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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

By J. B. Clark

In 1933, R. paris1 published a phase diagram for the entire Mg-Ca-Zn system. He reported the existence of a ternary intermetallic phase, Mg5Zn5Ca2, which formed a quasi-binary with the magnesium solid solution. Since then, no extensive work has been reported on the system. Recently however, indications of the existence of a second ternary in-termetallic phase in equilibrium with the magnesium solid solution led to a reinvestigation of the solid phase equilibria involving the magnesium solid solution. The 335°C (635°F) isothermal section was chosen for study. Metallography and powder X-ray diffraction were selected as the best mutually corroborating methods for determining the solid phase constitution. Seventy-six alloys were prepared from singly sublimed (99.95 pct) magnesium, electrolytic (99.95 pct) zinc and USP calcium. The analyses of some of the alloys are shown in Fig. 1. The alloys were melted in a graphite crucible under a chloride flu and chill cast in a split graphite mold. Dif-fusion couples (see Fig. 2) were prepared by casting molten Mg-20 wt pct Zn alloy around a cylinder of Mg-30 wt pct Ca alloy. The alloy samples and the diffusion couples were sealed under argon in Pyrex vials and annealed in a furnace capable of controlling within 1/2°C of 335°C. At the conclusion of the anneal, the alloy samples were quenched in ice water and the diffusion couples were air cooled. Procedures standard for magnesium alloys were used for metallographic examination. Powder X-ray diffraction patterns were taken on a G. E. camera with a 14.32 cm dim using nickel-filtered copper radiation. The redetermined 335°C (635°) section of the Mg-Ca-Zn phase diagram is shown in Fig. 1. TWO ternary intermetallic phases, designated w and

Jan 1, 1962

By R. E. Ogilvie, A. R. Kaufmann, S. H. Gelles

The solid-solubility limits of iron in beryllium were determined between 850o and 1200oC by analysis of differential type multiphase diffusion couples, using an X-ray absorption technique. The maximum value of the solubility limit was found to be 0.92 ± 0.02 at. pct (5.46 wt pet) at the eutectic temperature 1225°C. The solubilities of nickel and beryllium were determined between 900°and 1200°C by the same technique and the maximum solubility was found to be 4.93 + 0.01 at. pct (25.2 wt pet) at the eutectoid temperature, 1065°C. A previously unreported high-temperature phase which decomposes eutectoidally at 1065 °C was found to exist in the beryllium-nickel system at a composition of approximately 8 at. pct Ni (36 wt pet) by diffision-couple analysis. The presence of this phase was confirmed by thermal analysis and metallo-graphic analysis of the structure resulting from the eutectoid decomposition. G. V. Raynor1 has treated the solid solubilities of some of the elements in beryllium on the basis of the "Hume-Rothery" rules2 which have been modified to include ionic size and ionic distortion effects. It was predicted that the solubility of iron and nickel in beryllium should be slightly less than that of copper. The lowering of the solubility, according to Raynor, is due to a more unfavorable relative valency effect and an ionic size effect. Kaufmann and corzine3 have compiled data on the solubilities of elements in beryllium and have discussed them in the light of the Raynor paper. These authors feel that, because the elements having the greatest solubility in beryllium systematically fall in the Group VIII and IB Columns of the periodic table, the electronic structure greatly influences the maximum solid solubility of elements in beryllium. The solubility of iron in beryllium was determined by Teitel and cohen4 as part of the study of the beryllium-iron phase diagram. The determination was carried out by X-ray and thermal analysis and according to the phase diagram presented, the maximum solubility of iron in beryllium is 0.41 at. pct (2.5 wt pct) at 1225oC. However, it is estimated that the uncertainty in the position of the a-beryllium primary solid-solution boundary is about 0.5 at. pct (3wtpct). Losana and Goria3 in studying the beryllium-nickel phase diagram, determined the solid solubility of nickel in beryllium by thermal analysis. They found the maximum solubility to be between 1.65 and 2.65 at. pct (10 to 15 wt pct) at 1240°C. This value decreased rapidly with decreasing temperature. In determining approximate ranges of solubilities for different elements in beryllium, Kaufmann, et al,8 reported a value of between 1.3 and 1.7 at. pct (7.9 to 10.1 wt pct) for the solubility of nickel in beryllium. The value was obtained by metallographic examination of quenched alloys and lattice-parameter measurements. However, the authors also noted a single-phase structure for a 1.7 at. pct Ni alloy (10 wt pct) on cooling from the liquid. This would indicate a higher solubility range than was reported. ~isch,&apos; in his X-ray studies of beryllium-copper, beryllium-nickel, and beryllium-iron intermetallic compounds, reports the disappearance of a second phase (Ni,Be2) in the beryllium primary solid solution at approximately 4 at. pct (20 wt pct). THEORY The analysis of concentration gradients in diffusion couples has proven to be a useful tool in determining phase equilibria.8-14 In this particular study the diffusion couples were chosen to straddle the expected composition range of the phase boundary, then heat treated at a given temperature and the concentration gradient evaluated. The composition of the phase boundary for a given temperature appears at a point of discontinuity of the composition gradient. Examples of typical phase diagrams and the concentration gradients which should be found in such systems are shown in Fig. 1. In the present work, gradients of the form of Fig. l(c) were obtained in diffusion couples made of pure beryllium and two-phase alloys of beryllium with either iron or nickel. The composition at the point where the gradient becomes discontinuous, Cs, corresponds to the solubility limit of either iron or nickel in beryllium. The analysis of the concentration gradients was carried out by an X-ra absorption method developed and applied by Ogilvie and later used by Moll13 and Hilliard.l4 It depends on the fact that the absorption of X-rays by matter is determined by the concentration and type of the various atomic species present. The relationship for the intensity, I, of a monochro-

