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By F. L. Vogel

Densities and distributions of dislocations in plastically bent germanium crystals before and after annealing were studied. In the bent and annealed crystals, the theoretical relationship between radius of curvature and density of dislocations is confirmed. Before annealing, however, more dislocations are present than required, and these are distributed with a minimum at the neutral axis and maxima at the top and bottom surfaces. On annealing, three significant changes occur in the bent bars: 1) the average dislocation density is reduced, presumably by the annihilation of opposite signs; 2) dislocations migrate from the high density outside regions toward the low density neutral axis, thus producing the equilibrium distribution of dislocations; and 3) a polygonized structure is formed by movement of the dislocations into walls normal to the slip plane. ETCHING to reveal edge dislocations'-" offers a unique opportunity to study, in germanium single crystals, the functioning of dislocations in plastic deformation. Plastic bending is ideally suited to this purpose because it produces a static array of edge dislocations in an ideally plastic material which is predictable from the theory." " Nye" has given a rather complete treatment of the geometry of bent crystals in terms of dislocations. The density of excess dislocations is related to the radius of curvature of the slip plane of a bent crystal, as shown in Fig. 1, assuming the absence of macroscopic elastic stresses. The curvature of any plane in the zone of the bend axis is found by resolving the Burgers vector into that plane, so that for the usual case where the slip plane is inclined at some angle 0 to the neutral plane of the bar, the equation applies, with T equal to the radius of curvature of the neutral plane of the bent crystal. Two predictions which are relevant to this study may be made from Eq. 1. First, in crystals bent to various radii, the average excess dislocation density is inversely proportional to the radius and second, for a bar bent to a radius which is large compared with its thick-

Jan 1, 1957

By K. R. Janowski, H. Conrad

As part of a program to establish the effect of crystal imperfections on laser output, a detailed study was made of the dislocation structure of ruby crystals obtained from varioius sources. Using KHSO4 as an etchant, a detailed mapping of the dislocation structure on the (0001) and the (1150) planes was carried out. A less extensive study was made of the rhombohedral planes. The average dislocation density on the basal plane was 1.5 to 3 x 106 cm&apos;2 and on the (1120) planes -5 x 10&apos; cm-2. Howe7!er, considerable variation existed between areas on a given plane. The subboundaries in the basal plane tended to lie along [1100] type directions while those on the (1120) plane tended to lie: a) parallel to the basal plane, b) along truces of the (1101) plane, and c) along normals to the traces of the (0001) planes. The orientations suggest that mang of the dislocations lie on the (0001), (1701), and (1150) planes A Laue back-rellection nnalysis Of c-axis wander in one crystal revealed a gradual increase in misorientation from me end of the crystal to the other, with a maximum variation of 1°35&apos;. Slight etching of a mechanically polished crystal revealed mumerous polishing scratches, suggesting the existence of a plastic flow layer. It is recognized that chromium concentration, lineage, and other crystalline imperfections play a role in the efficiency of ruby lasers.&apos; Consequently, it is desirable to establish in detail the imperfection structure of ruby crystals to be used as lasers. Some exploratory work along this line had already been done by Barnes&apos; and Dils and Martin.&apos;" How- referred to as the morphological system, as described by Kronberg.4 The Miller indices given in the present paper will be based on this morphological system. In this system, the three common growth planes are the (0001), the {1120}, and the {1101}.5 Since these planes are of high stability, it is expected that they are most suitable for study by the etch-pit technique. In the present investigation, a detailed mapping of the dislocation structure was carried out for the (0001) and the (1130) planes. A less extensive study was made of the etch-pit structure on rhombohedra1 planes. An analysis of dislocations in sapphire (the ruby host) was first made by kronberg,4 who postulated that slip on the (0001) plane occurred by motion of dislocations with Burgers vector 1/3 <1150> and that these dislocations might dissociate into quarter partials. An etch-pit technique using boiling phosphoric acid was developed by Scheuplein and Gibbs6, 7 for identifying dislocations intersecting the (0001) planes and the planes near (2021) in sapphire. Al-ford and Stephens8 were able to reveal dislocations on the (0001) and {1120} planes in sapphire and ruby by etching in fused potassium bisulfate (KHS04) at 675°C. The termini of basal dislocations [(000l), <1120>] were revealed by Palmour and Kriegel9 on {1120} and {1010} by thermally etching sapphire. Bond and Harvey10 were able to decorate dislocations in sapphire by adding 1 pet by weight of ZrO2 to the Al2O3 powder used in the growth of the crystals by the Verneuil technique and then heating the grown crystals for 4 hr at 1500°C. They found, using optical transmission microscopy, that many dislocations had straight segments which were parallel to one of the six [1120] directions in the basal plane. Single dislocation lines curved smoothly from one alignment to the other, turning through angles of 60 to 120 deg. Also observed were isolated hexagonal loops, approximately 10 µ in diameter, the sides of which lay along the same directions as did the longer dislocations. MATERIAL AND ETCHING PROCEDURE The ruby specimens used in this study (all containing approximately 0.04 pet Cr) had been obtained from four different sources: Linde Co., Meller Co., Valpey Crystal Corp., and those grown at Aerospace Corp.* For the initial phase of this study, several specimens were cut from the slab shown schematically in Fig. 1 which had previously been cut from a disk boule obtained from the Linde

Jan 1, 1964

By Nicholas J. Grant, Oliver Preston

PUBLICATION of data by Irmann&apos; indicating outstanding thermal stability and elevated-temperature strength properties in a sintered aluminum powder product (SAP) stimulated interest in the strengthening of metals by means of a highly dispersed oxide phase. This material was produced by compacting, sintering and extruding a fine flake aluminum powder, resulting in a dense product containing 10 to 16 volume pct alumina, which arises from the thin oxide film on the starting powder, highly dispersed in a matrix of pure aluminum. The method of producing SAP, and the stable high-temperature properties obtained from it, introduced the possibility of developing superior materials for high-temperature service by powder metallurgy methods. The composition of the materials of this type which are presently under consideration is predominantly metallic, containing from less than 1 pct up to about 20 volume pct of the oxide. As a result, other properties of these materials, such as electrical and thermal conductivity, thermal shock resistance, and to a certain extent ductility, notch sensitivity, machinability and hot working properties, are also predominantly metallic. These alloys do not constitute, therefore, a completely new class of materials as do the cermets, but rather present a new means of modifying the properties of a metallic matrix. In this sense, SAP type alloys are more closely analogous to the alloy aging systems. While the strength properties realized in the oxide-pure metal systems at the lower temperatures are substantially less than the best attainable in the aging systems, the thermal instabilities in the latter make them of less value for use at temperature much above their optimum aging temperatures, as indicated in Fig. 1. While the means of obtaining an oxide dispersion are more difficult and costly than conventional aging processes, in the light of the as yet limited understanding of such systems there is considerable promise of improved high temperature properties arising from the thermal stability of these structures. Since the development of SAP in 1946, its properties and the properties of a series of similarAl,O,-A1 products containing from about 1 to 20 pct aluminum oxide have been fairly extensively studied and reported.&apos;- "articular attention has been given to

