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By R. M. Treco

The effect of small additions of oxygen on the hardness, density, and lattice parameters of high purity zirconium has been investigated. Precision parameter values and density of oxygen-free zirconium are given as: a,, - 3.2258 kX units, C = 5.1373 kX units, O23 = 6.507 g per cc. THE effect of oxygen on the hexagonal-phase lattice constants of pure zirconium is of interest for several reasons. The extensive hardening effect of oxygen noted by previous investigators' must be due in part to the lattice distortion accompanying its introduction and possibly to the preferred position of the oxygen atoms. De Boer and Fast' have shown that oxygen is taken into solid solution in large amounts, appearing to occupy an interstitial position. These authors also found that expansion of the lattice was not uniform since the a axis attained a constant value after the addition of 10 atomic pct while the c axis continued to expand. It was thought in the present work that this expansion of the zirconium lattice with increasing amounts of oxygen

Jan 1, 1954

By G. W. P. Rengstorff, L. E. Olds

PREVIOUS research has shown that cast molyb-denum has good ductility at room temperature if the metal is sufficiently pure. In practice, however, it is very difficult to reduce the impurity content of molybdenum to a few parts per million—the maximum content of impurities allowable to obtain reasonable ductility in coarse-grained molybdenum. Also, subsequent fabrication may reintroduce enough impurity to damage ductility greatly, unless extreme precautions are taken to protect the metal during working. The major impurities which embrittle molybdenum are oxygen, nitrogen and, to a lesser degree, carbon. It is reasonable to assume that the harmful effects of such impurities might be neutralized by small amounts of alloying additions. The desired alloying elements would be those which might fonn stable metal oxides or nitrides that would not be harmful to ductility, or thosewhich might increase the tolerance of molybdenum for oxygen and nitro-

Jan 1, 1958

By R. S. French, W. R. Hibbard

FOR tensile deformation, if the stress value is defined by the ratio of the load to the actual area, and the strain value by the natural logarithm of the ratio of the immediate length to the original gauge length of the sample, the resulting data when plotted with logarithmic coordinates result in a linear relation extending from the initial plastic yielding through the maximum load and in some cases, such as for copper and brass, to fracture'.'. The normal form of the curve, therefore, is a parabola, satisfying the following equation: S = Kem [1] where S and e represent true stress and strain, K is a constant, and m is a coefficient evaluating the slope of the curve and designated therefore as the strain hardening coefficient. Hollomon² has reported that this coefficient is a function of the yield strength as effected by grain size in alpha brass, and the carbon content and heat treatment of steel. In the case of steel, the value of the proportionality constant is a function of the carbon content. It was further shown in that paper2 and also by Low and Prater -hat, if the power relation between stress and strain remains valid, the strain hardening coefficient m, is equal to the strain at the maximum or necking load. Thus, an added significant feature is that this coefficient may serve as an index of the capacity for deformation prior to localized necking and subsequent fracture, e.g. drawability. The constant K may be defined as the strength of the material at unit strain. Together the K constant and the m coefficient describe an alloy's physical properties as far as strength and ductility are concerned for a particular yield strength and, in the case of steel, a particular carbon content. The work of Lacy and Gensamer4 gives data on alloy ferrites from which very close estimates of the strength of ternary and higher iron alloys have been made. It was the purpose of this investigation to develop similar data for commercial copper base alpha solid solution alloys. Preparation of Specimens Commercial binary alloys of copper with aluminum, beryllium, cadmium, nickel, silicon, tin and zinc were investigated. In addition, laboratory castings were made of a 1 pct silicon alloy and a 2 pct aluminum alloy to complete these particular series. Two types of copper were used, oxygen free OFHC and tough pitch FC. The analyses of these alloys are shown in table I. Standard ASTM sheet tensile specimens were used to determine the usual physical properties of each alloy and to obtain stress data at strains from 0.0005 to 0.01. Small (1/4 in. x 3 in. x tk.) samples

Jan 1, 1951

By R. S. French, W. R. Hibbard

M. Balicki—As one who some years ago spent much time searching for an alloy with high thermoelectric power that would be suitable for heat energy-electric energy converter based on the principle of a thermoelectric pile, I would like to make a few comments touching upon that aspect of this excellent paper. The specifications which one of the authors has formulated so aptly do not appear to me as complete, and I would like to add to them the following two demands: (1) Melting point not less than 500°C and (2) elongation on 2 in. at least 5 pct. Condition 1 would insure good thermodynamic efficiency of the pile by permitting a reasonably high temperature of the hot junction and at the same time making it suitable for such application as salvage of vast amounts of energy which is lost in the form of sensible heat of waste gas leaving industrial and other furnaces. Condition 2 is also a very important one; particularly, because of the fact that all elements and alloys which exhibit attractive thermoelectric properties seem to be as a rule very brittle. The design, erection, and maintenance of a pile incorporating a brittle material in one of the most crucial parts of the converter would be a very difficult and costly affair indeed. To illustrate this point I may mention great difficulties I encountered when trying to make thermocouples of silicon. Casting of 1/4-in. rods in iron or graphite mould yielded only broken pieces which had to be welded into a bar. Such a bar had to be handled very carefully during the determinations. Raising the molten silicon by vacuum into a silica tube and solidifying it there solved the problem of stability as well as that of electrical insulation of a bar. On tests this bar (98.5 pct Si) showed dE/dT = — 275.5 microvolts per degree C vs. copper. Thermocouples made of this grade of silicon and pure nickel developed 0.39 v when temperature of the hot junction was approximately 1000°C. The possibility Dr. Gonser has mentioned of trying to impart to metals good thermoelectric properties by alloying them with silicon, has been considered. The preliminary survey which I have made was, however, disappointing. Perhaps, it will be advisable to record the results obtained (table VI), as this might save repetition of search in fields already covered.