Jan 1, 1960

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

The solid solubility of holmium in copper, silver, and gold, determined using metallographic techniques, was found to be a maximum of 0.06, 2.4, and 3.2 wt pct, respectively, nt the eutectic temperatures. The temperatures of these eutectic horizontals, determined by differential thermal analysis, were found to be 868", 787", and 778"C, respectively. The first compound on the silver-rich side of the Ag-Ho system was found to be A&Ho2. In a study by Gold et al.,&apos; on the effect of small amounts of soluble impurities on the electrical resistance and thermoelectrical properties of copper, traces of iron were shown to have drastic effects on the thermoelectric power due to the magnetic scattering of the conduction electrons. Since little is known about the state of the iron ion when dissolved in copper, Gold had hoped to repeat these experiments using certain rare earths as the magnetic scattering centers; because of the tightly bound 4f electrons it is not expected that their spectroscopic state would be affected by alloying. In such an experiment the sole role of the trace rare-earth impurity would be to magnetically scatter the copper conduction electrons without producing appreciable distortions in the otherwise pure copper matrix. Since holmium was chosen for such a study because of its high magnetic moment, it was necessary to establish precisely its solid solubility in copper. The study was extended to include the solvents silver and gold, so that an evaluation of the data using classical solubility parameters could be made. No previously reported work on the systems of holmium with copper, silver, or gold was found in the literature. Information on the solubility of other rare earths in these three metals is also very meager. The solubility limit of yttrium in copper was reported as less than 1 wt pct2 and the solubility of praseodymium in gold as 0.35 wt pct,3 while no references were found on solubility studies of rare earths in silver. EXPERIMENTAL Materials. The materials used in this investigation were distilled holmium prepared in this Laboratory, OFHC* copper, electrorefined silver (99.9 *OFHC = oxygen-free high conductivity. pct), and gold with the analyses given in Table I. Sample Preparation. Samples weighing 1 to 2 g were prepared by arc-melting the components together on a water-cooled copper plate under purified argon atmosphere. The samples, turned and remelted at least five times to ensure homogeneity, showed weight changes of less than 0.1 wt pct after melting. Samples containing less than 0.7 wt pct Ho were prepared by diluting 1 to 2 wt pct Ho alloys with the pure solvent, while samples with concentrations greater than 0.7 wt pct Ho were prepared directly from the metals. The samples were wrapped in tantalum foil, sealed in outgassed Vycor tubing under helium, annealed at the desired temperature for at least 24 hr, and then quenched by shattering the Vycor tube in cold water. Metallographic Methods. Standard mechanical polishing methods were used in preparing the silver and gold alloys for examination, while the copper alloys were electropolished using a 40 pct solution

