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By T. E. Davidson, C. G. Homan

This report deals with a study of the effects of extreme hydrostatic pressure on a polycrystalline material which exhibits a high degree of elastic anisotropy. Metallographically prepared polycrystalline bismuth samples were subjected to pressure levels up to 20,000 atm in a piston type gas-liquid system. The observed structural changes consisted of deformation localized along the grain boundaries, which appears to be a boundary migration phenomena, and widespread slip and cross slip. Subsequent deformation is progressive with increased pressure, and is substantially different in type from the deformation characteristic of uniaxial compression. A large amount of interest has been shown in recent years in the effects of extreme pressure on metals. In this connection, many investigations into the effects of pressure on physical and mechanical properties and reactions have been undertaken. Due primarily to the experimental difficulties involved, there has not been a great deal of study of the structural changes resulting from exposure of metals to extreme pressure. It has been difficult, for instance, to produce a true hydrostatic stress state in many of the extreme pressure devices used, and also to directly examine specimens metallographically before and after pressurization. The results of a preliminary investigation into one of the effects of extreme pressure, and its associated true hydrostatic stress state, on the structure of metals are discussed herein. Metals exhibit varied degrees of anisotropy in their physical and mechanical properties. Significant among the anisotropic properties is the modulus of elasticity, which varies by a factor of as much as three in some of the fcc metals, and by much larger factors in some of the less symmetrical elements. As a result of the variation of the modulus of elasticity with crystallographic direction, shear stresses will be induced in polycrystalline samples under true hydrostatic compression. Depending on several factors, these shear stresses could become of suffi- cient magnitude to cause localized plastic deformation. Presented herein are some of the structural changes observed in polycrystalline bismuth as a result of exposure to hydrostatic pressures of up to 20,000 atm. Although several other materials in both bicrystal and polycrystalline form are also currently under investigation, bismuth was chosen initially, due to its low symmetry and moderately high degree of anisotropy in the elastic properties. The fact that extreme hydrostatic pressure can induce structural changes in some elements in the polycrystalline state is of importance in itself. However, it may also be a factor in many of the other extreme pressure effects. For instance, this localized plastic deformation could substantially lower the recrystallization temperature, thus altering the effects of extreme pressure on this phenomenon. Also, as commented on by ridgman,' the hysteresis in the polymorphic transition of polycrystalline bismuth under extreme pressures could possibly be attributed to structural changes brought about by unequal compressibility. Therefore, to accurately determine the true effects of extreme pressure on metallic substances, one first must determine if any structural changes due to hydrostatic compression are also simultaneously occurring. The purpose of this work has been to determine whether the shear stresses induced by hydrostatic pressure in an anisotropic polycrystalline material, such as bismuth, could reach a magnitude great enough to cause plastic deformation and to study the resultant structural changes and effects. THEORY In a metallic crystal, the elastic properties are dependent upon the crystallographic direction and plane. Depending upon the degree of anisotropy of the elastic properties and other factors, internal shear stresses of sufficient magnitude to cause localized plastic deformation may be induced in polycrystalline material by true hydrostatic compression. The following highly simplified two-dimensional model of two adjacent grains in a polycrystalline aggregate under hydrostatic pressure serves to schematically demonstrate how these shear stresses arise in an anisotropic material.

Jan 1, 1963

By H. H. Podgurski, Hsun Hu

Large cryslals of high-purily iron (99.996+ pcl) cannot be obtained by the usual strain-ameal technique. Repealed phase transformation by thermal cycling prior to crilical deformation improves the capabilily of cryslal growith but fails to yield satisfactory crystals. Imperfect columnar crystals can be produced by the phase-transition method. The addition of a small amounl (0.005 wl pet) of nilrogen IT was reported'-4 that for metals of very high purity the usual procedures of the strain-anneal technique fail to produce large crystals. This was eliminates practically all the difficulties in crystal growth by the normal strain-anneal procedures. An actual example is presented, showing that most parts of a strip have been converted into a large crystal by this method. The addition of nitrogen, which can be removed from iron with less difficulty, is believed to be more desirable than carbon additions for the same purpose. thought to be due to the strong tendency for recovery in the high-purity metals: hence the critically strained matrix undergoes polygonization instead of recrystallization upon annealing for crystal growth. Talbot1,2 succeeded in overcoming this difficulty by adding 0.03 to 0.04 pct C to high-purity iron before critical deformation. followed by annealing in hydro-

Jan 1, 1965

By E. Gregory, A. Coucoulas

Carbm dioxide, which exists metastably as a constantly subliming molecular solid in a normal room-temperature environment, has been shown to exhibit many microstructural features which are similar to those observed in metals and ceramics at temperatures approaching their melting points. An examinution has been made of some of the factors affecting a brittleness condition known as "smzdiness" which occurs in solid carbon dioxide. It has been shown to be highly dependent on manufacturing and storage conditions and is thought to be a marked intergranu-lay brittleness which results from a combination of excessive grain growth and a concentration of gas-filled pores at the grain boundaries. In the course of an investigation carried out to explain certain physical and mechanical properties of solid carbon dioxide, observations were made of the microstructure and fragmentation of samples after fabrication and storage under varying conditions. Many similarities between the structures of solid carbon dioxide observed under ambient conditions and those existing in metal and ceramic systems at temperatures approaching their melting points have been observed. There are, however, important differences between solid carbon dioxide and the common metallic and ionic solids which result primarily from the differences in the nature of the bonding forces. Despite these differences, the relative ease with which certain microscopic observations can be carried out, may make solid carbon dioxide a useful material model for the physical metallurgist and the ceramist. It could also prove a valuable material for educational purposes where it is required to illustrate grain phenomena which normally can only be seen with the aid of a high-temperature microscope, and after prolonged chemical or thermal etching.