Jan 1, 1951

By J. W. Rutter, K. T. Aust

THE object of the experiments to be described in this report was to determine, first, which grains, out of a large number introduced into a sample in which their growth could proceed, were able to grow most successfully; and, second, the factors which control their growth selection. In an earlier study of grain boundary migration in high-purity lead,' it was found that, under some conditions, the most successful grain in an experiment of this type bears a special orientation relationship to the grain into which it is growing. The present work represents an extension of the previous study, directed at the determination of the orientation relationships which appear under a wider variety of conditions. Experiments of this general type have been done using commercial aluminum2 with somewhat conflicting results. In the present experiments, in contrast to the above,' the sample composition was a principal controlled variable, resulting in a considerable clarification of the phenomena observed. EXPERIMENTAL TECHNIQUE The experimental technique used in this investigation was essentially that employed previously in studies of grain boundary migration kinetics.1'3 The material used was zone-refined lead with various small amounts of tin or silver added to study the effect of varying the sample composition. The specimens were in the form of single crystals, about 1/4 in. 1/4 in cross section by 3 to 4 in. in length, grown from the melt in a horizontal graphite boat. A crystal grown in this way generally contained a lineage or "striation" substructure of the type shown in Fig. 1 of Ref. 1. This substructure consists of an array of low-angle boundaries which partition the crystal into regions misoriented from each other by a few degrees. The substructure constitutes a source of driving energy, estimated to be about 4000 ergs per cc, for the migration of a large-angle grain boundary into the striated crystal. New grains were introduced by plastic deformation in compression of a localized reglon at one end of the sample. Re-crystallization occurs at room temperature in the deformed region, introducing a large number of new, striation-free grains into the specimen. These grains are able to grow into the melt-grown crystal on annealing at a suitable temperature. The stria-tions are removed by passage of the grain boundary, as the substructure energy is utilized to produce the boundary migration. Usually, only one or two of the recrystallized grains persist after growth has occurred for about 1 cm out of the deformed region. Laue back-reflection X-ray photographs were obtained to determine the orientation relationship between the melt-grown crystal and the recrystallized grain or grains which grew most successfully in each specimen. EXPERIMENTAL OBSERVATIONS Fig. 1 shows the results obtained for a series of specimens to which had been added tin in the range from 0 to 5 ppm by weight. In this figure is plotted, in the standard stereographic triangle, the single axis about which the most successful recrystallized grain or grains are related to the starting, melt-grown crystal by the smallest rotation. The range of angular rotations observed about such axes is indicated beside each stereographic triangle. The results shown in the stereographic triangle on the left, Fig. l(a), were obtained by localized deformation at one end of each specimen at room temperature, followed by annealing in a furnace which had been preheated to 300°C. The furnace atmosphere was argon. It will be seen that the orientation relationships observed between the melt-grown and recrystallized grains are quite random. On the right, Fig. l(b), are shown results for samples in which the de-

Jan 1, 1961

By J. Philibert, A. G. Guy

Diffusion in the presence of deformation was studied by the method of vacuum dezincification of copper-rich and silver-rich solid solutions containing 7 to 30 pct Zn. The specimens were designed to permit the study of diffusion in separate portions of a given specimen characterized by strain rates ranging from essentially zero to approximately 10 sec-. No effect of deformation on diffusion was observed. BEGINNING with the work of Buffington and Cohen: interest in the question of the effect of stress or strain on diffusion has largely been concentrated on the enhancement of diffusion in specimens subjected to Continuous plastic deformation. The present research is a contribution to this limited area. However, as a preliminary to focusing attention on this special topic, it will be desirable to make a broad survey of the larger question, especially since there has been considerable foreign work in areas outside those of current interest in the United States. Since most of the topics referred to in the following section are both complex and imperfectly understood at present, it has been expedient in most instances to offer only a guide to the general nature of the work rather than a critical evaluation. PREVIOUS WORK The effect of elastic stress on diffusion has received considerable attention, especially with regard to the thermodynamic driving force for diffusion. The thermodynamic treatments have been based on the work of Gibb, Voigt, Planck, and Leontovich.' Konobeevskii and Selisski6 made a first attempt at treating the problem in 1933, and Gorskii7 a few years later gave a solution applicable to single crystals as well as to polycrystalline specimens. In 1943 Konobeevski8 published treatments that have been the basis of much Russian work up to the present. For example, Aleksandrov and Lyubov used his work in explaining the velocity of lateral growth of pearlite. Early work in the United States was that of Mooradian and Norton, which showed that lattice distortion tends to be relieved before it can significantly affect the diffusion process. Druyvesteyn and Berghoutl1 observed a slight effect of elastic strain on self-diffusion in copper, while de Kazinczy12 found that both elastic and plastic deformation increased the rate of diffusion of hydrogen in steel. On the other hand, Grimes58 observed no effect of either elastic or plastic straining on the diffusion of hydrogen in nickel. High-frequency alternating stresses have been reported by various investigator s13-l5 to increase the rate of diffusion. A special form of elastic stressing is the imposition of hydrostatic pressure, a condition that is amenable to Conventional thermodvnamic analysis. Most of the experimental results in this area are consistent in showing a slight decrease in diffusion rates at high pressures.16-l8 Although Geguzinl reported a pronounced effect of relatively small pressures, Barnes and Mazey20 failed to Corroborate this finding, while Guy and Spinelli21 advanced an explanation of the phenomenon observed by Geguzin. It has been recognized that the thermodynamic treatment of diffusion phenomena in an arbitrarily stressed body is complicated by the fact that the desired state of quasi-equilibrium of the shear stresses cannot be maintained during a general diffusion process. However, attempts have been made by Meix-ner22-24 and Fasto to treat certain restricted cases, such as relaxation. FastovZ7 has also incorporated the general stress tensor into the thermodynamics of irreversible processes. The lattice strain that accompanies the formation of a solid solution has been the subject of much study,28-s0 and indirectly it has entered into many recent theories of diffusion. However, some Russian investigators31'32 have taken other views of this matter and have predicted large effects on diffusion rates because of concentration stresses.o In completing this brief resume of previous work involving elastic strains and before proceeding to a consideration of the effect of continuous plastic deformation, it should be pointed out that deformation of various additional types may also influence diffusion. The effect of cold-working on subsequent diffusion has been studied directly by AndreevaS and by Schumann and Erdmann-Jesnitzer, while indirect evidence has been obtained by Miller and Guarnieri and by Vitman.38 Thermal stresses may also influence diffusion, contributions to this subject having been made by Fastovs7 and by Aleksandrov and Lyubv. The work of Johnson and Martin,o Dienes and Damask,3Band DamaskS considered the question of radiation-enhanced diffusion. In considering previous work on the subject of plastic deformation and diffusion, attention will be directed to those studies concerned primarily with diffusion rather than with its relation to Creep, e.g., the work of Dorn, or to the acceleration of diffusion -controlled reactions. Observations of the effect of

