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By R. S. Craig, W. E. Wallace, G. S. Kamath

The destruction of long-range order in Mg-Cd, has previously been thought to occur as a second order process. In the present work a variety of X-ray diffraction techniques are employed to show that in this case and in Mg-Cd alloys containing frorn 20 to 33 1/3 at. pct Mg the process more nearly resembles an ordinary first order phase transformation. It is found that the ordered alloy, which occurs as a single phase at lgw temperatures, transforms with increasing temperature into a Cd-enriched disordered (P) phase and an Mg-enriched ordered (a) phase. The interval over which these two phases coexist under apparently equilibrium conditiovs extends from 25" to 87°C content. At 33 1/3 at. pct Mg the two-phase region has disappeared. Measurement of the long range order parameter, s, for the ordered phase in MgCd, (from superlattice line intensities) shows that it disorders only very slightly before it is transformed into the p form. Kinetic experiments reveal that when the a phase is formed frorn the p phase by suitably quenching so as to suppress phase segregation, the transformation is essentially complete and the a phase has acquired its equilibrium value of S in less than 1 min. In an earlier investigation carried out in this Laboratory1 the rate of disordering of the MgCd3 superlattice was determined over an extensive range of temperature. The process was found to be first order and the rate constant, k, was established. At the lowest temperatures studied k was found to increase exponentially with temperature. At higher temperatures, however, it passed through a maximum and for the 100 deg interval below the point at which the superlattice was destroyed k was observed to decrease with increasing temperature. The various theoretical treatments1-3 of order-disorder phenomena lead to a negative temperature coefficient for k near the superlattice critical temperature but this region for MgCdB seemed unusually wide. Detailed analysis of the point indicated' that either the long-range order parameter, S, for MgCda varied with temperature in an anomalous way or the disordering in this instance is more complex than in the cases which have been treated theoretically. The original objective in the present work was simply to determine S vs temperature for MgCda to ascertain whether the disordering was in conformity with theoretical expectations.2° In an early stage of the work it was noted, however, that under equilibrium conditions ordered (a) and disordered (0) phases coexist in MgCd, over a considerable range of temperature. Since on thermo- dynamic grounds the coexisting phases must differ in composition, the investigation was enlarged to permit the compositions of the two phases in MgCd3 to be determined and to establish the extent of the two phase regions for alloys containing from 20 to 33 1/3 at. pct Cd.

Jan 1, 1963

By G. P. Conard, R. C. Hall, J. F. Libsch

The 50 pct Fe-50 pct Co alloy undergoes a transformation from disorder to an ordered structure of the CsCl type reportedly in the vicinity of 732OC. During this process, the coercive force goes through a maximum, apparently as a result of strains associated with the coherent nucleation and growth reaction. This magnetic alloy also shows a marked increase in the ratio of residual to saturation induction, which is associated with annealing to a high degree of order with the continuous application of a magnetic field. The increase in ratio can be explained on the basis of a decrease in 90' domain boundaries and, perhaps, by an increase in anisotropy resulting from lattice distortion. THE 50 pct Fe-50 pct Co alloy undergoes a disorder-order transformation which has been reported to occur in the vicinity of 732°C1,2 The ordered structure is the CsCl type.' This magnetic alloy also shows a marked increase in the ratio of residual to saturation induction as a result of heat treatment in a magnetic field, sometimes called a response to magnetic anneal.'-' The purpose of this investigation was to study the course of the ordering reaction, the nature of the response to .heat treatment in a magnetic field, and the relation, if any, between ordering and the response. Procedure The method of approach in this investigation was to produce an initial structure as completely disordered as possible and then gradually to order the alloy by isothermal anneals at various temperatures under different conditions of the applied magnetic field. Magnetic, magnetostriction, and X-ray analyses were of primary importance in determining the property and structural changes resulting from the isothermal anneals. Rings of the 50 pct Fe-50 pct Co alloy were prepared from the elemental powders by a powder metallurgy technique, further details of which may be found in ref. 7. The initial structure was produced by annealing the specimens for ½ hr at 1000°C, cooling to and holding for ½ hr at 900°C (in the a range above the ordering temperature), and water quenching. Isothermal anneals were performed at 600°, 675°, 720°, and 740°C. For example, rings were heated to 600°C, held for a predetermined period of time, and cooled by natural cooling at a rate slightly slower than an air cool (average of 20" to 25°C per min). The tests (magnetic, etc.) were made after each heat treatment. All high temperature treatments were performed in a purified hydrogen atmosphere. The treatments at the various temperatures were carried out under one or more conditions of an applied field including 1—no field, 2—field of 20 oersteds applied on cooling only, and 3—field of 20 oersteds applied continuously during heating, holding, and cooling. Magnetic measurements were made using the standard Rowland ring technique8 with a maximum field strength of 100 oersteds. The magnetization curve, induction at 100 oersteds (B.), residual induction (Bt), and coercive force (Hc) were determined. All magnetic analysis data were based on an average of the results from three rings. A strain gage technique9 as used for the measurement of magnetostriction. The X-ray determination of the relative amount of ordered phase present was made on the ring specimen used for magnetic measurement. This was done by the back-reflection method using a rotating specimen (because of the large grain size) with unfiltered CoKa radiation and a 7 hr exposure time. Intensity measurements of the ordered line (300) were made by comparing visually the films so obtained with standard films prepared by exposing for different lengths of time a specimen given a long time anneal (high degree of order). Results In all instances the saturation induction (induction at 100 oersteds) was found to increase slightly with annealing time. This effect was small and appears to be the increase in saturation induction to be expected on ordering.10-13 The residual induction behavior was markedly influenced by the field condition during annealing, Figs. 1, 2. For the condition of no applied field, the ratio of residual to saturation induction remained essentially constant for short annealing times but showed a significant increase at longer times. With increasing annealing temperature, less time was required to produce this increase in the ratio. In the case of the 600°C anneals, the increase did not occur until approximately 20 hr, Fig. I, while on annealing at 740°C the increase was immediate, Fig. 2. Slight decreases in the ratio may be observed at 100 hr for specimens treated at 720°C and at 1 hr for those treated at 740°C. Specimens annealed in a field of 20 oersteds showed a residual to saturation induction ratio consistently higher than that for the specimens annealed without the field. The first anneal with the field (¼ hr) caused an abrupt increase in the ratio at all temperatures; thereafter, the increase in the ratio was generally similar for specimens annealed

