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Cracked carbides (non-propagating microcracks) were observed in M2 type high speed steels deformed in four-point bending in the temperature range 17–600°C to ∼0·7% of the yield strain, i.e. to below the stress for macroscopic yielding and ductile or brittle fracture. Although the sintered M2 steel underwent a brittle–ductile transition at ∼300°C, no associated significant change in the fracture mechanism was detected. In all samples studied, cracking of individual carbides (nucleation of failure) was followed by the subcritical linking of these microcracks through the matrix, i.e. growth, and only then by catastrophic propagation, as a Griffith crack from an approximately semielliptical edge (or corner facet) fracture initiating region. Fracture mechanics analyses indicate the depth of this region to be in excess of the Griffith–Irwin flaw size, except for the brittle specimens. The region of subcritical crack growth is compared to the process zone reported for ceramics.

Single crystals of LiF containing voids and of NaCl containing Na2SO4 precipitates were pressurized to introduce dislocations in the vicinities of the discontinuities and subsequently compressed along 〈100〉 at room temperature. The yield stress was raised in both materials; additionally, in LiF discontinuous yielding and easy glide were suppressed and work hardening rate increased by the pressurization-induced dislocations. Following pressurization at 0.85 GN m−2, for example, the 0.1 % shear flow stress of LiF was doubled to ∼ 4 MN m−2 and stage II work hardening rate quadrupled to ∼ 180 MN m−2. Pressurization of NaCl above 0.6 GN m−2 resulted in an increase in the 0.1 % flow stress from ∼ 1.2 to ∼ 2.0 MN m−2. If the slip bands in LiF were initiated by a precompression, pressurization prevented the broadening of these fresh slip bands during subsequent plastic flow. Deformation now took place at a higher stress both in LiF and NaCl. These effects resemble in some ways latent hardening in that oblique as well as conjugate dislocation intersections must take place to continue the deformation. In contrast to latent hardening data, the strain hardening rate was increased in LiF and was approximately proportional to the pressurization-induced dislocation density. This ratio, 5 to 6 dyne per dislocation, is in fair agreement with two sets of independent calculations reported by Gilman and Johnston. The results suggest, therefore, that in the present case also hardening may be due to defects left in the wakes of pressurization-induced moving dislocations.

Initiation and propagation of failure in uniaxially aligned 60% volume-fraction Type III carbon fibre-epoxide compressive specimens, strained parallel to the fibre axis, was studied at atmospheric and superposed hydrostatic pressures, H, extending to 300 MN m−2. The atmospheric axial compressive strength was approximately 1.5 GN m−2 and equal to the tensile strength, but mechanisms involving shear-operated failure of the fibres must be discounted since the failure process was very pressure sensitive above H∼ 150 MN m−2. The results also could not be satisfactorily interpreted by theories involving micro-buckling of individual fibres or laminae when the matrix shear modulus controls the compressive strength. For atmospheric tests and for H

An analysis of theoretical and experimental data relating to non-Newtonian filtration is presented. A filtration equation is proposed for suspensions.

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 tensile, compressive and shear yield strengths of two epoxides were measured under superposed hydrostatic pressure extending to 300 MN m−2. For both materials, the ratio of the moduli of the tensile,σT, to compressive,σC, yield stress at atmospheric pressure was approximately 3∶4, as has been reported previously for a number of thermoplastics. Theσ2=σ3 envelope in stress space was plotted according to these two-parameter (σC andσT) yield criteria: conical, paraboloidal and pyramidal; the best correlation was with the last. The experimental tensile and compressive data for tests under pressure, however, fit slightly better two straight lines which are consistent with a three-parameter single hexagonal pyramidal yield surface. For plane stress and shear under pressure yield envelopes of these surfaces, the correlation with experimental data is again best for the pyramidal criteria, except for biaxial or triaxial tension when these resins are brittle. The third independent parameter employed in the pyramidal criterion was the equi-biaxial compressive yield stress, determined by tensile experiments under appropriate superposed hydrostatic pressure; alternatively plane strain compressive yield stress,σPC, may be used.

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).

