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By Allan T. Gwathmey, J. Bruce Wagner, Kenneth R. Lawless

Between 250°and 550°C in oxygen pressures of 10 to 760 mm Hg, the relative oxide thicknesses formed per unit time on the (100), (111), (110), and (320), decreased in this order. The predominant oxide orientation was (001)Fe3O4 and [110] Fe304 A small amount of a-Fe2O3 which exhibited a fiber pattern was observed in all oxides except those formed at 250°c in low oxygen pressures. IN order to understand the properties and behavior of thin films, particularly their role in oxidation reactions, there is needed more experimental information on their structure, composition, and kinetics of formation on clean surfaces of known geometry. The present study was undertaken to obtain such information for thin oxide films formed on single crystals of a body centered cubic metal, iron. Crystals in the form of spheres were used in order to expose all possible crystal planes to the reactant gases simultaneously. Using this procedure the relative rates of oxidation of all planes could be estimated, and those exhibiting, maximum or minimum oxidation rates could be selected for a more intensive study of composition and expitaxy. Since the three types of information were obtained from the same experiment using a single crystal substrate, any differences between rate, composition, and expitaxy on the several planes may be attributed to the influence of the substrate structure. EXPERIMENTAL The starting material for this study was Armco iron rods, 6 in. long and 3/8 or 1/2 in. in diam. The analysis was reported to be 99.8 pct Fe, with major impurities listed as carbon 0.018 pct, manganese 0.027 pct, phosphorous 0.005 pct, sulfur 0.029 pct, silicon 0.005 pct, and copper 0.11 pct. These rods were cleaned with ether and heated in an atmosphere of moist hydrogen at 950°C for three days. The purpose of this treatment was to decarburize the iron and to provide a consistent, uniform grain size for the growth of single crystals. Crystals in the form of rods, up to 1/2 in. in diam and 2 in. long, were prepared by a method similar to that reported by Leidheiser and Buck.2 Spherical crystals, 3/8 and 1/2 in. in diam and having a handle about 1/4 in. in diam by 112 in. long, were machined from the single-crystal rods. After machining, each crystal was etched in aqueous nitric acid to remove the cold-worked metal at the surface. Before each experiment, the crystal was polished mechanically with metallographic papers through number 4/0, and electropolished in a perchloric acid-acetic anhydride mixture by the method of Jacquet and Rocquet.3 The crystal was washed in running distilled water for at least 5 min. The excess water was blotted with a soft paper tissue, and the crystal was placed in an all-glass reaction vessel (See Fig. 1). The crystal was heated at 550°C in hydrogen for at least 8 hr prior to the oxidation in order to reduce any surface oxides and to help relieve any stresses in the metal. For the structure and composition studies, flat faces were cut parallel to (loo), (111), and (110) planes on the iron crystal spheres. The orientation of these faces was checked by a back-reflection X-ray diffraction technique. The faces were within 2 deg of the desired orientation. The flat faces were treated as described above except that prior to electropolishing they were given an additional polish on metallographic felt saturated with twenty-minute levigated alumina. After the crystal was heated in hydrogen for at least 8 hr, the temperature was adjusted to the desired value, and the hydrogen was evacuated. Commercial oxygen which had been purified by successively passing through tubes containing magnesium perchlorate, Ascarite and magnesium perchlorate was admitted to the desired pressure as measured by a mercury manometer. A resistance-wound furnace surrounded the oxidation chamber, and a temperature controller maintained the temperature within ±3°C. The furnace was provided with a small window for observing the crystal. The changes in thickness of the oxide films were followed by observing the interference colors. The oxidation was allowed to proceed until the desired oxide thickness was attained on a given crystal plane. The experiment was terminated by evacuating the oxidation

Jan 1, 1962

By A. Bar-Or, G. Hirsch, P. Wynblatt

The rates of nucleation and growth at various temperatures in U-0.45, 1.1, and 1.35 at. pet Cr alloys were determined, providing a qliantitative proo-f that the transformation under study is controlled by nucleation and diffusional growth. The linear relationship, log NV US 1/T2, suggests that the nucleation of the alloys is probably controlled by the classical law of nuclei formation. The established linear relationship, log G us 1/T, in the three alloys suggests that growth in each alloy is controlled by bidimensional nuclei formation. It was established that both the rate of growth and the rale of nucleation are retarded by chromium. Using classical formulae, sApN and sApG (the ß/a apparent interphase energies) were calculated from the nucleation and growth results, respectively. It was found that their values are of the same order of magnitude. However, why sApN is unaffected by the addition of chromium, while sApG increases with it, needs further clarification. In the last few years, an appreciable amount of research has been conducted into the mechanism of the ß —a transformation in uranium and uranium alloys. Recently, Burke1 reviewed the published work concerned with this transformation. He concluded that in dilute uranium alloys it is firmly established that, at high temperatures in the a range, ß decomposes by a diffusional process. However, he based his conclusion either on experiments dealing with the over-all transformation,2 or on qualitative results.3 It was felt that a better insight into the mechanism of transformation could be gained from the quantitative measurements of growth and nucleation rates, which we aimed to carry out in the present work. In a previous work,4 we have shown that a rim of a grains is formed during isothermal transformation at the high a range. The thickening of this rim with time was used for measuring the rate of growth in U-0.45 at. pet Cr.5 This method was also used in the present work for two more alloys. For meas- uring the rate of nucleation, a method of quenching and macroetching was developed to allow differentiation between a grains formed during isothermal transformation at the upper a range and those formed after quenching, Fig. 1. EXPERIMENTAL WORK AND RESULTS The Rate of Nucleation. The measurement of the rate of nucleation was effected by counting the num-

