Carbon fibre compressive strength and its dependence on structure and morphology

The axial compressive strength of carbon fibres varies with the fibre tensile modulus and precursor material. While the development of tensile modulus and strength in carbon fibres has been the subject of numerous investigations, increasing attention is now being paid to the fibre and the composite compressive strength. In the present investigation, pitch- and PAN-based carbon fibres with wide-ranging moduli and compressive strengths were chosen for a study of fibre structure and morphology. A rayon-based carbon fibre was also included in this study. Structural parameters (L _c, L_a(0), L _a(90), orientation parameter Z, and the spacing between graphitic planes d(00, 2)) were determined from wide angle X-ray spectroscopy (WAXS). Fibre morphology was characterized using high-resolution scanning electron microscopy (HRSEM) of fractured fibre cross-sections. The mechanical properties of the fibres, including compressive strength, the structural parameters from WAXS, and the morphology determined from HRSEM are reported. The influence of structure and morphology on the fibre compressive strength is discussed. This study suggests that the width of the graphitic sheets, the crystallite size perpendicular to the fibre axis (L _c and L _a(0)), and crystal anisotropy play significant roles in accounting for the large differences in compressive strengths of various carbon fibres.     

Compressive strength of coated rigid-rod polymer fibres

One limitation to the use of high-strength/high-modulus rigid-rod polymer fibres like poly-(p-phenylene benzobisthiazole) (PBZT) and poly-(p-phenylene benzobisoxazole) (PBZO) in composite structures is their low compressive strength. Various theories have been developed to predict compressive strength of rigid-rod fibres. In this study the critical buckling stress for rigid-rod fibres with stiff external coatings has been theoretically modelled assuming that the failure mode in compression is the microbuckling of the fibrils in shear. Our model predicts that significant improvement in fibre compressive strength will occur only when relatively thick coatings, with thickness to diameter (t/D) ratios in excess of > 0.05, are used. Experimentally measured compressive strength of aluminium coated PBZT fibres shows values in good agreement to the theory at t/D ratios of 0.006 and below. Factors related to the selection of suitable coating materials and problems associated with establishing coating performance are identified.Nomenclature P axial compressive load - P _f axial compressive load on the fibre - P _c axial compressive load on the coating - P _cr ⁱ critical buckling load in the ith case - _cr critical buckling stress - _co compressive strength of the uncoated fibre - _c compressive strength of the coated fibre - v(x) lateral deflection of a buckled fibril or coating - V _m amplitude of the lateral deflection in the mth mode - m number of half-sine waves in the deflection mode - x coordinate distance along axial direction - y coordinate distance along radial direction - coordinate distance along circumferential direction - l length of the buckling unit - N number of fibrils in the fibre - D fibre diameter - d fibril diameter - t coating thickness - I _f moment of inertia of the fibril - A _f cross-sectional area of the fibril - E _f tensile modulus of the fibre - E _c tensile modulus of the coating material - E tensile modulus of the coated fibre - G torsional shear modulus of the fibre - v_c Poisson's ratio of the coating material - _f density of the fibre - _c density of the coating material - density of the coated fibre - U _f strain-energy change in the fibre - U _c strain-energy change in the coating - T _f external work done on the fibre - T _c external work done on the coating - d/D - t/D     

On the re-inforcement of thermoplastics by imperfectly aligned discontinuous fibres

Studies of deformation behaviour of short fibre reinforced thermoplastics are complicated by the facts that usually a wide range of fibre lengths are present in moulded test pieces and that the fibres are not systematically oriented with respect to any test direction. An equation has been derived for the stress/strain curve of such a material. This has been used to determine fibre/matrix bond strengths in two glass/nylon 6.6 and two glass/polypropylene composites from measured stress/strain curves and fibre length distributions.It is concluded that major improvements in the properties of these materials will only be achieved by modifying processing to retain longer fibres.List of symbols E _c Modulus of composite - E _f Modulus of fibre - E _m Modulus of matrix - V _f Volume fraction of fibre - _c Stress in the composite - _f Peak stress in a fibre - Average stress in a fibre - _uf Ultimate strength of fibres - _m Stress in matrix at fibre failure strain - _uc Ultimate strength of the composite - _c Strain in composite - _uc Ultimate strain of the composite - Shear strength of the fibre matrix bond - L Fibre length - L Critical fibre length at a composite strain e - L _c Critical fibre length for fibre failure - r Fibre radius     

