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By J. J. Casey, R. S. Busk, D. L. Leman

The mechanical properties and the phase-temperature relationships of magnesium-lithium-aluminum, magnesium-lithium-zinc, and magnesium-lithium-aluminum-zinc alloys are presented. It is shown that the strengths of some of these alloys can be tripled by heat treating and aging but that the high strength is accompanied by brittleness and lack of stability. MAGNESIUM-LITHIUM alloys are interesting because the hexagonal structure of magnesium can be changed into the more ductile and workable body-centered cubic structure by the addition of only 10 pct lithium.1 After Loonam2 determined that the ductility of the magnesium-lithium alloys is destroyed by even trace quantities of sodium, it was possible to prepare ductile alioys and study their mechanical properties. Jackson3 and coworkers have ably described their results with both the binary and polynary magnesium-lithium alloys. The present paper is a report of the mechanical properties of wrought magnesium-lithium-aluminurn, magnesium-lithium-zinc, and magnesium-lithium-aluminum-zinc alloys. Experimental Procedure The alloys were made by melting 7 lb of magnesium in a steel crucible under a molten flux of 75 pct LiCl + 25 pct LiF3 previously dried at 1600°F. The metal was heated to 1300°F for alloying aluminum and zinc and then cooled to 1250°F for adding lithium (sodium content: 0.01 to 0.04 pct). After the lithium had been alloyed, dry LiF powder was sprinkled on the melt to thicken the excess flux. After a 15 min holding period, the flux was skimmed off. Paraffin vapors were used to protect the metal as it was poured into 3 in. diam by 11 in. long magnesium molds. Billets were extruded into 3/8-in. diam rod and 1/16 x 3/4-in. strip; sheet was prepared from 1/8x 7/8-in. extruded strip, first by hot rolling and then by cold rolling 2 pct per pass. Hot rollability, or the temperature range through which the alloys could be rolled, was determined by rolling the strip two passes at 30 pct reduction per pass. Cold rollability, or the percentage of cold reduction to cause detrimental cracking, was determined on 0.064-in. hot-rolled sheet. Standard ASTM procedures were followed to obtain the mechanical properties, using 0.2 pct offset for the yield strengths. The level of properties produced with commercial magnesium alloys using the procedure described is given in a previous publication." Results Magnesium - Lithium - Aluminum System: The good workability of magnesium-lithium alloys is illustrated by the data in table I, which shows approximately how the extrusion pressure decreases as the amount of cubic phase increases. Binary magnesium-lithium alloys are soft and weak, but aluminum raises their strengths con-

Jan 1, 1951

By R. P. O’Shea

AS part of a program to investigate phase relationships in ternary systems containing semiconducting elements and compounds, an X-ray survey of alloys in the system Pb-Te-Bi was carried out which revealed that the pair PbTe-Bi probably formed a quasibinary system. This portion of the ternary was examined fully by thermal analysis, X-ray, and microscopic procedures. More than twenty-five alloys were made by melting together the proper combinations of the elements in quartz tubes at a temperature of 1000o C. The compositions which were high in PbTe content were prepared in evacuated, sealed, quartz tubes, allowed to freeze in vacuum, and then transferred to Crucibles for thermal analysis. The bismuth-rich alloys were made directly in the crucibles under a protective gas cover. The starting elements were obtained from the American Smelting and Refining Co. and had a purity greater than 99.99 pct. Cooling curves were made with about 50 g of alloy in graphite crucibles which were placed in a covered quartz container in which a protective atmosphere of helium or nitrogen was maintained. The thermocouple protection tube was made of mul-lite and served as a stirring rod when connected to a reciprocating stirring motor. Temperatures were