Jan 1, 1956

By F. R. Brotzen, D. L. Davidson

ALTHOUGH much effort has been devoted to the problem, the nature of plastic flow in bcc metals is still not fully understood. Some of the difficulties encountered stem from the scarcity of information on orientation effects and the uncertainty as to the glide system in these metals. In an attempt to provide some of this information, the present investigation was carried out, using zone-refined molybdenum single crystals deformed in tension and in direct shear. Considerable evidence exists to indicate that the primary slip system for molybdenum is (110) <111>.1,2 Figs. 1 and 2 indicate that this is, indeed, the case for direct shear at room temperature (see Ref. 3 for a description of the shear-testing equipment and procedure). Moreover, from Fig. 2, the shear stress on the (123) is lower than that on the (112). This would be the case if, as has been suggested,4,5 slip on these two planes is in reality com-

Jan 1, 1965

By L. F. Mondolfo, B. E. Sundquist

The crystallographic orientation relationships resulting when lead is nucleated from the liquid by Ni, Cu, Ag, and Ge were determined. For each nucleating agent several definite orientatioz relationships were found. These relationships seemed to be controlled by good symmetry relations and low crystallographic disregistry between mating planes. For any given nucleating agent the under colling for nucleation was found fairly constant and independent of the orientation relationship and consequent disregistry. It was also found that, upon re melting and refreezing the Pb, the orientation relationship was changed. These findings prove that crystallographic disregistry is not the controlling factor in heterogeneous nucleation from the liquid. The results of this investigation tend to confirm the theory presented in a preceding paper that heterogeneous nucleation starts with the formation of an adsorbed layer of nucleated metal on the nucleat-ing impurity. Evidence is given that cavities in the nucleating agent act as centers of nucleation. IT has long been known&apos; that solid extraneous particles are active in catalyzing phase transformations that occur in a system, particularly condensation and crystallization. It is well established that these heterogeneities act as catalysts by providing surfaces upon which nuclei of the precipitating phase can form with activation energies smaller than those required for homogeneous nucleation. Numerous investigations have shown that in this process of heterogeneous nucleation: a) the nucleus forms with one, or several, definite crystallographic orientation relationships with the nucleating phase2-4 and b) that there is a small range of undercoolings or super saturations characteristic of the nucleation of a given solid on a given Substrate.5-10 Turnbull and vonnegut11 have developed a theory based on theories developed by Volmer12 and Turn-bull and Fisher1= for heterogeneous nucleation from gases and liquids, that relates the super saturation or undercooling required for nucleation to the dis-registry between the lattices of the nucleus and the nucleating agent. This theory predicts that nucleation should occur with the orientation relationship between the nucleus and nucleating agent that minimizes the disregistry. Further, it predicts that the undercooling or super saturation necessary for nucleation should be a function of the disregistry. Numerous investigations have dealt with the orientation relationships resulting from the condensation of vapors onto crystalline solid substrates2,3 and a few with the nucleation of one phase by a second phase in solidification4,14. Others have dealt with the supersaturation8-10 and undercooling5-7 associated with nucleation in condensation and solidification respectively. However, there is virtually no report that gives both of these factors for the same system. In this investigation a study was made of the undercoolings and orientation relationships resulting when Pb is nucleated from the liquid by Ni, Cu, Ag, and Ge. It was the purpose of this investigation to check the Turnbull-Vonnegut theory, i.e., the importance of crystallographic disregistry between nucleating catalyst and nucleated metal. The results indicate that disregistry is not an important factor in nucleation and that the nucleation process is probably somewhat more complex than current theories suggest. EXPERMENTAL PROCEDURE Small single crystals of nickel, copper, silver, and germanium were prepared from materials of four to five nines purity, and the Pb used was also 99.999+ pet pure. Cu and Ag single crystals were prepared by sealing small chips of Cu or Ag in an evacuated quartz capsule and heating the capsule at 2000°F for 1 hr before cooling. Nickel crystals of 200 diam were also prepared in evacuated quartz capsules, but melting was done by heating the capsules in an oxy-acetylene flame for a few minutes. These spheres were invariably polycrystalline so

Jan 1, 1962

By R. H. Hiltz, R. W. Douglass

Large-grain specimens of iodide titanium prepared metal-lographically were stain etched using the technique of New York University as modified by Watertown Arsenal Laboratories. Orientations of grains showing an abnormal response to stain etching were determined by back-reflection X-ray techniques. X-ray results showed that a-titanium grains, oriented such that their basal planes were parallel to the polished surface, have a slower reactivity to stain etching than any other orientation. Grains oriented such that their basal planes made small angles with the polished surface exhibited reaction rates between that of the true basal orientation and the normal rate. As the angle with the surface increases the reactivity increases up to a maximum angle of 14 deg. Basal planes making an angle with the surface of greater than 14 deg show normal reactivity to stain etching. These data, plus work on polycrystalline specimens, indicate that the oxidation rate of titanium might be reduced by correct alignment of crystal planes. ThE earliest observed phenomenon in the micrography of titanium was the occurrence of a grains which exhibited no birefringency. Even though a titanium possesses a hexagonal-close-packed structure and should be optically anisotropic, it can appear optically isotropic, when the crystal is so oriented that its intersection with the polished surface is perpendicular to the axis of symmetry. During the develoment of color staining techniques for titanium: it was noted that the a phase did not stain uniformly. Staining of specimens of commercially pure titanium resulted in a micro-structure of blue-colored grains, among which were randomly oriented a few grains exhibiting colors between the initial yellow and the final blue color. Under polarized light, these off-colored grains appeared to be optically isotropic. More intense study with polarized light, however, showed only the yellow grains to be isotropic. The intermediate colored grains exhibited weak birefringency, which at first glance is not detectable due to the strong birefringency of the surrounding structure. Aminoff2 and Straumanis3 have previously shown that in zinc the basal orientation is less reactive than other orientations. Leidheiser4 showed the same thing to be true for the hexagonal-close -packed configuration of cobalt. Thus, it would appear that the basal-plane orientation is responsible for the abnormality in stained titanium. Later work by Leidheiser5 on cobalt has shown that other planes besides the basal plane exhibited reaction rates slower than the normal rate. Reported planes were the (100), the (1011) and several planes making small angles with the basal plane. The last observance appears not to be due to the reactivity of the reported plane but due to the reactivity of the basal plane which is very close to being parallel to the surface. There are, then, two possible explanations for the range of colors observed in stain-etched a titanium. Since extensive research on the effect of orientation on surface reaction has failed to provide any definite trend, these two explanations are not adequate in themselves. Planes showing the retarded reactivity differ not only for different crystal structures but for metals of the same crystal structure and even for different reactions of the same metal. Tragert,6 Leidheiser,7 Tammann,6 and wagner9 have shown copper to exhibit retarded reactivity on the (111) plane. However, Leidheiser10 has shown that under other conditions the retarded reactivity of copper is associated with the (100)

