Compressive failure and kinking in uniaxially aligned glass-resin composite under superposed hydrostatic pressure

The failure process in uniaxially-aligned 60% fibre volume fraction glass fibre-epoxide compressive specimens strained parallel to the fibre axis was investigated at atmospheric and superposed hydrostatic pressures up to 300 MN m^–2. The atmospheric strength was about 1.15 GN m^–2 (about 20% less than the tensile) and strongly pressure dependent, rising to over 2.2 GN m^–2 at 300 MM m^–2 pressure, i.e. by about 30% per 100 MN m^–2 of superposed pressure. The corresponding figure is 22% if the maximum shear stress and not the maximum principal compressive stress is considered. This is incompatible with atmospheric compressive failure mechanisms controlled by weakly dependent or pressure independent processes, e.g. shear of the fibres. The results also could not be satisfactorily interpreted in terms of microbuckling of individual fibres. Kinking, involving buckling of fibre bundles was proposed as the mechanism of failure propagation, but the critical stage (for this glass reinforced plastic) is suggested as being yielding of the matrix, which initially restrains surface bundles from buckling. A strong pressure dependent failure criterion, about 25% increase per 100 MN m^–2, was derived by modifying the Swift-Piggott analysis of deformation of initially curved fibres. It is postulated that it is the axial compression that causes bundle curvature. Other systems, particularly carbon fibre-reinforced plastic, in which there appears to be a transition in the critical stage of failure from bundle buckling to matrix yielding with increasing superposed pressure, are also considered.     

Fracture of a plasticized epoxide under superposed hydrostatic pressure

The deformation and fracture behaviour of rubber-coated and uncoated epoxy specimens has been studied under superposed hydrostatic pressures extending to 300 MN m^–2. Maximum shear stress at yield for this epoxy were about 25 MN m^–2 at atmospheric pressure and rose to about 48 MN m^–2 at 300 MN m^–2 superposed pressure. Yielding and failure of all specimens tested beyond pressures of 75 MN m^–2 took place when all the (macroscopic) principal stresses, though unequal, were compressive. Fractographic examination revealed three distinct zones of the failure surfaces at atmospheric pressure. The behaviour of all uncoated specimens and those coated and tested below 100 MN m^–2 was similar. A fracture-mechanics interpretation of failure could be applied to these tests assuming the deformation-produced first zone was the fracture initiating site. Coated samples tested beyond 100 MN m^–2 superposed pressure failed with no evidence of Zones II or III of failure; Zone I appeared to spread over the entire failure surface. An interpretation involving fluid penetration of Zone I failure nuclei, along the lines suggested by Duckett, can account for the failure stresses of the uncoated specimens but is not tenable for the coated samples. It appears that crack nucleation and (slow) growth, as opposed, perhaps, to (catastrophic) crack propagation, can take place in this polymer when all the principal stresses are compressive.     

The tensile properties of pultruded GRP tested under superposed hydrostatic pressure

The failure mechanisms in waisted tensile specimens of pultruded 60% volume fraction glass fibre-epoxide were investigated at atmospheric and superposed hydrostatic pressures extending to 350 MN m^–2. The maximum principal stress at fracture decreased from 1.7 GN m^–2 at atmospheric pressure to 1.3 GN m^–2 at 250 MN m^–2 superposed pressure and remained approximately constant at higher pressures, as had been observed with carbon fibre reinforced plastic (CFRP) and a nickel-matrix carbon fibre composite. In the high-pressure region the failure surfaces were fairly flat, consistent with the fracture process being solely controlled by fibre strength. Pre-failure damage, in particular debonding, was initiated at 0.95 GN m^–2 at atmospheric pressure and this stress rose to 1.2 GN m^–2 at 300 MN m^–2 superposed pressure, i.e. by about 9% per 100 MN m^–2. Unlike the pressure dependence in CFRP, this contrasts with the pressure dependence of the resin tensile strength, about 25% per 100 MN m^–2, but can be associated with that of the fibre bundle/resin debonding stress, about 12% per 100 MN m^–2 superposed pressure. Consistent with this interpretation, glass fibres of the failure surfaces were resin-free, again in contrast to CFRP.     