Studies on nickel coated carbon fibres and their composites

Uniform and continuous coating of nickel was given to the carbon fibres by cementation, electroless or electroplating techniques. The coating thickness was ranged between 0.2 and 0.6 m for all the three methods used. Coating thickness less than 0.2 m showed discontinuous coating of nickel over the fibre surface. Beyond 0.6 m thickness, nickel deposited in den-drite form over the continuous coating. For continuously coated fibres, the ultimate tensile properties of electroless coated fibres were near to uncoated carbon fibres suggesting adherent and defect free coating; while fibres coated by electrolytic and cementation process exhibited lower ultimate tensile strength (UTS) properties. The tensile fracture of the cementation coated fibres suggested degradation of the fibres. In composites, prepared by dispersing the coated fibres in pure aluminium matrix, no appreciable fibre-metal interaction was observed. NiAl₃ intermetallics were observed around and adjacent to the carbon fibres. Sometimes carbon fibres were found embedded in massive NiAl₃ intermetallics suggesting that fibre surface can also act as nucleating centre for these precipitates.     

Fragmentation of aramid fibres in single-fibre model composites

Raman spectroscopy has been used to monitor the state of axial stress along fragmented, high-modulus Kevlar 149 aramid fibres in an epoxy resin matrix by monitoring the peak position of the strain-sensitive 1610 cm^–1 aramid Raman band along individual fragments. It is shown that the interfacial shear stress along each fragment, derived from the strain distribution profiles, is not constant as assumed by conventional fragmentation analysis. The fragmentation process of as-received Kevlar 149 fibres is compared to that of irradiated Kevlar 149 fibres exposed to ultraviolet light where the tensile strength and modulus of the fibres have been reduced. It is found that the derived interfacial shear stress and interfacial shear strength values are higher for those fibres exposed to ultraviolet light compared with the as-received fibres. It is also clearly demonstrated that the values of interfacial shear strength calculated at high matrix strains from conventional fragmentation analysis are considerably lower than the maximum value of interfacial shear stress prior to fibre fracture that was found to be close to the shear yield stress of the resin matrix. Hence the determination of the interfacial shear strength following the saturation of the fragmentation process may give rise to misleading results.Nomenclature e _f Fibre strain - e _m Matrix strain - e _f ^max Maximum strain along each fragment - e _f ^* Failure strain of the fibre - E _f Fibre tensile modulus - l _c Critical fragment length - l _c Mean critical fragment length - l _f Fragment length - r Fibre radius - x Distance along the fibre - _f ^max Maximum stress along each fragment - _f ^* Fibre tensile strength - Interfacial shear stress - _s Interfacial shear strength     

Numerical solution to the solidification of aluminium in the presence of various fibres