Jan 1, 1960

By P. R. Gage, R. W. Bartlett

At high temperatues and reduced oxygen pressuves, molybdenum silicicles oxidize to form SiO(g) vathev than a passivating SiO2 film. This is a sevious problem for low-pressure applications of sili-cide-coated rejvactory tnetals. As oxiclalion pvoceeds, the higher oxidntion potential of silicon initially suppresses oxidation of molybdenum and silicon is depleted near the surface. Eventually, silicon depletion causes secondary passivation (SiO2) mid an active to passiue transition in oxidntion behavior at moderately low pressures. At very low pressures, silicon depletion continues until steady-state oxidation of both silicon and rnolyhdenum can proceed at rates propovtional to their initial stoicliiorrzetvies. Ohsel-vations of &apos;oxidatiou-rate belzcriliov at temperatures from 1101° to 1700°C and oxygen pressures from 10-7 to 1 at in were made using thermogravimetn&apos;c, X-ray diffraction, and elect?*on-rtric?ro/,vobe techniqrres. The experimental results are in good agreement with calculated lirnits for the various types of behavior that were based on tlievtnochernicnl atzd dffusion rate data for the Mo-Si system. A protective oxide film normally forms during the oxidation of elemental silicon. However, at elevated temperatures and reduced pressures this film is not present and oxidation is rapid. wagnerl has shown that this behavior occurs because SiO(g) is more stable than SiO2. The behavior of the silicides at high temperatures is similar but more complicated. Molybdenum and several silicide compounds are involved and the chemical activity of silicon, or the equivalent silicon vapor pressure, can vary considerably in these phases. Also, diffusion in the silicides resulting from reactions in which a higher silicide is reduced to a lower silicide, or molybdenum containing dissolved silicon, must be considered. EXPERIMENTAL AND ANALYTICAL METHODS Isothermal oxidation experiments were conducted using hot-pressed MoSi2 and Mo5Si3 wafers. Each experiment was conducted at a fixed oxygen pressure by controlling the flow of oxygen in and out of the furnace chamber. Most of the experiments were conducted in a molybdenum-wound tube furnace. Special construction features, including use of a high-purity dense A12O3 refractory tube, permitted operation in high vacuum to 1500°C. Higher temperatures could be obtained at higher working pressures and with a sacrifice of tube lifetime. Temperature was measured with three Pt 6 pct Rh/Pt 30 pct Rh thermocouples and controlled with a proportional controller. Weight changes were monitored with a quartz spring balance. The few experiments, at higher temperatures and low pressures, were conducted in a cold wall vacuum chamber using induction heating with direct susceptance of the sample. Because of the variable levitation effect of the induction coil on the sample, weight changes could not be monitored continuously in these experiments. All samples representing both experimental methods were weighed before and after oxidiation and examined with a microscope and analyzed by X-ray diffraction. The analytical treatment assumes equilibrium at interfaces and uses existing thermodynamic data. For convenience, equilibria are expressed as chemical potentials or as vapor pressures of the gases involved. The resulting differences in equilibrium vapor pressure between adjacent interfaces and between the surface and external atmosphere drive diffusion processes in the solid and gas boundary film, respectively. These diffusion processes govern the rate of oxidation. A bibliography of the sources of thermochemical data for the reactions pertinent to the Mo-Si-O system is given in Table I. Reactions are listed with the sources of free-energy data referenced. Since all of these reactions involve one or more gas species, usually oxygen, a plot of the product-gas chemical potential, -RT In Pg, vs absolute temperature for each reaction is given in Fig. 1. These curves can be used to determine equilibrium vapor pressures or to compare the stability of different compounds in the various gases: O2, Si, SiO, and Moo3,. The data of Table I and Fig. 1 were used to determine the analytical results presented in this paper. EXPERIMENTAL RESULTS The weight changes for three MoSi, samples oxidized at 1500°C are plotted against time in Fig. 2. The treatments were identical except for the indicated variance of oxygen pressure and each curve represents one of three characteristic types of behavior that were observed. The upper curve is an example of passive oxidation, which is defined as

Jan 1, 1965

By R. J. Pearce, I. Whittle, J. J. Stobbs

The oxidation kinetics of UCu5 in carbon dioxide have been studied over the temperature range 350° to 850°C. At any one temperature, two successive parabolic rate constants are obtained. Up to 650°C, the temperature dependence of the reaction is snzall and the first parabolic rate constant is higher than the second. At temperatures of 690°C and above, there is a very marked temperature dependence and the second rate constant is the higher. Metallo-graphic examination has shown that attack proceeds via internal oxidation. THE oxidation behavior of the U-Cu intermetallic compound, UCu5, in carbon dioxide was briefly examined during investigations aimed at selecting oxidation-resistant uranium-alloy coatings.&apos; It was found that an extremely thick, coherent scale was formed coated at the scale-gas interface with a thin, apparently continuous layer of copper. A more detailed investigation of the oxidation of UCus has now been made. Previous published work2 has been limited to a brief examination of UCu5 in connection with the oxidation of a U-1 wt pct Cu alloy in carbon dioxide. No attempt was made to explain its behavior, but the work indicated that attack proceeded by intern oxidation, with the formation of a layer of copper on the surface of the specimen. EXPERIMENTAL UCus was prepared by arc melting the stoichio-metric amounts of high-purity copper* and uran-

Jan 1, 1965

By Pol Duwez, Paul Pietrokowsky

Titanium and vanadium form a complete series of solid solutions at temperatures above 885°C. Below 885°C, vanadium is slightly soluble in a titanium (about 1.5 pct V at 650°C) and a two-phase a plus ß region extends to 18.5 pct V at 650°C. The alloys up to 15.0 pct V transform into a supersaturated a solid solution upon rapid cooling from the ß phase. Above 15 pct V the ß phase may be retained. The temperature at which the ß to a&apos; martensite-type transformation takes place decreases steadily from 850°C for 0.5 pct V to 325°C for 12.7 pct V. IN the transition series of the periodic table, the six elements that have a body-centered cubic structure at all temperatures, namely chromium, vanadium, molybdenum, columbium, tungsten, and tantalum, fulfill the conditions required for the existence of complete series of solid solutions with the high temperature (ß) form of titanium. Up to the present time, complete solubility at temperatures above the a to ß transformation of titanium has been observed in the systems Ti-MO,1,2 Ti-Cb,2 Ti-W,³ and Ti-Ta.³ The two remaining elements, vanadium and chromium, have atomic diameters which are not as near that of titanium as the atomic diameters of the four other elements. Both are smaller than titanium and the difference, expressed in percentage of the titanium atomic diameter, is 12.3 pct for chromium and 8.2 pct for vanadium. On this basis, the least favorable conditions for solid-solution formation are therefore encountered in the Ti-Cr system. The early work on the Ti-Cr system4 revealed that up to about 1100°C the solubility of chromium in ß titanium is limited, and an intermetallic phase exists near the stoichiometric composition TiCr2. It was found later&apos; that a continuous series of solid solutions extends across the diagram at higher temperatures, and TiCr2 forms by solid state reaction. It became apparent therefore that complete solubility would also be present in the Ti-V system. The present investigation confirms this possibility. The iodide-refined titanium used in this investigation was purchased from the New Jersey Zinc Co., New York. According to the manufacturer, a typical analysis of this product is 0.0065 pct Mn, 0.0022 pct Fe, 0.0015 pct Cu, and 0.0042 pct Pb. The Vickers hardness number (10 kg load) of the crystalline bar varies between 55 and 80. The vanadium metal was purchased from the Westinghouse Electric Corp., Lamp Div., Bloomfield, N. J. The purity of this metal is probably above 99 pct V. The oxygen content must have been quite low, since a bar about 1/4 in. square could be cold rolled into a ribbon 0.020 in. thick. To prepare the various alloys on the titanium-rich side of the diagram, the titanium rod was machined into cups 1 cm in diam and about 1.5 cm high. Pieces of vanadium, of the proper weight, were cut from the square bars and placed within the cups prior to melting. On the vanadium-rich end of the diagram, single pieces of the metals, in the proper weight pro-