Jan 1, 1963

By G. T. Horne, P. J. Wray, S. Kinoshita

CONSIDERABLE evidence exists for a strengthening mechanism operative at intermediate temperatures in bee metals containing interstitial solute atoms. The strengthening is indicated by a maximum in the ultimate tensile strength1 as well as the fatigue

Jan 1, 1965

By H. Sigurdson, T. V. Cherian, C. H. Moore

The phenomenon of recovery of cold-worked metals is interesting not only because of its practical importance but also because of its fundamental significance in solid state reactions. Although extensive investigations1,2 have already been made in an attempt to discover the mechanics of the recovery process, many of the observations have not yet been satisfactorily correlated to provide a completely consistent model for the process. The wide differences in the recovery rates of various properties can be cited as a typical example of one of the difficulties that are encountered. Frequently, for example, the electrical resistance will have almost completely recovered before any recovery in tensile strength can be detected. Of course, such differences in the recovery rates of different properties might be explained by assuming that each property is a unique function of the work-hardened state, and consequently each property exhibits its own unique recovery rate. The assumption that different properties are uniquely related to the work-hardened state cannot be denied. On the other hand, the properties that recover at different rates often exhibit more or less parallel changes upon work-hardening. This suggests that the microstructural changes attending recovery are not exactly the reverse of the changes attending work-hardening. Several types of imperfections must be postulated in order to account for this apparent anomaly. The different recovery rates for various properties, then, are due to the different recovery rales of the type of imperfection to which each property is most sensitive as well as the unique dependence of each property on the cold- worked state. This concept assumes that a simple model of the work-hardened state consisting only of one type of imperfection, such as Taylor's type of dislocation patterns, is inadequate to cope with the diversified phenomena attending work-hardening and recovery. Although current models for the work-hardened state are not useful for describing all aspects of the recovery process, the general trends of the recovery of each postulated type of imperfection as a function of time and temperature should be at least qualitatively deducible from the rather well developed laws of kinetics of reactions in the solid state. Consequently, recovery data might prove useful for elucidating some aspects of the complexities of the work-hardened state of metals. A preliminary attempt to study work-hardening by investigating recovery rates of cold-worked metals is outlined in the following pages of this report. Experimental Procedure Many properties recover when cold-worked metals are annealed below their recrystallization temperature. Therefore, electrical resistivity, thermal electromotive force, X ray diffraction line widths, X ray diffraction line intensities, elastic spring back, density and other physical and chemical properties have been used to study the recovery process. Major interest, however, has generally been directed toward the recovery of the mechanical properties such as hardness, yield strength, and tensile strength. But a search of the literature suggests that the effect of recovery on the true stress-true strain curve has been neglected, in spite of the current recognition of the fundamental importance of such an investigation. An investigation on the effect of recovery treatment on the true stress-true strain curves in tension, therefore, was undertaken in the present study. Commercially pure aluminum (99. + pet Al) in the form of 0.100 in. thick rolled sheet of 2S-O aluminum alloy was selected as the material for this investigation because rather extensive correlatable data are already available on the recovery of some of its properties. Tensile specimens having a 6 in. long gauge section and a uniform reduced section width of 0.500 in. were machined from the sheet in accordance with a design that has previously been reported.3 All specimens were selected with their axes aligned in the rolling direction. In order to eliminate the effects of previous work-hardening and the effects of machining, the specimens were annealed for 15 min. at 750°F before testing. During tensile testing the loads were measured by means of a proving ring (sensitive to 1/2 lb) in series with the specimen.4 Strains were determined from the extension of a rack and pinion strain gauge sensitive to a strain of + 0.0001. The stress was recorded as the true stress, namely

Jan 1, 1950

By Nicholas J. Grant, H. C. Chang

TRANSCRYSTALLINE fracture surfaces of the cleavage type have been examined by microscopy and X-rays for several metals.' These investigations revealed that the fractured surfaces were not flat and invariably had cleavage and other deformation markings on them. On the other hand, the study of inter crystalline fracture surfaces has been limited. Perryman, in a recently published paper on the inter crystalline fracture of ß-brasses,' found that a thin layer of heavily deformed metal was invariably associated with these surfaces. Furthermore, the thickness of this layer was found to depend on the orientation of the grain boundary surface with respect to the applied stress. The mechanisms of grain boundary sliding and intercrystalline fracture have been presented in two previous papers,which described certain aspects of the deformation of a 20 pct Zn-A1 alloy. It is the purpose of this paper to present some observations on the structure of intercrystalline fracture surfaces and to correlate the structure of such surfaces with the mechanism of intercrystalline fracture. Experimental Techniques—The high purity 20 pct Zn-Al alloy was generously supplied by the Research Laboratories of the Aluminum Co. of America. Two sizes of specimen were used; the first had two parallel flat surfaces of dimensions 1x0.17x0.09 in. and the second was round with a 0.25 in. diam. A grain diameter of 1 to 3 mm was obtained by annealing the specimens at 1000°F for 24 hr. The specimens were then air cooled. The specimens were subjected to creep at 500°F and 2300 psi. The rupture times ranged from % to 3 hr, with one room temperature tensile test. It has been mentioned previous1y" that the intercrystalline fracture surfaces of this alloy, when tested under the conditions given above, are sharply defined. In the optical examination of the fracture surfaces, whenever possible, magnifications of at least X250 were used. The locations of the regions examined, both optically and by X-rays, were selected to give several fracture surface orientations with respect to the stress axis and to the specimen surfaces. This was possible for the X-ray studies since the grain diameters were large compared to the X-ray beam. A film to specimen distance of 3 cm, a 1 mm pinhole and Cr radiation were used for the Laue back reflection pictures.