Jan 1, 1958

By E. W. Hart, W. R. Hibbard

The room temperature flow characteristics of a series of Cu-Cr alloys are found to be related to the amount and characteristics of the chromium-rich precipitate. The results are consistent with the theory of Fisher, Hart, and Pry, which considers the trapping to dislocation loops around precipitate particles. THE use of dispersed phases for hardening metals, particularly high temperature alloys, has been practiced for several years. Current theories describing the effects of dispersions on the basis of dislocation theory have been proposed by Mott and Nabarro.l-4 rowan," and Fisher, Hart, and Pry.8 These theories recently have been discussed by Hart,&apos; who pointed out that the experimental data available for comparison with theoretical calculations for overaged alloys are confined to studies by Gensamer, Pearsall, Pellini, and Low,8 Shaw, Sher-by, Starr, and Dorn,9 and some hitherto unpublished data on Cu-Cr alloys which will be described herein. Certain difficulties were associated with the interpretation of the first two studies. Gensamer, Pearsall, Pellini, and Low were concerned principally with interparticle spacing and did not document particle sizes. Shaw, Sherby, Starr, and Dorn included all the data in their report which was necessary for the theoretical evaluation, but their study is subject to some question1" due to difference in prior history of various specimens and due to systematic differences in particle characteristics" which appeared to be associated with differences in the magnification of their microscopic examination. The current study was undertaken to obtain critical data in a manner which avoids some of these limitations. Specimens of eight Cu-Cr alloys remaining from a previous investigation&apos;&apos; were available in the form of 100 mil sheet of the compositions shown in Table I. At 200 mil this material had been previously an- nealed ½ hr at 1000°C, water quenched, and cold rolled 50 pct. Samples were further cold rolled to a thickness of 20 mil for a total reduction of 90 pct. Sheet tensile specimens of standard geometry with a gage length 2½ in. long and 0.20 in. wide were machined from this sheet. These specimens were then annealed in a dried N2-H2 atmosphere for the times and temperatures indicated in Table I. This anneal served two purposes: 1—to obtain a fairly uniform and nearly constant grain size of 0.015 to 0.022 mm average diameter, and 2—to precipitate a fairly uniform, nearly spherical over-aged dispersion of essentially pure chromium. A typical microstructure illustrating the grain size and the uniformity of the dispersion is shown in Fig. 1. A typical dispersion illustrating the range of sizes and shapes of the particles is shown in Fig. 2. The tensile specimens were tested in the Instron equipment described in ref. 11 at room temperature at a strain rate of 0.04 in. per in. per min. Three tests were run on each alloy with sufficient repro-ducibility that a single true stress-true plastic strain curve could be drawn through the three sets of data for each alloy. Typical curves are shown in Fig. 3.

Jan 1, 1956

By N. J. Grant, K. M. Zwilsky

A series of copper-alumina dispersion strengthened alloys were prepared using three different copper and two different alumina powder sizes. Improvements in strength of up to ten times that of pure copper were found in room-temperature as well as elevated-temperature stress-rupture tests. Conductivity values remained from 70 to 90 pct that of pure copper. The best alloys were based on the finest metal powder (1 p) and contained from 7.5 to 10-0 vol pet of 0.018 or 0,033 p alumina. DISPERSION strengthening of metals consists of strengthening a metal matrix by the addition of oxides or other refractory constituents, followed by strain hardening, in order to reinforce the base material and extend its temperature range of application. Since the discovery of SAP by lrmannl and Von zeerleder2 in 1949, many investigators have reported data obtained on a variety of metal-inert oxide alloys.3-7 The present work deals with the mechanical mixing of copper and alumina powders. While several techniques of introducing the second phase have been investigated to date, mechanical mixing so far has been recognized as the simplest, least expensive, and most universally applicable technique that can be used to obtain an ultrafine dispersion of inert particles. In the present case, the effects of changing the metal particle size, the oxide particle size and the volume percent of oxide were investigated. This work is an extension of a previous study.6 EXPERIMENTAL PROCEDURE The starting materials were three different grades of copper powder and two different alumina powders. The three metallic powders were -74 µ(- 200 mesh) electrolytic copper powder consisting of irregular equiaxed particles having a normal screen distribution and analyzing 99.5 pct Cu, and two powders having a 5-µ and 1-µ particle size, average, respectively. The latter powders were produced by the hydrogen reduction of purified sulfate solutions at elevated temperature and pressure and were furnished by the Sherritt Gordon Mines, Ltd. The alumina used was obtained from Godfrey L. Cabot Co., under the trade name of Alon C, one batch of powder having a particle size of 0.018 µ, the other a particle size of 0.033 µ. The crystal structure of these alumina powders was found to be cubic gamma. Each batch of powder was mixed in a Waring Blendor for a total of 15 min, followed by a low-temperature reduction in hydrogen and vacuum compacting in cylinders which were hydrostatically pressed at 35,000 psi to give a compact that could be handled conveniently. Six of the compacts were hydrogen sintered at 900°C while the remaining four were sintered at 1000°C. Subsequently, all compacts were hot extruded at 760°C with a reduction of area ratio of approximately 20 to 1. Evaluation of alloys included density and resistivity measurements, room-temperature tensile tests, stress-rupture tests at 450°C plus one other temperature per alloy, hardness tests, microstruc-tural examination, X-ray analysis of the dispersed phase after dissolving the metal phase, and oxidation studies on two of the alloys. A vacuum-melted, wrought copper rod was tested together with these alloys for comparison purposes. RESULTS Table I is a summary of the compositions, sintering temperatures, densities and conductivity values of the ten alloys prepared. All alumina additions are given in volume percent, and will be referred to as