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 2.8 GN m−2 and the fracture toughness of the materials was in the range 18 to 25 MN m−3/2. Combining these data, the critical Griffith-Irwin flaw sizes were calculated to be typically of the order of 100 μm. 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 sub-critical cracking of carbides was detected (as in BT42), in contrast to BM2, BT6 and sintered and hot isostatically pressed BT1. 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 μm, cracks under monotonic loading or to associate the brittle fracture stress with that for crack nucleation: e.g. cleavage of a carbide cluster.

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 cold pressing and vacuum sintering characteristics of two batches of BTl high speed steel water atomized powder were investigated. The powder particle sizes, shapes, and microstructures before and after annealing were compared. Relationships were established between compacting pressure and green density, die fill, and linear shrinkage after sintering. Carbon was mixed with both batches before sirttering, 0·3% to batch 1 and 0'15%' to batch 2. The sintered densities of both batches with and without added carbon, are reported for a'range of sintering temperatures between 1230 and 1340°C together with typical sintered microstructures and grain and carbide size 'measurements. Carbon, oxygen, and nitrogen contents were estimated at major steps in the process. The differences found between the sintering characteristics of two batches of the same grade of high speed steel having only slight differences in alloy content were larger than expected. The addition of carbon, however, was found to promote sintering in several ways as well as bringing the final composition up to the specified value: the residual porosity in both batches was eliminated, the temperature range extended over which maximum density was achieved, and the microstructure at the optimum sintering temperature improved. PM/0328

Hardness and compressive strength, Young's modulus, bend strength, and fracture toughness were measured in wrought and commercially sintered T42 high-speed steel. After identical heat treatments, no significant differences were found between wrought and sintered materials in hardness (800–1010 HV50), Young's modulus (∼220 GN m−2), or fracture toughness (9−18 MN m−3/2). The four-point bend strengths in the wrought material were slightly higher in specimens cut longitudinal to the working direction than in transverse specimens (1·3–2·4 GN m−2), and both were markedly superior in strength to the sintered material (1·0–1·2 GN m−2). For three-point bend tests with laboratory-sintered material, the stresses for brittle fracture were in the range 1·9–3·0 GN m−2, a level comparable to the wrought specimens. Wrought material contained carbide stringers with a scattering of large Me carbides, whereas in the sintered material pores and incomplete bonding could be detected, albeit infrequently. Fracture initi...

A comparison was made of the secondary hardening characteristics of a conventionally wrought and two batches of directly sintered BT1 high speed steel. Minor differences in hardness between the types were observed after austenitizing between 1050 and 1280°C and tempering from 400 to 600°C. The recommended hardening temperature for BT1 was considered too high in at least two types by as much as 30 K. The wrought material had the finest microstructure throughout, e.g. 8 μm grain diameter for 1200°C austenitization treatment. Considerable grain refinement, however, was achieved in both directly sintered products by a post-sinter transformation anneal. The grain size after quenching was 2–3 times smaller than for the as-sintered, e.g. 15 and 19 μm compared to 46 and 36 μm respectively for the same austenitization temperature. PM/0331

Changes in microstructure in single crystals of MgO, containing elastic discontinuities, which result from pressurization treatments in the range 0.2 to 2 GN m−2 have been studied using an etch-pitting technique. Complex dislocation arrays have been observed around cubical and spherical cavities and precipitates. The observations are discussed in terms of stress and strain criteria. Production of {100} and {110} cracks at high pressure is also described. Following a 1 GN m−2 pressurization the flow stress was observed to decrease by < 10% and this has been related to the induced dislocation density. In polycrystalline MgO the fracture strength was unaltered by pressurization but the critical temperature difference of thermal shock,ΔTC, was increased by 55K. This has similarly been attributed to the movement of pressurization-induced dislocations.

Examination of both wrought and sintered T6 and M2 high-speed-steel cutting tools used in continuous-turning tests at cutting speeds of 15-60 m min−1 has revealed the presence of microcracking at the tool nose. In wet-cutting conditions, the maximum average steady-state flank-wear rates at cutting speeds of 20–30 m min−1 coincide with the appearance of cracks extending across the matrix through several carbides. At higher cutting speeds only individual carbides were broken, with little extension into the matrix. In dry-cutting conditions the wear land is smaller than in wet cutting, but consists of a rough cracked area. The cracks were present beneath the built-up edge, or just below it on the flank face. The generation of cracks under both wet- and dry-cutting conditions is discussed in terms of cutting speed, built-up-edge stability, and cutting temperature.

