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By Y. Nakada, U. F. Kocks, B. Chalrners

THE orientation dependence of work hardening has previously been studied over the entire range, i.e., including special orientations of high symmetry, in aluminum1-3 and silver.* The differences between various orientations are substantial, and the trend is the same in all fcc crystals. However, there are quantitative differences in behavior between aluminum and silver at room temperature, particularly in the (100) orientation. While many experiments on gold have been reported,5'9 none included the special orientations. We therefore undertook tension tests on gold crystals of the axial orientations (100),(110), (Ill), (211), and, as a representative of single slip, one whose Schmid factor was 0.5 (hereafter referred to as m = 0.5). Single crystals of dimensions 1/4 by .1/4 by 6 in. were grown from the melt1' at a rate of 4 in. per hr, using gold of 99.99 pct purity obtained from Handy and Harman. A growth rate of 1/8 in. per hr, or a purity of 99.999 pct,ll produced no difference in results. The crystals were annealed at 1000°C in air for 24 hr and furnace-cooled down to room temperature, after which they were electro-polished in a solution of potassium cyanide (40 g), potassium ferrocyanide (10 g), soda (20 g), and enough distilled water to make 1 liter of solution, at a current density of 0.02 amp per sq mm. After 2 or 3 hr, a very smooth surface was obtained by this method. Nine m = 0.5, two (loo), one (110), three (Ill), and one (211; crystals were tested at room temperature in a floor-model Instron machine at a tensile strain rate of 0.5 pct per min. The accuracy of the stress measurement was k10 g per sq mm up to 500 g per sq mm, k2 pct for higher stresses. The strain measurement was accurate to t2 pct. The scatter of stress at a given strain among the crystals of the same orientation was small, *5 pct being the largest. The representative stress-strain curves for various orientations are shown in Fig. 1. Table I summarizes the work-hardening parameters as used by seeger.12 Results of other investigators are also included in this table for comparison. There are no previous data for the corner orientations. Values for m = 0.5 crystals agree fairly well with those of Berner. Tm is defined as the stress at which the stress-strain curve begins to deviate from linearity of Stage 11. However, in practice this is a very difficult value to estimate because each investigator has a different idea of linearity. Therefore, the comparison with the values of other investigators may not be valid. In the present experiment, (100) crystals had the highest 111, followed by (111) crystals. The work-hardening rate in Stage I1 was highest for the (111) crystal followed by the (100) crystals. Our value for 0x1 of m = 0.5 orientation agrees very well with those of other investigators. 1) Single-Slip orientation (m = 0.5). These crystals were oriented so that the primary-slip vector was contained in one side face. The dimension perpendicular to this side should then not change if single slip indeed takes place. Within the accuracy of measurement, this dimension did not change during the deformation. Since the accuracy is k0.2 pct, the amount of secondary slip is less than 0.7 pct of the slip on the primary-slip system at 30 pct tensile strain. This is in conformity with the results obtained by Kocks" for aluminum crystals. The tensile axis moved toward (211) during the deformation. Slip bands, Fig. 2(a), were very fine and closely spaced. Some deformation bands were observed. There were no clear-cut cross-slip traces such as the ones observed on aluminum m = 0.5 crystals. 2) (111) Crystals. The orientation of the tensile axis was stable during the deformation. From this observation, one can deduce that at least three slip systems must have operated, and probably all six because the remaining three all have cross-slip relation to one of the first three.' It was very difficult to observe the slip markings. Consequently, we could not confirm by this method that six systems were operative during the deformation. At high strains (above 50 pct shear strain), this orientation had the highest stress-strain curve. At 80

Jan 1, 1964

By N. K. Chen, R. Maddin

Single crystals of molybdenum were extended at temperatures from 1300° to 2500°C. It was found that with increasing temperatures, the yield becomes more pronounced and the number of slip bands for equal amounts of elongation decreases. Slip at high temperatures fragments the structure. OTHER than the investigations on sodium and potassium by Andrade and Tsien,&apos; the researches of Steijn and Brick&apos; on a iron single crystals at low temperatures are about the only studies of the influence of temperature on glide of body-centered cubic single crystals. The effect of test temperature on glide has been rather widely investigated with hexagonal crystals" and less widely with face-centered cubic crystals. ~~ a result of these studies it can be said that, in general, the onset of plastic flow as measured by the initial rate of strain hardening becomes more abrupt with increasing temperature, and the value of the yield decreases with increasing temperature. Strain hardening, however, is markedly affected by changes in temperature except at very low and very high temperatures where the strain hardening rate becames constant. In the present investigation it is shown that temperature affects the glide of molybdenum single crystals in much the same manner as crystals of other classes, it., for comparable strains the number of slip bands decreases with increasing temperature; shear along the slip bands increases with increasing temperature; the yield becomes more pronounced with increasing temperature of test; and the strain hardening approaches a constant at very high temperatures. In addition, it is shown that, in agreement with the work of Tsien and Chow, Qlip occurs on (110) planes in <Ill> directions. It is further shown that slip at high temperatures fragments the structure into either irregular sized cells or bent atomic planes which are roughly aligned in crystallographic directions. Experimental Procedure Single crystals 3 mm in diameter by about 25 mm long were prepared using the technique previously described.&apos; The furnace in which the crystals were grown served as the extension apparatus for the tension tests. It consisted of a 4 in. water-cooled brass tube 20 in. long which was continuously evacuated at the bottom. Two water-cooled copper electrodes held the specimen by means of spherical collets secured in place in the electrodes by set screws. These electrodes served also to transmit power and stress to the specimen. The lower electrode was clamped to a stand, while the upper electrode was fastened to one side of a lever arm flexibly supported at the top through a slightly offset pinion. A metal bellows operating elastically was attached to the upper electrode to affect a vacuum seal and allow freedom of motion. The level of the arm was adjusted by rotating a graduated nut. Rotation of the electrode was prevented by a key fitted into a groove. After the specimen was mounted and the vacuum secured, the lever arm was adjusted so that the bellows was at its relaxed position. This insured that the total weight of the specimen, the upper electrode, and the downpull of the vacuum were balanced by the horizontal beam. Since the specimen was heated by its own resistance, expansion of the specimen proceeded freely by this arrangement. During the time of heating the specimen, the beam was maintained in a horizontal position by turning the nut at the top. By graduating the lever arm and using a known dead weight on the lever arm to supply a tensile stress, the graduated nut could be read to 0.0001 in. extension. Since

Jan 1, 1955

By R. A. Wood, R. I. Jaffee, H. R. Ogden, D. N. Williams

Strain-induced porosity has been found to occur in titanium and other materials at tensile strains greater than the uniform elongation of the material. Porosity in titanium increases with increasing strain and temperature, being most severe at 300° to 800°F. Strain-induced porosity is believed to be a normal part of the mechanism of deformation in ductile materials. DURING the past few years there have been numerous references to the occurrence of voids in laboratory test specimens of metals subjected to strain.l-7 These voids have been referred to as intercrystalline cracking, strain-induced porosity, intergranular cav-itation, voids, and porosity. They appear to occur at grain boundaries or interfaces, and are generally associated with large amounts of strain and have been found mostly in creep specimens and to a lesser extent in tensile specimens. Because the phenomenon does not occur unless the material has been strained,