Jan 1, 1962

By W. A. Backofen, G. Y. Chin, W. F. Hosford

The ductile fracturing process was studied in single-crystal and poly cvystalline aluminum deformed in tension over a temperature range from 295° to 4.2°K. At temperatures as low as 77°K, the fracture of "inclusion-free" material, including zone-refined aluminum, was by rupture (-100 pct RA). At 4.2 OK, fracture was brought on by adia-batic shear. Metallographic examination did not disclose any voids or slip-band microcracks, thus negating for inherently ductile metals any mechanism of void nucleation by vacancy condensation or of cracking due to dislocation pile-ups. In Izigh-purity aluminum not treated to be inclusion-free, fracture at temperatures as low as 45°K was of the double-cup type and a result of void formation. The reduction-of-area decreased as temperature was lowered, corresponding to the earlier appearance of voids. Such behavior was rationalized in terms of a larger increase, with decreasing temperature, in the .flow stress relative to the strength of the inclusion-matrix interface. Evidence for low-temperature adiabatic shear was found in discontinuous flow at 4.2"K, in the transition to a localized shear fracture at low temperatures, and in the suppression of shear fracture with an elastically hard pulling device. A simple analysis for the initiation of adiabatic shew permitted a general correlation of the various contributing factors. It has been pointed out that the duration of shear depends upon effective mass and elastic stiffness of the deformation system. IT has long been recognized that fracture* may Throughout this paper, the term "fracture" is taken to mean any process that results in the separation of a material into two (or more) parts. Thus rupture as it may be encountered in a tension test leading to 100 pct reduction-of-area is included in this category. occur in a ductile mode, and that the process can be of great practical as well as general interest. Much information about ductile fracture has also been accumulated over this period, but only recently has an understanding of mechanism begun to appear. Ludwik,&apos; in 1926, first reported fracture in a tensile specimen starting with a central crack in the necked section. Since then, other studies have disclosed that such cracks may form by the coalescence of voids nucleated in this region where hydrostatic tension is highest.2-4 Rogers and Crussard et al.&apos; have emphasized void formation and reori-entation along localized shear bands as a mode of crack propagation. pines6 has considered the tensile rod as a bundle of fibers joined by weak interfaces, which subsequently separate to allow individual fiber contraction. The notion of cavity growth and coalescence by purely plastic processes was discussed by Cottrell: who added that the tensile reduction-of-area ought not to be sensitive to temperature. On the other hand, it has been observed that the reduction-of-area is greatly increased if tests are carried out at high temperaturesa or under high hydrostatic pressure.&apos; Fracturing anisotropy in wrought products lends support to the idea of void formation from preexisting flaws strongly aligned by earlier processing.&apos;&apos;-l2 There is evidence that many voids result from the fracturing of inclusions or separation at the inclusion-matrix interface Another possibility is that voids grow out of pore volume produced in the initial solidification and never fully removed in later working. In general, a structure 3f particles, pores, and weak interfaces can be expected, at least in materials of engineering interest. Vacancy condensation has been suggested as an alternative mechanism of void formation for materials considered to be inclusion-free.13 Yet experience has shown that tensile reduction-of-area increases with purity, to the extreme of rupture as so often observed in single crystals. Adiabatic shear has an important bearing on ductile fracture. It occurs when the decrease of flow stress, as a result of local temperature rise from heat generated during straining, becomes larger than the increase due to strain and strain-rate hardening. As demonstrated by experiments on punching of plates,14 a large temperature rise may be brought about by rapid straining. Adiabatic flow as a result of the high strain rate reached in an ordinary tensile specimen just prior to separation may account for the cone formation in cup-and-cone fracture;14 evidence of such local heating has been presented.15 For geometrical reasons, however, pure sliding along the conical surfaces is unlikely, and separation under tensile forces is probably an important accompanying feature of the shear.7 In deformation processing operations, a high shear-strain rate may exist at boundaries between plas-

Jan 1, 1964

By J. Greenspan

The anisotropy of fracture and slip, that is, the brittleness and ductility of the beryllium single crystal, is characteristic also of po1ycrystalline beryllium in which the grains are oriented in a preferred manner. Beryllium with grain orientations resulting from hot working either unidirectionally or bidirectionally is ductile or brittle in directions given by the anisotropy and orientation of the grains. Textures developed from various hot-working sequences are given, and tensile results are correlated with textures. Not only texture, but fine grain size is necessary for obtaining high tensile elongation. Ultimate tensile strength and elongation are both correlated with grain diameter. Of particular interest is the case in which unusually high tensile elongations of 30 to 40 pct are obtained when basal planes are oriented to carry THE beryllium crystal (of nominal purity) is probably unique among metals with respect to its unusually acute anisotropy of slip and fracture. While extensive slip can occur on the system (1010) [ll20] starting at about 19,000 psi, the (0001) plane is particularly vulnerable to fracture at 4500 psi or less. Fracture occurs also on the (11zo) plane at about 26,000 psi.1"3 Further, there seems to be no other slip System which contributes significantly to

Jan 1, 1960

By Erhard Hornbogen

Twins were propagated into large, well-annealed crystals of a, iron-phosphorous and a, iron-molybdenum solid solutions. Strain fields caused by interaction of these twins were made visible by precipitation. These strain fields can be explained by assuming that transmitted and reflected plastic waves are created when a twin strikes an obstacle with high velocity. The twin front itself can be regarded as a shear wave that may be able to develop a one-dimensional shock front. The occurrence of {WO} fracture by the reflected stress is also discussed. The described effects were not found when twins propagated at lout velocity into disturbed crystals, for example, into crystals previously deformed by slip. DEFORMATION twinning leads to shearing of a crystal lattice under external load. The shear appears to be homogeneous over a large number of crystal layers. In a iron, shear occurs on the (112) planes and in the <1ll> directions. The formation of a twin proceeds by the movement of partial dislocations that change the stacking at the twinning plane.l, In the bcc lattice, this partial dislocation can be created by dissociation of a sessile a/2[111] dislocation that lies in the (112) plane: Cottrell and Bilby&apos;s spiral mechanism can explain the homogeneous shear, but has not been proven experimentally. Nevertheless, this mechanism can explain the possible high velocity of formation of a twin. In Cottrell and Bilby&apos;s model, as in other models of twinning or phase transformation, a much higher energy is required to spread the first stacking fault than to add new layers by further movement of the partial dislocations. The dislocations will therefore become accelerated proportional to the difference between the stress, tN, to nucleate and tp to propagate the twin. It has been assumed3,4 that tn is the stress required to make two partial dislocations of apposite sign pass for the first time. After the first layer of new stacking has been formed, the maximum stress is TN=Gb/4pd where G is the shear modulus, b is the Burgers vector of the partial dislocation, and d is their perpendicular spacing so that the amount of shear is S = b/d. From this formula, a value of 25 kg per mm2 has been estimated for twinning in Zn3. The measured values are smaller, but local internal stresses may help to provide the necessary stress. For a iron, a higher value can be expected, mainly due to the higher value of G. That agrees with the observation that "burst" twinning occurs if the yield strength is in-