Jan 1, 1965

By Klaus Detert, Joseph Ziebs

The progress of primary recrystallization has been studied in alloys of nickel with 0.01, 0.1, and 1.0 at. pct Ta. It can be shown that in one particular alloy containing 0.1 Ta and 0.02 02 a large grain structure (0.5 mm diameter) could he developed by primary recrystallization in thin sheetswhich were cold-worked to 97.5 pct reduction in thickness. Electron transmission microscopy demmstrates that this phenonmena was linked to the occurrence of fine precipitates at subboundaries. A discussion of how fine precipitates can exert under certain circumstances a rather distinct retarding effect on nucleation rate is included. IN a former investigation, the recrystallization and texture formation in a series of binary nickel alloys including pure nickel were studied.' Nickel alloys with some tantalum added displayed peculiarities during primary recrystallization which have been studied in more detail by the present investigation. The alloys were prepared by powder-metallurgy techniques similar to the previous investigations.' Pure metal powders were thoroughly mixed to obtain alloys with 0.01, 0.1, and 1.0 at. pct Ta, referred to as alloy A, B, or C, respectively. The metallic impurities were kept below 0.005 at. pct except iron and cobalt with slightly higher values, by using highly purified nickel and tantalum powders.' The mixtures were pressed into bars and sintered at a temperature of 1300°C. The sintering atmosphere was either high vacuum, dry hydrogen, or wet hydrogen. The sintered rectangular bars were rolled 60 pct to a thickness of 3 mm, then annealed under dry hydrogen at 1000°C and again cold-rolled to about 97.5 pct reduction of thickness. After this heavy deformation, the sheet thickness was 0.07 5 mm. In order to study primary recrystallization which occurred in measurable times between 400" and 600°C, samples of the rolled sheets were annealed in a salt bath at various temperatures for various times in order to study the progress of recrystallization by hardness measurements and microscopic examination. For optical observation, the samples were polished and etched in a solution of 65 cu cm HNO3, and 18 cu cm acetic acid in 20 to 100 cu cm H2O, depending of the particular sample and annealing treatment. For transmission electron microscopic observations, thin foils were prepared by the combined window and Bollmann techniques.273 The electrolyte consisted of 23 pct perchloric acid and 77 pct acetic acid held at 15°C. Polishing was performed at a current density of 0.4 amp per sq cm. The examination of the thin films was accomplished by a Siemens Elmiscope-type electron microscope operating with an accelerating voltage of 100 kv. RESULTS The progress of primary recrystallization was very accurately followed in these alloys in order to establish the time when recrystallization started and ended during an isothermal anneal. Table I gives some of the data obtained by these studies. The time for beginning is a mean value, where somewhat less than 5 pct of the matrix has recrys-tallized. The end of recrystallization is the time when at least 95 pct of the matrix has recrystal-lized. Listed are three different B alloys, with 0.1 Ta distinguished by different sintering atmospheres which influence the rate of primary recrystallization. By the above method, the state of the grain structure when primary recrystallization was completed, but grain coarsening had not started, could be clearly defined. The average grain diameter of the recrystallized matrix is also given in Table I. The grain size did not depend on the temperature of the isothermal anneal. With one exception the grain size of the primary recrystallized matrix was found to be rather fine

Jan 1, 1965

By F. J. Plecity, W. C. Leslie, F. W. Aul

During isothermal recrystallization, the rate of grain growth in dilute Fe-hlo alloys decreased rapidly with increasing Mo content, up to -0.04 at pct, and less rapidly therea-fter. Rate of growth also decreased sharply with time at constant temperature. The observed rates of boundary migration did not cor,espond to predictions of the Lucke-Detert theory. It is proposed that the growth of recrystallized grains is inhibited by clustering of solute atoms at inzperfections in the unrecrystallized matrix. It is now firmly etablished" that the rate of re-crystallization of cold-worked metals, i.e., the rate of growth of recrystallized grains into the deformed matrix, is greatly inhibited by the presence of alloying elements in solid solution. The effect is greatest with the first very small quantity of alloying elements, and additional increments cause much less pronounced retardation of growth. The early papers on this subject were reviewed in Ref. 6. The mechanism of retardation of growth is not clear; only one theory, due to Lucke and Detert? has been advanced. This assumes an elastic interaction between grain boundaries and solute atoms which tends to increase the concentration of the latter in the grain boundaries. Except at very low solute concentrations or at high temperatures, the mobility of the boundary is assumed to be controlled by the rate of diffusion of the accompanying solute atoms. Abrahamson and Blakeney2 have noted a correlation between "recrystallization temperature of binary alloys of iron and the number of d-shell electrons in the solute atoms. There is need for considerably more experimental data on the effect of solute elements on the rate of migration of grain boundaries into the deformed matrix, to provide the basis for an understanding of the annealing process in metallic solid solutions. Only a start has been made in this direction. To avoid unnecessary complications, alloys in which precipitation does not occur should be used, but to produce the maximum effects it is desirable to select alloys with a large difference between the sizes of the atoms of the solvent and solute. For these reasons, alloys of iron and molybdenum sen for this study. MATERIALS AND PROCEDURES A 30-lb. ingot of iron and four similar ingots of Fe-Mo alloys were vacuum melted in magnesia crucibles and poured into cast-iron molds. The molybdenum contents were planned according to the results of Abrahamson and Blakeney2 sothattwoalloyswould fall in the region of rapid change of recrystallization kinetics with change in composition (0.006 to 0.015 and 0.030 to 0.045 at. pct Mo) and two alloys would fall in the region of gradual change (0.18 to 0.24 and

Jan 1, 1962

By R. E. Cech

Metallographic observations on hydrogen-reduced nickel oxide crystals suggest that nucleation of nickel occurs at structural singularities in the oxide. The fully reduced structure contains micron-diameter, worm-like voids occupying 40 to 50 pct of the volume. The crystallographic orientation relationships between nickel oxide and nickel resulting from oxide reduction are recorded and discussed. THE process of reduction of oxide to metal is one of considerable complexity involving a chemical reduction reaction, mass transport of reducing agent to the point of reaction and of reaction product away from it, and finally a transformation whereby metal atoms released from the oxide rearrange to form a stable reduced phase. In addition, nucleation of the reduced phase affects the overall reaction kinetics. Thus far only the thermodynamic aspect of reduction has received the attention warranted by the importance of

