. The surface treatment to modify natural fibre properties can be done via chemical treatments or physical treatments.

2.6.1 Chemical Treatments

2.6.1.1 Alkaline Treatment

Alkaline treatment (or mercerization) is commonly reported in the literature and is perhaps the most effective method to improve interfacial bonding in natural fibre polymer composites at minimal cost. Alkaline treatment can be performed by immersing natural fibres in an alkaline aqueous solution such as sodium hydroxide (NaOH) and sodium bicarbonate for a given time duration and temperature [126][127]. Alkaline treatment removes portions of the fibres constituents such as hemicellulose, lignin, and wax from the external surface of fibres. The chemical reaction between fibre and alkaline solution is shown below [124]:-

Fibre – OH + NaOH → Fibre – O- Na+ + H2O + impurities 2.3

Alkaline treatments disrupt the hydrogen bonding in the network structure and, hence increases surface roughness which provides better mechanical interlocking at the fibre/matrix interface. Several studies have shown that mechanical properties (strength and stiffness) in natural fibre reinforced composites increased after alkaline fibre treatment [128]–[130].

2.6.1.2 Silane Treatment

Coupling agents such as silanes, isocyanates, anhydrides and maleated polymers are used to enhance the mechanical properties of natural fibre composites. Coupling agents bear two reactive groups; one to react with the hydroxyl (OH) rich surface of the fibre and another to react (copolymerise) with the matrix, generating covalent bonding between the fibre and matrix [131]. Amongst coupling agents, silanes are an attractive option because they are widely commercially available [131]. Silanization begins with hydrolysis in the presence of water, forming reactive silanol groups, which subsequently are condensed or absorbed onto the fibre surface to form hydrogen bonds with hydroxyl groups. Upon heating, this hydrogen bonding can be converted into covalent (Si-O-C-) bonds upon liberation of water [132][133]. In their review, Xie and co-workers [133] suggest that proper selection of the organofunctionality of a silane to match specific polymer matrices is important in 57

attaining the desired coupling. The common organofunctionalities for epoxy matrices are aminosilane and glycidoxysilane. Aminosilane coupling agents such as aminopropyltriethoxysilane (APS) [134]–[136] and aminopropyltrimethoxysilane (APTMS) [126][131][132] are commonly reported to be used with natural fibre composites and results indicate that the fibre/matrix interface and mechanical properties of natural fibres are improved after the silane treatments. Then et al. [132] showed that the tensile strength values of oil palm mesocarp fibre/ poly(butylene succinate) biocomposites were improved from 13.86 MPa to 27.12 MPa when the fibre were treated with APTMS (2 wt.% solution) Significant improvement was also observed for tensile modulus, which increased from 94.8 MPa to 868 MPa, an ~815% increase.

2.6.2 Silane on Alkalized Fibre Treatment

Many previous investigations have applied either alkaline or silane treatments to modify natural fibres characteristic. However, combinations of these treatments are also possible due to fact that these treatments have different mechanism on natural fibres. The treatment begins initially with alkaline treatment followed by the silane treatment. One of the earliest studies of silane on alkalized fibre treatment was reported by Franco and Gonzalez [139] in 2005. They studied various chemical treatments in henequen/high density polyethylene (HDPE) composites. They showed IFSS (pull-out test), Iosipescu shear strength and the tensile strength of combined treatment samples (2% NaOH and 1% silane) was highest compared to untreated and others treated fibres, including the standalone treatments. For example, the IFSS of untreated henequen fibre is 5.4 MPa. NaOH and silane treatment increased IFSS to 9.2 and 11.9 MPa, respectively whereas the combined treatment showed the highest IFSS of 16 MPa. The improvement of mechanical properties of these composites was proposed to be due to the benefits of having both mechanical interlocking (alkaline) and chemical bonding (silane) at the fibre-matrix interface. Hence, studies of silane on alkalized fibre are increasing since these combined treatments have resulted in improved mechanical properties on composites of natural fibres and polymer matrices such as kenaf/polypropylene [136] and coconut/unsatured polyester [140].

2.6.2.1 Acetylation

Acetylation (acetic anhydride) treatment is a method of esterification for plasticizing natural fibres. The acetyl functional group (CH3COO–) reacts with reactive hydroxyl (OH) groups on cellulose and prevents any other OH group such as water interacting with them. Consequently, the hydrophilicity of natural fibres is reduced while aiding dimensionally stable dispersion of fibres into polymeric matrices [49]. Kabir and co-workers [124] explained that acetylation treatment improves mechanical interlocking of the fibres with the matrix as a result of rough surface topography. Following acetylation, it is reported that the flexural properties of flax/epoxy composites showed a 10% increase in flexural strength and 13% increase in flexural modulus [141].

2.6.3 Physical Treatments

A number of physical treatments can be applied to fibres prior to composites fabrication; such as corona, plasma and ultraviolet (UV) treatments with no modification of the chemical composition of the fibre [142]–[145]. However, these kinds of treatments require specific equipment.

Corona treatment - This treatment involves surface oxidation of the fibre and an etching effect which improves interfacial compatibility between the hydrophilic fibres and hydrophobic matrices [142][146]. Improvement of tensile properties was reported by Rogoubi et al. [147] for miscanthus fibres/PLA composites after corona treatment.

Plasma treatment - This treatment interacts with surface materials by radicals, energetic particles crosslinking, and deposition reactions [143]. The effectiveness of surface modification depends on the type of gases used and the improvement of interfacial adhesion was proposed to result from increased surface roughness of the fibres [143].

UV treatment - This treatment is intended to increase the polarity of the fibre surface and consequently, improving fibre wettability and composites strength [148]. Abdullah et al. [144] showed an improvement of 28% in tensile strength and

21% in tensile modulus was achieved after UV treatment over the untreated samples of jute/unsaturated polyester composites.

2.7 Effect of Humidity on Composite Mechanical Performance

The mechanical performance of natural fibre composites with different RH during fabrication was reported in several studies [119][121][149][150], all agreed that the performance of the composites was affected by humidity (or fibre MC) during composites fabrication. Fuentes et al. [121] studied the flexural properties of unidirectional flax/unsaturated polyester composites produced at 0% and 100% RH, and their results are shown in Table 2.7. The significant longitudinal properties loss of composites was observed at 100% RH fabrication. However, they believed the interface properties is not drastically affected by humidity evidently from the same transversal properties values (Young’s modulus and strength) between 0% and 100% RH samples which is slightly reduced. The insensitivity of the interface to presence of moisture in the fabrication can be explained by the presence of lignin at the flax fibre surface which is relatively hydrophobic in nature.

Table 2.7: Flexural longitudinal and transversal properties of flax/unsaturated polyester composites produced at 0% and 100% RH [121].

Longitudinal Transverse RH Young’s modulus Strength Young’s modulus Strength (%) (GPa) (MPa) (GPa) (MPa)

0 20.3 ± 0.3 230.7 ± 5.6 2.3 ± 0.2 16.4 ± 0.6

100 15.3 ± 1.1 188.6 ± 5.8 2.1 ± 0.4 17.5 ± 2.6

Islam and Mio [149] developed a novel biocomposite using Acrodur acrylic resin and woven flax fabric. They reported that the tensile strength of this biocomposite was not significantly affected by the RH during fabrication (40%, 60% and 80% RH) as shown in Figure 2.19. However, the tensile modulus reduced consistently and larger tensile strain was observed with increasing RH. The authors did not explain the state of fibre/matrix interface in the study. Increased tensile

strength at 40% RH over dried (0% RH) samples was observed. They suggested it has to do with curing acrylic resin in the presence of moisture where the bonding is activated by an esterification reaction between the acid groups of the resin and the hydroxyl groups of the flax fabric constituents i.e cellulose, hemicellulose and lignin.

Figure 2.19: Influence of humidity on flax/Acrodur biocomposites; (a) tensile strength and (b) tensile modulus [149].

Zhang et al. [119] reported flexural properties of unidirectional flax/unsaturated polyester composites produced using dried fibres and under humid conditions (40%, 60% , 70% 80%, 90%) during fabrication as shown in Figure 2.20, they dried fibre composites started at a low flexural strength 140 MPa but then improved at 40% RH (160 MPa). With increasing fibre MC, the flexural strength consistently reduced. In contrast, the flexural modulus was highest at dried conditions and the flexural modulus reduced consistently across the RH.

Figure 2.20: Flexural properties of flax/unsaturated polyester composites fabricated at different humidity condition [119].

Recently, Moudood[AHA10] et al. [150] reported both tensile and flexural properties of unidirectional flax/epoxy composites of dried (C-D) and humidified flax (35%, 50%, 70%, 90% RH) during fabrication as shown in Figure 2.21. Both tensile and flexural data showed different trends with variation of humidity. It can be seen in Figure 2.21(a) that the flexural strength and flexural modulus consistently decreased with increasing humidity which is consistent with tensile modulus in Figure 2.21(b). They showed that the reduction of tensile and flexural modulus was attributed to the poor fibre/matrix bonding For tensile properties in Figure 2.21(b), the dried samples had the lowest tensile strength (~273 MPa) which slightly increased up to 50% RH (~288 MPa). They believed that this is due to increase fibre strength with increasing fibre MC. After the threshold (50% RH), the tensile strength reduced slightly to ~285 MPa .

Flexural Flexural Modulus

Strength

Flexural Strength (MPa) Flexural Modulus (GPa)

Young’s Modulus Tensile Strength

(b)

Tensile Strength (MPa) Young’s Young’s Modulus (GPa)

Figure 2.21: Influence of humidity during fabrication of unidirectional flax/epoxy composites; (a) flexural and (b) tensile properties [150].

2.8 Fracture Toughness and Impact Properties of Composites

2.8.1 Interlaminar Fracture Toughness (ILFT)

Delamination or interlaminar fracture is one of the important aspects when designing a composite. Delamination occurs due to impact loading during service or possibly due to contamination or non-optimized curing during manufacturing manufacture prior to service [151]. A delamination is described by a crack that initiates and grows in between the different plies of a composite in the resin rich area. [152]. Interlaminar fracture behaviour can be explained by fracture mechanics and is measured as a strain energy release rate, GIc. There are three modes of fracture toughness that correspond upon loading as shown in Figure 2.22. Mode I

(GIC) is often characterized by a double cantilevered beam (DCB) test and Mode II

(GIIC) is with an end notch flexural (ENF) test.

Mode I: Opening Mode II: In-plane shear Mode III: Out-of-plane shear

Figure 2.22: The three modes of fracture toughness [153].

The fracture energy absorption depends on the interaction between the constituent components, including the fibre/matrix interface. Both fibres and matrix absorb the strain energy until one of the components started to fail. In a unidirectional fibre polymer composite, the elastic, viscoelelastic and plastic deformation is performed by the matrix whilst the role of fibre is to absorb the energy or limits the plastic zone by breakage or fibre bridging [154]. Figure 2.23 shows schematic interlaminar failure (delamination) modes of composites where the crack path is parallel to the fibres between two plies of a laminate. Various damage

modes can be seen related to the fibres such as fibre bridging, fibre surface fracture and fibre breakge. For the the matrix, the damage is presented by side crack and matrix microcracking (plastically deformed) around the crack tip.

Figure 2.23: Failure modes of composites due to delamination [153].

There are numerous studies attempting to evaluate and improve interlaminar fracture toughness (Mode I and Mode II) of natural fibre polymer composites through fibre surface treatment - silane [155][156], hybridization of reinforcements [157], different fabric architecture and orientation [32][151][152], toughening the matrix [101] and optimizing the manufacturing process [101]. Recently, the effect of moisture absorption on cured woven flax/vinyl ester composites and hybrid composites (with woven basalt reinforcement) was conducted by Almansour et al. [157]. The results indicated that moisture absorption caused a reduction in the interlaminar fracture toughness of flax/vinyl ester composites by 27%. In contrast, the MC had mixed effects on hybrid composites but still the hybridization composites could provide a better shield to the swelled flax fibre. They also proposed that delamination and crack growth could be affected significantly by the hygrothermal conditioning.

2.8.2 Low Velocity Impact[W11]

Impact properties of a material is the ability to absorb and dissipate energies under impact loading, which can be initiated by a crash, a falling object or by debris [159]. Common tests to characterize impact properties include Izod and Charpy pendulum tests and drop-weight impact tests. A drop-weight impact test has advantages because variety geometries can be tested and it is possible to use different impactor shapes and sizes. Impact damage from low velocity impact is one of the severe threats to composite structures because the damage is difficult to see in visual inspection (“barely visible impact damage”) and consequently, the residual strength of composites would be reduced [160]. Low velocity impact refers to an impact at velocities below 10 m/s [161]. The absorbed energy during impact test can be characterized by load-deflection curves as shown in Figure 2.24(a-c). The description of each curve is explained below [162][163]:-

Rebounding: When a sample is impacted at low level energy (less than penetration threshold), impactor rebounding occurs. This is a closed curve.

Penetration: This curve is considered an open curve as the impactor has penetrated into the specimen, although partial rebounding may happen.

Perforation: The sample is completely perforated with no rebounding of the impactor. This is an open curve.

(a) (b) (c)

Figure 2.24: Typical impact response of composites (a) rebounding (b) penetration (c) perforation [162].

A schematic of typical failure modes after an impact event is shown in Figure 2.25. In case of full perforation of the specimen, the first failure mode of the impacted specimens is due to matrix damage (matrix cracking, fibre matrix interface 66

damage) which occurs parallel to the fibre [164]. The matrix cracking is due to bending, compression and shear because of the property mismatch between the fibres and the matrix. The second failure mode is delamination which results from the bending stiffness mismatch between adjacent layers [161]. Caprino et al. [165] described that the starting points of delamination can be noticed when there is a change in slope in the load-displacement curve. The third failure mode is fibre breakage and in-compression fibre buckling [161].

Figure 2.25: Schematic of failure mode of laminated composites [166].

According to Bensadoun et al. [164], in addition to the interlaminar fracture toughness, many factors contribute to the impact properties of composites such as fibre architecture, fibre orientations, fibre-matrix interface quality, stacking sequence, fibre type, matrix type, matrix ductility . The study of impact properties of dry and salt water conditioned basalt and flax reinforced vinyl ester composites was reported by Živković and co-workers [167]. Both type of composites showed different impact properties. The data showed that the dry basalt composites have higher absorbed energy than the salt water conditioned samples, arising from basalt and cured vinyl ester matrix becoming more brittle after aging. In case of flax composites, the salt water conditioned flax composites had higher absorbed energy than the dry flax composites. They believed that flax fibre become more ductile after moisture absorption and hence, increasing the toughness of the flax composites.

2.8.3 Compression After Impact (CAI[AHA12])

The residual strength of composites after low-velocity impact can be measured using an uniaxial in-plane compression test, commonly known as compression after impact (CAI) [168]. CAI is the most common method for measuring the damage tolerance of composites in order to establish confidence in the materials [169]. To avoid global buckling during compression loading, the CAI test must be conducted with a designated test rig. [168]. Cantwell and Morton [170] suggested there is relationship between Mode II interlaminar fracture toughness and CAI properties, but not with Mode I interlaminar fracture toughness and CAI. Cartié and Irving [171] explained that CAI failure is a matrix-dominated failure process and resin toughness is the major influence on the CAI strength rather than fibre strength and stiffness. A study by Ismail et al. [172] on kenaf/glass hybrid composites showed that larger damage area and delamination on the impacted specimens contributes to lower CAI strength.

2.9 Key findings of The Literature[W13]

Chapter 2 presents a comprehensive review of relevant studies from past to present research work on the influence of MC on the natural fibres reinforced polymer composites during manufacturing. The review also includes a brief explanation on the origin of hygroscopic behaviour and their negative implication to the physical and mechanical properties of composites including their individual elements such as natural fibre, epoxy matrix and fibre/matrix interfacial bonding. The technique for measurement such as fibre density, fibre MC and interfacial bond strength is reviewed thoroughly for which help to develop experimental procedure and hence, obtaining accurate data. The review highlights the importance of controlling environmental humidity in the composite fabrication facilities where natural fibres are stored and processes into composites. Here, the role of surface treatment such as alkaline, silane, silane on the alkalized fibre and acetylation to enhance interface bonding of composites and reduce moisture regain in fibre is presented and discussed. Importance key information pertaining on the interlaminar

fracture toughness and low velocity impact properties of natural fibre composites has been reviewed.

Flax/epoxy composites could offer very good mechanical properties that suited various applications such as automotive sector. However, there are limited studies on how fibre MC could affect to the fibre/matrix interface bonding and mechanical performance of flax/epoxy composites. The presence of moisture in flax fibre is a natural occurrence and understanding moisture interaction at the interface during flax/epoxy composites manufacturing is essential to the composites design and application.

Chapter 3 Experimental

3.1 Materials

The fabric reinforcement used was a commercial 2x2 twill flax fabric (shown in Figure 3.1(a), supplied by Composites Evolution under the Biotex Flax trade name. The manufacturer’s datasheet states the fabric has similar warp and weft yarn count, 7 yarns per cm. The Biotex Flax fabrics are based on twistless yarns and have 250 tex for both warp and weft. The ply thickness of the fabric was between 0.60 and 0.80 mm and the fibre areal density 400 g/m2. The fabric was cut into 290 mm (warp direction) x 220 mm (weft direction) sheets prior to fibre surface treatments and composites fabrication as shown in Figure 3.1(b).

Figure 3.1: Photograph of (a) 2x2 twill flax fabric and (b) 250 tex flax yarn and (c) flax fabric ready for surface treatment and composites fabrication.

The epoxy resin matrix system used was a mixture (100:35) of Araldite LY 564, a low molar mass diglycidyl ether of bisphenol A (DGEBA)/ butane diol- diglycidyl ether resin blend of equivalent weight 165 ± 5 g/eq., and Aradur 2954 (2, 2’-dimethyl-4.4’-methylenebis (cyclohexylamine)) curing agent (both Huntsman Advanced Materials). The datasheet properties for this system are shown in Table 3.1. Prior to resin infusion, preparation began by stirring the two liquids using a 70

mechanical stirrer for 5 minutes in an open container followed by degassing in a vacuum chamber for about 1 hour to remove entrained air.

Table 3.1: Resin properties from manufacturer’s datasheet.

Mixed Viscosity at 25 °C Gel Time at 60 °C Pot Life at 25 °C Density at 25 °C (min) (min) (gcm3) (mPas) 500-700 90-120 400-600 1.2

3.2 Fabric Surface Treatments

3.2.1 Alkaline treatment

NaOH pellets were supplied by Sigma-Aldrich. The alkaline treatment began with the immersion of flax fibre in solutions of different concentrations; 1.5, 3.0, and 4.5 w/v% at room temperature. 20 minutes immersion time was chosen to minimize structural damages to the fibre [173]. The fibres were then removed and washed thoroughly with distilled water. The excess alkaline was removed by immersion in an acidified solution (10 drops of 0.1 M HCl for every 1 L of distilled water). Again, the fibre was rinsed thoroughly with distilled water and checked using indicator paper until pH 7. Finally, fibres were dried in a convection oven at 50 ºC for 16 hours.

3.2.2 Silane Treatment

For silane treatment, 3-Aminopropyltrimethoxysilane (APTMS – Figure 3.2) supplied by Sigma-Aldrich was used as a silanization agent. Silane preparation began with mixing of acetone and distilled water at a volume ratio of 50/50. Silane was added to the acetone/water mixture at one of three volume % (1.0. 1.5 and 2.0% v/v) and stirred thoroughly for about 10 minutes using a mechanical stirrer. The fibres were then soaked in the silane solution for 2 hours at room temperature. Finally, fibres were dried in a convection oven at 50 ºC for 16 hours.

OCH3

NH2 H3CO Si

OCH 3

Figure 3.2: Chemical structure of APTMS silane agent.

3.2.3 Alkaline & Silane Treatment

The procedure of silane treatment on alkalized fibre was the same as those described in Section 3.2.1 and Section 3.2.2.

3.3 Conditioning Flax Fibres

In order to get the desired moisture content (MC), flax fibres were conditioned at a selected relative humidity (RH) in a climate chamber (Binder KMF 115) prior to fibre characterization and composites manufacturing. Flax fibres were left in the climate chamber for 24 hours and the conditioning temperature was fixed at 23 °C to ensure that equilibrium MC had been achieved. To avoid the hysteresis effect of moisture content between absorptions and desorption [69][71][174], the conditioning of flax fibre began on absorption line, starting from 10% RH and progressing to next RH level.

In case of dried fibres corresponding to the lowest level of MC 0% RH (designated as 0% MC, but unlikely to be exactly zero – this value is discussed later in Section 4.1.3), a different method was employed because the climate chamber has a limited relative humidity and temperature service range, 10% to 98% RH for the former and -10 to 100°C for the latter. Flax fibres were dried to remove water content in the cell walls [59]. A temperature of 105 °C was selected and flax fibres were dried for 24 hours (giving constant weight) in a convection oven (Gallenkamp Hot Box).

3.4 Composites Manufacturing

The untreated flax/epoxy composites were fabricated via a resin infusion (RI) technique. The flax fabrics were vulnerable to moisture uptake outside the humidity chambers and care was needed to minimize their exposure. Therefore, RI was done next to the humidity chamber. For RI, a layer of mould cleaner and release agent (700NC, Frekote) was first applied to the surface of a flat tool plate and allowed to dry. The lay-up configuration was prepared as shown in Figure 3.3(a) with peel plies and distribution mesh on top of it. The inlet spiral tube and outlet spiral tube pipes were connected with the inlet and outlet tubes. The preform stack was immediately inserted in between two peel plies once taken out from the humidity chamber; the nylon bag was immediately laid on top and sealed along its perimeter using tacky tape. The air was drawn out from the set-up using the vacuum pump (<0.01 mbar) and the outlet pipe was clamped. The process was completed in less than 1 minute. Then, the set-up was checked for any sign of leakage for about 5 minutes by monitoring the pressure changes indicated by the pressure gauge as shown Figure 3.3(b).

The RI process was done without the aid of additional vacuum pumping to ensure any fibre MC remains in the lay-up as shown in Figure 3.3(c). The fabric was infused until the resin filled the outlet spiral tube and the vacuum was then applied again to ensure no entrapped air was left in the laminates. Finally, the inlet tube was clamped to stop the infusion process. RI was completed within 40-50 minutes. The infused fabric lay-up was then cured in a circulating-air oven. The curing cycle involved two stages, initial curing and post curing as specified by the manufacturer. The initial curing stage was a gradual heating from room temperature to 80 °C within 1 hour, which was maintained for next 2 hours. The post cured was began with heating from 80 °C to 140 °C within 1 hour and which was maintained for another 8 hours. The cured laminate was allowed to cool down to room temperature in the oven after the post curing. Figure 3.3(d) shows the cured composites.

Figure 3.3: Resin infusion of flax/epoxy composites showing (a) lay-up configuration, (b) vacuum check for sign of leakage after sealing (c) resin infusion process without vacuum pump aid and (d) cured composites prior to cutting.

3.5 Sample Cutting and Post-Conditioning

The cured laminate was cut into the required test specimen geometries and dimensions using a diamond cutter (Benetec) machine without the use of cooling water. Prior to testing, the flax fibre and flax/epoxy composites were conditioned for at least 24 hours in a temperature and humidity controlled laboratory, 20 ± 2 °C and 50 ± 5 % RH respectively. All the samples were sealed in plastic bags for storage and transport purposes before cutting.

3.6 Characterization Methods

3.6.1 Dynamic Vapour Sorption (DVS) Analysis

The moisture absorption and desorption behaviour of the untreated and treated flax fibres was studied using a Dynamic Vapour Sorption Analyser manufactured by Surface Measurement Systems Ltd (Figure 3.4). Approximately 16 µg of flax fibre was placed in a conical steel sample holder before being hung in the DVS chamber. The DVS employed a dry nitrogen purge, steam and saturated water to produce the required RH. The method began by drying the samples at 0% RH for 3 hours. The moisture absorption isotherm was measured in steps of 10% RH, in the range of 10% to 90%. The moisture desorption isotherm setting was similar but in reverse order. The specimen was considered to have achieved equilibrium MC if the weight changes are <0.01 % over 15 minutes, before progressing to the next RH level. All the parameters (temperature, gas flow, time, RH) are controlled by the DVS control software.

Figure 3.4: The Dynamic Vapour Sorption Analyser (DVS) used in this study.

3.6.2 Simple Weight Gain Method

A simple weight gain method is the easiest way to determine the MC of the flax fibre. Following the ASTM D 3766 standard [175], a standardized flax fabric specimen was created by utilizing fibre areal density. Fabric was cut into squares

following a plastic template with dimensions of 81 mm x 81 mm as shown in Figure 3.5(a). The use of the plastic template reduced errors as without it the flax fabric easily distorted during the cutting process. 5 specimens were prepared for averaging and data accuracy. The fibre areal density, 휌푠푎푐 was calculated using:-

푀 6 푓푎푏푟푖푐 3.1 휌푠푎푐 = 10 (푙푓푎푏푟푖푐 × 푙푓푎푏푟푖푐)

Where 푀푓푎푏푟푖푐 = weight of fabric, g

푙푓푎푏푟푖푐 = length of fabric, mm

Initially, the fabrics were dried in a convection oven (105 °C, 24 hr) since the climate chamber has a limited RH range (10% to 98%). The weight of dried fabric

(푀푑푟푖푒푑) and the weight of conditioned fabric (푀푐표푛푑푖푡푖표푛푒푑) following the method in Section 3.3, were both measured using a sensitive electronic balance (Sartorius Waage Analytical CP153) with 0.001 g accuracy as shown in Figure 3.5(b). MC (wt.%) was calculated using:-

푀 − 푀 3.2 MC (wt. %) = 푐표푛푑푖푡푖표푛푒푑 푑푟푖푒푑 푥 100 푀푑푟푖푒푑

(a) (b)

Plastic template

Produced standardized flax fabric

Figure 3.5: Photograph of (a) produced MC fabric samples and (b) weighting the fabric on the sensitive balance (Sartorius Waage Analytical CP153).