Jan 1, 1962

By Harry L. Brown, Philip E. Armstrong

Young&apos;s modulus doto ore presented for W, Mo. Ta. V, Cr. Ni, Ti, and Zr as a function of temperature up to about 0.7 of the melting points. A plot of reduced temperature us reduced modulus produces a general correlation except for titanium and zirconium. Young&apos;s modudlus data for some representatiue commercial graphites are given up to 2400°C. All results were obtained from longitudinal resonance-frequency measzrrements of 4-in. long samples, corrected for simultaneously measured thermal expansion. Comparison of poly crystalline and single-crystal tungsten behavior indicates that all the observed modulus values are frequencv -dependent at the higher temperatures. STATIC measurements of elastic properties of materials at elevated temperatures are extremely difficult to make since the elastic behavior is obscured by relatively large first-stage creep effects. Dynamic methods which avoid this difficulty by utilizing the measurement of the resonance frequency of a simple vibrational mode have been described in the 1iterature.l-3 The authors have extended the operating temperature range of this type of measurement to above 2000°C and have described the equipment elsewhere.~ The Young&apos;s modulus data presented here were obtained with this equipment, using samples about 4 in. long and 1/4 in. diam, supported horizontally at the center and excited to the fundamental longitudinal resonance frequency. Young&apos;s modulus was calculated by the formula: E = 4f2f2p where E = Young&apos;s modulus, / = frequency, 1 = length, and p = density. Dilatometric data were obtained simultaneously by optical measurements of the change of length of the sample. For the metals tested, an attempt was made to obtain modulus values that would be representative of the pure, strain-free, randomly oriented material. All of the graphites studied exhibited a high degree of elastic anisotropy, while the method assumes an isotropic homogeneous material. However , when the values for Young&apos;s modulus obtained for AUC graphite were compared with those taken from static measurements at room temperature they were found to be only about 4 pet higher. This is excellent agreement considering the low precision of modulus calculations from stress-strain curves. The difficulty in using data of apparent Young&apos;s modulus for anisotropic materials arises when the attempt is made to use these data and shear-modulus data to obtain Poisson&apos;s ratio. The results are usually meaningless and sometimes absurd.&apos; RESULTS Table I lists the sources, impurity content, thermal history, and dilatometric constants of the metal samples that were investigated. All the metal samples were tested under a vacuum of 1 to 5 x 10-5 Torr. Where our dilatometric data agreed with previously published data, the table lists the previous values and the literature reference. Where no reference is given, the constants were determined from our dilatometric data. Table II presents the modulus data in units of millions of psi as a function of temperature in degrees centigrade. The numbers were obtained from curves fitted to the experimental data. For some types of calculations, such as those attempting to correlate creep data, the slope of the modulus curve is of interest, and this information also is reported in Table 11. It is always tempting to try to fit physical-property data into some general pattern and then use this pattern to predict the behavior of other similar materials. The general pattern for an elastic modulus of a metal seems to be a linear deoro- with increasing temperature Over the range between 0.1 and 0.4 of the melting temperature, then a progressive1y decreasing value up to the melting point. Metals appear to retain finite moduli

Jan 1, 1964

By R. L. Fleischer, W. F. Hosford

Tensile data for coarse grained aluminum Polycrystals suggest that the "grain size" effect is not due to dislocations piled up at grain boundaries but rather is primarily a relative size effect due to surface crystals being weaker and less confined. STUDIES directed at interpreting hardening of poly-crystalline metals normally identify their strain hardening properties with those in some particular type of single crystal.1"4 The recent recognition in face centered-cubic metals of a nearly linear stage with rapid hardening occuring at comparable rates for both polycrystals and single crystals, suggested that the same process or processes determine both cases and hence that there exists some justification for the use of single crystals to understand polycrystals. Further evidence for the above view may be found by an approach initiated by Chalmers:5 By using bicrystals of controlled orientation it is possible to begin to assemble a polycrystal and also to study grain boundary effects in detail. In this way it has been found that a single grain boundary affects easy glide but not the subsequent stage II hardening.6 This result suggests that a sensitive way to observe grain boundary effects in polycrystals would be to vary grain size and measure easy glide. As will be seen, easy glide is only possible for coarse-grained samples, and hence the results will serve to fill in the gap in measurements between single crystals and bicrystals on one hand and fine-grained polycrystals on the other. One problem inherent in comparing single crystals with polycrystals is the uncertainty as to what slip systems are acting in a polycrystal. To compare the two types of samples, rates of shear hardeninn---L. on the acting -planes are needed. and these may be computed only if it is known what particular systems are active. The acting systems were examined for a coarse-grained polycrystal and it will be shown that the systems supplying the preponderance of slip can be determined with little ambiguity. EXPERIMENTAL PROCEDURE Twelve samples of aluminum were prepared by chill casting into a heated graphite mold, followed by annealing at 635° ± 5°C for 24 hr with an 8-hr fur- nace cool, and finally either etching7 or electropol-ishing.&apos; The samples, with a 7 to 10 cm length between grips and 4.4 by 6.6 mm in cross section, were deformed at a strain rate of about 3 10 -3 . per min in a tensile device which has been described elsewhere.5 The composition was reported by Alcoa as 99.992 pct Al, 0.004 pct Zn, 0.002 pct Cu, 0.001 pct Fe, and 0.001 pct Si; nine samples were deformed while immersed in liquid helium and three in air at room temperature. The stress-strain curve for one of the samples (P-1) deformed at 4.2 "K has been reported previ~usl~.~ This sample was selected for determination of active slip systems. Eighteen of the crystals were examined by optical microscopy to determine the angles of slip line traces and by X-ray back reflection to determine orientation. By this means the slip planes were determined and the resolved shear stress factors for possible slip systems could be computed. Finally each sample was sectioned so that after etching, the number of crystals could be counted for each of ten newly exposed surfaces. The average of these ten values will be termed n, the number of crystals per cross section. Values of 11, varied from 1.9 (nearly bamboo structure) to 12.7. Sketches of typical cross sections appear in Fig. 1. RESULTS AND DISCUSSION: SLIP SYSTEMS 1) Determination of Acting Slip Planes—The stress axis orientation and operative slip planes in eighteen crystals of sample P-1, as determined by slip line traces and crystal orientation, are summarized in Fig. 2. For one of the crystals two planes had a common trace. so that the traces alone did not distinguish which plane or planes were slipping. However it was found that the stress resolving factor for the primary system was 0.386, .while that for the most stressed system in the other plane (indicated bv the dotted arrow) is 0.138. It will be assumed tgerefore that only the primary plane acted. Since the orientations were determined after extending the samples 4 pct, the stress axes may be rotated from their original value by as much as 2 deg in some cases. It is interesting to note that in five crystals only one slip plane acted, in eight two acted, and in five three planes were observed—an average of two slip