Jan 1, 1960

By A. U. Seybolt

RECENTLY, cases of coexisting ordered and disordered phase equilibria have been reported among face-centered cubic metals,1"3 but not on body-centered cubic systems.4 McQueen and Kuczynski5 showed a single line as the boundary between ordered solid solution of Fe3Al type and corresponding disordered solution in the Fe-A1 phase diagram based mainly on dilatometric and magnetic measurements. Such a single line would indicate a second or higher order transition. or an absence of coexisting order and disorder. The writer found previously a two-phase equilibrium region at 800°C consisting of ordered and disordered solid solutions of Fe3Si in a matrix of approximately 1-6 wt pct Si (Fe3Si is of 14.3 wt pct Si). A similar two-phase region of ordered and disordered Fe3Al solid solutions might be expected if the similarity between Fe3Si and Fe3Si structures hold true. In the present investigation, a wire of 22 at. pct Al, balance Fe and 0.0513 cm in diameter was subjected to electrical resistivity measurement in vacuum as a function of temperature between 300 and 800°C, see Fig. 1. There was excellent agreement between data taken on heating and on cooling: no hysteresis was observed. The Curie temperature was determined to be 640 °C which agrees with the finding by McQueen and Kuczynski.5 But in addition, there were two transitional points in the resistance-temperature plot at 510oC and 400 "C. The former is thought to be a connection with the transition from disordered state to a coexisting order and disordered equilibrium state; while the latter, from such coexisting state to ordered state. Fig. 2 shows a phase diagram where the suggested two-phase region lies between dotted lines. The solid lines are those of McQueen and Kuczynski.5 Since these authors found only a single-phase boundary between ordered and disordered Fe3 Al, it would seem likely that the dilatometric method used by these authors was not sufficiently sensitive to resolve the ordering reaction into the two-phase boundaries. In addition, their phase diagram conflicts with the statement they make on P. 1 of their paper: "The sharpness of the peaks indicates that ordering to Fe3A1 is a first order reaction." The suggestion of coexisting order plus disorder in Fe3A1 must be considered tentative until confirmed by high-temperature X-ray examination. If

Jan 1, 1961

By Roswell P. Angier, Herbert S. Kalish, Henry H. Hausner

POWDER metallurgical methods as applied to zirconium are of great interest because they permit not only the fabrication of parts directly to shape with a minimum loss of material but also the utilization of scrap, and the metal may be processed at temperatures well below the melting point. There are two difficulties involved in the melting of zirconium: 1—High vacuum is necessary for melting, and 2—zirconium reduces nearly all oxides so that it is difficult to find a suitable crucible material. The purest material, zirconium crystal bar, is expensive, and the methods of subsequent fabrication are limited. Cold rolling or extrusion of zirconium crystal bar is suitable for making certain shaped parts, but where melting is involved, the purity of the metal is invariably decreased, and a method of casting to a definite shape does not seem to be feasible. Where it is necessary to hot roll, zirconium must be covered with a steel or copper jacket to protect it from the air. This is necessarily a costly operation. Where annealing of zirconium is necessary it must be conducted in a high vacuum. In an investigation of powder metallurgical methods for zirconium, the zirconium powder used should be characterized by the following properties: 1— Highest degree of purity, 2—uniformity in particle shape and particle size distribution, 3—low compression ratio, and 4—satisfactory strength of the compacted powder (green strength). Zirconium powders have been made by two methods: By the reduction of zirconium tetrachloride (ZrC1,) with sodium,2&apos; *•B "nd by the reduction of zirconium oxide with calcium and calcium chloride." The first method represents one of the earliest attempts to make pure zirconium metal, and the sintered product from this powder is extremely brittle. The second method of making zirconium powder is far more successful, but the powder made in this way has an extremely low apparent density and, therefore, a high compression ratio which makes it difficult to use it for powder metallurgical purposes. For these reasons the above-mentioned zirconium powders were not used in this investigation. To obtain sintered zirconium compacts of highest &apos;purity and good ductility, it was first necessary to produce a suitable powder. Pure zirconium has been made chiefly by two methods. The iodide method, from which the product is "crystal bar" zirconium, was developed first.&apos;." By this purification process very soft, pure, and ductile zirconium metal can be produced. The second method, developed by the U. S. Bureau of Mines, is the reduction of zirconium chloride with magnesium." The sponge zirconium thus produced is also soft, pure, and ductile.

Jan 1, 1952

By R. Taggart, D. H. Polonis, W. C. Gallaugher

In the Ti-Cu system, the a&apos; phase can be produced over a wide range of alloy composition witJwut the retention of measurable amounts of the ß or ? phases. This paper reports on the decomposition of this hexagonal martensite phase front the standpoint of solute corzcentratiorz, tempering temperature, and coherency strain condition. The mechanism of precipitation during tempering comprises localized precipitation which is followed by discontinuous precipitation if sufficient coherency strains are present. The localized precipitation process in hypoeutectoid alloys is described by the generalized Johnson-hiehl equation and an activation energy of 46,000 cal per mole. In hypereutectoid alloys the corresponcEing activation energy is 51,000 cal per mole. The rate-controlling process has been proposed as the difision of copper along a&apos; platelet interfnces. SEVERAL investigations have been reported concerning the kinetics of decomposition of the mar-tensitic a&apos; phase in titanium binary alloys.&apos; In the Ti-Mo system,&apos; two modes of a&apos; decomposition have been observed in the presence of /3 phase; one of these involves the precipitation of fine a from a&apos;, and the other involves diffusion across the a&apos;-ß interface. In studies of the Ti-Cr system, Rostoker2 did not detect the formation of TiCr2 during tempering. In Ti-Ni alloys,3 the intermediate phase Ti2Ni has been found to precipitate directly from a&apos; along the interfaces between plates. With the exception of the study of Ti-Ni alloys the previous investigations of tempering phenomena in substitutional martensites are mainly qualitative and do not present a detailed description of the precipitation processes. The following limitations restrict the correlation of the microstruc-tural processes and reaction kinetics during tempering in most binary-alloy systems. 1) A mixture of at least two phases characterizes the constitution of quenched alloys. 2) Difficulties have been encountered in obtaining uniform structures throughout quenched samples. 3) The reaction products, and in particular their morphology. have not been clearly resolved. 4) The martensitic a&apos; phase forms over only a limited composition range for most titanium-base binary alloys. Gomez and polonis4 showed that the Ti-Cu system provides an excellent basis for investigating the tempering of a&apos; over a range of composition without the complication of retained P or the transition w phase. In the present work the effects of solute concentration, tempering temperature, and coherency strain conditions are considered with reference to the over-all precipitation process in quenched alloys of the Ti-Cu system. Both microstructural observations and kinetic data are correlated to define the rate-controlling processes which govern the observed localized and discontinuous precipitation reactions. An earlier paper5 was devoted to a discussion of the modes of heterogeneous nucleation of Ti2Cu from the martensitic a&apos; structure. The present paper will therefore emphasize the progress of the tempering reaction beyond the initial stages. EXPERIMENTAL METHODS The Ti-Cu alloys for this study were prepared by conventional arc-melting procedures. Chemical analysis revealed that the alloy buttons contained approximately 0.02 wt pct 0, 0.04 wt pct N, and within 0.1 wt pct of the intended copper concentration. The progress of the tempering reaction was followed by means of electrical-resistance measurements utilizing a Kelvin double bridge, microhard-ness readings with a 400-g load, and X-ray dif-fractometry patterns to reveal line-broadening effects. Metallographic specimens examined at magnifications greater than X1000 were shadowed with germanium to reveal fine structural details. Direct carbon replicas were prepared for the electron-microscopy studies. EXPERIMENTAL RESULTS Property Changes. The changes of microhardness that accompany the precipitation of Ti2Cu from a&apos; are shown in Figs. 1 and 2. As the solute concentration increases the peak hardness, for a given tempering temperature, increases. In alloys of given composition the time to reach the peak hardness agreed with the time to attain maximum X-ray diffraction peak breadth, as shown in Fig. 3. The maximum hardness and the maximum line breadth increased with lower tempering temperatures, and the time to reach the maximum also increased. The a&apos; peaks broadened during the initial stages of tem-