Jan 1, 1953

By Harry C. Gatos, Frank J. Bachner, Mario D. Banus

The portion of the Nb-Sn phase diagram between 75 and 79 at. pct Nb at temperatures near the liquidus has been investigated by melting alloys of known composition and examining the microstmc-tzlres resulting when these melts were quenched. The Nb-Sn samples were contained in magnesia sleeves, and were mounted in a high-pressure apparatus. It was shown that Nb3Sn me1ts by a peri-tectic decomposition and that the composition of the compound Nb3Sn at the peritectic horizontal lies between 77 and 78 at. pct Nb. Analysis of the phases present in the samples was accomplished by anodic staining techniques and an electron-microbeam analyzer. DURING the past several years a great deal of work has been devoted to the Nb-Sn system, and the bulk of this work has been primarily concerned with the intermetallic phase Nb3Sn. However, the melting characteristics of Nb3Sn have not been investigated because of the difficulty in maintaining known compositions due to tin losses by vaporization, and the added difficulty of containing Nb-Sn alloys without sample-container reaction at the high temperatures involved. This paper reports on the development of a technique for containing Nb-Sn alloys under pressure high enough to prevent tin losses at temperatures in excess of the liquidus. Using this technique, melts of known composition were made and a phase diagram near the liquidus is proposed based on the microstructures resulting when these melts were quenched. EXPERIMENTAL The Nb-Sn alloys weighing 0.7 g each were prepared in the form of cylinders. Niobium powder, -200, +325 mesh, and tin shavings with a minute amount of tin powder for a binder were mixed in the desired proportions and cold-pressed at 50,000 psi in a 0.150-in. die. This cylinder was then enclosed as shown in Fig. 1, and the entire arrangement mounted in a 2.6-in. pyrophyllite tetrahedron.&apos; The desired temperatures (-2500°C) presented several problems in high-pressure container materials. Boron nitride, the usual sample holder, was found to react with niobium to form the nitride. However, MgO cans and plugs machined from refractory rods did not melt nor react with the sample. Thick tantalum contact disks were used to carry the large currents (~450 amps) to the graphite heater, since tantalum did not melt and had enough strength to prevent extrusion of the sample. It was further observed that the graphite heater and tantalum disks could not be in direct contact with the pyrophyllite because it melts and decomposes at these temperatures to kyanite and coesite (SiO2),2 the former reacting with graphite and the latter with tantalum to form the silicide. The use of a MgO outer sleeve

Jan 1, 1965

By M. van Swaay, C. E. Birchenall

Palladium has a large capacity for the dissolution or occlusion of hydrogen; the gas also diffuses very rapidly through the metal. Palladium thimbles are widely used in the laboratory for purification of hydrogen by diffusion. The use of palladium filters could be extended to several large-scale applications if greater stability of the permeation process could be attained. In this paper a number of aspects contributing toward this goal are discussed, and new data on solubility and diffusivity are given. Palladium sheets of 0.005, 0.010, or 0.020 in. thickness were prepared by three procedures: casting and rolling, sintering palladium powder, and sintering palladium powder to which 0.1 pct thorium oxide had been added. Half-inch diam circular specimens were cut and clamped in a stainless steel holder provided with knife edges to seal the palladium sheet with soft copper gaskets. The seals were tested for leaks at l0-5 mm Hg. No difficulties from this source were encountered even after many temperature cycles in the system. The holder was silver-soldered to Kovar-glass seals and attached to the system through glass ball joints. Each joint was sealed with Apiezon wax and was adjacent to a cold trap. In several experiments the holders were fused directly into the system, eliminating the ball joints, to show that the wax did not affect the filter. Three filter holders were supported in holes in a stainless steel cylinder, which also acted as a temperature equalizing block inside the resistance wound furnace. The measuring thermocouple was inserted in one of the holders. A potentiometric controller with control thermocouple on the furnace heater winding maintained ± 0.2oC. Electrolytic hydrogen was purified and stored as shown in the flow diagram, Fig. 1. It could be admitted to either or both sides of the palladium specimens. Only a mechanical pump was used, for even high pumping speeds would not bring the pressure on the off-gas side of the filter below 10 to 50p Hg owing to the rapid hydrogen flow through the filters. A series of steady-state permeability studies allowed identification of several conditions of permeability loss. A constant fore pressure was maintained while the low-pressure face was pumped. After enough time had elapsed for a steady flow to be established the pump was cut off, and the rate of pressure increase in a thermostated bulb was measured to determine the flow rate. The time delay method of Daynes1,9 was adopted since it offered the possibility of determining solubilities and diffusivities. For each measurement both sides of the filter were evacuated for 15 min, which gave about 10-4mm Hg residual pressure. Hydrogen was admitted to one side or the other of the filter at a fixed pressure. The pressure on the off-gas side was observed as a function of time. A