Jan 1, 1963

By Walter Showak

The published phase diagrams for the copper-zirconium system1,2 indicate that the solid solubility of zirconium in copper is not accurately known. Studies made on relatively impure alloys by Pogodin, Shumova, and Kugucheva3 yielded a maximum solubility value of 1 wt pct Zr at the eutectic temperature. Allibone and sykes5 reported that an as-cast 0.2 wt pct Zr alloy is single phase at room temperature. Limited results obtained by Raub and Enge14 indicated that zirconium solubility in copper is less than 0.13 wt pct Zr at 940°C. Tn 1960, on the basis of a microstructural examination of several copper-rich Cu-Zr alloys, saarivirta6 placed the maximum solubility of zirconium in copper at 0.15 wt pct at the eutectic temperature, 980°C. The present note describes results obtained during a re-investigation of the Cu-Zr system for the purpose of accurately establishing the solidus and solvus curves and the eutectic temperature in copper-rich alloys. High-purity copper (99.99 pet) and zirconium (99.64 pet) consolidated reactor grade sponge were used to prepare several alloys, containing up to 2 wt pct Zr, by induction melting in graphite crucibles in vacuum and chill casting into a split steel mold. The melting procedure consisted of first melting the copper charge and then adding the bulk zirconium wrapped in pure copper foil. Chemical analyses were made of both the cold-rolled alloy ingots after homogenization at 925°C for 450 hr in evacuated quartz capsules and of small specimens which had been used in the phase diagram studies. &apos;The carbon content of each alloy was also determined and was found to be less than 0.001 wt pct in every case. Metallographic and electrical resistivity measurements were used as the primary methods for this investigation, although some thermal analysis studies were also performed. Heat treatments of all the alloys were carried out with the specimens wrapped in molybdenum or alloy foil (of the same composition as the specimen) and sealed in quartz tubes either in vacuum or under a purified helium atmosphere. Following high-temperature anneals for long periods of time, slight depletion of zirconium from the specimen surface was observed. Therefore, careful preparation of samples for microscopic examination and chemical analysis of the samples themselves was necessary to check the composition of samples used in determining phase boundaries. Points on the solvus curve were determined on quenched specimens by both isothermal reaction after solution heat treatment at 950°C for 8 hr and by isothermal annealing of alloy specimens at successively higher temperatures for times ranging from 24 to 500 hr. Liquation studies for positioning of the solidus and eutectic temperatures were performed by holding a pre-

Jan 1, 1962

By F. Weinberg, E. H. McLaren

Dilute binary alloys have been solidified under controlled thermal conditions, and solute distributions, temperatures during freezing and melting, and the position and morphology of the solid-liquid interfaces have been examined. Binary Zn, Sn, and Pb based alloys were investigated with solute additions of AgU°, AU°, Sb&apos;24, TlZo4, snH3, and Zns5. Approximately 150 cc of metal was used in each ingot. Appreciable solute segregation occurred in most of the ingots, depending on the solute concentrations, distribution coefficients. and freezing conditions of the alloys. Measurements were made of the general solute distribution by sectioning layers of material from the ingot circumference and counting the activity of each layer. Local segregation was detected by autoradiographic techniques. Freezing and melting curves for the same alloys were obtained by precision resistance thermometry. These curves are related to the corresponding measured solute distribution curves and those calculated from theoretical expressions. The position and morphology of solid-liquid interfaces during freezing were determined by both decanting and quenching techniques. The results indicate the manner in which freezing and remelting occurred under the conditions investigated. If an isomorphous dilute binary alloy is solidified or melted under equilibrium conditions, the concentration of solute in both solid and liquid would be uniform at all stages in the freezing process. The solute concentrations during freezing could be determined from the phase diagram of the alloy, and the freezing curves would have an alloy slope defined by the solidus and liquidus points. In practice, however, solidification under equilibrium conditions rarely, if ever, occurs, primarily because of low rates of diffusion in the solid and incomplete mixing in the liquid. As a result, the solute distribution and the shape of the freezing and melting curves obtained in a given system depends on a number of factors. These include, in addition to the extent of diffusion and mixing, the initial solute distribution and concentration, the distribution coefficient, the rate of freezing or melting, supercooling of the liquid, the morphology of the solid-liquid interface, and the effect of annealing. The purpose of the present investigation was to determine, experimentally, actual solute distributions and thermal curves for a series of dilute binary alloys, using tracer techniques to determine the solute distributions and precision resistance thermometry for the temperature measurements. In addition, observations were made of the positions and morphologies of the solid-liquid interfaces during freezing and remelting. The distribution of solute in a dilute alloy, assuming complete mixing in the liquid, and negligible diffusion in the solid, is given by1 where Co is the mean solute concentration of the alloy, C, is the solute concentration in the solidg is the fraction of the ingot solidified, and ko is the equilibrium solute distribution coefficient (assumed to be independent of concentration for dilute alloys). For the case of partial mixing, ko can be replaced by an effective distribution coefficient k. If solute transport in the liquid is assumed to be entirely due to diffusion, with negligible diffusion in the solid, then the solute distribution in the solid is given by Tiller et a1.&apos; as where R is the rate of advance of the solid-liquid interface, D is the solute diffusion coefficient in the liquid, and x is the position of the interface from the start of freezing. The temperature at the interface TI is related to the solute concentration of the solid at the interface C, by the expression where Tm is the freezing temperature of the pure solvent and ml is the slope of the liquidus for the binary alloy being considered. Accordingly, by relating TL to the actual temperatures measured in the present experiments, solute distributions can be deduced from the temperature measurements. Using the above expressions, the measured solute distribution of a given alloy may be compared to the theoretical distributions, as well as to the distribution calculated from the observed freezing curve. EXPERIMENTAL The solvent materials used in making up the binary alloys were C. P. New Jersey zinc (99.999 pct), Vulcan Extra Pure Tin (99.999 pet), and Tadanac lead (99.998 pct). The solute elements were of 99.99 pct purity or better, and were activated in the