The effect of hydrostatic pressure on the tensile properties of pultruded CFRP

The failure process in waisted tensile specimens of pultruded 60% volume fraction carbon fibre-epoxide was investigated at atmospheric and superposed hydrostatic pressures up to 300 MN m^–2. The maximum principal stress at fracture decreased from ~ 2.0 GN m^–2 at atmospheric pressure to ~ 1.5 GN m^–2 by 200 MN m^–2 superposed pressure and then remained approximately constant. These latter failures were fairly flat and no damage preceding the catastrophic fracture was detected, which indicates that composite strength is solely controlled by fibre strength. Fracture of fibres at lower pressures appeared to commence also in the range 1.5 to 1.6 GN m^–2, but, as it did not result in catastrophic failure, account has to be taken of the resin and the fibre bundles. Debonding was initiated at ~ 1.2 GN m^–2 at atmospheric pressure and this stress increased to ~ 1.5 GN m^–2 when 150 MN m^–2 superposed pressure was applied; the pressure dependence was related to that of the resin tensile strength. This process is described as the first stage, straightening and debond initiation of curved surface bundles, on our model of tensile failure. The second stage, delamination, i.e. the growth of transverse cracks leading to the detachment of these bundles, was impeded by the transverse pressure, being suppressed beyond 150 MN m^–2. Only below this pressure was load redistribution between bundles possible, but, as the pressure was increased from atmospheric, it become more difficult, resulting in a decrease in the composite tensile strength and reduced fibre pull-out.     

Transverse (interlaminar) cracking under tensile loading in pultruded CFRP

The examination of microstructure of tensile specimens of pultruded 60% V _f carbon fibre-reinforced epoxide of up to 6 mm unreduced diameter shows that transverse cracking precedes the tensile failure of groups of fibres. In material whose strength is 2 GN m^–2, the process can commence in waisted specimens at stresses as low as 1 GN m^–2; in those of unreduced section it was not detected below 1.5 GN m^–2. This failure initiation stage can be associated with the decrease in the slope of the load-extension curve. With increasing load the inter-tow cracks were observed to grow and some surface fibre bundles detached. It is suggested that misaligned fibres in these surface bundles were straightened out and contributed to the load-carrying capacity of the rod. Only following detachment of numerous bundles (for the specimens with unreduced section) or growth of interlaminar cracks into the specimen shoulders (for those with a reduced gauge diameter) did tensile failure of fibre bundles lead to catastrophic fracture. It is to this last propagation stage that statistical models of failure of bundles at different cross-sections should refer.     

Mechanical behaviour of bamboo and bamboo composite

The tensile, flexural and impact strengths of bamboo and bamboo fibre-reinforced plastic (BFRP) composite have been evaluated. The high strengths of bamboo, in the fibre direction, have been explained by its anatomical properties and ultra structure. Bamboo fibres and bamboo orthogonal strip mats (bamboo mat) have been used to reinforce epoxy resin. BFRP composites with unidirectional, bidirectional and multidirectional strengths have been made. In bamboo mat composites, the fibre volume fraction,V _f, achieved was as high as 65%. The tensile, flexural and impact strengths of bamboo along the fibres are 200.5 MN m^–2, 230.09 MN m^–2 and 63.54 kJ m^–2, respectively, whereas those of bamboo fibre composites and bamboo mat composites are 175.27 M N m^–2, 151.83 MN m^–2 and 45.6 kJ m^–2, and 110.5 MN m^–2, 93.6 M N m^–2 and 34.03 kJ m^–2, respectively. These composites possess a close to linear elastic behaviour. Scanning electron microscopy studies of the fractured BFRP composite specimens reveal a perfect bonding between bamboo fibres and the epoxy. Furthermore, high strength, low density, low production cost and ease of manufacturing make BFRP composite a commercially viable material for structural applications.     

Jute-reinforced polyester composites

Raw jute fibre has been incorporated in a polyester resin matrix to form uniaxially reinforced composites containing up to 60 vol% fibre. The tensile strength and Young's modulus, work of fracture determined by Charpy impact and inter-laminar shear strength have been measured as a function of fibre volume fraction. These properties all follow a Rule of Mixtures relationship with the volume fraction of jute. Derived fibre strength and Young's modulus were calculated as 442 MN m^–2 and 55.5 GN m^–2 respectively. Polyester resin forms an intimate bond with jute fibres up to a volume fraction of 0.6, above which the quantity of resin is insufficient to wet fibres completely. At this volume fraction the Young's modulus of the composite is approximately 35 GN m^–2, the tensile strength is 250 MN m^–2, the work of fracture is 22 kJ m^–2 and the inter-laminar shear strength is 24 MN m^–2. The properties of jute and glass fibres are compared, and on a weight and cost basis jute fibres are seen in many respects to be superior to glass fibres as a composite reinforcement.     