Jan 1, 1961

By C. C. Clark, J. S. Iwanski

The purpose of this investigation was to study mi-crostructural changes that take place in a commercial nickel-chromium-iron alloy, such as Incoloy "901," over long periods of time at temperatures up to 13 50°F and to correlate such changes with creep-rupture properties. of the alloy. In the course of the investigation, some microstructural changes were observed which correspond to microstructural alterations found in previous studies of Inconel "X." Microstructural phases have been identified in both alloys with the aid of electron and X-ray diffraction and an electron probe microanalyzer. The materials used in this investigation were from commercial heats of the following chemical composition: Pct Cr Pct Ni Pct Mo Pct Fe Pct Cb Incoloy "901" 12.51 43.67 5.72 Bal. -Inconel "X" 14.61 73.44 - 6.61 1.02 Pct A1 Pct Ti Pct Mn Pct Si Pct C Incoloy "901" 0.10 2.46 0.46 0.24 0.05 Inconel "X" 0.70 2.43 0.51 0.36 0.04 Long-term creep-rupture tests were conducted on Incoloy "901" at temperatures of 1000°, 1200°1 and 1350°F. A summary of these results is shown in Table I along with the results of a long-term creep-rupture test conducted on Inconel "X" at 1500°F. The heat treatments given the specimens are also shown in Table I. When the creep-rupture results on Incoloy "901" are compared with shorter-tertn data, there is no evidence of any undue deterioration of the creep-rupture strength with time, nor is the elongation decreased more than would be expected from the influence of time alone. INCOLOY "901" PHASE IDENTIFICATION Microscopic examination of the Incoloy "901" specimens reveals a dark etching grain boundary constituent in the specimen which was under test for about 11,500 hr at 1000°F. This grain boundary constituent was seen to increase in size as the test temperature was increased to 1200" and 1350°F, Figs. 1, 2, and 3. Simultaneously, an acicular structure of definite crystallographic orientation appears within the grains of the 1350°F specimen, Fig. 3. Under examination with the electron microscope, a fine precipitate is observed in the 1200 °F specimen. Fig. 4. It becomes considerably larger in the 2850-hr 1350°F specimen, Fig. 5. In a few areas the precipitated particles have been replaced by, or transformed to, small needles or platelets. After 7380 hr at 1350°F, Figs. 3 and 6, the needles are much larger and have replaced nearly all of the fine precipitate. X-ray diffraction studies detected the presence of hexagonal Laves phase and hexagonal n phase (Ni,Ti) in the 1350°F specimens, Table 11. Only the Laves phase was detected by X-ray diffraction in the 1200°F specimen and none was detected in the 1000°F specimen. In this latter specimen, however, the small quantity of precipitate in the grain boundaries (presumed to be Laves phase) precludes its detection by X-ray diffraction. Composition of the hexagonal Laves phase is not yet known, although its diffraction

Jan 1, 1960

By K. J. Gill, P. Chiotti

Thermal, metallographic, and vapor pressure data were obtained to establish the phase diagram for the thorium-zinc system. Four compounds corresponding to the stoichiometric formulas Th2Zn, ThZn2, ThZn4, and Th2Zn17 were observed. The melting points of these compounds under constrained vapor conditions were found to be 1055°, 1105°, 1095°, and 1015°C respectively. Four eutectics exist in this system with the following eutectic temperatures and thorium contents in wt pct: 1040°C, 89 Th; 945°C, 79 Th; 1045°C, 55 Th; and 995°C, 35 Th. The standard free energy, enthalpy, and entropy of formation for the compounds were determined from zinc vapor pressure data. ALTHOUGH the thorium-zinc system has been of technical importance in the commercial production of thorium and more recently of interest to some proposed pyrometallurgical separations processes for reactor fuels or blanket materials no detailed investigation of this system has been made in the past. Some previous work has been reviewed by Hanssen.1 The existence of four compounds, Th2Zn17, ThZn4, ThZn2, and Th2Zn has been established and their crystal structures have been studied by X-ray diffraction techniques.2-4 Smirnov et al.,5 determined the solubility of thorium in zinc and investigated the two-phase region between Th2Zn17 and zinc. They calculated some of the thermodynamic properties for these alloys from emf measurements on the cells Th/ThCl4(fused)/Cl2, C and Th2Zn17(s)+ Zn (l)/ThCl4(fused)/Cl2, C. The possibility of reaction between the anode materials and fused thorium tetrachloride to produce a lower valent thorium salt would have to be resolved in order to evaluate the results obtained. The purpose of the present investigation is to establish the phase diagram for the thorium-zinc system and to determine the standard free energy, enthalpy, and entropy of formation for the compounds formed. MATERIALS AND EXPERIMENTAL PROCEDURES The metals employed in the preparation of alloys were Bunker Hill slab zinc, 99.99 pct pure and Ames Lalboratory thorium. The major impurities in the thorium in parts per million were oxygen—1250, carbon—350, iron—100, nitrogen-75, and silicon—70. The metals in the form of several small massive pieces, total weight about 30 g, were cleaned with acid, washed with water and acetone, and dried just prior to their use for alloy preparation. Tantalum containers were found to be suitable for all the alloys studied and were used throughout this investigation. The metals were sealed inside of small 1-in.-diam tantalum crucibles by welding on preformed tantalum covers. A small thermocouple well made from 1/8-in.-diam tantalum tubing was welded in the bottom of each crucible. These crucibles were enclosed in

Jan 1, 1962

By P. G. Sprang, S. Rosen

The Y-Al- C ternary phase diagram for the com -position range from 55 to 100 at. pct Y and for a temperature of 950°C has been constructed from metallographic and X-ray diffraction data. The significant features of the isothermal section which is presented are the small solubility of aluminum in the binary Y3C phase and the presence of a ternary carbide phase with a small homogeneity range around the composition Y3AlC. The similarity in structure and composition between these two carbide phases is discussed in the light of the misci-bility gap which exists between room temperature and 1300°C, which was the temperature range studied. THE occurrence and stability criteria of ternary T3BC, carbide phases having the L&apos;12 structure and the general formula T3BCx have been discussed in detail by stadelmaier1 and more recently by the present authors., Stadelmaier found many representatives of these phases formed between the T metals Mn, Fe, Co, and Ni and the B metals Mg, Al, Zn, Ca, Ge, Cd, In, Sn, Hg, T1, and Pb. The present authors reported on a new series of iso-structural carbide phases formed between the T metals of the Sc group and the B metal, aluminum. These new phases are unique in that many of the Sc group elements form stable binary carbide phases which occur about the composition T3C and which crystallize with a structure similar to that of the T9BCx ternary carbide phases. Furthermore, these binary and ternary phases are found to coexist in equilibrium with each other. The similarity in structure and composition between, for example, Y3C and Y3AlC in the light of the mis-cibility gap which had been found to exist between these two phases indicated a need for a more thorough knowledge of the phase relations which occur in alloys based upon these compositions. Consequently, an investigation of the phase equilibria in the yttrium-rich portion of the Y-A1-C phase diagram was undertaken. The limiting Y-A1 binary diagram determined by Lundin and Klodt3 and the results of Spedding et al.4 governing the Y3C phase were used as guides in , constructing the diagram. EXPERIMENTAL Alloys weighing about 10 g were prepared by arc-melting powder compacts made from the various components on a water-cooled copper hearth using a tungsten electrode in a gettered argon atmosphere. Good homogeneity was obtained by melting the alloys three times. The purity of the yttrium used was 99.9 pct (0.07 pct rare earths). The purity of the aluminum and the carbon was 99.99 pct. The results of chemical analysis on five representative alloys showed that losses during the alloying process were predominantly carbon losses and that in no instance did the measured concentration of this component differ by more than 3 at. pct from the intended composition. For this reason the alloy compositions which are reported are those of the charges prepared for melting. Three alloys which were analyzed for copper, tungsten, nitrogen, and oxygen compositions showed that no contamination had occurred either during casting or during the subsequent annealing treatments. For the major part annealing treatments were carried out on specimens which were placed in