Jan 1, 1962

By W. J. Rogers, L. J. Dwiggins, R. O. Williams

THERE has been comparatively little work on the importance of strain rate and temperature as variables in the formation of deformation textures. For this reason the present work was started, the choice of compression being in part that of convenience and in part because of the absence of quantitative data in this area. These do not turn out to be important variables within the range studied which permits the results to be condensed into this note. Samples 1/4 in. in diam and 1/2 in. long were machined from 99.9 pct pure commercial aluminum and recrystallized to a fine grain size. A set of five samples were deformed between a spring-driven weight and an anvil resulting in roughly 20 pct strain each time. The strain rate varied because of decle-ration and shortening of the sample but would average about 10&apos; per min. Deformation at temperatures over the range of -76" to 150°C were obtained by heating or cooling the samples. Complete lubrication could be obtained by oil since the short time of contact did not permit its escape. The samples were compressed to a total true strain of 1.4 to 1.6. A larger sample was prepared similarly with spiral grooves in the ends to hold paraffin which provided essentially complete lubrication. This sam ple was compressed slowly at room temperature at a rate of about 2 per min to a total strain of 1.0. Cylindrical samples were carefully machined from these discs with a radial direction as an axis. They came from the center half of the samples and were roughly 1/16 in. in diam. The samples were etched down 15 pct to remove any surface effects. Diffraction data were then obtained for six different kinds of planes by the use of the Norton technique as applied to a fiber texture. Direct comparison between the raw data failed to reveal any significant difference between any of the samples and consequently the data from only one sample was completely reduced. The particular sample had been deformed at O°C The background was subtracted from the intensity data which were averaged over the four quadrants and then normalized to random. These data were then used to obtain fiber-axis distribution charts as previously described.2 For this sample the (422) plane was also run to assist in this inversion. The final result, Fig. 1, provided a comparable fit to that previously obtained from fiber data. About 85 pct of the material forms an elliptical texture around the <110> direction. The radius (at half height) is 5.5 deg along the (111) boundary and 8 deg along the (100) boundary. The remaining 15 pct forms a ridge running from this texture to near the <311> direction. The regions around the (100) and (111) planes are devoid of material. The earlier qualitative results of ~arrett~ are in complete agreement with the present results. The possibility has been dis-cused that the material around the <311> results from friction between the sample and compression plates.4 The present results are considered proof that such is not the case because of the uniformity of sample deformation and the requirement of axial symmetry before fiber-axis distribution charts can be accurately obtained. Recently results have been obtained to show that strain rate and temperature are important variables in the extrusion of aluminum.5 Perhaps the most important observation was that the partial recrystal-lization occurred even at room temperature. Strain rate was considered to be a factor only in that it produced a temperature change. There are many differences between the two experiments, the most important possibly being that their material was roughly ten times the purity of the present material. The present experiment, however, may not be as definitive as desired in this respect since it has been shown that compressed aluminum recrystallizes without a texture change6 although supplementary

Jan 1, 1961

By W. L. Phillips

Stress-strain data are reported for sodium chloride, lithium fluoride, and magnesium oxide tested in compression at diffevent strain rates in the telrlperature range from 70° to 873°K. Quantitative equations are given for the stvain-rate dependence of the critical resolved shear stress and the initial rate of work havdening. The properties so determined are compared with results previously reported for metals. ALTHOUGH the effects of strain rate and temperature on the stress-strain characteristics of metals have been extensively investigated1,2 these effects have not been studied in detail in ionic crystals. In the only previous investigations on the effect of strain rate on the deformation of ionic crystals. Gilman and gohnston3 and Hulse and pask4 have shown that the yield stress of lithium fluoride and magnesium oxide at room temperature is proportional to the ninth or tenth root of the strain rate. CRYSTAL PREPARATION Lithium fluoride, sodium chloride. and magnesium oxide single crystals were purchased from the Har-shaw Chemical Co. The compression fixture and sample preparation have been described in detail previously.5 The detectable impurities in MgO were eight parts in 105 of Fe and six parts in 105 of Si; in LiF, six parts in 105 of Mg; and in NaC1, two parts in 105 of Fe and four parts in 10&apos; of Mg. The cleaved samples measured - 0.1 in. by 0.1 in. by 0.3 in. Prior to testing all samples were annealed at -0.8 of their melting point and cooled to room tempera- ture at 50 C per hr.&apos; Samples were strained at rates (<) of 0.1, 0.01, 0.002, and 0.0002 in. per in. per min. EXPERIMENTAL RESULTS Fig. 1 is a plot of the critical resolved shear stress a,. (extrapolation of slope of easy glide region to zero strain) of sodium chloride, lithium fluoride, and magnesium oxide tested at different strain rates as a function of temperature. In all crystals at all strain rates a, decreases with increasing temperature according to an equation of the type397 where uo is a stress constant, K is a constant, and T is the absolute temperature. The points in Fig. 1 are the average of six tests. It is evident that there is a break in the log 0,- temperature plots at a characteristic temperature T,. The constant K of Eq. [I] is dependent on strain rate, Fig. 2. The relationship