Structure and properties of some vegetable fibres

The stress-strain curves for pineapple leaf fibre have been analysed. Ultimate tensile strength (UTS), initial modulus (YM), average modulus (AM) and elongation of fibres have been calculated as functions of fibre diameter test length and test speed. UTS, YM, and elongation lie in the range of 362 to 748 MN m^–2, 25 to 36 GN m^–2, and 2.0 to 2.8%, respectively for fibres of diameters ranging from 45 to 205m. UTS Was found to decrease with increasing test lengths in the range 15 to 65 mm. Various mechanical parameters show marginal changes with change in speed of testing in the range of 1 to 50 mm min^–1. The above results are explained on the basis of structural variables of the fibre. Scanning electron microscope studies of the fibres reveal that the failure of the fibres is mainly due to large defect content of the fibre bo1h along the fibre and through the cross-section, The crack is always initiated by the defective cells and further aggravated by the weak bonding material between the cells.     

Structure and properties of some vegetable fibres

The stress-strain curve for sisal fibres has been experimentally determined. Ultimate tensile strength (UTS), initial modulus (YM), average modulus (AM) and per cent elongation at break of fibres have been measured as function of fibre diameter, test length and test speed. UTS, YM, AM and per cent elongation lie in the range 530 to 630 MN m^–2, 17 to 22 GN m^–2, 9.8 to 16.5 GN m^–2 and 3.64 to 5.12 respectively for fibres of diameters ranging between 100 and 300m. No significant variation of mechanical properties with change in diameter of the fibres was observed. However, with increase in test length of the fibres, the UTS and per cent elongation are found to decrease while YM and AM increased in the test length ranging from 15 to 65 mm. With the increase in speed of testing from 1 to 50 mm min^–1, YM and UTS are found to increase whereas per cent elongation and AM do not show any significant variation. At a test speed of 500 mm min^–1 the UTS value decreases sharply. The above results are explained in terms of the internal structure of the fibre such as the cell structure, microfibrillar angle, defects, etc. Scanning electron microscope (SEM) studies of the fractured tips of the sisal fibres reveal that the failure of the fibre is due to the uncoiling of microfibrils accompanied by decohesion and finally tearing of cell walls. The tendency of uncoiling seems to decrease with increasing speed of testing.     

Tensile properties of directionally solidified Al-Al3Ni composites with off-eutectic compositions

Ingots (1/8 in. diameter) of Al-Al₃Ni with off-eutectic compositions were directionally solidified at two growth rates. At 1 cm h^–1 fibres exhibited a blade-like morphology, and at 11 cm h^–1, a rod-like morphology. Speciments were mechanically evaluated in tension. The composite modulus at stresses above the yield strength of the aluminium matrix obeyed the rule of mixtures, assuming ideal plastic behaviour of the matrix. An extrapolation of these data for composites with rod fibres gave a value of 146 GN m^–2 for the modulus of Al₃Ni. Tensile strength of composites with rod-like fibres followed the rule of mixtures, whereas those with non-uniform blade fibres showed a lower strength. In composites with blade-like fibres the extrapolated aluminium matrix strength was 86.3 MN m^–2, a high value attributed to dispersion hardening, and that of the fibres was 2.69 GN m^–2. Composites with blade-like fibres failed at lower strains than did those with more uniform rod-like fibres.     

Microhardness and brittle fracture of garnet single crystals

Hardness and fracture toughness were measured using the Vickers microhardness test in the low load range from 25 to 100 g near to the fracture threshold for near-perfect single crystals of garnets. The influence of crystal growth parameters, calcium impurity content and crystallographic orientation of Gd₃Ga₅O₁₂ (GGG) and Ca₃Ga₂Ge₃O₁₂ (CaGeGG) samples was investigated. Fracture starts with radial cracking from indent corners followed by lateral fracture of two distinct modes. The mean hardness of [111] oriented GGG isH=13 GN m^–2, for [111] oriented CaGeGG it is 12 GN m^–2, the average fracture toughness beingK _c=1.2 and 0.8 MN m^–3/2, respectively for the two crystals. Impurity doping slightly increases the strength of the material. Among the investigated crystals (111) faces are the least strong, the (100) face has maximumH andK _c values for CaGeGG. The constraint factor,, and yield stress,Y, were deduced from the measured hardness data, giving=2.2 andY about 7 GN m^–2.     