Jan 1, 1963

By Robert A. Rapp, Jan E. Verfurth

From the measurement of internal-oxidation band widths in Ag-In and Cu-A1 alloys, products for the solubility No and the diffusivity Do of oxygen in silver and apparent NODo products in copper have been obtained over a range of temperature. The NODO products for silver are in agreement with the product of the individually determined No and Do values of Eichenauer and Müller. Although interstitial diffusion of oxygen in silver is expected, the temperature-independent factor Do of the diffusion coefficient does not agree well with the theory of Wert and Zener. For Cu-A1 alloys internal oxidation does not follow the expression derived for ideal kinetics, since the apparent NODo product calculated from internal-oxidation measurements is a function of the alloy content in the specimen. Thus, unequivocal values for No or DO or even the product NODO in copper are not available from the literature. RELIABLE values for the saturation solubility NO and diffusivity DO of oxygen in fcc metals are lacking and difficult to determine, partially because internal-friction measurements are not applicable. Permeability experiments and Sieverts7-type adsorption -desorption methods3 have met with difficulty; however, Eichenauer and Müller4 have recently reported on the solubility and diffusivity of oxygen in silver as determined by the latter technique . Rhines, Johnson, and Anderson,5 and independently Meijering and Druyvesteyn,&apos; introduced the possibility of deriving values for the product NODO in a metal from measurements of the widths of internal-oxidation zones (subscales) formed during the internal oxidation of alloys. In addition to these authors, several other investigators7-9 have made such measurements in various silver and copper alloys at various temperatures, mostly using me- chanically polished specimen surfaces. The purpose of this investigation was to determine over a range of temperature the calculated products NODO in silver and copper from systematic measurements of the widths of internally oxidized zones in Ag-In and Cu-A1 alloys, using specimens with electro-etched surfaces. Consider the internal oxidation of a dilute solid-solution alloy in which the base metal is not oxidized. If diffusion of oxygen is the rate-controlling step, then the depth 5 of the internal-oxidation zone is given by where y is a dimensionless parameter and t the oxidation time. Based on the solution of the diffusion equation, wagnerl0 has shown that under certain conditions several mathematical simplifications may be used in the evaluation of y, so that Eq. [I] may be expressed for the case that y « 1, and y(Do/DB) 1, where N(so)the is the equilibrium solubility (mole fraction) of oxygen in the solvent metal for the given oxidizing atmosphere, is the mole fraction of the solute metal in the bulk alloy, 4 is the diffusivity of the solute in the base metal, and v is the ratio of oxygen atoms to solute atoms in the precipitated oxide BO,. Rearranging Eq. [2] Maak&apos; has derived an expression analogous to Eq. [3] for the case that an external-oxide scale, in addition to an internal-oxidation zone, is formed. For this combined external and internal oxidation for the case that 5/2 where 5 is the distance of the internal-oxidation front from the original external s.urface, and X is the distance of the external oxide/metal interface from the original external interface. The function F is defined by the expression