Jan 1, 1962

By N. Jaffe, J. E. Dorn

This investigation is concerned with the possibility that the creep resistance of metals might be reduced as a result of recovery at the creep temperature when the applied stress is reduced. For this purpose the effect of stress on the recovery of the creep behavior of pure polycrys talline aluminum was investigated in the cross-slip temperature range where the activation energy for creep is 28.000 cal per mole. Before recovery all specimens were given the same substructure by creeping them under a stress of 7800psi to a strain of 0.218. Recovery was then accomplished by holding the specimens at the creep temperatures of 327" and 337°K under a series of reduced stresses for various periods of time. When the original stress was reapplied the specimens crept more rapidly than when the original stress was first removed, illustrating that recovery did take place. The extent of recovery could be indexed in terms of the UNDER certain conditions of test, the creep rate decreases over the primary stage of creep often by as much as a factor of 10~B. Since the external variables of stress and temperature are held constant, such pronounced and significant changes in the creep rate must be attributed to various substructural modifications that are introduced as a result of creep itself. At the secondary stage of creep, over which the creep rate remains constant and the substructure no longer changes, a balance must be achieved between the rate of increase in creep resistance due to creep straining and the rate of recovery of the substructural modifications. The general validity of these concepts on substructural changes has already been shown,&apos; including the corollary that a more creep-resistant substructure is introduced under a more rapid creep rate arising from a higher applied stress.&apos; In spite of these well-documented facts, very little progress has been additional strain undertaken by the recovered specimen to achieve the same strain rate as the unrecovered specimen. It is significant that the higher recovery stresses gave the most rapid rates of recovery. Over the high range of reduced stresses, where some creq took place during recovery, the activation energy for recovery equaled that for creep, namely 28,000 cal per mole. But for stresses during recovery below 0.80 of the original creep stress, no creep was detected during recovery and an activation energy for recovery of about 72,000 cal per mole was obtained. Whereas the activation energy of 28,000 cal per mole suggests that cross-slip of dislocations is responsible for this recovery, the mechanism for the 72,000 cal per mole process is not known. This activation energy, however, is about twice that for self-difhsion. made in determining, in satisfactory detail, precisely what substructural changes so effectively modify the creep rate.&apos; Up to the present, electron microscopy has failed to uncover, in a satisfactory quantitative manner, the mechanistic processes involved. Since, however, creep, at least at moderate and high temperatures, arises from one or more types of dynamic recovery processes, it should be possible to shed additional light on the total problem by investigating the recovery of creep resistance. The importance of such an investigation becomes apparent when the fact is recognized that there has yet been no direct experimental proof that the rate of softening due to substructural changes as a result of recovery equals the rate of hardening due to creep-induced substructural strengthening over the secondary stages of creep. Furthermore the problem of what happens to the creep-strain induced strengthening when the material recovers either at zero or some reduced operating stress has significant engineering consequences. It was for these reasons that the present investigation on the recovery of creep behavior was initiated. Inasmuch as a substantial background has now developed on the creep of high-purity aluminum, this material was again selected for the current investigation. Furthermore extensive evidence is now

Jan 1, 1963

By N. Jaffe, J. E. Dorn

The effect of stress on the creep rate of high-purity poly crystalline aluminum in the intermediate temperature range was investigated by conducting creep tests at a series of different stresses and by decreasing the stress abruptly during the course of primary and secondary creep. Three series of decreases in stress tests were conducted: the stress was decreased in primary and secondary regions of creep at 347.5°K and the stress was decrensed in primary creep at 327°K. The results demonstrate the effects of structure and temperature on the stress dependence of the creep rate. The trends in the creep rate after a drop in stress are complicated and show that recovery takes place. The experimental results were compared with Friedel&apos;s theory for cross-slip, the experimentally determined energy for a constriction being in good agreement with the theoretical estimate proposed by Stroh. PREVIOUS investigations have shown that the ap-- parent activation energy for creep of polycrystalline aluminum, over the range from about 250° to 380°K, is approximately 27,400 cal per mole,&apos; a value that is to a first approximation insensitive to the applied stress, the creep rate, and the creep strain. Several facts strongly support the thesis that creep, in this temperature range, is controlled by the rate-determining mechanism of cross-slip: 1) Recovery of the cold-worked state can take place by a number of mechanisms prominent among which is climb and cross-slip. Astrom2 reported rapid recovery of cold-worked polycrystalline aluminum at 340° to 370°K giving an activation energy of 28,000 cal per mole and additional recovery at the higher temperatures of 455° to 478°K giving an activation energy of 36,000 cal per mole. Whereas the latter process was clearly identified with the climb mechanism, the former process. which has an activation energy in good agreement for that of creep in the same temperature range, is most likely due to recovery by thermally-activated cross-slip. 2) Single crystals of aluminum exhibit sharp slip-

Jan 1, 1962

By R. Frenkel, O. D. Sherby, J. E. Dorn, J. Nadeau

A method is shown for the study of the creep rate dependence of metals on the applied stress under the condition of constant structure. The method was applied to pure aluminum and to dilute solid solution alloys of magnesium, copper, germanium, zinc, and silver in aluminum. A linear relationship was found to exist between the true stress and the logarithm of the creep rate. ALTHOUGH extensive creep data are now avail-Aable, the fundamentals of the creep phenomena are as yet only vaguely understood. Many attempts have been made to formulate theories for creep. Eut their agreement with fact has been disappointing. It appears that merely pyramiding more of the same type of data that is now available will not provide the essential aid toward formulating a better understanding of the complicated process of creep. Perhaps attempts to construct better models for the mechanism of creep might have to be tentatively abandoned until new types of definitive creep data are uncovered. The question therefore arises regarding what type of data might be needed to provide the essential knowledge for formulating better and more realistic theories of creep. Perhaps one source of the failure to uncover an adequate theory arises from the fact that practically all theories for creep and nearly all analyses of experimental data on creep disregard the well-known fact that the structure of metals changes during creep. For example, it is customary in evaluating the effect of stress on the creep rate to correlate the secondary creep rate with the applied stress. But, as has been demonstrated recently,&apos; the structures obtained during secondary creep tend to exhibit deformation banding at high stresses and polygoniza- tion at the lower stresses. Consequently the stress-secondary creep rate relationships found in this way are not fundamental, inasmuch as they reveal not only the effect of stress but also the effect of differences in structure on the secondary creep rates. It is the purpose of the study reported here to isolate the effect of stress alone on the creep rate in an attempt to provide new data for possible formulation of a better theory of creep. In addition an attempt will also be made to ascertain how structure, temperature, and alloying might affect the stress-creep rate relationship. The procedure that was adopted was simple: A specimen was precrept at a given engineering stress U, (initial engineering stress during creep under constant load, namely the load divided by the initial area of the specimen) and temperature T to a selected engineering strain, e (engineering strain, namely, the instantaneous gage length minus the initial gage length divided by the initial gage length), at which time the true stress was reduced to some lower value of the true stress, ^ (the true instantaneous stress, namely, the load divided by the instantaneous area). The instantaneous true creep rate, i, (the true strain rate; with E, the true strain, namely, the natural logarithm of the instantaneous over the initial gage length) following reduction of the true stress to U, was then determined. Second, third, etc., tests were conducted under identical conditions except the stress was reduced to yet lower values, us, u,, yielding yet lower instantaneous creep rates, ;, and i,. Inasmuch as the precreep conditions were identical in each series of tests, the instantaneous structures obtained immediately after reducing the stress were presumed to be identical.