Jan 1, 1962

By M. J. Marcinkowski, A. Szirmae, R. M. Fisher

Room -temperature hardness measurements obtained from single and polycrystalline samples of a 47.8 at, pet Cr-Fe alloy which were aged for various times al 500°C show a two-fold increase over that of the unaged alloy after annealing for 1000 hr. A detailed examination of the deformation markings in the neighborhood of the hardness imbressions reveals that twinning becomes an increasingly more important mode of deformation as aging proceeds. this observation is shown to be inconsistent with the Proposals that the transformation is eve of order-disorder On the other hand, the results are in agreement with previous observations of Fisher et al. of a coherent chromium-rich precipitate which forms in an iron-rich matrix as a result of a miscibility gap in this alloy system as proposed by Williams . Using the coherent precipitate hypotlzesis as a model, a detailed analysis of the various possible strengthening contributions to both the slip mid twinning stresses is made. In the case of&apos; slip, lattice friction within the chromium-rich phase and the chemical energy associated with the interface between the two phases contribute about 60pet to the total strength. The contributions from coherency strains and modulus differences are thought to contribute the remaining 40 pet of the strength but are difficult to evaluate because of uncertainties regarding the flexibility of the dislocation line. All of the factors have nearly the same or else a smaller effect on the twinning stress in the aged alloy. Twinning is never observed in the unaged alloy because the twinning stress is much higher- than that for slip. With increased aging times, the various contributions to the total stress for both twinning and slip increase, but most of those for slip increase much more rapidly, so that, in the fully aged alloy, it surpasses the stress for twin propagation. When a twin is nucleated by the chance occurrence of an internal stress concentration during a test, the subsequent twin burst results in a low hardness reading. FeRNTIC Fe-Cr alloys from about 15 to 80 at. pet Cr content show considerable change in properties such as hardness, electrical resistivity, saturation magnetization, and so forth, after aging in the vicinity of 500°C.* The most marked change is a large increase in hardness accompanied by a sharp decrease in ductility, so the phenomenon has often been referred to as the 475°C (885°F) em-brittlement problem. Changes in microstructure are rather subtle and most investigators have not been able to observe any X-ray diffraction or metal-lographic changes during aging. There is little doubt that the phenomenon is inherent to the Fe-Cr binary system and is not directly related to the formation of a phase. Two distinct schools of thought have developed during the past decade concerning this problem. Masumoto, Saito, and sugiharal have concluded, from their own specific-heat measurements, that atomic ordering occurs in this temperature range. Pomey and Bastien2 have also attributed the changes in physical properties with aging in the neighborhood of 500°C to atomic ordering. In addition, Takeda and Nagai3 claimed to have found X-ray verification for superlattices corresponding to the compositions FesCr, FeCr, and FeCr3. However, all known attempts to observe superlattices in Fe-Cr alloys using neutron diffraction, which should be much more sensitive, have been unsuccessful. These have been reported by Shull et al.4 for 25 at. pet Cr, williams5 for 75 at. pet Cr, and Tisinai and samans6 for a 28.5 at. pet Cr. steel. The second and more prevalent opinion ascribes the 475°C embrittlement phenomenon to the formation of a coherent precipitate due to the occurrence of a miscibility gap in the Fe-Cr system below about 550°C. This latter concept was first indicated by the results of Fisher, Dulis, and Carroll,7 who were able to extract fine particles about 200Å in diameter from samples of 28.5 at. pet Cr steels aged from 1 to 3 years at 475°C. The extracted material was found to be nonmagnetic, to have a bee structure with a lattice parameter between that of iron and chromium, and to contain about 80 at. pet Cr. Williams and paxton8 and williams5 confirmed these results and first explicitly proposed the existence of a miscibility gap below the a region in the equilibrium diagram. Alloys aged within this gap

Jan 1, 1964

By R. Raring, W. J. Harris, J. A. Rinebolt

The effects of additions of alloying elements on the true-stress, true-strain characteristics of 0.30 pct C, 1.00 pct Mn, 0.30 pct Si pearlitic steel were studied. The alloying elements investigated were C, Mn, Si, P, The alloys are compared on the basis of fracture stress, fracture strain, energy to fracture, rate of strain hardening, strain hardening exponent, height of the flow curve, and lower yield point. IN the past, nearly all studies of the effects of alloying elements on the tensile properties of metals have compared the conventional engineering parameters, namely, tensile strength, yield strength, elongation, and reduction of area. Recently, however, the true-stress, true-strain method of recording results from the tension test has been used; Lacy and Gensamer,&apos; Hollomon," and French and Hibbard- are used this method to study the effects of composition on strength and ductility of metals. In this paper, the effects of certain alloying elements on true-stress, true-strain curves of steels with nearly constant grain size and pearlite spacing are presented. The notation employed in this paper is given in Table I. Since the definitions of stress and strain and the analytical geometry of the stress-strain curve have been explained previously, for example by Gensamer,&apos; they will not be reviewed here. However, the method used for the determination of energy to fracture, E,, requires mention. These energy values were calculated on the assumption that the flow curve followed the idealized equation:&apos; corresponds to the specific energy absorption to that strain value. There is justification for preferring the use of eq 1 for the determination of the energy to a particular strain rather than the area under the experimental stress-strain curve, since Gensamer showed that the use of the Bridgman correctione for stress at the neck of a tensile specimen brought the corrected experimental curve nearer to the idealized power-law curve than to the uncorrected experimental curve. However, it should be noted that the difference between the areas under the idealized curve and the experimental curve is not great for strain of less than 1.0, as can be seen from Fig. 1. In this figure, the true-stress, true-strain curve for a 0.01 pct C, 1.00 pct Mn steel is plotted through the experimental data points. From the value of a,, from this curve and the value of n as determined by the Inm-