Aluminium alloy beams selectively reinforced on one or both surfaces with uniaxially aligned continuous carbon fibre-reinforced plastic were tested in flexure. Three manufacturing routes were employed; two of which involved the use of a room temperature curing adhesive and in the third pre-impregnated tapes were hot-pressed directly onto the alloy. Determined stiffness and strengths of the beams were in good agreement with the predictions of strength of materials theory. A 46% increase in stiffness, from 69 to 101 GNm−2, and a 141% increase in yield strength, from 485 to 1170 MNm−2, were obtained, for example, by incorporating 40% of cfrp by volume to make a sandwich beam. Although smaller increases in strength and stiffness result from less reinforcement, only 10% is sufficient to increase fatigue life by a factor of four.

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.

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.

The effect of temperature on the fracture stress of as-recrystallized cast polycrystalline molybdenum specimens and those containing spark- and slip-induced cracks has been investigated in the range 4·2–230 K at strain rates of ∼ 1·5 × 10−4 s−l. The method of cracking that gave the most reproducible results was a 1% precompression at 77 K. Spark-cracks did not alter the ductile–brittle transition temperature of the material; the effect of the precompression was to lower it by ∼ 14 K. In all cases microscopic plastic deformation preceded failure. The brittle-fracture stress of the spark-cracked specimens decreased from ∼ 660 MN/m2 at 210 K to ∼ 200 MN/m2 at 4·2 K. The fracture (effective surface) energy calculated according to the Griffith–Orowan formula from the 4·2 K test is 5·7 J/m2, compared with estimates of true surface energy using Gilman's model of 2·6–4·5 J/m2. The brittle-fracture stress of precompressed specimens remained approximately constant at 650 ± 50 MN/m2 and was identified with ...

Fracture of polycrystalline tungsten has been studied by two methods: by investigating the tensile properties of bulk material in the temperature range 77°–490°K and directly by stretching thin foils inside the electron microscope. The ductile-brittle transition temperature was found to be ~350°K at a stress level of ~40 kg. mm−2 and the brittle fracture stress to increase markedly with decreasing test temperature. In the temperature range 140°–490°K the metal fails in tension by slip-induced cleavage, whilst at 77°K twinning occurs and is probably responsible for fracture. During the deformation of thin foils inside the electron microscope dislocations have been found to nucleate within the grains and move in both the {110} and the {112} planes. The dislocations generally did not “pile-up” against the grain boundaries but formed arrays of a different configuration. Microcracks were seen to be nucleated within these dense dislocation arrays in the {110} planes only. They propagated in these planes to the grain boundaries and then continued in transcrystalline and/or intergranular fashion. A possible interpretation of the results is given in terms of the observed distribution of dislocations. (The stress field in a typical array is considered in the Appendix.) Although the material and experimental conditions for the two parts of this investigation were quite different, the measured stresses for bulk material in the region of the ductile-brittle transition and the calculated stresses for thin foils, are within 50% of each other.

The change in the ductile-brittle transition temperature of polycrystalline chromium produced by subjecting specimens to a hydrostatic pressure, P, in the range 5–31 kbars and then testing at atmospheric pressure was investigated. The transition temperature, Tt, for the as-recrystallized material was ∼550°K and only treatments above 6 kbars resulted in a significant lowering of Tt. For the pressure range 6–12 kbars the transition temperature was sensitively dependent on P; the 12 kbar treatment reduced Tt to ∼290°K. For these pressurizations Tt occurred at an approximately constant (critical) stress, at the upper yield point and brittle fracture, of 20 kg mm−2. Treatment of 31 kbars resulted in a further decrease in TT of only ∼20°K and the transition took place at a significantly higher stress. For all pressurizations there was a continuity, in the variation with temperature, between the values of the upper yield stress and the brittle fracture stress, which was identified with the crack nucleation stress. In some specimens fractured above TT non-propagating cracks were observed, demonstrating that then crack nucleation is easier than crack propagation. Treatments at 6 kbars and above resulted in decrease and eventual suppression of the yield point drop and lowering of the lower yield stress. A simple analysis indicates that the mobile dislocation density at the upper yield point is raised by an order of magnitude as the pressurization treatment is increased from 5 to 12 kbars.