Jan 1, 1965

By R. W. Fountain, W. D. Forgeng

A new constitutional diagram is presented for the cobalt-rich end of the cobalt-titaniurn system. The modifications result from the presence of a new, intermediate, fcc phase, ?, the existence and homogeneity limits of which were established by metallo-graphic and X-ray studies of alloys containing from about 3 to 30 pct Ti. The precipitation of the ? phase from supersaturated solid solution was studied by hardness and electrical resistivity measzsrements, and two distinct stages in the process were observed. COBALT forms the base for a number of precipitation -hardenable alloy systems which may be divided into two distinct categories of practical interest, a) those hardened by intermetallic compounds and b) those hardened by carbide formation. The precipitation of intermetallic compounds from solid solution in cobalt-rich alloys has, however, received very little attention. Although the phase diagrams for some binary systems capable of precipitation have been determined,&apos; there is an almost complete lack of data on the property changes associated with the precipitation, and even less information on the kinetics of the reactions or morphology of the products. More information is available on the precipitation of carbides because of the practical significance in superalloys. A survey of cobalt binary phase diagrams suggested that the cobalt-titanium alloys might provide interesting and useful precipitation-hardenable alloys. The equilibrium diagram as proposed by Wallbaum2 is shown in Fig. 1. Köster and wagner3 have indicated that the maximum solubility of titanium in cobalt is about 10 pct at the eutectic temperature (1135oC), and this decreases to about 7.2 pct at room temperature. The temperature of the allotropic transformation in cobalt is lowered by the addition of titanium, about 5 pct being sufficient to retain the high-temperature fcc modification to room temperature. Wallbaum and Witte 4,5 have indicated that the precipitating phase in alloys containing up to about 29 pct Ti is Co2Ti, a hexagonal Laves phase of the MgNi, type. With slightly higher titanium contents, they also report a cubic modification of Co2Ti of the MgCu2-type Laves phase. Duwez and Taylor6 confirmed the existence of the hexagonal (MgNi2) modification but not the cubic (MgCuz) modification of Co2Ti and suggested that existence of the cubic form may have resulted from impurities in Wallbaum&apos;s alloys. In their work on Laves-type phases, Elliott and Rostoker 7 reported the cubic modification of Co,Ti, but did not confirm the existence of the hexagonal modification. However, Dwight8 in a discussion of the work of Eliott and Rostoker again showed the existence of both modifications of Co2Ti, a result which was confirmed at that time by Elliott and Rostoker. As a result of a study on the iron-cobalt-titanium system, Köster and Gellers suggested the existence of a Co3Ti phase isomorphous with Fe3Ti. Witte and Wallbaum,&apos; however, established the fact that no compound, Fe3Ti, exists in the iron-titanium system, and in a later publication, wallbaum2 stated that Kitster and Geller&apos;s reasoning was not valid, and no compound, Co3Ti, exists. This conclusion was later acknowledged by Köster.10 Preliminary experiments by the present authors to determine the precipitation-hardening characteristics of the cobalt-rich, cobalt-titanium alloys re-

Jan 1, 1960

By P. Chiotti, J. T. Mason

Metallographic, thermal, X-ray, and vapor-pressure data were employed in establishing the Ce-Zn phase diagram. Nine compounds and three eutectics were observed. The eutectic compositions in weight percent zinc and temperatures are 10 pct and 495"C, 37 pct and 795°C, and 56 pct and 810°C. Three of the compounds, CeZn, CeZnn, and CeZne.5, melt congru-ently at 825°, 87.505°, and 980°C, respectively; the other six compounds, CeZns, CeZns.67, CeZn4.5, CeZn5.25, CeZn,, and CeZnll, decompose peritectically at 820°, 840°, 870°, 885", 960°, and 795"C, respectively. The compound CeZn5.25 is an AB5 type with a small range of composition which lies on the zinc-rich side of the 1 to 5 stoichiometry. Zinc vapor-pressure data obtained by the dewpoint method were employed in calczilating the thermodynamic properties of both the liquid alloys and solid compounds. The data for the A critical review of the information available on the Ce-Zn system up to 1960 has been given by K. Gschneidner, Jr.&apos; Of particular interest is the work by J. schramm2 who investigated the region between 55 and 100 wt pct Zn and reported a CeZn,, compound which decomposes peritectically at 785"C, a CeZn, compound which melts congruently at 972"C, and five temperature horizontals at 79@, 817", 840", liquid alloys were fitted to the relation from which the relation was obtained by integration. In these relations F XS is the excess partial molal free energy, N is atom fraction, and the constants c and d have the values c = (-22,000 + 7.0T) and d = 6171,000 + 74.0T). Relations for the standard free energy of formation for the solid compounds were derived from the data. At 973°K the enthalpy in kilocalories per gram-atom z)aries from —11.3 for CeZnz to —8.21 for CeZn~l and the free energy in the same units varies from —7.1 joy CeZn, to -3.82 for CeZn11. and 942°C in the region between 55 and 75 wt pct Zn. No attempt was made to identify other features of the system. I. Johnson and R. yonko3 used a fused-salt cell to measure the electromotive force between pure cerium and an alloy consisting of zinc-rich liquid in equilibrium with CeZnll. Solubility data for cerium in zinc has been reported by J. Knighton et u1 .4 Johnson and Yonko give the following equa-. tions for the mole fraction of cerium in zinc and the standard free energy of formation of CeZnll:

Jan 1, 1965

By W. D. Forgeng, P. F. Wieser

MnSi alloys containing from 2 to 24 at. pct Si have been investigated by metallographic and X-ray methods. Contrary to published data, the temperature of the ß-manganese to a-manganese transformation is lowered by the addition of silicon. The existence of two additional intermediate phases E and ( is demonstrated. E, ranging in composition from 12 to 15.5 at. pct Si, forms by a peritectoid reaction between ß manganese and the second new phase, (. The latter phase is homogeneous at silicon contents from 16.2 to 18 at. pct and forms peritectically from the 0 phase and liquid. The eutectic reaction occurs between the ( phase and MnsSi rather than the ß phase and Mn3Si. As a result of this investigation, a modification of the manganese-rich section of the MnSi equilihrium diagram is proposed. The binary diagram of Mn-Si, according to Hansen,&apos; Fig. 1, exhibits in the low-silicon region (below 24 at. pct Si) a peritectic transformation at 900°C of 0 manganese + Mn3Si to a manganese as determined by R. Vogel and H. Bedarff2 as well as a thermal effect at 980°C between 16.2 and 22.6 at. pct Si reported by the same authors. The nature of the latter transformation was unknown. K. Amark, B. Boren, and A. westgren3 reported indications that possibly four more phases existed in the range of 10 to 27 at. pct Si. The present investigation deals with the phase relations in the range of 2.1 to 17.7 at. pct Si. A few spot checks were conducted up to 23.84 at. pct Si. EXPERIMENTAL PROCEDURE Small experimental heats were prepared by induction melting in a vacuum furnace using argon as a protective atmosphere. The crucible was of commercial-quality magnesia. The furnace charge consisted of silicon metal and dehydrogenated electrolytic manganese. The typical analysis for these materials is shown in Table I. The molten alloy was tapped into 2-in.-square, 7-in.-long ingots. After visual inspection, a small quantity of alloy in the middle portion of the ingot was separated for use in this investigation.MnSi alloys containing from 2 to 24 at. pct Si have been investigated by metallographic and X-ray methods. Contrary to published data, the temperature of the ß-manganese to a-manganese transformation is lowered by the addition of silicon. The existence of two additional intermediate phases E and ( is demonstrated. E, ranging in composition from 12 to 15.5 at. pct Si, forms by a peritectoid reaction between ß manganese and the second new phase, (. The latter phase is homogeneous at silicon contents from 16.2 to 18 at. pct and forms peritectically from the 0 phase and liquid. The eutectic reaction occurs between the ( phase and MnsSi rather than the ß phase and Mn3Si. As a result of this investigation, a modification of the manganese-rich section of the MnSi equilihrium diagram is proposed. The binary diagram of Mn-Si, according to Hansen,&apos; Fig. 1, exhibits in the low-silicon region (below 24 at. pct Si) a peritectic transformation at 900°C of 0 manganese + Mn3Si to a manganese as determined by R. Vogel and H. Bedarff2 as well as a thermal effect at 980°C between 16.2 and 22.6 at. pct Si reported by the same authors. The nature of the latter transformation was unknown. K. Amark, B. Boren, and A. westgren3 reported indications that possibly four more phases existed in the range of 10 to 27 at. pct Si. The present investigation deals with the phase relations in the range of 2.1 to 17.7 at. pct Si. A few spot checks were conducted up to 23.84 at. pct Si. EXPERIMENTAL PROCEDURE Small experimental heats were prepared by induction melting in a vacuum furnace using argon as a protective atmosphere. The crucible was of commercial-quality magnesia. The furnace charge consisted of silicon metal and dehydrogenated electrolytic manganese. The typical analysis for these materials is shown in Table I. The molten alloy was tapped into 2-in.-square, 7-in.-long ingots. After visual inspection, a small quantity of alloy in the middle portion of the ingot was separated for use in this investigation.

Jan 1, 1964

By H. I. Aaronson

ONLY limited studies have been made of pro-eutectoid a morphology in hypoeutectoid Ti-Cr alloys during previous investigations, 1-3 The nature of the eutectoid reaction, ß?a + TiCr2, has been considered in somewhat more detail.1,2,4,4 Inasmuch as the proeutectoid ferrite reaction in plain carbon steels has recently been extensively investigated in this laboratory, advantage has been taken of the opportunity to compare the morphologies developed in two rather different eutectoid systems in order to permit a preliminary estimate to be made of the degree of generality which may be ascribed to the various results obtained. Experimental Procedure The chemical analyses of the alloys used are given in Table I. The 7.22 pct Cr alloy was prepared by the Battelle Memorial Institute from iodide titanium and high purity chromium. The initial ingot was made by arc melting a master alloy and iodide titanium. The resultant ingot was twice remelted. by the consumable electrode technique. Upset forging at 950°C reduced the final ingot to 3/4 in. sq bars. The 10.94 pct Cr alloy was prepared by the Titanium Metals Corp. of America from sponge titanium and chromium of ordinary purity. This alloy was melted twice by the consumable electrode technique, forged at 1200° to a 3/4 in. sq bar, and rolled at 815°C to a 1/2 in. round. Specimens 1/4 in. x 1/4 in. x 1/16 in. were cut from the 7.22 pct Cr alloy; halves of disks 1/2 in. in dlam and 1/16 in. thick served as specimens of the 10.94 pct Cr alloy. These specimens were individually encapsulated in Vycor, Each capsule was evacuated three times to a pressure of 2 to 6 x lo-&apos; mm Hg and flushed twice with a reduced pressure of argon prior to sealing. The inclusion of two small packets of zirconium chips permitted the capsules to be outgassed and internally gettered, without heating the specimen, prior to heat treatment. The specimens were solution annealed for 40 min at 1000°C, isothermally reacted in deoxidized lead pots, and quenched in iced water. Most of the isothermal reaction treatments were carried out at 25" intervals between 725° and 550°C on specimens of the 7.22 pct Cr alloy and at 650°C on specimens of the 10.94 pct Cr alloy. Results and Discussion Morphologies of Proeutectoid a*—Kinetics of the ß?a Reaction: Fig. 1 shows the TTT-curves for the beginning of the proeutectoid a and the eutectoid reactions in the temperature range investigated in the 7.22 pct Cr alloy. The M, temperature in this alloy is below 0°C. The curve for the beginning of the proeutectoid a reaction is in good agreement with that of Frost, Parris, Hirsch, Doig, and Schwartz&apos; for a 7.54 pct Cr alloy, appears at considerably earlier reaction times than the corresponding curve presented by Rostoker4 for a 7 pct Cr alloy, and lies athwart the curve of Spachner and co-workers11 for an 8 pct Cr alloy. Differences in criteria for the beginning of the a reaction and the relatively high rate of this reaction at lower temperatures presumably account for some of these divergences and may also explain, in part, the apparent failure of an iodide-base 7.22 pct Cr alloy to react more slowly than sponge-base 7 to 8 pct Cr alloys. Grain Boundary Allotriomorphs: Crystals of this morphology nucleate at and grow preferentially and more or less smoothly along the matrix grain boundaries.5-7 Allotriomorphs are the first morphology to precipitate at most grain boundaries at reaction temperatures from 725°C, the highest temperature studied, at least through 575°C. Fig. 2 Shows typical allotriomorphs, formed during an early stage of transformation at 725°C. Grain boundary allotriomorphs are not the predc=iaant morphology at late reaction times at any temperature studied. However; from 725" through