In an attempt to understand the experimentally observed solidification microstructures in metal matrix composites, the influence of SiC, graphite and alumina fibres on the solidification of aluminium has been studied numerically. Irregular geometries of the composite material were mapped into simple rectangles through numerical conformal mapping techniques to analyse the influence of a single fibre or a row of fibres on a unidirectionally advancing planar solid-liquid interface. The fibres were assumed to be circular in cross-section and the direction of the interface movement was perpendicular to the length of the fibres. The study showed that for fibres with lower thermal conductivity than aluminium, the interface first goes through acceleration as it approaches and ascends the fibre and then deceleration as it descends the fibre. The acceleration and deceleration phenomena of the interface increases as the thermal conductivity ratio of fibre to liquid aluminium decreases. With low thermal conductivity ratios (K _f/K _L1), the interface is orthogonal to the fibre surface. When the conductivity of the fibre is lower than that of the melt, the interface becomes convex facing the fibre; this mode would lead to pushing of the fibre ahead if it was free to move, as has been experimentally observed in cast microstructures of metal matrix composites. The temperature versus solidification time plots of two points, one in the fibre and the other in aluminium, show that the fibre with a conductivity lower than the matrix is at a temperature higher than the melt; the temperature difference between the two points increases with increasing solidification rate for all the positions of the interface before it touches the fibre. The three-fibre study shows that as the number of fibres increases, the curvature of the interface increases upon approaching the subsequent fibres. The relationship between these numerical computations and experimental observations has been discussed.Nomenclature a reference length = diameter of the fibre - h - K thermal conductivity; in Equation 4 it is defined as K = (K + K _f)/2 for the common boundary between fibre and the freezing medium. For all the rest of the points K = K in Equation 4 - L latent heat of fusion - r non-dimensional variable in radial direction - S non-dimensional distance travelled by the interface - Ste Stefan number = - T non-dimensional temperature - t non-dimensional time - x a non-dimensional spatial coordinate of physical plane - y a non-dimensional spatial coordinate of physical plane - thermal diffusivity - non-dimensional axial coordinate of the mapped plane - non-dimensional vertical coordinate of the mapped plane - a polar coordinate - l liquid - m melting - s solid - O constant wall temperature - i initial - f fibre - * dimensional variables     

TEM characterization of some crude or air heat-treated SiC Nicalon fibres

Commercial Nicalon fibres were prepared by thin transverse sectioning and studied by transmission electron microscopy. A progressive tilting of the incident beam allows us to explore the selected-area diffraction (SAD) pattern along two orthogonal directions, increasing the tilting angle (dark-field (DF) imaging). The lattice fringes technique was also used. The samples were Nicalon 001, 101 and 201 fibres, the latter also being studied after heat treatment in air at 1300° C for 48 h. The SAD pattern of the 001 fibre only shows the SiC, 1 1 intense halo whereas the other samples show all the SiC (1 1 1, 2 2 0 and 31 1) strongly scattered beams, indicating a microcrystalline state. Correspondingly, DF imaging does not indicate any localized measurable scattering domain for 001. Only bright dots can be seen, less than 1 nm in size. The other fibres show SiC microcrystals respectively 2 nm (1 01 ), 3 nm (201 ) and up to 7 nm (heat-treated 201) in extent. Free aromatic carbon, shaped in small units less than 1 nm in size fills up the interstices between SiC. These units tend to lie flat on SiC. In heat-treated fibres, they form incomplete layers around the edges. In addition, the heat-treated 2 01 fibre show a 1m thick layer of cristobalite at the fibre surface. These crystals are polytypes.     

Effect of short fibres on critical cut length in tensile failure of rubber vulcanizates

Variation of critical cut lengthI _c in tensile failure of rubber vulcanizates has been studied with respect to the following variables: addition of short silk fibre, fibre concentration and orientation, ageing, reinforcing carbon black filler and elevated temperature. Strain crystallizing rubbers, e.g. natural (NR) and polychloroprene (CR), show higherI _c values than non-strain crystallizing nitrile rubber (NBR). The addition of short fibres was found to cause an increase inI _c in all cases. The increase is more prominent in the case of NBR than for NR and CR. TheI _c values for unfilled NBR vulcanizates are low and a marginal increase is noted on the addition of carbon black. Addition of short fibres leads to a significant improvement in theI _c values, which show a gradual increase with increase in fibre concentration in the composites.I _c exists only in the composites wherein the fibres are oriented along the direction of application of tensile stress rather than across it, and the decrease in tensile strength is marginal at the initial stages but a sharp fall is observed with increasing size of cut lengths. On ageing,I _c values for composites increase while those for unfilled vulcanizates decrease. Critical cut length values for the fibre reinforced composites at a higher temperature (e.g. 100° C) remained unchanged, but dropped in the case of unfilled vulcanizates.     