Jan 1, 1965

By R. Stickler, A. Vinckier

An optical and electron microscope study was made on a commercial 302 stainless steel heat treated for massive carbide precipitation. Convincirg evidence was obtained that the grain boundary precipitate does not grow away from the grain boundary into the matrix, but grows as a network of large thin dendrites in between the grains. IN austenitic stainless steels the carbon in supersaturated solution precipitates when the steel is heated in the temperature range 450" to 1000°C. During such heat treatments the chromium-rich carbide occurs primarily at the grain boundaries and has adverse effects on the corrosion resistance and low-temperature ductility of the material. It has been shown that the morphology of the precipitated particles differs widely depending on the time and temperature of the heat treatment, the nature of the boundary, the misfit between the grains, and so forth.1"5 It was assumed by Mahla et al.,&apos; Streicher,&apos; and Mc Nutt&apos; that the carbide particles which nucleate in the grain boundary grow in various geometric or dendritic forms away from the boundary into the matrix. However, Kinzel,&apos; Hatwell et a, and Plateau et aL4 found that the particles nucleate in the grain boundary and grow as a network of thin dendrites in the grain boundary interface. We have made a comprehensive studyg of the precipitation occurring in a commercial 302 austenitic stainless steel heat treated for times ranging from 0.15 to 1500 hr at various temperatures from 480" to 1065"~. This study was made to elucidate the influence of heat treatments on both the mechanical and the corrosion properties of this steel. Part of this investigation which is pertinent to the morphology of the grain boundary carbides referred to above is separately reported here. We found that the particles which nucleated in the boundaries grew in the interface of these boundaries, thus supporting the viewpoints of Kinzel,&apos; Hatwell et al.,3 and Plateau et al.4 Furthermore, when carbide particles nucleated in the matrix very close to the boundaries, they possessed a morphology distinctively different from the particles growing in the grain boundaries. MATERIAL A study of metallographic samples and of fracture surfaces of a 302 stainless steel is reported in this paper. The composition and heat treatments are listed in Table I. Polished and etched samples were examined und813r the optical microscope, and carbon extraction replicas of etched micrographic samples and various fracture surfaces3&apos;8&apos;9 were examined under the electron microscope. RESULTS AND DISCUSSION A representative optical micrograph of the 302 steel, condition A, is shown in Fig. l(a). Particles can be seen on the grain boundaries but no traces can be found of particles which grow from the grain boundaries into the matrix. A carbon extraction replica of this sample shows numerous large thin dendrites, Fig. l(b). The lateral dimension of such carbide particles vary up to 100 p, and their thickness estimated from the amoun! of electron penetration i:; between 500 and 1000A. Such large dendrites are extracted from almost all grain boundaries and, if they had grown into the grains, they would definitely be revealed by etching traces in the matrix, Fig. l(a). The reason why these large dendrites on the extraction replica appear to be grown into the grain is due to the mechanism of the extraction replica process. Thus, these dendrites are supported by the carbon replica film only at the edges originally along the grain boundary trace A-A, Fig. l(b), in the polished metal section and have subsequently fallen over onto the carbon replica film during handling. Although such dendrites generally fall only to one particular side of the grain boundary trace, one finds occasionally that parts of a dendrite have fallen to both sides, as shown in Fig. l(b). Further evidence for the growth of carbide in the grain boundary can be obtained from the appearance of fracture surfaces. A small impact specimen of the 302 steel, condition A, was broken at liquid nitrogen temperature. The fracture path follows the

Jan 1, 1962

By G. R. Speich

The precipitation of the Feab and Fe,Ti Laves phases (MgZn, type, C14) from Fe-Nb and Fe-Ti solid solutions, respectively, has been studied in the temperature range 500" to 800°C using hardness measurements, light and electron microscopy, and X-ray and electron diffraction. The Fe-Nb alloys contained 0.96, 1.67, and 3.80 pct Nb and the Fe-Ti alloys 3.5, 5.8, and 8.8 pct Ti. In both systems only the equilibrium Laves phase is formed, with nucleation occurring at grain boundaries, at dislocations, and in the matrix. No cellular precipitation or transition phases were observed. The precipitation of the Laves phase in a fine dispersion leads to significant hardening in both systems. The hardness of the aged alloys is proportional to the reciprocal of the square root of the interparticle spacing. The maximum solid solubility of titanium in a! iron is 9.0 pct Ti at 1280°C. THE precipitation of the FezNb and FezTi Laves phases (MgZn2 type, C 14) from Fe-Nb and Fe-Ti solid solutions, respectively, has been investigated as part of a continuing study of the mechanisms of precipitation from ferritic binary iron-base alloys. The iron-rich portions of the Fe-Nb and Fe-Ti phase diagrams are shown in Fig. l(a) and l(b). The Fe-Nb diagram is essentially that given by ansen&apos; with the more recent work of Hume-Rothery and Gibson2 added. Goldschmidt3 indicates that the FezNb Laves phase has a solubility extending from 32 to 55 wt pct Nb. The Fe-Ti diagram is also essentially that reported by Hansen.l The recent results of Nishimura and kamei, Murakami, Kimura, and Nishimura,&apos; indicate that the FezTi Laves phase has a solubility extending from 22.5 to 30.5 wt pct Ti. The L/6 + FezTi eutectic temperature of 1298°C indicated by Kornilov and oriskina&apos; has been adopted. The addition of niobium or titanium to iron lowers the 6 — y transformation temperature and raises the y — a! transformation temperature. This results in a closed y field in the Fe-Ti system, but the lower solubility of niobium in a! (6) iron results in a eutec-toid, 6 — y + FezNb, at 1200°C, and a peritectoid, y + FezNb —-a, at 989°C in the Fe-Nb system. The maximum solubility of niobium is reported to be 4.5 wt pct in 6 iron,&apos; 1.0 wt pct in y iron,7 and 1.8 wt pct in a iron.&apos; Some results from the present work indicate that the maximum solubility in a! iron is actually less than 0.96 wt pct Nb. The maximum solubility of titanium in a! (6) iron is somewhat uncertain since the early work of To faute and Butting-haus indicated a value of 6.0 wt pct while later work of Kornilov and Boriskina6 yielded a value of 14.0 wt pct. The present work indicates that the maximum solubility is actually 9.0 wt pct. In both systems the phase in equilibrium with the iron-base solid solutions is a Laves phase (Mg znztype, C14), FezNb or FezTi. Both of these Laves phases have a wide range of Solubility. The precipitation-hardening behavior of Fe-Nb or Fe-Ti alloys has not been studied extensively. Genders and Harrison&apos; reported that binary Fe-Nb alloys hardened upon aging. Peter and Fisher7 studied the mechanical properties of Fe-Nb and Fe-Ti alloys quenched from different phase fields.