Jan 1, 1964

By J. F. Enrietto

Internal friction measurements were used to determine the effect of manganese on the solubility and precipitation kinetics of nitrogen. Manganese, in concentrations up to 0.75 pct, has little effect on the solubility at temperatures above 250°C. On the other hand, at Concentrations as low as 0.15 pct, manganese inhibits the formation of iron nitrides, especially Fe4N, even though it may not form a precipitnte itself. The precipitation and solubility of carbides and nitrides have been extensively investigated in the pure Fe-C and Fe-N systems.1-3 In recent years, some effort has been ispent in studying the influence of substitutional alloying elements on the behavior of carbon and nitrogen in ferrite.4 -7 In particular Fast, Dijkstra, and Sladek have investigated the effect of 0.5 pct Mn on the internal friction and hardness during the quench aging of Fe-Mn-N alloys.&apos;, &apos; They found that at low temperatures (below 200°C) the presence of 0.5 pct Mn greatly retarded quench aging. For example, after 66 hr at 200°C very little precipitation had taken place in the iron alloyed with manganese, whereas precipitation was complete after a few minutes in a pure Fe-N alloy. The effect of varying the manganese content and the details of the precipitation process were not mentioned in these papers. Fast&apos; postulated that manganese causes a local lowering of the free energy of the lattice with a resulting segregation of nitrogen atoms to these low energy sites. The segregated nitrogen atoms are bound so tightly to the manganese atoms that they cannot form a precipitate. The internal friction measurements of Dijkstra and Sladek tended to confirm the concept of segregation of nitrogen around manganese atoms, and the increase in free energy on transferring a mole of nitrogen atoms from a segregated to a "normal" lattice site was computed to be - 2800 cal. Dijkstra and Sladek9 distinguished between two types of precipitates: ortho, a nitride of appreciably different manganese content than that of the matrix, and para, a nitride with a manganese content essentially that of the matrix. With each type of precipitate a solubility, again designated ortho or para, can be associated. Since the internal friction maximum in alloys which were aged several hours at 600" C dropped almost to zero, Dijkstra and Sladek9 concluded that the ortho solubility must be very low. The effect of temperature on the ortho and para solubilities has no1: been investigated. There are obviously several gaps in our knowledge concerning the influence of manganese on the behavior of nitrogen in a-iron. It was the purpose of the experiments described in this paper to determine the following: 1) The ortho and para solubilities of nitrogen as a function of temperature. 2) The details of the precipitation process at elevated temperatures. 3) The effect of varying the manganese concentration on the above phenomena. EXPERIMENTAL PROCEDURE Internal friction is conveniently employed in studying the precipitation of nitrides and/or carbides from a -iron because it is one of the few parameters, perhaps the only one, which is not affected by the presence of the precipitate itself. For this reason, internal friction techniques were heavily relied upon in the present experiment. A) Preparat of -. All specimens were prepared from electrolytic iron and electrolytic manganese. Alloys containing 0.15, 0.33, 0.65, and 0.75 wt pct Mn were vacuum melted and cast into 25 lb ingots. After being hot rolled to 3/4 in. bars, the ingots were swaged and drawn to 0.030 in. wires. The wires wen? decarburized and denitrided by annealing at 750° C for 17 hr in flowing hydrogen saturated with warer vapor. To obtain a medium grain size, - 0.1 mm, the wires were then heated to 945oC, allowed to soak for 1 hr, furnace cooled to 750°C, and water quenched. Subsequent internal friction measurements showed that this procedure reduced the nitrogen and carbon concentrations of the alloys to less than 0.001 wt pct. The wires were nitrided by sealing them in pyrex capsules containing anhydrous ammonia and annealing them for 24 hr at 580°C, the nitrogen being retained in solid solution by quenching the capsule into water. Immediately after quenching, the wires were stored in liquid nitrogen to prevent any precipitation of nitrides. By varying the pressure of ammonia in the capsule, it was possible to produce any desired nitrogen concentration. B) Internal Friction. The internal Friction measurements were made on a torsional pendulum of the Ke type,&apos;&apos; a frequency OF 1. or 2 cps being used. For

