Powder Metallurgy

Liquid phase sintered and spheroidised Fe-1.4C-0.65Si-0.85Mo specimens were warm forged to discs at 750 and 700°C. The former experiments were conducted on a screw press and the latter, using a Gleeble instrumented tester, at strain rates of 10 -3 , 10 -2 , 10 -1 and 1 s -1 to ~ 1.15 natural strain. Ferrite grain size of the spheroidised PM steel, ~30 μm, diminished as a result of the forging to 6-7 μm, with a fine distribution of sub-micron carbides. The discs were tested in diametral compression and a procedure is presented for the determination, in this testing geometry, of the (compressive) yield strength. These values, above 740 MPa, compare favourably with 350-410 MPa, determined directly in tension, for the as-spheroidised material.

Dilatometry, gravimetric and differential thermal analyses were used to study the sintering response of Al at 600C in atmospheres of pure nitrogen, nitrogen -hydrogen and argon. Shrinkage took place only when nitrogen was present; the best results were obtained in pure nitrogen. Accordingly attempts were made to interpret its positive role utilising, inter alia, theoretical phase diagrams for alloys of Al, O and N. It is proposed that, as the specific volume ratio Al 2 O 3 / Al is ~1.2, oxide layer on Al particles cracks during heating, exposing areas of Al to the atmosphere, thereby allowing self-gettering of oxygen. This reduces the oxygen partial pressure within the connected pores of the compact, in turn allowing oxide transformation to non-stoichiometric oxy-nitride(s) implying presence of atomic nitrogen. Creation of clean Al alloy surfaces within the pores can allow nitrogen gas to react with Al to produce AlN and also to go into solution, which reduces the melting point at surfaces of powder particles, which adjoin the pores. The surface liquid will promote further disruption of any remaining oxide film and keep oxygen partial pressure low, allowing the nitrogen gas to further react with Al. Metallographic and EDX examinations are in accord with the model now postulated.

Metallurgy Fe-1.5Cr-0.2Mo-0.7C steel specimen fatigued in bending with R = −1 at 24 Hz and a stress amplitude of 312 MPa. The fatigue limit was ∼240 MPa, at which stress level no microcracks were detected in static loading. Testing was interrupted at 100 cycles and at further 29 intervals until failure after 49 900 cycles. For each arrest, surface replicas were made in the two regions where maximum stress was applied. Microcracks could nucleate below 100 cycles, when their sizes ranged from <5 to ∼20 μm. Fractographic examination identified the failure-originating site, which was then associated with the crack system observed on the 'last' pre-failure micrograph. Detailed examination of the eventual failure region showed nucleation, at various cycle intervals, of 18 microcracks, their subcritical growths, arrests and coalescences with continuing cycling to form a critical crack 2.25 mm deep. Stepwise microcrack growth was probably rapid -to the next arrest or coalescence. For each (micro)crack size stress intensity factors, K a s, were estimated and, at the end of Stage II, for the coalesced crack, K a reached K 1C , independently estimated to be ∼36 MPa m 1/2 .

Thermo-Calc modelling was employed to predict liquid phase amounts for Fe-0?85Mo-(0?4-0?6)Si-(1?2-1?4)C in the temperature range of 1285-1300uC and such powder mixes were pressed and liquid phase sintered. In high C steels, carbide networks form at the prior particle boundaries, leading to brittleness, unless the steel is heat treated. To assist the break-up of these continuous carbide networks, 0?4-0?6% silicon, in the form of silicon carbide, was added. After solution of processing problems associated with the formation of CO gas in the early part of the sintering cycle, and hence large porosity, densities in excess of 7?75 g cc 21 were attained. A spheroidising treatment resulted in microstructures having the potential of producing components, which are both tough and suitable for sizing to improve dimensional tolerance. Yield strengths up to 410 MPa, fracture strengths up to 950 MPa and strains up to 16% were attained.

Liquid phase sintered and spheroidised Fe-1.4C-0.65Si-0.85Mo specimens were warm forged to discs at 750 and 700°C. The former experiments were conducted on a screw press and the latter, using a Gleeble instrumented tester, at strain rates of 10 -3 , 10 -2 , 10 -1 and 1 s -1 to ~ 1.15 natural strain. Ferrite grain size of the spheroidised PM steel, ~30 μm, diminished as a result of the forging to 6-7 μm, with a fine distribution of sub-micron carbides. The discs were tested in diametral compression and a procedure is presented for the determination, in this testing geometry, of the (compressive) yield strength. These values, above 740 MPa, compare favourably with 350-410 MPa, determined directly in tension, for the as-spheroidised material.