Jan 1, 1962

By C. G. Dunn, W. W. Martin

Various X ray diffraction techniques have been developed for determining orientations of crystals. Transmission and back-reflection Laue methods1,2,3 in particular have been found to be very useful. In many applications, however, a need still exists for a simplification of the details of analysis so that orientations can be determined more rapidly, and operators with little or no training in crystallography can be easily instructed in the necessary procedures. Filling this need in part is a system of analysis that has been in use for several years in the Metals Section of the Laboratory, General Electric Co., Pittsfield, Mass. This method is an extension of the one of Majima and Togino,4 which uses fifty-five indexed standard transmission Lauemams to cover the entire range of crystal orientations. Greater rapidity in determining the orientation of a crystal is achieved by using a slightly larger set of standard Lauegrams and by systematizing the steps followed. This system of analysis in its present form constitutes the subject matter of the present paper. Although the method uses transmission X ray Laue patterns, many of the operations are applicable to the analysis of back-reflection Lauegrams. Very briefly the method is based on the fact that Bragg and azimuthal angles for two or more identified Laue diffraction spots completely determine the orientation of a crystal, and standard Laue photographs of the Majima and Togino type provide a ready means for identifying Laue spots. After a preliminary analysis, which consists of identifying the pattern in terms of one of the standard patterns, a correctly marked Lauegram is selected to carry out the final analysis. From measured Bragg and azimuthal angles, the normals (poles) of the reflecting planes for the selected Laue spots are located on a stereographic projection and the orient.at.ion of the crystal determined (in terms of the positions of cube poles for example) by a few additional operations on a \I1ulff net. The actual procedure followed will be illustrated in detail after a brief treatment or X ray problems connected with orientation determinations and a discussion of useful principles and techniques. X Ray and Plotting Problems RELATIVE ADVANTAGES OF BACK-REFLECTION AND TRANSMISSION LAUE DIFFRACTION METHODS In the back-reflection Laue method it is easy to make a direct measurement of the tilt of a zone of Laue spots. Specimen thickness also is generally no problem. In the case of thin specimens, however, the time of exposure is less for the transmission Laue method. For example, with silicon ferrite this advantage holds up to a thickness of 0.6 mm if tungsten radiation at 40 kv peak voltage is used. Furthermore, the transmission photograph represents the entire thickness of the crystal whereas the back-reflection picture represents only a surface layer (a thickness of less than 0.05 mm for silicon ferrite), the deeper layers supplying a negligible amount to the diffraction pattern. If a crystal is distorted, the transmission Laue pattern gives more information particularly with regard to the orientation of the crystal and the overall amount of spread in orientation. Finally a 3 1/4 X 4 in. film suffices for transmission photograms; a 4 X 5 in. or 5 X 7 in. film with a central hole is needed for back-reflection patterns Not only is the manipulation of film easier because of the size of film, but

Jan 1, 1950

By G. R. Speich, R. A. Oriani

The rate of coarsening (Ostwald ripening) of copper precipitate Particles in an a-iron matrix has been studied in three Fe-Cu alloys containing 2.3, 4.0, and 5.4 wt pct Cu in the temperature range 730° to 830°C. The kinetics of Ostwald ripening ohevs the law as deduced by Lifshitz and Wagner so that matrix diffusion controls the process. The equilibrium shape of the particles is that of rods with hemispherical caps, and is due to the BOTH Lifshitz and slyozov and wagner3 have recently solved the complex problem of determining the rate of coarsening of an assembly of precipitate particles. Both workers, although using somewhat different mathematical approaches, arrive at the same equation for the relation between the mean radius 7 of a distribution of spherical particles and the bulk solubility co, diffusivity D in the matrix, large anisotropy in the copper/a iron interfacial free energy (y,/y, = 4). A Lifshitz -Wagner equation, modified for both the nonsphericity of the diffusion field and the reformulation of the Gibbs-Thomson equation for the case of nonsphencal particles, is derived. Using the modified equation, the calculated values of the mean radius are smaller than those observed by a factor of about one -half. absolute temperature T, gas constant R, surface free energy y of the precipitate/matrix interface, and the time t of aging: where V, is the molar volume of the precipitate and v is a stoichiometric factor. Wagner's treatment is restricted to the case of spherical particles growing in a liquid matrix where strain energy is not a factor. Lifshitz and slyozov have discussed the effect of anisotropic particle shapes and internal strains on the rate of coarsening. Oriani4