3.6.3 Crimp Measurements

The measurement of crimp before and after surface treatment was done according to the ASTM D 3883 standard [176]. Fabric was cut into squares using a plastic template with dimensions of 81 mm x 81 mm as described in Section 3.6.2.

The straightened fibre yarn length, 푙푠 values were obtained by manually applying sufficient force to remove the fibre waviness against a measurement ruler. The % crimp was calculated using:-

푙푠 − 푙푓푎푏푟푖푐 3.3 퐶푟𝑖푚푝 (%) = 푙푓푎푏푟푖푐

3.6.4 FTIR-ATR Spectroscopy

IR spectroscopy was carried out using a Nicolet FTIR spectrometer operating with a Diamond Attenuated Total Reflectance (ATR) fixture at room temperature. 64 scans per minute were performed with a band range from 500-4000 cm-1. Before commencing the measurement, a measurement of background noise was collected and then the bundle of fibres was placed and secured on the ATR cell using the screw module. At least 7 samples of each system were tested. Spectra were analysed using Universal Analysis 2000 software (TA Instruments).

3.6.5 Single Fibre Tensile Testing (SFTT)

SFTT was conducted on individual elementary fibres separated from the fibre yarn using tweezers and carefully attached onto a cardboard window-frame with a 10 mm gauge length (Figure 3.6). The fibre was secured at both ends using cyanoacrylate adhesive. The measurement of diameter of fibre was performed using a Keyence VHX-5000 optical microscope fitted with a 700x magnification objective lens. The variability of cross sectional area along the fibres length is expected for natural fibres [29], hence 5 measurements were taken along the fibre length. These values were used to calculate the average cross-sectional area without taking into account the lumen. Fibres with defects such as twisted fibres or irregular cross- sectional area were systemically removed as shown in Figure 3.7.

The determination of mechanical properties of single flax fibres was done with accordance with ASTM D 3822 [9] using a Zwick (Zwick/Roell Z2,5 TN Zwick-Line l) tensile testing machine equipped with a 20 N load cell in a controlled environment laboratory (55 ± 5 % RH, temperature 20 ± 3 ºC). Both ends of the frame were clamped into the mechanical grip, and then both sides of the window- frame were carefully cut in the middle as shown in Figure 3.6. Extension proceeded at a crosshead speed of 0.5 mm/min up to failure. Specimens that fractured very near the super glue were rejected. At least 30 specimens of each system were evaluated for tensile properties.

Figure 3.6: Single fibre tensile test configuration.

Figure 3.7: Excluded samples from SFTT; (a) twisted fibre (b) irregular cross- sectional area.

3.6.6 Density Measurements

3.6.6.1 Density of Flax Fibre

Measuring the density of flax fibres is considered complicated, owing to the internal structural of natural fibres. As discussed in the literature review, the Archimedes method was chosen because this is only applicable method to measure fibre density across the MC. However, the accuracy of this method can be dependent on the immersion liquid used e.g. methanol, water or acetone [74]. Following the recommendation of Truong and co-workers [74], Rapeseed oil (Canola oil) was used as the immersion liquid. Initially, the density of Rapeseed oil was measured using a density meter. The density of flax fibre was measured using a Metler Toledo XP504 analytical balance fitted with a density measurement kit (Figure 3.8). The analytical balance has a resolution of 0.0001 g. Flax samples were cut using scissors 20 x 20 mm squares. The flax fibres were quickly weighed in air and then immersed in the Rapeseed oil. The immersion time was between 4-6 minutes (giving constant weight) before the weight of the immersed fibres was recorded. At least 7 specimens were tested for each sample. The calculation of the density of flax fibre (휌푓) was given by:-

푚푎푓 3.4 휌푓 = 휌푙 푚푎푓 − 푚푙푓

-3 Where: 휌푙 = density of Rapeseed oil at 20 ºC (0.916763 gcm ); 푚푎 = weight of fibre in air; 푚푙 = weight of fibre immersed in the Rapeseed oil.

Figure 3.8: Density measurement configuration set-up on a Metler Toledo analytical balance.

3.6.6.2 Density of Composites

The densities of the composites were determined by Archimedes’ method using the same Metler Toledo XP504 Analytical balance (Figure 3.8) except that the immersion fluid was change to distilled water. Distilled water was used because there was found to be no variation of weight within 15 seconds of immersion; hence the use of Rapeseed oil was not necessary. Composites were cut into specimens 15 mm x 15 mm in dimensions. Four specimens were measured for each system. The density of composites (휌푐) was calculated using:-

푚푎푐 3.5 휌푐 = 휌푤 푚푎푐 − 푚푤푐

-3 where 휌푤 = density of distilled water at 20 ºC (0.9982 gcm ), 푚푎푐 = weight of composite sample immersed in air and 푚푤푐 = weight of composite sample immersed in distilled water.

3.6.7 Determination of the Fibre Volume Fraction of Composites

The volumes fraction of fibre, 푉푓 play important roles in determining the final mechanical properties of a fibre composite. The 푉푓 of the composites were calculated from the following equations [177]:-

휌푠푎푐 푁 3.6 푉푓 = 휌푓 h

where 푉푓 = fibre volume fraction, 휌푠푎푐 = fibre areal density, 푁 = number of fabric layers, 휌푓푙푎푥 = density of the flax fibre, h = thickness of the sample.

3.6.8 Microbond Shear Test Measurements of Fibre-Resin Interfacial Shear Strength (IFSS)

In this experiment, single elementary flax fibres were separated from the fabric using tweezers and then attached carefully to cardboard window-frames with a fixed gauge length of 10 mm. Both ends of a fibre were bonded using cyanoacrylate adhesive while being straightened with negligible force. The specimens were then conditioned in a climate chamber (Binder KMF 115) as described in Section 3.3. A droplet of the epoxy matrix system (after degassing) was placed on a flax fibre using a single carbon fibre (5 µm in diameter). After many unsuccessful trials, it appeared that the microdroplet length or embedded length, 푙푒, should be less than 180 µm in order to achieve a successful debonding or the fibre will break, invalidating the IFSS test.

Manufacturing microdroplets requires the conditioned fibres to be taken out from humidity oven for a short period of time (less than 5 seconds) in order to maintain MC. Hence, the placement of the microdroplet had to been done quickly with correct droplet size. Therefore, a reference for droplet size was made to aid in microdroplet manufacturing. The reference was several cured epoxy microdroplets (100-180 µm embedded length) on a single flax fibre. Once a microdroplet was positioned on a fibre, the microdroplet size produced was estimated against the reference as shown in Figure 3.9. Microbond specimens with microdroplets of the

appropriate size were then placed immediately in the climate chamber (Binder KMF 115) for curing at a desired % RH. The temperature was set at 80ºC and specimens were left to cure for 14 hours.

Carbon fibre with 100 µm 푙푒 microdroplet

180 µm 푙푒

Figure 3.9: Estimating the produced microdroplet size using a size reference.

A Keyence VHX-5000 optical microscope fitted with a 700x magnification objective lens was used to measure the dimensions of fibres and microdroplets and to examine the shape of the microdroplets. The diameter of each flax fibre was measured at least 6 times at both ends of a microdroplet as well as the embedded length, 푙푒 as shown in Figure 3.10(a). Microdroplets with defects, i.e. incomplete development of the microdroplet as shown in Figure 3.10(b), were systematically excluded.

The microbond specimens were tested on a Zwick tensile testing machine (Zwick/Roell Z2,5 TN Zwick-Line model) equipped with a 20 N load cell in a controlled environment laboratory (55 ± 5 % RH, temperature 20 ± 3 ºC). An in- house manufactured microbond test rig equipped with a Vernier gauge was used in this study (Figure 3.11). The two parallel blades were adjusted to 40 µm. The displacement rate was 0.2 mm/minute during the test and force-displacement curves were recorded. Despite best efforts at controlling the microdroplet size, fibre breakage still occurred (on occasion up to one-third of the specimens). Therefore, 70 specimens were produced and at least 30 specimens were expected to debond successful. The IFSS was calculated from the maximum force of debonding, assuming the shear force is uniform along the fibre using [112]:-

퐹푑푒푏표푛푑푖푛푔 IFSS = 3.7 휋퐷푙푒

Where 퐹푑푒푏표푛푑푖푛푔 = Debonding load

퐷푓 = Diameter of fibre

푙푒= Embedded length

Figure 3.10: Optical images of (a) accepted microdroplet and (b) rejected microdroplet.

Load cell

Microbond test rig

Figure 3.11: Microbond shear test configuration.

3.6.8.1 Weibull Statistical A[W14]nalysis

The Weibull statistical analysis has been employed by many investigations to describe the variability of interface shear strength (IFSS) of natural fibres composites [120]. Weibull distribution is based on the weakest link theory and it

implies that only kind of flaw lead to failure [178]. The two parameter Weibull cumulative distribution function is given by:-

흉 α 푃(흉) = 1 − exp [− ( ) ] , α ≥ 0, 흉 ≥ 0, 흉0 ≥ 0 3.8 흉0

Where 푃 is the probability of failure at any given of shear stress, 흉0 is characteristic shear strength and α is Weibull modulus respectively. The parameters of the distribution function are estimated from observation.

1 ln [ln ( )] = 훼 ln 흉 − 훼 ln 흉0 3.9 1 − 푃(흉)

Then, the 푃 values are formed in the sample on the basis ith position of 푛 ordered and -values which are estimated from observed values: order 푛 observations from smallest to largest, and let 흉(i) denote the ith smallest observation. Hence, a good estimator of Pi for median ranked values is used

𝑖 − 0.3 3.10 푃 = 푖 (푛 + 0.4)

Using equation 3.8, the reliability in which the probability of shear strength at any given stress, given by;

흉 α 푅(흉) = exp [− ( ) ] 3.11 흉0

Where 푅 is the probability of survival of material at any given of shear stress. The Weibull statistical analysis also has been adopted in single fibre tensile strength due to strength variability of natural fibres [179][180].

3.6.9 Dynamic Mechanical Analysis (DMA)

In this study, DMA was conducted using a Perkin Elmer Q8000 in dual cantilever mode (Figure 3.12). Rectangular samples with dimensions of 12 mm in width, 60 mm in length and 3-4 mm in thickness were tested. The oscillating

frequency was fixed at 1 Hz and testing was performed between 25 ºC and 220 ºC at a heating rate of 5 ºC/minute. The data of storage modulus, E’, loss modulus, E” and tan delta (damping factor) were recorded against temperature. The data were analysed using Universal Analysis 2000 software (TA Instruments). The glass transition temperature (Tg) of samples were determined from the peak of the tan delta curves.

Figure 3.12: DMA fixture set-up.

3.6.10 Tensile Testing

Tensile testing of composites was carried out in accordance with ASTM D 3039 [181] at 20 ± 2 ºC and 50 ± 5% RH using a Instron 430 Universal Testing Machine fitted with a 10 kN load cell. The tensile test specimens were rectangular in shape, having dimensions of 25 mm in width, 250 mm in length and 3-4 mm in thickness as shown in Figure 3.13(a). Strain measurement was done using a 50 mm mechanical extensometer attached to the specimen. Preliminary tensile tests were conducted to ensure no slipping in the grips occurred and failure must initiate within the gauge length for data validity as specified by ASTM D 3039 as shown Figure 3.13(b). Therefore, the use of tabs for gripping was not necessary in this study. Samples were loaded at a cross-head speed of 2 mm/min up to failure. A total numbers of 6 specimens were tested for each system. The Young’s modulus (in the linear strain range between 0.05-0.10%), tensile strength and tensile strain were calculated by the Instron Bluehill 3 software.

Figure 3.13: Photograph of (a) tensile samples and (b) fractured composites within gauge length after tensile test.

3.6.11 Flexural Testing

The flexural properties of composites were determined according to ASTM D790 [182] using an Instron 430 Universal Testing Machine fitted with a 10 kN load cell in a temperature and humidity controlled lab (20 ± 2 ºC and 50 ± 5 % RH). Specimens were rectangular, having dimensions of 13 mm in width, 105 mm in length and 3-4 mm in thickness. Specimens were tested using a three point bending fixture with an upper loading nose of radius in 6 mm and lower supports also of 6 mm in radius (Figure 3.14). The support span was set-up at 16x the specimen thickness. Through calculation, the crosshead motion was determined accordingly, varying between 1.2-1.9 mm/min. A least 5 specimens were tested for each system. The measurement of flexural modulus (in the linear strain range between 0.05- 0.10%), flexural strength and flexural strain were calculated by the Instron Bluehill 3 software.

Figure 3.14: Photograph of (a) flexural samples and (b) three point bending test configuration.

3.6.12 Interlaminar Shear Strength (ILSS)

ILSS (also known as short-beam shear strength) testing was carried out in accordance with ASTM D2344/D2344M [183] using an Instron 430 Universal Testing Machine (UTM) fitted with a 5 kN load cell. Testing was conducted in a temperature and humidity controlled laboratory at 20±2ºC and 50±5% RH, respectively. The size of specimens is dependent on the thickness of the composite, as required by the ASTM D2344/D2344M standard. The length and width shall be 6x and 2x the thickness. The test set-up was similar to 3 point bending except the lower nose was changed to a 3 mm radius. The crosshead speed was maintained at 1 mm/minute up to failure. The support span was fixed at 4x the thickness. A load drop-off of 30 % was applied after the specimens reached their highest peak load. The ILSS was calculated using:-

퐹 3.12 ILSS = 0.75 푚푎푥 푏 × ℎ

Where 퐹푚푎푥 is maximum load during the test, 푏 and ℎ is specimen width and thickness, respectively.

3.6.13 Mode I - Interlaminar Fracture Toughness (ILFT)

Mode I ILFT testing was conducted in dual cantilever beam (DCB) mode in accordance with ASTM D 5528 [184]. However, preliminary testing was a failure due to the occurrence of excessive bending, causing rupture in the beam arms rather than crack propagation. Therefore, reinforcement was applied to the outside of the arms as suggested by Bensadoun[W15] et al. [34] to increase their stiffness as shown in Figure 3.15. Initially, four layers of flax fibres were reinforced using 5 layers of woven carbon fibres but excessive bending in the beam arms still occurred without any successful crack propagation. Finally, the flax/epoxy specimens were produced with 5 layers of unidirectional carbon fibre at each face (one layer is 0.30 mm in thickness). Composite manufacturing for Mode I ILFT is similar to that described previously except that a layer of release film (Cytec ETFE) of 13 μm thickness was inserted in the mid-plane to stimulate a pre-crack. The specimens had a thickness of 6.8 mm on average. The edges of the specimens were painted white and marked 87

from the end of the release film with vertical lines every 1 mm up to 50 mm. This facilitates measurement of crack length by visual observation on a large screen of the magnified feed from a video camera.

The Mode I - DCB test requires a pair of piano hinges to be attached at the end of the specimens; fixed using epoxy adhesive (3MTM Scotch-WeldTM Epoxy Adhesive) after sand blasting of the specimen. The specimen configuration is shown in Figure 3.15. The nominal specimen dimensions were 130 mm in length, 21 mm in width and the initial crack (release film) was 50 mm. The hinges on the specimen were fitted into the grips of an Instron 4031 Universal Tensile Machine and an initial pre-crack was initiated in the specimen at a crosshead speed of 1 mm/minute between 1-3 mm in length from the insert (release film), followed by deloading and reloading for delamination measurements. All the specimens were delaminated up to 40 mm a constant crosshead speed of 1 mm/minute. Figure 3.16 shows a Mode-I

ILFT specimen during testing. The Mode I critical energy release rate, GIC, values were calculated from Modified Beam Theory (MBT)) using:-

3퐹훿 3.13 퐺 = 퐼퐶 2푏 (푎 + |∆|)

Where 퐹 = Load

훿 = Displacement

b = Specimen width

푎 = Delamination length

|∆| = Correction factor for delamination length

Figure 3.15: Specimen geometry and dimension for DCB test.

Figure 3.16: Mode I-ILFT specimen during testing.

The correction factor for delamination length, ∆ is introduced in the MBT calculation as rotation may occur at the delamination front which can overestimate

GIC. Therefore, a slight longer delamination should be applied (푎 + |∆|) in the equation. ∆ can be experimental determined by constructing a least-squares plot of the cube root of compliance, C1/3, as a function of delamination length. The compliance, C, is the ratio of the load point displacement to the applied load, 훿⁄푃. The experimental determination of the correction factor for delamination length is shown in Figure 3.17. ASTM D 5528 [184] describes 3 methods for identifying initial delamination; the visual observation (VIS), the deviation from nonlinearity (NL) and the 5% offset from nonlinearity (N) as shown in Figure 3.18.

1) The visual observation (VIS) – The VIS is the most straight forward method where the operator identifies the initial crack propagation from

the edge of composite during the testing. In this study, this method was observed through the magnified feed from a video camera. 2) The deviation from nonlinearity (NL) –was determined from the load- displacement curve. The initial delamination was considered at a point where the curve started to deviate from the linearity. 3) The 5% offset from linearity (5%/Max offset) – Similar to that NL method, an imaginary linear line were created on the load-displacement curve but the slope was 5% lower than the original initial slope. The intersection between the 5% offset imaginary linear line and the curve was considered as the crack initiation point.

0.8

0.7

0.6 y = 0.0083x + 0.0217

2 )

0.5 R = 0.9964

-1/3 N

1/3 0.4 (mm

1/3 0.3 C

0.2

0.1

0.0 -20 0 20 40 60 80 100 ∆ Delamination, a (mm)

Figure 3.17: Example of determination correction factor for delamination length, Δ for DCB test.

140 VIS Propagation Marker 120 NL 5% 100

Load (N) Load 60

0 0 10 20 30 40 Extension (mm)

Figure 3.18: Examples of determination of crack initiation for the DCB test.

3.6.14 Mode II- Interlaminar Fracture Toughness (ILFT)

Mode II-ILFT samples were reinforced with 5 layers of unidirectional carbon fibre similar to that Mode I ILFT. The Mode II ILFT testing employed four point end-notched flexure (4ENF) that followed the procedure outlined by Kuwata and Hogg [185] as shown in Figure 3.19. The 4ENF specimens were 150 mm in length and 21 mm wide. The initial crack (release film) was 50 mm and the crosshead speed used was 0.5 mm/minute. The pre-crack in the specimens were done initially between 1-3 mm crack lengths from the insert (release film). All the specimens were delaminated up to 40 mm. The Mode II critical strain energy, GIIC values were calculated from Modified Beam Theory (MBT) using:-

퐹2 휕퐶 3.14 퐺 = 퐼퐼퐶 2푏 휕푎

Where 퐹 = Load

훿 = Displacement

b = Specimen width

푎 = Delamination length

휕퐶 퐺 [W16] = Slope of the compliance to the delamination in the Compliance 퐼퐼퐶 휕푎 Calibration

Figure 3.19: Specimen and test configuration for 4ENF test.

3.6.15 Low Velocity Impact Testing

The impact testing of composites in this investigation followed the method proposed by Prichard and Hogg [169]. Their test method used smaller test specimens (89 mm in length and 55 mm wide) as compared to the ASTM D7136 standard (150 mm in length and 100 mm wide) [186]. The low velocity impact test configuration is shown in Figure 3.20. The specimens were clamped between the two plates, having a 40 mm diameter circular opening. Subsequently, the specimens were struck by a 20 mm loading nose with a mass of 1.780 kg using a Instron CEAST 9350 High Velocity Impact machine. The drop height of the loading nose was automatically adjusted to achieve the desired impact energy level. The selection of impact energy levels was based on user experience by conducting preliminary tests. Three different energies were finally selected, 1.5 J, 2.5 J and 5.0 J for all specimens. From impact testing, the peak force, contact time and absorbed energy were measured.

Hemisphere 20 mm diameter impactor

Specimen

Diameter 40 mm

Figure 3.20: Low velocity impact test configuration [169][163].

3.6.16 Ultrasonic C-Scanning Inspection

Ultrasonic C-scan is a non-destructive technique (NDT) typically employed to observe damage and defects in a composite. The impacted specimens were scanned using a Midas NDT C-scan. Test samples were positioned horizontally between two nozzles. The 5 MHz transducer probe and signal receiver were coupled to the specimen via two high velocity water jets which hit both front and back simultaneously and raster over the surface of the specimen at 100 mm/min.

3.6.17 Compression after impact (CAI) Test

The damage resistance of impacted specimens was evaluated from a post- impact compression test. This compression after impact (CAI) test followed the test method developed by Prichard and Hogg [169] with the use of an anti-buckling guide CAI fixture (Figure 3.21) to avoid out-of-plane buckling during the test. Samples of rectangular shape, 89 mm in length and 55 mm wide were positioned into the CAI fixture. The compression load was applied at 0.5 mm/minute at room temperature using an Instron 5989 equipped with a 600 kN load cell.

Figure 3.21: (a) Anti-buckling guide CAI fixture and (b) CAI experiment.

3.6.18 Scanning Electron Microscopy (SEM)

Fibre morphology and fracture surfaces of composites were studied initially using a Philips XL-30 FEG SEM. However, the School had to decommission the Philips XL-30 FEG because it was no longer being supported by the supplier and it was replaced with a Tescan Mira 3 FEG SEM. To prepare specimens for SEM, a carbon adhesive tape was first applied onto the stub onto the surface of which the specimen was placed. The specimen was then coated with platinum-gold alloy for about 5 minutes using a sputter coater (Quorum Q150T ES) to give a 9 nm thickness. Flax fibres are a biological substance and prone to damage from the electron beam of an SEM; hence, the acceleration voltage was limited to between 2 kV and 6 kV.

Chapter 4 Results and Discussions

4.1 Characterization of Untreated Flax Fibre

4.1.1 Moisture Content

Figure 4.1 shows moisture content, MC (absorption and desorption modes) for untreated flax fibres from 0% to 90% RH using the DVS analyser. The figure shows characteristic sigmoid IUPAC Type II absorption and desorption curves with hysteresis loops that are normally found for lignocellulose plant fibre such as cotton, sisal, hemp, jute and sisal [52][65]. Ceylan et al. [72] explained that a sigmoid shaped curve is a sign of maturity of a flax fibre as the secondary cell has thickened and the cellulose microfibrils almost locked into axial direction along the fibre length [26].

18 Untreated - Absorption 16 14.8 Untreated - Desorption 14 12.4 14.8 12 10.6 10 9.2 11.3 7.9 8 6.7 9.2

MC (wt.%) MC 5.5 7.9 6 4.3 6.8 4 2.9 5.6 4.5 2 3.5 0.1 2.2 0 00.0 10 20 30 40 50 60 70 80 90 100 RH (%)

Figure 4.1: MC as function of RH of untreated flax fibres at 23°C.

Figure 4.1 shows the absorption MC of the flax fibre to increase proportional to the RH between 10% and 70% but it then rises quickly to 90% RH. This behavior can be attributed to monolayer and polylayer water formation in flax fibres [71]. During moisture absorption, water molecules are directly bonded to the hydrophilic

groups in the flax fibre and this causes monolayer water formation. However further moisture absorption at > 70% RH results in more layers (polylayers) of water molecules forming on top of the existing monolayer. The average wt.% MC for each RH is given in Figure 4.1. MC increased accordingly with the RH increments from as low as 2.2 wt.% MC to as high as 14.8 wt.% MC on the absorption line. The behavior of our flax fibres is comparable with a previous study employing a DVS analyzer [69]. In term of composites manufacturing, the MC of flax fibres measured here are significantly high and could affect the interfacial properties of fibre composites. The curves in Figure 4.1 also show that the MC of the untreated flax fibres varied between absorption and desorption. This means the MC of flax fibres depends not only on the ambient RH but also on their history of exposure to ambient humidity. For instance, it appears that the MC at the 50% RH absorption line is similar to that at the 40% RH desorption line, 6.8 wt.% and 6.7 wt.% MC, respectively.

4.1.2 Relative Humidity (RH) Sensitivity

Moisture removal from natural fibres before composite fabrication is often reported in the literature [53][95][96]. Typically, the fibre is dried in an oven at 80- 105 ºC for several hours. Although it seems to be a good practice but it must be considered that MC will begin to rise immediately after drying during handling whilst producing a composite. To study this, isotherm moisture absorption tests were conducted using the DVS analyser between at 10% and 90% RH; the result of which are presented in Figure 4.2. In general, the curves exhibit typical Fickian diffusion behaviour where the MC rose quickly in a linear manner and then slowed down before levelling off as the MC reaches its equilibrium value. Using Equation 2.1, Fickian diffusion coefficients were calculated and these are tabulated in Table 4.1. The Fickian diffusion coefficients increased from 1.23 x10-9 cm2/s to 34.17 x10-9 cm2/s, reflecting that moisture absorption accelerated at higher RH. The measuring the point where half of the equilibrium MC was gained (MC0.5) showed that the 90%

RH exposed samples requires 7 minutes to achieve MC0.5 whereas 15 minutes is required for the 10% RH exposed sample. This suggests fibres will increase their MC between drying and being used to produce composites. Figure 4.2 also indicates that conditioning flax fibres for 100 minutes is more than sufficient to obtain a

desired MC as moisture uptake was minimal once it approaches equilibrium. All the data indicates that manufacturing of natural fibre composites should be done preferably under conditions of controlled humidity. In addition, creating a strategic guideline for automotive industry should be developed because every country has difference range of RH and it is useful for controlling fibre MC[AHA17] and maintaining the quality of produced composite part.

