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By J. T. Norton, H. Blumenthal, S. J. Sindeband

The interstitial phases formed by the transition elements with carbon, nitrogen and boron constitute a unique class of substances which are of considerable technical interest because of their well developed metallic properties and their high hardness and melting points. They are also of interest from the point of view of structure. It was pointed out by Hagg1 some years ago that the principal factor determining these structures was the relative sizes of metal and metalloid atoms. He predicted that if the ratio of the radii of metalloid to metal atoms was less than 0.59, the metal atoms would be arranged in a close-packed fashion of relatively simple type. This prediction has been amply confirmed. It appears, however, that this specification applies most generally to those structures in which the metalloid atoms occupy isolated positions in the lattice and there is no tendency for strong binding between the metalloid atoms themselves. The series of borides of the type MeB2 formed from the transition elements of the fourth (Ti, Zr) and fifth (V, Cb, Ta) groups of the periodic table provide an opportunity to investigate this situation further. If one takes as a basis the radius of 0.87 A for the boron radius and the Goldschmidt values for the metal atoms in 12 coordination, then the radius ratio varies from 0.54 for zirconium to 0.64 for vanadium, which latter is considerably greater than Hagg's limit of 0.59. It was of interest, therefore, to see if these five elements would form borides of the MeB2 type and have simple isomorphous structures. When this investigation was initiated, relatively little information had been published on the structure of this group of borides. McKenna2 reported the preparation of ZrB2 by a carbon reduction process. Its structure was hexagonal and the approximate lattice constants were given as a = 3.15 A, c = 3.53 A. Ehrlich3 reported the structure of TiB2 as being of the C 32 type with a = 3.02, c = 3.21 and c/a = 1.06 A. Recently, Kiessling4 published the results on ZrB2, showing it to be one in which the metal atoms are arranged in a simple hexagonal lattice with the boron atoms in flat sheets, midway between the layers of metal atoms. The structure is C 32 type with a = 3.169, c = 3.530 and an axial ratio =1.11 A. He also states that CbB2 and TaB2 are isomorphous with ZrB2. Preparation of Borides The borides of titanium, zirconium, columbium and tantalum were made by the electrolysis of fused salt baths after the method described by Andrieux.5 The bath, totalling about 6500 g in each case had the following general composition: Metal Oxide + B2O3 + CaO + CaF2. The actual proportions of the ingredients were determined for each individual case. Although Andrieux reports the preparation of vanadium boride by the same electrolytic method, several attempts proved unsuccessful. However, it was possible to prepare VB2 by the carbon reduction method used by McKenna to produce ZrB2. The mixture consisted of V2O5 + 6B2O3 + 11C and the reduction was carried out in an induction heated graphite crucible. This method proved satisfactory for ZrB2 and TiB2 and undoubtedly the others could be made in the same way. The borides, which were all in the form of fine gray metallic crystals from the electrolytic bath or gray powders from the reduction process, were carefully separated from other constituents and analyzed chemically for metal and boron content. The densities of the boride powders were measured using a standard 10 cc pyenometer at 18°C and ethyl benzene proved to be the most satisfactory liquid for this purpose. X Ray Technique The boride crystals, as prepared, were too small for single crystal examination so that powder techniques were employed. Patterns of the finely ground borides were prepared, using the Norelco recording X ray spectrometer. The lines on these patterns were indexed with the aid of Hull-Davy charts and the approximate dimensions

Jan 1, 1950

By J. J. Gilman

Zinc monocrystals, when compressed nearly parallel to their basal planes (within 20), deform in a new way; and certain deformation markings, called "-bands," are typical of this new mode. Characteristics of these bands are described. Mechanical behavior of low angle zinc monocrystals was investigated in some detail. Recrystallization did not occur upon annealing of banded specimens that were free of deformation twins. -Bands do not seem to be the result of nonbasal slip, but rather, a result of a microbuckling mechanism. u-Bands differ qualitatively and quantitatively from ortho kink bands. A dislocation model, consistent with some of the features of u-bands, is described. THE plastic deformation of zinc monocrystals with slip planes (0001) nearly parallel to the stress axis was studied in this investigation. In a previous investigation, it was shown that some low angle zinc crystals (those with slip planes making angles of more than about 2.5+1° with the compression axis), deform by ortho kinking while others (those with slip planes making angles of less than 2.5+1° with the compression axis) deform in a new way.' This new mode of deformation will be described in some detail here. Some of its striking features are: 1—it operates only for compressive stresses, 2—its temperature dependence is anomalous, and 3—it exhibits a new type of "yield-point" behavior. The mode of deformation to be discussed is not new in the sense that no previous evidence of its existence had been reported. The markings that it creates on the cleavage surfaces of zinc crystals have been described by several investigators.'" However, this is the first time that it has been made to operate under controlled conditions so that its response to mechanical variables, temperature, and aging treatments could be investigated and its geometry completely described. Apparently. Jillson' as well as Holden3 and Pratt" believed that the small bands they observed were simply small versions of ortho kink bands. In fact, two of these authors3,6 "alled the markings "micro kink bands." In spite of the precedence, it is felt that this term is not entirely suitable because these bands do not have the same genesis as kink bands, and they are qualitatively different in structure. As will be indicated by later discussion, the bands seem to be related to the tensile deformation bands observed in aluminum. The present author has chosen to call the bands "micro deformation bands," or simply ",ubands," and the deformation process, ''p-banding." Experimental Methods Crystals were grown from 99.999+ pct Zn obtained from the New Jersey Zinc Co. The details of the technique are given elsewhere." It should suffice to say that the crystals were seeded for orientation control, and precision pyrex tubes were used for molding them. About 40 crystals were used, yielding some 130 specimens. Fig. 1 illustrates the typical form of the specimens and shows how the orientation angles are defined. The crystal specimens were usually 1 3/4 in. long and about 150 mil in diameter. The angle between the slip or cleavage plane (0001) and the specimen axis is designated by xo, while a is the angle between the maximum slope of the (0001) plane and the nearest close-packed direction [1210], vary ing from 0°to 30°. The crystals were cut into specimens with a jeweler's saw and the ends were squared by filing (specimens held in V-block). Crystals of this orientation are stiff toward bending, so this procedure introduced only deformation localized at the ends. Then the specimens were polished either chemically in 50 pct HNO³ or electrolytically in 20 pct chromic acid. The mechanical tests were made with an Instron testing machine. This machine weighs by means of electric strain gages on elastic bars and is quite stiff, since the full-load deflection of its load cells is at most 0.002 in. Most of the testing was done at a