Fatigue properties of PM Fe-1.5Cr-0.2Mo-0.7C steel were determined in bending at 24Hz with R=-1 and the fatigue limit found to be ~ 240 MPa, at which stress level no microcracks were detected in static loading. Microcracks, whose sizes ranged from <5 to ~20 microns, could nucleate below 100 cycles. Stress amplitudes, S, of 260 MPa was selected to investigate the progressing failure mechanism; specimens fatigued for 1000, 5000, 15000, 24000 and 79300 [failure] cycles were examined in detail by replica microscopy in the regions where maximum stress was applied. For each specimen, fractographic examination identified the failure-originating site, which could be associated with the crack system[s] observed on the "last" pre-failure micrograph. Detailed examination showed microcrack nucleation, subcritical growth and coalescence with continuing cycling to 0.5 mm size. This subcritical crack extended conventionally to the critical size of ~ 3mm, when the then estimated stress intensity factor, K a , appeared to reach K 1C ,, independently estimated to be ~36 MPa.m 1/2 and thus the presented crack growth sequence illustrates good correlation between experimental data and the theoretical analysis.

Nucleation of microcracks, their growth and coalescence are analysed in powder metallurgy (PM). Fe-1 . 5Cr-0 . 2Mo-0 . 7C steel by fractography allied to surface replica microscopy -at several stress levels as the maximum tensile stress in three-point bend specimens was raised to 99 . 6% of the transverse rupture strength TRS of 1397 MPa. The fatigue limit in this material is y240 MPa, at which stress level no microcracks were detected in static loading. Numerous microcracks, ranging in size from ,5 to y20 mm, however, were nucleated above ,800 MPa, i.e. beyond the yield strength of ,620 MPa. With increasing stress, some microcracks became dormant, whilst others grew subcritically, stress step-wise, to some 400 mm. Of particular importance are observations of the coalescence of two and three of such microcrack systems to produce a critical, propagating crack. The then estimated stress intensity factor K a , could reach K 1C , independently estimated to be ,36 MPa m 1/2 . Microcrack coalescence was associated with easy paths for crack growth, principally prior particle boundaries linking pores. Ways of making subcritical crack growth more difficult and hence improving both static and dynamic mechanical properties, are considered.

Mn was introduced, with graphite, as fine ferromanganese particles to form compacts of Fe-3%Mn-0?5%C. Each was sintered in dry 25%H 2 z75%N 2 for 3 min at 770, 1040, 1080, 1170 and 1220uC respectively, held for 3 min and quenched. The surface of a successively heated specimen was similarly examined. Specimens were sectioned and, especially the reacting ferromanganese particles and adjoining regions, investigated using light and scanning electron microscopy and EDX. Development of microstructure and microcompositions during sintering was related, from 740uC, to diffusion and condensation of Mn vapour on iron particle surfaces and subsequent chemical reactions. Above 1080uC microstructures included features resulting from a transient liquid phase, in accord with a y45 wt%Mn ternary eutectic calculated by ThermoCalc. The thicknesses of the highly Mn enriched regions were substantially higher than those resulting from bulk Mn diffusion in the Fe lattice; our interpretation invokes predominant operation of another type of mechanism: diffusion induced grain boundary migration.

Thermo-Calc modelling was employed to predict liquid phase amounts for Fe-0?85Mo-(0?4-0?6)Si-(1?2-1?4)C in the temperature range of 1285-1300uC and such powder mixes were pressed and liquid phase sintered. In high C steels, carbide networks form at the prior particle boundaries, leading to brittleness, unless the steel is heat treated. To assist the break-up of these continuous carbide networks, 0?4-0?6% silicon, in the form of silicon carbide, was added. After solution of processing problems associated with the formation of CO gas in the early part of the sintering cycle, and hence large porosity, densities in excess of 7?75 g cc 21 were attained. A spheroidising treatment resulted in microstructures having the potential of producing components, which are both tough and suitable for sizing to improve dimensional tolerance. Yield strengths up to 410 MPa, fracture strengths up to 950 MPa and strains up to 16% were attained.

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 adhesive fracture energy (G/c) of several adhesive cement systems has been measured at temperatures extending to 700°C using alumina adherends in the double torsion test geometry. Two commercially available low-temperature curing systems were evaluated together with three laboratory-formulated cements, two of which were cured at ~ IO00°C. The room temperature values of Gic ranged from 0.6 to 4.9 J m -2 for the former and from 6. 7 to 11.4 J m -2 for the latter. Despite the adoption of a linear compliance test geometry it was essential to precrack the specimens in order to obtain "valid" Gic data. Generally the fracture energy of the adhesive cement joints increased with temperature up to a peak between 400 and 600°C, and then decreased. The maximum GIc value recorded was 40.2 J m -2. The increase in toughness can be associated with viscous effects in the glassy phases present in the adhesive cements.