16 90% RH 14

8 MC (wt.%) MC 6 MC05 = ~7 minutes

4 MC = ~15 minutes 10% RH 2 05

0 0 20 40 60 80 100 120 Time (Minutes)

Figure 4.2: Moisture absorption curves of untreated flax fibres subjected to 10% RH and 90% RH at 23°C as a function of time.

Table 4.1: Fickian diffusion coefficient of flax fibre exposed to 10% RH and 90% RH.

Exposed RH (%) Equilibrium MC, Fickian Diffusion Coefficient, D x 10-9, (wt. %) (cm2/s)

10 14.1 ± 0.11 1.23

90 2.2 ± 0.82 34.17

4.1.3 Graphical Method for Estimating MC of Flax Fabric

Evaluating and monitoring MC is a crucial factor during fabrication of natural fibre composites. However, MC measurement is often a difficult task and time consuming because natural fibres are extremely hygroscopic (as evidently presented in Section 4.1.1 and 4.1.2). In fact, the prediction of fibre MC by measuring RH is problematical because each natural fibre has different equilibrium MCs at a given RH value. For instance, bamboo fibres absorb more moisture than 96

flax fibres due to voids inside the bamboo structure which water can occupy [123]. In this study, the MC during composites manufacturing was determined. Described in methodology, using a standardized flax fabric (with dimensions of 81 mm x 81 mm) and the MC was determined by weighing the standard fabric sample and using Equation 3.2. Data of measured fabric areal density and MC are given in Table 4.2. Alternatively, the MC can be estimated using a simple graphical method. Figure 4.3 shows such a graphical presentation for estimating the MC of untreated flax fibre (derived from Table 4.2) from which MC can be estimated by weighing the standardized flax samples to determine the weight increase during composite processing.

Table 4.2: Example of MC calculation. UT1 to UT5 is standardized fabric samples (81 mm x 81 mm).

RH RH RH 0% 10% 50% 90% 0% 10% 50% 90% 0% 10% 50% 90% 2 Samples Weight Increases (g) Fibre Areal density (g/cm ) MC (wt.%) UT-1 2.451 2.507 2.608 2.736 373.6 382.1 397.5 417.0 0 2.3 6.4 11.6 UT-2 2.539 2.597 2.701 2.878 387.0 395.8 411.7 438.7 0 2.3 6.4 13.4 UT-3 2.559 2.618 2.716 2.880 390.0 399.0 414.0 439.0 0 2.3 6.1 12.5 UT-4 2.588 2.650 2.752 2.932 394.5 403.9 419.4 446.9 0 2.4 6.3 13.3 UT-5 2.603 2.670 2.765 2.940 396.7 407.0 421.4 448.1 0 2.6 6.2 12.9 Avg 2.54 2.608 2.708 2.873 388.6 397.5 412.8 437.9 0 2.3 6.2 12.7 Std 0.06 0.06 0.06 0.08 9.07 9.66 9.4 12.4 0 0.13 0.13 0.73 Avg = Average Std = Standard Deviation

6 MC (wt.%) MC 4 UT-1 UT-2 UT-3 2 UT-4 UT-5 0

-2 2.4 2.5 2.6 2.7 2.8 2.9 3.0 Weight Increases (g)

Figure 4.3: Graphical presentation for estimating MC of untreated flax fibre. UT1 to UT5 is standardized fabric samples (81 mm x 81 mm).

4.1.4 Changes of Fibre Density

Table 4.3 presents the experimental data of fibre density with respect to MC. From this table, it can be seen that the fibre density reduced as MC increases due to swelling of the fibre or/and possible absorbed water molecules in the cell wall internal nanopores [58].The data was validated with student two-tail t-test (γ = 0.05 level of significance) which shows that the differences in the mean fibre density for each MC are statistically significant. It is interesting to note that although flax fibres can absorb MCs up to 12.8 wt.%, the decease of flax density due to moisture absorption is only up to 4.1%. The fibre density results show good agreement with those of previous studies which measured flax fibre density in the range of 1.4-1.6 g/cm3 [74].

Table 4.3: Influence of MC on flax fibre density.

RH (%) MC (wt.%) Fibre density (g/cm3) Changes (%)

0 0 1.474 ± 0.008 -

10 2.7 ± 0.39 1.464 ± 0.003 -0.7

50 6.6 ± 0.29 1.451 ± 0.005 -1.6

90 12.8 ± 0.82 1.413 ± 0.005 -4.1

As mentioned in the literature review, an attempt to describe fibre density as a function of MC using a mathematical expression was proposed by Messiry [73] but no experimental work was done to support the claim. Therefore, the experimental data in this investigation was compared with Messiry’s proposed model (Equation 2.2). To calculate the flax density against MC, the initial value of flax density was determined by applying simple linear-regression analysis on the experimental data. The linear-regression (R2 = 0.9733) was obtained for flax fibre density as a function of MC and gave an initial fibre density of 1.4775 g/cm3. Messiry’s model is plotted against MC in Figure 4.4. Interestingly, there was reasonable correlation between the experimental data and the model especially at lower RH values.

1.49

1.48 Untreated (Experimental Data) Messiry's Proposed Model

1.47 ) 3 1.46

1.45 y = -0.0049x + 1.4775 R2=0.9733 1.44

1.43 Fibre Density (g/cm Density Fibre 1.42 y = -0.0059x + 1.4753 R² = 0.9977 1.41

1.40

0 2 4 6 8 10 12 14 MC (wt.%) Figure 4.4: Comparison between current experimental data and Messiry’s model [73]. The dashed line is a linear-regression plot of the experimental data.

4.2 Interfacial Shear Strength (IFSS) of Flax/Epoxy Composites

4.2.1 Microbond Force-Displacement Curves

Figure 4.5 shows example force-displacement curves of untreated flax/epoxy composites prepared at different MC. There are two regions noticeable in the curves, fibre-matrix debonding and the pull-out region. The curves begin with a linear response that is attributed to the elastic properties of the flax fibres, signifying 99

effective load transfer at the fibre/matrix interface. The force drops abruptly after the curves reach a peak. Here, fibre-matrix debonding occurs and this debonding force is used for IFSS calculations. A sudden drop in load is a signature of brittle fracture at the interface [111]. The pull-out regions is observed because microdroplet sliding along the fibre axis with a frictional force ~50% less than the debonding force. The force fluctuations indicate that there is variation in the frictional resistance between the contacting wall inside the droplet and the fibre surface during the pull-out process. Variation of pull-out force is expected due to differences of fibre diameter, cross-sectional shapes, surface roughness and impurities along the flax fibres axis.

The specimens with 2.4 wt.% MC, 6.1 wt.% MC and 12.7 wt.% MC during composites manufacturing showed no significant differences in the shape of their force-displacement curves but with higher MC, the specimens failed at lower debonding forces. At high MC, the failure behaviour in this study is in contrast with the findings of Zhang et al. [119] on flax/unsaturated polyester composites and of Chen et al. [18] on bamboo/vinyl ester composites. These curves were shown in Figure 2.16 and both showed a broad peak in their load-displacement curves at highest MC where a gradual transition from debonding to pull-out had occurred. The curves were then decreased steadily in the pull-out region. It is worth to note that the slope of the highest MC specimens had a lower gradient compared to others MC specimens, suggesting a reduction on the longitudinal stiffness that is due to softening of flax fibre [121].

100

0.20 Untreated-2.4 wt.% MC Untreated-6.1 wt.% MC Untreated-12.7 wt.% MC 0.16

0.12

Force(N) 0.08

0.04

0.00 0.00 0.04 0.08 0.12 Displacement (mm)

Figure 4.5: Typical microbond force-displacement curves of untreated flax/epoxy composites.

4.2.2 Effect of Embedded Length

The distribution of individual calculated IFSS data with respect of their embedded length of microbond specimens is plotted in Figure 4.6. The large spread in embedded lengths formed was expected due to ‘natural variability’ during manual preparation of the microbond samples. In their pioneering works on the microbond technique, Miller et al. [112] found that IFSS was independent of the embedded length but the claimed was disputed in later studies [102][115][181–183]. In Figure 4.6, the IFSS data are highly scattered on applying simple linear regression analysis to the data very low regression R2-values of 0.123, 0.019, 0.046 were obtained for the untreated 2.4 wt.% MC, 6.1 wt.%. MC and 12.7 wt.% MC samples respectively. Hence, it is clear that the IFSS was independent of the embedded length in the present investigation.

101

12 IFSS (MPa) IFSS

8 2.4 wt.% MC 6.1 wt.% MC 4 12.7 wt.% MC

80 100 120 140 160 180 200 Embedded Length (µm)

Figure 4.6: Plot of IFSS as function of embedded length for flax/epoxy composites.

4.2.3 Interfacial Shear Strength (IFSS) and Weibull Analysis

Due to the large data scattering, the IFSS may not be easily interpreted for comparison and Weibull statistics are normally adopted for this kind of consideration especially in natural fibres [79][18][190]. This is not surprising because the IFSS calculation assumes a circular cross-section of fibres whereas flax fibres display a non-uniform geometrical characteristic [191]. The IFSS was reported by means of two-parameter Weibull distribution and average IFSS. The data is given in Table 4.4 and average IFSS plot against MC is shown in Figure 4.7.

It can be seen the Characteristic IFSS, τ0 had similar trend with average IFSS for which they become lower upon MC increases, indicating weaker fibre/matrix interfacial. At the lowest level of MC (2.6 wt.%), the average IFSS was 18.4 MPa which reduced to 16.7 MPa and 14.7 MPa for MC values of 6.3 wt.% MC and 12.9 wt.% MC, respectively.

As shown in Figure 4.7, the IFSS trend in present investigation is comparable with results for bamboo/unsaturated polyester composites [18] and flax/unsaturated polyester composites [119] where the it can be seen that the IFSS consistently reduced with increasing fibre MC (within 2 and 15 wt.% MC range).

102

Table 4.4: IFSS data for specimens manufactured with different fibre MC.

Weibull Parameters Diameter of

Samples Fibre, Characteristic Weibull Average IFSS (MPa) 푫 (µm)[AHA18] 풇 IFSS, τ0 (MPa) Modulus

Untreated-2.6 wt.% MC 19.0 ± 3.9 22.7 7.47 18.4 ± 3.4

Untreated-6.3 wt.% MC 19.6 ± 4.9 18.8 7.32 16.7 ± 2.9

Untread-12.9 wt.% MC 20.1 ± 4.3 16.9 5.53 14.7 ± 3.4

25 Flax/Epoxy Composites (Experimental Data) Flax/Vinyl Ester Composites [119] 20 Bamboo/Vinyl Ester Composites [18]

10 IFSS (MPa) IFSS

0 4 8 12 16 20 24 28 MC (wt.%) [AHA19]

Figure 4.7: Average IFSS of flax/epoxy composites as a function of MC.

In term of Weibull modulus, the results indicates that the Weibull modulus of 2.3 wt.% MC and 6.3 wt.% MC samples are about the same, at 7.47 and 7.32 respectively, but the Weibull modulus of 12.9 wt.% MC samples is reduced greatly to 5.53, suggesting a greater IFSS data scattering when the samples were produced at the highest MC. The advantage of Weibull distribution is that it can estimate IFSS with a certain probability of failure which is shown in Figure 4.8. The sigmoid curves indicate that the two-parameter Weibull distribution function provided good fitting with the experimental IFSS data. It is shown that the untreated 2.6 wt.% MC curves shifted furthest to the right followed by the 6.3 wt.% MC and 12.9 wt.% MC to left side of the graph, proposing for a certain of IFSS value i.e. 15 MPa, the probability of failure of untreated 2.3 wt.% MC, 6.3 wt.% MC and 12.9 wt.% MC specimens are 0.18, 0.32 and 0.58 respectively.

103

1.0 Untreated-2.6 wt.% MC Untreated-6.3 wt.% MC Untreated-12.9 wt.% MC 0.8

0.6

0.4 Probability of Failure of Probability 0.2

0.0 0 5 10 15 20 25 30 IFSS (MPa)

Figure 4.8: Weibull fitting on the IFSS data distribution as a function of MC

SEM imaging was conducted on specimens after IFSS testing as shown in Figure 4.9 to study the surface topography of flax fibre. Figure 4.9(a) shows the low MC sample where a much rougher surface is noticeable as compared to the relatively the relatively smooth surface of the high MC sample (Figure 4.9b(a), signifying that interfacial bonding was compromised.

104

(a)

(b)

Figure 4.9: SEM images showing (a) the rough fibre surface of low MC specimen and (b) smoother surface of high MC specimen.

105

4.3 Mechanical Performance of Untreated Flax/Epoxy Composites

4.3.1 Influence of Moisture Content on Fibre Volume Fraction

The determination fibre volume fraction, 푉푓 of a composite can be challenging due to variations in MC that can affect measurements of 푉푓. To study this, values of the theoretical 푉푓 were calculated using Equation 3.6. In Table 4.5, the theoretical 푉푓 increased with increasing MC although the composites had similar thickness. This can be misleading because MC in composites is subject to the environmental RH and the composites in this investigation were post-conditioned before testing. Hence, the MC of composites was assumed to be identical and calculation of 푉푓 should consider the same fibre density and fibre areal density. A fibre density of 1.464 g/cm3 and fibre areal density of 412.8 g/cm2 were chosen.

Table 4.5: Theoretical 푉푓 of composites based on the experimental data of fibre areal density and fibre density.

MC Fibre Density, 흆 Fibre Areal Density, (g/m2) 풇 Theoretical 푉 (wt.%) (g/cm3) 푓 0.0 388.6 1.474 0.26 2.6 397.5 1.464 0.27 6.3 412.8 1.451 0.28 12.9 437.9 1.413 0.31 It is interesting to observe in Table 4.6 that the panels manufactured with higher MCs had greater thickness after being cured. The untreated 2.6 wt.% MC samples had a thickness of 2.98 ± 0.033 and this reduced to 2.91 ± 0.019 mm for untreated 6.9 wt.% MC samples. The untreated 12.9 wt.% MC samples were the thinnest panels at 2.83 ± 0.030 mm. This result however is in contrast with Gamsted [65] and Chen et. al [18]. They showed that moisture absorption causes swelling of the natural fibres, which is larger in the radial direction than in the fibre axis direction. Hence, the possible cause is related to the composite manufacturing process. As discussed in the literature, during resin infusion (RI) the preform is laid out on a one-sides mould covered with a vacuum bag. Vacuum is applied to force resin into the preform and it also provides external compaction pressure, causing the

106

laminates to reduce in thickness. Several authors showed that compaction pressures were different for dry and wet fibre conditions (after filling with resin) [177][192]. In this investigation, a consistent method of composites manufacturing was applied and the possible reason for the variation in thickness is that fibre swelling due to the absorbed water produces reduced resistance to the vacuum-only compaction force applied to the preforms during RI, hence causing the thickness of the cured composites to reduce.

In Table 4.6, the differences in composites thickness produce slight changes in composite density and theoretical values of 푉푓. The untreated-12.9 wt.% MC samples had the highest 푉푓 at 0.30 whilst the untreated 6.3 wt.% MC and 2.9 wt.%

MC samples had 푉푓 values of 0.29.

Table 4.6: Thickness, density and theoretical 푉푓 of untreated flax/epoxy composites.

Theoretical Composite Thickness, h Composite Density, 흆풄 Samples 3 (mm) (g/cm ) 푉푓 Untreated-2.9 wt.% MC 2.98 ± 0.03 1.2303 ± 0.003 0.29 ± 0.003

Untreated-6.3 wt.% MC 2.91 ± 0.02 1.2333 ± 0.002 0.29 ± 0.002

Untreated-12.9 wt.% MC 2.83 ± 0.03 1.2335 ± 0.002 0.30 ± 0.003

4.3.2 Tensile Properties

Figure 4.10 shows typical tensile stress-strain curves of untreated flax/epoxy composites with 2.8 wt.% MC, 6.3 wt.% MC and 12.9 wt.% MC. The curves was normalised at 30% 푉푓 for comparison purposes. All tested specimens failed within the gauge length of the specimen. The stress-strain curves display two distinct of zones corresponding to the changes of slope. The first is a straight line elastic region; the stress and strain are directly proportional corresponding to longitudinal modulus. The curves drop to lower gradients in the second zone over a small strain range (0.3-0.4%) and this second zone covers the final ~60% of the tensile curves until sudden failure occurred, indicating brittle fracture. This behaviour is commonly found for woven fabric reinforced composites [123][187–189]. The two different zones suggest that the composites lose some rigidity after a certain point. In the

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second zone, it is believed that the composites sustain irreversible damage due to straightening of the interlaced fibre yarns under tensile loading [196][194].

Tensile(MPa) Stress 20 Untreated-2.8 wt.% MC 10 Untreated-6.3 wt.% MC Untreated-12.9 wt.% MC 0 0.0 0.4 0.8 1.2 1.6 2.0 Tensile Strain (%)

Figure 4.10: Typical tensile stress-strain curves of untreated flax/epoxy composites with various fibre MC during composites fabrication. The curves were

normalised at 30% 푉푓.

With increasing MC, it is interesting to observe that the gaps between curves become wider in the second zone. The inflection point for which this started to change occurs at lower strain at higher MC. Apart from weak interface properties, such behaviour can be explained by the softening of the cellulose structure in the fibres with increasing MC. The bound water forces cellulose molecules apart, reducing the rigidity of the cellulose structure [64]. This finding is in agreement the findings of Moudood et al. [150], for unidirectional flax/epoxy composites, who attributed this behaviour to enhanced sliding and delamination of cellulose fibrils due to moisture. Based on X-ray photoelectron spectroscopy (XPS) results, Fuentes et al. [121] explained the possibly of that humidified elementary fibres (lignin and pectin) may dissolve, causing fibre defibrillation and ultimately a further reduction in longitudinal strength.

Table 4.7 shows the derived tensile properties for each composite. Tensile strength and tensile modulus were normalised at 30% 푉푓. An increase of MC causes a consistent reduction in tensile strength and tensile modulus. This may be attributed

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to the effect of MC on the fibres, as discussed above, but also that moisture reduces the interfacial properties in these fibre composites as shown by the microbond shear test. SEM micrographs of fractured surfaces in Figure 4.11(a) and Figure 4.11(b) show evidence of a weakening of the fibre/matrix interfaces. In Figure 4.11(a), there is evidence of peeling fibre surfaces resulting from a strong fibre/matrix interface for the untreated 6.3 wt.% MC samples. In contrast, there are gaps in the matrix surrounding the fibres corresponding to weak interface for the untreated 12.9 wt.% MC samples shown in Figure 4.11(b). These results are in good agreement with previous findings in the literature, in which MC caused reductions in tensile properties especially at high MC (>50% RH) [149][150].

Table 4.7: Tensile properties of untreated flax/epoxy composites with various fibres MC during composites fabrication.

Normalised Tensile Normalised Tensile Tensile Strain at Samples Strength (MPa) Modulus (GPa) Break (%)

Untreated-2.8 wt.% MC 67.4 ± 2.1 6.19 ± 0.2 1.47 ± 0.1

Untreated-6.3 wt.% MC 62.6 ±.1.6 6.18 ± 0.1 1.62 ± 0.1

Untreated-12.9 wt.% MC 57.3 ± 1.5 5.75 ± 0.2 1.74 ± 0.1 Increased fibre MC allows the specimens to elongate further before failure, and the 12.9 wt.% MC composites samples stretched about 18% further than the 2.9 wt.% MC samples. This is an interesting finding because the majority of literature studies [129][140][197] found stronger composites to have higher tensile strain at failure, unlike the present investigation. Higher failure strains for natural fibre composites after being humidified was reported by Thuault et al. [77] for flax fibres and by Chen et. al [18] for bamboo strips. It is worth noting, however, that their tensile tests were performed under the desired humidity conditions in a climate chamber, whereas the composites specimens in this investigation were post- conditioned as described in Section 3.5, indicating that the MC of composites can be assumed to be identical prior testing. A possible explanation is that the softening effect of water on fibre cellulose structure is suggested to be an irreversible process that occurs during the cure, increasing the tensile strain at break of composites. This argument is supported by Baley and co-worker [59] studies of dried and non-dried flax fibre reinforced epoxy composites. Although they did not specify the samples’

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conditioning, the experimental data was collected in an environmentally-controlled laboratory in accordance to NF T 57-101 ISO 527. They demonstrated that the dried flax fibres composites gave lower composite tensile strains in tensile testing compared to the non-dried flax composites.

(a)

Fibre surface peeling

(b)

Figure 4.11: SEM images showing (a) strong fibre/matrix interface of a 6.3 wt.% MC composite and (b) poor fibre/matrix interface of a 12.9 wt.% MC composite.

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4.3.3 Flexural Properties

Figure 4.12 shows typical flexural stress-strain curves for flax fibre composites loaded in 3 point bending. The flexural curves were normalised to 30%

푉푓 for comparison purposes. Similar to tensile behaviour, two regions are noticeable in the stress-strain curves, beginning with a linear, elastic region. The second regions started around 0.4-0.5% strain after which stress-strain behaviour is non- linear. Here, the composites with higher MC have a lower gradient change, starting at 0.025% flexural strain which suggests the onset of damage. The widening effect between 2.9 wt.% and 12.9 wt.% MC curves in the second region is due mainly to a combination of poor interfacial properties and the softening effect of MC on flax fibre as explained previously. The composites finally fail abruptly upon their peak stress, signifying brittle fracture. All failed samples showed a single tensile crack at the bottom face at the mid-span.

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100

FlexuralStress (MPa) Untreated-2.1 wt.% MC 20 Untreated-6.2 wt.% MC Untreated-12.6 wt.% MC 0 0.0 0.5 1.0 1.5 2.0 2.5 3.0 3.5 Flexural Strain (%)

Figure 4.12: Typical flexural stress-strain curves of untreated flax/epoxy composites with various fibre MC during composites fabrication. The curves were

normalised at 30% 푉푓.

Table 4.8 shows the normalised flexural properties of the composites. As MC was increased from 2.9 wt.% MC to 6.3 wt.% their flexural strength was consistent at ≈95 MPa. Flexural modulus, however, decreased significantly by about 10 MPa as MC was increased. For flexural modulus and flexural strain, there was a

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consistent decreasing trend of flexural modulus accompanied with increasing strain across the MC range, similar to that found for tensile properties. To corroborate the finding, the fractured surfaces of composites examined using SEM is shown in Figure 4.13. It is evidence that there was more fibre fracture for the untreated 2.1 wt.% MC sample in Figure 4.13(a) whereas, in contrast, fibre pull-out was predominantly found on the fracture surfaces of the untreated 12.9 wt.% MC sample, suggesting poor fibre/matrix interface and corresponding to reduce flexural properties. These results are in good agreement with literature results [18][119][150] all of which propose a negative effect of MC on the fibre-matrix interface.

In term composites performance under influence of MC during fabrication, the experiment data from Moudood et al. [150] showed the flexural strength and flexural modulus losses about 32% for unidirectional flax/epoxy composites that produced with highest MC (~14 wt.% MC) whereas in present investigation, both flexural strength and modulus has minimal losses (~11%) at the same MC. This suggests the detrimental effect of poor interface properties could be limited with the use weave fabric.

Table 4.8: Flexural properties of untreated flax/epoxy composites with various fibres MC during composites fabrication.

Normalised Normalised Flexural Flexural Strain at Samples Flexural Strength Modulus (GPa) Break (%) (MPa)

Untreated-2.1 wt.% MC 95.40 ± 3.2 4.93 ± 0.3 2.81 ± 0.2

Untreated-6.2 wt.% MC 94.34 ± 3.8 4.68 ± 0.1 3.08 ± 0.2

Untreated-12.6 wt.% MC 85.55 ± 3.8 4.38 ± 0.1 3.15 ± 0.1

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Figure 4.13: SEM images showing fracture surface of (a) an untreated 2.1 wt.% MC sample and (b) an untreated 12.6 wt.% MC sample.

4.3.4 Interlaminar Shear Strength (ILSS)

The effect of fibre MC during composites manufacturing on the laminate strength was assessed by a short beam shear test used often to assess the interface strength of fibre composites [192]–[194]. Figure 4.14 shows typical load-deflection curves for untreated flax/epoxy composites manufactured at 6.3 wt.% MC and 12.8 wt.% fibre MC. The curves in Figure 4.14 do not show similar trends to those from the flexural tests, as these curves exhibit non-linear behaviour from the beginning and continue to increase until peak load is reached. The curves become irregular after substantial drops occurred, signifying sequentially failure of the composites. 113

0.5

0.4

0.3

Load(kN) 0.2

0.1 Untreated 6.3 wt.% MC Untreated 12.8 wt.% MC 0.0 0.0 0.2 0.4 0.6 0.8 1.0 1.2 1.4 1.6 1.8 Extension (mm)

Figure 4.14: Typical load versus deflection curves of untreated flax/epoxy composites manufactured at 6.3 wt.% and 12.8wt.% fibre MC.

Unfortunately, further investigation on the failure mode of the fractured samples revealed that for these composites the short beam shear test was invalid. All the samples failed in tension, as described by ASTM D2344/D2344M [183], with cracks visible on the bottom face of the composites specimens as shown in Figure 4.15. This is corroborated by the ILSS results (Table 4.9) which the difference between the two means is not statistically significant using student two-tail t-test (γ = 0.05 level of significance), although the ILSS results reduced from 154.7 ± 18.2 MPa to 147.0 ± 8.4 MPa. To explain this, it is essential to understand the mechanism of short beam shear strength during the test. The determination of ILSS is based on the classical beam theory of a homogenous body. This describes that maximum shear stress occurs in the neutral plane where the normal stress are zero [201]. In composites, on the other hand, this can be complicated due to other stresses [202] and the failure mode can be combination of fibre breakage, micro buckling and interlaminar shear cracking [159]. The interlaminar failure will occur when the mid- plane shear stress exceeding the interfacial shear stress or the matrix shear stress [203]. The presence of fibre crimp in the flax fabric could create an interlocking mechanism between the laminates creating greater sliding motion resistance. Low bending modulus of the flax/epoxy composites accompanying enhanced interlaminar shear strength, can impart excessive stress to the fibres which will break rather than delaminate.