Jan 1, 1956

By David A. Thomas, John F. Peck

Fibrous microstructures and their development have been studied by metallography and by hardness and quantitative metallographic measurements. Thin, curved grains were observed in transverse sections of commercial tungsten wire and were developed in laboratory-drawn iron wire at reductions greater than about 70 pct. The unusual grain shapes are attributed to the nonuniform mode of deformation of the grains associated with the characteristic [I101 preferred orientation that develops in bcc metals during drawing. Transverse sections of heavily swaged iron wire showed a ospiral nebulao structure that appears to result from the elliptical shape of the cross section of the swaging dies and their rotary motion. Hardness and quantitative metallographic results show unusual trends that parallel the changes in microstructure. Other phenomena are considwed in light of the results, including the strength of fine wires of bcc metals, ocylindricalo textures, and the ductility of fibrous tungsten. 1 HE microstructure of tungsten wire in a longitudinal section shows grains highly elongated in the working direction, giving the appearance of a bundle of parallel fibers, Fig. l(a). Wires of tungsten and other bcc metals show a pronounced preferred orientation in which a [I101 direction is parallel to the wire axis.l The microstructure and texture of tungsten develop during warm working (working above room temperature but below the recrystallization temperature) into wire by swaging and drawing. Fibrous tungsten so produced has limited but useful room-temperature ductility, whereas annealed tungsten is brittle at room temperature. During an investigation of the effect of recovery annealing on the substructure of commerical 0.030 in. diam tungsten wire, it was found that transverse sections of the as-received material showed thin, curved grains, Fig. 1(b), instead of the expected equiaxed grains. The fact that Figs. l(a) and (b) show different sections of the same material at the same magnification is at first disconcerting, because the long dimension of the curved grains in the transverse section is several times the average fiber width in the longitudinal section. However, a longitudinal section cuts most of the curved grains across their narrow dimension, giving the appearance of a multitude of narrow fibers. On the other hand, the plane of a longitudinal section occasionally includes a long dimension of a curved grain, resulting in a broad ofibero, such as A in Fig. l(a). EXPERIMENTS AND RESULTS To investigate the origin of fibrous microstructures and particularly the unusual transverse-section structure of tungsten, rods of iron (also bcc) were drawn and swaged in the laboratory. One rod of annealed Ferrovac E was drawn in six steps to a total reduction in area of 87 pct (0.10 to 0.036 in. diam), and a second rod was swaged to the same reduction, using approximately the same reduction steps. Longitudinal and transverse sections of the swaged and drawn iron wires are shown in Fig. 2, for reductions in area of 39, 73, and 87 pct. A fibrous microstructure in longitudinal section is first evident at 73 pct reduction, and it is further developed at 87 pct reduction. The grains in transverse section remain equiaxed until about 73 pct reduction, when they become somewhat elongated and curved. At 87 pct reduction, the drawn wire shows thin, curved grains in the transverse section, similar to those observed in tungsten. The transverse sections of the iron wires swaged 83 and 87 pct reduction in area show the striking ospiral nebulao structure in Fig. 2(d). This struc-

Jan 1, 1962

By R. H. Atkinson

With the object of improving the melting, casting, and working of palladium and 95.5 pct Pd-4.5 pct Ru, the effects of different melting atmospheres (reducing, oxidizing, and neutral), crucible linings, and annealing atmospheres have been studied. Some interesting observations are included about the solubility and diffusion of gases in palladium. THE object of the present work was to ascertain the effect of oxidizing, inert, and reducing ambient atmospheres on the melting and casting of palladium and a Pd-Ru alloy commonly used for jewelry. The need for such a study was indicated during World War I1 when jewelers experienced difficulty in their attempts to make small castings of the torch-melted Pd-Ru alloy by the lost-wax process. When the metal was melted under oxidizing conditions similar to those used for platinum, the castings were gassy; on the other hand, if the melting atmosphere was reducing, the castings were hot-torn and brittle. General Experimental Procedure Pure palladium sponge (99.9 pct Pd) and pure ruthenium powder (99.8 pct Ru) were used for making small ingots (50 g) of palladium and 95.5 pct Pd-4.5 pct Ru by induction melting. The coil, 3 in. high by 2 in. ID with 16 turns, was energized by a 20 kva Ajax high frequency converter. The melting crucibles, so-called sand crucibles, which were made of a highly refractory fireclay, were lined with different materials, including proprietary magnesia, alundum, zircon, and zirconia

Jan 1, 1957

By D. A. Thomas, M. F. Comerford

Poly crystalline wires of Fe-3.2 wt pct Si were cold-drawn to 31, 66, and 87 pct reductions of area. Mechanical properties and tnicrostructures were studied after recovery and re crystallization. An initial increase in strength during recovery was ascribed to strain aging, and subsequent softening was largely due to polygonization rather than re-crystallization. No unique actioation energies could be determined for the recovery process. Rather, activation energies ranged from 60,600 cal per mole at 10 pct recovery to 83,200 cal per mole when recovery was 50 pct complete. Sites for the nucleation of recrystallization were found to depend on time and temperature. In the early stages of re-crystallization, nucleation occurred in deformation hands. Later, or at high temperatures, recrystal-lized grains were nucleated primarily al grain boundaries. Polygonized structures did not seem to be preferred sites for nucleation. MUCH of the research on deformation, recovery, and recrystallization processes in bcc materials has been performed on single crystals or lightly worked polycrystals. Little study has been devoted to the relation of mechanical properties to micro-structure in polycrystals after severe deformation and subsequent annealing. This paper presents the results of a study of the tensile properties and microstructures of Fe-Si wires drawn varying amounts and annealed at a series of temperatures. EXPERIMENTAL PROCEDURE The experiments were carried out on an alloy of the following composition: 3.2 wt pct Si, 0.08 Mn, 0.016 C, 0.005 0, 0.0002 N, 0.0001 H, 0.008 P, 0.034 S. The material was received as l-in.-sq bars cut longitudinally from a hot-rolled slab provided by the Armco Steel Corp. The bars were forged to 3/8-in. rounds, centerless-ground to 0.250 in., and swaged to '0.108 in. All subsequent reductions were performed by room-temperature wire drawing at the rate of 2 in. per min. Each preliminary draft gave approximately 30 pct reduction of area and was followed by a l-hr anneal at 1010°C in dry hydrogen. Three lots of wire were made with ultimate cold reductions of 31, 66, and 87 pct reduc-