Jan 1, 1960

By J. L. Lytton

The flow-stress recovery of high-purity aluminurn following a 10 pct tensile prestrain was studied in terms of a fractional flow-stress recovery parameter fr. The flow-stress recovery behavior was related to the dislocation arrangements using transmissiort electron rnicvoscopy. During recovery at 120° and 160°C, nearly stable strength levels were attained, the strength level for 120°C recovery being greater than that for 160°C. This stabilization of flow stress was related to the formation of meta-stable dislocation networks bounding nearly perfect cells. Both network formation and the annealing of dislocation loops were found to be recovery mechanisms, and the stronger metastable strength level joy 120°C recovery was related to a smaller average cell size. Analysis of the networks was consistent with the reaction of sets of dislocations with a/2 (110) Burgers vector to give product Burgers DURING isothermal annealing of cold-worked metals, a degree of softening occurs without re-crystallization, and the processes responsible for this change are termed recovery processes. During recovery annealing, the dislocations generated during cold work become rearranged into low-energy configurations, and as much as 50 pct of the stored energy is released."' The relationship between this rearrangement of dislocations and mechanical behavior is of considerable importance, since the dislocation substructures formed during recovery provide a significant degree of low-temperature strength3j4 and can favorably alter the course of primary creep at high temperatures.4 The purpose of the present study was to relate the flow-stress recovery behavior of high-purity aluminum to the substructural modifications that occur. Electron transmission microscopy was therefore employed for direct observation of dislocation substructures at various stages of recovery following 10 pct prestrain. Several investigators5-' have proposed that dislocations can become rearranged or lost during the preparation of thin foils. It was found early in the present study that the dislocation structures in the thicker regions of the foils changed in a systematic way during the course of recovery. Therefore, it is unlikely that substantial losses or rearrangements occurred during thinning. This point will be discussed in more detail later. The results of various investigators indicate that recovery processes in fcc metals can result in release of 50 pct or more of the total energy stored in the as-prestrained metal, this percentage decreasing with increasing prestrain.' Alloying elements or impurities tend to increase this percentage. The release of stored energy of fcc metals during recovery appears to occur by 1) annealing of point defects if the deformation occurred at sufficiently low temperaturesg and 2) rearrangement of dislocations into low-energy Configurations.1,2In bcc metals, the annealing of dislocation loops has also been observed during recovery after prestrain,10-12 and it has been proposed that this is a recovery mechanism for unidirectionally strained fcc metals.' The density of dislocations in bcc metals has not been observed to decrease appreciably during recovery,10 although the formation of regular dislocation networks has been observed. has shown that the formation of such networks can result in elimination of long-range stresses and can lead to significant release of stored energy without dislocation annihilation. It might thus be expected that network formation is a significant recovery mechanism in fcc metals, although extensive network formation has apparently not been reported in controlled recovery experiments. The formation of sub-boundaries in aluminum has been observed by various investigators13'17 during recovery studies of aluminum, but the structure of these boundaries was not resolved. In this regard, it is significant to note that Bailey and Hirsch,' using transmission electron microscopy, detected no network or sub-boundary formation during recovery of 99.99 pct Ag. Furthermore, Clarebrough et al.2 found no such rearrangements during recovery of copper. EXPERIMENTAL PROCEDURE The high-purity aluminum used in this investigation was obtained in the form of 0.125-in.-thick sheet. Spectrographic analysis revealed that the sheet was 99.995 pct A1 with the following weight-percent impurities: 0.001 Cu, 0.002 Fe, 0.001 Si, and 0.0004 Mg. The as-received sheet was in the heavily cold-rolled condition. The as-received sheet was cold-rolled at room temperature to 0.040-in.-thick sheet in the longitudinal direction, which was the original rolling

Jan 1, 1965

By W. J. Babyak, F. N. Rhines

It has long been supposed that, regardless of the mechanism by which grain boundaries increase the hardness of a ductile metal, the effect should be linearly proportional to the grain boundary area. At the present time two methods for determining grain boundary area in an isotropic aggregate are available. The first, derived by Angus and Summers,' presumes that the grain boundary area is equal to 3vn, where n is the number of grains per sq mm in a random section through the metal. This estimate is based upon the assumption that the grains are equivalent to cubes whose faces correspond with the grains as observed in the micro-section. Measurements on copper, bronze, and a brass showed that a nearly linear relation exists between 3vn and hardness. This is illustrated, for the case of brass, in Fig. 1. The second method, developed independently by Smith and Guttman2 and by Duff in, Meussner, and Rhines,3 is absolute and employs a simple linear intercept count for measuring grain boundary area. It involves no assumptions with respect to the shape and distribution of the grain boundaries. The total surface in unit volume is 2 NL , where NL is the number of intercepts between grain boundary and unit length of a test line passed randomly through the sample. In practice, NL,, the average number of grain boundaries intercepted by unit length of traverse across a prepared microsection, is ob-

Jan 1, 1961

By M. Kaufman, A. E. Palty

THE mechanical properties of an alloy under given test conditions are in the main determined by the grain size and the phases present, their distribution and temperature-time dependence. Alloys intended for use at elevated temperatures, especially multiphase alloys depending upon precipitation reactions for strengthening, are particularly susceptible to changes in the phases present. Unfortunately, for modern superalloys, very little published information exists which relates the mechanical properties to the basic structural variables. The present work is an effort to do this for an alloy widely used in aircraft gas turbines and related applications, Udimet 500.* The mechanical properties inves- two heats is within the normal specifications for the alloy. However, the elements which most strongly influence the phase reactions, C, Ti, Al, and Mo, are comparable in the two heats. There was no intentional addition of boron, but normal pickup is about 0.002 pct. The vacuum-melted material was received in the form of centerless ground wrought bar stock % in. diam. PHASE STUDIES A) Methods—Two methods were used to determine the phases present: X-ray diffraction measurements on electrolytically digested residues, and microscopic examination. No single commonly used electrolyte will reveal all the phases that may be present; therefore, two complementary acids were selected: HCl and H3PO4. The necessity for this will be shown below in the results. The acid concentrations used were 10 pct HC1 in alcohol and 10 pct H3 PO4 in water. The current density was maintained below 0.1 amp per sq in. Residues were rinsed and separated by centrifuga-tion. X-ray diffraction patterns were obtained in a 114 mm Debye-Scherrer type camera using cobalt radiation. A temperature of 2250°F was selected to obtain as complete solutioning of phases as possible without encountering incipient melting. Specimens were