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Figure 4.15: Photograph of ILSS samples (untreated 6.3wt.% MC) showing tension failure,

Table 4.9: Average ILSS data of untreated flax/epoxy composites with various fibre MC.

Samples ILSS (MPa)

Untreated-6.3 wt.% MC 154.7 ± 18.2

Untreated-12.8 wt.% MC 147.0 ± 8.4

4.3.5 Dynamic Mechanical Analysis (DMA)

4.3.5.1 Storage Modulus

In DMA, the storage modulus, E´ is a measurement of elasticity of the specimen, corresponding to the recoverable energy stored during a loading cycle. The dynamic mechanical properties of epoxy and untreated flax/epoxy composites manufactured with different fibre MC were evaluated. E´ as a function of temperature is shown in Figure 4.16. E´ of all specimens decreased steadily with increasing temperature, starting in the glassy region of the epoxy matrix of the epoxy matrix but then there is a sharp drop of modulus as it goes through its glass transition as it goes through its glass transition prior to it to it reaching the rubbery plateau. The decreased modulus across the temperatures is due to dilution of polymers, corresponding to the increased molecular mobility of the polymer chains [204]. A lower E´ was obtained for the epoxy across the temperatures, indicting the influence of reinforcement to increase the stiffness and thermal stability of the composites. Other authors have also reported similar findings [199]–[201].

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Previous studies reported that E´ is sensitive to the interface properties of composites [208][209]. Figure 4.16 showed that with increasing temperature from the glassy region passing through the glass transition (Tg) that composites manufactured at higher MC exhibited lower E´. This can be attributed to a decrease in the constraint of polymer segmental mobility at the fibre-matrix interface [101][204][205]; or more simply, composites produced at lower MC tend to have greater fibre/matrix adhesion. It is interesting to observe that the E´ of the composites are very similar above 180 ºC. A possible reason for this behaviour was discussed by Jawaid et al. [207] and Jabbar et al. [212] in that in the rubbery region, the intermolecular forces become weak and the polymer chain segments chain segments become more mobile and lose their close packing as the temperature rises. This means factors such as fibre rigidity and fibre-matrix interface degradation due to higher MC are as dominant in the rubbery region.

3500

3000

2500

2000

1500

1000 Epoxy StorageModulus (MPa) 500 Untreated-2.9 wt.% MC Untreated-6.3 wt.% MC 0 Untreated-12.9 wt.% MC

30 60 90 120 150 180 210 o Temperature ( C) Figure 4.16: Temperature dependence of storage modulus, E´ of untreated flax/epoxy composites at different fibre MC.

4.3.5.2 Loss Modulus

Loss modulus, E” is the stress 90º out-of-phase with the strain divided by the strain. It is associated with the viscous response of a material and molecular motions [106][213]. Figure 4.17 shows the variation of E” as a function of temperature for composites with different fibre MC during fabrication. As the temperature increases, it is observed that a peak in E” occurred associated with matrix glass transition for

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all samples. It is seen that the peak E” increases with MC, which is which is associated with an increase in internal friction [214].

Epoxy 300 Untreated-2.9 wt.% MC Untreated-6.3 wt.% MC 250 Untreated-12.9 wt.% MC

200

150

100 LossModulus (MPa)

0 30 60 90 120 150 180 210 Temperature (oC)

Figure 4.17: Temperature dependence of loss modulus, E” of untreated flax/epoxy composites at different fibre MC.

4.3.5.3 Tan Delta

Loss tangent, or tan delta loss tangent, or tan delta (ratio of E´ and E”) is a measurement of damping properties corresponding to the energy dissipation of the materials. Figure 4.18 shows the tan delta of epoxy and untreated flax/epoxy composites as the temperature increases, it is observed that a peak in E” occurred associated with matrix glass transition for all samples. It is interesting to note that the height of tan delta peak of the unfilled system is very high which represents higher energy dissipation. In case of the composites, lower tan delta is expected as compared to neat epoxy because the decreased volume fraction of the matrix upon the incorporation of fibres [23]. Ornaghi et al. [213] explained that in the event of dynamic mechanical loading, fibres carry most of the load and only a only a small part of it generates strain at the interface. This means the energy dissipation will occurs in the polymer matrix at the fibre interface.

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0.9 0.30 0.8 0.25 0.7 0.20

0.6 0.15

0.5 0.10

0.05

TanDelta 0.4

0.3 120 140 160 180 200 Epoxy 0.2 Untreated-2.9 wt.% MC Untreated-6.3 wt.% MC 0.1 Untreated-12.9 wt.% MC

0.0 30 60 90 120 150 180 210 o Temperature ( C)

Figure 4.18: Temperature dependence of tan delta of untreated flax/epoxy composites at different fibre MC.

The glass temperature (Tg) obtained from the peak in tan delta gives an indication of the degree of crosslinking and the strength of the fibre/matrix interface in the system [106][204][215]; i.e. a higher Tg indicates higher crosslink density and a stronger fibre-matrix interface. In Table 4.10 , the Tg of neat epoxy is 139.6 ºC and the Tg is shifted to higher temperature with incorporation of fibres (156.6 ºC - 162.0 ºC). This is due to the restriction of chain segmental movement at the interface induced by interaction with the stiff fibre. For a stronger composite with higher interface adhesion, the movement of the polymeric chain become more restricted and a higher Tg is expected. The influence of MC during composites fabrication was seen here with the Tg of composites reducing with increasing MC; 2.9 wt.% MC (162.0 ºC) > 6.3 wt.% MC (161.4 ºC) > 12.9 wt.% MC (156.6 ºC), suggesting a lower degree of cross-linking and weaker fibre-matrix interface. In other words, the MC of the fibre has plasticised the matrix.

Many previous studies have demonstrated that a higher tan delta peak height is indicative of poor interfacial adhesion [210][212][215]. However, it is interesting to note that the tan delta values of composites in this study were consistently lower with increasing MC; 2.9 wt.% MC (0.28) > 6.3 wt.% MC (0.27) > 12.9 wt.% MC

(0.25). Such behaviour (increases in Tg and reduces tan delta with higher MC) were also reported by Costa et al. [216] who studied the influence of hygrothermal

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conditioning of cured carbon/epoxy and carbon/bismaleimide composites. They described the reduced tan delta (epoxy-based composites) of moist samples compared to the dried samples as resulting from moisture absorption behaviour in the composites during hygrothermal conditioning. They also reported that the

reduced Tg was consistent with the weak interlaminar shear strength (ILSS) results of moist composites, indicating weak interface properties. In the present study, however, the all flax/epoxy composites were post-conditioned and it is assumed all composites had similar MC prior to DMA testing. Comparison of our experimental data and the reported literature [216][217], suggests there is a more complex behaviour in our composites systems due to the presence of moisture in the fibre during curing which could affect the properties of the fibre and the molecular structure of the forming polymer matrix. These parameters might be evaluated simultaneously in the case of DMTA analysis and interpretation of this behaviour can be challenging.

In attempt to explain this behaviour, further analysis was carried out on full width at half maximum (FWHM) of the tan delta curves and results were tabulated in Table 4.10. The outcome results is clear, higher tan delta FWHM values were obtained with increasing MC, suggesting a broader interface (interfacial area per unit volume of material) [210]. The result is consistent with fibre volume fraction of composites that were slightly increased with increasing MC (Table 4.10). This means although the changes of fibre volume fraction was relatively small, it will have direct impact on the tan delta peak height.

Table 4.10: The 푉푓, tan delta peak height, tan delta FWHM, Tg and interfacial strength indicator of untreated flax/epoxy composites.

Interfacial Tan Delta Tan Delta Samples 푉 T (ºC) Strength 푓 FWHM (ºC) g Peak Height Indicator, B Epoxy - 0.83 ± 0.008 28.52 ± 0.64 139.6 ± 0.4 -

Untreated-2.9 wt.% MC 0.304 0.28 ± 0.005 22.18 ± 0.49 162.0 ± 0.1 2.36

Untreated-6.3 wt.% MC 0.296 0.27 ± 0.003 23.30 ± 0.64 161.4 ± 0.2 2.32

Untreated-12.9 wt.% MC 0.283 0.25 ± 0.006 25.66 ± 0.13 156.6 ± 0.1 2.29

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According to Shinoj et al. [218], tan delta can be used to characterize interfacial properties if only matrix and fibre volume fraction and fibre orientation are identical. The characterization of fibre/matrix interface using DMA assumes that the composite’s dissipation is attributed not only to the matrix, but also governed by the fibre/matrix interaction at the interfaces that will tend to form layers of immobilised interface [219]. Based on this consideration, an interfacial strength indicator, B was introduced to correct the fibre volume fraction variations in the composites [218][220][221]. B is given by:

푇푎푛 푑푒푙푡푎푐표푚푝표푠푖푡푒 퐵 = (1 − )푉푓 4.1 푇푎푛 푑푒푙푡푎푚푎푡푟푖푥

Higher B value indicates greater adhesion strength. The value interfacial strength indicator, B is tabulated in Table 4.10. It can be seen the interfacial strength of composites decreases with increasing MC; 2.36, 2.32 and 2.29 respectively

4.3.5 Mode I DCB - Interlaminar Fracture Toughness (ILFT)

Mode I DCB testing was performed to determine the interlaminar fracture between the laminates. Typical load-extension curves for untreated 2.3 wt.% MC, 6.6 wt.% MC and 12.7 wt.% MC DCB specimens are shown in Figure 4.19. In general, the curves begin with a linear region that is attributed to the elastic coefficient of the composites. There was a slightly deviation of linearity in the curves for which the nonlinear crack initiation, GIC initiation NL was considered. The curves then followed a declining slope, signifying the crack propagating into to the specimens. It is interesting to notice that the untreated 12.7 wt.% MC sample exhibits a relatively smooth curve, whereas the untreated 2.3 wt.% MC and 6.7 wt.% MC samples show extensive jagged shape curves. According to Elias [154], such behaviour is consequences of a sudden release of strain energy because of the sudden failure of constituent components. This can be either crack initiation at the crack tip of a matrix or the fibre reinforcement. In this study, it is likely the sudden release of strain energy is due to improve fibre/matrix interface of low MC samples, giving unstable crack propagation in the untreated 2.3 wt.% MC and 6.7 wt.% MC samples.

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Figure 4.20 shows typical resistance fracture curve (R-curve) as a function of crack length from the from the Mode I DCB test. The GIC for crack propagation (GIC propagation) value was calculated using Modified Beam Theory (MBT) as suggested by ASTM D 5526. All the curves display the same rising profile where the GIC propagation- initially starts at low values and rises as crack length increases. This is attributed to the fibre bridging mechanism as reported by other researchers [222][223].

200 180 160 140 120 100 Load(N) 80 60 40 Untreated-2.3 wt.% MC 20 Untreated-6.6 wt.% MC Untreated-12.7 wt.% MC 0 0 5 10 15 20 25 30 35 40 Extension (mm)

Figure 4.19: Typical load extension curves of untreated 2.3 wt.% MC, 6.6 wt.% MC and 12.7 wt.% MC DCB specimens.

3.2

2.8

2.4

) 2.0 2

1.6

(kJ/m ic

G 1.2

0.8 Untreated-2.3 wt.% MC 0.4 Untreated-6.6 wt.% MC Untreated-12.7 wt.% MC 0.0 0 5 10 15 20 25 30 35 40 45 Crack length, a (mm)

Figure 4.20: Typical resistance curves (R-Curves) of untreated 2.3 wt.% MC, 6.6 wt.% MC and 12.7 wt.% MC DCB specimens.

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Figure 4.22(a) shows the average GIC initiation for the NL, VIS and 5%

Offset while Figure 4.22(b) displays GIC propagation of the untreated flax/epoxy composites. The interlaminar factures toughness is attributed by many factors and it is reported the strong fibre/matrix interface has also play important role to improve

GIC [157][223][224]. Contrary to those findings, the GIC propagation (Figure 4.22(b)) in the present investigation increased with increasing MC from 1.81 kJ/m2 to 2.48 kJ/m2 but then reduced to 2.18 kJ/m2. Nathan [86] also observed an increase 2 2 of GIC propagation from 249.9 J/m (without MC) to 269.6 J/m for carbon/epoxy prepreg manufactured with 0.5 wt.% MC during composites fabrication. When the composite was manufactured with 2.0 wt.% MC, the GIC propagation decreased 2 again to 231.0 J/m . He attributed this behaviour (Increases GIC propagation with MC increasing) to the plasticization effect of the epoxy matrix. For natural fibre composites, it is known that MC has direct consequences for the fibre, matrix and fibre/matrix interface and any changes to these parameters could affect the final properties of a composite [150]. If the matrix was responsible for the gain in GIC, perhaps the reason could be that high MC during composites manufacturing had toughened the matrix. Toughening of the of the matrix, one of the main methods to increase the interlaminar GIC of composites [163][222][225], as a results in a larger plastic zone opening at the crack tip.

According to Wu [163], a brittle matrix system generally has similar NL and VIS values whereas for a tough matrix, the NL initiations occur before the VIS (Figure 4.21). Figure 4.22(a) shows the NL and VIS values were values were almost similar for the untreated 2.3 wt.% MC composites. In contrast, the NL was higher than the VIS for the 6.6 wt.% MC and 12.7 wt.% MC samples, suggesting a toughening of the matrix. Another interesting finding was the colour change in in the composite samples observed as shown in Figure 4.23 which can possibly corroborate the argument. The untreated 2.3 wt.% MC sample had a dark brown colour compared to the 12.7 wt.% MC sample which exhibited a lighter brown appearance. Colour changes may indicate there was a chemistry change within the epoxy matrix [199][226]. The possible explanation of the toughening of the epoxy matrix can be due to reduced crosslink density [199][227] which correlates with the reduced Tg of composites from DMA. Therefore, it is suggested the toughening

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effect of the epoxy matrix had surpassed the negative effect of weak fibre/matrix interface in interlaminar fracture toughness of untreated flax/epoxy composites.

Figure 4.21: Force-displacement curves of (a) brittle and (b) ductile matrix [228].

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(a) 2.2

2.0

) 2 1.8

1.6 Initiation (kJ/m Initiation

IC 1.4 G NL 1.2 VIS 5% Offset 1.0 0 2 4 6 8 10 12 14 MC (wt.%)

(b)

2.8 2.48 ) 2 2.18 2.4

1.81

2.0

Propagation (kJ/m IC

G 1.6

1.2 0 2 4 6 8 10 12 14 MC (wt.%)

Figure 4.22: Influence of MC on Mode I (a) GIC initiation (NL, VIS, 5% Offset) and

(b) GIC propagation as a function of MC of flax/epoxy composites.

Figure 4.23: Photograph of delaminated composites showing colour changes between untreated 2.3 wt.% MC and 12.7 wt.% MC samples.

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Figure 4.24 shows qualitative observation [AHA20]of fibre bridging that was visible on the specimens from the Mode I DCB test. Previous studies have shown fibre bridging was found on flax composites during Mode I DCB tests [34][158][229]. Fibre bridging occurs due to unaligned fibres across the crack plane and a crack trip that has grown in into many layers and planes [154]. Johnson and Mangalgiri [230] described fibre bridging associated with a weak fibre/matrix interface similar to the effect shown in Figure 4.24. In Figure 4.24(a), the untreated 2.2 wt.% MC samples developed less fibre bridging across the deformation plane as compared to the untreated 6.6 wt.% MC and 12.6 wt.% MC samples in Figure 4.24(b) and Figure 4.24(c). In other words, weak fibre/matrix adhesion causes bridged fibres which were [W21]easily pulled-out from the matrix as the delamination advances. This might explain why the GIC values (initiation and propagation) of the untreated 13.7 wt.% MC samples in Figure 4.22 were reduced to 2.18 kJ/m2 from 2.48 k/m2 although there was evidence of epoxy matrix toughening.

Apart from the fibre bridging phenomenon, an additional damage mechanism during Mode I DCB testing can be seen in Figure 4.25 that can be attributed to the presence of MC during composites fabrication. In Figure 4.25(a), there was a long peeling failure on the fibre surface of the untreated 2.3 wt.% MC sample. This is not surprising because flax fibres have a complex layered internal structure (i.e primary and secondary walls) and the peeling was possibly due to primary wall detachment. Baley and co-workers [59] found the same observation of flax fibres peeling on the dried flax fibres after tensile tests of composites. In Figure 4.25(b), however, there was a unique failure where fibre splitting along the fibre axis was observed for the untreated 12.7 wt.% MC sample, signifying the flax fibre was damaged internally due to presence of high MC. Closer observation the fibre in Figure 4.25(b) indicates microfibril bridging inside the flax fibre. As described earlier, natural fibres swell with the presence of moisture particularly in the radial direction [18][65][66]. It has been has been shown that both thermosetting matrix and flax fibre experienced shrinkage during cure [231]. Due to the to the thermal expansion (CTE) mismatch between flax fibre and the epoxy, this could create residual stresses in the cured composites. Therefore, one possible explanation is that fibres swollen due to high MC induced greater residual stress development, ultimately damaging the fibre structure in the radial direction. 125

Figure 4.24: Qualitative observation [AHA22]of fibre bridging (white arrows) on Mode I DCB test showing fibre bridging of untreated (a) 2.3 wt.% MC (b) 6.6 wt.% MC and 12.7 wt.% MC samples.

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Microfibril bridging

Figure 4.25: Damage within a fibre of untreated (a) 2.2 wt.% MC and (b)12.7 wt.% MC samples.

4.3.6 Mode II ENF - Interlaminar Fracture Toughness (ILFT)

The study proceeds with mode II interlaminar fracture toughness for which the manufacturing process is described. Similarly to the interlaminar shear strength (ILSS) test, Mode II ILFT tests could not be determined although specimens were reinforced with unidirectional carbon fibre to enhance their bending stiffness. Preliminary tests on the untreated 6.6 wt.% MC composites and untreated 12.7 wt.% MC composites showed that the crack propagation deviated severely (after a pre- crack of 3 mm) from the mid-plane into adjacent flax-flax laminates, thus

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invalidating the result [232]. The evidence of crack deflection during the Mode II ENF test is shown in Figure 4.26. Crack deviation is due to additional energy of the crack propagation [233]. In the worst cases, the crack tip continued to progress into flax-carbon laminates when delamination were growing as shown in Figure 4.27. In fact, the deviation of crack path caused inconsistent crack growth between laminates i.e. the crack tip reached the carbon fibre fabric at different crack length as shown in Figure 4.27.

Figure 4.28(A1) and Figure 4.28(A2) show the load-displacement curves of untreated 6.6 wt.% MC and 12.6 wt.% MC samples. These showed similar behaviour and it was hard to notice any sign of crack deviation into other adjacent laminar in the curves. However, there was a sudden decreased of loading as shown in Figure 4.28(A2), indicating that the crack path had progressed from the flax-flax laminates into flax-carbon laminates. The GIIC propagation was calculated and presented in Figure 4.28(B1) and Figure 4.28(B2) for untreated 6.6 wt.% MC and untreated 12.6 wt.% MC samples. Although they were the same samples, there was a large variation of GIIC propagation and this possibly was due to crack deflection during the Mode II ENF test. As a result, getting the correct GIIC propagation values is questionable and unreliable. Hence, the Mode II 4ENF testing on other samples (different MC and fibre treatment) was not carried out.

It is hard to explain why the crack path deflects from the mid-plane in the Mode II ENF test based on current data except the enhanced fracture toughness resistance of flax/epoxy composites. The following factors could contribute:

The presence of yarn crimp in the fabrics - The mode I-DCB test employs tension loading to create a crack path between the laminar where the presence of crimps might be limited, resulting in crack growth in the mid-plane. In contrast, Mode II-ENF is in-plane shear where parallel loadings are created by bending. It is reported that Mode I is higher than Mode II values [34][234]. The yarn crimp can create additional mechanical interlocking [235] and result in crack deviation.

Insufficient bending stiffness - The Mode I-DCB and Mode II-ENF tests were developed for unidirectional fabric laminates. For a woven laminates, in most

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cases when not using very high modulus fibres such as flax, modification of specimens is necessary by employing reinforcement of the bending arms with stiffer material to enhance bending modulus. Due to the toughness of the weave fabric architecture, the current reinforcement with UD carbon fabrics might be in insufficient to generate the delamination resistance for the Mode-II-ENF test.

Figure 4.26: Photograph of side view of composites after Mode II ENF test of untreated (a) 6.6 wt.% MC and (b) 12.6 wt.% MC samples.

Carbon fibre fabric

Flax fibre fabric 20 mm

Figure 4.27: Photograph of delaminated of the untreated 12.7 wt.% MC sample surfaces showing the consequences of crack deflection from the mid- plane where the crack path eventually reached the carbon fibre fabric.

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Load-extension curves R-Curves

(A1) 3.5k (B2) 30

3.0k 25 2.5k

) 20 2.0k 2 15 Untreated-6.6 wt.% MC-1 Untreated-6.6 wt.% MC-2

1.5k (kJ/m Load (N) Load IIC 10 1.0k G

500.0 Untreated-6.6 wt.% MC-1 5 Untreated-6.6 wt.% MC-2 0.0 0 0 1 2 3 4 5 6 7 0 5 10 15 20 25 30

Untreated 6.6 wt.% MC wt.% 6.6 Untreated Crack length, a (mm) Extension (mm)

(A2) 3.5k (B2) 30 3.0k 25 2.5k

) 20 2 2.0k Untreated-12.6 wt.% MC-1 15 Untreated-12.6 wt.% MC-2

1.5k (kJ/m Load (N) Load Delamination between flax IIC 10 1.0k and carbon flax laminates G

500.0 Untreated-12.6 wt.% MC-1 5 Untreated-12.6 wt.% MC-2 0.0 0 0 1 2 3 4 5 6 7 0 5 10 15 20 25 30

Untreated 12.6 wt.% MC wt.% 12.6 Untreated Crack length, a (mm) Extension (mm) Figure 4.28: Typical load-displacement curves and R-curves of Mode II ENF of untreated flax/epoxy composites.

4.3.7 Low Velocity Impact

4.3.7.1 Impact Behaviour

Typical load-deflection curves at two MCs during composites fabrication at various impact energies is presented in Figure 4.29(a-c). The behaviour of the curves is suggested to be a closed form because there is no penetration or perforation occurred in the composites at 1.5 J, 2.5 J and 5.0 J impact energies. The curves exhibited a smooth descent concurrent with decreasing force and displacement, indicating the rebounding effect of the impactor from the samples [236][162]. Nevertheless, the ascend part of the loading phase showed two different characteristics. It begins with a linear portion until the onset of damage [80][237], represented by the significant changes of gradient, which implies that bending stiffness has reduced and damage may result from matrix cracking, fibre fracture and delamination [154][159][231][232].

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For a lower impact energy at 1.5 J in Figure 4.29(a), the bending stiffness at the linear region of 6.5 wt.% MC is higher than that 12.5 wt.% MC samples and this correlates with the flexural properties. It is interesting to observe that the untreated 12.5 wt.% MC samples experienced a slight reduction in reduction in peak force but it regained this up to the peak. This could be indication of load redistribution upon impact loading, leading to a broader plateau region compared to the untreated 6.5 wt.% MC samples. Poor fibre/matrix interfacial properties cause further damage i.e. greater fibre separation from the matrix and ultimately, fibre breakage and reduced bending stiffness. As impact energy increases to 2.5 J shown in Figure 4.29(b), the 6.5 wt.% MC samples still showed higher bending stiffness (in the in the linear region), maximum peak force and maximum deflection upon completion compared to that 12.5 wt.% MC.

Figure 4.29(c) shows force-deflection curves for 5.0 J impact energy. In contrast with Figure 4.29(a) and Figure 4.29(b), there is no significant difference in in the bending between the 6.5% wt. MC and 12.5% wt. MC samples and in fact, they have the same peak force and deflection.

(a) 2000 1.5J

1600

1200 plateau region

Force(N) 800

400 Untreated 6.5 wt.% MC Untreated 12.5 wt.% MC 0 0 1 2 3 4 5 6 Displacement (mm)

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(b) 2000 2.5J Untreated 6.5 wt.% MC 1600 Untreated 12.5 wt.% MC

1200

Force(N) 800

400

0 0 1 2 3 4 5 6 Displacement (mm)

1600 Fibre Fracture

1200

Force(N) 800

400 Untreated 6.5 wt.% MC Ubtreated 12.5 wt.% MC 0 0 1 2 3 4 5 6 Displacement (mm)

Figure 4.29: Typical load-deflection curves with variation impact loading at different MCs during composites fabrication of 6.5% wt. MC and 12.5% wt. MC.

4.3.7.2 Peak Force, Contact Time and Absorbed Energy

For a better understanding of impact behaviour of flax/epoxy composites with respect of MC, the average values of peak force, contact time and absorbed energy are presented in Figure 4.30. Figure 4.30(a) shows the peak force increases with increasing impact loading because the composites sustained more damage. Although the 6.5 wt.% MC sample has higher peak force than 12.5 wt.% MC at 1.5 J corresponding to better interface properties, the difference is relatively small at 2.5

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J impact energy and then it became insignificant once the impact energy increased to 5.0 J. Figure 4.30(b) displays the contact time profiles. It can be seen that that the contact time between the impactor and the composites increased with increasing impact energy. Higher contact time is associated with an increase in the damage growth as a result of higher impact energy [239]. It is interestingly to observe that the untreated 12.5 wt.% MC samples shows higher contact time than the untreated 6.5 wt.% MC samples in both 1.5 J and 2,5 J impact energy. This indicates poor interfacial bonding causes higher contact time. At 5.0 J, however, the contact time remained similar for both samples which suggest that that damage resistance is similar for those composites.