Jan 1, 1955

By Norman Brown, J. V. Rinnovatore

It was found that a shear stress applied during the recovery had no effect on the amount of recovery if the stress was less than the instantaneous yield point irrespective of the direction of the stress. Large ajplied stresses during recovery had a marked effect, which depended on the direction of stressing. Single strain reversals at -72°C not only produced the Bausch-inger effect but also increased the flow stress. The increase inflow stress was attributed to vacancies, which have appreciable mobility at -72 °C. THEORIES I-&apos; of recovery suggest that the internal stress field of the dislocations plays an essential role. Most investigations of recovery use time and temperature as the independent variables; however, here stress has been introduced as a variable in two ways: 1) by varying the internal stress field according to the amount of strain prior to recovery and 2) by applying an external stress during recovery. Tannen- baum and Kauzrnann studied the effect of applying a stress during recovery by measuring the change in creep rate after recovery; they found that stresses as high as 95 pct of the creep stress had no effect on recovery. Earlier, Wood and scrutton6 showed that an applied stress during creep increased the rate of poly-gonization. The experiment by Wood involved poly-crystalline aluminum whereas Tannenbaum and Kauzmann used a single crystal of zinc. In this investigation, zinc single crystals were used and deformed in a manner to prevent polygonization. Thus, the main difference between this investigation and that of Tannenbaum and Kauzmann is the use of the stress-strain curve instead of creep rate as the index of recovery. In addition, the applied stress during recovery varied

Jan 1, 1961

By A. L. Mincher, W. F. Sheely

Mechanical properties of columbium have been studied over the temperature range of -196 to 1093oC. The decreased strengthening influence of cold-work at temperatures below ambient has been interpreted in terms of the Peierls-Nabarro effect. Maxima in the rate of strain hardening observed during tensile testing in the range 250-600°C. have been correlated with interstitial impurities to indicate the temperature ranges at which carbon, oxygen, and nitrogen, respectively, are responsible for strain aging. THE growing need for structural materials for use above the useful service temperatures of the iron-, nickel-, or cobalt-base alloys has caused the refractory metals to be considered as potential engineering materials. These metals, which include columbium, tantalum, molybdenum, and tungsten, are called refractory because the lowest melting point among them,that of columbium, is about 1000°C higher than the average melting temperatures of conventional high-temperature alloys. They are all body-centered cubic transition metals and, as such, their mechanical properties have basic characteristics which distinguish them from the face-centered cubic metals. For example, all show a much steeper rise in strength with decreasing temperature below room temperature than do the face-centered cubic metals, and their mechanical properties are strongly influenced by interstitially dissolved impurities. In order that these new metals may be used efficiently, it is necessary that their characteristics of behavior be fully known. In this paper, the mechanical properties of columbium will be examined over a wide range of temperatures. In particular, the influences of cold-work and individual species of interstitial impurity atoms on mechanical properties will be described, and basic mechanisms which may control the observed characteristics will be explored. EXPERIMENTAL The material used in this investigation was Union Carbide Metals Co. columbium roundels consolidated to four 4-in. diam ingots, three by consumable-electrode arc melting and one ingot by electron beam melting. Impurity contents of the ingots and methods of ingot conversion and treatment are summarized in Table I. The only metallic impurity occurring in any significant quantity was tantalum at about 0.1 pct. Iron, silicon, titanium, and zirconium were each less than 0.015 pct; boron was 1 ppm or less. This should have no appreciable influence on properties. The electron beam melted material, being the purest, will be used as the basis for comparison in the discussions to follow. Tensile tests were conducted from-196 to 1093oC, on both cold-worked and fully recrystallized arc-melted and electron-beam melted columbium using standard 1/4-in. diam, 1-in. long gage length test specimens. A strain-rate of 0.005 in. per in. per min was employed until the 0.2 pct yield strength was achieved and then the strain-rate was increased to 0.05 in. per in. per min for the balance of the test. Samples were protected in an inert atmosphere at tests above 300°C. The tensile properties obtained on the electron-beam melted columbium, E, in both the cold-swaged and recrystallized conditions are given in Fig. 1. The yield strength data of Dyson, et al.,&apos; obtained on recrystallized electron beam melted columbium and the tensile strength data reported by Tottle2 on powder metallurgy columbium are included in Fig. 1. The material used by Tottle had been purified by vacuum sintering. There is excellent agreement between Dyson&apos;s data and those obtained in the present investigation. The tensile strengths obtained by Tottle were slightly greater than those obtained in this investigation on electron-beam melted columbium but varied with temperature in a similar manner. Tottle&apos;s data showed a maximum in tensile strength near 500°C, as did our data on electron-beam melted material, and also showed a small maximum at 300°C. The significance of these maxima will become evident later in the discussion. The tensile properties of cold-swaged and recrys-tallized arc melted columbium are plotted in Fig. 2. It was found that the properties of the recrystallized arc-melted columbium from all three heats showed very close agreement except at temperatures between about 500" and 800°C. A reason for this range of disagreement will be suggested in the discussion. The generally good agreement, however, attests to the ability of cold-working and subsequent recrystal-lization to erase the effects of the three different primary breakdown procedures and to produce nearly equivalent structures in the samples derived from the three different heats. wesse13 reported tensile data on columbium having interstitial impurity contents between those of the

Jan 1, 1962

By W. A. Wood, H. M. Bendler

Cylindrical fatigue specimens of OFHC* copper were subjected to alternating torsion while under axial tension. The superimposed tension strongly influenced the fatigue life of the specimens as well as causing characteristic changes in the structure of the fissures that form and develop throughout the life of the specimen. A further effect of the superimposed tension is a continuous extension of the specimen throughout its life in a manner formally similar to creep at elevated temperatures. RECENTLY, studies of the structural changes leading to fracture of metals by fatigue have been made by a technique which reveals, in section, the lattice distortion arising from slip movements inside the metal and the disturbances built up by these slip movements at the metal surface.1&apos;2 According to this approach, fatigue develops as follows: a) fine slip movements concentrate in bands, producing little strain-hardening and thus permitting abnormal totals of plastic strain before fracture and, therefore, abnormal amounts of lattice distortion along each slip band; b) both forward and backward slip movements occur in each band which, according to their distribution across the band, may build up surface disturbances that include sharp notches; c) these notches early in the life of a specimen penetrate the distorted slip zones and transform them into deep, fine fissures; d) the fissures on continued cycling locally saturate the metal and cause local collapse in the form of a microcrack. Thus, the significant consequence of fatigue or cyclic