Jan 1, 1952

By R. Raring, W. J. Harris, J. A. Rinebolt

G. W. Geil (National Bureau of Standards, Washington, D. C.)—The authors state that the degree of accuracy realized in the experimental determination of a, is likely rather low. This inaccuracy is attributed to the difficulty in reading the load of the tension machine at the instant of fracture of the test bars. I believe there is another factor which may have added considerably to the inaccuracy in the determination of a,, namely, the determination of 8,. Fracture of a cylindrical tensile specimen of a ductile metal tested in tension at room temperature is generally initiated at the axis and rapidly propagates radially to the periphery of the necked section. During the propagation of the fracture, the metal at the rim of the minimum cross-section extends longitudinally and contracts radially. Thus values of 81, based on diameter measurements made after complete fracture of the specimens do not necessarily represent the strain at the initial fracture." Data obtained at the National Bureau of Standards with a 0.3 pct C steel in the condition as-quenched from 1575°F and tempered at 1000°F for 1 hr and tested in tension at room temperature indicated a considerable "rim effect" and as shown in Fig. 13, resulted in an apparent increase in Sf from 0.90 to 0.98. As the fracture strain of this steel is approximately equal to those of the steels reported by the authors, it seems probable that a similar "rim effect" may have added to the inaccuracies in the reported values of Si and accompanying values of El. R. Raring (authors&apos; reply)—The authors are grateful to Mr. Geil for having called attention to a possible source of uncertainty in the determination of fracture stress, and they are in agreement with the principles upon which his discussion is based.

Jan 1, 1952

By H. T. Green, R. M. Brick

True stress-strain data on alloys of pure iron with up to 2.4 pct Al were obtained in the temperature range +100° to —185°C. Alumi-num was found to reduce yield and flow stresses of iron at low temperatures but to have little or no effect on ductility. The effects of temperature and composition on strain hardening are discussed. SEVERAL independent studies of the behavior of high purity iron binary alloys at low temperatures are now in progress in attempts to evaluate systematically the variables affecting the low temperature brittleness of ferritic steels. This paper reports the results of one such investigation in which the tensile properties of aluminum and aluminum plus silicon ferrites were measured from 100" to —192°C. True stress-natural strain data have been obtained in order to evaluate as many as possible of the parameters which describe the behavior of the materials involved. In comparable studies at the National Physical Laboratory in England, iron and iron alloys of high purity have been produced&apos; and tested at subat-mospheric temperatures.&apos; True stress-natural strain curves were obtained there also. The purest iron contained 0.0025 pct C and 0.001 pct O and N. Even this, as normalized at 950°C following hot rolling, showed little ductility at -196°C. The grain size was ASTM No. 3, and the room-temperature yield strength was 17,800 psi (which seems too high for pure iron). Some of the NPL irons contained considerably more oxygen and demonstrated intergran-ular fracture at —196°C. The authors2 carefully differentiated between intergranular fractures associated with excessive oxygen content and transcrys-talline cleavage with little ductility encountered at —196°C in the purer material. The cleavage stress was half again as great as that associated with inter-granular fracture. Test Material, Preparation, and Procedures Of a number of Fe-A1 alloys produced, eight were considered to be sufficiently pure for testing. Partial chemical analyses (Table I), low observed yield points, and high ductilities indicate these alloys to be comparatively pure for vacuum-melted irons of sizable ingots, 5 Ib or more. To produce the binary Fe-A1 alloys, electrolytic iron was melted in air, cast into slabs, and rolled to strips 0.010 in. thick. These strips, joined into a continuous ribbon and wound into 2 1/2 in. diameter spools, were subjected for four weeks to a moving atmosphere of purified dry hydrogen in a stainless-steel tube at 1050" to 1150°C. Charges of these spools were melted in beryllia crucibles under good vacuums (1 micron), and aluminum (99.97 pct Al) was added to the melts. Compositions of these alloys are recorded in Table I. The ingots were hot forged and then cold rolled at least 65 pct to 3/8 in. rods which were vacuum annealed to the desired grain size, approximately ASTM No. 4, prior to machining into tensile test bars. All tensile specimens had gage sections 1 in. long, with a fillet of 1.5 in. radius to the shoulder. Gage diameters were 0.250 in, except for a few rods where additional cold work required use of a 0.200 in. gage section. After machining, 0.002 in. was removed from the gage diameter using 240, 400, and 600-grit metallo-graphic papers. The final polish with 600 grit left the fine scratches running in the longitudinal direction. By this means, surface metal strained during machining was removed. A few specimens heat treated after machining were similarly reduced 0.004 in. to remove any material affected chemically by the atmosphere during heat treatments, as is discussed in a later section. Tensile tests of the eight alloys at constant temperatures from +100° to —185°C were performed in apparatus which has been described." The essentials include a double-walled insulated metal vessel which contained the liquid heat-transfer medium surrounding the test specimen. A constant temperature was maintained by means of a pyrometer which regulated the pressure of dry air driving liquid air through a copper coil. Temperature variation was less than ±2°C during a specific test. For axial straining, two lengths of case-hardened chain, terminating in simple shackles, loaded the specimen through threaded grips. The lower grip bar passed through a hole in the bottom of the test vessel to which it was joined by a thin-walled

Jan 1, 1955

By B. L. Averbach, Morris Cohen, S. A. Kulin

The martensitic transformation can be initiated by elastic stresses at temperatures above M. in a steel containing 20 pct Ni and 0.5 pct C. Shear strains and normal tensile strains acting on a potential habit plane promote the transformation but compressive strains oppose it. THE strain sensitivity of the martensitic transformation has long been recognized. It is now well known that martensite formation can be induced by plastic deformation at temperatures below, and not too far above, the M, point.&apos; However, the exact role of elastic vs. plastic strains and of the concomitant stresses has proved to be a very elusive matter. Many investigators2-&apos; of the steel hardening reaction have made use of the stresses resulting from volume changes to account for such observations as the existence of retained austenite and its variation across the section of a bar, the change in the amount of retained austenite with quenching rate, and the self-stopping nature of the transformation when the cooling is stopped. The effect of applied strain as an independent variable has been studied in 70 pct Fe-30 pct Ni alloys by Scheil7 and McReynolds.8 Some data are also available for lithium and Li-Mg alloys,9 Cu-Zn,10 Cu-Sn and Cu-A1 alloys," and aus-tenitic stainless steels."-" It has been found that martensite formation can be induced isothermally, even at temperatures above M,, by mechanical deformation of the parent phase. The competitive nature of plastic yielding by slip and by the martensitic reaction has been depicted by Scheil.7 Basically, Scheil postulated that: 1—at temperatures not too far above M, (where austenite is less stable than martensite), a critical resolved shear stress (within the elastic range) is required to promote the transformation to martensite; 2—at and below M,, the austenite lattice becomes mechanically, as well as thermodynamically, unstable and shears over spontaneously into martensite without the application of external stress; and 3—at temperatures sufficiently far above M8, the critical resolved shear stress for martensite formation increases to a level above that required for slip, and plastic flow then supersedes the transformation when external stress is applied. Scheil&apos;s concepts were tested by McReyno1ds8 ho found no change in the elastic moduli (measured both statically and dynamically) of 71 pct Fe-29 pct Ni alloys in the vicinity of the M, temperature. Since the moduli did not approach zero or become negative on cooling to M,, it was concluded that the austenite lattice does not become mechanically unstable at M2, as postulated by Scheil. McReynolds also reported that M, is not raised by elastic stresses. Accordingly, plastic deformation was considered to stimulate the transformation by generating martensite nuclei in the distorted regions of the parent phase. The above issues with regard to the role of applied stress require clarification, if a basic understanding of the martensitic transformation is to be obtained. For example, in the reaction-path theory,&apos;"," it is postulated that strain embryos exist in the austenite that provide part of the activation energy for the nucleation process. These embryos are visualized to be regions of localized strain in which the atoms in the austenite lattice are displaced part way along the path to their ultimate positions in the martensite. Consequently, it would