Jan 1, 1962

By Irving Cadoff, Kurt Komarek, Edward Miller

THE lead-telluride phase diagram as compiled by ansen&apos; has one intermetallic compound, PbTe, with solid-solubility limits extending from 22 to 45 wt pct Te (31.4 to 57.1 at. pct). Pelzel2 redetermined the solubility limits on the lead-rich side of the compound by measuring the relative areas of the two phases in samples annealed in the liquid + compound region, and obtained data that agreed with that of Hansen. Recent studies of the semiconducting properties of PbTe3 however, have indicated that addition of either lead or tellurium to PbTe does not produce samples having carrier concentrations greater than about 10 19 cm-3. Assuming that each impurity atom results in one ionized charge carrier, this carrier concentration corresponds to a solubility range of less than 0.04 at. pct. Fritts and Karrar4 have also observed two phases in samples containing from 0.15 to 3.0 pct excess Pb. To resolve the discrepancy between these data, the solid solubility limits of the compound PbTe were reinvestigated, using several other techniques. Ten 25-g samples with compositions of 1/2, 1, 3, 5, and 10 wt pct excess Pb and Te and two 100-g samples of 15 and 65 wt pct Te were prepared, sealed in evacuated quartz capsules, and melted at 950oC. The 1 pct excess Pb and Te and the 15 and 65 wt pct Te alloys were homogenized for 96 hr at 450°C; the remaining specimens were annealed 50°C below the solidus temperature reported in Hansen. The specimens were polished and etched with dilute HNO,. All the samples were two phase, the second phase appearing at grain boundaries or as spheroidized inclusions in the grains. In the case of the tellurium rich specimens, the second phase could be identified as a eutectic. Debye-Scherrer powder diffraction patterns of the i, 1, 3, 5, and 10 pct excess samples were obtained using CuKa filtered radiation. The films were indexed as the B1 or NaCl structure. The calculated lattice parameter of PbTe was the same for all the specimens, with an average of a0 = 6.460 ± 0.0005A. Lines due to excess lead or tellurium were very faint but could be clearly seen in all the films. The relative areas of the two phases in the 15 and 65 wt pct Te samples were corrected for the differences in density of Pb, Te, and PbTe, and taken as proportional to the relative volumes of the two phases. The corrected relative volumes agreed with a phase diagram having negligible solubility of both lead and tellurium in lead tellurium. The density of the two specimens was obtained to give a third independent determination of the phase limits. The specimens were broken into small pieces to eliminate any internal porosity, and adherent air bubbles were removed by vigorous shaking under water for several hours. The density of lead (11.34 g per cm3) and of tellurium (6.25 g per cm3) was taken from the International Critical Tables and that of stoichiometric PbTe (8.25 g per cm3) calculated from the lattice parameter. For the 15 wt pct Te samples, the density assuming no solubility was calculated to be 9.89 g per cm3; the observed density was 9.86 g per cm8. For the 65 wt pct Te alloy, the density calculated for the case of no solubility was 7.24 g per cm3; the measured density wag 7.22 g per cm3. These results are consistent with those of planimetry and X-ray. SUMMARY Data have been presented which show that the solid solubility of both lead and tellurium in PbTe is