Jan 1, 1965

By S. L. Craig, C. Orr, H. G. Blocker

WITHIN recent years gas adsorption methods have been developed for measuring the surface area of finely divided materials and have become extremely valuable in research on the corrosion and the catalytic activity of metals. Rather elaborate apparatus is required, and a single determination is so time-consuming that these methods have not been utilized to the fullest extent; the methods are un-suited for most routine control work such as that encountered in powder metallurgical operations and in processes employing metal catalysts. These difficulties are largely eliminated, and surface area is reduced to a routine determination if the liquid-phase adsorption of a surface-active agent such as a fatty acid can be used. When the affinity of the fatty acid carboxyl group for the solid surface is greater than its affinity for the solvent, a unimolec-ular layer of orientated fatty acid molecules will be formed at the solid-liquid interface in a manner similar to that of a compressed fatty acid film on a water surface. The measurement of surface area is then reduced to a measurement of fatty acid adsorption. This propitious circumstance, first investigated by Harkins and Gans,¹ has been employed with somewhat inconclusive results by a number of investigators in evaluating the surface properties of metals, metal catalysts, and metal oxides. The specific surface area values for nickel and platinum catalysts, determined from the adsorption of a number of fatty acids from various solvents, were found by Smith and Fuzek² to agree with values calculated by the gas adsorption technique of Brunauer, Emmett, and Teller," he so-called BET technique. And recently Orr and Bankston4 have also reported good agreement between nitrogen gas and stearic acid adsorption results in the measurement of the surface areas of clay materials. On the other hand, Ries, Johnson, and Melik5 found only order-of-magnitude agreement between these two methods in studying supported, cobalt catalysts having specific surface areas as great as 420 sq m per g; the reason is partially attributable to the very porous nature of the materials. Greenhill,6 investigating the adsorption of long-chain, polar compounds in organic solvents on a number of metal powders, concluded that a uni-molecular layer of stearic acid was formed on exposure of the solid to the acid solution and that the presence of an oxide or another film did not alter this result. Furthermore, the adsorption process appeared to be the same whether or not the sample was degassed prior to exposure to the solution. Greenhill estimated the surface area of one of the powders he investigated from microscopic diameter measurements, and obtained a rough check with surface area evaluation. Russell and Cochran7 found moderate agreement for alumina surface area results by fatty acid and gas adsorption methods. In addition, they also found that the prolonged heating and evacuating pretreatments previously used by investigators were unnecessary. The present work, however, considerably extends these previous investigations, shows that fatty acid adsorption can be used to determine the surface area of a variety of metals and metal compounds, offers further confirmation of the correctness of gas adsorption methods, and presents a simplified technique for the determination of the metal surface area which is suitable for routine work. Experimental Technique Basically, the fatty acid adsorption method is quite simple. It consists of exposing a sample of the material of which the surface area is desired to a fatty acid solution of known concentration. By analysis of an aliquot of the solution, the concentration after adsorption has occurred may be determined. The difference between the initial quantity of acid in solution and the final quantity is that quantity of acid adsorbed by the sample. The specific surface area of the adsorbent material may be calculated from the quantity adsorbed and the weight of the sample. In agreement with the findings of others as outlined above, it was found entirely unnecessary to degas or pretreat the nonporous materials employed other than by drying them thoroughly. However, precaution was necessary so that the dried sample entered the fatty acid solution with little exposure to moisture. The effect of moisture on the interaction of stearic acid with finely divided materials has been thoroughly investigated by Hirst and Lancaster." They found the presence of water merely reduced the amount of acid adsorbed by powders such as TiO2, SiO2, Tic, and Sic. With reactive materials such as Cu, Cu2O, CuO, Zn, and ZnO, however, water was found to initiate chemical reaction. Only with ZnO was reaction observed when the solid and the solu-

Jan 1, 1953

By J. J. Gilman, T. A. Read

S URFACE effects in the cleavage of brittle crystals have been known for some oftime,1, 2 but our knowledge of surface effects in the plastic deformation of crystals is of relatively recent origin. In 1936, Roscoe9 eported his observations comparing oxidized and clean cadmium crystals. He found that cadmium crystals 1/2 mm in diam when covered by an oxide film less than 1200 atoms thick had critical shear stresses about twice those for clean crystals. This effect is anomalous in that it cannot be caused by the inherent shear strength of a normal oxide film. Either an interaction between the film and the crystal causes the large increase in strength, or the oxide film has an abnormal mechanical strength. Several authors since Roscoe have studied surface effects in crystal plasticity. A complete review of the literature will not be presented, but most of the results may be summarized as follows: 1—The critical shear stress for slip is increased markedly by very thin oxide films on zinc4 and cadmium crystals;" ' also, the entire stress-strain curve is raised by an oxide coating;3,4 both effects increase with decreasing specimen radius. 2—Surface active substances such as oleic acid destroy the hardening effect of oxide films.4, 6-8 The time required for the reduction of the hardening effect to begin increases with the viscosity of the solvent4~ 3—There has been dispute in recent years about the effects of surface active substances themselves on plastic deformation. Rehbinder, Lichtmann, and Maslenikov9 state that surface active substances reduce the strength of metal monocrystals markedly and increase the number of slip lines per unit length at a given strain. The effect varies with orientation for tin crystals being smallest for x, = 0, largest for x, = 45", and intermediate as x, + 90°.10 Several workers claim to have disproved the results of Rehbinder')"." but Rehbinder and Lichtmann9 have defended themselves and Masing with his collaborators" has confirmed some of Rehbinder's results. 4—Immersion of metal specimens in electrolytes such as sodium chloride reduces their strength12,13' and subsequent cathodic polarization increases the strength reduction. For noble metals, the strength reduction occurs for both anodic and cathodic polarization," increasing with increasing polarizing voltage. 5—Electroplated copper coatings effect the creep rates of zinc monocrystals. Poly crystals are not affected." 6—Immersion in dilute acid destroys the hardening effect of an hydroxide film. When an acid solution is applied to a steadily creeping crystal, the creep rate jumps to a very high value and then drops gradually to a steady value which is higher than the initial steady rate.*' " If the hydroxide film to be removed is very thin (about 100A thick), the creep rate does not change measurably, but a small rapid strain-increment is observed when dilute acid is applied." 7—Crystal plasticity depends on specimen dimensions. This was observed by Lichtmann and Ven-stroml" for monocrystalline tin and by Andrade and Kennedy" for polycrystalline lead. 8—Immerson of copper crystals in mercury has no effect on their stress-strain curves.'" 9—Bombardment of creeping cadmium crystals with a particles which penetrate the surface to only 0.005 mm may increase their creep rates as much as five times.''' The experiments which will be described in this paper pertain to: (a) the effect of a crystal's shape on its plastic deformation, and (b) the effect of metallic surface films on plastic deformation. For the work pertaining to (a), crystals of zinc, tin, and lead were grown which had various shapes. These were tested in simple tension. Macroscopic rotations about the tension axis which are unpredicted by the classical theory of slip were observed, and serrated edges were observed in tin crystals. For the work pertaining to (b), the effects of evaporated and electrodeposited films of copper, nickel, gold, and zinc on the mechanical properties of zinc, tin, and lead crystals were determined. The films influenced both slip and twinning of the crystals in creep and tensile tests. Experimental Details Chemically pure zinc, tin, and lead were used in this investigation. The zinc was 99.999t pct pure;