Jan 1, 1957

By G. R. Wilms

A study has been made of the structural changes in polycrystalline high purity aluminum during the tertiary stage of creep under conditions of constant tensile load. It appears that there is no basic modification in the particular mechanism of deformation that characterized the preceding secondary stage of creep, but that the changes in crystalline structure produced during the tertiary stage are only a consequence and not a cause of the accelerated strain. Intercrystalline fissures were observed on entry into the tertiary stage, and these might supplement the accelerated rate of strain owing to local stress concentrations. There is evidence that the formation of these fissures may be associated with the presence of surface imperfections, and it is suggested that a more generalized treatment of the results may explain certain creep characteristics of commercial alloys. IN previous work'-4 a study was made of the changes produced when specimens of annealed aluminum were deformed at various temperatures and rates of strain, especially under conditions leading to creep under constant tensile load. It was shown that, as the temperature was increased or as the rate of strain was decreased, elements of the substructure produced within the grains became coarser, and with the strain rate constant at a given temperature, their size tended to an equilibrium value. The preliminary process of the breaking down of the grains into a substructure was identified with primary creep.' The subsequent equilibrium, condition was associated with secondary creep. The question then arose whether the tertiary stage corresponded to any further significant change in the substructure. The question is also of interest in view of the opinion, expressed for instance by Andrade,5 that the tertiary stage is merely an accelerated rate of strain due to the reduced area of section of a specimen under continuous tensile load, and the other opinion, as suggested by the observation by Sully, Cale, and Willoughby6 of a tertiary stage in com-pressive loading, that this third stage is a real effect in its own right. The object of this paper is to record some observations on the tertiary stage of high purity aluminum. Experimental Procedure The material was aluminum of 99.98 pct purity as previously used. The specimens had a rectangular section of l/8 x l/2 in., and extensions were measured over a 1 in. gage length. After machining to shape, the specimens were annealed to give a standard grain size of 0.2 mm and then electropolished. The general procedure involved metallographic and X-ray diffraction examination at progressive stages of elongation (marked by arrows on the curves shown in Figs. 1 and 2), the specimen being extended to fracture under constant load at a given temperature. The X-ray tests were made by the back-reflection method from a stationary specimen, using CoK radiation. The longer edge of the micrographs in the paper corresponds to the direction of applied stress. Results The previous work showed that, in the secondary stage of creep in aluminum, three basic types of structure could be developed according to the temperature and rate of strain. These were: 1—slip, which predominated at the lower temperatures and higher rates of strain and produced the finest and least perfect substructure; 2—cell structure, which superseded slip as the temperature was raised and the rate of strain decreased and which was characterized by a coarser and more perfect structure and an absence of slip lines; and 3—a stage termed "boundary microflow," which predominated at still higher temperatures and lower rates of strain when the cell size had increased and become comparable

Jan 1, 1955

By E. H. Edwards, J. Washburn, E. R. Parker

The mechanism of strain hardening was discussed in connection with some recent observations on the stress-induced motion of dislocation boundaries and on the simple shear deformation of zinc, cadmium, and copper crystals. THE nature of work hardening accompanying plastic deformation of metals has long been the subject of speculation. As early as 1925, Becker' suggested a thermodynamic process to explain slip based on the statistical probability for the occurrence of local glide steps. The theory that plastic glide was due to the presence of lattice defects known as dislocations was first introduced in 1934 by Taylor: and Orowan, and Polanyi.4 Taylor presented a rather complete analysis of the behavior and interaction of dislocations. He assumed a certain type of stress field to be associated with a single dislocation, and postulated that work hardening would result from plastic flow in two ways. First, dislocations following one another across a slip plane might encounter a barrier which would impede their movement, thereby gradually diminishing and eventually stopping plastic flow on that plane. The second mechanism involved the concept of dislocations traveling in opposite directions on nearby planes attracting one another to form a meta-stable lattice. The interaction of dislocations in such an array was shown to be large enough to require an increase in stress before plastic flow would continue. Kochendorfer5,6 extended the concept of work hardening involving a back stress due to an accumulation of like dislocations on a slip plane. He suggested that the formation of a new dislocation is hindered by those dislocations already present. The hardening resulting from the interaction of newly forming dislocations with bound dislocations was designated as "formative hardening," and indicated to be the only type of hardening necessarily connected with the slip process. In another detailed picture of work hardening, Mott and Nabarro7 assumed that sufficient dislocations are primarily present in the crystal as a consequence of growth irregularities (mosaic structure) to initiate slip. Work hardening resulted from the interaction of migrating dislocations with localized internal stresses produced during deformation. More recently, however, Mott8 as modified his treatment to include the generation of dislocations at Frank-Read sources and he accounts for the hardening by the formation of sessile dislocations. It has been widely accepted that metals of hexagonal structure, usually possessing only one set of glide planes, exhibit mechanical behavior markedly different from those having cubic structures. However, the differences in behavior of the two classes may be less fundamental than suggested. Rohm and Kochendorfer9 have shown that a single crystal of aluminum subjected to approximately simple shear gives the type of stress-strain curve associated with the hexagonal metals. Similar behavior has been demonstrated for gold and silver crystals by Andrade and Henderson.'" Thus when glide in cubic metals is restricted to one plane, the strain hardening characteristics closely resemble those of the hexagonal metals. It would appear, therefore, that an investigation of some of the strain hardening properties of single crystals sheared in simple glide might provide a more complete understanding of the phenomenon of work hardening. This report relates a number of recent observations made on single crystals of copper, zinc, and cadmium tested in simple shear. In addition, pertinent observations of effects accompanying stress-induced movement of dislocation boundaries are reported. Experimental Procedure and Results Details of the production and advantages of the type of single-crystal shear specimen employed for the following tests have been presented in a previous publication.n The method of testing makes