Longitudinal shear properties of human compact bone and its constituents,and the associated failure mechanisms

Human compact bone specimens were tested in longitudinal shear at two different strain rates. The maximum stress and energy absorption capacities were 50.40±14.08 MN m^–2 and 20720±9310J m^–2 respectively for 14 embalmed specimens tested at a cross head speed of 2.1×10^–6 m sec^–1. The maximum stress was found to be 75% of the transverse shearing strength. Bone specimens were also tested after selectively dissolving the organic and mineral components. The results showed that the composite strength of bone was much higher than the summation of strengths of its organic and mineral phases. Fractographic examination of the fracture surfaces showed that debonding of the interfaces between the osteons and the surrounding bone matrix and between the osteonal lamellae were the main mechanisms of longitudinal shear failure.     

Synthesis and fabrication of β-tricalcium phosphate (whitlockite) ceramics for potential prosthetic applications

A novel process is described for preparing dense, polycrystalline tricalcium phosphates. Single-phase compositional integrity is achieved by introducing catalytic amounts of sulphate ion and this pore free material has close to theoretical density. Preliminary mechanical properties include a compression strength of 687 MN m^–2 and a tensile strength of 154 MN m^–2. The relationship between processing variables and phase composition, microstructure, strength and translucence is described. The material has promise for bone implant applications.     

Preparation and properties of cast aluminium-ceramic particle composites

A casting technique for preparing aluminium-alumina, aluminium-illite and aluminium-silicon carbide particle composites has been developed. The method essentially consists of stirring uncoated but suitably heat-treated ceramic particles of sizes varying from 10 to 200 m in molten aluminium alloys (above their liquidus temperature) using the vortex method of dispersion of particles, followed by casting of the composite melts. Recoveries and microscopic distribution of variously pretreated ceramic particles in the castings have been reported. Mechanical properties and wear of these composites have been investigated. Ultimate tensile strength (UTS) and hardness of aluminium increased from 75.50 MN m^–2 and 27 Brinell hardness number (BHN) to 93.15 MN m^–2 and 37 BHN respectively due to additions of 3 wt % alumina particles of 100 m size. As a contrast, the tensile strength of aluminium-11.8 wt % Si alloy decreased from 156.89 MN m^–2 to 122.57 MN m^–2 due to the addition of 3 wt % alumina particles of the same size. Adhesive wear rates of aluminium, aluminium-11.8 wt % Si and aluminium-16 wt % Si alloys decreased from 3.62×10^–8, 1.75×10^–8 and 1.59×10^–8 cm³ cm^–1 to 2.0×10^–8, 0.87×10^–8 and 0.70×10^–8 cm³ cm^–1, respectively, due to the additions of 3 wt % alumina particles.Formerly with the Department of Mechanical Engineering, Indian Institute of Science, Bangalore 560 012, India.     

The mechanical properties of ceramics from bauxite waste

This paper summarizes the development of engineering ceramics made from bauxite waste (red mud) produced in the alumina industry in Jamaica. Test specimens are fabricated from powders by sintering. For a particle size distribution of less than 75m in the starting powders and firing temperatures in the range of 1000 to 1100° C various mechanical properties are measured. These include fracture toughness (K _Ic), modulus of rupture (MOR), compression strength ( _c) and Brinell hardness. While apparent porosity varies between 40 and 48%, K _Ic is found to vary between 0.39 and 0.68 MN m^–3/2. The values are compared with those measured for commercial ceramic tiles and also with ceramics of magnesia and calcium zirconate. Within the fabrication temperature range studied the MOR ranges between 17.23 and 27.09 MN m^–2, compressive strength between 42.0 and 83.9 MN m^–2 and Brinell hardness between 26.2 and 59.9 kg mm^–2. With the aid of scanning electron microscopy and a basic knowledge of the physicochemical properties of the mud an attempt is made to explain the high strength and toughness achieved. The ready availability of raw material and the relatively low firing temperatures required for suitable engineering products should keep the production costs low for red mud ceramics.     

Fabrication of cellular structure composite material from recycled soda-lime glass and phlogopite mica powders

A homogeneous composite material with different physical structures has been fabricated from recycled colourless soda-lime glass powders and phlogopite-type mica powders by mixing the two powder components and sintering the mixture at a temperature above 850° C for a period of 30 min or longer. The physical structure of the composite material can be fabricated into either a cellular structure consisting of both closed and open cells or a highly densified ceramic body. The cellular structure composite material is found to have a compressive strength of about 0.877 MN m^–2 and thermal conductivity values in the range of 0.290 to 0.306 W m^–1 °C^–1 when measured at temperatures in the range of 25 to 100° C. The highly densified composite material, on the other hand, is found to have a compressive strength of about 53.0 MN m^–2 and thermal conductivity values in the range of 0.198 to 0.250 W m^–1 °C^–1. The composite material, when compared with other common building materials, is found to be potential material for construction applications because of its superior mechanical and thermal properties.     