Crystallization and properties of Li-Al-B-Ti-Zn-silicate system glass-ceramic fibres

Phase separation, nucleation, crystallization and micro-crack extension and their affects on the tensile strength and alkaline resistance of TiO₂ nucleating Li₂O-Al₂O₃-B₂O₃-ZnO-SiO₂ system glass-ceramic fibres are studied by DTA, XRD, TEM and SEM. Phase separation, temperature range of nucleation, T _g, sequence and kinetics of crystallization, sizes of microstructure and surface microcracks, tensile strength and weight loss of alkali corrosion of glass-ceramic fibres are also studied. The mechanism of crystallization and the process of microcrack extension during the preparation of glass-ceramic fibres are discussed in detail. The glass prepared for glass-ceramic fibres should be characterized by the temperature of phase separation, nucleation and crystallization of the glass ought to be low, as near T _g as possible, corresponding to its basic properties and the rate of dense bulk crystallization must be closely controlled. The microstructure of small and concentrated crystallites, about, 25 nm in size while the diameter of the glass-ceramic fibres is 16 m, produced in the glass-ceramic fibres increases its tensile strength and alkaline resistance. A suitable coupling agent covering the surface of the glass-ceramic fibres and tensile stress exerted on them during heat treatment benefit its mechanical and chemical properties.     

Inviscid melt-spun high-temperature alumina-magnesia fibres

Al₂O₃-MgO (AM) fibres containing 98.16 wt% Al₂O₃ and 1.84 wt% MgO, were produced via inviscid melt spinning. By using scanning electron microscopy it was found that the as-spun AM fibres were hollow and their surfaces were very rough. The X-ray diffraction pattern of the as-spun AM fibre showed -Al₂O₃ as a major phase and -Al₂O₃ as a minor phase. The DTA curve of the as-spun AM fibre showed a single endothermic peak representing the phase transformation of -Al₂O₃ to -Al₂O₃. This phase transformation was readily confirmed by analysing the X-ray diffraction pattern of heat-treated AM fibres.     

Fracture of aluminium-coated carbon fibres

A degradation in the ultimate tensile strength (UTS) of aluminium-coated carbon fibres was associated with the formation of a reaction layer of aluminium carbide during annealing treatments 475° C for high tensile fibres (HT) and 550° C for high modulus fibres (HM). It was established that for a given annealing treatment, the UTS depended on the square root of the original coating thickness and proposed that fracture was controlled by cracks in the aluminium carbide, with a specific surface energy () and intrinsic crack length (c ₀) of 2.33 J m^–2 and 30 nm for HT fibres, and of 0.64 to 0.77 J m^–2 and 20 nm for HM fibres.     

High-T c superconducting Bi(Al)-Ca-Sr-Cu-O glass-ceramic fibres drawn from glass preforms

Bi(Al)-Ca-Sr Cu-O glass-ceramic fibres over 100 cm in length were successfully drawn from a glass preform above the crystallization temperature,T _x. The diameters of the uniformly drawn fibres with circular cross-section could be controlled in the range from 25–200 m and the drawing speed was as high as 200 cm min^–1. In this work Al₂O₃ was used to modify the properties of the glass. It increased the glass transition and crystallization temperatures but did not significantly increase the glass working range. Shrinkage and increase of density during heat treatment of the glass fibres were observed. The annealed (825°C/12 h in air) Bi₄Al_0.1Ca₃Sr₃Cu₄O_y glass-ceramic fibre showed aT _c(onset) of 82 K and aT _c(zero) of 71 K.     

Copper coating on carbon fibres and their composites with aluminium matrix

A uniform and continuous coating of copper was given to carbon fibres by cementation or electroless techniques. In both cases, when coating thicknesses were less than 0.2 m, copper deposition was discontinuous over the fibres, and above 0.2 m, coatings were continuous. In electroless coating, about 75% of the continuously coated fibres had a coating thickness range 0.2–0.5 m and above this showed isolated dendrite deposits of copper. In the cementation process, about 75% of the continuously coated fibres had a coating thickness range 0.2–0.6 m, and above this thickness, fine crystallite-type copper deposition was found over smoothly coated copper. The ultimate tensile strength of continuously electroless-coated fibres were nearer to the uncoated fibres, suggesting defect-free coating, while fibres coated by the cementation process exhibited lower ultimate tensile strength values. The tensile fracture of both electroless- and cementation-coated fibres showed delamination of the coating, suggesting poor bonding between coating and the fibre. In composites, prepared by dispersing the coated chopped fibres in a pure aluminium matrix, uniform and random distribution of the fibres were observed without appreciable fibre-metal interaction. The CuAl₂ intermetallics were largely found in the matrix and only very small amounts were observed at fibre/matrix interfaces. Additions of about 2 wt% Mg to the matrix prior to the fibre dispersion did not appreciably change the distribution pattern of the fibres, but in addition to CuAl₂ phase, Mg₂Si phases were observed in the matrix as well as at the interface.     