Jan 1, 1962

By C. H. Li, R. J. Stokes, T. L. Johnston

Dislocation half-loops, artificially introduced by sprinkling with carborundum, were subjected to stress using three-point loading. The different stages of loop expansion and multiplication were then correlated with the macroscopic stress-deflection curve. Loop expansion started at approximately one-third of the "yield stress," defined here as the motivation stress. Above one-half of the "yield stress" simple multiplication commenced while about two-thirds of the yield stress gross multiplication developed short slip-line segments. These values were not affected by the age of surfaces exposed to both dry and wet air. Slip was not detected on unsprinkled crystals below two-thirds of the "yield stress." Thus the stress to activate intrinsic sources, defined here as the activation stress, was approximately twice the motivation stress. This relationship was independent of the strain rate used. It is concluded that lattice resistance takes the form of a viscous drag rather than a static type of force. THE importance of surface condition and environment as one of the more significant factors controlling the mechanical properties of ionic solids has recently received renewed attention.1-4 In some materials, notably the alkali halides, exposing freshly cleaved specimens to the atmosphere results in a progressive hardening and embrittlement of the crystals. This has been attributed to the surface layer becoming "aged," thereby restricting the motion of dislocations in such a way that pileup and eventually fracture results. It is necessary therefore in studying the mechanical properties of ionic crystals to inquire first into the density and distribution of dislocations in the normal cleaved surface and to understand the role these dislocations play in the onset of plastic flow. Knowledge of the effect of environment on this behavior will then assist in understanding the mechanism of embrittlement on aging. In the work reported here a method is described for studying the initiation of plastic flow from dislocation sources in the surface of magnesium oxide crystals. Preliminary results are described of the influence of surface age on this behavior. EXPERIMENTAL TECHNIQUES The magnesium oxide used in this investigation was obtained from the Norton Co. and had a nominal purity of 99.5.pct. Single-crystal specimens, having dimensions typically 1/4 by 1/8 in. cross section and 1 in. in length, were prepared by cleavage over (100) planes from the parent crystal. They were deformed under three-point loading with 3/4 in. spacing between the two end supports and a constant rate of deflection at the center of 0.001 in. per min. The maximum strain rate therefore was 1.5 X 10-5 sec-1. Reasons for choosing this particular form of loading will become apparent in a later section. In addition, a number of confirmatory type tests have been performed at higher loading rates where the maximum strain rate was increased to 1 x 10-3 sec -1. Following deformation the crystals were etched to reveal the dislocation distribution. In a previous paper5 we described a satisfactory etching solution consisting of 5 parts of saturated ammonium chloride, one part of sulfuric acid, and one part distilled water. Immersion of the crystals for 1/4 hr at room temperature developed reproducible conical etch pits at the points where dislocations emerged through the surface. It should be mentioned that dislocations freshly introduced by room temperature deformation gave sharply pointed etch pits which generally distinguished them from the almost flat-bottomed pits obtained on "grown-in" dislocations. The distinction is obvious in the photomicrograph included in Fig. 5(b). Annealing crystals at 1000°C for 12 hr on the other hand gave sharper etch pits on the grown-in dislocations and the distinction was no lower -possible; a similar observation has been reported for lithium fluoride. We have succeeded furthermore in chemically polishing magnesium-oxide crystals. Boiling in fresh 85 pct ortho-phosphoric acid for only 1 min resulted in a microscopically smooth surface which could again be etched to reexamine the dislocation distribution. This boiling time was sufficient to remove the surface damage introduced by cleavage (see next section). SURFACE DISLOCATION SOURCES There are a number of ways by which dislocations may be introduced into cleaved surfaces to act later

Jan 1, 1960

By N. Jaffee, J. E. Dorn

An activation energy of 27,400 5 1000 cal per mole was obtained for the creep of poly crystaLline aluminum over the temperature range of 273° to 350°K, at strains varying from 0.003 to 0.230. Stresses varied fronz 2250 to 6000 psi, strain rates from 0.1145 x 10-5 to 29.5 x 10~s min.-1. A strain temperature-compensated time correlatzon was oblained indicating that the structure was depedent solely upon the strain at a given stress; X-ray diffraction analyses and tensile tests on precrept specimens verified this. PREVIOUS investigations1 have revealed that the apparent activation energies for creep of high-purity polycrystalline aluminum are insensitive to the applied stress, strain, and strain rate. On the other hand, the apparent activation energy is dependent on the temperature1,2 as shown by the data reproduced in Fig. 1. Each plateau of this curve is associated with a unique rate-controlling dislocation mechanism for thermally activated creep. It is now very well-established that over the high-temperature range, where the activation energy is about 35,500 cal per mole, the creep rate is controlled by the dislocation climb mechanism. 3 Over the intermediate range of temperatures, from about 260° to 370°K, another unique activation energy, namely 28,000 cal per mole, is obtained. Current evidence strongly supports the contention that cross slip is the rate controlling mechanism for creep in this case: 1) Crystal recovery can occur by a number of mechanisms prominent among which are climb and cross slip. Astrom,4 using the Borelius micro-calorimeter, reported rapid crystal recovery of cold-worked polycrystalline A1 at 340° to 370°K giving an activation energy of 28,000 cal per mole and additional recovery at 455° to 478°K giving an activation energy of 3 6,000 cal per mole. Whereas the latter recovery process was clearly associated with the dislocation climb mechanism, the former, 28,000 cal per mole process could have arisen as a result of the cross-slip mechanism. 2) Single crystals of high-purity A1 favorably oriented for octahedral glide exhibit an activation energy for creep of about 3400 cal per mole over the range from 0° to 400°K and that of about 28,O:O cal per mole over the range from 600° to 775°K . Whereas the lower activation energy process is accompanied by the development of sharp slip-band traces, the 28,000 cal per mole process was uniquely characterized by extensive cross slip. 3) Theoretical calculations by Seeger and schoeck6 suggest an apparent activation energy for cross slip in A1 of about 24,000 cal per mole and another theoretical estimate by Friedel 7 gives the observed 28,000 cal per mole as the theoretical activation energy. It is therefore consistent with this evidence to, at least tentatively, ascribe the 28,000 cal per mole creep process in A1 to the operation of cross slip as the rate controlling mechanism. Although creep has been rather well-explored in the dislocation climb range, very little, excepting perhaps the activation energy, is known about creep in the cross slip range. This investigation was undertaken, therefore, to reconfirm the activation energy for creep in polycrystalline A1 in the intermediate temperature range and to study the strain-time relationship for creep at a constant stress in the cross slip region. EXPERIMENTAL PROCEDURE AND TECHNIQUE The three types of high-purity A1 alloy specimens described in Table I were used in this investigation. Specimens were milled from 0.100 in. thick cold-rolled sheets with their tensile axes in the rolling direction following which they were given a recrys-tallization and grain growth heat treatment at temperatures well above those which were subsequently employed in the creep tests.