Jan 1, 1962

By D. G. Westlake, E. S. Fisher

The habit planes for zirconium hydride precipitation in crystals of a zirconium have been determined at various hydrogen concentrations. The (10 • 0)planes are the predominant habit planes; some (10 • 5) and (10 -1) planes were observed after charging conditions which pevmitted formation of a surface layer of hydride. No platelets zrlere observed parallel to any of the twinning planes. Mechanisms for em-brittlement due to hydride precipitation are discussed. LANGERON and Lehr1 found, in 1956, that zirconium hydride precipitates in zirconium in platelets lying parallel to (10.0) slip planes. In 1960, Kunz and Bibbz reported observing the platelets in three different families of twinning planes, (10.2) {11.1), and {11.2), but never in the slip planes. The present study was undertaken to ascertain the cause of these seemingly contradictory results. EXPERIMENTAL Two specimens of iodide zirconium were used in the experiments. Specimen I was a single crystal, while specimen II contained three grains. Grain growth was accomplished by the technique of Lange-ron and Lehr. Three faces were ground on specimen I; the first was 2 deg from the (00.1) basal plane, a second was 2 deg from the (10.0) prism plane, and a third was 4 deg from the (12.0) plane. One face was ground on specimen II such that the poles of the surfaces of crystals A, B, and C were as shown in Fig. 1. After the faces were ground on 600 grit silicon carbide paper, they were chemically polished by dipping in a solution of 8 parts hydrofluoric acid, 50 parts nitric acid, and 50 parts water. This removed all cold worked metal from the surface. The samples were charged to various levels of hydrogen content by heating in hydrogen gas which had been purified by reaction with uranium powder. The various charging conditions are shown in Table I. Specimen I was heated to the reaction temperature at the indicated hydrogen pressure and allowed to react until essentially all the hydrogen was used up. Then the temperature was raised for homogenization. Specimen II was charged by making periodic addi-

Jan 1, 1962

By R. W. Fountain, M. Korchynsky

The precipitation phenomena occurring in cobalt-tantalum alloys have been investigated in the temperature range frm 500" to 1050°C by correlating the results of metallographic, X-ray, micro-and macrohardness, and electrical resistivity studies. The property andmacrohardness,changes were found to depend on 1) general precipitation, and 2) lamellar precipitation. Two new intermetallic phases have been identified: 1) a Co3Ta, a metastable ordered face-centered-cubic compound, and 2) a stable ß Co3Ta phase of hexagonal structure. In addition, the previously reported Co2Ta phase was found to exist in two allotropic modifications: the hexagonal MgZn,-type and the cubic MgCu2-type Laves phases. SINCE a large variety of structures can result as a consequence of the decomposition of a solid solution, predictions on the nature of property changes are difficult, if not impossible, to make. For any rational attempt to correlate properties and structures of a precipitation-hardenable alloy, a detailed understanding of the kinetics of decomposition and morphology of phase separation, as well as knowledge of phase relationships, appears to be prerequisite. Information of this type has been accumulated in the past for many alloy systems, both of theoretical and pastforpractical importance.1,2 Although the presence of intermetallic compounds has been reported in cobalt-base alloys,3 the amount of published information on precipitation-hardenable cobalt-base systems is very limited. A survey of the binary phase diagrams of cobalt indicates that cobalt-tantalum alloys might be of interest as typical of other cobalt-base systems in which Laves phases of the A,B type can be precipitated from solid solution. The present work has been undertaken, therefore, to study the kinetics and morphology of the precipitation reaction in this system and to establish a base for a correlation between the structural aspects and properties in this class of alloys. PREVIOUS WORK The only available phase diagram of the cobalt-tantalum system is based on the work of Koster and Mulfinger. According to these authors, the maximum solubility of tantalum in cobalt is about 13 pct (at 1275°C) and. less than 7 pct at room temperature. Tantalum additions lower the temperature of allotropic transformation of cobalt (about 420°C), and at 7 pct Ta, the high-temperature face-centered-cubic modification (ß cobalt) is retained at room temperature. The precipitating phase was originally designated as Co5Ta2 compound (55.2 pct Ta, about 1550°C melting point), but subsequent investigations by wallbaum5" identified this constituent as the A,B-type Laves phase. Wallbaum&apos;s data indicate that there are two modifications of this intermetallic compound: one richer in cobalt (Co2.2 Tao.8)of the hexagonal MgNi, type; and another of a higher tantalum content (Co2Ta) of the cubic MgCu, type. On the other hand, Elliott7 found that the cobalt-rich alloy (CO2.10,Tao.~l) was predominantly the cubic MgCu, type at 800°C and a mixture of both the MgCu2 and the hexagonal MgZn,-type Laves phases at 1000°C. At 1200°C, Elliott found only the MgZn, type while at 1400°C, he observed only the MgCu2 type. At the stoichiometric composition, Co2Ta, Elliott reported only the cubic MgCu2-type Laves phase in the temperature range of 600oto 1600°C. The precipitation of the cobalt-tantalum intermetallic compound is accompanied by a marked increase in hardness. According to Koster&apos;s4 data, the Brinell hardness of an 8 pct Ta-Co alloy increases from 230 to 340 upon short-time aging at 800°C. EXPERIMENTAL PROCEDURE The binary cobalt-tantalum alloys investigated contained 5, 10, and 15 pct Ta. The range of tantalum additions was thus slightly broader than the reported minimum and maximum solid solubility limits of tantalum in cobalt (7 and 13 pct, respectively)4 The alloys were vacuum-induction melted in a magnesia crucible using cobalt rondelles and technically pure tantalum sheet as raw materials. Deoxidation of the melt was accomplished with carbon, and the chemical analysis of the alloys is given in Table I. The effect of isothermal aging treatments on the progress of precipitation was studied on samples cut from cast ingots. These samples were solution treated for 2 hr at 1250°C and water-quenched. Aging was conducted in the temperature range from 500" to 1050°C for periods between 15 min and 1000 hr and followed by water-quenching. To prevent contamination from the atmosphere, all samples were sealed in evacuated Vycor or quartz tubes for heat-treatments. For solution treatment, argon at 0.2 atmospheric pressure was introduced prior to sealing of the capsule to prevent collapse at high temperature, and titanium sponge was placed at one end of the capsule to act as a getter. MACROHARDNESS The effect of aging on Vickers hardness (Dph) of