Jan 1, 1954

By Martin Epner, Eric Gregory

DURING the past decade, considerable work has been carried out on various cermet systems in an effort to produce materials suitable for high-temperature applications in gas turbines. Most of the materials have not, as yet, met the requirements of the engine manufacturers, primarily because of their lack of room-temperature ductility and thermal shock resistance. In the field of tool materials, however, although ductility is still of prime importance, much has been accomplished by improving design parameters and the quality of carbide materials with a view to using cermet compositions successfully. The systems described in this paper were designed primarily for application in this field, although their development was a natural outcome of work on similar compositions designed for high-temperature service. In steel-bonded titanium carbide, the steel binder imparts heat treatability to the system while the carbide provides abrasion and wear resistance. Combining the carbide and the steel matrix in carefully controlled proportions results in a composite material which, when the steel is hardened, has man of the properties of cemented tungsten carbide. In the annealed condition, this material can be readily machined by the conventional chip-removal techniques. Titanium carbide has certain distinct advantages over many of the other carbides when combined with a steel matrix. It is among the hardest of the carbides and is one of the most stable. Liquid iron wets titanium carbide, and this wetting assists in the formation of a dense cermet. Once the composite structure has been formed, the titanium carbide acts as a relatively inert phase during the heat treatment of the matrix. PREPARATION OF ALLOYS The steel-bonded titanium-carbide materials described in this paper were prepared by the liquid-phase sintering technique. The titanium carbide used had an average particle size of 6 to 7 p, a total carbon content of 19.4 pct and 0.35 pct free carbon. This was blended with the elemental metal powders by ball milling them together in a hexane medium. After drying, paraffin wax was added to the blend to provide green strength in the pressed briquettes, and compacting was performed at 15 tsi. Liquid-phase sintering was carried out in vacuum at temperatures which depended on the chemical composition of the system involved, but which were usually around 2640° F. The resultant materials had a density approaching 100 pct of theoretical. The

Jan 1, 1961

By Webster Hodge

MALL generators and motors are required to Soperate, in some critical applications, at temperatures where cold-worked silver-bearing copper re-crystallizes. Copper containing up to 30 oz Ag per ton has been used for many years to build the highest grades of commutators and slip rings. With failure of commutators impending because silver-bearing copper cannot withstand the stresses imposed at these higher temperatures, Cr-Cu was tried as a substitute with varying degrees of success. Cr-Cu proved difficult to fabricate because a critical heat treatment was necessary to obtain sufficient ductility. Besides, its rupture strength was shown to be low at operating temperatures, its electrical conductivity lower than desired, and its cost high. Furthermore, it often contained hidden flaws. The Nippert Electric Products Co. set out to find a better material by undertaking a research project with the stated objective of providing a material suitable for commutators that would have 1) a softening temperature above 750°F (400°C), 2) a minimum electrical conductivity of 90 pct of that of pure copper, 3) a strength not less that that of cold-worked silver-bearing copper, and 4) a materials cost of not over 1.5 times the cost of copper containing 25 to 30 oz Ag per ton. No commercial alloy met these specifications. However, it was found that copper containing from 0.18 to 0.35 pct Zr could be produced quite easily in the laboratory, and promised to fulfill the major requirements noted above. Interest naturally centered in the lower part of this composition range, since such alloys gave somewhat better electrical conductivities, and were lower in

Jan 1, 1958

By C. Y. Ang, C. Wert

Quench aging of supersaturated solid solutions of nitrogen in columbium takes place in reasonable times in the temperature range 300&apos; to 500°C. Changes in internal friction, hardness, and electrical resistivity are found to accompany this change. THERE is practically no detailed information in the literature concerning interstitial solid solutions in either columbium or tantalum. The fact that both of these metals are used extensively commercially in both pure and alloyed states makes studies of their properties of value. While the plan is to study both these metals and their alloys extensively, so far only a small number of measurements have been made. This paper presents data obtained in the initial investigations of columbium. The principal interstitial impurities of columbium are likely to be hydrogen, carbon, oxygen, and nitrogen. The first of these, hydrogen, is outside the scope of internal friction measurements at present and nothing quantitative is known about the initial hydrogen content of the specimens. However, it is believed that it is completely absent after the preliminary heat treatments are carried out. The carbon content of the original samples is low and can be reduced still further during sample preparation. Oxygen is normally present to some extent in columbium and can also be added or removed easily. Nitrogen is also normally present in small amounts and can also be controlled within limits. Since it is not possible yet to remove all the nitrogen from columbium and since it is possible to remove essentially all the oxygen and carbon, initial measurements were made on the alloy system nitrogen in columbium. In this sense then, the alloys are true binary alloys. Little is known about the effects of nitrogen on the physical and metallurgical properties of columbium, so it was necessary to perform some rather elementary measurements on the specimens. These measurements are described briefly in the next section and are discussed in detail in later sections. Scope of the Work That internal friction is associated with diffusion of nitrogen in columbium has been known for some time. It is similar in nature to that for carbon and nitrogen in a-iron. In these latter cases, the information obtained from internal friction has been useful metallurgically.&apos; and it is not too much to expect useful information from similar studies on columbium containing nitrogen. The technique of internal friction measurements is adequately described by Ke2 At a frequency of 1 cps the internal friction peak for nitrogen in columbium occurs at 300°C and that for oxygen at 160°C." So that all the measurements will be meaningful metallurgically, the internal friction has been calibrated by making a measurement of the damping at the two peaks and a measurement of oxygen and nitrogen content by chemical analysis. A comparison of the two gives the following calibration factor: To go from damping capacity expressed in terms of 1/Q to weight percent of oxygen and nitrogen in solid solution in columbium, multiply the peak values of internal friction by about 4 and 3 respectively." Data from which these factors were calculated are presented in Table I. The increase of electrical resistance of columbium caused by the addition of oxygen and nitrogen as interstitial impurities has been quantitatively determined. At room temperature the percentage increase is small for small additions of impurities (<0.1 wt pct); as the temperature is lowered it becomes an increasingly larger factor of the total resistance. An empirical relation has been established to express the resistivity at — 195°C in terms of weight percent of oxygen and nitrogen in solid solution. It is easy to make supersaturated solid solutions of nitrogen in columbium and to observe aging phenomena. The optimum aging temperature, as in the case for carbon and nitrogen in a-iron, occurs at the temperature for which the time necessary for a single place-change in diffusion is about 1 to 10 millisec. This means for nitrogen in columbium that significant property changes should take place for this alloy at about 400" to 500°C in reasonable time intervals. Results of such aging experiments are