Jan 1, 1962

By Hans Conrad, W. D. Robertson

FORMAL theories of yielding and flow have been developed to describe thermally activated plastic flow in metal crystals.&apos;-" hey are all based on Becker&apos;s&apos; original idea that thermal fluctuations aid the applied stress in overcoming internal barriers to flow and lead to the following relation. where j is the strain rate; C is a constant which includes a frequency factor, the entropy, the number of flow units, and the amount of strain associated with each successful fluctuation; E(T) is the activation energy associated with flow, and is a decreasing function of the applied stress, 7; k is Boltzmann&apos;s constant; and T is the temperature in OK. The various theories differ mainly in the model chosen and in the nature of the relationship between the applied stress and the activation energy. Beckerl and Orowan&apos; suggest a relationship of the form E a(Ti — t)2, where .ri represents the stress required for flow without the aid of thermal fluctuations; Kauzmann,3 Cottrel1,5 and Seeger6,7 suggest E a(t — t) ; Fisher8 suggests E while Cottrell and Bilby4 suggest a more complex relationship between E and 7. Additional data are needed to test these various theories., esweciallv- in view of Dorn&apos;s9 ork. which indicates that the activation energy for creep is independent of the stress and strain. The present paper presents the stress-strain data derived from incremental loading of magnesium single crystals in tension and direct shear, in the temperature range of 78" to 364°K. Incremental loading was employed in order that the creep component of the strain could be determined. Stress-strain curves of aluminum10-13 and copper" single crvstals. derived from incremental loading tests, have been found to be in agreement with those obtained by the more conventional constant strain rate or loading rate tests. Two types of tests were conducted in the present investigation: 1) tests in which the temperature was maintained constant throughout the entire

Jan 1, 1958

By R. S. Busk

A PREVIOUS paper1 has summarized the effect of A all metallic elements on the lattice parameters of magnesium. The present paper deals with the effect of temperature on the lattice parameters and the way in which that effect is modified by alloyed elements. The experimental techniques for sample preparation, exposure to the X-ray beam, measurement of line diameters, and calculation of the parameters were the same as reported in the previous paper. The temperature of the specimen was raised by surrounding the mounted compact with a conical Glas-Col,* the base of which was sealed with Saran film. In this way the temperature of the specimen could be raised to 200 °C, which was the maximum desired in the present study. Control of temperature was effected by maintaining a constant, desired voltage to the resistance coils of the Glas-Col through a Variac. Temperature was measured by an iron constantan thermocouple embedded in the body of the specimerl so that the thermocouple bead was within 1/16 in. of the surface exposed to the X-ray beam. Variation of temperature during an exposure was less than + 1/4 oF. A series of eight alloys, consisting of pure Mg, Mg+Ag, Mg+Al, and Mg+Sn were studied. All percentage values are in atomic percent. The results are given in Table I. The density values given in Table I were calculated from the lattice parameters, using an equation due to Foote and Jette:"

Jan 1, 1953

By R. S. Busk

Niels Engel (University of Alabama, University, Ala.)— In this paper it was pointed out that the electron-gas and energy-band theory accounts for the fact that the lattice parameters exhibit a sudden change when the electron concentration (number of bonding electrons per atom) exceeds a certain number around two. This statement is said to support and prove the electron-gas theory. But this theory is not able to account for a series of experimental data. Also several expectations, deduced from this theory, are not found to exist. In Figs. 6 and 7 the energy bands of the second and third periods are given as they must be assumed in order to account for the electrical properties of the elements in these periods. In Figs. 6 and 7 the electron-gas and energy-band theory is compared with the electron-oscillator hypothesis in accounting for the properties of the elements in the second and third periods. Fig. 6 shows the second period, The energy-bands are overlapping and separated to be in agreement with the electrical conductivity of the elements. The oscillator hypothesis explains conductivity due to electron vacancies. In graphite there is a closed s-shell in every other atom and two vacancies in the others. Conductivity is therefore only maintained by migration of s-electrons in graphite. In boron there are no s-electrons. The diatomic molecules of nitrogen and oxygen and the paramagnetism of oxygen can be accounted for by a similar behavior as the s-electrons of the bonding electrons. But this explanation will deviate too much for the purpose of this discussion. Fig. 7 shows the third period. In the energy-band picture about two s-electrons are assumed in magnesium and aluminum, but only one s-electron is assumed in silicon. The diamond lattice is assumed to be controlled by a sp3 hybrid. However the electron distribution develops ideally according to the oscillator hypothesis. Only sodium, magnesium, and aluminum exhibit electron vacancies and conductivity. To account for the insulator properties in Si, P, and S in the third period it must be assumed that the four last added p-electrons must be taken up in bands containing only one electron per band.&apos; (Compare the electron band picture in Hume-Rothery.&apos; Hume-Rothery does not consider the insulator properties of the nonmetals.) In the second period already the first p-electron must have entered a single electron band. Based on the energy-band picture in Figs. 6 and 7, the following questions must be asked: 1—Is it consistent with the energy-band idea that electrons of the same kind (p-electrons) can be divided into separated bands? 2—Is it consistent with the energy band idea that single electron bands can exist? 3—Why are the first two p-electrons (in boron and diamond) separated into two single electron bands in the second period, but overlapping in the third period (aluminum)? 4—Why are s-electrons and d-electrons taken up in continuous overlapping bands, while p-electrons are divided into single electron bands? 5—Why do the peaks and valleys (y and w and further x and z) of the energy band below four electrons per atom not show up in the electrical conductivity of alloys? For example consider the Li-Mg system or the alloys between Mg and three electron metals where the mentioned discontinuity in the lattice parameter is found. 6—Why does the beginning of the p-electron band (x) not show up in the lattice constants similar to the filling up of the s-electron band (z) ? In magnesium alloys the electron-gas theory postulates the first Brillouin zone to be filled at about two electrons per atom. This is claimed to explain the sudden change in lattice spacing and c/a values of several magnesium alloys when the electron concentration exceeds a few percentage points over two electrans per atom. This was emphasized in the paper by Busk. If the electron-gas energy-band theory is correct a sudden change in electrical conductivity and possibly other properties .should be expected when the same electron-concentration or temperature is exceeded. A sudden change in lattice spacing or other properties should also be expected when the filling degree is such that p-electrons are introduced into the p-band, for example at x in Figs. 6 and 7. Such phenomena are at found by experiment. and If the number of electrons should vary with the energy level depending on the average number of bonding electrons per atom, the electrical conductivity should be expected to vary in accordance with the energy band layout (Figs. 6 and 7) caused by different numbers of conducting electrons at different filling up degrees. Nothing indicating such a behavior is observed. In addition to these discrepancies between the electron-gas and energy-band theory and measured data, the theory violates the principles developed along with the Bohr theory of atomic structure. According to these principles a filled shell is saturated and therefore unable to form bonds. Therefore two S-electrons per atom should form a closed or saturated shell, which has been pointed out as accounting for the inability of helium to form bonds. Beryllium, magnesium, or calcium atoms with two s-electrons should be expected to form inert atoms with properties almost like the helium atoms. Several other inconsistencies and disagreements with measured data of the energy-band theory can be mentioned. Some of these are discussed with reference to other papers. 8 Because the electron-gas and energy-band theory seems to fail on several points, I have developed another theory which can account for all the phenomena the electron-gas theory is able to account for. This new theory is further able to account for things which are impossible to explain by the electron-gas theory at the present state.