Jan 1, 1961

By E. S. Anolick, Joseph Singer

The magnetic after-effect, in the form of time decay of permeability (l/µ), has been used to obtain independent data on the solubility of&apos; carbon in pure iron. The results differ slightly from the solubilities obtained through internal -friction experiments. The mechanism of the time decay being the magnetic analogue of internal friction, it appears necessary to explain the observed discrepancies, and some suggestions are made. THE purpose of this work was to establish the reliability of the time decay of permeability as a measure of dissolved interstitial in the body-centered-cubic lattice, and further, to establish techniques suitable for such measurements with samples of a form convenient for various other measurements. This particular sample form was the "epstein" strip, i.e., 25 cm by 3 cm by 0.035 cms. While the general conclusions are not limited by this geometry, there are two aspects of the samples that may be relevant: a) texture, and b) grain size, which ranged generally from 1 to 10 mm. Effects due to these two possible parameters were thought to be small and were not investigated. Calibrating the time decay technique and the other experimental procedures by measuring the solubility of carbon in pure iron suggested itself because of the accumulation of data on this subject through the related technique of the internal friction&apos;; The essentials of this report have been published in a somewhat condensed form.5 Since then some additional data have been taken. It is the purpose of this paper to add these data and to provide more complete description of the experimental techniques than was possible then. The time decay of permeability is a form of the magnetic after-effect.6,7 Permeability is measured immediately after demagnetization (µº) and also long after (µ8). The time decay, ?(l/µ), is the difference 1/µ8 - 1/µ0 and has been considered to be proportional to the concentration of a dissolved interstitial6,8,9,12 The relaxation time! 7, is characteristic of a particular diffusion, and affects the permeability as follows: where he being the activation energy for diffusion of the interstitial. Reference is made to the works of snoek6 and Neel7 for a complete discussion. A brief exposition of the principal ideas may be in order here, as follows. On the basis of the simplest model, resistance against domain wall motion increases with time after demagnetization due to the diffusion of interstitials to preferred sites created by the magnetostrictive tetragonolization under the strong magnetic field. Thus permeability, measured with a small field after the abrupt removal of the large field, will continuously fall with the process of diffusion of the interstitials, and the total "time decay" (or "time decrease") should be proportional to the amount of interstitial in solid solution. In the present report, use is made only of the total time decay rather than of the nature of the progress of the decay; the only problem involved in this approach is the need for assurance that only one interstitial is contributing to the effect as used. It is felt that the special use made here of the total time decay warrants the conclusions. Other work has been going on elsewhere on the detailed study of the nature of the time decay.8&apos;9 EXPERIMENTAL A) Materials—12-lb ingots were made in a vacuum furnace from high-purity iron and by a series of hot and warm reductions were brought down to 14-mil strip. This was cut into 25-cm epsteins. The pure iron was decarburized at 750 °C by hydrogen. The time decay measurement for the decarburized iron (see below) indicated a satisfactorily low starting level of interstitials. Grain growth to grain sizes of 1 to 10 mm was produced in the pure iron by a 700" anneal in hydrogen after a critical reduction of about 7 pct as suggested by previous trials. No measurement of texture in the pure iron was made. The grain size was taken to be large enough to be beyond any sharply critical size regarding solubility. B) Infusion of Carbon—Since the permeability measurement required about two dozen epsteins, a special method of carbon infusion had to be developed to ensure homogeneity of infused carbon. This was accomplished by setting the samples edgewise in a

Jan 1, 1961

By K. K. Rao, M. E. Nicholson

IN a recent paper on the solubility of carbon in Ni-Cu alloys,&apos; Nicholson reported that the carbon solubility appeared to be limited by the electronic structure of the alloys and that the solubility was propor- tional to the density of states at the Fermi surface of the "3d" band. In order to determine whether or not this relationship is a general one, the study has been extended to other systems. One of these is the CO-Ni system. This system was chosen for study because 1) it is adjacent to Ni-Cu system in the periodic table, 2) the alloys of cobalt and nickel do not form stable carbides at 1000°C, and 3) the system exhibits no pronounced change (less than 1 pct) in lattice parameter from cobalt to nickel,&apos; see Fig. 1. The alloys were prepared from -200-mesh 99.92 pct Ni furnished by the International Nickel Co. and spectroscopically pure -200-mesh cobalt sponge obtained from Johns on-Matthey and Co. The powders were weighed, mixed, and pressed using a pressure of 40,000 psi. The samples were sintered for 24 hr at 900° C in dry hydrogen and then cold rolled to a thickness of 0.010 in. (Samples containing more than 70 wt pct Co had to be annealed after cold reductions of about 20 pct to avoid cracking.) The alloys were carburized at 1000°C in an alun-durn boat containing Acheson AUC powdered graphite in an atmosphere of carbon monoxide cylinder gas which had been passed through Ascarite to remove carbon dioxide. After carburizing for either 12 or 24 hr the samples were moved to a cool zone of the furnace to cool. Cooling required about 5 min. The samples were filed to remove about 0.005 in. from the surface and then analyzed using a conventional combustion-gravimetric analytic method. The results of this investigation are shown in the lower portion of Fig. 1 and Tabulated in Table I. The nickel percentages for these alloys were determined from the weights of the powders from which the alloys were made. The solubility of carbon in pure cobalt at 1000°C (fcc) is 0.25 wt pct. This is somewhat less than the 0.33 wt pct measured by Boecker.3 From his data, it would appear Boecker&apos;s samples were supersaturated before the equilibrating treatment and that a 20-min treatment was not long enough to achieve equilibrium. Since in the present investigation the 12 and 24 treatments produced essentially equal carbon contents, our figure is presumed to represent more nearly the equilibrium value. The solubility of carbon in nickel of 0.19 wt pct agrees quite well with the value of 0.21 wt pct

Jan 1, 1963