Jan 1, 1953

By C. G. Dunn, J. L. Walter

Either (100) (001] or (110) (0011 oriented secondaries, or both, depending on annealing atmosphere and material composition, grew in a matrix of 2-dimensional grains. The growth dependence of secondaries of(hkl) orientation in the iron and 0.6 pct Si-Fe is explained in terms of surface energy driving forces and the dependence of these driving forces on both (hkl) orientation and amount of oxygen at the surace. (100) [001] grains grow when suflicient oxygen is present at the gas-metal interface to make y100 the lowest (hkl) surface energy. When little or no oxygen is present at the gas -metal interface, making Y110 the lowest surface energy, (110) (001 1 grains grow. When Y100 and y110 are approximately equal but lowest among (hkl) surface energies, both (100) (0011 and (110) (0011 grains grow simultaneously. An excess of surface oxygen produces complete grain growth inhibition in high-purity iron; i.e., a 3-dimensional grain structure is retained over long annealing periods. However, a small amount of silicon in the iron reduces the growth inhibiting tendency of oxygen without altering the (hkl) surface energy relationships. Some information on annealing twins is presented. SECONDARY recrystallization in high-purity a iron, cold rolled to strip a few mils thick, was reported recently.' The present paper provides additional information on this subject, also furnishes similar information on both zone-refined iron and a high-purity 0.6 pct Si-Fe alloy, and treats the more important surface phenomena that determine to a large extent whether one or more secondary recrystallization textures will appear during annealing. Previous studies on high-purity 3 pct Si-Fe2,3 have shown the importance of (hkl) and impurity-atom dependent gas-metal surface energies on the formation of secondary recrystallization textures in high-purity material. For example, either (100) or (110) oriented grains may have the lowest surface energy depending on the amount of oxygen at the gas-metal interface, and hence the highest driving force for growth if other growth factors are the same. Another important phenomenon in high-purity material is stabilization of the matrix structure; neither impurities nor a strong matrix texture seem to be necessary when recrystallization and normal grain growth occur, developing a stable two-dimensional type grain structure. The average grain size in such a structure is usually between one and two times the specimen thickness. Consequently, most of the grains have grain boundaries terminating at exposed surfaces, and this results in gas-metal interfaces (or hkl surfaces) as well as

Jan 1, 1962

By G. R. Speich

RECENT studies by Smith and Olney,1,2 Olney,3 Greifer and Salli,4 Rys etal., and Olney and smith 6 have established that graphite is the first decomposition product to format the surface of hypereutectoid plain-carbon steels when they are austenitized in a vacuum or inert atmosphere and slowly cooled. This graphite evidently forms directly from the austenite and only at the surface of the specimen. The interior of the specimen is found to be free of graphite. The graphite layer can almost completely cover the specimen surface and the rate of its formation is markedly sensitive to the orientation of the austenite. The present work was undertaken to study this phenomenon further, especially with regard to the effect of temperature and austenite orientation. The material studied in this work was a high-purity iron-carbon alloy containing 0.88 pct C, 0.005 pct Mn, 0.016 pct Si, 0.005 pct S, and 0.002 pct P. Specimens were heat-treated in a commercial hot stage modified for rapid quenching.7 This permitted direct observation of the formation of the surface graphite. A vacuum of 10"" to 10-5 mm of Hg prevented oxidation or decarburization of the specimens. Specimens were austenitized for 30 min at 1100°C and quenched to a series of lower temperatures where they were isothermally transformed. Surface graphite was observed to form at 700°, 750°, and 780°C but not at 825°, 900° or 1000°C. Since the maximum solubility of austenite for graphite is 0.88 pct C at 800°C,8 surface graphite only forms at tem-

Jan 1, 1962

By R. L. Dietrich

Magnesium alloy sheet has less ability to accept bending at room temperature than most of the heavier metals. In work designed to improve the bend properties, the preferred orientation of the sheet is of major importance as it is in all studies of the properties of wrought magnesium products. When rolled into sheet, all of the common magnesium alloys form an orientation texture having the basal (002) planes approaching parallel to the surface of the sheet. This texture is only slightly affected by annealing. Magnesium single crystals are highly anisotropic, and, as might be expected, so are magnesium alloy wrought products in which a strong preferred orientation is developed. It is therefore not surprising that bend properties are affected by orientation. Ansel and Betterton1 reported that the orientation of AZ3lXt sheet varies from surface to center and that bend properties are improved by etching away the sharply oriented material at the surface of the sheet to reach the more broadly oriented structure below. This paper covers a study of that orientation, either during the rolling process or by treatment of the finished sheet, in an effort to improve the bend properties and toughness of sheet. Literature The orientation texture of magnesium and magnesium alloy sheet has been studied extensively. Early determinations2 showed that pure magnesium sheet has a preferred orientation in which the basal planes are parallel to the sheet surface within very narrow limits. J. C. hIcDonald3 and J. D. Hanawalt4 reported that sheet containing a small amount of calcium develops a "double" texture, that is, the majority of the basal planes are a few degrees from parallel to the surface and there is a noticeable vacancy at the parallel position. Bakarian5 made careful quantitative pole figures of both pure magnesium sheet and MI alloy SEPTEMBER 1949 sheet which show these features. In all of these studies, however, the orientation was determined by transmission methods in which the resulting pattern is an average through the thickness of the sheet. The tendency of wrought metal to exhibit a different orientation at the surface from that in the center has been reported by many investigators. G. von Vargha and G. Wasserman6 found that with copper, aluminum, iron, and brass the textures of rolled compared to drawn wires were the same at the center but differed markedly at the surface. It was also reported by investigators7 that the orientation of rolled aluminum varies from surface to center. Har-greaves8 found that the surface texture of AM503 (magnesium alloy similar to MI) sheet was different from the center texture. It is reported by Edmunds and Fuller9 that zinc alloy sheet sometimes had a thin layer at the surface with a strong orientation of the basal planes parallel to the surface, which, if present, impaired the bend properties of the sheet. Part1 Surface Orientation ofMag- nesium Alloy Sheet and the Effect on Properties Attempts to correlate the bend properties of magnesium alloy sheet with tension ductility over short gauge lengths proved unsuccessful and the subsequent investigation showed that nonuniformity in orientation is a con- tributing factor as the properties of the surface material have a much more important effect in bending than in tension. A program to study the relationship between surface orientation at the surface and bend properties was then undertaken. First, the effect of etching away the surface of sheet on the bend properties and the orientations at the various depths were studied. Sheet samples of M1, AZ31X, and AZ61X were etched in dilute nitric acid to remove the surface material for various depths to 0.015 in. As may be seen in Table 1, the minimum bend radius improved considerably as the surface layers were etched away but it was necessary to etch the sheet quite deeply, much more so than was found necessary by Edmunds and Fuller9 on zinc sheet. It is also apparent that the amount of etching required is a function of the sheet thickness. In all of this work, radii were measured as R/t, the radius divided by the sheet thickness, in order to eliminate the effect of the reduction in sheet thickness produced by the etching. To determine the orientation texture of the sheet, X ray reflection patterns were taken using copper radiation with the bearn striking the specimen at an angle of 17' to the surface, which is the Bragg angle for the (002) planes of magnesium. Two exposures were made of each specimen, one with the beam perpendicular to the rolling direction and the other with the beam parallel to the rolling direction. The symmetry of the preferred orientation in magnesium sheet is such that these two photographs gave an approximation of the pole figure sufficiently accurate for qualitative work and it was not thought worthwhile to make complete pole figures. These X ray patterns show that the orientation has a much narrower spread at the original surface of the sheet than below the surface. The narrow spread is found in sheet having the majority of the basal planes (002) parallel to the surface, and since this is an unfavorable position for slip, it is