Figure 4.30(c) shows the absorbed energy profiles. The absorbed energy is transferred from the impactor and dissipated by the composites in the event of impact loading [163]. The diagonal line in Figure 4.30(c) represents the Equal Energy Line corresponding to the ratio of impacted energy and absorbed energy. The plots shifted further below the Equal Energy Line as impact energy increases, indicating reduced energy recovery. In term of MC, the untreated 6.5 wt.% MC sample has higher absorbed energy than the untreated 12.5 wt.% MC but again to the different reduced above the difference reduced above 2.5 J impact loading. As described earlier, MC has the largest effect at 1.5 J but differences started to diminish at 2.5 J impact energy, becoming insignificant when 5.0 J impact energy is reached. The possible explanation is that as impact energy increases, more damage is created i.e. matrix cracking and fibre fracture that surpassed the degradation of fibre/matrix interface effect. However, it is unclear the reason at 5.0 J impact energy, there are no differences between untreated 6.5 wt.% MC and untreated 12.5 wt.% MC.

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(a)

1800

1600

1400 PeakForce (N)

1200 Untreated-6.5 wt.% MC Untreated-12.5 wt.% MC 1000 1 2 3 4 5 Impacted Energy (J)

(b) 8.0

7.5

7.0

6.5

ContactTime(ms) 6.0

Untreated-6.5 wt.% MC 5.5 Untreated-12.5 wt.% MC

1 2 3 4 5 Impacted Energy (J)

Equal Energy Line (Eabsorbed= Eimpacted)

2 AbsorbedEnergy (J)

1 Untreated-6.5 wt.% MC Untreated-12.5 wt.% MC 0 0 1 2 3 4 5 Impacted Energy (J)

Figure 4.30: Evolution of (a) peak force, (b) contact time and (c) absorbed energy as a function of impact energy.

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4.3.7.3 Observation of Damage in C-Scanning

Visual observation cannot describe the internal damage of the flax/epoxy composites upon low velocity impact. It is because damages such as delamination, matrix crack or debonding progress beneath the surface unnoticeable. This is called barely visible impact damage (BVID) [240][241] and therefore, ultrasonic C-Scan was employed. Figure 4.31 show the ultrasonic C-Scanning images of the untreated 6.5 wt.% MC and 12.5 wt.% MC samples. Internal damage represented by the dark areas is attributed to delamination [164]. Even at 1.5 J impact energy, damages had occurred, corresponding to the brittle nature of epoxy polymer and low strength of flax fibre. A greater area damage is visible on the 12.5 wt.% MC samples as compared to 6.5 wt.% MC samples. A cross-shaped pattern began to appear on the 12.5 wt.% MC samples. This feature, also reported in the literature [80][242][164], indicates that the delamination followed the warp and weft direction of the fabric plies [243]. As impact energy increases to 2.5 J, the dark area began to intensify particularly with 12.5 wt.% samples, indicating greater damages. For 6.5 wt.% MC, the C-Scan images showed the crack began to grow longer in the warp and weft direction as compared to 1.5 J impact energy. When the composited were subjected to 5.0 J impact energy, the C-Scan images become more intensified and bigger damage area is detected.

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Impact Energy

Samples 1.5J 2.5J 5.0J

6.5 wt.% MC Untreated 30 mm 30 mm 30 mm

12.5 MC wt.%

Untreated 30 mm 30 mm 30 mm

[AHA23]

Figure 4.31: Ultrasonic C-Scan of 6.5 wt.% MC and 12.5 wt.% MC impacted specimens.

4.3.8 Compression After Impact (CAI)

The average CAI strength as a function of impact energy for composites is plotted in Figure 4.32(a). All samples showed the same trend where CAI strength reduced significantly after being impacted at 1.5 J, reflecting the weakening of composites structures. It is reported the decreasing CAI strength is accompanied by progressing delamination [74][162][233][238]. However in this investigation, the CAI strength trend remains quite consistent at 1.5 J. 2.5 J and 5.0 J impact energy. This suggests the size of the delaminated area was not critical enough to affect the CAI strength. As shown in Figure 4.31, the delaminated area was small and confined along the fibre yarn direction. Liang et al. [239] also found the CAI strength of

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flax/epoxy composites remained consistent up to 6 J. Similarly, Bensadoun et al. [164] showed that delamination in woven and cross laminates of flax/epoxy composites hardly grew although the samples were intentionally perforated (at 12.4 J and 18.7 J) in the low velocity impact test.

Figure 4.32(b) shows the normalized CAI strength as a function of the incident impact energy. This figure gives a better comparison of the effect of impact damage on compression strength of the composites. It is the ratio between the average CAI strength of impacted samples and of the strength of the unimpacted samples. The normalized CAI strength indicates the reduction of composites was merely 8-10%. This can be explained by the non-severity of the pre-existing damages. It is observed that the untreated 12.5 wt.% MC samples shows the greatest reduction of normalized CAI strength for all given impact energy. This means the untreated 12.5 wt.% MC has greater stress concentration at the damage zone due to weak interfacial properties, creating further delamination during compression test and lowering the compression strength.

(a) (b) 68.9 Untreated 6.5 wt.% MC 1.02 70 Untreated 12.5 wt.% MC Untreated 6.5 wt.% MC Untreated 12.5 wt.% MC 0.99 64.0 63.4 65 62.4 0.96 59.3 60 0.93 54.7 53.5 53.4

CAI Strength CAI (MPa) 55 0.90 NormalisedStrength CAI

50 0.87

0 1 2 3 4 5 6 0 1 2 3 4 5 Impact Energy (J) Impact Energy (J)

Figure 4.32: The (a) CAI strength and (b) normalised CAI strength as a function of the incident impact energy.

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4.4 Mechanism of Moisture Interaction at the Interface during Flax/Epoxy Composites Manufacturing

Based on the previous results, it can be said moisture could enter into the hydrogen bond formation with hydroxyl group on surface of flax fibre [245] which otherwise would have contributed to the lower fibre-matrix interactions. It is also known that MC has an effect to the polymerization of epoxy resin during cure [19][91] [93]. The same outcome could occur with the presence of flax fibre where the moist flax fibre could affect the polymerization of the epoxy at the interface during cure. Hence, moisture interactions at the fibre-resin interface during flax/epoxy composites manufacturing can be proposed in 2 stages.

1- Fibre Wetting Stages

In the uncured epoxy-flax fibre lay-up when the resin infusion process is completed, few epoxide chains and amine hardener molecules will have reacted before the curing process is process is started. Therefore, water molecules at fibre surface might be absorbed to the ample hydrogen bonding sites of the hygroscopic amine hardener [86] as shown Figure 4.33(a).

2-Curing Stages

Curing of an epoxy resin is an exothermic reaction where heat is generated during the polymerisation process. In the present study, curing was initiated in a curing oven and where the temperature was raised gradually over and hour over an hour up to 80 °C from room temperature (≈1 °C / minute). Water molecule is expected to migrate from the core to the surface of flax fibres at a greater diffusion rate as shown schematically in Figure 4.33(b) [246]. Hence, the fibre/matrix interface is saturated with water molecules. However, water-epoxy solubility has its limitation and any excess of water molecules will not fit any longer in the free volume available in the epoxy resin. Therefore, the epoxy curing reaction may not properly taking part, creating a nanophase separated region at the fibre-matrix interface [86]. It is suggested that due to this excess water present at the interface during curing, polymeric chain mobility becomes less restricted because of

138

disruption to the crosslinking formation between reactive groups on the surface of the fibre and the epoxy matrix, leading to a reduction in chemical bonding [119].

When curing is completed in the composites, this means the moisturized fibre had possibly plasticised the epoxy matrix at the interface. A schematic illustrating this is shown in Figure 4.33(c). Excess water might also diffuse beyond the fibre/matrix interface region and moisture absorption through the bulk polymer is possible as illustrated in Figure 4.33(b).

Figure 4.33: Moisture interaction during composites fabrication during (a) wetting stages (b) curing stages. (c) Weak chemical bonding due to excess of MC on the fibre surface.

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4.5 Characterization of Chemically Treated Flax Fibre

4.5.1 Moisture Content of Chemically Treated Flax Fibre

NaOH Treated Flax Fibre

Figure 4.34 show the fibre MC as a function of RH for 3.0% NaOH and 4.5% NaOH treated fibre during absorption modes. In general, sigmoid type II curves were observed for NaOH treated fibre (3.0% and 4.5%) samples similar to that of untreated flax. In term of MC, MC of 3.0% NaOH and 4.5% NaOH samples reduced when compared with untreated flax fibre. However, the MC at highest humidity (90% RH) was similar. The reduction of MC after NaOH treatment indicates that that amounts of hemicellulose or/and lignin were successfully removed from the fibres, causing less hydroxyl (OH) groups in the flax fibres and ultimately, making the flax fibres less hydrophilic to moisture uptake.

It is interesting to observe that MC capacity was not consistent for each RH between the untreated and NaOH treated. In Figure 4.34(a) and Figure 4.34(b), the gap between the untreated and NaOH treated fibres curves began to move apart with increasing RH, indicating different moisture absorption capacity but this effect was limited to up to 60% RH. The untreated and NaOH treated curves began to shift closer again with higher humidity (≥70% RH). Therefore, it is suggests that the effectiveness of NaOH treatment was limited between 10% and 60% RH in the monolayer water formation.

In the polylayer formation at high humidity (≥70% RH), moisture absorption was no longer associated with the -OH group concentration (due to removal of hemicellulose), but rather closely to capillary condensation [58]. This is possibly why MC at 90% RH for untreated and NaOH treated fibres was relatively the same. In term of NaOH concentrations, however, the MC of 4.5% NaOH was slightly reduced over 3.0% NaOH samples. For instances at 50% RH, MC of 4.5% NaOH samples was reduced from 5.7 wt.% MC to 5.5 wt.% MC, representing merely a 3.5% reduction. The data suggests that most of hemicellulose or/and lignin constituents were removed at 3.0% NaOH concentration.

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(a) 18 16 Untreated - Absorption 14.8 3.0% NaOH - Absorption 14 15.0 11.3 12 9.2 10 8.0 11.0 8 6.8 8.6 MC (wt.%) MC 5.6 6 4.6 6.9 3.5 5.7 4 2.2 4.8 3.9 0.02 2.9 1.7 0 00.0 10 20 30 40 50 60 70 80 90 100 RH (%)

(b) 18

16 Untreated - Absorption 14.8 4.5% NaOH - Absorption 14 14.6 12 11.3

10 9.2 10.7 8.0 8 MC (%) MC 6.8 8.3 6 5.6 4.6 6.7 3.5 5.5 4 4.7 2.2 3.8 2 2.8 0.0 1.7 0 00.0 10 20 30 40 50 60 70 80 90 100 RH (%)

Figure 4.34: MC of (a) 3.0% NaOH and 4.5% NaOH treated fibre as a function of RH.

1.0% Silane Treated Fibre

Figure 4.35 shows MC comparison as a function of RH between untreated and 1.0% silane treated fibres in absorption mode. The sigmoid type II curves were again observed (as for untreated flax) and after 1.0% silane treatment. Except for 80% and 90% RH where the MC was considered similar, that MC were decreased although silanization generally does not involves removal of removal of any fibre’s constituents such as hemicellulose or pectin. These results are consistent with Chen and co-workers [123] who found there was no reduction of MC at higher humidity (80% RH) on a silanized bamboo strip.

141

Similar to the NaOH treated fibre, MC capacity was different for each RH where the 1.0% silane curves began to shift away from 10% RH and then getting closer above 70% RH with respect to untreated flax, indicating polylayer water formation in the 1.0% silane treated fibre. Moisture reduction after silanization can be explained by the interaction of silane with natural fibre. Firstly, silane serves as a water repellent (hydrophobic) layers layer on the surface of cell walls by creating hydrogen and/or covalent bonding between silanol (Si-OH) and the hydroxyl groups of the fibre. This would reduce water absorption to the hygroscopic hydroxyl site of fibres after silane treatments. Secondly, due to bulking silane treatment (fibres were immersed in in silane/acetone solutions) the side effect could be to be to possibly reduce the size of nanopores with access for water molecules being restricted and silane treatment could also deactivate or mask the hydroxyl functionalities thereby water absorption could be reduced [133].

16 Untreated - Absorption 14.8 1.0% Silanesilane - Absorption 14 15.1 12 11.3

10 9.2 11.1 7.9 8 6.8 8.5

5.6 MC (wt.%) MC 6 4.5 6.7 3.5 5.6 4 4.8 2.2 4.1 2 3.2 0.0 2.0 0 0.00 10 20 30 40 50 60 70 80 90 100 RH (%)

Figure 4.35: MC of 1.0% silane treated fibre as function of RH.

Combined 3.0% NaOH & 1.0% Silane Treated Fibre

Figure 4.36 shows MC comparison across RH between untreated and 3.0% NaOH & 1.0% silane treated fibres in absorption mode. Sigmoid type II curves were observed before (untreated flax) and after the NaOH-silane treatment. The curves

142

indicate the MC reduced with the increasing RH similar to that for the individually treated NaOH and silane treated fibres.

18 Untreated - Absorption 16 3.0% NaOHNaoH & 1.0% Silanesilane - Absorption 14.8 14 14.3 12 11.3 10 9.2 10.9 7.9 8 6.8 8.8 5.6 MC (wt.%) MC 6 7.4 4.5 6.3 3.5 4 5.3 2.2 4.3 2 3.3 0.0 2.1 0 00.0 10 20 30 40 50 60 70 80 90 100 RH (%)

Figure 4.36: MC of 3.0% NaOH & 1.0% silane treated fibres as a function of RH.

Comparison between the Chemically Treated Fibres

Figure 4.37 shows a comparison between all the chemical treatments for different humidity level (10%, 50% and 90% RH). Except at 90% RH, it can be seen that all the chemical treatments caused a modest MC reduction of the flax fibres at 10% and 50% RH. Among the surface treatments, 3.0% NaOH and 1.0% silane were the best in reducing MC in flax fibre. In contrast, the silane on alkalized fibres (3.0% NaOH & 1.0% silane) had the least MC which is unexpected based on the standalone surface treatments (NaOH and silane). [AHA24]

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Untreated 15.1 16 14.8 3.0% NaOH 15.01 14.3 14 1.0% silaneSilane 3.0% NaOHNaOH-1.0% & 1.0% Silane silane 12

8 6.8

MC (wt.%) MC 6.3 6 5.7 5.6

2.2 2.0 2.1 2 1.7

0 10 50 90 RH (%) [AHA25] Figure 4.37: Effect of chemical treatments on the MC of flax fibres.

4.5.2 Fourier Transform Infrared (FTIR) Analysis

Untreated Flax Fibress

The main constituents of flax fibres are cellulose (~78%), hemicellulose (~20.6%), lignin (~2%), pectin (~1.7%), moisture and some trace waxes. Cellulose can be identified by several bands in FTIR; the -OH stretching vibration of hydrogen bonded between ~3600-2995 cm-1 (where the water molecule attracted to the hydrogen bond, making them hydrophilic), the -CH ~2906 cm-1 & ~2844 cm-1 , -OH -1 -1 -1 bending ~1575 cm , -CH2 bending and CH bending at ~1409 cm and 1377 cm respectively [135][171][241]. The band at ~1735 cm-1 corresponds to the carbonyl stretching bands of the acetyl group in hemicellulose [141]. Lignin can be detected with C–O stretching of the acetyl groups at ~1595 cm-1 respectively. The peak ~2917 cm-1 and ~2850 cm-1 is a unique identification of flax fibres as compared to other type of natural fibres that is attributed to the vibrational C-H stretching and

CH2 symmetrical stretching [135][194][242]. These bands are associated to the impurities and wax composition [141]. The band at 1635 cm-1 in natural fibre is assigned to O-H bonding which belongs to absorbed water. The higher the peak, the more water is absorbed [249].

144

NaOH Treated Fibres

As shown in Figure 4.38(a), the effect of NaOH solution concentrations, 1.5%, 3.0% and 4.5% in flax fibre treatment is not conclusive due to the similarities of the NaOH treated spectra. However, the differences between NaOH treated fibres and untreated flax fibre spectra were noticeable. Consistent with the literature [147][177], the NaOH treatment gives partial disappearance of the 1735 cm-1 band, related to the presence of hemicellulose, as the as the hemicellulose dissolves in the NaOH solution. In Figure 4.38(b), a minor peak at 1595 cm-1 was observed for the untreated fibres. The peak, however, disappeared on the NaOH treated fibres, indicating that lignin was removed from the fibre surface [141].

(a) 0.25 Untreated 1.5% NaOH 1735 cm-1 3.0% NaOH 0.090 Hemicellulose 4.5% NaOH 0.20 0.085

0.080

0.15 0.075

0.070

1800 1780 1760 1740 1720 1700 1680 Absorbance(%) 0.10

0.05 4000 3500 3000 2500 2000 1500 1000 500 Wavenumber (cm-1)

(b) 0.100 1595 cm-1 C-O stretching in lignin

0.095

0.090 Absorbance(%)

0.085

0.080 1650 1600 1550 Wavenumber (cm-1)

Figure 4.38: FTIR-ATR spectra of NaOH treated fibres with different NaOH concentration that shows (a) hemicellulose and (b) lignin removal.

145

Silane Treated Fibres

As shown in Figure 4.39, the effect of increased silane solutions (1.0%, 1.5% and 2.0%) in flax fibre treatment was not pronounced in the spectrums. The possible explanation is that silane treatment in the present investigation was done for the same duration (2 hours) and temperature (room temperature) [250]. However, the changes between the untreated and silane treated flax fibre can be distinguished clearly. The 3-aminopropyltriemethoxysilane (APTMS) is detected in the spectra at 1000 cm-1 and 982 cm-1 as peak intensity decreased after the silane treatment, which implies the presence of Si-O-C bonds [27].

The most significant changes before and after silane treatment are that the silane treatment appears to disrupt the 1636 cm-1 band which holds the water [249], suggesting lower moisture absorption. The hydroxyl group peak intensity in the 3600-2995 cm-1 band of silane treated fibres was weaker as compared to the untreated fibres [249]. This argument is corroborated by the reduced MC of the 1.0% silane treated fibres measured using the DVS analyser as compared to the untreated fibres. The newly emerged peak at 1558 cm-1 can be explained by the presence of amine groups to form hydrogen bonds between silanol group (Si–O–H) and flax fibre [133][134]. This probably can be explained by hydrolysis of APTMS in the water/acetone mixture that successfully absorbed and reacts chemically to the hydroxyls group on the flax fibre.

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0.30 Untreated 1.0% Silane 1000 cm-1 & 982 cm-1 1.5% Silane Si-O-C -1 0.25 2.0% Silane 1636 cm Absorbed water

0.085 1558 cm-1 0.20 0.080 Si-O-H 0.075

0.070 -1 0.15 3600-2900 cm0.065

Absorbance(%) Hydroxyl group 1700 1650 1600 1550

0.10

0.05 4000 3500 3000 2500 2000 1500 1000 500 Wavenumber (cm-1)

Figure 4.39: FTIR-ATR spectra of silane treated fibres with different silane concentrations.

3.0% NaOH & 1.0% Silane Treated Fibres

Figure 4.40 shows the effect of the combination of 3.0% NaOH & 1.0% silane treated fibre compared to the untreated flax fibres. The same decreased intensity was found at 1000 cm-1 and 982 cm-1, which indicates Si-O-C bonds were present. Similarly, the peak at 1558 cm-1 indicates hydrogen bonding between silanol groups and flax fibres.

The combination of NaOH and silane treatment resulted in two important characteristics in the spectra. The 1735 cm-1 band disappeared, which suggests that hemicellulose was almost completely removed from the fibre. It also disrupts the 1636 cm-1 band that corresponds to the absorbed water. This is followed by a weak band at 3600-2995 cm-1 for hydroxyl groups as compared to the untreated fibres. The most interesting aspect of these spectra that could not be found for the standalone NaOH (Figure 4.38) and silane treatments (Figure 4.39) is the reduction in the peaks at 2917 cm-1and 2850 cm-1 that possibly relate to the removal of wax and pectin [251][247]. This means the combination of alkaline and silane treatment has successfully modified the fibre surface by removing pectin, wax and hemicellulose but also to reduce hydrophobicity.

147

0.25 Untreated 3.0% NaOH & 1.0% silaneSilane

0.100 2917 & 2850 cm-1 Wax & Impurities 0.20 0.095 0.090

0.085

0.080

0.075 -1 0.15 1636 cm 3000 2950 2900 2850 2800 2750Absorbed water 3600-2900 cm-1 Hydroxyl group -1 Absorbance(%) 1735 cm 0.10 Hemicellulose

0.05 4000 3500 3000 2500 2000 1500 1000 500 Wavenumber (cm-1)

Figure 4.40: FTIR-ATR spectra of combined 3.0% NaOH & 1.0% silane treated fibres.

4.5.3 Changes of Fibre Density

Figure 4.41 presents fibre density with respect to MC for untreated and treated flax samples. In general, the fibre density of treated flax fibres had a similar trend with the untreated flax fibre samples where they decreased in density with increasing MC. It is observed that the treated fibres had higher fibre density across the measured MC. Increasing fibre density after chemical treatment was reported in the literature [252][253][248] where mostly the fibre density measurement was done at ambient RH, presumably at 50% RH. In general, the increased fibre density after chemical treatments was attributed to cell wall densification, signifying permanent changes to the cellular structure [254]. In the 3.0% NaOH sample, the increased of fibre density is due to removal of the less dense non-cellulosic constituents constituents/impurities (hemicellulose, fats, waxes) [252][255] and this is likely to create voids in the inner structure of the fibres. These voids can be filled by the low viscosity Rapeseed oil. As cellulose fraction was increased, the measured density for NaOH treated fibre is higher than the untreated fibre. Another possible explanation is that the increased density after NaOH treatment was possibly due to increased crystallinity in flax fibre [256]. Meanwhile, the densification of silane treated fibre can be attributed to the filling of pores by grafted silane molecules [252]. It is interesting to observe that the combination of 3.0% NaOH & 1.0% silane treatment

148

gave similar fibre density to 1.0% silane at dried conditions but then higher than 3.0% NaOH when measured at medium and high MC.

1.52 Untreated 3.0% NaOH 1.0% Silanesilane 1.50

3.0% NaOH & 1.0% Silanesilane

) 3 1.48

1.46

1.44 FibreDensity (g/cm 1.42

1.40 0 2 4 6 8 10 12 14 MC (wt.%)

Figure 4.41: Fibre density as function of MC after chemical treatments.

Once again, using Equation 2.2, the experimental data of chemical treated fibres was compared with Messiry’s model; plotted against MC as shown in Figure 4.42(a-c). Figure 4.42 (a) and Figure 4.42 (b) indicate high applicability of the model for 3.0% NaOH and 1.0% silane treated samples. In contrast, the model data for model data for the 3.0% NaOH & 1.0% silane treated samples (Figure 4.42(c)) did not fit with the experimental data because the model deviates severely as MC increases.

149

(a) 1.52 (b) 1.50

1.50 1.48

)

) 3 1.48 2 R = 0.9713 1.46 R² = 0.9214

1.46 1.44

1.44 FibreDensity (g/cm FibreDensity (g/cm 1.42 1.42 3.0% NaOH (Experimental Data) 1.0% silane (Experimental Data) 3.0% NaOH (Messiry's Proposed Model) 1.0% silane (Messiry's Proposed Model) 1.40 1.40 0 2 4 6 8 10 12 14 0 2 4 6 8 10 12 14 MC (wt.%) MC (wt.%)

1.48

) 2 3 R = 0.8307

1.46

1.44 FibreDensity (g/cm 1.42

3.0% NaOH & 1.0% silane (Experimental Data) 3.0% NaOH & 1.0% silane (Messiry's Proposed Model) 1.40 0 2 4 6 8 10 12 14 MC (wt.%)

Figure 4.42: Comparison between current experimental data and Messiry’s model for (a) 3.0% NaOH, (b) 1.0% silane and (c) 3.0% NaOH & 1.0% silane treated fibre. Dashed plot is a linear-regression of experimental data.

4.5.4 Single Fibre Tensile Test

The tensile properties of elementary flax fibre were evaluated by the single fibre tensile test (SFTT). The SFTT was done to determine the changes of mechanical properties of flax fibres after chemical treatments.

4.5.4.1 Stress-Strain Behaviour of Elementary Flax Fibres

Figure 4.43(a-c) shows typical stress-strain behaviour of untreated and treated flax fibres. The curves of the untreated sample show two distinct regions. The first is non-linear where sudden changes in in the slope in the curves occur and a pronounced of ‘knee’ is noticeable about 0.25%-0.5% strain [257]. This can be

150

explained by the reorientation of microfibrils inside a S2 layer or/and sliding of microfibrils in a cell wall caused by cell wall deformation. When the applied load increases, microfibrils that are oriented at 10° from the fibre axis are forced parallel with the fibre axis, causing rearrangement of the core surrounding the amorphous matrix (lignin, pectin and hemicelluloses) [53][103][128][251]. The second region covers the final half of the curves appear to fit a linear line that represents the true elastic properties of the fibre. Hence, the longitudinal Young’s modulus was determined from the slope of this region, assuming aligned microfibrils cellulose was the only material to deform in the fibre axis [257]. Once they reached their peak, the curves dropped abruptly, signifying the elementary fibres fractured in a brittle manner.