Jan 1, 1958

By J. Gordon Parr, D. H. Polonis

The formation and subsequent decomposition of metastable phases in Ti-Ni alloys containing up to 11 pet (atomic) Ni have been studied. The decomposition of a completely retained ß phase and of a completely transformed 8 phase in a 6 pet alloy has been followed by X-ray diffraction and metallographic methods and by hardness determinations made during the process. The possible mechanisms of the reactions are discussed. PREVIOUS investigations of the Ti-Ni system have usually been limited to phase diagram studies; but little attention has been given to metastable phases and tempering processes. These are important, however, since equilibrium conditions are rarely achieved—nor, perhaps, are they desired—in practice. Up to the present time the available transformation data for titanium alloys appear to be restricted to isothermal transformation curves for selected binary and ternary systems. No quantitative studies have been made of the mechanism by which equilibrium is approached. Several independent investigations have been carried out to determine the Ti-Ni phase diagram. Early work by Wallbaum&apos; and later by Long et al.&apos; produced tentative phase diagrams which were not representative of true binary conditions, since the alloys were contaminated with oxygen and nitrogen. The most recent diagram (Fig. 1) is due to Margolin et al.," and a slight modification of the a + ß/ß boundary (shown dotted) has been proposed by McQuillan.&apos; The structure of the phase Ti,Ni was shown by Duwez and Taylor" to be face-centered-cubic, with 96 atoms per structure cell. The formation of a martensitic ß phase (close-packed-hexagonal) in Ti-Ni alloys has generally been observed when specimens of low nickel content are rapidly cooled from the 0 (body-centered-cubic) range. Margolin et al. reported an increasing tendency for P to be retained in lump specimens as the nickel content increases. McQuillan&apos; has studied the effect on the microstructure of delay time in quenching, and has reported the formation of pro- eutectoid a precipitate in a quenched hypoeutectoid alloy. In a previous study of Ti-Fe alloys," experiments were made on powder specimens produced by filing alloy ingots. Techniques were devised for the preparation and heat treatment of alloys, and X-ray diffraction methods were used both for structure determinations and phase ratio estimations. A similar approach has been made in this study of Ti-Ni alloys. Experimental Methods Ten alloys ranging from 0.25 pet to 18 pet Ni* * All compositions are in atomic percentages unless otherwise stated. were prepared from iodide titanium bar stock (hardness Rf 72) and Johnson-Matthey spectrographic standard nickel by levitation melting.&apos; As a check

Jan 1, 1957

By R. I. Jaffee, C. T. Sims, C. M. Craighead

The fabrication of rhenium metal by powder metallurgy techniques is discussed. The following physical and mechanical properties have been measured and are reported: lattice constants, melting point, electrical resistivity, thermal expansion, spectral emissivity, modulus of elasticity, tensile properties and ductility at room and elevated temperatures, work hardening, recrystallization, grain growth, and oxidation resistance. IT is only occasionally that a metal reaches present-day technology with as little known about its properties as element No. 75, rhenium. Relatively scarce to date, it has been studied only infrequently and used less. Previous work on rhenium has been carried out primarily by European scientists. However, realization that rhenium occurs in commercial quantities in the United States recently has focused attention on this metal. Rhenium is a very dense high melting hexagonal-close-packed metal. Like tungsten and molybdenum, its highest oxide is volatile. However, when protected by reducing or inert atmospheres, its strength at high temperatures is greater than that of tungsten. Rhenium work hardens more than any other known pure metal, but, on annealing, becomes quite soft and ductile. Its vapor pressure is approximately that of tantalum, but it does not enter into the so-called "water cycle" nearly as readily as does tungsten. (The water cycle is a deleterious phenomenon causing blackening of lamp bulbs and vacuum tubes with subsequent failure of the filaments.) These and other properties, some already known and some found as a result of the current researches, have indicated that rhenium has many potential uses, primarily of an electrical or electronic nature. Rhenium or its alloys soon may be found in service as electron-tube filaments, cathodic emitters, or as electrical contact materials. In addition, rhenium or its alloys have potential use in such devices as high temperature thermocouples, and high wear-resistant parts such as precision instrument points. Rhenium, in the form of potassium perrhenate, was made available for this study by the Kennecott Copper Corp. The potassium perrhenate was reduced to rhenium metal in two stages, accompanied by leaching to remove potassium in the form of hydroxide. This impure metal was then oxidized to rhenium heptoxide and dissolved in water. The solution was neutralized with ammonium hydroxide to produce highly purified ammonium perrhenate, a typical analysis of which is shown in Table I. The ammonium perrhenate was ground to —325 mesh in a rubber-lined ball mill with Burundum balls. This operation caused a certain amount of impurity pickup, as reflected in the rhenium analysis shown in Table I. The fine ammonium perrhenate was reduced to metal with hydrogen. Particle size of the powder ranged from 1 to 25 microns, as illustrated by Fig. 1.

Jan 1, 1956

By R. J. Reynik, F. R. Meeks

IN the classical paper1 of Mehl and Rhines entitled "The Rates of Diffusion in the Alpha Solid Solutions of Copper", the authors state: "It has been mathematically demonstrated that the treatment of the data as though the diffusion were taking place between two flat plates instead of radially in a cylinder has no appreciable effect upon the results. Were the bars of a much smaller diameter, a modification of the analytical treatment would have been required." Unfortunately, the errors in the reported concentrations are much too large to neglect, as shown in the results below. A planar treatment of diffusion data assumes that the reference planes separated by a nominal distance (AX) have identical areas. An adjustment must be made on the concentration data should the successive diffusion planes differ in area. The Mehl-Rhines&apos; diffusion samples were cylindrical with copper electroplated onto copper-based alloys. Successive surface areas through which diffusion occurs are in the ratio ri/r1, where ri is the radial distance between the center of the sample and the center of the i th shell to which the host (alloy) atoms diffuse, and r1 is the radial distance between the center of the sample and the center of the innermost (i = 1) experimentally reported shell from which the alloy atoms diffuse. Hence, ri/r1 (i = 1, ..., n layers) is a dimensionless correction factor to be multiplied by the reported concentration data before planar-data analysis is attempted. In the forty-seven samples reported, the original interface between the copper alloys and the electroplated copper was located at a mean distance of 8.255 mm from the center of the samples. The successive layers, machined off the samples, were approximately 0.0762 mm thick. The distance (xi) at which concentration (ci) was reported was always a distance measured with respect to the original interface, that is, with respect to the fixed distance of 8.255 mm. Hence, the as yet undiscovered Kirkendall Effect was precompensated. The values of ri and r1 are therefore given by ri = 8.255 + xi and Y, = 8.255 + x,. The illustrative table below lists the position xi (mm), ri (mm), the value of ri/r1, the reported concentration (ci) in weight percent (only valid for radial analysis of data), and the corrected concentration [c&apos;i = (ri/r1)ci] in weight percent (only valid for planar analysis of data). Similar tables can be constructed for the remaining forty-six samples. The corrected data can be used to re-evaluate the diffusion coefficients for the reported systems.

Jan 1, 1965