Jan 1, 1954

By J. A. Hren, C. M. Wayman

RECENT investigators have noted anomalous behavior in the electrical resistivity of vanadium below room temperature. Rostoker and Yamamoto1 in studying the vanadium-oxygen system hypothesized that the body-centered-tetragonal phase existing from 15 to 34 at. pet O at approximately 800°K was really a low-temperature allotrope of vanadium that was stabilized by oxygen. In measuring the electrical resistance of cold-swaged vanadium rod over a range of temperatures, a small discontinuity was observed at approximately 240°K, supporting their notion of allotropy. However, the magnitude of the observed discontinuity was extremely small. Loomis and Carlson2 in studying the ductile-brittle transition in vanadium also made resistance measurements and reported a discontinuity near 200°K. More recently, White and Woods3 showed that the electrical resistivity from room temperature to approximately 100 °K was not linear and at variance with the Bloch-Griineisen law. Instead of finding a discontinuity, they observed a deviation from linearity below 200°K. In discussing their results they pointed

Jan 1, 1961

By J. T. Michalak, H. W. Paxton

The recovery of the initial flow stress of poly-crystalline iron is characterized by a) a logarithmic time dependence; b) an increasing activation energy with increasing recovery; c) an increased ?,ate and fraction of recovery with decreased temperature of deformation; and d) an increased rate with increased deformation. Isochro)ml studies of the recovery of single crystals revealed more sluggish recovery. The original flou3 stress of single crystals could be regained by recovery anneals at 800"or 900°C. ThE great majority of the theoretical treatments and experimental studies of strain-hardening have been concerned with the close-packed lattices.&apos; Very little effort has been expended on the bcc structure so that the theory of work-hardening of this lattice is essentially nonexistent. The limited amount of experimental data resulting from the direct studies on silicon-iron,2 low-carbon steel,3 and purified irons4"7 does not as yet permit a detailed analysis of the strain-hardening. Studies concerned with the effects of annealing following deformation have been few in number," have been somewhat limited in extent, and have had the disadvantage of using irons containing a substantial impurity content. The present research was undertaken to augment the limited knowledge we have of the characteristics of a bcc metal as regards strain-hardening and recovery from the strain-hardened state. The recovery of the flow stress has been employed since it was felt that the stress associated with the hardened state gives at least a fair qualitative description of the state of the imperfections causing the hardening. The study was concerned primarily with the recovery of flow stress following strain-hardening of polycrys-talline, zone-melted iron as a function of the amount of deformation, the temperature of deformation, and the time and temperature of the recovery anneal. A limited investigation of the recovery of the initial flow stress of single crystals of this same material as a function of the recovery temperature was also conducted in conjunction with the Lambot X-ray technique." MATERIAL The zone-melted iron used in this investigation was obtained through the courtesy of the American Iron and Steel Institute and Battelle Memorial Institute. The analysis, as given by the supplier indicated a maximum impurity content of 0.02 wt pct excluding those impurities listed in Table I. (The actual impurity content is very probably less than this since detection limits were considered as the amount of a particular impurity level.) Also listed in Table I are the interstitial impurity contents of the iron at various stages of experimental work indicating that these im-

Jan 1, 1962

By G. F. Bolling

STUDIES of grain boundary migration in zone-refined metals have all shown that the rate of migration is greatly reduced by small added solute concentrations. However, it is apparent that a difference exists between boundary migration during normal grain growth and single boundaries migrating in a bicrystal to consume a substructure. To effect the same reduction in velocity in the two cases, much more solute is required for grain growth than for the single boundary experiments. One case is available for direct comparison; both Bolling and winegardl and Aust and utter&apos; added silver and gold to zone-refined lead to study grain growth and single boundary migration, respectively. For comparable reductions in migration rates, about 500 times more solute was required to retard grain growth than to retard the single boundaries. A reason for this difference is suggested here. The rate of grain boundary migration is dependent on solute concentration and must therefore also depend on the solute distribution; i.e., regions of higher solute concentration encountered by a moving boundary must produce greater retardation and thus could determine any observed rate. A dislocation substructure can be the source of a nonuniform solute distribution since it can attract an excess concentration of certain solutes. In fact, it is probable that the solutes which impede grain boundary migration most would segregate most severely to a substructure for the same reasons. Thus a dislocation substructure present in a crystal being consumed could locally magnify the concentration of solute confronting an advancing grain boundary. In the single boundary experiments a low-angle substructure, within single crystals obtained by growth from the melt, was used to provide the driving force to move a grain boundary; in grain growth, no substructure of this magnitude was present. The increased solute concentration at subboundaries should be given approximately by C, = G e c,/kT, where t, is a binding energy and CO the bulk concentration. To account for the difference between the two experiments in the Pb-Ag and Pb-Au cases, C, must be the concentration impeding the single boundary migration, and a value of t, = 0.25 ev is necessary. This is reasonable, even though calculation on a purely elastic basis gives t, = 0.12 ev. because electronic effects must enter for silver and gold in lead. The compound AuPbz forms3 and the metastable compound AgrPb has been reported to nucleate at dislocations prior to the formation of the stable, silver-rich phase.4 Other observations support the hypothesis that a magnified solute concentration impedes the single boundary migration. For example, some crystals were grown by Aust and Rutter at concentrations of ~ 0.1 wt pct Sn and 2 x X at. pct Ag or Au which exhibited a cellular substructure, and in these crystals no boundary migration was observed. It is therefore evident that the higher concentrations at cell boundaries drastically inhibited migration. Inclusions would not have been responsible for this inhibition since according to recent work on cellular segregation,5 no second phase should have occurred in the segregated regions at the cell boundaries for the conditions of growth used, at least in the Pb-Sn system. In the purest lead, only the "special" boundaries observed by Aust and Rutter gave rise to the same activation energy as that obtained in grain growth. It is reasonable to suppose that the structure of special boundaries does not favor segregation at low concentrations and thus solute, or an inhomogeneity in its distribution, would have no effect. Random boundaries, on the other hand, are affected by solute and the substructure would enhance residual concentrations in the zone-refined lead, leading to a higher activation energy. It is clear, even without a detailed theory, that the apparent activation energies and exact solute dependence in the two experiments must be different as long as the non-uniform solute distribution produced by the substructure is important. Recrystallization experiments should also be susceptible to the same kind of local segregation at subboundaries or disloca tion cell walls; a suggestion similar to this has been made by Leslie et al.&apos; Following the arguments presented here, the effects of a given solute concentration would be like those observed by Aust and Rutter if segregation occurred, and like those of grain growth otherwise. This work was partially supported by the Air Force Office of Scientific Research; Contract AF-49(638)-1029.