The axial compressive strengths of three 0.52 V~ glass-fibre/polyester pultrusions were determined under superposed hydrostatic pressures extending to 300 MPa. Atmospheric strengths were in the range 380-790MPa, and all materials showed a strong, about 3, pressure-dependence. Kinking of fibre bundles or sub-bundles preceded fracture in all materials at all pressures, but failure to detect non-propagating transverse cracks, which allow such groups of fibres to act in unison, indicates that bundle-debonding controlled the critical mechanism of failure. The effects of shape, size, and eccentricity of the bundles on buckling and debonding are considered. It is tentatively concluded that the critical parameter is the transverse tensile strain of the resin, 61, causing bundles (rather than individual fibres) to debond. The pressuredependence of the compressive strength is well predicted by the E1 criterion. To obtain fair correlations with the measured values, the bundle-bundle decohesion strengths would have to be to the order of 223 MPa, rather than approximately 50-100 MPa estimated for the fibre-fibre debond strengths.

The tensile properties of five types of pultrt~ded 0"52 Vf glass fibre-polyester rods were #westigated by extending waisted round specimens at atmospheric and superimposed hydrostatic pressures, -H. to 300 MPa. The maximum principal stress at fracture, -700 MPa, decreased, with the superimposition of -H, approximately by its magnitude. As -H increased the failure surfaces became flatter, the amount o f fibre pull-out decreased and transt'erse cracks became shorter or were eliminated. Glass fibres in the failure surfaces were resin free, and failure of the glass fibre bundles appeared to control the fracture process in the entire pressure range for all materials. The decrease in maximum principal tensile stress with increasing -H indicates that the glass .fibre failure process is not controlled by a critical tensile stress. Failure criteria are discussed, and in the tension-compression-compression octant of stress space the relevant criteria appear to be strain energy and deviatoric tensile stress, strain and strain energy for these GRPs and glass itself.

Cast, wrought, and directly sintered smooth and precracked beam specimens of BT1 steels were studied in three-and four-point bending at room temperature. Following austenitization at 1250 ° C and tempering between 500 and 560 ° C, brittle fracture strengths varied between 1.1 and 2.8GN m -~ and the fracture toughness of the materials was in the range 18 to 25 MN m -3/2. Combining these data, the critical Griffith-lrwin flaw sizes were calculated to be typically of the order of 100 #m. This is in reasonable agreement with the observed sizes of some failure-initiating sites, particularly pores in sintered material, but generally several times larger than the carbide and grain sizes. In wrought specimens, failure frequently originated from groups of carbides, apparently fracturing on contiguous planes. No evidence of subcritical cracking of carbides was detected (as in BT42), in contrast to BM2, BT6 and sintered and hot isostatically pressed BTI. Only inter-powder particle parting occurred in this sintered material. Conventional fracture mechanics thus successfully interprets results on sintered specimens, but only on several of the wrought specimens. For the majority of the latter it appears necessary to invoke operation of propagation mechanisms involving "short", --~ 10#m, cracks under monotonic loading or to associate the brittle fracture stress with that for crack nucleation: e.g. cleavage of a carbide cluster.

The tensile and compressive strengths of three polyester resins were measured under superposed hydrostatic pressure extending to 300 MPa, in an attempt to establish yield criteria. The polyesters were brittle in uniaxial tension at all pressures, and accordingly, a third testing geometry, diametral compression of a disc, was employed to complete the two or three necessary parameters to establish the yield surface in stress space. From the biaxial (disc) and axial compressive test data, the atmospheric tensile yield strength (higher than the fracture strength) was computed to be ∼67 MPa in comparison with the compressive strength of ∼120 MPa, their ratio 0.56 being significantly less than the more common 0.75 found for thermoplastics and epoxides. The data for compressive yield strength under superposed pressure were compared with the predictions of the two-parameter pyramidal, conical and paraboloidal criteria and the fit, though reasonable for the latter, could be significantly improved if a further independent material parameter was employed to give a three-parameter pyramidal criterion (the principal stresses σ1, σ2 and σ3 being measured in MPa) of the form 0.0150σ1−0.0039σ2−0.0083σ3=1

A methodology, involving fretting tests, to develop wear and crack resistant materials for tribological applications for automotive valve train parts (e.g. cams, tappets) has been recently reported for high speed steels. Modifications to one of these sintered steels, M3 Class 2, were effected by additions, singly and in combination, of 5 wt.% of wear resistant titanium carbide and of solid lubricant manganese sulphide. In our fretting tests alternate displacements were imposed between the test material (plane) and a chromium steel or alumina ball. Running conditions fretting and material response fretting maps were constructed for the four materials. Two types of fretting damage were detected and analysed: cracking or particle detachment and wear through the tribologicaly transformed structure (TTS). Crack initiation, associated with porosity and interfaces, was detected when the maximum tensile stress in the contact reached 1.2 GPa. Cracking analyses were also carried out using static and fatigue mechanical tests and replica scanning electron microscopy. Crack growth and propagation were influenced by details of the microstructure, e.g. TiC was observed to arrest crack growth, whereas MnS made it easier. Wear analysis included the determination after each test of the wear volume, which could be related to the coefficient of friction and the cumulative dissipation energy during the fretting test.