Jan 1, 1953

By W. Rostoker, D. W. Levinson, A. Yamamoto

The influence of carbon on tensile strength, tensile ductility, transformation kinetics, and grain growth characteristics of selected Ti-Mo base alloys was studied. No systematic influence of carbon in tensile strength or significant deleterious effect on tensile ductility was observed. There was a marked increase in transformation rate due to carbon. Dispersed carbides exert a marked inhibiting effect on grain boundary migration, and thus effectively prevent grain coarsening during solution treatment in the ß phase field. AS part of a research program directed toward the illumination of factors governing modes of transformation, transformation kinetics, and mechanical properties obtainable through these transformations by heat treatment, Ti-Mo base alloys were selected for study. The binary system is a characteristic phase diagram type which has been studied in some detail&apos;, &apos; in relation to transformation kinetics and mechanical properties. Although carbon bearing alloys have acquired a bad reputation, the possibilities of melting titanium under conditions which occasion some carbon pickup are very attractive. It is therefore of considerable importance to know whether carbon, present as a dispersed carbide phase, is actually deleterious when all other detrimental factors are minimized. Basic alloy compositions containing nominally 3, 5, 7, and 11 pet Mo were prepared with systematic carbon addition up to 1 pet by weight. Various of these alloys were subjected to a series of critical and definitive experiments to delineate the effect of carbon on the transformation behavior and associated mechanical properties. Experimental Procedure Alloys were prepared from 115 Bhn magnesium-reduced sponge titanium. This sponge analyzed ap- proximately 0.05 pet C. Ti-Mo and Ti-C master alloys were made by arc-melting the sponge with high purity (99.9 + pet) molybdenum and spectro-scopic grade graphite, respectively. Alloys were then prepared from the master alloys and sponge by double are-melting. One kg ingots were prepared by melting in a nonconsumable electrode furnace. These ingots were forged at 1000°C to &apos;/z in. rods, and cleaned and processed into electrodes for the second melting operation. The final ingots were ground to remove surface defects, then forged again to 1/2 in. or 3/8 in. diam rods for processing to test samples. It is significant that all of the alloys at all carbon levels forged satisfactorily. For convenience, the alloys will be referred to throughout the text by their nominal analyses. Samples for the study of transformation rates were prepared from 1/2 in. diam round stock. They were then solution treated in a helium filled tube furnace at 1000°C for 30 min prior to the isothermal transformation treatments in lead bath furnaces. Temperature control in both cases was within ±3oC. Following the isothermal transformation treatments the samples were water quenched. These experiments were carried out before the general adoption of the resistometric procedure for kinetic studies. The data were therefore obtained entirely by metal-lographic analysis. M, temperature was also determined by the metallographic method. Shoulder-type tensile test pieces were machined from heat treated, 1/2 in. diam slugs to provide a 0.252 in. test diam and a 1 in. gage length. The isothermal-type heat treatments were chosen from previous data&apos; to provide three levels of strength and ductility.

Jan 1, 1957

By H. Conrad

The strain hardening associated with the creep of magnesium single crystals at room temperatu.Je was investigated by shear tests in which the direction of stressing was reversed a number of times after creep in a <100> direction. The strain hardening was strongly anisotropic; the largest amount occurring in the direction of flow and the least 180 deg to this. The time law for creep changed from a parabolic type to logarithmic for the first and subsequent reversals of stress. For a given stress the creep strain decreased significantly for the first and second reversals and then only slightly for subsequent reversals. The logarithmic creep constant a changed in a similar manner. A linear relationship was observed between the creep strain for a given reversal and the stress. The slope for the third reversal yielded a strain hardening coefficient 12 to 16 times greater than that observed for unidirectional flow. Al-though no recovery was observed upon resting at room temperature, the hardening resulting from stress reversals could be completely removed by heating to 450" C. PREVIOUS work1.&apos; indicated that the creep rate of magnesium single crystals tested in the temperature range of 78" to 364°K decreased with time. This was attributed to an increase in internal stress resulting from an accumulation of dislocations at internal obstacles. The object of the present investigation was to further evaluate the nature of this hardening by conducting tests in which the direction of stressing is changed after previous creep in shear in a given slip direction. Previous to 1950 three investigations indicated that a Bauschinger effect3 occurred for single crystals as well as polycrystals brass in tension compression,