Jan 1, 1962

By S. Storchheim

MORE and more attention is being paid to the bonding of metals in their solid states. For a better understanding of this technique for joining metals and how it is affected by changes in temperature, pressure, and time at temperature and pressure, a detailed report concerning nickel to aluminum bonding has been published.&apos; In order to broaden the knowledge accrued, some additional work concerning solid state joining of aluminum to copper and aluminum to zirconium was performed. The investigation of the Al-Cu system was considerably more extensive than the investigation of the Al-Zr system. For the A1-Cu system, not only were tensile sudies made but intermetallic penetration rate investigations also were carried out. The effect of temperature on intermetallic penetration rate for the A1-Cu system was determined at 11 tsi pressure, held 2 min. Procedure Apparatus: The hot pressing technique was the means of solid state reaction used and required the equipment depicted in Fig. 1. The following procedure was involved: The two metals to be reacted were placed in an aquadag-lubricated 18-4-1 tool steel die, 16 in. high by 1.440 in. ID, between punches of 1.366 in. diam made of the same material. A thermocouple well was located in the die body 3½ in. down from the top of the die, while another well was located centrally in the bottom punch 8½ in. from the bottom of the die. This die assembly was located in three cylindrical ceramic heating furnaces placed in tandem. Each furnace was controlled individually by a Variac power transformer. In turn, the die and furnaces were placed in a water-cooled stainless steel pot which could be evacuated. A cover, which contained a centrally located Wilson seal with an 18-4-1 1 in. diam ram running through it, was bolted on the pot. After sealing, the pot was evacuated by a roughing pump to 200 microns pressure, after which a diffusion pump was used to bring the pressure down to 5 to 15 u. At this pressure, the furnaces were turned on. As soon as they started to heat, out-gassing of the entire unit raised the pressure to 30 to 400 p. By the time the specimens were at temperature ready to be pressed, approximately 4/2 hr, the vacuum pumps had re-established the 5 to 15 u pressure. Once the desired temperature was reached, the required pressure was applied for a predetermined length of time to the 1 in. ram, through to the top punch, and to the specimen. When the time for keeping the specimen under pressure had elapsed, the pressure was released, the energizing coil current turned off, and the assembly allowed to cool. After cooling, the die was removed from the pot and the specimen was ejected. Specimen Preparation: Two different types of specimens were made for this investigation. One was for subsequent tensile testing, while the other was for determination of intermetallic alloy zone penetration into the parent metals. Tensile Bars—Commercially rolled copper pieces in. thick or zirconium sheet pieces 1/32 in. thick and 1.366 in. diam were placed between commercial 2-S aluminum rod 1 in. thick and 1.366 in. diam. This sandwich in turn was slipped into a 2-S aluminum sleeve 1.438 in. OD and 1.370 in. ID. This sleeve lined the couple up and prevented the aquadag lubricant from getting in between either the A1-Cu or Al-Zr interfaces. Immediately prior to the specimen assembly, the copper or zirconium was abraded on the flat surfaces with 320 grit silicon carbide paper, producing clean smooth surfaces. The aluminum was chemically cleaned just before assembly by: l—degrease in acetone, 2—distilled water rinse, 3—immersion for 3 min in 5 pct NaOH at 70" to 80°C, 4—distilled water rinse, 5—immersion for 2 min in 50 pct HNO3 solution at room temperature, 6—distilled water rinse, and 7—drying in a blast of gas. After the A1-Cu sandwiches were hot pressed and ejected, the specimens were machined such that the aluminum sleeve was removed, and the remaining aluminum was then threaded; the bar so produced was tested later for tensile strength. In all the instances where Al-Cu couples were tested, the specimens broke during the test at the Cu-A1 interface and never within the aluminum or copper. The ultimate tensile strength values at times showed considerable scatter for a set of given reaction conditions. Because of this, as many as three to five specimens were made for a particular set of conditions. The trend of the average tensile strengths obtained was not as conclusive as was the trend of the maximum tensile strengths, the latter values being obtained under optimum reaction conditions. Therefore, the values of ultimate tensile strength given herein are maximums.