Jan 1, 1960

By A Guinier

RECIPITATION in alloys is undoubtedly one of the most essential phase transformations in metallurgy and, besides, it is a phenomenon of great interest to physicists. It seems then that it can be chosen as a topic for an Annual IMD Lecture. Of course, the subject is a very wide one and I shall discuss here only a very restricted part of it. I should like to show why it is now considered that there is, at the beginning of the process, a stage distinct from the real precipitation, which is the pre-precipitation. We shall be chiefly interested in the atomic structure of the metal in this preprecipi-tation stage. Numerous reviews on this subject have been published recently.&apos;, &apos; The latest was given by H. K. Hardyh t the last meeting of the Institute of Metals in London. One of the most complete articles has been written by the late Dr. Geisler." I should like, at the beginning of this lecture, to pay a tribute to the memory of our friend Arthur Geisler. Some of the ideas which will be developed here are different from those advocated by him. But everybody knows the considerable progress in experiment and in theory which we owe to him and, personally, I deeply feel how fruitful have been the long discussions which we have had during many years. Experimental Basis of the Precipitation Problems Let us recall the essential facts pertaining to the problem of the precipitation. The starting point and the final one are well known. The initial stage is the supersaturated solid solution before quenching. The dissolved atoms replace some of the atoms of the solvent at the points of a definite lattice. It is generally considered that the distribution of the atoms is completely random, and that the lattice distortions are relatively weak. These assumptions are only an approximation. It has now beenproven", that, in some cases, the dissolved atoms have a tendency to segregate to form very small nuclei. These fluctuations of composition are greater than the normal statistical fluctuations in a purely random alloy. We shall see the role played by these nuclei later. The final state is obtained by a very long annealing at such a temperature that an equilibrium is reached in a reasonable time. The alloy is then composed of a matrix which is a saturated solid solution, much poorer than the initial one. Imbedded

Jan 1, 1957

By E. F. Sturcken, J. W. Croach

A grain orientation distribution function, P(u,F), was developed for use in predicting physical properties in oriented metals. Examples are given of the use of the function to predict thermal expansion coefficients in rolled uranium rods and plate. The P(u,F) function was synthesized from X-ray difpaction intensity measurements, Ii, norrnally used to plot inverse pole figures. The symbols u and F specify the angles that the crystallographic axes of the grain make with the normal to the sample surface and P(u, F) gives the number of grains oriented in each direction, (u,F). P(u,F) is obtained from a least squares fit of the Ii to an orthonormal set of spherical harmonic functions. The P(u, F) function was also used to define a quantity, J, which is a measure of the departure of the orientation distribution from unifomzity (or "randomness"). The J parameter may be used to follow, quantitatively, the amount of preferred orientation induced as a function of mechanical working or heat treatment. An example is given of the application of the J parameter to prefevred orientation studies of hot- and cold-rolled uranium rods and hot-rolled uranium plate. Because P(u,F) is derived from relative intensity measurements by the powder diffraction technique, it is applicable to material with large absorption coefficients and/or grain size and is particularly suitable for characterizing the preferred orientation of large numbers of samples such as fuel elenzents for a nuclear reactor. IN general the physical properties of an oriented metal depend on the direction in which the properties are measured. Hence to predict physical properties the grain orientation must be known as a function of the direction of the specimen. The conventional methods of expressing preferred orientation measurements are the pole figure1 and the inverse pole figure.2-8 Both methods are difficult to use for predicting physical propertiess because they are graphical and because they are stereographic projections and hence are not "area-true." In the present report a quantitative orientation distribution function, P(u,F), is synthesized from pole density measurements normally used to plot inverse pole figures. For inverse pole figure plots the pole densities, pi, are obtained by powder dif-fractometer intensity measurements of their reflections from a flat surface of the sample, normal to the direction of interest. Each pole density, pi, represents a single orientation, i, and is proportional to the volume of material with orientation, i; hence a set of pi gives an approximation for the general orientation for the sample direction measured. The expression8&apos;7 for calculating the pole densities, pi, from the diffraction measurements was previously

Jan 1, 1963

By James T. Waber, Allen C. Larson, Karl Gschneidner, Margaret Y. Prince

The original graphical method proposed by Darken and Gurry for predicting which elements are soluble in a given element was slightly modified and was used to predict terminal solubilities in 1455 known binary combinations. it was found that the modified Darken and Gurry method had a reliability of 61.7 pct in predicting extensive solid solubility (greater than 5 at. pct) and of 84.8pct in predicting limited solid solubility (less than 5 at. pct), an over-all reliability percentage of 76.6. Using the same experimental data, it was found that using only the simple Hume-Rothery size-effect criterion was 50.1 pct reliable in predicting extensive solid solubility, only slightly better than a purely random choice. and 90.3 pct reliable for predicting limited solid solubility, an over-all reliability percentage of 67.6. THE first rational rules for determining a priori whether two metals are not freely soluble in the solid state were announced by Hume-Rothery and coworkers.1,2 These well-known rules have been accorded a special place in metallurgical literature. Since the "electronegative" and "relative valence effect" rules of Hume-Rothery were only qualitative in nature it was difficult to use them in a quantitative analysis. The next major step was made by Darken and gurry3 and apparently went unexplored for a number of years. They prepared graphs illustrating the influence of both electronegativity and size difference on the solubility of elements in silver, magnesium, and aluminum, and they drew elliptical figures, whose maximum width was ±15 pct of the solvent radius and the maximum height was ±0.4 electronegativity units. The purpose of the present paper is to assess and compare the effectiveness of using the simple size-factor and the Darken-Gurry criteria to predict which elements will be soluble and insoluble in a given solute. In 1957, waber4 prepared Darken-Gurry plots for determining the solutes that were likely to be soluble in 6 and plutonium and introduced an inner ellipse as an aid to determining which elements should be soluble to a considerable extent. As illustrated in Fig. 1, the major axis of the outer ellipse corresponded to ±0.4 electronegativity units (OD) while the minor axis corresponded to ±15 pct of the radius of the solvent phase (OB). This simultaneous "action" of the two Hume-Rothery criteria was used to indicate which elements would have little or no solubility in the host lattice, namely those elements whose points lay outside this ellipse. An inner ellipse, "centered" over the same point of the solvent was drawn with ±0.2 electronegativity units (OC) and ±8 pct (OA) as the major and minor axes. Solutes with points lying within the inner ellipse were expected to be capable of forming extensive solid solutions with the host element. Use of this method of applying the earlier Hume-Rothery criteria was moderately successful for plutonium. For 6-phase plutonium, only one element out of the eight predicted to be extensively soluble was known to be almost insoluble, and of the sixteen elements predicted to have more restricted solubility, five were known to be essentially insoluble. The success of discerning solubility in the E phase by this method was comparable. One prediction made on the basis of this analysis was that the heavier rare earth metals, such as holmium, erbium, and thulium, should be soluble in the 6 phase, whereas the lighter rare earth elements, such as lanthanum or