Jan 1, 1965

By A. K. Mukherjee, John E. Dorn

The effects of strain rate and temperature on the critical resolved shear stress for (321)[111] slip were determined for the silver-rich CsCl type of intermetallic compound AgMg. The flow stress increased only slightly as the test temperature was first decreased below room temperature; a rapid increase in the flow stress was obtained with yet further decrease in temperature from about 250' to 4OK. The effect of both the temperature and strain rate on the flow stress over the low-temperature range could be rationalized satisfactorily in terms of the Peiwls mechanism when the deformation was controlled by rate of nucleation of pairs of kinks. IN spite of the well-documented interest of metallurgists and engineers concerning the mechanical behavior of intermediate phases and intermetallic compounds, very little definitive information is available on this important subject. As recently summarized by westbrook,1 only limited and modest progress has been made thus far in elucidating the details of the mechanism of plastic deformation for the intermetallic compounds; the available information largely concerns phenom-enological descriptions of empirical observations and experimental facts, principally with reference to poly crystalline aggregates. The present report on the elucidation of the rate-controlling mechanism for slip in the bcc structure of AgMg is part of the comprehensive program of a systematic investigation on the mechanical behavior of intermetallic compounds. The intermetallic compound AgMg has a CsCl type of lattice structureZ and is completely ordered up to its melting point of 820c3 This material was selected for our present investigation because of its simple crystal structure, moderate and congruent melting point, ordered structure persisting up to the melting point, and some solubilities of the constituent elements, which could be expected to help in the growing of single-crystal specimens. Whereas previous investigations on the properties of AgMg include hardness,*"6 slip systems,7 tensile flow stress of polycrystalline specimens,' and electrical resistivity,4 the current investigation will be directed principally at elucidating the rate-controlling dislocation mechanism responsible for slip in single crystals at low temperatures. It will be shown that the strain rate is consistent with a model involving the rate of nucleation of pairs of kinks by the Peierls mechanism for plastic deformation. EXPERIMENTAL TECHNIQUES Single-crystal specimens of the AgMg intermetallic compound were prepared and tested as follows: 1) A master alloy ingot of AgMg was produced by melting and chill casting the high-purity silver (99.995 wt pct) and high-purity magnesium (99.997 wt pct) in an induction furnace under a helium atmosphere. 2) Sections of the above-mentioned ingot were placed in a graphite mold containing a spherical cavity in which a single-crystal sphere of 1 in. diam was grown under helium by the Bridgman technique. 3) The operative slip systems were investigated at room temperature on a singlecrystal of AgMg, from measurements of the angles made by the slip traces on the two surfaces of the specimen, which were at 90 deg to each other. The investigation confirmed that the AgMg compound undergoes slip primarily in the (321) plane and in the [Ill] direction, but a small amount of secondary slip was also noticed in the (112) (111) system. No slip was observed in the (110) planes. 4) The single-crystal sphere, mentioned earlier, was oriented in a graphite mold containing a cylindrical cavity of 3/8 in. diam above the spherical receptable to give the angle xo = 45 deg between the axis of the cylindrical bar and the normal to the slip plane (321), and the angle AO = 45 deg between the slip direction [Ill] and the axis of the bar. Oriented single-crystal bar seeds were produced by melting, under a helium atmosphere, a polycrystalline bar in the cylindrical cavity above the oriented spherical seed and by growing an oriented single-crystal bar from the seed. 5) Finally, oriented single-crystal specimens were grown from these cylindrical seeds. 6) The cylindrical specimens were machined in a High-Tension Spark Cutting unit to give a 2-in. gage length. The spark-machined gage section was elec-tropolished in a bath containing 200 ml of H3PO4,

Jan 1, 1964

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 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 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 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 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