Tensile properties of “flash-annealed” and conventionally recrystallized niobium were studied at room temperature and at 77°K. Twinning at 77° at stresses lower than the predicted value of the Peierls-Nabarro force proved that it is not this which causes the large temperature dependence on the yield stress. Electron microscopic examination of foils loaded at room temperature to < 3 kg/mm2 showed that edge dislocations are mobile before the apparent proportional limit of the “flash-annealed” niobium, ∼ 3.5 kg/mm2, is reached. At this temperature and 1% min−1 strain rate the material was characterized by an absence of an upper yield point, high work hardening rate and high U.T.S./yield-stress ratio. At 77°K even specimens taken from the same foil did not have reproducible properties due to twinning starting at or continuing to (500 kg/mm2 was recorded) various stress values.

The fracture stress at the ductile-brittle transition temperature of polycrystalline molybdenum was found to be constant, 64± 2 kg mm−2, as the strain rate was raised from 5 × 10−6to 5 × 10−2sec−1. It is suggested that only above the transition temperature is crack nucleation easier than crack propagation and that brittle fracture takes place at the crack nucleation stress. At the transition temperature the stresses for yielding, crack nucleation and crack propagation are thought to be equal. Assuming simple phenomenological correlations between yield stress and temperature, strain rate and grain size, respectively, an analysis of the dependence of the ductile-brittle transition temperature on strain rate, neutron irradiation, pressurization and grain size is presented. This analysis should be applicable to the other Group Via b.c.c. transition metals: chromium and tungsten, which appear to exhibit similar behaviour.

Polycrystalline tensile specimens of molybdenum of average grain diameter 75 mu were irradiated at approximates 40 deg C with an epithermal neutron dose of approximately 2 x 10/sup 18/ nvt. Specimens were annealed for one hour at temperatures of unirradiated specimens, irradiated and annealed specimens, and irradiated but unannealed specimens were tested at room temperature in a tensometer at a strain rate of 0.9 x 10/sup -4/ sec.sup -1/. The effect of the radiation was to increase the yield stress without significantly altering the rest of the load-elongation curve. (C.E.S.)

The initiations of slip, fracture and rupture in tension were studied in sintered and cast molybdenum specimens in the temperature range 77–474°K using optical and scanning electron microscopy. The ductile—brittle transition temperatures, T T , were ~212°K for the sintered material and ~160°K for both coarse and fine-grained cast material. For all batches, at all test temperatures, slip lines were first observed at stresses between 0.35 and 0.50 of the yield stress (in tension, σ Y , or compression, σ Y c ). In cast materials brittle fracture took place at stresses equal to σ Y c , but in the sintered specimens at any stress between the microyield stress, σ M , and σ Y c . Slip lines crossing grain boundaries, cross-slip, slip on more than one system were observed in ductile and brittle, sintered and cast, tensile specimens strained to stresses between σ M and σ Y c . Non-propagating surface microcracks were detected only in ductile specimens at stresses greater than σ Y and only in one sintered specimen at a stress lower than σ Y . The majority of these cracks were intergranular and approximately normal to the tensile axis. In sintered material they were more numerous and were nucleated concurrently with increasing plastic deformation (up to ~30 per cent nominal strain). Although crack nucleation was predominantly intergranular, crack propagation was mainly transgranular in cast specimens and both transgranular and intergranular in sintered specimens. Above the fracture mode transition temperature, in sintered material, voids were formed predominantly at grain boundaries and rupture took place by the pulling out of the fibre-like grains and the shear of the inter-void fibres.

The strain rate sensitivity of the stress at the apparent proportional limit, [sgrave]A, and the variation of dislocation density with strain have been investigated for polycrystalline tungsten deformed at 200°c. The results are consistent with Hahn's (1962) theory of yielding and the dislocation velocities at [sgrave]A evaluate to ∼ 10−4 and ∼ 10−2 cm sec−1 at strain rates of 1 and 50% min−1, respectively.