Jan 1, 1963

By W. A. Tiller

SEVERAL authors1-6 have shown that, during solidification from the melt, the direction of formation of substructure boundaries depends upon the direction of heat flow and the rate of solidification of the metal. The most prominent of these observations was that a preferred orientation developed during the solidification of the columnar zone of an ingot which for face-centered-cubic metals was the <100> direction. Previous explanations1-&apos; given to account for this preferred orientation and the other boundary phenomena are all based on the premise that the phenomena are characteristics of the pure element. However, recent experiments by Rosenberg and Tiller7 have demonstrated that the preferred orientation in pure lead is the <111> orientation, and that the heretofore observed <100> orientation is due to the effect of solute on the mode of solidification of the metal. Teghtsoonian and Chalmers,4 studying striation boundaries in tin, and Chalmers and Rutter,5 studying corrugation boundaries in tin, showed that these substructure boundaries were aligned parallel to the axis of heat flow for slow rates of growth, and that the alignment varied from this direction as a function of the rate of growth. to approach the den-drite direction at high rates. Chalmers,6 studying the direction of formation of a grain boundary separating two crystals of different orientation, showed that at slow growth rates the boundary is parallel to the axis of heat flow, while at higher speeds it will deviate from this direction, sloping into the crystal whose dendrite orientation departs the most from the axis of heat flow. It is this phenomenon that allows certain crystallites in the chill zone region of an ingot to encroach on their neighbors and ultimately produce the preferred orientation of the columnar zone of the ingot. Since the previous explanations of this preferred orientation are unable to account for the recent observations,&apos; a new mechanism will be considered. This will lead to an explanation of the other substructure boundary phenomena as well. Investigations by Graf,8 Billig,9 Rsenberg,10 and others have demonstrated in a striking fashion that crystal growth takes place by the deposition of individual layers oriented along the close-packed planes of the material considered. This lamellar or

Jan 1, 1958

By S. E. Bronisz, R. E. Tate

s-plutonium samples possessing a strong growth texture have been produced by allowing them to transform under pressure from p to a. A fiber texture with [010] parallel to the pressure axis results. The textured samples are expected to be useful in determining the effects of crystallographic orientation on the physical properties of a plutonium, and initial results of measurements of resistivity and mechanical properties are described. THE crystal structure of a plutonium1 is mono-clinic, P21/m, and many of the metal&apos;s properties should reflect the crystallographic anisotropy. Unfortunately, it has not yet been possible to grow useful sizes of single crystals of a plutonium. Therefore, few of plutonium&apos;s orientation-sensitive properties have been determined as a function of crystallographic orientation. One way to gain information about at least some of the orientation-dependent properties is to study a highly textured sample. Such samples have been produced by cold rolling a plutonium,2 but they are necessarily very small in one dimension. We have now successfully used two other techniques to produce highly textured samples, not limited to small thicknesses. METHODS OF PRODUCING PREFERRED ORIENTATION The texture is produced by causing the plutonium to transform from ß to a while under pressure. It was observed by Gardner3 that applying a compres-sive load during the transformation resulted in the formation of columnar grains of a plutonium, and Cramer4 observed a similar structure in some diffusion couples rolled in the P phase. Neither of these workers, however, investigated the subject beyond these observations. The more flexible of the two methods developed is that of compressive loading. The technique is to enclose a sample of plutonium in a die assembly before applying the load. It is possible, of course, to dispense with the die and simply press the sample between two platens. In this case, however, the unrestrained sample deforms during the pressing, making it difficult to maintain the load and to control the final sample thickness. This compressive loading technique has been used on specimens ranging from 1/4 to 3 in. diameter, with height-to-diameter ratios ranging from 5 : l to 1:6. The procedure begins with loading the sample, a right cylinder of a plutonium, into the pressing die. Plastic deformation of the sample while it is being pressed is minimized by machining the sample so that its diameter is nearly the same as that of the die barrel. Then the loaded die assembly is placed in a vacuum pressing can that fits on a hydraulic press. After a vacuum is achieved the die assembly is heated to 180°C and equilibrated. The load is then applied and the specimen and die are allowed to cool to about 40° to 50°C before the pressure is removed. The resulting sample is composed of rod-shaped a -plutonium grains, the rod axes being parallel to the cylinder axis and the pressing direction. Photomicrographs of the oriented structure are shown in Fig. 1. Although most of the specimen is composed of grains whose rod axes are parallel to the cylinder axis, a thin, peripheral layer consists of rod-shaped grains whose rod axes are perpendicular to the cylinder axis. Presumably, the highly plastic nature of p plutonium leads to approximately hydrostatic conditions during the initial stages of the P — a transformation under pressure. At this time, the grains in the peripheral layer would have grown parallel to the local force direction. During the latter stages of the transformation, the accompanying volume contraction would have eliminated the radial forces, leaving only the axial force and resulting in the growth of the observed structure. The growth texture did not form without the application of pressure. Attempts were made to produce the texture by repeating the standard temperature cycle without

Jan 1, 1965

By L. K. Jetter, J. C. Ogle, C. L. McHargue

The preferred orientation developed in extruded aluminum rods has been studied as a function of extrusion temperature, extrusion speed, and position in the rod. Duplex <111> - <001> textures were developed for slow extrusion speeds at temperatures of 75° to 850° F and for fast extrusion speeds at temperatures of 75° to 450°F. The <001> texture component was shown to be primarily due to re crystallization occurring during fabrication. At fast speeds and billet temperatures of 650° to 850°F, complete recrystallization occurred and <115>or<118> textures were found. The nature of the structure has proved to be an important variable in the production and use of extruded products. Data on the flow of aluminum and aluminum alloys summarized by Smith,&apos; Hardy,&apos; and Locke3 show that the extruded product may possess a deformed, a partially recrystallized, or a completely recrystallized structure, and there may be important variations in the structure from surface to center and from front end to back end of an extrusion. This inhomogeneity in grain structure and preferred orientation affects both the mechanical properties of the extruded section and its response to subsequent heat treatment. Among the fabrication variables causing these differences are temperature, extrusion ratio, and extrusion speed. Preferred orientation which develops during the extrusion of rods is of the fiber-type and, in general, has been assumed to be similar to that developed in rods or wires fabricated by drawing through dies, rolling, or swaging.4 Because of the nature of the flow of metal being extruded, this assumption may be valid for material near the center of the rod but probably is not valid for the outer layers. In most face-centered-cubic metals a duplex fiber texture has commonly been reported in which the grains have either <Ill> or <001> directions parallel to the longitudinal axis. Sachs and fichiebold5 reported cold-drawn aluminum (99.6 pct pure) to have a single <Ill> fiber texture, and this observation was confirmed by Schmid and Wassermann.6 Hibbard, on the other hand, observed a minor <001> and a major <111> texture in 99.994 pct A1 cold drawn to 98 pct reduction in diam.7 He concluded that the <001> component resulted from a small amount of room-temperature recrystallization. Dayal8 concluded that the <001> component was an intermediate component which disappeared at high reductions in 99.5 pct pure Al. Kostron and schippers9 found duplex <111> — <001> textures in extruded 99.5 pct Al, as did Gow and cahnl0 and GOW11 for extruded rods of commercial and superpure aluminum. Similar duplex fiber textures have been reported for aluminum alloys by Kostron,= Unckel,13 and Van Horn.14 The present paper is concerned with the variation in texture with fabricating conditions and with the nature of the texture components in extruded aluminum rod (99.99 pct pure). EXPERIMENTAL PROCEDURE Ingots were chill-cast in cast-iron molds and machined to 3-in. diam by 6-in. long extrusion billets. These billets were upset to 3.125 in. diam and extruded through a flat-face 0.690-in. diam die. Quenching of the rods was accomplished by a water spray at the die face. The combinations of fabricat-