(b) 500 (a) 500

400 400

Knee Knee 300 300

200 200

Stress(MPa) Stress(MPa)

Untreated Untreated 100 1.5% NaOH 100 1.0% Silane 3.0% NaOH 1.5% Silane 4.5% NaOH 2.0% Silane 0 0 0.0 0.5 1.0 1.5 2.0 2.5 0.0 0.5 1.0 1.5 2.0 2.5 Strain (%) Strain (%)

400

300 Knee

200 Stress(MPa)

100 Unreated 3.0% NaOH &1.0% Silane 0 0.0 0.5 1.0 1.5 2.0 2.5 Strain (%)

Figure 4.43: Typical stress-strain curves of elementary fibres after (a) NaOH, (b) silane and (c) 3.0% NaOH & 1.0% silane treatments. The dash line is purposely added, showing the non-linear behaviour of the curves.

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The stress-strain curves of NaOH treated flax fibres were somewhat at different from the untreated flax. The ‘knee’ was still visible in the curves but almost disappears as compared to the untreated flax samples, making the entire curves almost a straight line. In Figure 4.43(a), there were no significant changes between the curves when NaOH concentrations were increased. The result is consistent with Raj et al. [79] for their for their experimental work on the single fibre tensile test of elementary flax fibre. Their result results showed the almost disappearance of the non-linear part of their stress-strain strain curves for the flax fibres treated with 1.0% and 3.0% NaOH treatment. However, However 5.0% NaOH and 10.0% NaOH treated flax fibres produced biphasic stress-strain curves, behaviour attributed to the severity of the NaOH attack on the hemicellulose and pectin macromolecules in the secondary layer.

In the case of silane-treated fibres, the shape of the stress-strain curves also different from the untreated flax samples as shown in Figure 4.43(b). These curves still had both non-linear and linear regions but the ‘knee’ was less pronounced, similar to the NaOH-treated samples. It is observed that increasing silane concentration did not change the behaviour of the curves. Figure 4.43(c) shows the stress-strain curves of 3.0% NaOH and 1.0% silane treated flax samples. These curves also had different behaviour to the untreated samples but had similar shapes to the NaOH and silane treated samples with an almost disappearance of the ‘knee’.

4.5.4.2 Tensile Properties of Elementary Flax Fibres and Weibull Analysis

As shown in Figure 4.43(a-c), the flax fibres were brittle and this satisfying the weakest link theory (fibre flaws and defect) according to the Weibull distribution [258]. Anderson and co-workers [109] suggested that the tensile properties of elementary fibres is well approximated by the two parameter Weibull distribution because of the relatively high scattered data. The analysis of tensile strength using the two two-perimeter Weibull distribution of untreated and treated elementary flax fibres is tabulated in Table 4.11 for of average tensile strength, Young’s modulus and tensile strain at break.

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Figure 4.44(a-c) shows the two-perimeter Weibull distribution for tensile strength data of treated flax fibre samples. It can be seen that at all the experimental tensile strength data has a good fitting with the two-perimeter Weibull distribution. It is interesting to see all the treated curves were shifted to the left of the untreated samples, suggesting that the tensile strength of flax fibre reduced after the chemical treatments.

(a) (b) 1.0 1.0

0.8 0.8

0.6 0.6

0.4 0.4

Untreated Failure of Probability

Probability of Failure of Probability Untreated 0.2 1.5% NaOH 0.2 1.0% Silane 3.0% NaOH 1.5% Silane 4.5% NaOH 2.0% Silane 0.0 0.0 0 100 200 300 400 500 600 700 800 900 1000 0 100 200 300 400 500 600 700 800 900 1000 Tensile Strength (MPa) Tensile Strength (MPa)

0.8

0.6

0.4 Probability of Failure of Probability 0.2 3.0% NaOH-1.0% Silane Untreated 0.0 0 100 200 300 400 500 600 700 800 900 1000 Tensile Strength (MPa)

Figure 4.44: Two-parameter Weibull distribution fitting on the tensile strength data of flax fibres after (a) NaOH (b) silane and (c) silane on alkalized treated fibre.

Table 4.11 shows the Weibull characteristic strength had a similar trend to the average tensile strength which is expected. The calculated Weibull moduli are between 2.3 and 3.9, which are in the same range in comparison with other natural fibres in the literature. These authors reported the Weibull modulus of cellulosic fibres were between 1.4 to 6.64 [46][145][180][259]. It is interesting to observe that untreated samples had the highest Weibull modulus (3.9) than the treated fibres (2.3-

153

3.1). Higher Weibull modulus indicates there was a lower variability of the failure strength. The lower Weibull modulus obtained for the treated samples may be explained by the fact that those fibres more likely contained a greater number of defects along the fibre length [260].

Tensile Properties of Untreated Fibre

The characteristic strength and average Young’s modulus of untreated flax fibres were 541.61 MPa and 40.75 GPa, respectively, which are within the range for flax fibres reported in the literature of 343-1035 MPa and 28-100 GPa [45]. The variability in the tensile properties of flax fibres is possibly due to the secondary processing parameters (i.e extraction, hackling, spinning) prior to being woven into fabric and hence, damaging the fibres [29]. From Table 4.11, it is found that all chemical treatments; NaOH, silane and silane on alkalized fibres decreased the tensile strength and Young’s modulus of flax fibres. However, the treated samples showed greater deformation before break.

Tensile Properties of NaOH Treated Fibre

The reduction of tensile properties after NaOH treatment possibly can be can be explained by looking to the overall constituents in the flax fibre. In general, NaOH treatment promotes surface roughness and mechanical interlocking [261]. However, NaOH treatments may has taken out some portion of the hemicelluloses and lignin that hold and secure the cellulose microfibrils in place [118][167][247]. This argument is consistent in our FTIR-ATR analysis where hemicellulose was partly removed after the NaOH treatment. The removal of the non-cellulose constituents could create voids in the fibre and hence result in lower tensile properties [253]. Among the NaOH concentrations, the 4.5% NaOH samples had the lowest tensile strength compared to 1.5% NaOH and 3.0% NaOH samples. The decreased of tensile strength with increasing NaOH concentration was similarly reported by Aly et al. [173] and Raj et al. [79] for elementary flax fibres.

Tensile Properties of Silane Treated Fibre

In case of silane treatments, the decrease of strength was in descending order with increasing silane concentration; 1.0% silane (502.11 MPa) > 1.5% silane

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(483.79 MPa) > 2.0% silane (445.89 MPa). This result suggests that that furthering silane treatment above 2.0% may not be beneficial since the 2.0% silane treated samples lost around 18% of its tensile strength compared to the untreated fibre, and hence interfacial shear strength measurement of this sample was not conducted. Silane treated fibres were reported to have lower tensile strength as compared to untreated fibres [117][252][253][261]. Although silane generally does not induce damage to the fibre, the treating system could consists of fibre-damaging elements such as an acidic medium or high temperature [133]. In the present investigation, the reduced tensile strength and Young’s modulus of flax fibres after being silane treated was likely due to the fact that the flax was hydrolysed in an alkaline medium (pH 9) and hence, damaging the fibres.

Tensile Properties of NaOH & Silane Treated Fibre

The strength of combined alkaline and silane (3.0% NaOH and 1.0% silane) treatment was 512.84 MPa. Interestingly, this values was somewhere in between those of in between those of the standalone treatments of 3.0% NaOH and 1.0% silane, at 521.03 MPa and 502.11 MPa respectively. The possible explanation is that fibres were initially damaged during the first NaOH (pH 14) treatment and subsequent silane treatment may not affect the fibre properties as the alkaline medium used for hydrolysis had lesser pH (pH 9). A possible explanation of the tensile strength can be related to the study done by Sawpan et al. [252]. Using X-ray diffraction (XRD), the authors suggested there was a positive relationship between composite tensile strength and the cellulose crystallinity index of untreated and treated hemp/PLA composites. They measured a decrease of cellulose crystallinity after alkaline & silane treatment is responsible for the reduced tensile strength.

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Table 4.11: Weibull parameters and experimental tensile properties of untreated and treated flax fibre.

Weibull Parameters Samples Average Tensile Strength Average Young’s Tensile Strain at Characteristic Weibull (MPa) Modulus (GPa) Break (%) Strength, σ0 (MPa) Modulus Untreated 541.61 3.9 490.09 ± 141.01 40.8 ± 8.06 1.23 ± 0.42

NaOH Treated Fibre

1.5% NaOH 510.85 2.7 459.70 ± 172.60 30.9 ± 9.14 1.58 ± 0.57

3.0% NaOH 521.03 2.4 461.70 ± 206.94 32.0 ± 10.80 1.56 ± 0.57

4.5% NaOH 468.77 2.9 417.23 ± 157.23 32.1 ± 14.55 1.70 ± 0.57

Silane Treated Fibre

1.0% silane 502.11 2.5 444.46 ± 191.64 33.1 ± 11.64 1.49 ± 0.53

1.5% silane 483.79 2.7 437.60 ± 160.52 32.1 ± 11.55 1.63 ± 0.49

2.0% silane 445.89 2.3 394.82 ± 191.58 32.6 ± 10.66 1.52 ± 0.57

NaOH & Silane Treated Fibre

3.0% NaOH & 1.0% silane 512.84 3.1 456.40 ± 161.35 28.7 ± 11.64 1.81 ± 0.51

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4.5.5 Physical Characteristic of Chemically Treated Flax Fabric

Figure 4.45 showed the influence of the chemical treatments on the % crimp and fabric areal density. In Figure 4.45(a), the data shows that the yarn % crimp is greater for the chemically treated flax samples as compared to the untreated samples. The NaOH samples had the highest % crimp followed by the silane and combination of NaOH and silane treatment. The increase in % crimp was likely due to the swelling of the cell wall of the flax fibres after the NaOH and silane treatments which is similarly reported in the literature [15][173][262][177]. As a result, the swelled fibre increased the fibre areal density (measured at dried condition) which is shown in Figure 4.45(b). All the NaOH treated fabrics had fibre areal density values around 437 g/cm2, around 12% higher than higher than the untreated samples while the 1.0% silane treated samples was the lowest at 422.2 g/cm2. The combination of NaOH-silane samples had similar values with 3.0% NaOH standalone treatment at 437.6 g/cm2.

A photograph before and after 3.0% NaOH treatment of fibre yarn is shown in Figure 4.46(a). It can be seen that the swelled 3.0% NaOH treated yarn had a larger diameter and greater waviness in the yarns structure as compared to the untreated sample. As a result, the thickness of NaOH treated fabric had increased considerably, causing shrinkage predominantly in the warp direction as shown in Figure 4.46(b). It is reported that the waviness structure in fabric was found to have a detrimental effect on the toughness and strength of the composites [196][263][264]. The fibre waviness is suggested to induce non-linear behaviour in continuous-fibre reinforced composites where it produced a matrix dominated failure as a resulted of the shifted fibre orientation in the primary loading direction [263].

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(a) (b) 500 16 15.0 14.3

) 436.2 438.0

13.1 -2 437.6 437.6 14 12.7 422.2 426.1 12.1 12.9 12 400 387.4

10 8.5 8

Crimp (%) Crimp 6 300 4

2 FibreDensity Areal (gcm 0 200

Untreated Untreated 1.5% NaOH 3.0% NaOH 4.5% NaOH 1.0% Silane 1.5% Silane 1.5% NaOH 3.0% NaOH 4.5% NaOH 1.0% Silane 1.5% SIlane 3.0% NaOH & 3.0% NaOH & 1.0% Silane 1.0% Silane

Figure 4.45: Influence of the chemical treatments to the (a) % crimp and (b) fibre areal density of flax fabric.

Figure 4.46: Photographs of (a) fibre yarn and (b) flax fabric before and after 3.0% NaOH treatment.

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4.6 Interfacial Shear Strength (IFSS) of Treated Flax/Epoxy Composites

4.6.1 Microbond Force-Displacement Curves

Figure 4.47(a-c), Figure 4.47(d-e) and Figure 4.47(f) show typical force- displacement curves of the NaOH (1,5% NaOH, 3.0% NaOH and 4.5% NaOH), silane (1.0% silane and 1.5% silane) and combination of 3.0% NaOH- 1.0% silane treated flax/epoxy composites prepared at different MC. In general, they show similar behaviour in the curves to the untreated sample where the applied load increased with increasing displacement and fail abruptly after the peak, indicating fast interfacial cracking of the interface [191]. From this point, the pull-out process begins with variation in the force that is due to the frictional resistance between the droplet and fibre surfaces. The linear portion of the curve also describes the stiffness of the longitudinal fibre. The loss of stiffness of the treated fibre is seen clearly when the sample was produced at high MC. The loss of stiffness is probably related to the irreversible softening effect of flax fibre with MC [121] and chemical treatments could cause more damage to the elementary flax fibre. It can be seen that at higher MC, the treated samples failed at lower debonding force which is consistent with behaviour reported in the literature [18][119] and hence, lower IFSS values were expected. The IFSS data is presented and discussed in the next section.

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(a) 0.20 (b) 0.20

0.15 0.15

0.10 0.10

Force(N) Force(N)

0.05 0.05

0.00 0.00 0.0 0.1 0.2 0.3 0.4 0.5 0.0 0.1 0.2 0.3 0.4 0.5 Displacement (mm) Displacement (mm) Untreated-2.4 wt.% MC 1.5% NaOH-2.3 wt.% MC Untreated-2.4 wt.% MC 3.0% NaOH-2.4 wt.% MC Untreated-6.1 wt.% MC 1.5% NaOH-6.5 wt.% MC Untreated-6.1 wt.% MC 3.0% NaOH-6.7 wt.% MC Untreated-12.7 wt.% MC 1.5% NaOH-12.7 wt.% MC Untreated-12.7 wt.% MC 3.0% NaOH-12.9 wt.% MC

0.15 0.15

0.10 0.10

Force(N) Force (N) Force 0.05 0.05

0.00 0.00 0.0 0.1 0.2 0.3 0.4 0.5 0.0 0.1 0.2 0.3 0.4 0.5 Displacement (mm) Displacement (mm) Untreated-2.4 wt.% MC 4.5% NaOH-2.5 wt.% MC Untreated-2.4 wt.% MC 1.0% silane-2.6 wt.% MC Untreated-6.1 wt.% MC 4.5% NaOH-6.5 wt.% MC Untreated-6.1 wt.% MC 1.0% silane-6.3 wt.% MC Untreated-12.7 wt.% MC 4.5% NaOH-13.1 wt.% MC Untreated-12.7 wt.% MC 1.0% silane-12.9 wt.% MC

(e) 0.20 (g) 0.20

0.15 0.15

0.10 0.10

Force(N) Force(N)

0.05 0.05

0.00 0.00 0.0 0.1 0.2 0.3 0.4 0.5 0.0 0.1 0.2 0.3 0.4 0.5 Displacement (mm) Displacement (mm) Untreated-2.4% MC 1.5% silane-2.4 MC Untreated-2.4 wt.% MC 3.0% NaOH & 1.0% silane-2.9 wt.% MC Untreated-6.1% MC 1.5% silane-6.4 MC Untreated-6.1 wt.% MC 3.0% NaOH & 1.0% silane-6.9 wt.% MC Untreated-12.7 wt.% MC 3.0% NaOH & 1.0% silane-12.9 wt.% MC Untreated-12.7% MC 1.5% silane-12.9 MC

Figure 4.47: Typical force-displacement curves of (a) 1.5% NaOH, (b) 3.0% NaOH and (c) 4.5% NaOH, (d) 1.0% silane, (e) 1.5% silane and combined NaOH-silane treated samples: (f) 3.0% NaOH & 1.0% silane.

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4.6.2 Interfacial Shear Strength (IFSS) and Weibull Analysis

The IFSS data of chemically treated samples was reported by means of a two-parameter Weibull distribution and average IFSS. The data is tabulated in Table 4.12. It can be seen that the Weibull characteristic IFSS has a similar trend to the average IFSS in that they become lower as MC increases. Depending on chemical modification, variation of the Weibull modulus of chemically treated samples was relatively large and varied between 3.3 and 7.7. This Weibull modulus is difficult to compare with any previous studies since the uses of Weibull distribution for microbond analysis on elementary fibres are rarely reported in the literature for natural fibres. However, in most cases, it can be seen that the Weibull modulus in Table 4.12 decreased with increasing MC which means the IFSS data were scattered greatly. Therefore, when the Weibull modulus is lower, the presence of inhomogeneous distributed flaws on weak fibre-matrix interfaces is proposed [178]. A Weibull probability plot for IFSS of treated flax fibres is presented in Figure 4.48, which shows the two-perimeter Weibull distribution had a good fit with all IFSS experimental data.

Interfacial failure can be described in two ways; cohesive failure of the matrix and adhesive failure due to fibre-matrix debonding at the interface. Cohesive failure indicates strong interfacial bond strength in the composite system. If the composite’s strength is limited by the shear strength of the matrix and also assuming the matrix is isotropic l, the matrix shear strength, 휏푚 is estimated using Von Mises criterion using [117][265]:-

휎푚 4.2 휏푚 = √3

where 휎푚 is the tensile strength of the matrix

Assuming there is no influence of fibre MC on the matrix tensile strength 휎푚 and taking 60 MPa as the strength of an epoxy [45], the calculated 휏푚 using Equation 4.2 is around 35 MPa, which suggests the maximum IFSS value. In can be seen that all the values in Table 4.12 are lower than 35 MPa, with the highest being 26 MPa (74% of the max.) and the lowest (14 MPa) being only 40% of the 161

maximum. This weak adhesion produces adhesive interfacial failure in the microbond test.

Table 4.12: Weibull parameters and average IFSS of treated flax fibre.

Weibull Parameters Main Characteristic Average Types Samples Weibull IFSS, σ IFSS (MPa) of 0 Modulus (MPa) Failure NaOH Treated Fibre

1.5% NaOH-2.3 wt.% MC 24.1 4.6 22.1 ± 5.3

1.5% NaOH-6.5 wt.% MC 19.1 4.9 17.5 ± 4.1

1.5% NaOH-12.7 wt.% MC 14.0 3.3 12.5 ± 4.3 Adhesive

3.0% NaOH-2.4 wt.% MC 26.0 5.1 24.1 ± 5.4

3.0% NaOH-6.7 wt.% MC 21.5 5.1 20.1 ± 4.3

3.0% NaOH-12.9 wt.% MC 18.2 3.8 17.1 ± 5.5 Adhesive

4.5% NaOH-2.6 wt.% MC 23.7 5.9 22.0 ± 4.3

4.5% NaOH-6.6% wt.% MC 20.5 7.1 19.3 ± 3.1

4.5% NaOH-13.1 wt.% MC 14.3 5.9 13.3 ± 2.6 Adhesive

Silane Treated Fibre

1.0% silane-2.5 wt.% MC 22.6 5.9 21.0 ± 3.9

1.0% silane-6.5 wt.% MC 22.3 5.9 20.7 ± 4.2

1.0% silane-12.7 wt.% MC 18.7 4.4 17.2 ± 4.6 Adhesive

1.5% silane-2.4 wt.% MC 24.3 6.1 22.6 ± 4.1

1.5% silane-6.6 wt.% MC 23.5 5.3 21.7 ± 4.9

1.5% silane-13.1 wt.% MC 19.1 4.1 17.3 ± 5.0 Adhesive

NaOH & Silane Treated Fibre

3.0% NaOH & 1.0% silane-2.4 wt.% MC 22.8 7.4 21.4 ± 3.4

3.0% NaOH & 1.0% silane-6.5 wt.% MC 18.8 7.4 17.6 ± 3.0

3.0% NaOH & 1.0% silane-13.0 wt.% MC 16.7 5.1 15.6 ± 3.4 Adhesive

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(a) (b) (c) 1.0 1.0 1.0

0.8 0.8 0.8

0.6 0.6 0.6

0.4 0.4 0.4 Untreated-2.6 wt.% MC Untreated-2.6 wt.% MC

Untreated-2.4% wt. MC Untreated-6.3 wt.% MC

Probability of Failure of Probability Probability of Failure of Probability Probability of Failure of Probability Untreated-6.3 wt.% MC Untreated-6.7% wt. MC Untreated-12.9 wt.% MC Untreated-12.9 wt.% MC 0.2 0.2 Untreated-12.9% wt. MC 0.2 1.5% NaOH-2.3 wt.% MC 4.5% NaOH-2.6 wt.% MC 3.0% NaOH-2.4 wt.% MC 4.5% NaOH-6.6 wt%. MC 1.5% NaOH-6.5 wt.% MC 3.0% NaOH-6.7 wt.% MC 4.5% NaOH-13.1 wt.% MC 1.5% NaOH-12.7 wt.% MC 3.0% NaOH-12.9% wt.% MC 0.0 0.0 0.0 0 5 10 15 20 25 30 35 40 45 0 5 10 15 20 25 30 35 40 45 0 5 10 15 20 25 30 35 40 45 IFSS (MPa) IFSS (MPa) IFSS (MPa)

(d) (e) (f) 1.0 1.0 1.0

0.8 0.8 0.8

0.6 0.6 0.6

0.4 Untreated-2.6 wt.% MC 0.4 0.4 Untreated-2.6 wt.% MC

Untreated-6.3 wt.% MC Untreated-2.6 wt.% MC

Probability of Failure of Probability Failure of Probability Probability of Failure of Probability Untreated-6.3 wt.% MC Untreated-12.9 wt.% MC Untreated-12.9 wt,% MC Untreated-6.3 wt.% MC 0.2 0.2 0.2 Untreated-12.9 wt.% MC 1.0% silane-2.6 wt.% MC 1.5% silane-2.6 wt.% MC 1.0% silane-6.5 wt.% MC 3.0% NaOH & 1.0% silane-2.4 wt.% MC 1.5% silane-6.3 wt.% MC 3.0% NaOH & 1.0% silane-6.5 wt.% MC 1.0% silane-13.1 wt.% MC 1.5% silane-13.1 wt.% MC 3.0% NaOH & 1.0% dilane-12.9 wt.% MC 0.0 0.0 0.0 0 5 10 15 20 25 30 35 40 45 0 5 10 15 20 25 30 35 40 45 5 10 15 20 25 30 35 40 45 IFSS (MPa) IFSS (MPa) IFSS (MPa)

Figure 4.48: Two-parameter Weibull distribution fitting on the IFSS data; (a-c) NaOH, (d-e) silane and (f) 3.0% NaOH &1.0% silane treated fibres.

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IFSS of NaOH and Silane Treated Fibres

Comparisons between IFSS data for untreated and treated (NaOH and silane) fibres manufactured at different MC are presented in Figure 4.49(a) and Figure 4.49 (b). It can be seen that the average IFSS values of chemically treated samples decreased steadily with increasing MC, a trend consistent with the untreated samples. Figure 4.49(a) shows the influence of the NaOH concentrations on the IFSS across the MC range. All the NaOH treated samples had higher IFSS values as compared to the untreated samples when the samples were manufactured at low and medium MC. This can be explained by the removal of non-cellulose components such as hemicellulose and lignin as observed in the ATR-FTIR, which can increase the cellulose content creating more chemical bonding with the epoxy system [129][141]. The surface topographies of the untreated and treated flax fibre were studied and the SEM images are presented in Figure 4.51. It can be seen from Figure 4.51(a) that the surface untreated fibre is quite smooth, probably due to presence of wax and fats [27][123]. The possible fraction of middle lamella which cements the elementary fibres was also observed, containing pectin and lignin [114].

The benefits of alkalization can be seen in Figure 4.51(b) where microfibril traces are seen, resulting in a rougher surface compared to the untreated samples. This provides better mechanical interlocking at the fibre/matrix interface. Among the NaOH treatments, the 3.0% NaOH samples had the highest IFSS across the MC, suggesting an optimum NaOH concentration was achieved as shown in Figure 4.49(a). The 4.5% NaOH samples didn’t increase the adhesion strength at the interface as expected and can be attributed to excessive degradation of the flax fibre surface which can be seen in Figure 4.51(c) [114].

In regard of the silane treated samples, it can be seen that that IFSS of these samples were higher than the untreated samples for all MCs as shown in Figure 4.49(b). The influence of silane concentration was clearly seen here. The IFSS improved from 22.6 MPa to 24.3 MPa when the silane concentration increased from 1.0% to 1.5% when the samples were manufactured at low MC. Here, the IFSS improvement was around 7.5%. At a high MC of 13% MC, it turns out the effect of silane concentration was minimal because both samples had very similar IFSS

165

values, ~19 MPa. The IFSS improvement after silane treatment is due to the hydrophobic coating on the fibre and the crosslinking reactions between the epoxy matrix and the silane grafting on the flax fibres [133][135]. SEM images for 1.0% silane and 1.5% silane fibres are presented in Figure 4.51(d) and Figure 4.51(e). Traces of microfibrils were also seen in both images but the surface roughness was less than for the NaOH treated fibres.

(a) 30

15 IFSS (MPa) IFSS

Untreated 10 1.5% NaOH 3.0% NaOH 4.5% NaOH 5 0 2 4 6 8 10 12 14 16 MC (wt.%)

(b) 30

15 IFSS (MPa) IFSS

10 Untreated 1.0% Silane 1.5% Silane 5 0 2 4 6 8 10 12 14 16 MC (wt.%)

Figure 4.49: Influence of MC on the IFSS after (a) NaOH and (b) silane treatments.

IFSS of Combined Silane and NaOH Treated Fibres

The combination of NaOH and silane treatments, the 3.0% NaOH & 1.0% samples, aimed to increase the IFSS values over the standalone treatments and the IFSS results are presented in Figure 4.50. Although it can be seen that the IFSS

166

values are slightly higher than the untreated samples across the MC range, they are generally weaker than their standalone treatments, 3.0% NaOH and 1.0% silane respectively. This suggests that the combined treatments produced less effective load transfer at the fibre/matrix interface. In the literature, some authors also reported that combined NaOH and silane treatment was not beneficial to improving the mechanical properties of natural fibre composites [180][252]. The SEM image of the 3.0% NaOH & 1.0% fibre topography is shown in Figure 4.51(f). Unfortunately, it appears there is no appreciable difference with the 3.0% NaOH fibres except that the fibres are relatively clean and traces of microfibrils can be seen.