Electrical resistivity of Si-Ti-C-O fibres after rapid heat treatment

Two types of Si-Ti-C-O fibres were heat treated in a preheated graphite furnace at temperatures between 1273 and 1973 K, and the change in the electrical resistivity was measured after removing the fibres from the furnace. The resistivity of the fibres decreased monotonically with increasing heat-treatment temperature, but showed a significant increase of the order of 10¹–10² in the temperature range of gas evolution from the fibres. The resistivity of the fibre which has an amorphous character began to increase at a lower temperature than that of the fibre with a crystalline character. This increase in resistivity did not occur during heat treatment in a pure oxygen atmosphere, because the oxide layer formed on the fibre surface suppressed gas evolution from the fibres. The X-ray diffraction patterns of heat-treated fibres in nitrogen or oxygen atmospheres revealed that -SiC crystals began to precipitate from the amorphous state as the heat-treatment temperature increased. The -SiC crystal growth, however, did not always correspond with the decrease in the fibre resistivity.     

Compatibility between alumina fibres and aluminium

An investigation is carried out on the interfacial wetting behaviour and reactions between aluminium and alumina fibres (85mass% Al₂O₃ and 15mass% SiO₂). Aluminium is coated onto alumina fibres by a vacuum evaporation technique and the surface of the fully coated fibres and the edge of the partially coated fibres are examined by scanning electron microscope after heat treatments at various temperatures. Within a temperature regime between 943 and 1273 K, occurrence of such interfacial reactions as 4Al(I) + Al₂O₃(s) 3Al₂O₃(g) and 4Al(I) + 3SiO₂(s) 2Al₂O₃(s) + 3Si(s) are detected. It is found that molten aluminium can cover the alumina fibre surface but it peels off near the edge of the coating film on a partially coated fibre, showing the very weak interface cohesion. This is ascribed to the lack of a stable compound formation at the interface. Results of tensile test show that the strength of the coated fibres is degraded after heat-treating at above the melting point of aluminium. The culprits for the tensile failure of alumina fibres are evaluated by the Weibull distribution theory.     

Influence of carbide formation on the strength of carbon fibres on which silicon and titanium have been deposited

Silicon or titanium was deposited on the filaments of carbon fibres by chemical vapour depositions and the reactions between the deposited silicon or titanium and the carbon fibres were investigated below 1300° C. Between the silicon and the carbon fibres, -SiC layers formed at rates of 1.5 to 3 nm in 3 h at 1300° C. These rates were 10^–4 times that of the TiC formation by the reaction of titanium with carbon fibre. Furthermore, the effect of the reaction on fibre strength was investigated. By reaction with silicon, the carbon fibre at a carbonized stage decreased in strength at the beginning of the reaction, but afterwards it recovered to the original level. The carbon fibre at a graphitized stage maintained its original strength after heat treatment for several hours at 1300° C. With the TiC-coated carbon fibres, the carbon fibres decreased in strength following the relation _m d ^–1/2, where d is the thickness of the TiC layer.     