Jan 1, 1950

By F. H. Buttner, H. Udin, J. Wulff

Using a modified Udin, Shaler, and Wulff technique, the surface tension of gold Udin, purified helium was found to be 1400 ± 65 dynes per cm for the temperature range 1017° to 1042°C. IN the original Udin, Shaler, and Wulff technique for measuring the surface tension of copper: variously weighted wires were allowed to extend or contract in a copper cell held at elevated temperatures in vacuum. By plotting stress vs. strain for a wire array in one test, the stress at zero strain is obtained. This is the point where the contractile forces resulting from surface tension are balanced by the applied load, according to the expression: y = e=o r [1] where y is the surface tension in dynes per cm; a,,,, the stress at zero strain in dynes per cm; and T, the radius of the wire in cm. The assumption that the wires deform viscously permits the drawing of a straight line through the points on the stress-strain plot. Justification of the assumption has received further experimental support recently.'-' The presence of grain boundaries in the wires requires a correction to the original expression used." Thus: y = d4=T [l- (dl) (ar)Y1 [2] where, n/l is the number of grain boundaries per unit length, and a, the ratio of grain boundary tension to free surface tension. Alexander, Kuczynski, and Dawson in studying the creep of gold wire in vacuum were unable to obtain reproducible values of the surface tension of gold. In plotting stress vs. strain for progressively longer times, they found that the stress at zero strain drifted with time from positive stress values to negative values. Similarly, for the surface tension of silver, reproducible values were obtained only when a purified helium atmosphere was substituted.' Evidently the evaporation rate of silver in vacuum is too high at the temperatures employed to obtain solid-gas equilibrium even in a similar metal enclosure. Thus reproducibility of results is lost. Experimental Procedure The experimental procedure was much the same as that originally developed by Udin, Shaler, and Wulff with a few modifications and improvements. For greater accuracy in strain measurements, knots gave way to cut gage marks as shown in Fig. 1. These were made with a hand-driven lathe in which razor blades serv'ed as cutting tools. Also a more precise cathetometer with a screw accurate to 0.00015 cm was used. The tests were conducted in an atmosphere of purified .helium rather than in vacuum in order to avoid possible evaporation difficulties. Five mil wire of high purity gold (99.98 pct) was used. After cutting in the gage marks, each wire of a series of about 12 was differently loaded by welding a gold ball to one end. This was done by dipping the end of the wire in a cooling gold droplet, previously melted on a charcoal block with a No. 2 acetylene torch. The other end of the wire was strung through a hole in a gold lid and twisted over the edge to hold the wires fixed and in suspension from the lid. The lid and mounted wires were then dipped in pure ethyl alcohol to dissolve any skin oils and dirt on the surface of the wires due to handling. Finally the lid was put in place on an alundum crucible lined with gold so that the wires hung freely within the gold-lined chamber. This whole assembly was next heated in a quartz nichrome wound tube furnace and heated for a few minutes at 600°C to soften the wires. After this anneal the wires were easily straightened with tweezers. The wire assembly was finally annealed 10" to 25°C above the subsequent test temperature for 2 hr. This treatment allowed the grains to grow to equilibrium size and shape. After the anneal, the lid was mounted in front of the cathetometer. The gage length was measured by sighting the 40 power microscope on the upper lip of the lower gage mark for the first reading, then traveling up to the lower lip of the upper gage mark for the final reading. This procedure was repeated four times to give an average gage length value. In this manner the annealed gage length and the final gage length could be measured to determine the strains. During all measurements, grain counts were made.

Jan 1, 1952

By W. W. Mullins, F. A. Nichols

Solutions are developed, assuming surface diffusion and both internal and external volume diffusion, for the relaxation of bodies slightly perturbed from spherical and cylindrical geometries. Combined with those previously published for the nearly planar case, these results provide a means of gaging the relative contributions of the two diffusional proc-cesses in any given case. It is shown that in all sintering experiments to date, and probably in any attainable in practice, surface diffusion has played the dominant role, although most previous authors have assumed otherwise. It is also shown that surface diffusion predominates in normal field-emission tip blunting and also for the coalescence of gas bubbles introduced into metals by a bombardment. The surface-diffusion solutions for a perturbed sphere are combined with previous results for volume diffusion to show that the inclusion of interface diffusion permits considerably larger spheres to develop in diffusion-controlled precipitate growth before the onset of instability. A mechanism is also proposed for the spheroidization of precipitate platelets as well as rods. In a previous paper1 the relaxation of a nearly plane surface to flatness by the combined action of the transport processes of viscous flow, evaporation-condensation (in a closed system), volume diffusion, and surface diffusion has been analyzed under the assumption that all surface properties are independent of orientation. In this treatment, criteria were developed for deciding which process predominates, and solutions valid in the latter stages of the sintering of small wires and particles to a plane were obtained. A numerical solution, valid throughout the entire particle-sintering process for the case of surface diffusion, was subsequently obtained by the present authors.' It was found that the analytic solution (which assumed small slopes everywhere) is accurate to within -10 pct when the maximum slope of the profile is less than 0.3; the wire-sintering problem has also been solved nu- merically for the case of surface diffusion and here again the results converge to the analytic small-slope solution at late stages of the process, the two solutions agreeing in this case to within 10 pct when the maximum slope of the profile is less than -0.6. The purpose of this paper is to extend the perturbation solutions to nearly spherical and nearly cylindrical geometries. We treat only the two principal diffusional processes, i.e., surface and volume, but for these geometries we discuss volume diffusion both inside and outside of the solid. Our results, when coupled with Mullins' solutions1 for nearly planar surfaces, provide criteria for gaging the relative contributions of surface and volume diffusion to the over-all transport process in three basic geometries. A very interesting feature in the cylindrical case is the occurrence of instability for longitudinal perturbations with wavelengths greater than the cylindrical circumference, a classical result. This instability of the cylindrical surface is applied to give a quantitative explanation for the often-observed "erratic" pore closure in the late stages of the sintering of wire compacts; also, the theory previously presented for the spheroidization of rod-shaped precipitates by surface (interface) diffusion' is expanded now to include volume diffusion inside and outside of the particle. The results for circumferential perturbations on a long cylinder allow quantitative estimates for gaging the relative importance of surface or volume diffusion in the early stages of the sintering of spheres or wires. The results here demonstrate clearly that surface diffusion has played a very important, if not dominant, role in all sintering experiments discussed in the literature, although the surface-diffusion contribution to the kinetics has usually been ignored. The results for the sphere (surface-diffusion case) are added to the results obtained previously by Mullins and sekerka3 concerning instabilities of a growing spherical precipitate particle (with interface diffusion disallowed) to obtain a general solution to this problem including interface diffusion. The inclusion of interface diffusion is found to increase significantly the range of stability of a growing spherical precipitate for typical metallurgical cases. The following assumptions are made: (i) the initial surface lies everywhere near and has a slope differing only slightly from that of the reference