Jan 1, 1956

By Volker Ruth

DURING an experimental investigation of the Kirken dall effect in the Ag-Au system, cracks along grain boundaries in connection with a shift of adjoining crys tallites are observed on the gold-rich side of the couple. Also on the gold-rich side the formation of a zone of small protrusions is found.&apos; The growth law of the width of this zone indicates that the protrusion formation is a secondary effect of diffusion. The protrusions are not found within a distance of approximately 60 from a grain boundary. Unaffected regions along grain boundaries were also observed by Doherty and Davis during precipitation of vacancies in aluminum2 and by Wyon and Lacombe during precipitation of iron in aluminum.3 Thus we may suppose that we have an effect analogous to precipitation in which the crystal tends to maintain its equilibrium vacancy concentration by transporting excess atoms to the surface. At an early stage (i.e., before sintering has influenced the shape of the protrusions) the protrusions have similar shapes (three sided pyramids and roof shapes) and similar orientations within a single grain but vary from one grain to another. Further experiments were carried out on cylindrical Au-Ag single crystal couples. During the weld-

Jan 1, 1963

By Richard Borg

The partial molal free energy of Cd in each of the four intermediate phases, Cd3 Cu, Cd8 Cu5, Cd3Cu4, and CdCu2 is determined using the Knudsen vapor pressure technique. Measurements are made also within the two-phase regions, and thus the variation in FCd is now known across the entire phase diagram. The absolute values are all less the 2.0 kcal and increase regularly with increasing Cd concentration. The small absolute value of the free energy precludes the possibility of obtaining highly accurate values; nevertheless, the results are compatible with those previously reported for the liquid phase. Additional results are reported concerning the lattice parameters of the 6, E, and b phases. THE elements Cu and Cd combine to form four intermediate phases which are designated* and E and which include respectively the following inter metallic compounds: CdCu2, Cd3Cu4, Cd3 Cu5, and Cd3 Cu. The extent of each phase with respect to temperature and composition is shown in Fig. 1 which is the equilibrium diagram drawn according to Hansen.1 All four phases are brittle, and most probably are true intermetallic compounds. Only the crystal structure of the 6 phase has been thoroughly investigated.2,3 The structure is cubic and isotypic with y brass. The structure of Cd3 Cu has not been previously reported, It is hexagonal with a = 8.11 and c = 8.76 0.02Å and the structure is presently being determined.* The compounds Cd3 Cu4 and CdCu, have been each previously the subject of a single investigation. Cd3Cu, is presumed to be cubic4 with the space group Fd3m, whereas CdCu2 is hexagonal2 with the C14-type structure. Cu2 Cd is found by X-ray analysis to be hexagonat with lattice parameters a = 5.00A and c = 16.17Å. The material appears to be pure Cu2Cd as all reflections could be indexed on the basis of the single hexagonal structure. These values disagree with those of Kripyakevich and coworkers5 who report a = 4.96Å and c = 7.98Å. Their investi-

Jan 1, 1962

By H. A. Lequear, J. D. Lubahn

A sudden change from one constant strain rate to another during a tensile test causes an unusual transient in aluminum alloy 61ST. A sudden change from one constant stress to another during a creep test does not produce an unusual transient in this material. A specially devised test shows that the nature of the previous steady-state conditions, which the sudden change disturbs, is one of the factors affecting the transient behavior. Some of the observations may be explained in terms of strain aging. USUALLY a sudden change during a tensile test from one constant strain rate to another causes a small, abrupt vertical shift from one stress-strain curve to another, parallel one. This behavior is apparently characteristic of all metals at very low temperatures and of many metals for wide ranges of temperature. It is illustrated in Fig. 1 for some early data on 61ST tested at — 195oC, and in Fig. 2 for some recent data on copper at room temperature. The greater accuracy of the results of Fig. 2 over previous ones&apos; is due to the equipment shown in Fig. 3. Three levers apply the bulk of the load, while the changes in the small remainder load are indicated sensitively by a low capacity load scale. Refs. 2 and 3 describe the recording equipment. A test like those in Figs. 1 and 2 provides the information necessary to determine the rate sensitizrity, n,4 of the metal. ( a log s ) ? log s n = ------------------------ ----------------------------= ( a logs e, ? log e logS2-logS1 log S2/S1 log i2 — log e1 log e2/e1 where S equals stress: e is logarithmic strain or natural logarithm of (final gage length/initial gage length);" i is strain rate or de/dt: and n is rate sensitivity, rate of change of log stress with log strain rate at a given strain, or (a log S/a log i),. Normally, the rate sensitivity, n, is positive, which means that an increase in stress causes an increase in rate. However, certain exceptions, attributed to strain aging,5 have been observed. Fig. 2 indicates how rate sensitivity values are derived from the measurements. Table I summarizes the rate sensitivity values derivable from Fig. 1. In contrast to the behavior of copper in Fig. 2, the aluminum alloy 61ST, which strain ages when tested near room temperature, exhibits an unusual transient when the strain rate is suddenly changed,&apos; as shown in Fig. 4. Initially, there is a comparatively sudden vertical shift of the stress-strain curve in a direction corresponding to the rate change, but then the curve gradually shifts back again. The initial shift is not as sudden as might be expected, because the testing machine is unable to achieve the impressed rate change instantaneously, as illustrated in Fig. 2. The behavior shown in Fig. 4 will be referred to as a tensile transient, the word tensile implying that the behavior was observed in a tension type of test. The corresponding test, where the strain rate is maintained constant except for the sudden change, will be called a tensile transient test. A creep transient, by contrast, will be defined as a temporarily abnormal behavior in a creep transient test, where there is a sudden change from one constant load to another.&apos; In Fig. 4, a naive look at the gradually decreasing and then increasing load after a sudden rate incre-