Phase transformation of inviscid melt spun (IMS) alumina-zirconia eutectic fibres

Al₂O₃-ZrO₂ (AZ) eutectic fibres containing 41.05 wt% ZrO₂ and 2.03 wt% Y₂O₃ were produced via inviscid melt spinning (IMS). Scanning electron microscopy (SEM) was used to examine the morphology of the as-spun AZ fibre. Differential thermal analysis (DTA) and X-ray diffraction (XRD) were used to investigate the phase transformations in this fibre. The XRD pattern of the as-spun AZ fibre showed tetragonal ZrO₂, -Al₂O₃, and non-equilibrium -Al₂O₃ phase formation. The DTA curve of the as-spun AZ fibre showed only one endothermic peak representing the phase transformation of -Al₂O₃ to -Al₂O₃. This phase transformation was confirmed by analysing the XRD pattern of heat-treated AZ fibre.     

Effect of fibre concentration and strain rate on mechanical properties of single-gated and double-gated injection-moulded short glass fibre-reinforced polypropylene copolymer composites

The effect of fibre concentration, strain rate and weldline on tensile strength, tensile modulus and fracture toughness of injection-moulded polypropylene copolymer (PPC) reinforced with 10, 20, 30 and 40% by weight short glass fibre was studied. It was found that tensile modulus of single- and double-gated mouldings increased with increasing volume fraction of fibres, ϕ_f, according to additive rule-of-mixtures, and increased linearly with natural logarithm of strain rate . The presence of weldlines in double-gated mouldings led to reduction in tensile modulus which for composite containing 40% by weight short fibres was as much as 30%. A linear dependence was obtained between fibre efficiency parameter for composite modulus and for both single- and double-gated moulding. Tensile strength of single-gated mouldings, σ _c, increased with increasing ϕ_f in a nonlinear manner. However, for ϕ_f in the range 0–12% a simple additive rule-of-mixtures adequately described the variation of σ _c with ϕ_f. A linear dependence was obtained between fibre efficiency parameter for tensile strength and The presence of weldlines in double-gated mouldings reduced tensile strength by as much as 70%. Tensile strength of both single- and double-gated mouldings increased linearly with Fracture toughness of single-gated mouldings increased linearly with increasing ϕ_f. The presence of weldlines in double-gated mouldings reduced fracture toughness by as much as 60% for composite containing 40% by weight short glass fibres.     

High modulus/high strength poly-(p-phenylene benzobisthiazole) fibres

Wide-angle X-ray diffraction studies of poly-(p-phenylene benzobisthiazole) fibres from Part 1 of this work were undertaken to examine fibre structural changes associated with the heat treatment process and which contribute to the observed significant enhancement of mechanical properties. Crystallite size perpendicular to the fibre axis increases from approximately 2 nm in as-spun fibres to 10 to 12 nm in fibres heat treated at temperatures above 600 C. Fibre tensile strength was found to increase with this increase in the extent of the lateral molecular order. However, tensile modulus and tensile strength did not depend directly on heat treatment parameters but rather indirectly through the effect of applied tension during heat treatment on the overall axial orientation. Higher values of fibre tensile modulus and tensile strength were exhibited by the more highly oriented fibres.     

High-temperature compatibility of carbon fibres with nickel

A study has been made of the elevated temperature degradation of a number of carbon fibre types coated with nickel by a variety of methods (electroless, electrolytic, carbonyl and physical vapour deposition). At high temperatures, Ni-coated fibres undergo a transformation of structure to crystalline graphite with a consequent loss of strength and elastic modulus. Resistance to this recrystallization is related to the fibre type and structure and increases in the order HTS PAN-based, HM PAN-based, HM rayon-based. For PAN-based fibres the resistance increases with the degree of structural order and orientation. The recrystallization of HTS fibres is consistent with a simple model of dissolution and reprecipitation controlled by diffusion of carbon in nickel. To explain the higher stability of HM fibres an additional factor must be introduced. For example, their behaviour can be explained in terms of a highly stable surface layer between about 0.1 and 0.5m thick. Rapid recrystallization occurs when the nickel breaks through this layer e.g. by dissolution. The recrystallization was not greatly affected by the type of nickel coating but the recrystallization temperature of HM fibres was considerably reduced by a small proportion of air in the heat-treatment atmosphere. HTS fibres were not affected in this way but the fibres were severely weakened through surface attack by both air and hydrogen at temperatures well below the recrystallization temperature.