Jan 1, 1960

By S. L. Lopata, E. B. Kula

A variation of yield and tensile strength with direction has been noted in heat-treated 4340 steel which had been warn-worked by rolling in the austenitic condition prior to quenching. Measurements by X-rays showed a preferred orientation of martensite crystals. Various ideal preferred orientations of martensite were calculated by assuming a preferred orientation for austenite and transforming this to martensite by the Kurdjumov -Sachs and Nishiyama relationships. Comparison with the measured preferred orientation showed that the ideal rolling texture for austenite in 4340 steel could be expressed as a (112) [111] texture with a component of (110) [112]. No evidence was found for any fiber axis in the rolling direction, or fop- a (123) [412] texture. The observed martensite pole figures can be adequately accounted for by these ideal austenite textures and by the Kurdjumou-Sachs or Nishiyama transformation relationships between austenite and martensite. ALTHOUGH preferred orientations are found in cold-rolled steel sheet and in many nonferrous metals, they are not generally encountered in wrought, heat-treated high-strength steels, since such steels are generally not cold worked sufficiently to develop a strong texture, and any texture that is present is removed during heat treatment by the phase transformations taking place. This report gives details on a preferred orientation found in warm-rolled and hardened 4340 steel. In a recent study, a special combination heat-treating and working technique was used in which 4340 steel was deformed in the austenitic condition, and before re crystallization of the austenite could occur, quenched to form martensite.&apos; Since the austenite was unrecrystallized, it was cold worked and hence could have contained a preferred orientation which could be transmitted to the martensite on quenching, and this in turn could be manifested as a directionality of mechanical properties. Such a directionality of properties was found as a result of working, and this prompted an X-ray study of the preferred orientation. Literature Survey— Rolling Texture in Face-Cen-tered-Cubic Austenite—M a preferred orientation is found in a rolled and hardened steel, it is first appropriate to consider what texture existed in the austenite as a result of the rolling but prior to the transformation to martensite. The changes taking place as a result of the austenite-martensite transformation are then considered. In view of the instability of austenite in 4340 steel at low temperatures, no direct determination of the austenite texture was made. However, an estimate of the texture for austenite in 4340 steel may be made by considering the textures found for other face-centered-cubic metals. Barrett&apos; has reviewed the rolling textures for many materials. The most common orientation in face-centered-cubic metals has the (110) planes parallel to the rolling plane and the [112] directions parallel to the rolling direction, or (110) [112]. Other orientations found are (112) [111], as well as (100) [001], (110) [001], and others. Beck and his co-workers, using quantitative pole figures, report that (123) [121],3 later corrected to (123) [412],4 better describes the ideal orientation. smallman5 reported that the ideal preferred orientation for high-purity face-centered-cubic metals is (123) [121], but that there is a transition to (110) [112] as solid solution elements are added. The latter texture tends to occur more readily at lower temperatures. This transition between the two textures was explained by differences in the slip characteristics with concentration or temperature. More recently, Jones and Fell6 concluded that the (123) [412] texture for face-centered-cubic metals could equally well be represented by a mixture of (110) [112] and (112) [111] textures, and hence it had no physical reality. From the above results it can be seen that there is no agreement on the "ideal" texture for fcc metals and, hence, no firm basis for assuming a texture for austenite in 4340 steel. Therefore, the orientations (110) [112], (123) [412], and (112) [111] have been considered in the following.* *The (123) [4121 is identical to the (123) [121] rotated 90 deg. Orientation Relationship between Austenite and Martensite—When martensite forms from austenite, certain crystallographic planes of the martensite are parallel to planes in the austenite. This orientation relationship was first determined by Kurdjumov and

Jan 1, 1960

By R. K. McGeary, B. Lustman

The textures produced in zirconium by cold and hot rolling, and by recrystallization above and below the transformation temperature were determined. Thermal expansivities were measured in the thickness, transverse, and rolling directions of preferentially oriented zirconium and were correlated with the texture scatter in these directions. REVIOUS investigations have indicated that minor differences between hexagonal close-packed metals of similar axial ratio may appear with respect to the textures produced both on cold rolling and on subsequent recrystallization. In the case of magnesium, beryllium, and titanium, metals of axial ratio similar to that of zirconium, the ideal orientations produced by rolling are fundamentally the same, although marked variance is reported in the degree and type of scatter about the mean orientation; in those instances where recrystallization textures were observed, they were reported to be similar to the rolling textures. Measurement of the anisot-ropy of thermal expansion of both rolled and re-crystallized zirconium could not be correlated satisfactorily with the textures reported for the above metals, and therefore a study was made of the preferred orientations produced in zirconium. Reported below are the textures produced in zirconium by cold and hot rolling, and recrystallization above and below the transformation temperature, together with the results of thermal expansion measurements. Determination of Preferred Orientation Two types of zirconium were investigated: 1— "crystal bar" zirconium obtained from the Foote Mineral Co., produced by the thermal decomposition of zirconium tetraiodide, and 2—zirconium ingot obtained from the Bureau of Mines prepared by melting sponge zirconium in a graphite resistor vacuum furnace in a graphite crucible. The major impurities present in the two materials used are listed in Table I. Several of the pole figures were later checked with 0.03 pct hafnium crystal bar material and the results were identical with those to be shown for the 1.5 pct hafnium material. The materials were cold rolled to 0.014 in. in thickness as shown in Table 11. Specimens were cut from the 0.014 in. thick rolled sheets and etched to thicknesses of 0.002 to 0.010 in. Such specimens were used for exposures up to a 50&apos; to 60" angle between the beam and plane of the specimen; for higher angles a wire shape, similar to that described by Bakarian,&apos; was formed on an end of the original 0.014 in. sheet. A fine-bladed abrasive cut-off wheel was used to slot the sheet and to form the cylindrical cross-section. The wire shaped ends were then etched to 0.006 to 0.010 in. in diam. Although absorption of X-rays in the wire-shaped specimens does not vary with angle of rotation, the line width around the diffraction rings was not uniform, because the wire was narrower than the X-ray beam, and this condition caused some uncertainty in the estimation of azimuthal intensities. Furthermore, scanning was not practicable with this type of specimen so that spottiness of the rings due to large grain size was excessive for specimens which had been heated above about 650°C. Nevertheless, satisfactory information could be obtained for high angle exposures from the negatives by the use of both types of specimens. Transmission Laue photograms were taken using unfiltered molybdenum radiation (47.5 kv, 18 ma) and a 0.025 in. pinhole. With the film 8 cm from a 0.005 in. thick specimen exposures of about 30 min were adequate. For specimens with a coarse grain size, a device that scanned about 0.15 sq in. of sheet surface was used. An attempt was made to plot the pole figures by use of an X-ray spectrometer as described by Norton.&apos; However, for the particular technique used, the intensity variations obtained were not considered definite enough to give reliable results, especially for the large grained recrystallized and transformed specimens. This method was therefore abandoned in favor of the standard photographic method. Nine exposures were taken of each specimen: seven exposures with the beam perpendicular to the rolling direction and at 0°, 10°, 20°, 35", 50°, 65", and 80" to the transverse direction, and two exposures with the beam perpendicular to the transverse direction and at 60" and 80" to the rolling direction. Additional exposures were then made where necessary. The intensity variations of the diffraction rings were estimated by eye. It was usually possible to estimate 3 degrees of intensity from the photograms but in some cases 2, 4, or 5 degrees were estimated. Experimental Results The preferred orientation was determined for the following treatments: 1—cold-rolled, 2—low temperature rolled, 3—cold-rolled surface layer, 4— cross-rolled, 5—hot-rolled, 6—recrystallized below the transformation temperature, and 7-—recrystallized above the transformation temperature. I—Cold-Rolled Textures: The slip plane in hexag-