Jan 1, 1965

By H. I. Aaronson, C. Wells

Configurations of ferrite crystals have been found in a plain carbon steel which appear to have resulted from the nucleation of new ferrite crystals at the interphase boundaries of previously formed crystals despite the high carbon concentrations which necessarily develop at these boundaries. This phenomenon has been termed sympathetic nucleation. An attempt has been made to reconcile the occurrence of sympathetic nu-cleation with current nucleation theory. THIS investigation is one of a series on the formation of proeutectoid ferrite from austenite. From the viewpoint of chemical composition, this reaction consists of the nucleation and diffusional growth of crystals of carbon-poor ferrite within a matrix of carbon-rich austenite. The austenite adjacent to the austenite-ferrite boundaries will be greatly enriched in carbon, approximately to the value of the y/(a + y) equilibrium curve or its metastable extrapolation at the temperature of transformation. Those areas of austenite appreciably farther removed from the growing ferrite, on the other hand, will be relatively unaltered in composition, especially at the earlier stages of transformation. Since rates of nucleation are considered to decrease exponentially with decreasing supersaturation,' the frequency with which ferrite nuclei appear at austenite-ferrite boundaries should be negligible in relation to that at which they form in other regions of the austenite. During this investigation, however, many groupings of ferrite crystals have been found which appear to have resulted from the nucleation of ferrite at austenite-ferrite boundaries. This phenomenon has been given the name of sympathetic 71.1tcleation. A number of micrographs of morphological configurations caused by sympathetic nucleation will be presented, after which an explanation for this reaction will be proposed in terms of current nucleation theory. Some of the structures to be considered are composed of bainite, an aggregate of ferrite and carbide, rather than of ferrite. Since ferrite and bainite differ only in that bainite forms under conditions which result in the nucleation of carbides behind the advancing austenite-ferrite boundaries,' it will usually be unnecessary, for the purpose of this paper, to distinguish between the two reaction products. All studies were performed on an electric furnace steel (obtained from the Vanadium Alloy Steel Co.) containing 0.29 pct C, 0.76 pct Mn, 0.25 pct Si, 0.005 pct P, and 0.007 pct S. The alloy was cast as a 150 Ib, 7x7 in. cross section ingot and forged into bars 2x2 in. in cross section. These bars were homogenized for 48 hr at 1250°C in an Endo-Gas atmosphere. The depth to which decarburization penetrated during this heat treatment was determined by chemical and microscopic analyses and the affected metal was removed by machining. Specimens for isothermal transformation studies were cut from the remaining material; most of these specimens were 1/2x1/4X1/16 in., though some with a thickness of 1/32 in. were prepared for use at the shorter reaction times and lower reaction temperatures. Specimens were austenitized for 30 min at 1300°C, isothermally reacted for various times at temperatures ranging from 775" to 475 "C, and then quenched in iced water. The austenite grain sizes within individual specimens ranged from ASTM Nos. 1 through —4. A commercial heat-treating salt which was continuously deoxidized by an immersed graphite crucible served to minimize the loss of carbon during austenitizing; thick covers of powdered graphite and immersed graphite rods effectively prevented decarburization in the lead pots employed for the isothermal reaction treatments. The heat-treated specimens were sectioned and mounted in Bakelite. Following the completion of standard grinding and mechanical polishing procedures, the specimens were electrolytically polished with a Buehler-Waisman apparatus and etched in 2 pct nital. Experimental Results Rules of Evidence for Sympathetic Nucleation—On the basis of observations made on a single plane of polish, one precipitate crystal may be considered to have been sympathetically nucleated at the inter-phase boundary of another precipitate crystal when the following conditions are fulfilled: 1) The sympathetically nucleated crystal is not in contact with a grain boundary or a subboundary in the matrix phase. 2) The shape, size, and location of the crystal at whose boundary sympathetic nucleation occurred (hereafter termed the base crystal) and the crystal formed by sympathetic nucleation substantially pre-clude the possibility that the plane of polish em-ployed may have concealed the fact that both crys-tals actually nucleated at a grain boundary or a sub-boundary in the matrix phase.

Jan 1, 1957

By B. Post, F. W. Glaser

Three borides of zirconium have been reported: ZrB,l ZrB2,2,3 and ZrB12, 4 The phase relationships, ranges of stability, and some physical properties of these compounds are described. THE zirconium borides differ considerably from one another, particularly with respect to their relative stabilities at high temperatures in the presence of carbon.4,5 ZrB2 is extremelv stable at all temperatures (to the melting point) in contact with carbon. Both ZrB and ZrB12 tend to decompose at elevated temperatures when carbon is present in excess of about 0.5 pct. ZrB forms ZrB2 and ZrC, and ZrB12 decomposes to form ZrB2 and a boron-carbide phase. In the ZrB12 preparation, boron present in excess of that needed to form ZrB2 reacts with carbon, if present, to form carbides, in preference to forming the higher boride. Similar instability of borides, other than diborides of the transition metals of the 4th and 5th groups of the periodic table, has been reported.5,6 special precautions had, therefore, to be taken to minimize carbon contamination in the preparation of ZrB and ZrB12. Reactions below 1400°C were carried out by heating powdered mixtures of zirconium hydride and boron in quartz tubes in a hydrogen atmosphere. For temperatures above 1400°C, reactions were carried out in graphite molds. Carbon contamination from the molds was minimized by employing extremely rapid heating rates to reduce carbon diffusion into the body of the specimen, and by subsequent "shaving" of sample surfaces that had been in contact with the graphite dies. A heating rate of about 2000°C per min was obtained by heating the dies (by direct conduction), using a current of 40,000 amp. Although this type of heating cycle apparently allowed little time for "homogenization," close approximations to equilibrium conditions were usually obtained as a result of the very strongly exothermic reactions between zirconium hydride and boron at elevated temperatures. Reproducibility of results in repeated runs under varying conditions was taken as a reasonable indication that equilibrium conditions had been effectively achieved. Quenching was accomplished by rapidly immersing the entire die, containing the reaction products, in water. Specimens could also be cooled at controlled rates by adjustment of the rate of flow of water through the water-cooled copper electrodes in contact with the graphite dies.