Jan 1, 1957

By R. G. Aspden

The cold rolling and annealing textures were studied for 3 pct Si-Fe grains initially (001) [hkl]. A concentration of (001) [loo] primaries were observed only in the annealing textures of crystals initially having [loo] within 8 deg of the rolling direction. Eke annealing textures of the (001) [I101 cold rolling textures were sensitive to the initial orientation of the grains. Single crystal data were used to explain the formation of (001) [loo] in poly-crystalline malerial. CUBE texture formation in 3 pet Si-Fe sheet occurs during the final high temperature anneal. Its formation is dependent on the proper distribution of (001) plane primaries with reference to the rolling direction&apos; and on the growth of these primaries by secondary recrystallization. The selective driving force for these primaries with (001) within 7 deg of rolling plane is derived from the difference in surface energy between the annealing atmosphere and the crystallographic planes exposed at the surface of the sheet.&apos;-= The alignment of [loo] directions of (001) secondaries with the rolling direction is required for optimum magnetic characteristics7 and is dependent on processing.&apos; A high degree of alignment has been observed when the final cold rolled texture has a strong (111) [ll2] type component3 and the normal grain growth texture prior to secondary recrystallization has components with a [loo] parallel to the rolling direction and planes from (001) to (110) parallel to the rolling plane.1,3,8,9 Generally, this (001) component is much weaker than the (110) component. The growth rate of (001) plane primaries to secondaries has been found to be independent of the orientation of the primaries with reference to the rolling direction, i.e., the secondaries have the same degree of alignment of [loo] directions with the rolling direction as the (001) primaries.&apos; Hence, the rolling and annealing textures of individual grains or single crystals are of interest in understanding the development of the (001) [loo] secondary recrystallization texture. (001) components have been observed in the annealing textures of cold rolled grains initially having a [loo] parallel to the rolling direction and an (001) rotated up to 30 deg from the rolling plane. Crystals initially near the (001) [loo] orientation when rolled under the influence of constraints imposed by neighboring grains form deformation bands and ro- tate in both directions toward (001) <110>.10,11 Deformation bands have been reported also for crystals initially with an (001) within 1 deg of the rolling plane and a [loo] within 1 deg of the rolling direction when rolled as a free single crystal.12 These crystals have weak recrystallization textures and contain near an (001) [loo] component.10-12 When near (001) [loo] crystals are rolled as free single crystals no deformation bands form and the crystals rotate by a single rotation toward (001) [110]. Recrystallization after a 70 pct cold reduction yielded near an (001) [loo] component13 and after a 90 pct cold reduction an (001) [I201 component." Crystals initially near (210) [001] had a texture after a reduction of 70 and 84 pct which was similar to the (111) [li2] but rotated 10 to 15 deg in the transverse direction. These crystals recrystallize to (210) [OOl] or (410) [001] and contain an (001) [loo] component.10,13 An (001) component has not been detected in the normal grain growth textures of other single crystals. Crystals initially having orientations between (110) [001] and (111) [ll2] have a cold rolled texture of principally (111) [ll2] .10>16 Other crystals with a [I101 parallel to the cross direction rotate to (111) [112] and/or (001) [110] stable end orientations. Crystals initially having from (001) [110] to (111) [lie] retain this orientation after cold rolling.15 The cold rolled textures having [I101 directions parallel to the rolling direction had components in the annealing textures related to the deformation textures by 25 to 30 deg rotations about common (1 10) poles. Cold rolled textures of the (111) (112) type recrys-tallized to (120) [001] or (110) [00l]. The purpose of the present work was to further the understanding of the alignment of [loo] directions of (001) secondaries with the rolling direction. The cold rolling and annealing textures of grains initially (001) [hkl] were studied. These data were applied to the formation of (001) [loo] by secondary recrystallization. PROCEDURES AND EXPERIMENTAL TECHNIQUES The (001) poles near the sheet normal and rolling direction are given in Fig. 1 for the 10 grains studied. Each of these grains was located in the center of a polycrystalline specimen approximately 25 mm wide and 50 mm long. The lower case lettered grains, a through dl were about 1.2 mm in diam and near the (001) [loo] orientation. Specimens containing these grains were from a commercial 3 pct Si-Fe alloy with the principle orientation of (110) [001] as obtained by impurity inhibition of growth with manganese sulfide inclusions.17,18 Upper case lettered grains, A through F, were about 6 mm in diam and had (001) planes near the rolling plane and [hkl] di-

Jan 1, 1963

By J. E. Hilliard

It is proposed that shape anisotropy in the internal structure of materials be defined by distribution functions which specify what fraction of the total line length or boundary area lies in a given direction. Methods are descriibed for determinirg these distribution functions from measurements on sections through an opaque material. New derivations, that have certain advanifages over those given previously, are Provided for well-known expressions used in estinzating the total density of lineal or areal features. INCREASING attention is being given to the quantitative description of imicrostructures in metals and other materials. This interest stems from the recognition that the structure often provides a convenient half-way point between the properties on the one hand and the processing variables on the other. Such cross correlations become increasingly desirable as materials become more complex and the demands on them more stringent. One of the principal aims in structure analysis is therefore to find parameters that are readily measurable and specific enough to be correlated with other properties of the material. Certain properties, such as magnetic and mechani cal behavior, are strongly influenced by structural anisotropies; in order to relate such properties to the structure it is obviously necessary to have means of specifying and measuring the type and degree of anisotropy. For the determination of preferred orientations in crystal structure, X-ray diffraction methods are available and are widely used, but scant attention has been given to the estimation of anisotropy in the shape of grains and particles. It is the latter problem that will be considered in this paper. It is evident that shape anisotropy can be describe in terms of preferred orientations of lineal and area elements of the structure. It will be shown that these preferred orientations can, in turn, be specified by distribution functions that define what fraction of the total line length or boundary area lies in a given direction. Methods will be given for determining the distribution functions from measurements on sections through the structure. Even when the orientation distributions are of no interest per se, they still have an important bearing on the measurement of the total line length or boundary area by means of an intercept analysis. The treatment in terms of orientation distribution functions leads to new derivations, that have certain advantages of those given previously, for the well-known expressions used in estimating the total density of lineal or areal features. The treatment is divided into two main sections. The first deals with the orientation distribution of a two-dimensional array of lineal features such as occurs in the projected image of dislocations in electron transmission microscopy. The other section treats the estimation of the distribution of lineal and areal features dispersed in three dimensions.

Jan 1, 1962