Jan 1, 1952

By J. J. Kramer, W. A. Tiller, G. F. Bolling

The solidification of poly crystalline zone-refined lead has been examined. A novel casting technique was used, with several advantages such as unidirectional heat flow, atmosphere control, and decanting of the liquid during any stage of growth. The experimental results show the separate existence of two textures, a (111) surface texture at the chill surface and a texture in the columnar zone, in the absence of conditions which would produce dendritic growth. WALTON and chalmers have most recently shown that the dendrite direction is the preferred axial orientation for the solidification of many metals in casting. Experimentally, for aluminum and lead, they showed that in the columnar zone the dendrite direction, <l00>, predominated, developing from an equiaxed chill cast layer that exhibited no preferred orientation. A reasonable description of the way in which grains closer to the dendrite orientation were able to encompass other grains was then set forth. For their materials, it was also shown that high superheats could inhibit a preferred axial orientation from developing, presumably by inhibiting dendrite growth. The observations of Rosenberg and iller, on the other hand, showed that zone-refined lead, cast under conditions similar to those of Walton and halmers,&apos; had a definite <111> texture in the columnar zone. However, no examination was made of the chill surface to find out if the columnar zone texture was independent of the orientations at the surface. Since the existence of preferred growth in Upure" lead is thus in doubt, a reexamination of the casting of zone-refined lead is presented here, using a novel casting technique. EXPERIMENTAL PROCEDURES AND RESULTS Zone-refined lead was prepared from two starting materials;99.999+pct supplied by ASARCO, and 99.999pct supplied by COMINCO. The manner of zone-refining and the phenomenological tests for purity after refining, have been describedelsewhere.&apos; The casting apparatus is shown schematically in Fig. 1. The following general procedure was used: The metal was cut directly from the best parts of

Jan 1, 1963

By A. H. Geisler, J. H. Keeler

Preferred orientations in rolled and annealed titanium sheets were determined by the Geiger counter spectrometer X-ray diffraction technique. Five annealing textures dependent upon the temperature range of annealing were found, and in order of increasing annealing temperature pendent upon the temperature range are: 1—a deformation like texture, 2—a rotated inorder a-recrystallization temperature texture, 3-a retained u-recrysraIlization texture, on annealing at lower temperatures of the ß-region, 4—a transformation texture based on recrystallized a and predicted by the Burgers&apos; relationship, and 5—a ,ß-cube texture. These results are examined in terms of current theories of recrystallization textures. UMEROUS investigators have described the tex- ture obtained by cold rolling the hexagonal metals, titanium, zirconium, and beryllium, which have c/a ratios less than that of ideal packing, 1.633. The basal planes are rotated out of the rolling plane, about the rolling direction, so that the basal poles are tilted toward the transverse direction as shown schematically in Fig. la. In all instances but one,&apos; it was also reported that the [1010] direction was parallel to the rolling direction (see Fig. lb). Hot rolling has been reported as causing a similar tilt of the basal poles in the transverse direction (see Fig. la) and causing the [1010] direction also to be parallel to the rolling direction as shown schematically in Fig. lb. Annealing after deformation does not appreciably change the tilt of the basal poles in the transverse direction." Beryllium2-7 continues to have the [1010] direction in the rolling direction after annealing, and similar observations for titanium and zirconium&apos; . have been reported for annealing at fairly low temperatures, again as in Fig. lb. At higher annealing temperatures, however, the recrystallized grains of titanium" and zirconium have an orientation such that the [1120] direction is approximately in the rolling direction, although the basal poles are still inclined in the transverse direction. Figs. la and lc show the resulting orientations schematically. This change in orientation has been described as a nominally ±30° rotation of the hexagonal crystallites about the basal poles of the cold rolled texture and is apparent from the results which are summarized in Table I for investigations with the X-ray diffraction technique employing film. The angles y, , and ß are indicated in Fig. 2 which represents the stereographic projection of (1070) poles for the mean orientation of a pole figure. Texture determinations for titanium using the Geiger counter spectrometer have provided similar results except that in some instances additional components of the texture were proposed, as shown by the summary of data in the upper half of Table 11. On the other hand, the spectrometer technique, when applied to zirconium,* has revealed a splitting Recently completed studies of the textures of annealed zirconium", show zirconium to possess textures very similar to those reported here for titanium. Therefore, much of this discussion will include zirconium by virtue of its close similarity to titanium in pref erred orientations. of the intense areas of the pole figure for samples annealed at 600°C. This splitting could be described by a 7" rotation of the tilt axis about the normal to the rolling plane. Such a splitting for the annealed texture relative to the cold rolled texture was not observed in other determinations for either zirconium or titanium using the less sensitive film X-ray methoe and makes the relationship between the two types of texture more complex than the simple rotation about the (0001) pole based on film work. The more precise investigations on zirconium permit the descriptions in the lower part of Table 11, which show that the texture depends quantitatively on the temperature of annealing. When zirconium is annealed at temperatures up to 400°C, the texture is similar to the cold rolled texture, while annealing in the range 500" to 900°C produces a texture which is only approximately described as [11%] in the rolling direction. More precisely described results for zirconium show that the two types of splitting ( 1—about an axis in the rolling plane through an angle given in the second column in Table II and 2—about the normal to the rolling plane through an angle given in the third column of Table 11) depend on annealing temperature. The [1120 is the rolling direction only when the annealing temperature is in the vicinity of 900°C

Jan 1, 1957