Jan 1, 1954

By R. F. Domagala, M. Hansen, D. J. McPherson

Iodide zirconium was combined with calculated amounts of nitrided zirconium sponge and arc melted to prepare alloys in the 0 to 6 wt pct N region. Annealing treatments were carried out at 21 temperature levels. Metallographic examination of the heat-treated specimens permitted construction of the binary phase diagram from 0 to 6 pct N. Features of the diagram include the peritectic formation of both a and ß solid solutions. The maximum solubility of nitrogen is 0.8 pct in 8 zirconium and 4.8 pct in a zirconium. An X-ray study of nitrided materials was made in the range 6 to 13 wt pct N region because serious nitrogen losses were experienced when attempts were made to arc melt these high nitrogen alloys. PHASE relationships in the Zr-N system have been determined. Due to the inability to retain nitrogen in nitrogen-rich alloys subjected to arc melting, definitive work was possible only in the range 0 to 6 wt pct (0 to 30 atomic pct) N. Sufficient complementary X-ray work was done to permit construction of the binary phase diagram up to 13.3 wt pct (50 atomic pct) N (ZrN). Small ingots were prepared by arc melting master alloys with enough zirconium to produce an alloy of the desired composition. Following pretreatment, alloys were annealed at temperature levels between 600" and 2020°C. Determination of the phase boundaries was then accomplished by metallographic evaluation of specimens quenched from the various temperatures. Incipient melting techniques were used to corroborate solidus curves, and X-ray diffraction was employed to study the nitrogen-rich region of the a + ZrN phase field. Materials Westinghouse Grade 1 iodide zirconium crystal bar served as the base material for this investigation. The as-received bars were lightly sandblasted, pickled in a 20 pct HN0-5 pct HF aqueous solution, rinsed in water and acetone, and dried. They were rolled to about 1/32 in. strip and acid-pickled, followed by water and then acetone rinses. The material was sheared to approximately 1/4 in. squares, cleaned with acetone, and stored for use. Pure zirconium nitride is not commercially available, according to correspondence with possible suppliers. A literature survey indicated that nitrogen may be introduced into zirconium metal by passing nitrogen or ammonia gas over zirconium at an elevated temperature. After considerable experimentation, a train was devised whereby high quality nitrogen was passed through a series of bubblers to remove the last traces of oxygen, through an H,SO, bubbler and a cold trap to remove moisture, and finally over zirconium within a resistance furnace. Hand-picked Bureau of Mines magnesium-reduced zirconium sponge proved more amenable to nitriding than the crystal bar. The nitrogen used was extra pure gas purchased from Linde Air Products Co. The first two bubblers through which the gas was passed contained a solution suggested by L. F. Fieser" as being ideal for removing the last traces of oxygen from nitrogen. The solution contains 20 g KOH, 2 g sodium-anthraquinone ,9-sulphonate, and 15 g NaHSO, dissolved in 100 ml water. It is blood-red and turns brown when contaminated with oxygen. A saturated lead acetate solution, next in the series of bubblers, removed any H2S which might have formed in the 0, removal stage. An H,SO, bubbler and cold trap removed moisture before the nitrogen was passed over the zirconium. The zirconium was contained in a stainless steel screen within a Vycor or quartz tube in a resistance furnace. Unreacted nitrogen passed out of the system through a mercury bubbler. A nitriding run was performed in the following manner: Zirconium was placed in the screen within the furnace tube. The unit was assembled, suitably clamped off, and evacuated to remove the air and moisture within the tube and that trapped by the zirconium sponge. The unit was flushed with nitrogen and evacuated twice. After nitrogen was allowed to pass over the zirconium for about 30 min, the furnace was turned on; several hours were required to reach temperature. The reaction was allowed to continue at temperature for about 4 hr and the screen was then withdrawn from the hot zone of the furnace by means of a Nichrome or Kanthal rod. The screen and its contents were allowed to cool at the end of the tube and were then removed. At 800°C no more than about 5 pct N could be introduced into the sponge zirconium. This material was designated M-1 (master alloy No. 1) and reserved for the preparation of alloys in the dilute nitrogen region. A second set of experiments was conducted at 1000°C; no more than 7.2 pct N could be introduced into the zirconium at this temperature. Enough material (M-2) of this nitrogen content was prepared to produce a second set of alloys to complement the first and to provide alloys in the 0 to 6 wt pct N range. It was hoped that a diffusion anneal plus a re-nitriding might yield a material of higher nitrogen content. The material containing about 7 pct N was therefore given a homogenization treatment at 1000°C for 6 hr under an argon atmosphere. It was

Jan 1, 1957

By C. A. Siebert, S. H. Bush

Isothermal temper-embrittlement studies were conducted on a 5140 steel at various temperatures for times as long as 3000 hr. Specimens from the embrittled steel were subjected to impact tests, metallographic examination, microhardness, and lattice-parameter measurements. SEVERAL of the steels in the low alloy group are subject to a pronounced decrease in impact energy if held within, or cooled slowly through, a range of temperatures below the lower critical. This phenomenon, known as temper embrittlement, is quite pronounced in the chromium steels of which AISI 5140 is typical. This study considered the response of one heat of 5140 steel when embrittled isothermally at various temperatures for times as long as 3000 hr. Excellent reviews of the work done in the field of temper embrittlement have been written by Hollomonl and Woodfine.' Both authors list the various theories which have been advanced to explain temper embrittlement. Some of the mechanisms considered were: 1—a transformation below 700°C; 2—an allotropic modification of iron; 3—precipitation from the ferrite of such compounds as carbides, nitrides, phosphides, or chromium oxides; 4—decomposition of retained austenite; 5—modification of special carbides; 6—grain-boundary segregation; and 7—selective distribution of carbides. The present study investigated the changes in hardness, Meyer's number, impact properties, micro-structure, and lattice parameter which occurred on embrittlement. Previous studies considered the influence of microstructure and of prior austenitic grain size3 on the response of 5140 steel to temper embrittlement. Equipment and Procedure All samples were taken from one heat of 5140 steel; the chemical composition of this steel was: 0.43 pct C, 0.86 pct Mn, 0.020 pct P, 0.019 pct S, 0.27 pct Si, 0.84 pct Cr, 0.10 pct Ni, 0.06 pct Mo, 0.10 pct Cu, and the remainder Fe. The material was received as % in. rounds. This stock was cut down to blanks measuring approximately 0.425~0.425~2.196 in. Heat Treatment: The rough blanks were normalized at 1600°F (870°C) for 1 hr, then austenitized at 1600°F (870°C) for 1 1/2 hr, and quenched into an agitated oil bath. The specimens were tempered at 1275°F (690°C) for 5 hr and water quenched. These tempered specimens were designated as the toughened state. Embrittling Treatments: The tempered bars were divided into sets and subjected to the various embrittling treatments listed in Table I. Temperatures from 750" to 1050°F were investigated for times of 1 to 3000 hr. Special tests to determine the effect of heating a short time at a higher temperature after long-time embrittlement at some lower temperature were also conducted. All specimens were water quenched at the end of the embrittling cycle, then ground and notched to obtain standard Charpy V-notch specimens. Impact Tests: The bars were fractured in a 260 ft-lb Charpy impact unit; specimens were fractured at temperatures ranging from —315°F to room temperature. Fixed temperatures used consisted of —315°F with liquid nitrogen, —115°F with dry ice and acetone, and 32°F with ice and water. Intermediate temperatures were obtained by suspending the bars in a jig over the liquid nitrogen, dry ice and acetone, or ice and water.

Jan 1, 1955

By R. A. Bosch, F. V. Lenel, G. S. Ansell

RECENTLY, Ansell and aenel' proposed a dislocation model to account for the yielding behavior of dispersion-strengthened alloys. The criterion for yielding used in this model was that yielding occurs when the shear stress due to arrays of piled-up dislocations fractures or plastically deforms the dispersed second-phase particles. Calculations based on this model predict that the yield strength, uysof a dispersion-strengthened alloy containing particles whose geometry permits them to be considered as straight, i.e. not curved, barriers to dislocations follows the relation where and are the shear moduli of the matrix and dispersed phase respectively, b is the Burger vector of a dislocation in the matrix, A is the mean free path between dispersed phase particles and C is a constant. This yield stress, ys, is closely approximated in polycrystalline alloys by the proportional limit. The offset yield stress, uoy, is the yield stress, uys, plus the increase of stress due to strain hardening in the offset strain increment. Previous investigation1 has shown that Eq. [I] describes the yielding behavior of several different dispersion-strengthened alloys as a function of the mean free path between dispersed phase particles at a constant temperature. The purpose of this investigation was to investigate the temperature dependence of the yielding behavior of dispersion-strengthened alloys. The alloys studied, MD2100 and MD5100, are SAP-type alloys containing flake shaped aluminum oxide particles dispersed in a matrix of commercial purity aluminum. The proportional limit and 0.2 pct offset stress of both of these alloys were determined as a function of temperature over the temperature range from -190" to 500 C. All tests were conducted on an Instron tensile testing machine. Stress was measured by means of the Instron load cell. Strain was measured by means of SR-4 strain gages attached to the tensile specimens. Stress sensitivity was 80 psi. Strain sensitivity was 2 x 10F Fig. 1 shows the values received for the proportional limit and 0.2 pct offset yield stress plotted as a function of homologous temperature for both the MD2100 and MD5100 alloys. Inspection of Eq. [I] shows that of the terms which determine the theoretically predicted yield stress, uy,, only the shear moduli have an appreciable temperature dependence. From Eq. [I] the yield stress at any temperature T, UT, may be predicted from the yield stress at any arbitrary reference temperature, in this case 25 OC, u25, by the relation where the subscripts T and 25 refer to the values of the property at test temperature and at 25"C, respectively. The temperature dependence of the yield stress predicted by Eq. 121, based on the observed value for the proportional limit at 25OC, is also shown in Fig. 1. In evaluating Eq. [2] the following data was used. The values for the shear modulus of