The ductility transition temperature, TT , of polycrystalline cast chromium, tungsten, and molybdenum was lowered by 280, 230, and > 90 K, respectively, by a 4% precompression at room tempe...

The energy equation is solved for the case of laminar flow of a compressible Newtonian liquid with allowance for heat of friction. Theoretical temperature profiles are presented for flow in channels with constant and linearly varying capillary wall temperatures. The theoretical values of the mean temperature are in good agreement with the experimental results.

Single crystals of NaCl, of known surface dislocation substructures and containing precipitate particles of Na2SO4 approximately 10μm in diameter, were subjected to hydrostatic pressures in the range 1 to 10 kbars and subsequently re-etched at atmsopheric pressure. Rearrangement and nucleation of dislocations in the vicinities of the precipitates was observed to result from pressurisation treatments above 3 kbars. After pressurisation at 8 and 10 kbars the dislocation groups were in general complex; linear arrays on 〈110〉 {110} and “rosettes” were, however, observed. The observations are discussed in terms of recent theories (developed for metal matrices) of generation of dislocations at misfitting particles and are in fair agreement with the model of Ashby and Johnson.

By etch-pitting NaCl polycrystals which had been subjected to a hydrostatic pressure of 8 or 10 kbar it was established that dislocations are generated during the pressurization treatment. The effect of these freshly nucleated dislocations on the stress-strain relationship at atmospheric pressure is established and the relevance of these observations to the interpretations of experiments carried out under pressure is discussed.

A series of experiments on polycrystalline “flash-annealed” niobium are described, which were designed to relate the room temperature yielding and work-hardening properties to purity and microstructure, as observed by transmission electron microscopy. It was found that, for the same low strain rate, whereas the shape of the load-elongation curve appears to depend markedly on purity, the dislocation density and structure does not and that dipoles, tangles and eventually cells are formed for all purities of the material studied. Qualitatively and quantitatively the electron microscope results are very similar to some reported earlier for less pure niobium, for iron and even for face-centred metals, such as copper. The mechanical properties of the purest material resemble those recently found for zone-melted niobium single crystals, namely yield stress of ~ 3.5 kg/mm2 and a region of rapid work hardening. Results are discussed in terms of a model in which yielding is governed by the dislocation multiplication and velocity characteristics in the material. A new and suprising feature of the results is the observation that the strain rate, which apparently does not influence the dislocation density, has a pronounced effect on the dislocation arrangement.

Conversion‐electron Mossbauer spectroscopy, energy‐dispersive x‐ray spectroscopy, and x‐ray diffraction were used for surface studies of the as‐quenched and hydrogenized metallic glass Fe90Zr10. Surface layers of as‐quenched ribbon, paramagnetic at room temperature, were found to be overlaid by a ferromagnetic thin film. The overlayer is composed of α‐Fe microcrystals and a crystalline compound (Fe3Zr). Magnetic α‐Fe microcrystals are about 500 A in diameter and are oriented with the (200) plane parallel to the Fe90Zr10 ribbon plane. Mossbauer spectra show that the magnetic moments of the precipitated small particles lie in the ribbon plane. A hydrogen‐induced magnetic phase transformation, previously reported for bulk Fe90Zr10, was observed in the amorphous surface layers. At the same time, the magnetic crystalline overlayer disappears on the ribbon shiny surface.

Upon charging Fe 90 Zr 10 amorphous ribbons with hydrogen, the Curie temperature increases by more than 70° C. The transition from paramagnetism to ferromagnetism has been investigated by Mossbauer spectroscopy at room temperature. Low-temperature annealing (∼100° C ) is sufficient to release the hydrogen and reverse the transformation.

Molybdenum specimens with differing grain sizes but similar sub-grain sizes were irradiated to an epithermal neutron dose of approximates 2 x 10/sup 18/ nvt and their yield stress measured. The plot of the data versus (2d)/jsup - 1/2/, where d is the grain diameter, shows that the yield stress is not independent of grain size; hence, the hypothesis of sub-boundaries being strengthened by irradiation is not correct. It is noted that the irradiation hardening is approximately the same as that observed previously for a dose ten times as large. (D.L.C.)