Jan 1, 1961

By J. L. Moriarty, Irving Johnson, H. M. Feder

The crystallization of ternary alloys has been studied as a coprecipitation phenomenon by measuring the substitutional carrying of trace metals with CeCd11 precipitating from liquid cadmium. The Doerner-Hoskins logarithmic distribution law has been found to apply in all cases; the coprecipitation coefficients obtained were: 1.48 (La), 1.06 (Th), 0.63 (Pr), 0.24 (Gd), 0.17 (Sm), 0.13 (U), 0.10 (Er), 0.09 (Sr). Li, Na, K, Ba, Y, Sc, and Zr were not measurably carried. It has been observed that coprecipitation was most extensive for those tracers which form with cadmium an intermetallic compmdnd isostructural with CeCd11 or closely related thereto. For these tracers an inverse relation between Solubility in cadmium and coprecipitation coefficient was noted. THE partition of impurities between the liquid and solid phases of a metal or a congruently melting intermetallic compound has been extensively studied by the methods of zone melting.' However, when the composition of a liquid alloy and of the solid phase crystallizing from it differ grossly, partition of an impurity cannot be investigated in this manner. Partition of the impurity under equilibrium conditions may be studied by investigation of the appropriate phase diagram and partition under nonequi-librium conditions may be studied by observation of coring in the solidified alloys.2 The nonequilibrium case is of special interest in metallurgical practice because many complex alloys are cooled too rapidly to permit the uniform diffusion of impurities throughout the crystallizing phase(s). With this in mind, the present study was undertaken to obtain quantitative data on the phenomenon of coprecipitation via substitutional solid solution in certain metallic systems. It will be noticed that rather than adopting metallurgical ter- minology we have chosen to refer to the partition of impurities during alloy crystallization as coprecipitation.3 In the terminology of coprecipitation the impurity is referred to as the tracer and the precipitating phase as the carrier. From aqueous solution, coprecipitation may occur as a result of adsorption on the surface of the carrier, by occlusion of the liquid within the crystals, or by solid solution of the tracer in the carrier. By slow cooling of liquid metallic solutions it is relatively easy to form large, regular crystals of the carrier; the effects of adsorption and occlusion may be minimized and the influence of substitutional solid-solution formation may be made to predominate. Previous studies of the crystallization of some intermetallic compounds4 and of copper, silver, and gold5 have shown that the coprecipitation or partition behavior of trace metals is entirely analogous to that observed in aqueous systems when ionic substitution occurs, e.g., coprecipitation of Ra++ with BaSo4. The precipitation of CeCd11 from liquid cadmium was chosen for study. This choice was influenced by the following consideration. If the relative atomic sizes of the tracer atom and the atom for which it substitutes determine the tendency to coprecipi-tate, then a lattice having a substitutional site of well-defined size would be desirable. In CeCd116 (structure type cubic BaHg117) the cerium atom is completely surrounded by a cage of twelve nearly equidistant cadmium atoms. Thus, atoms much larger or much smaller than cerium should be poorly accommodated in this cage. The choice of a lanthanide central atom allowed the testing of the size hypothesis with a number of lanthanide tracers which differed in size but not in valence or electro-negativity. EXPERIMENTAL The general technique of this investigation was to prepare a slightly undersaturated solution of carrier in liquid cadmium to which a tracer was added. The concentration of tracer was always so low that saturation over the entire temperature range investigated was never exceeded. The solution was cooled in a stepwise fashion. After each step a por-

Jan 1, 1964

By E. P. Abrahamson

The effect of transition element binary solid solution additions upon the recrystallization temperature of titanium has been investigated. All additions, except columbium and tungsten, raised the recrystallization temperature. A correlation is obtained between the logarithm of the absolute rate of change of recrystallization temperature with atomic percent solute and the number of uuter d shell electrons attributed to the free atom of the solute element in its ground state. The work of Abrahamson, et aL.1-6 on dilute binary solid solution alloys indicated a correlation between the free atom ground state electron configuration of the solute and both the brittle-ductile transition and recrystallization temperature. It is shown in these studies that the number of outer s and d electrons in the solute are of prime importance in determining the absolute rate of change of recrystallization or transition temperature with composition. Furthermore, the limit of initial linearity (in the curve of recrystallization temperature vs at. pct solute) has been shown to be a function of the number of d shell electrons in the solute for dilute iron- and vanadium-base alloys. To date no complete systematic recrystallization study of dilute binary titanium base alloys has been made. This study together with that of zirconium base alloys 6 will provide the data for the comparison of the effect of the same solutes upon two solvents with the same number of outer d shell electrons.