Impact-modified epoxy resin with glassy second component

A resorcinol-based epoxy resin was modified by incorporating a glassy second component. The mixture showed a heterogeneous morphology with two clearly defined phases, one phase rich in oligomer, the other phase composed mainly of resorcinol epoxy resin. The fracture toughness measured asG _1c andK _1c values showed an increase from 174J m^–2 and 0.89 MN m^–1.5 S in pure epoxy resin to 431 J m^–2 and 1.36 MN m^–1.5 in 30% oligomer modified resins. The scanning electron micrographs showed that the oligomer-rich phase exhibited ductile failure behaviour and formed the dispersed phase at low concentrations while it was the continuous matrix when the concentration was 30%. Optical observations on the failure mode of thin films of the oligomer-modified epoxy resin showed the existence of both inter face failure and considerable distortion in both phase.     

Axial compression fracture in carbon fibres

Axial compression fracture of carbon fibres was studied by embedding single fibres in epoxy resin and compressing the specimens parallel to the fibre axis. By careful optical monitoring of the fibre surface the earliest stages of fracture were identified leading to estimates of the fibre axial compression failure strengths. Compression strength decreases markedly from about 2.2 GN m^?2 for moderately oriented fibres to <1 GN m^?2 for highest modulus filaments. The trend towards decreasing compression strength with increasing anisotropy is explained on the basis of an increasing fibre microfibrillar nature. However fracture morphology studies show that the unduly rapid strength decrease results from an increasing degree of fibre outer layer ordering which accompanies increasing axial anisotropy in carbon fibres since cracking occurs first on the more highly aligned filament surfaces. It is suggested that fibre compression fracture changes from a shear to a microbuckling or kinking mode with increasing fibre anisotropy, where the latter initiates in individual, well-aligned but uncoupled microfibrils. The similarity of fine axial compression fractures in oriented carbon fibres to those found in elastica loop experiments is noted as are the possible implications which the low strain-to-failure in compression of very high modulus fibres might have for practical composites.     

Kinking and tensile,compressive and interlaminar shear failure in carbon-fibre-reinforced plastic beams tested in flexure

The first stage of the failure process in pultruded, 60% volume fraction, type III carbon fibre-epoxide beam specimens with span-to-depth ratios of 5, 15 and 40 deformed in flexure at atmospheric pressure was the initiation of kinking by the compression roller. Kink growth during the non-linear part of the load-deflection curve was followed by kink propagation at peak load. Acoustic emission and load-unload tests to detect irrecoverable deflection supported direct microscopic observations of damage. Kink growth with decreasing load, increasing deflection and accompanying redistribution of stresses led to two types of failure, commonly referred to as flexural and interlaminar. In the former, tensile failure was concurrently initiated to give the characteristic tensile and compressive zones on the failure surfaces. In the latter, the growing kink initiated interlaminar cracks in resin-rich zones as it propagated (with decreasing load) towards the convex surface. Kinking was associated with triaxial compressive stresses in the contact zone of the compressive roller or rollers (in the case of four-point bend specimens). When hydrostatic pressure was superposed on flexure, at pressures between 150 and 300 MNm^–2 depending on the type of specimen, kinking was inhibited and eventually suppressed to give tensile failures, even in the so-called interlaminar shear strength type of specimen. When non-linear deflections were not large, the maximum principal tensile stress in the beams was close to the tensile strength of the carbon-fibre-reinforced plastic (1.8 GN m^–2).     

Load and directional effects on microhardness and estimation of toughness and brittleness for flux-grown LaBO3 crystals

Results of microhardness measurements on (100) and (110) planes of flux-grown LaBO₃ crystals, in the applied load range of 10–100g, are presented. The microhardness was found to decrease with increasing load in a non-linear manner. By applying Hays and Kendall's law, the materials resistance pressure and other constants of the equation could be calculated. Hardness anisotropy, showing periodic variation of H _v with the maxima and minima repeating at every 15° change in orientation of the indentor, is described and discussed. H _max/H_min are estimated as 1.14 and 1.06 for (100) and (110) planes, respectively. The fracture toughness values, K _c, determined from measurements of crack lengths, are estimated to be 1.6, 1.7 MN m^–3/2 (for (100) planes) and 1.2, 1.5 MN m^–3/2 (for (110) planes) at 90 and 100g loads, respectively. The brittleness index, B _i, is estimated as 4.6, 4.0 m^–1/2 (for (100) planes) 6.0, 4.6 m^–1/2 (for (110) planes) at 90 and 100g, loads respectively.