Jan 1, 1962

By R. P. Carreker

GERMANIUM is a member of that group of ele-nents—carbon, silicon, germanium, and tin-— that are currently of particular interest because of their interesting electrical properties. Near room temperature these materials of diamond cubic crystal structure are characteristically brittle. However, there have been several recent reports that germanium is not brittle at elevated temperatures These observations may be summarized as follows: 1—Germanium single crystals were observed to creep in bending under a maximum fiber stress of 6 kg per mm&apos; (8500 psi) at 500°C and above.&apos; The creep curves were sigmoidal, with an induction period that decreased with increasing temperature. 2—Germanium crystal bars have been bent 90" about radii ten times the bar thickness at 550°C.&apos; 3—-Audible clicks were heard when l/4 in. thick crystals were bent drastically at 450" to 500°C. No microstructural evidence of mechanical twinning was observed 4—Slip lines have been observed in tensile specimens deformed at 600°C, and are coincident with (111) planes, the close-packed planes of the diamond cubic structure.] The slip direction has not been established. 5—A germanium crystal has been compressed approximately 20 pct along a <110> axis at 700°C." The results reported subsequently were obtained in exploratory experiments intended to indicate some aspects of the behavior of germanium single crystals when tested in simple tension. Specimen Preparation A single crystal bar of n-type germanium was used. It was of purity corresponding to a resistance range of 2 to 20 ohm-cm. Tensile specimens were prepared by cutting small rectangular bars

Jan 1, 1957

By R. P. Carreker, R. W. Guard

True stress-true strain data are given for nominally pure molybdenum (99.95 pct) over the temperature range -196° to 1540°C (0.027 to 0.63 T/T). Strain rate sensitivity was determined by rate change tests and stress relaxation tests. Inhomogeneous yielding and strain aging effects were observed. The yield stress and tensile strength depend markedly on temperature below 400°C and are insensitive to temperature in the range 400&apos; to 800°C. The ductile to brittle transition range is +25O to -25OC. Strain rate sensitivity and stress relaxation effects are very large near room temperature. THIS report presents data which were obtained as a part of a continuing program designed to document the deformation behavior of nominally pure metals over a wide temperature range.&apos;" This research program is rooted in the conviction that sufficient tensile and creep data of appropriate quality and detail does not exist to provide precise phe-nomenological descriptions of the deformation behavior of pure metals as a function of temperature, strain rate, and grain size. Such knowledge is needed to provide a firm base for alloy design and an experimental framework to which theories of deformation must conform. These data also complement a parallel exploratory program investigating the properties of refractory metals, including molybdenum, and their alloys. Several important reviews and research reports dealing with the tensile properties of molybdenum as a function of temperature have appeared."-lo To a great extent the work to be reported in the present paper parallels that of Bechtold and of Pugh. The authors&apos; results are in general agreement with those of Bechtold and Pugh. The principle differences be- tween the present experiments and those published previously are as follows: 1—Wire specimens prepared from arc cast molybdenum were used in the present experiments, bar specimens were used by others. Pugh used arc cast material while Bechtold&apos;s most extensive results" were obtained on powder metallurgy specimens. 2—Several different grain sizes were tested in the present experiments. 3—The testing range was extended to cover from -196°C to 1570°C (0.027 to 0.63 T/Tm). 4—The combination of wire specimens, high speed autographic recording, and a stiff testing machine permitted more detailed study of initial yielding. The principle contribution of the present paper is the demonstration of the relatively consistent mechanical properties of nominally pure arc cast molybdenum specimens of differing metallurgical history. Bechtold and Scotts have shown that arc cast and powder metallurgy molybdenum samples have comparable tensile properties in the range —75" to 200°C when processed to have the same re-crystallized grain size. Material The molybdenum used in this investigation was supplied by the Climax Molybdenum CO. Its metallurgical history is outlined in Table I; the final products were 0,030 in, diam wire specimens annealed at 1100&apos; and 1375°C and 0.060 in. diam wire

Jan 1, 1957

By John Nunes, Frank R. Larson

Temperature-dependent functions of various ten-sile flour stress and fracture parameters were investigated on iron and low composition alloys of Fe-C, Fe-Cr, Fe-Mn, and Fe-Ni. Data were obtained over a range of test temperatures from 200" to -196"C at an initial strain rate of 0.01 min". It was found that the strain-hardening exponent, n, for the Fe-Ni alloys increased significantly at very low testing temperatures with increased alloy addition. Also, the stress and strain at fracture followed a predictable transitional type of behavior. The flour stress-temperature dependence was evaluated employing the relation: M where it is shown that composition can strongly influence the constant M. An empirical formula describing this effect is as follows: M = MI (equivalent at. pct Ni)-&apos;" Some preliminay, tests were run on several commercial steels to test the generality of this relationship. One of the phenomenological explanations of the ductile-to-brittle transition encountered in bcc metals is based upon the difference in the temperature dependence of the flow and fracture stresses.&apos; It states simply that the fracture stress is relatively insensitive to temperature and that the yield or flow stress rises rapidly with decreasing temperature so that, when the flow stress reaches the fracture stress, brittle fracture occurs without appreciable plastic deformation. This behavior is illustrated schematically in Fig. 1 where the ductile-to-brittle transition is shown to occur at two possible temperatures. The lowest temperature, TB, represents essentially completely brittle behavior while the higher temperature, TN, represents "notch" brittle behavior. Support for this theory is obtained when it is observed that fcc metals, which do not commonly have a ductile-to-brittle transition, also do not commonly exhibit any strong temperature de- pendence of the yield stress. This relationship of yield or flow stress behavior with the ductile-to-brittle transition has been generally observed for most metals representative of these two crystallo-graphic structures. One exception to this generalization appears in the case of tantalum where no apparent ductile-to-brittle transition has been observed. However, this may simply be an indication that this type of behavior is not directly related to the flow stress and that some other mechanism such as twinning must also be operative. A comparison of the typical temperature dependence of the yield stress of some representative metals is shown in Fig. 2. In the study of the temperature dependence of the strength of metallic materials, considerable ef-

Jan 1, 1963