Jan 1, 1962

By O. N. Carlson, S. A. Bradford

Discontinuous yielding in tensile tests was observed in V-O alloys in the temperature ranges of 150° to 175°C and also 350° to 400°C. The magnitude and intensity of the serrations were found to vary considerably with oxygen content. Maxima were observed in tensile and yield strengths and in the strain-hardening coefficient at the higher temperature only. The strain rate sensitivity was observed to be negative between 150° and 400°C. THIS investigation was undertaken to study the effect of oxygen on the tensile properties of iodide vanadium in the temperature range of 25o to 450°C. Brown1 observed an increase in strength between room temperature and 400°C in vanadium metal, and found that oxygen and nitrogen had a rather pronounced effect on the strength and ductility. A maximum in the tensile strength was observed by Rostoker et al.2 near 300oC and by Pugh3 around 450°C for calcium-reduced vanadium. Pugh also found a maximum in the yield strength and in the strain-hardening exponent, and minima in the elongation and strain rate sensitivity at the same temperature. Eustice and Carlson4 reported the appearance of serrations in the stress-strain curves between 140° and 180°C in iodide vanadium containing 600 ppm O. These anomalies in the mechanical properties indicate that strain aging occurs in vanadium, but the impurity or impurities responsible for the above-mentioned effects have not been identified. The phenomenon of strain aging is usually characterized by the return of the yield point after interruption of a strength test. In the temperature range where strain aging occurs, the yield and tensile strengths attain maximum values, elongation and strain rate sensitivity exhibit minima, and discontinuous yielding is generally observed in the stress-strain curve. Cottrell5, 6 has postulated that strain aging is due to the migration of solute atoms to dislocation sites to produce locking after the dislocations have broken free from their impurity atmospheres during the initial yielding. At the strain-aging temperature the process is a dynamic one in which the solute impurity atoms diffuse to the vicinity of the moving disloca- tion producing "locking" which gives rise to maxima in the tensile strength and serrations in the elongation curves. Cottrel17 has noted that discontinuous yielding in iron occurs when the diffusion coefficient of nitrogen, D, and the strain rate, i, are related by D = 10-9 €. EXPERTMENTAL PROCEDURE The vanadium metal employed in this study was prepared by the iodide refining process as described by Carlson and owen.8 A representative analysis of the vanadium used in this investigation was: 150 ppm O, <5 ppm N, <1 ppm H, 150 ppm C, 150 ppm Fe, 70 ppm Cr, <50 ppm Si, 30ppm Cu, 20 ppm Ni, <20 ppm Ca, <20 ppm Mg and <20 ppm Ti. Alloys containing from 200 to 1800 ppm O, all of which lie in the solid solution range of the V-O system, were prepared by arc melting vanadium together with portions of a high-oxygen master alloy. The master alloy was prepared by tamping pure V2O5 into holes drilled in a vanadium ingot and arc melting this five or six times in an inert gas atmosphere, inverting the button between each melting step. The oxygen content of the master alloy was then determined by vacuum fusion analysis. Vanadium containing less than 150 ppm O was prepared in the following manner. A bar of iodide vanadium was deoxidized by sealing it in a tantalum crucible with a few grams of high-purity calcium. This was held at 1100°C for 4 days to allow time for the oxygen to diffuse to the surface and to react with the calcium vapors. The calcium oxide product was later dissolved from the surface of the bar with dilute acetic acid. In this way vanadium containing from 20 to 50 ppm O was prepared. Sample Preparation. The are-melted ingots were cold swaged into 3/16-in. diam rods and these were machined into cylindrical tensile specimens with a reduced section of 1.00-in. length and 0.120-in. diam. The test specimens were annealed for 4 hr at 900°C in a dynamic vacuum of mm of Hg to remove hydrogen from the metal. This recrystal-lization treatment produced a uniformly fine-grained structure with a mean grain size of approximately 0.06-mm diam. The oxygen contents reported in this paper were determined by a vacuum fusion analysis of the tensile specimens after testing. Analyses for other interstitial or metallic impurities showed no significant changes from that of the original material. Tension Tests. Tension tests were performed on a screw-driven tensile machine at a constant cross-head speed of 0.01 in. per min. Tests at elevated temperatures were carried out by heating the

Jan 1, 1962

By ZJ. H. Brophy, A. L. Prill, H. W. Hayden

The effect of Group VIII transition-element additions on the sintering of tungsten powders is shown to be dependent on the phase relationships of the particular tungsten-Group VIII metal alloy system. The acceleration of sintering with nickel additions is diminished by alloying nickel with copper which is known to diminish both the tungsten solubility in nickel and the solid-vapor surface energy. In work previously published on the effects of additions of Ni, Pd, Pt, Rh, Ru, and Ir, which form solid solutions or intermediate phases of wide composition width with tungsten, sintering behavior is independent of element concentration beyond that concentration necessary to form a continuous layer on the surface of the tungsten powders employed. The sintering behavior of tungsten with additions of iron or cobalt, which form intermetallic compounds with tungsten, is highly dependent on addition-element concentration. Additions in excess of amounts necessary to form continuous layers lead to successively decreased sintering rates. The effects of certain Group VIII transition metals on the sintering rates of tungsten at temperatures below those at which a liquid phase would be predicted have recently been investigated by the present authors1-4 It was found that small additions of palladium, nickel, rhodium, platinum, and ruthenium accelerate the rate of tungsten sintering. Additions of iridium decelerated the rate of densification compared to that of pure tungsten in the same temperature range. In all cases the rate of densification changed with increasing amounts of these additions, up to those compositions at which a continuous layer was formed on the surfaces of the tungsten-powder particles. The rate of densification was then independent of further increases of the compositions of any of these addition elements. These observations were consistent with a sintering model in which the addition element formed an interface with tungsten in which tungsten dif-fusivity was either greater or less than its grain boundary self-diffusivity. Kinetic analysis of this model indicated that linear shrinkage should vary according to the relation:

Jan 1, 1964

By A. R. Troiano, E. A. Steigerwald, F. W. Schaller

The initiation of localized cracking in hydrogenated high strength steel was dependent on the developmeni of a critical hydrogen conceniration and relatively insensitive to the magnitude of the applied stress. The stress was believed to influence the deluyed failure process by providing the means for grouping the hydrogen. On the basis of an assumed distribution law the observed changes in the lower critical stress as a function of notch acuity, yield strength, and temperature were predictable over a significant range of these variables. THE phenomenological characteristics of delayed failure in high-strength steel have been comprehensively investigated and the behavior has been attributed to the strong embrittling action of hydrogen.1-3 Delayed failure or, as it is also termed, static fatigue, is the most sensitive means of detecting hydrogen embrittlement since it allows the maximum time for hydrogen diffusion during a test.3 The nature of delayed failure cail be aptly described by four parameters which are schematically depicted in Fig. 1. The upper critical stress corresponds to the rupture stress in a conventional notch tensile test. The lower critical stress or static fatigue limit is the stress below which no delayed failure will occur. The incubation period is the time required for the formation of the first crack and the fourth parameter is the failure time. An increase in hydrogen content, notch acuity, or strength level of the material increases the susceptibility for delayed failure by decreasing the lower critical stress, the incubation period, and the time to failure.lY3 A decrease in test temperature prolongs the incubation period in accordance with the lower diffusion rate of hydrogen, and also de-