Jan 1, 1960

By D. Walton

The effect of small solute additions of Cu on the activation energies for creep A1 single crystals were determined over the range from 78° to 850° K. Below 240°K and above 800°K activation energies were unchanged. Between 240°K and 400°K strain aging causes the creep rate to become vanishingly small and the flow stress was independent of the temperature. Between 500° and 50°K the activation energy was increased to 41,000 cal per mole. NUMEROUS investigationsL have shown that solid solution alloying invariably increases the flow stress. A typical example is documented in Fig. 1. Such increases in flow stress can arise from two general effects: Alloying can increase the long-range stress fields through which dislocations must move by causing appropriate clustering; modification of dislocation densities, and so forth. Alloying can also modify the short-range stress fields that determine the activation energies for thermally stimulated migration of dislocations, such as occurs in some solute atom pinning processes, and so forth. Actually both long- and short-range stress-field effects can occur simultaneously. It was the purpose of this investigation to evaluate the effect of alloying on the short-range stress fields by determining the effect of alloying on the activation energies for creep. The resulting observations also shed some light on the effect of long-range stress fields on dislocation processes. EXPERIMENTAL TECHNIQUE Single-crystal bars (6 in. long 3/8 in. wide and l/4 in. thick) of dilute solutions of Cu in A1 were seeded, so as to provide extensive single slip, with the pole of the tensile axis located in the standard stereographic projection as shown in Fig. 2. Each crystal was grown under an argon atmosphere in a graphite mold containing a large reservoir of molten alloy for the purpose of obtaining rather uniform composition of Cu along the length of the bar. The nominal compositions of the alloys which were produced from 99.99 pct Cu and high-purity A1 containing 0.001 pct Fe and 0.001 pct Si, are given in Table I. Strains, measured by linear variable transformers attached by quartz rods to a 3 in. gage section were sensitive to and the stresses were ap- plied by direct loading of the specimen. Temperatures were determined by thermocouples attached to the gage section of the specimens. Below 400°K the apparent activation energies for creep, Q, were determined by the effect of small, rapid changes in temperature of about + 10°K from T1 to T2 on the corresponding creep rates ?1 and ?2 during primary creep according to where R is the gas constant.&apos; Above 400°K, Q was evaluated from Eq. [1] for the secondary creep rate using temperature changes of about ±20K. A typical set of results is shown in Fig. 3. RESULTS AND DISCUSSION The previously obtained effect of temperature on the apparent activation energy for creep of single crystals of high-purity A12 is shown by the broken curve of Fig. 4. The datum points on the same graph illustrate the effect of dilute a solid solution alloying with Cu on the activation energy between 600" and 750°K. Each plotted point is the average of not less than five independent test values between 600° and 750°K and not less than three at other tempera-

Jan 1, 1962

By R. J. Quigg, G. S. Doble

Commercial stress -relieved TD Nickel bar was shown to retain room- and elevated-temperature tensile strength after exposure up to 2501°F. Cold swaging increased both room -temperature and 2000°F tensile strength. After annealing at 2500°F the strengthening due to swaging was retained at 2000°F hilt eliminated at room temperature. Annealing produced changes in hardness and X-ray half peak width indicating the presence of recovery processes as 1071° as 1000°F but no recrystallization of swaged material was noted at any temperature. In contrast to swaging, rolling induced a recrystal-lized structure after annealing. The tensile strength of this recrystallized structure was comparable to the cold-worked condition. High-temperature tensile properties of as -received bar were highly anisotropic with the transverse tensile strength being one fifth the longitudinal value. The aniso-/ropy is affected by working direction and may he partially reversed by changing the direction of metal flow. The divergence in tensile strength, beginning above 550°F or- approximate one third the absolute melting point, is strain-rate dependent. T D Nickel is a commercial dispersion-hardened alloy containing 2 vol pct thoria. The dispersion, consisting of spherical particles 0.030 to 0.050 in diameter with an interparticle spacing of less than 1 is reported to maintain stable properties up to 2400°F.1 This stability of second phase results in maintenance of elevated-temperature strength at higher temperatures and longer times than normal nickel-base superalloys.2 The purpose of this paper is to describe the effect of various mechanical working operations on the strength and microstructural stability of TD Nickel bar. MATERIAL AND PROCEDURES All material used in this study was commercial TD Nickel bar. This material is produced by a chemical process which results in a nickel powder containing a dispersion of thoria particles. The powder is then pressed, sintered, extruded, and subsequently worked to bar stock. Three starting materials from three separate lots were used: as-extruded bar. 1-in.-diam stress-relieved bar, and 1/2-in.-diam stress-relieved bar. Since TD Nickel bar stock is cold-worked without a recrystalliza-tion anneal, the accumulated deformation or degree of cold work in these materials increases as the bar diameter decreases. Some of the 1/2-in.-diam bar was swaged to 1/4 in. diameter at room temperature in order to produce an as-worked condition for comparison with the stress-relieved stock. Tensile testing was performed on an Instron Tensile Tester at a constant crosshead speed of 0.020 in. per min. The 0.2 pct offset yield strength was measured from the load-crosshead movement plot. Specimen gage dimensions were 0.080 in. in diameter by 0.40 in. in length, producing a strain rate of approximately 0.050 min-1. Specimens were tested in air and heated within a resistance-wound furnace controlled to ±5°F. A 15-min soak at temperature was employed before testing. Samples were heat-treated in air and specimens machined after exposure with sufficient material removed to eliminate any surface contamination. X-ray line breadth was determined from a diffractometer scan of a transverse face of a bar using copper K, radiation. Metallographic preparation utilized standard mechanical polishing with Carapella&apos;s etch. Specimens for electron microscopy were electropolished in a 1:7 sulfuric/methaol solution and chromium-shadowed collodion replicas prepared. RESULTS AND DISCUSSION I) Effect of Elevated-Temperature Exposure on the properties of TD Nickel. The effect of various elevated-temperature exposures on the room-temperature tensile strength of TD Nickel is shown in Table I. In the unexposed condition the tensile strength, reduction in area, and hardness all increase with increasing accumulated deformation while the elongation decreases. After exposure at 2500°F the strength of all three materials decreases to about 65,000 psi and the elongations increase slightly. The apparent increase in strength of the 1-in. bar after a 2000°F treatment is probably experimental scatter. These results indicate that the room-temperature strengthening due to additional deformation is annealed out upon elevated-temperature exposure. In contrast to the room-temperature data, the 2000°F tensile strengths increase in order of increasing deformation even after exposure. Although some loss of strength occurs after exposure, much of the effect of working is still maintained at 2000°F. The room-temperature hardness after various

Jan 1, 1965

By B. Wilshire, P. W. Davis

It has been shown that grain-boundary migration during high-temperature creep can reduce or even prevent the formation of intercrystalline voids, giving a considerable increase in ductility.&apos; A similar observation has been made in a recent investigation of the creep properties of various grades of nickel and nickel alloys. The results can be illustrated by the observations made on: a) Pure vacuum-cast nickel (99.993 pct Ni) b) Impure vacuum-cast nickel (99.97 pct Ni) c) Pure nickel containing internal voids (produced by a process of intering). d) A substitutional solid solution alloy of pure nickel containing 1.16 pct Sn. All the materials were given a high-temperature anneal (>850°C) to produce a stable grain size of 7 to 10 grains per linear mm. The apparatus and techniques used to obtain the present results have been described elsewhere.&apos; In this series, grain growth was observed only in the highest-purity, vacuum-cast metal. Thus we may consider that in the other materials grain-boundary migration was prevented by either preexisting voids or solute atoms. The effect of migration on the creep ductility of the above materials can be seen from Fig. 1, which

Jan 1, 1962