IFSS (MPa) IFSS 15

Untreated 10 3.0% NaOH 1.0% Silane 3.0% NaOH & 1.0% Silane 5 0 2 4 6 8 10 12 14 16 MC (wt.%)

Figure 4.50: IFSS comparison between 3.0% NaOH, 1.0% silane and combined treatment as a function of MC.

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(a) (b)

Middle lamella

Surface degradation

(e) (f)

Figure 4.51: SEM images of flax fibres for different conditions: (a) Untreated, (b) 3.0% NaOH, (c) 4.5% NaOH, (d) 1.0% silane (e) 1.5% silane and (f) 3.0% NaOH & 1.0% silane treatments.

168

The surface topography of IFSS samples after microbond testing are shown in Figure 4.52, Figure 4.53 and Figure 4.54 for composites produced at low and high MC following a 3.0% NaOH ,1.0% silane and combined treatment (3.0% NaOH & 1.0% silane). In Figure 4.52(a), it is interesting to observe that there is a large peeling of the 3.0% NaOH treated fibre for the samples manufactured at low MC, suggesting stronger interaction at the fibre/matrix interface than the strength between the fibre layers. Another possible explanation of the peeling is that the strength of the fibre layers was compromised after the removal of the pectin and hemicellulose during NaOH treatment [250].

Peeling of the fibres was also observed in the silane and combined NaOH- silane treatment when the samples were manufactured at low MC, as shown in Figure 4.53(a) and Figure 4.54(a), suggesting an improvement in fibre-matrix adhesion. In the case of silane treatment, the fibre was hydrolysed in alkaline medium (pH 9) and this may be responsible for the weakening of the fibre layers [250]. However, the peeling of the fibres was less severe as opposed to the NaOH treated samples as shown in Figure 4.52(a). This might explain the high IFSS data of the 3.0% NaOH sample as compared to the others chemical treatment at low MC (Figure 4.50). For treated samples that were manufactured at high MC, as shown in Figure 4.52(b), Figure 4.53(b) and Figure 4.54(b), the failure surfaces of the fibre were smooth along the pull-out region indicating weak fibre/matrix adhesion, resulting in lower IFSS values.

169

Figure 4.52: SEM images showing surface morphology after microbond test of (a) 3.0% NaOH-2.4% wt. MC and (b) 3.0% NaOH-12.9% wt. MC.

170

Figure 4.53: SEM images showing surface morphology after microbon test of (a) 1.0% silane-2.5% wt. MC and (b) 1.0% silane-12.7% wt. MC

171

Figure 4.54: SEM images showing surface morphology after microbond test of (a) 3.0% NaOH & 1.0% silane-2.4% wt. MC and (b) 3.0% NaOH & 1.0% silane-13.0% wt. MC.

172

4.7 Mechanical Performance of Chemically Treated Flax/Epoxy Composites

To study the effect of MC during composites manufacturing on the mechanical performance of treated flax/epoxy composites, the composites were subjected to tensile, flexural and Mode I DCB interlaminar fracture toughness tests and comparison were made with the untreated flax/epoxy composites.

4.7.1 Influence of Moisture Content on Fibre Volume Fraction of Composites

The determined fibre volume fraction, 푉푓 of both treated and untreated composites is plotted against MC in Figure 4.55. It can be seen the Vf of treated composites was lower than the untreated composites across all MC. Such changes of

푉푓 after chemical treatment was reported in the literature [129][177][266][220]. It is not surprising because the all chemical treatments induced swelling of the fibre as shown earlier. The increase swelling would require would require additional compaction pressure to attain similar 푉푓 as the untreated. Additionally, the absorbed water caused variation of 푉푓 of these treated composites by creating more compaction force on the laminates, causing 푉푓 to increase with MC as shown in Figure 4.55.

0.32

0.28 f

0.24

0.20 Untreated 3.0% NaOH 0.16 1.0% Silane FibreVolume Fraction, V 3.0% NaOH & 1.0% Silane

0.12 2 4 6 8 10 12 14 MC (wt.%)

Figure 4.55: Influence of MC on fibre volume fraction, 푉푓 of composites.

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4.7.2 Dynamic Mechanical Analysis

4.7.2.1 Storage Modulus

Figure 4.56(a-c) showed the comparison of storage modulus, E’ between untreated and treated flax/epoxy composites produced with different MC as a function of temperature. The E’ of treated composites decreases gradually in the glassy region. The E’ then falls rapidly in the transition regions and remains consistent in the rubbery region. In all cases, it can be seen that the chemically- treated composites had lower E’ than the untreated composites over the full temperature range. The decrease in E’ for the chemically treated composites is probably due to a low stiffness of the treated flax fabric resulting from the measured loss of tensile strength and Young’s modulus of treated elementary fibres. To support the argument, the work of Onaghi et al. [213] on sisal/glass hybrid polyester composites is referenced. Increasing glass fibre content in the hybrid composites resulted in higher E’, which is associated with high elastic modulus of the glass fibre. Additionally, the increased % crimp of the treated fabric due to fibre swelling, causing slightly lower 푉푓 these composites also contributes to the lower E’. Nevertheless, the presence of MC during composites fabrication is manifest in Figure 4.56(a-c) where low MC samples had highest storage modulus, showing the stronger interface properties of these composites [208][209].

(a) 3500

3000

2500

2000

1500 Untreated-2.9 wt.% MC Untreated-6.3 wt.% MC 1000 Untreated-12.9 wt.% MC StorageModulus (MPa) 3.0% NaOH-2.5 wt.% MC 500 3.0% NaOH-6.6 wt.% MC 3.0% NaOH-12.7 wt.% MC 0 30 60 90 120 150 180 210 Temperature (oC)

174

(b) 3500

3000

2500

2000

1500 Untreated-2.9 wt.% MC Untreated-6.3 wt.% MC 1000 Untreated-12.9 wt.% MC StorageModulus (MPa) 1.0% silane-2.6 wt.% MC 500 1.0% silane-6.7 wt.% MC 1.0% silane-12.9 wt.% MC 0 30 60 90 120 150 180 210 Temperature (oC)

(c) Untreated-2.9 wt.% MC 4000 Untreated-6.3 wt.% MC Untreated-12.9 wt.% MC 3500 3.0% NaOH & 1.0% silane-2.8 wt.% MC 3.0% NaOH & 1.0% silane-6.4 wt.% MC 3000 3.0% NaOH & 1.0% silane-13.1 wt.% MC

2500

2000

1500

1000 StorageModulus (MPa) 500

0 30 60 90 120 150 180 210 Temperature (oC)

Figure 4.56: Storage modulus of treated composites produced at different MC; (a) 3.0% NaOH, 1.0% silane and 3.0% NaOH & 1.0% silane.

4.7.2.2 Loss Modulus

Figure 4.57(a-c) shows comparisons of loss modulus, E” between untreated and treated composites as a function of temperature for composites produced at different MC. The E” height peak also indicates the glass transition, Tg. The influence of MC can be seen where the peaks shift to lower temperatures with increasing MC which means a decrease in Tg, signifying that that plasticization occurred in these composites. The peak E” of treated composites decreased with increasing MC suggesting lower energy dissipation [214].

175

(a) 450 Untreated-2.9 wt.% MC Untreated-6.3 wt.% MC 400 Untreated-12.9 wt.% MC 3.0% NaOH-2.9 wt.% MC 350 3.0% NaOH-6.3 wt.% MC 250 3.0% NaOH-12.9 wt.% MC 300

200 250

150 200

100 150 120 150 180 LossModulus (MPa) 100

0 30 60 90 120 150 180 210 o Temperature ( C)

(b) 450 Untreated-2.9 wt.% MC Untreated-6.3 wt.% MC 400 Untreated-12.9 wt.% MC 1.0% silane-2.6 wt.% MC 350 1.0% silane-6.7 wt.% MC

250 1.0% silane-12.9 wt.% MC 300

200 250

150 200 100

150 120 150 180

LossModulus (MPa) 100

0 30 60 90 120 150 180 210 o Temperature ( C)

(c) 450 Untreated-2.9 wt.% MC Untreated-6.3 wt.% MC 400 Untreated-12.9 wt.% MC 3.0% NaOH & 1.0% silane-2.8 wt.% MC 350 3.0% NaOH & 1.0% silane-6.4 wt.% MC 3.0% NaOH & 1.0% silane-13.1 wt.% MC 300 250 250 200

200 150

150 100

120 150 180 LossModulus (MPa) 100

0 30 60 90 120 150 180 210 o Temperature ( C)

Figure 4.57: Loss modulus of treated composites produced at different MC; (a) 3.0% NaOH, 1.0% silane d and 3.0% NaOH & 1.0% silane.

176

4.7.2.3 Tan Delta

Figure 4.58(a-c) shows the comparison of tan delta of the untreated and treated composites produced at different MC. An increase of tan delta peak height is observed for the treated flax/epoxy composites compared to the untreated samples. The average tan delta peak height of modified composites is given in Table 4.13. The observed increase of the tan delta peak height can be explained by slight lower fibre volume fraction of these treated composites resulting from swelling of the treated flax fabric as explained in Section 4.5.5. The glass transition, Tg measured from the tan delta peak is also listed in Table 4.13. With increasing MC, it can be seen that the Tg shows a similar trend to the untreated samples where the Tg values reduced. Lower Tg indicates that movement of the polymer chain segments at the interface is less constraint with the presence of moisture and hence, weaker fibre/matrix interface bonding occurred. In other words, the MC in the treated fibre had possibly plasticised or affecting polymerisation of the epoxy matrix.

The interfacial strength indicator, B proposed by Shinoj el al. [218] was calculated quantitatively using Equation 4.1 and comparison of B values between untreated and treated composites is presented in Figure 4.59. A higher B value indicates greater interfacial adhesion strength [221][220]. It is interesting to observe that in most cases, the B values of the treated composites were higher than the untreated composites, showing the fibre/matrix interface was actually improved after the surface treatment and the result was consistent with IFSS. It can be seen that the B values of 3.0% NaOH and 1.0% silane decreased steadily with increasing MC and this corresponds to weaker interfacial adhesion. In the case of the of 3.0% NaOH and 1.0% silane composites, a slight increase of the B value was observed at medium MC and then it then reduced again at highest MC.

177

(a) 0.5 Untreated-2.9 wt.% MC Untreated-6.3 wt.% MC Untreated-12.9 wt.% MC 3.0% NaOH-2.5 wt.% MC 0.4 3.0% NaOH-6.6 wt.% MC 3.0% NaOH-12.7 wt.% MC

0.3

TanDelta 0.2

0.1

0.0 120 150 180 210 o Temperature ( C)

(b) 0.5 Untreated-2.9 wt.% MC Untreated-6.3 wt.% MC Untreated-12.9 wt.% MC 0.4 1.0% silane-2.6 wt.% MC 1.0% silane-6.7 wt.% MC 1.0% silane-12.9 wt.% MC 0.3

TanDelta 0.2

0.1

0.0 120 150 180 210 o Temperature ( C)

(c) Untreated-2.9 wt.% MC Untreated-6.3 wt.% MC 0.5 Untreated-12.9 wt.% MC 3.0% NaOH & 1.0% silane-2.8 wt.% MC 3.0% NaOH & 1.0% silane-6.4 wt.% MC 0.4 3.0% NaOH & 1.0% silane-13.1 wt.% MC

0.3

TanDelta 0.2

0.1

0.0 120 150 180 210 o Temperature ( C)

Figure 4.58: Tan delta of treated composites produced with different MC; (a) 3.0% NaOH treated, 1.0% silane treated and 3.0% NaOH & 1.0% silane.

178

Table 4.13: The average tan delta peak height, Tg (peak max) and tan delta FWHM of untreated and treated composites produced at different fibre MC.

Tan Delta Tan Delta Composites Tg (ºC ) Peak Height FWHM (ºC )

Untreated

Untreated-2.9 wt.% MC 0.28 ± 0.005 162.0 ± 0.1 22.18 ± 0.49

Untreated-6.3 wt.% MC 0.27 ± 0.003 161.4 ± 0.2 23.30 ± 0.64

Untreated-12.9 wt.% MC 0.25 ± 0.006 156.6 ± 0.1 25.66 ± 0.13

3.0% NaOH 3.0% NaOH-2.5 wt.% MC 0.34 ± 0.008 161.2 ± 0.7 23.00 ± 0.23 3.0% NaOH-6.6 wt.% MC 0.32 ± 0.007 159.6 ± 0.9 24.40 ± 0.54 3.0% NaOH-12.7 wt.% MC 0.27 ± 0.003 157.3 ± 0.7 24.34 ± 0.46 1.0% Silane 1.0% silane-2.6 wt.% MC 0.31 ± 0.004 164.3 ± 0.1 21.55 ± 0.17 1.0% silane-6.7 wt.% MC 0.30 ± 0.001 158.8 ± 0.3 23.98 ± 0.46 1.0% silane-12.9 wt.% MC 0.25 ± 0.004 157.1 ± 0.6 26.73 ± 0.29 3.0% NaOH & 1.0% silane 3.0% NaOH & 1.0% silane-2.8 wt.% MC 0.36 ± 0.003 165.8 ± 0.7 21.24 ± 0.55 3.0% NaOH & 1.0% silane-6.4 wt.% MC 0.32 ± 0.013 162.2 ± 0.9 23.27 ± 0.84 3.0% NaOH & 1.0% silane-13.1 wt.% MC 0.26 ± 0.018 160.5 + 0.9 28.03 ± 0.59

2.7

B 2.6

2.5

2.4

2.3

2.2 Untreated 3.0% NaOH 1.0% silane Interfacial strength Interfacial indicator, 2.1 3.0% NaOH & 1.0% Silane 2.0 0 2 4 6 8 10 12 14 MC (wt.%)

Figure 4.59: Interfacial strength indicator, B of untreated and chemically treated as a function of MC.

179

4.7.3 Tensile Properties

Figure 4.60(a-c) shows typical tensile stress-strain curves of untreated and treated flax/epoxy composites manufactured at different MC. The tensile stress- strain curves were normalised at 30% 푉푓 for fair comparison. It can be seen there is no appreciable difference in the curves between the untreated and treated flax/epoxy composites. The curves begin with a proportional linear line, corresponding to the elastic properties of the composites. The curves then show non-linear behaviour where changes of gradient were observed around 0.25% strain. All the treated flax/epoxy composites failed in a brittle manner without apparent yielding. Although the curves’ behaviour remains similar, it can be seen that all the chemical treatments caused significant changes to the tensile properties. The increase of tensile strain at break with increasing of MC probably results from plasticization of the flax fibre which similar to that untreated composites and the matrix [150].

(a) 70

30 Untreated-2.8 wt.% MC Untreated-6.3 wt.% MC

20 Untreated-12.9 wt.% MC Tensile(MPa) Stress 3.0% NaOH-2.4 wt.% MC 10 3.0% NaOH-5.9 wt.% MC 3.0% NaOH-13.1 wt.% MC 0 0.0 0.5 1.0 1.5 2.0 Tensile Strain (%)

(b) 70

30 Untreated-2.8 wt.% MC Untreated-6.3 wt.% MC 20

Untreated-12.9 wt.% MC Tensile(MPa) Stress 1.0% silane-2.7 wt.% MC 10 1.0% silane-6.4 wt.% MC 1.0% silane-12.9 wt.% MC 0 0.0 0.5 1.0 1.5 2.0 Tensile Strain (%)

180

30 Untreated-2.8 wt.% MC 20 Untreated-6.3 wt.% MC

Tensile(MPa) Stress Untreated-12.9 wt.% MC 10 3.0% NaOH-1.0% silane-2.6 wt.% MC 3.0% NaOH-1.0% silane-6.1 wt.% MC 3.0% NaOH-1.0% silane-12.9 wt.% MC 0 0.0 0.5 1.0 1.5 2.0 2.5 Tensile Strain (%)

Figure 4.60: Typical flexural stress-strain curves of (a) 3.0% NaOH, (b) 1.0% silane and (c) 3.0% NaOH &1.0% silane samples. The curves were

normalised at 30% 푉푓.

For better understanding of the influence of MC between the untreated and treated flax/epoxy composites, tensile strength and tensile modulus of samples were plotted against MC and are presented in Figure 4.61(a) and Figure 4.61(b). The average tensile strength, tensile modulus and tensile strain are also given in Table 4.14. Generally, it is observed that the tensile strength was reduced after the chemical treatments. Depending on the MC and type of chemical treatments, the percentage reduction can be as high as 22% for tensile strength. It is known that tensile properties of composites are fibre-dominated and hence, the properties of the fibres strongly affect the final performance of composites. The reduction of tensile properties in these composites is mainly due to the reduction in mechanical properties of flax fibres after chemical treatments and this argument is supported by the single fibre tensile results between untreated and treated flax fibre. These findings are also in agreement with those of Li et al. [220], who investigated the influence of surface treatment (silane, permanganate) on sisal textile/vinyl-ester composites. They found the tensile strength of silane treated reduced on treatment although there was improvement of interfacial properties between untreated and silane treated sisal.

181

Table 4.14: Tensile properties of untreated and treated composites

Normalised Normalised Tensile Composites tensile strength tensile modulus strain at (MPa) (GPa) break (%) Untreated Untreated-2.9 wt.% MC 67.4 ± 2.1 6.19 ± 0.21 1.47 ± 0.05 Untreated-6.3 wt.% MC 62.6 ±.1.6 6.18 ± 0.12 1.62 ± 0.09 Untreated-12.9 wt.% MC 57.3 ± 1.5 5.75 ± 0.23 1.74 ± 0.10 3.0% NaOH 3.0% NaOH-2.4 wt.% MC 55.2 ±1.4 5.55 ± 0.33 1.46 ± 0.03 3.0% NaOH-5.9 wt.% MC 54.6 ± 1.1 5.22 ± 0.15 1.68 ± 0.12 3.0% NaOH-13.1 wt.% MC 53.5 ± 2.6 4.98 ± 0.21 1.72 ± 0.20 1.0% Silane 1.0% silane-2.7 wt.% MC 57.2 ± 1.1 6.34 ± 0.18 1.34 ± 0.05 1.0% silane-6.4 wt.% MC 53.7 ± 2.1 5.79 ± 0.25 1.46 ± 0.10 1.0% silane-12.9 wt.% MC 53.1 ± 3.0 4.85 ± 0.39 1.72 ± 0.09 3.0% NaOH & 1.0% silane 3.0% NaOH & 1.0% silane-2.6 wt.% MC 52.6 ± 1.4 6.18 ± 0.24 1.21 ± 0.09 3.0% NaOH & 1.0% silane-6.1 wt.% MC 53.5 ± 1.0 5.53 ± 0.22 1.67 ± 0.08 3.0% NaOH & 1.0% silane-13.1 wt.% MC 51.2 ± 2.0 4.72 ± 0.28 2.00 ± 0.09

(a) 70

Untreated Tensile(MPa) Strength 45 3.0% NaOH 1.0% silane 3.0% NaOH-1.0% silane 40 0 2 4 6 8 10 12 14 MC (wt.%)

182

(b) 7.0

6.5

6.0

5.5

5.0 Untreated TensileModulus (MPa) 3.0% NaOH 4.5 1.0% silane 3.0% NaOH-1.0% silane 4.0 2 4 6 8 10 12 14 MC (wt.%)

Figure 4.61: Influence of MC on the (a) tensile strength and (b) tensile modulus of

treated composites. Data points were normalised at 30% Vf.

As shown in Table 4.14, the tensile strength of 3.0% NaOH reduced consistently with MC, from 55.2 MPa to 54.6 MPa and finally 53.5 MPa. The tensile strength of 1.0% silane was the highest (57.2 MPa) when the composite was produced at low MC. However, MC had a limited effect at medium and high levels and both composites had similar tensile strength, ~53 MPa. In the case of the combined 3.0% NaOH and 1.0% silane treatment, there was a slight tensile strength reduction between low MC and high MC from 52.6 MPa to 51.2 MPa. Another interesting observation found from Figure 4.61(a) is that the treated composites had the lowest downturn behaviour as compared to the untreated composites. The percentage of composites’ tensile strength loss between the lowest and highest MC was calculated and is given as follows: Untreated (14.9%) > 1.0% silane (7.2%) > 3.0% NaOH (2.9%) > 3.0% NaOH & 1.0% silane (2.7%). This suggests surface treatment could minimize the reduction of tensile strength due to increasing MC but the reduction in fibre strength following treatment needs to be addressed.

Comparing the treated flax/epoxy composites, it can be seen the 1.0% silane had the highest tensile strength compared to the others treatments. However, when the composites were manufactured at high MC, both 1.0% silane and 3.0% NaOH had similar tensile strength. In the case of combined 3.0% NaOH & 1.0% silane treatment, these composites were relatively weak because they had lowest tensile strength across all MC compared to other treatments. 183

Similar to the tensile modulus of untreated composites, the tensile modulus of treated flax/epoxy composites consistently reduced as the MC in the fibres increased as shown in Table 4.14, proposing that the weak fibre/matrix interface was the main factor in reducing the elastic properties of the composites and this argument supported by the decreased Tg of the composites in DMA testing and the IFSS. As shown Figure 4.61(b), it is interesting to note that the 1.0% silane treatment had the highest tensile modulus (manufactured at low MC) but the rest of the treated composites had similar or lower tensile modulus than untreated composites. At high MC, the tensile modulus improvement of 3.0% NaOH composites over of 1.0% silane composites was minimal at 4.98 GPa and 4.85 GPa respectively, while combination of surface treatment (3.0% NaOH & 1.0% silane) performed the worst, giving 4.72 GPa.

As shown in Figure 4.60(a-c), it is interesting to observe that the tensile modulus of treated composites had a lower slope compared to the untreated composites with increasing MC. This may be explained by the fact that the fibre that was damaged by the surface treatment and the presence of water could additionally weaken the internal the flax fibres due to swelling, particularly in the radial direction [150][173]. Scida et al. [267] observed the internal bonding between microfibrils layers in the S2 region was compromised when the cross-section of flax/epoxy composites was polished with water. The same observation was thought to happen in this investigation. The SEM images of tensile fracture samples for low and high MC untreated MC samples are shown in Figure 4.62(a1) and Figure 4.62(a2). There is a slight bonding failure between microfibrils layers in the S2 region shown in Figure 4.62(a2) as compared to Figure 4.62(a1). In the case of surface treated composites, the separation of microfibril layers in the S2 regions was greater, causing large voids within the fibre as seen in Figure 4.62(b) and Figure 4.62(c) for 1.0% silane and combined 3.0% NaOH & 1.0% silane treated composites that were produced at medium MC.

184

(a1) (a2) Untreated-12.9 wt.% MC

Bonding failure between microfibrils Untreated-2.8 wt.% MC layers

(b) (c)

1.0% silane-6.4 wt.% MC 3.0% NaOH & 1.0% silane-6.1 wt.% MC

Figure 4.62: SEM imaging of untreated and treated composites showing internal failure in the fibre due to MC.

4.7.4 Flexural Properties

Figure 4.63(a-c) shows a comparison of typical of typical flexural stress- strain curves of untreated and treated flax fibre composites MC manufactured with different MC, loaded in three point bending configuration. The flexural stress-strain curves were normalised at 30% Vf. All the curves began in linear elastic

185

deformation, followed by non-linear behaviour up to failure. The samples failed in a brittle manner with no evidence of yielding during the flexural test. From Figure 4.63(a-c), it can be seen that combinations of chemical treatments and MC caused significant changes to the flexural properties of the composites.

The flexural properties of untreated and treated flax/epoxy composites are plotted against MC in Figure 4.64(a) and in Figure 4.64(b). The average data of flexural properties are given are given in Table 4.15. Interestingly, in some cases (depending on the MC and type of surface treatment) the flexural strength and flexural modulus of treated composites were somewhere higher than the untreated composites. It is not surprising because flexural properties are often related to the matrix-dominated properties and the fibre-matrix interfacial strength. However, any loss of fibres strength due to MC and chemical treatments can reduce the flexural strength of these composites. The evidence of loss of strength and modulus in flax fibres after chemical treatments was presented in Section 4.5.4.

With increasing MC, common behaviour can be established for the treated composites in Figure 4.64(a). In comparison with untreated composites for which consistently flexural strength reduced across the MC, treated composites showed a showed a slight increased at medium MC but which dropped later at high MC. This can be related to the experimental work of Baley et al. [21], who investigated dried and non-dried flax fibre reinforced epoxy composites. They explained that at MC removal could induce mechanical stresses in the fibre and ultimately, lowering the mechanical properties of the fibre. In the present investigation, the fibre was humidified at 10% RH and the measured MC was around 2.5 wt.% which was closer to their dried condition. These fibres were already physically deteriorated by the action of surface treatment and decreasing MC could induce more damages within the fibre (as evidently in the SEM images in Figure 4.62). Hence, weak flexural strength of treated composites resulted at low MC compared to medium MC.