Jan 1, 1953

By K. Tanaka, J. D. Mote, J. E. Dorn

It) an attempt to ulLcoce.lP the operative strain-rate-contl-olliy: dislocation nieclzanistns, specially oviented sizgle clystals of the intel-nzediate 1zexagonal phase containing Ag plus 33 at. pct A1 were tested in tension over a wide range of temperatures. Slip was observed to take place by the {0001} <1120> {l100} mechani fracture took place across the(i100) plane and winning occurred by the (i01Z) ?lechanisn. Basal slip exhibited a strong yield point over the -alzge from 77 to 450°K, the upper ,esolved shear st]-ess having the exceptionally high value of 10,500 psi over this entire ?-a?zge of tenzpei,atuves. The critical 9-esolved shear stress for prismatic slip decreased f7-om 48,000 psi at 4.3"K to 23,000 psi at 170°K (Region 1) follozcirg zt:lzich it decl>eased sloz&ly to 21,500 psi at 475°K (Res&apos;on II); from 475" to 575°K (Regioz III), the c7-itical esolced shear stress dec&apos;-eased precipitously to 2000 psi; and from 575" to 750°K (Region IV) it decreased less afi&apos;dly to a low value of about 500 psi. Pvistintic slip in Region I was pobably controlled by the tliel-nally activated riecharzisui of nucleation and g,-ozcth of kinks in dislocations lying in Peierls potential troughs. In Region II for prismatic slip the critical 1-esolved shear stress was slzocn to be deteemined by sh0l.t-range 01-dering, Overall the forgiorz fo basal slip, 7.c.lre1-e a Strong yield-point phenorlienu ia7as observed, the critical vesolved slzea?-stress was shoztn to be determined by n conibirzation 0-f Szizuki locking and short-range-order Izavderzizg, The precipitous decrease in the critical resolved shear stress with increase in ter,/pe7-atrir-e over Region HI was tentatively ascribed to a decrease in the degree of slort-)ange 07-del;iqq (0)- clusteing) and also the effect of fluctuations the degree of o?der, It is at pgreser2t zrtzce)taitz as to 1t1hethe1- these or other possi1)le effects are also ,esponsible. fo- the data obsel-ved 172 Region IV. 1NTEREST in inter metallic compounds stems not only from their role in dispersion hardening of polyphase alloy ystems but equally from their potentialities for high strength, hardness, and stability not only at atmospheric temperatures but especially at elevated temperatures. As summarized in a re- cent symposium of the Electrochemical Society on "Mechanical Properties of Inter metallic Compound", most of the experimental evidence regarding the mechanical behavior of intermetallic compounds centers about the effect of temperature on the hardness and ductility of polycrystalline specimens. The available data reveal that the plastic behavior of intermetallic compounds might be rationalized in terms of the usual dislocation mechanisms appropriate to a solid solutions providing the additional complexities arising from crystal structure, long-range ordering, short-range ordering, and defect lattices are taken into consideration. It is apparent, however, in terms of the history on a solid solutions, that a complete detailed mechanistic rationalization of dislocation processes may not be possible until the deformation processes are studied in single crystals of intermetallic compounds. The present paper contains a preliminary report on the plastic behavior of single crystals of the hexagonal Ag-A1 intermetallic phase over a wide range of temperatures. The results confirm the thesis that single crystal data provide a most effective method of identifying operative dislocation mechanisms in intermetallic compounds. EXPERIMENTAL TECHNIQUES Several factors prompted the selection of the hexagonal Ag-A1 intermetallic phase for this preliminar investigation on the plastic properties of single crystals of intermetallic compounds: 1) This phase has a wide solubility range5 which would permit future investigations on the effect of composition and axial ratios on slip mechanisms. 2) Although it undoubtedly exhibits short-range ordering (or clustering) this intermetallic phase is free from complexities arising from long-range ordering.6 3) Since the atomic radii of aluminum and silver are practically identical, the possible complications due to Cottrell locking are minimized. 4)Whereas the dislocations on the basal planes are expected to dissociate into Shockley partials and are thus susceptible to Suzuki locking, those on the prismatic planes probably remain complete. 5) The axial ratio, being 1.61, is almost ideal, suggesting that short-range ordering may be almost spherically symmetrical. The present investigation was conducted exclusively with the hexagonal Ag-A1 alloy containing 33 at. pct Al. Preliminary investigations revealed that this alloy undergoes basal slip by the (0001)

Jan 1, 1962