Jan 1, 1961

By A. H. Daane, B. J. Beaudry

The Sc-Y and Sc-Zr systems were studied by thermal and X-ray methods. Both systems are characterized by complete solid solubility in the low temperature hexagonal form and in the high temperature bcc form. In the Sc-Y system, there is a minimum in the solidus at 50 at. pct and 1365°C, while the solidus in the Sc-Zr system follows a nearly straight line relationship between the melting point of scandium and zirconium. There is a minimum in the temperature of transition from hexagonal to bcc in the Sc-Y system at 43 at. pct Y and 1175°C; while the temperature of transition is raised in the Sc-Zr system to a maximum of 1415 °C at 30 at. pct Zr. INTEREST in the metals of the first transition series has usually stemmed from their useful properties; however, there is now a great deal of interest in determining the fundamental nature of these metals. Since scandium is the first of these transition metals, a study of the preparation and Properties1 and the alloying behavior of scandium2 was initiated. In a continuing study of the alloying behavior of scandium, the systems of scandium with zirconium and yttrium have been established. Extensive solid solubility in both the Sc-Zr and Sc-Y systems was expected from application of the classical rules of Hume-Rothery. Also, the modified Darken and Gurry plot as suggested by Gschneid-ner3 showed ideal conditions for solubility in the Sc-Zr system with less favorable conditions in the Sc-Y system. The effect of scandium on the transformation temperature of zirconium and yttrium was of interest, since a recent empirical observation by Uly and Burr4 predicted that the solute scandium would raise the transformation temperature of zirconium and lower the yttrium transformation. In addition to noting the validity of the above predictions, the SC-Y and Sc-Zr systems were established. EXPERIMENTAL METHODS Materials. The metals employed in this investigation were crystal bar zirconium, distilled scandium prepared by the method described by Spedding, et al.1 and yttrium prepared by the method described by Carlson, et a1.5 The major impurities present in these metals are listed in Table I in ppm. Alloy Preparation. All the alloys examined were initially formed by comelting the two metals in a nonconsumable electrode arc-melting furnace under an atmosphere of purified helium or argon. The buttons were inverted at least five times and re-melted each time to ensure homogeneity. During preparation of Sc-Zr alloys, approximately 1 wt pct was lost which was attributed entirely to the loss of scandium by distillation. When preparing Sc-Y alloys, approximately 1.5 wt pct was lost, but it was assumed that the intended composition was maintained, since the relative weight lost when arc melting the component metals was nearly the same. Samples weighing 15 to 20 g were prepared for time-temperature thermal analysis, while only 1- to 2-g samples were prepared for the optical pyrometer and X-ray diffraction studies. Thermal Methods. The solidus and the melting points of the metals were determined by an optical pyrometer method described previously.2 The temperatures at which the transformation from hexagonal to bcc occurs were determined by differential thermal analysis. Several points on the solidus and liquidus in the Sc-Y system were also determined by this method and are in good agreement with those obtained by the optical pyrometer method. The procedure followed in the preparation of samples for thermal analysis and the furnace used have been described elsewhere.2&apos;6 Since a limited quantity of scandium was available, most of the alloys were not melted during thermal analysis, thus avoiding contamination by the crucible. The alloy could then be reused and the composition changed by adding the desired quantity of solute to obtain the next alloy. Heat Treatment and X-Ray Methods. In order to ensure homogeneity, the "as arc-melted" alloys Were wrapped in tantalum foil, sealed in quartz capsules and heated at 1100°C for 15 hr and then cooled at a rate of 8 deg per min. Samples for X-ray diffraction were filed from these heat-treated alloys, sealed in 1/8 in. by 1/2 in. tantalum capsules under

Jan 1, 1963

By B. F. Oliver

Equilibrium and kinetic segregation coefficients for the dilute high-temperature system tantalum in iron have been experimentally determined in a levitating zone melter. The equilibrium segregation coefficient, ko, is 0.43 * 0.03. Planar interface solidification is possible up to a freezing velocity of 4 in. per hr but not at 6 in. per hr. Precise concentration control of the solid is possible over this range of freezing velocities. The momentum boundary layer thickness, and the diffusion boundary layer thickness, 6, are such that either 6, > 6 or ALTHOUGH considerable data exists in the literature concerning the stability of planar-solidification fronts, little data exists for high-temperature metallic systems. The study of fundamental aspects of solidification in high-temperature iron-base systems is complicated by the extensive solvent power of liquid iron. In order to minimize such effects a levitating zone melter was used to study the segregation of tantalum in iron.&apos; The solute tantalum was selected because: 1) the ko value was, according to the literature, initially thought to be less than 0.1, thus increasing the precision of the method used;2,3 2) the vapor pressure and is a constant, fixed by induction stirring, 6, the diffusion boundary layer, is 5150 u. At 400 rpm and 6 in. per hr freezing velocity the cast structure exhibits circular macroporosity and banding concentric with the axis of the rotation of the interface. Electron-beam microprobe analyses indicate that the tantalum content is highest at the center of rotation (1.2/1) and decreases with increasing radius as the fine-grained bands become sharper. This entire structure exhibits microseg-regation of tantalum. of tantalum in 1 atm of H2 is small, assuring conservation of solute during redistribution; and 3) spec-troscopic analyses were available for the appropriate range of tantalum considered. Additional objectives were 1) to determine kinetic segregation coefficients for well-stirred zones with constant boundary conditions and thus quantitatively calibrate solute-segregation behavior in the levitating zone melter, and 2) to examine the range of freezing velocities over which smooth solid-liquid interfaces could be controlled in an iron-base alloy under the restrictions of the thermal gradients imposed while levitating a zone. The single-pass zone-melting method was used to calculate individual k values.&apos;