186

(a) 120

100

Untreated-2.1 wt.% MC 40 Untreated-6.2 wt.% MC FlexuralStress (MPa) Untreated-12.6 wt.% MC 20 3.0% NaOH-2.5 wt.% MC 3.0% NaOH-5.9 wt.% MC 3.0% NaOH-12.6 wt.% MC 0 0 1 2 3 4 Flexural Strain (%)

(b) 100

40 Untreated-2.1 wt.% MC Untreated-6.2 wt.% MC

Untreated-12.6 wt.% MC FlexuralStress (MPa) 20 1.0% silane-2.6 wt.% MC 1.0% silane-6.4 wt.% MC 1.0% silane-12.6 wt.% MC 0 0 1 2 3 4 Flexural Strain (%)

100

40 Untreated-2.1 wt.% MC

FlexuralStress (MPa) Untreated-6.2 wt.% MC Untreated-12.6 wt.% MC 20 3.0% NaOH & 1.0% silane-2.4 wt.% MC 3.0% NaOH & 1.0% Silane-6.1 wt.% MC 0 3.0% NaOH & 1.0% Silane-12.9 wt.% MC 0 1 2 3 4 5 6 Flexural Strain (%)

Figure 4.63: Typical flexural stress-strain curves of (a) 3.0% NaOH, (b) 1.0% silane and (c) 3.0% NaOH & 1.0% silane samples. The curves were

normalised at 30% 푉푓.

At medium MC, the flexural strength of 3.0% NaOH and combined treatment (3.0% NaOH & 1.0% silane) composites were equally the highest while the 1.0% silane had similar values to the untreated samples. This result emphasized the greater fibre/matrix adhesion achieved by these surface treatments although the strength of treated fibres was greatly reduced. For a composites that were

187

manufactured at high MC, both 1.0% silane and 3.0% NaOH composites exhibited higher flexural modulus than untreated composites.

Table 4.15: Flexural properties of untreated and treated composites

Normalised Normalised Flexural Composites Flexural Flexural Strain at Strength (MPa) Modulus (GPa) break (%) Untreated Untreated-2.1 wt.% MC 95.4 ± 3.2 4.93 ± 0.31 2.80 ± 0.17 Untreated-6.2 wt.% MC 94.3 ± 3.8 4.68 ± 0.14 3.08 ± 0.19 Untreated-12.6 wt.% MC 85.5 ± 3.8 4.38 ± 0.15 3.15 ± 0.14 3.0% NaOH 3.0% NaOH-2.4 wt.% MC 89.3 ± 5.0 4.59 ± 0.17 2.81 ± 0.13 3.0% NaOH-5.9 wt.% MC 97.7 ± 4.3 5.29 ± 0.50 2.76 ± 0.09 3.0% NaOH-12.6 wt.% MC 88.5 ± 3.1 3.72 ± 0.13 3.68 ± 0.11 1.0% Silane 1.0% silane-2.6 wt.% MC 89.4 ± 2.9 4.86 ± 0.31 2.73 ± 0.10 1.0% silane-6.4 wt.% MC 95.0 ± 6.7 4.99 ± 0.31 3.06 ± 0.34 1.0% silane-12.6 wt.% MC 91.0 ± 4.8 4.56 ± 0.26 3.00 ± 0.20 3.0% NaOH & 1.0% silane 3.0% NaOH & 1.0% silane-2.4 wt.% MC 92.9 ±2.9 4.95 ± 0.22 2.61 ± 0.17 3.0% NaOH & 1.0% silane-6.1 wt.% MC 97.1 ± 4.8 4.56 ± 0.23 3.45 ± 0.23 3.0% NaOH & 1.0% silane-12.9 wt.% MC 74.0 ± 2.3 2.16 ± 0.07 4.07 ± 0.03

In Figure 4.64(b), the majority of treated composites had lower flexural modulus across the range of range of MC, implying reduced adhesion strength between fibres and the matrix. At low MC, the flexural modulus of 1.0% silane and 3.0% NaOH & 1.0% silane samples were essentially the same. At medium MC, the flexural modulus of untreated composites was 4.68 GPa. The 3.0% NaOH treated composites showed improved flexural modulus of up to 13% over the untreated composites. The 1.0% silane treated composites also showed increased flexural modulus of 4.99 GPa, an increment about 0.31 GPa. For composites that were manufactured at high MC, the 1.0% silane treated composites were the only treated composites showing increased flexural modulus over the untreated composites. The flexural modulus improved slightly from 4.38 GPa (untreated) to 4.56 GPa. In

188

contrast, the combined 3.0% NaOH & 1.0% silane composites had the lowest flexural modulus similarly to their flexural strength shown in Figure 4.64(a).

(a) 110

100

80 Untreated

FlexuralStrength (MPa) 70 3.0% NaOH 1.0% silane 3.0% NaOH & 1.0% Silane 60 0 2 4 6 8 10 12 14 MC (wt.%)

(b) 6

3 Untreated 3.0% NaOH 2 FlexuralModulus (GPa) 1.0% Silane 3.0% NaOH & 1.0% silane 1 0 2 4 6 8 10 12 14 MC (wt.%)

Figure 4.64: Influence of MC on (a) flexural strength and (b) flexural modulus of

treated composites. Data points are normalised at 30% 푉푓.

SEM images of the fracture surfaces of treated composites are shown in Figure 4.65, which emphasize the morphology aspect between low and high MC during composite manufacturing. For composites produced at low MC, fibre peeling is observed in all specimens in which indicate better interface properties. In contrast, for composites manufactured at high MC, a visible gap between fibre/matrix observed signifies the fibre/matrix interface adhesion has had been compromised.

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(a1) 3.0% NaOH-2.5 wt.% MC (a2) 3.0% NaOH-12.6 wt.% MC Gap between fibre/matrix

Fibre peeling

(b1) 1.0% silane-2.6 wt.% MC (b2) 1.0% silane-12.6 wt.% MC

Fibre Peeling Gap between fibre/matrix

(c1) (c2) Gap between Fibre peeling fibre/matrix

3.0% NaOH & 1.0% silane-12.9 wt.% 3.0% NaOH & 1.0% silane-2.4 wt.% MC MC

Figure 4.65: SEM images of the fracture surfaces of treated composites: (a) 3.0% NaOH, (b) 1.0% silane and (c) combined 3.0% NaOH & 1.0% silane.

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4.7.5 Mode I Interlaminar Fracture Toughness

The influence of MC on Mode I interlaminar fracture toughness performance is presented in Figure 4.66(a-c) with untreated samples as a reference. The load- extension curves began with linear behaviour which is attributed to the elastic coefficient of the treated composites. It can be seen that all treated composites had lower linear gradients compared to the untreated composites due to the loss of strength of the flax fibre after the chemical treatments. In case of treated composites, a lower linear gradient was expected for high MC samples attributed to the weakened fibre-matrix interfacial [223]. The curves then showed non-linear behaviour towards the peak. Inspection of the test specimens (Figure 4.70 and Figure 4.71) indicated that crack propagation ion in the delamination plane was accompanied by crack branching and side cracking.

In Figure 4.66(a-c) it is worth noting that the load required for crack initiation was lower for treated samples but initiation occurred at higher extension (around 10 mm) as compared to the untreated samples (around 7 mm). The influence of MC on crack propagation behaviour is also shown in Figure 4.66(a-c). With increasing extension, the crack propagation of low MC samples was unstable, showing sudden loads rises and falls during crack opening. In contrast, in the high MC samples had very stable crack propagation. The transformation from unstable to stable crack propagation between the low and high MC samples was also observed for the untreated sampled and was attributed to the changes in fibre/matrix adhesion.

(a) 180

160

140

120

100

80 Load(N) 60 Untreated-2.3 wt.% MC 40 Untreated-12.7 wt.% MC 3.0% NaOH-2.4% wt. MC 20 3.0% NaOH-12.7% wt. MC 0 0 5 10 15 20 25 30 35 40 45 Extension (mm)

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(b) 180

160

140

120

100

80 Load(N) 60 Untreated-2.3 wt.% MC 40 Untreated-12.7 wt.% MC 1.0% silane-2.1 wt.% MC 20 1.0% silane-13.1 wt.% MC 0 0 5 10 15 20 25 30 35 40 45 Extension (mm)

80 Load(N) 60 Untreated-2.3 wt.% MC 40 Untreated-12.7 wt.% MC 20 3.0% NaOH &1.0% silane-2.2 wt.% MC 3.0% NaOH &1.0% silane-12.9 wt.% MC 0 0 5 10 15 20 25 30 35 40 45 Extension (mm)

Figure 4.66: Typical load-extension curves for (a) 3.0% NaOH (b) 1.0% silane and (c) 3.0% NaOH & 1.0% silane.

The resistance curves (R-Curves) of the untreated and untreated flax/epoxy composites are presented in Figure 4.67(a-c). It can be seen that all the R-curves of the treated samples have the same curves shape as the untreated samples where the

GIC increases at the beginning. The differences between untreated and treated are that the R-curves of treated samples are somewhat higher than those of the untreated samples. In fact, a sudden increase in GIC propagation for treated composites is seen in the first 5-15 mm of crack growth, signifying the improvement of fibre/matrix interfacial adhesion. The result is an agreement with other authors which indicates high GIC was attributed to enhanced interfacial bond strength of chemical treated fibre [155][156][220]. In a treated composite, the influence of MC was not obviously noticeable since the curves were overlapping with each other as they delaminated at higher crack length.

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(a) 3.2

2.8

2.4

) 2

2.0

(kJ/m

IC G 1.6 Untreated-2.3 wt.% MC Untreated-12.7 wt.% MC 1.2 3.0% NaOH-2.4 wt.% MC 3.0% NaOH-12.7 wt.% MC 0.8 0 5 10 15 20 25 30 35 40 45 Crack length, a (mm)

(b) 3.2

2.8

2.4

) 2

2.0

(kJ/m

IC G 1.6

Untreated-2.3 wt.% MC 1.2 Untreated-12.7% wt. MC 1.0% silane-2.5 wt.% MC 1.0% silane-13.3 wt.% MC 0.8 0 5 10 15 20 25 30 35 40 45 Crack length, a (mm)

2.8

2.4

) 2

2.0

(kJ/m IC G 1.6 Untreated-2.3 wt.% MC 1.2 Untreated-12.7 wt.% MC 3.0% NaOH &1.0% silane-2.5 wt.% MC 3.0% NaOH &1.0% silane-12.9 wt.% MC 0.8 0 5 10 15 20 25 30 35 40 45 Crack length, a (mm) Figure 4.67: Influences of MC on the R-Curves behaviour on (a) 3.0% NaOH, 1.0% silane and (c) 3.0% NaOH and 1.0% silane treated samples.

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Influence of Surface Treatment at Low MC on GIC initiation

Figure 4.68 shows a comparison of the Mode I GIC initiation-VIS values of the untreated and treated composites. A comparable reference data experiment from

Bensadoun et al. [34] was included[AHA26]. The average data for GIC initiation-VIS of treated samples is in Table 4.16. At low MC, there is a significant improvement of

GIC initiation-VIS values between 52% and 73% for the treated samples (1.83-2.07 3 2 kJ/m ) compared untreated samples (1.28 kJ/m ). The improvement in in GIC initiation-VIS indicates better interface adhesion was achieved by the treated composites. Among the chemical treatment, the highest GIC initiation-VIS values were in the order; 3.0% NaOH & 1.0% silane > 3.0% NaOH > 1.0% silane > untreated.

Influence of Surface Treatment at High MC on GIC initiation

At high MC, except for the combined 3.0% NaOH-1.0% silane composites, it was found that chemical treatments still increased the GIC initiation-VIS values over the untreated composites but increments were marginal. In this situation, GIC initiation-VIS of 3.0% NaOH composites was higher than 1.0% silane composites, 1.96 kJ/m2 and 1.82 kJ/m2 respectively. According to Liu and Hughes [268] toughness is most probably limited by the intrinsic properties of the fibre and hence, the weak GIC initiation-VIS of combined 3.0% NaOH-1.0% silane composites as compared to untreated samples was probably due to the loss of strength of flax fibre. The results are consistent with the low flexural properties and stiffness of the elementary fibres reported previously for these materials.

194

) 2.4 2

2.0

1.6

1.2 Initiation-VIS (kJ/m Initiation-VIS

IC 0.8 Untreated G 3.0% NaOH 0.4 1.0% silane 3.0% NaOH & 1.0% silane 0.0 0 2 4 6 8 10 12 14 MC (wt.%)

Figure 4.68: Comparison of Mode I GIC initiation-VIS values for untreated and treated flax/epoxy composites as a function of MC.

Table 4.16: The average of GIC initiation-VIS and GIC propagation of composites.

G initiation- G propagation Composites IC IC Ref. VIS (kJ/m2) (kJ/m2)

Flax/Epoxy Composites (Twill Low Twist[AHA27]) 0.75 1.60 [34] Flax/Epoxy Composites (Twill High Twist) 0.61 1.15 Untreated

Untreated-2.9 wt.% MC 1.28 ± 0.10 1.81 ± 0.20

Untreated-12.7 wt.% MC 1.76 ± 0.25 2.18 ± 0.33

3.0% NaOH

3.0% NaOH-2.5 wt.% MC 1.96 ± 0.13 2.32 ± 0.14

3.0% NaOH-12.7 wt.% MC 1.95 ± 0.14 2.30 ± 0.18 Present

1.0% Silane data

1.0% silane-2.6 wt.% MC 1.82 ± 0.39 2.24 ± 0.22

1.0% silane-12.9 wt.% MC 1.83 ± 0.11 2.31 ± 0.11

3.0% NaOH & 1.0% silane

3.0% NaOH & 1.0% silane-2.8 wt.% MC 2.07 ± 0.29 2.50 ± 0.18

3.0% NaOH & 1.0% silane-13.1 wt.% MC 1.66 ± 0.16 2.29 ± 0.19

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Influence of Surface Treatment at Low MC on GIC propagation

Figure 4.69 shows a comparison of the Mode I-GIC propagation values for the untreated and treated samples. The averages of GIC propagation are given in

Table 4.16. In general, greater GIC propagation values were exhibited by the treated composite compared to the untreated composites, in the order; 3.0% NaOH & 1.0% silane (2.50 kJ/m2) > 3.0% NaOH (2.32 kJ/m2) > 1.0% silane (2.24 kJ/m2) > untreated (2.18 kJ/m2). This indicates an improvement of the fibre/matrix interface. According to previous studies [224][163][230][269], low fracture energy is usually associated with the presence fibre bridging because of low interface quality and ultimately, promoting fibre/matrix debonding. Qualitative observation of fibre bridging was done on samples during the Mode I test and shown in Figure

4.70[AHA28], it is clear that fibre bridging was almost completely absent in all treated composites as compared to untreated samples, indicating stronger fibre/matrix interfacial in the low MC composites.

3.0

2.8 )

2 2.6 2.4 2.2 2.0

1.8 Propogation (kJ/m

IC Untreated G 1.6 3.0% NaOH 1.4 1.0% silane 3.0% NaOH & 1.0% silane 1.2 0 2 4 6 8 10 12 14 MC (wt.%)

Figure 4.69: Comparison of Mode I GIC propagation values for untreated and treated flax/epoxy composites as a function of MC.

Influence of Surface Treatment at High MC on GIC propagation

In case of high MC, there was slight improvement of GIC propagation of the treated samples over the untreated samples. This data was statically analysed using two-tailed t-tests at confidence level of 95% (γ = 0.05 level of significance).

According to the two-tailed results, there is enough evidence to show that the GIC

196

propagation of the treated samples is not significantly different. In term of surface treatments, none of those composites were actually performing better than the others. All the treated samples had similar values of ~2.3 kJ/m2. Figure 4.71 shows qualitative observation of fibre bridging phenomena where it can be seen that fibre bridging is dominant in the untreated and treated composites at high MC[AHA29].

With the results from GIC initiation-VIS and GIC propagation, improvement of the fibre/matrix interface is the main reason for the higher GIC of treated composites compared to the untreated when the samples was fabricated at low MC .

The influence of MC on the interface on delaminated surface of treated composites is shown in Figure 4.72. All the samples produced at low MC showed typical peeling of cell wall layers with traces of peeled cell wall visible on the matrix in the pull out region, signifying a strong interface. In contrast, for composites manufactured at high MC, it can be seen most fibre/matrix debonding regions had smooth surfaces either on the fibre itself or on the matrix surface. Hence, the fibres were easily pulled out from the matrix, showing weak interface bonding.

197

5 mm

Untreated-2.9 wt.% MC

5 mm

3.0% NaOH-2.5 wt.% MC

5 mm

1.0% silane 2.6 wt.% MC

5 mm

3.0% NaOH & 1.0% silane-2.8 wt.% MC

Figure 4.70: Qualitative observation [AHA30]of fibre bridging (white arrows) on Mode I DCB test between untreated and treated composites when the composites were produced with low MC.

198

5 mm

Untreated-12.9% wt. MC

5 mm

3.0% NaOH-12.7% wt. MC

5 mm

1.0% silane-12.9% wt. MC

5 mm

3.0% NaOH & 1.0% silane-13.1% wt. MC

Figure 4.71: Qualitative observation [AHA31]of fibre bridging (white arrows) on Mode I DCB test between of untreated and treated composites when the composites were produced at high MC.

199

(a1) (a2)

3.0% NaOH-2.5% wt. MC 3.0% NaOH-12.7% wt. MC

(b1) (b2)

1.0% silane-2.6% wt. MC 1.0% silane-12.9% wt. MC

(c1) (c2)

3.0% NaOH & 1.0% silane-2.8% 3.0% NaOH & 1.0% silane-13.1% wt. MC wt. MC

Figure 4.72: SEM images of the delaminated surfaces of treated composites: (a) 3.0% NaOH (b) 1.0% silane and (c) combined 3.0% NaOH & 1.0% silane treatments. White arrows locate the peeling of the cell wall on the fibre or matrix surface. 200

Influence of MC on Interlaminar Fracture Toughness of Treated Composites

As explained in Section 4.3.5, the increased of GIC of untreated composites was probably attributed to the toughening effect of matrix with presence of MC; behaviour which is agreement with Nathan [86], who investigated the addition of water during cure on carbon/epoxy composites. The moisturized fibre might plasticize at the interface but also in the matrix nearby the fibre interface during cure.

Among the treated composites, the relation between MC and interlaminar fracture toughness is seen in Figure 4.68 (GIC initiation-VIS) and Figure 4.69 (GIC propagation). It can be seen that GIC initiation-VIS and GIC propagation for 3.0% NaOH and 1.0% silane composites remain consistent which is in contrast with the untreated composites. The results suggest the interface after surface treatment can compensate for the reduced toughening of the matrix at low MC and hence, consistent interlaminar fracture toughness values were obtained at high MC. This argument is supported by the reduced glass transition, Tg of treated composites which also suggests the plasticisation of the matrix due to moisture. Additionally, colour changes of the delaminated treated composites can be seen in Figure 4.73. Similarly to the untreated samples, the changes of colour from dark brown to light brown were observed clearly with increasing MC, suggesting a chemistry changes within the interface or/and epoxy matrix with the presence of moisture [199][226].

Figure 4.73: Photograph of delaminated treated composites showing colour changes between (a) NaOH (b) silane and (c) combined 3.0% NaOH & 1.0% silane.

201

Chapter 5 Conclusions

The aim of this study was to evaluate the influence of fibre MC on flax fibre properties, fibre/matrix interface properties and the properties of the composites. The effect of surface treatment such as NaOH, silane and combined silane-alkalized on moisture content were studied[AHA32]. The fibre/matrix interfaces and their flax/epoxy composites properties were also evaluated. The overall objectives of this work were successfully achieved.

5.1 Effect of Moisture Content

 Using a standardized flax fabric specimen, the MC of flax fibres can be estimated during composite fabrication. The fibre density decreases with increasing MC and there is a positive correlation between experimental data and Messiry’s proposed model.

 Flax fibre is an extremely hygroscopic material where moisture uptake followed the Fickian diffusion behaviour and it accelerated quickly at higher

RH[AHA33] with higher MC.

 The microbond test showed that IFSS is consistently reduced with increasing MC (18.4 MPa to 14.7 MPa). The strong interface was characterized by a rough surface in the debonding area as shown in SEM images.

 Increased MC caused a thinner composite that is corresponding to the increase fibre volume fraction. The absorbed water created extra compaction force to the preforms upon the composites manufacturing (vacuum infusion).

202

 The tensile and flexural properties were reduced accordingly with increasing

MC. DMA testing also revealed that the matrix glass transition, Tg reduced with increasing MC for which suggests lower cross-link density, corresponding to weak interface bonding.

 Interlaminar shear strength (ILSS) and Mode II ENF could not be determined because the test results were invalid. In the Mode II ENF test, the main crack deviated from the mid-plane severely. The low bending stiffness of flax fibre, high interlaminar fracture toughness and the presence of yarn crimp in the fabric are among the contributing factors.

 With increasing MC from 2.3 wt.% to 6.6 wt.% MC, the GIC initiation and

GIC propagation was found to increase which is likely due to toughening or plasticization of epoxy matrix with presence of MC. A slight decreases of

GIC initiation and GIC propagation occurred afterward at 12.6 wt.% MC, corresponding to the weak fibre/matrix interface.

 Impact testing (1.5J and 2.0J impact loading) showed that the untreated 12.5 wt.% MC samples absorbed higher energy than the untreated 6.5 wt.% MC samples which is possibly due to weak fibre/matrix interfacial bonding. High MC in composites during fabrication resulted higher contact time and peak force and lower bending stiffness.

 The CAI results indicate that little delamination occurred in the impacted flax/epoxy composites due to the high interlaminar fracture toughness of these composites.

5.2 Effect of Surface Treatments

 MC was modestly reduced after surface treatment (3.0% NaOH, 4.5% NaOH, 1.0% silane and 3.0% NaOH-1.0% silane) but their effectiveness was limited between 10% and 60% RH in the monolayer water formation.

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 The FTIR results indicate that the MC reduction after NaOH was attributed to the hemicellulose removal, whilst in case of silane treatment; the hygroscopic hydroxyl sites in the fibre are reduced after the formation of silanol bridges. The evidence of moisture reduction in 1.0% silane can be located at 1636 cm-1 band since the position related to the hold water.

 The density of treated fibres was higher than untreated due to cell wall densification. The density of treated fibres decreased with increasing MC which is consistent with untreated samples. The fibre density of 3.0% NaOH & 1.0% silane treated samples did not fit with Messiry’s proposed model.

 Microbond testing shows improvement of IFSS for the treated samples compared to untreated samples due to mechanical interlocking and silane coupling at the interface. 3.0% NaOH is the optimized NaOH concentrations for which improvement of IFSS is seen for all MC as compared 1.5% NaOH and 4.5% NaOH. The 4.5% NaOH was thought to be an excessive concentration for which causing surface degradation on surface of fibre.

 A slight improvement of IFSS for the 1.5% silane over 1.0% silane was achieved at low and medium MC. The combined 3.0% NaOH & 1.0% silane treated composites showed weaker IFSS compared to the standalone treatment at all MC.

 Tensile strength and tensile modulus of treated composites are lower than untreated composite in all MC. This is likely due to the loss of fibre strength and stiffness after surface treatment. These results were supported by single fibre tensile testing between untreated and treated flax fibre. With presence of moisture in treated fibres, this could additionally weaken the internal fibre properties where failure internal bonding between microfibrils layers which can be found in SEM images. In case of flexural testing, an improvement of

204

flexural modulus of treated composites was achieved only at medium MC compared to the untreated composites.

 The calculated interfacial strength indicator, B showed the treated composites had higher values than untreated composites which mean there was an improvement of fibre/matrix interface. The B values decrease with

MC increasing for which consistent with reduction in Tg, signifying lower cross-link density that lead to poor fibre/matrix interfacial bonding.

 It is observed that the improvement of GIC initiation-VIS and GIC propagation of treated composites over the untreated composites is significant at low MC which is due to increase fibre/matrix interface bonding. Strong interface was characterized by the almost disappearing of fibre bridging in the treated composites compared to untreated composites. In case of high MC, however, there was no improvement interlaminar fracture toughness was observed for the treated composites.

5.3 Suggestions for Future Work

Based on the current findings, the following suggestions for future research are proposed to increase understanding of the effects of MC in flax fibre during composites fabrication:-

 The DMA results of this study suggest that the MC in flax fibre has plasticised the epoxy matrix. Therefore, it would be interesting to study the curing behaviour of flax/epoxy composites with different fibre MC.

 Further investigation could be carried out with different chemical treatments

such as acetylation and potassium permanganate (KMnO4) for which possibly offer better MC reduction. Recently, a hydrophobic treatment

[AHA34]by coating the fibres with metal oxide nanoparticles such as titanium

205

oxide and zinc oxide was proposed and this could be a promising solution [270].

 It was found that the silane treatment reduces the fibre strength and stiffness. Alternative methods can be proposed which is applying silane to the epoxy matrix, rather surface treatment of the fibres [271]. This would be the best solution.

 As flax fibre absorbs moisture, it swells and there is possibility of fibre shrinkage during cure. Hence, further investigation can be carried out to study the internal residual stress of flax/epoxy composites with different MC.

 Permeability studies of flax/epoxy composites in the resin infusion should be assessed in order to study the resin impregnation with moisturized fibre. A series compaction tests can be conducted (i.e dry compaction – before resin filling, wet compaction – after resin filling) to study the consequences of MC in the context of composite processing.

 Qualitative observation of fibre bridging in Mode I test was conducted in this

investigation. Quantitative study [AHA35]between fibre bridging and strength of fibre/matrix interfaces can be further investigated via various statistical, mathematical and computational techniques.

 Further study can be done by using other types of fabric architecture

[AHA36]such as aligned fibre, higher weaving density or hybrid between man- made and natural fibres for which the composites will have greater mechanical properties and fair comparison can be made with current investigation.

 The evaluation of fibre MC in composites during manufacturing also could be explored with the uses of thermoplastic matrices. High density

polyethylene (HDPE) which [AHA37]is known for its hydrophobicity

206

characteristic due to the absence of polar functional groups in polyethylene molecular chains is one of the potential candidates [272].

207

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