Jan 1, 1964

By R. P. Abendroth

This study is concerned with the kinetics of selective oxidation of chromium in a commercial Fe-Cr-Ni alloy. Selective oxidation of chromium in this alloy, by use of a low oxygen-potential atmosphere, leads to the formation of a compact, protective layer of Cr2O3. This layer serves to protect this alloy from gross scaling when it is subsequently exposed to severe oxidizing conditions at high temperature. The interaction of low oxygen-potential atmospheres close to equilibrium with Fe-Cr and Ni-Cr alloys has been studied by others."&apos; These studies werk concerned with the surface-structure variations under slightly oxidizing conditions. NO detailed study was made of the oxide scale formation kinetics, however. The alloy samples were cut from 0.012-in.-thick sheet, with an apparent surface area of 7.1 sq cm. These sheet samples were abraded through 4/0 metallographic paper, and washed in alcohol and acetone. The analysis of the sheet alloy is (in weight percent): 42 pct Ni, 5.5 pct Cr, 0.09 pct C, 0.18 pct Al, 0.36 pct Si, 0.26 pct Mn, balance Fe. The weight gain vs time data were obtained with a 2-g capacity fused-silica spring—cathetometer system. The spring deflection was optically magnified ten times before being read by the cathetometer. A sensitivity of about 0.01 mg was attainable. The spring was enclosed in a water jacket maintained at 60°C to minimize the effect of temperature changes. The sample was suspended from the spring with a fused-silica hangdown and was positioned in the thermal center of a mullite furnace tube. Sample temperature was read with the aid of a thermocouple placed outside the mullite tube and an inside vs outside temperature calibration. Temperature change during the course of a run was ±0.25°C, with a temperature gradient of less than l.O°C over the length of the sample. Total temperature uncertainty was no more than ±3.0°C. Alignment difficulties between the hangdown and radiation shields in the furnace tube required that the sample be positioned in the furnace when cold, and heated with the furnace until the temperature stabilized at the desired point. This required 5 to 6 hr, and was carried out in Matheson ultrahigh-purity hydrogen. Oxidation was started, after evacuation, by introducing a hydrogen-water vapor mixture, obtained by saturating hydrogen with water vapor at 31.00o ± 0.02oC. Oxidation was continued for 90 min. Gas flow was 300 ml per min during heat up and oxidation. Since only several milligrams of oxide are formed on each sample, chemical analysis of the oxide is difficult. A representative analysis is: 80 pct Cr2O3, 5 pct Fe2O3, 3 pct Al2O3, 4 pct MnO, 7 pct SiO2, 1 pct or less NiO. X-ray diffraction analysis of the oxide as formed on the alloy gives rhombohedra1 Cr2O3, and barely distinguishable amounts of a cubic spinel phase, and possibly AlZOs and SO2. The identification of these latter two compounds is by no means certain. The spinel phase could be based on iron or manganese as these elements are present in significant amounts in the oxide. The results of the kinetic studies using 31°C dew-point hydrogen-water vapor mixtures were found to conform to a parabolic rate law. In many cases the parabolic plot consisted of two intersecting straight lines, defining an early and a late rate for a particular run, and in the other cases the parabolic plot consisted of one straight line for the entire run. The slopes of the various straight lines were determined by the method of least squares. Reproducibility of the data was good enough for multiple runs at the same temperature such that the value given for the rate constant is the average for two or more closely similar values, rather than widely varying values of the rate, where more than one determination is indicated in Table I. The exhibition of only one or of two rate constants during a run can happen at the same temperature. Thus, Table I shows that at 11&apos;74°C a single rate constant was obtained from one sample, while other samples oxidized at the same temperature gave an early and a late rate constant. It should be noted that the single rate constant corresponds very closely with the early rate constant. This is also true at 1153°C. The time at which the late rate started to appear was variable, usually occurring 20 to 40 min after oxidation had started.

Jan 1, 1964

By L. S. Castleman, R. Resnick

The self-diffusion coefficient of columbium has been measured over the temperature range 1535o to 2120°C, using the radioactive isotope Nb95 as the tracer. The data for the temperature dependence of the self-diffusion coefficient are best fitted by the equation D = (12.4 ± 0.8) exp [(-105,000 ± 3000)/RT] cm2/sec. THE increasing need for high-temperature structural materials has stimulated considerable interest in the properties of columbium, since it is a strong, ductile metal with a high melting point—in the vicin- ity of 2450oc.l One fundamental property which has not heretofore been measured is the self-diffusion coefficient. The self-diffusion coefficient plays an important role in theories of oxidatioq2 creep,&apos; sintering4 and other phenomena. Accordingly, the investigation reported in this paper was undertaken to determine the value of the self-diffusion coefficient and its temperature dependence by the use of Mb95.

Jan 1, 1961

By J. E. Sandor, P. F. Schmidt

A general review of the present knowledge of the semiconductor/oxide interface is given, with emphasis on silicon and germanium. The electrical properties of clean surfaces and the changes that occur upon oxidation are described. A brief disczcssion of the methods and mechanisms of oxidation of silicon is given, and the interaction between the semiconductor substrate and the oxide phase is considered. Physical properties of the oxide film, charge storage effects due to the dielectric, and some device applications of oxide films on semiconductors are discussed. THIS paper will be heavily preoccupied with the elemental semiconductors germanium and silicon. Undoubtedly some injustice is done thereby to other semiconductor-oxide systems, but it is not readily possible to cover all aspects of so vast a field in the present review. In considering the semiconductor/oxide interface, we may start out by discussing briefly an idealized clean semiconductor surface, and may then follow the changes brought about by oxidation, as the oxide is allowed to grow progressively thicker. The references given are not a complete list, but are to be taken as guides to the literature. Tamml suggested in 1933, based upon the study of a one-dimensional asymmetrical potential model, that localized surface states may appear at the surface of a crystalline lattice; it was pointed out by Fowler that these states should occur in pairs. shockley3 examined the origin of such surface states on the basis of a more general one-dimensional model and found that surface states may appear in the forbidden region between the highest filled and the lowest vacant band. This happens when the two bands overlap, as, for instance, the "s" and (&apos;p" bands of the diamond structure. In a three-dimensional model these levels would produce a surface band which is half-filled for diamond, and the number of surface states should be approximately equal to the number of surface atoms. This picture of the surface has been refined by more recent theoretical work, like Handler&apos;s4 and Koutecky and Tomasek&apos;s~ without introducing essential modifications. When these concepts were applied to semiconduc-

Jan 1, 1965

By K. M. Busen, J. Lindmayer

A heterojunction model is suggested for descrihzng the electronic conditions at oxide -semiconductor interfaces. Detailed study of the silicon oxide-silicon interface shows that the heterojunction model explains many experimental facts; however, it is necessary to assume that a transition layer exists at the interface. It will he demonstrated that ionic contrbution way also he present. Extended low-temperature heating stabilizes the oxide to a charge carrier density in the order of 10&apos; cln-: determined by the heterojuncton model. OXIDIZED semiconductor surfaces are widely used, since they offer technological advantages and "passivate" the surfaces of semiconductor devices. However, it has been recognized that there are always n-type inversion layers formed in the silicon oxide-silicon system, which must be regarded as a definite disadvantage. An added complication is the "instability" of these inversion layers;&apos; the inversion layer appears to vary with the processing and they change at temperatures as low as 100" to 200°C This paper will point out that the inversion layer is a result of the dissimilarity of the contacting materials and therefore must always be present. The dissimilarity is the essence of the heterojunction or electronic model and involves directly the structure of the materials involved. In addition, one finds the contribution from ions; this disturbance, however. can be eliminated for the case of silicon oxide-silicon interface. The strength of the inversion layer or channel will be expressed in terms of the density of elementary charges per square centimeter, N*, maintaining the channel. The next section contains the formulation of the heterojunction model by assuming a contact between two homogeneous materials. This calculation should not imply that silicon oxide is homogeneous; it serves only to point out the general properties of the electronic model. Following the electronic model, the practical case of the silicon oxide-silicon interface will be discussed. I) THE HETEROJUNCTION MODEL Contacting an oxide with a semiconductor, a work-function difference must be accommodated. Fig. 1 represents the oxide and the semiconductor before and after contact. When the contact has been made, one finds that qVi + qV, = A@; i.e., the work-function difference is distributed between the oxide and the semiconductor. Normally, the work-function difference appears on the wide-bandgap material, but in the present case the oxide is thin and far from being a single crystal. Accordingly, a significant fraction of the work-function difference must appear on the semiconductor. The work function of silicon2 is about 4.8 ev, but we shall not assign any particular work function for the oxide as it may vary. However, oxides usually have lower work functions and therefore the potentials must build up in the manner shown in Fig. 1. Such an argument predicts that the inversion layer must be n type, as opposed to the clean surface which is p type. Another important factor to be considered is that the oxide is disorderly and therefore contains many trapping levels. As a result, the charges in the oxide arise primarily from the change in trap occupation, while on the semiconductor side the space charge is composed of mobile free electrons. The consequence of this fact is that a change to another thermal equilibrium will require some time, dictated by the traps in the oxide. The conditions can easily be calculated on the semiconductor side. Assuming an intrinsic semiconductor, Poisson&apos;s equation is

Jan 1, 1965

By P. Greenfield, P. A. Beck

The phase is a hard and extremely brittle material with a tetragonal crystal structure, containing 30 atoms per unit cell&apos; It occurs in many binary and ternary alloys of the transition elements. The existence of twelve a phases has been reported in binary alloys of metals from the group vanadium, chromium, molybdenum, and tungsten with metals from the group manganese, iron, cobalt, and nickel. It is a characteristic fact that the composition of the o- phases in the various binary systems is not the same. Rather, there is a gradual shift in composition, for instance in the series of chromium a phases, approximately as follows: CrMn,, FeCr, Cr3Co,. In recent years several ideas have been put forward to account for this variability of composition. SullyZ proposed that a was a type of electron compound, characterized by 1.7 electrons per atom in excess of the number required to fill the Pauling atomic orbitals. These excess electrons were supposed to just fill the first Brillouin zone of the u structure. However, later work indicated that the phase is ferromagnetic at low temperatures, so that a- is unlikely to be a full zone compound.~ Another proposal,* which is also based on the structure of the transition elements advanced by Pauling, and in fact leads to essentially the same a compositions as Sully&apos;s theory, suggests that the u phase is characterized by a constant number of electron vacancies per atom. The electron vacancy numbers N,. used were originally the same as those in the Pauling theory, but were later modified empirically vn order to bring about better agreement with data for the p phase in certain ternary systems. For alloys, N,, was calculated on the assumption of simple additivity from the electron vacancy numbers of the components, as follows: N,. = 4.66 (Mo + Cr + V) + 3.2 (Mn) + 2.2 (Fe) + 1.71 (Co) + 1.6 (Ni). On this basis a reasonably constant value of N, = 3.4 was obtained for all of the binary u phases and a fairly constant value of N,, = 3.1 for the various p phases. More recently, Bloom and Granta pointed out that a similar correlation with the total number of 3d and 4s electrons per atom is also fairly well fulfilled and concluded that no reference need be made to the Pauling theory. This proposal does not lead to the same binary a compositions as the one described above, so that the relative merits of the two proposals may be determined on the basis of direct comparison with experimental composition data. The purpose of the present paper is to provide such a comparison. Since the experimental data for the composition ranges of several of the known u phases were rather incomplete, it was decided to carry out new experimental determinations for many of the binary systems. Experimental Procedure The alloys were prepared by induction melting in recrystallized alumina crucibles, either in vacuum or in helium atmospheres. Some aluminum was picked up from the crucibles in the alloys of high vanadium content, but this was minimized by keeping the alloys in the molten state for as short a time as possible. The maximum aluminum impurity was 0.13 wt pct by chemical analysis, in an alloy containing over 50 wt pct V. In general, the as-melted compositions agreed accurately with those obtained by chemical analysis. In the case of manganese, however, losses inevitably occurred; consequently all manganese alloys were analyzed. Specimens of all alloys were annealed at temperatures ranging from 1000" to 1300°C in an atmosphere of purified 92 pct helium and 8 pct hydrogen mixture. All manganese alloys were sealed in quartz tubes under vacuum before annealing. The specimens were held at temperature for between four and eight days before being quenched into cold water. Alloys were prepared for microscopic examination by mechanical polishing, usually followed by electrolytic etching. X-ray diffraction specimens were prepared by crushing the homogenized alloys, and by reanneal-ing the obtained powders in evacuated quartz tubes for a few hours at the temperature of the lump anneal. X-ray diffraction patterns were made with unfiltered chromium radiation, using an asymmetrical focusing camera of high dispersion. Using microscopic and X-ray techniques jointly, the accuracy

Jan 1, 1955