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By Earnshaw Cook, J. A. Fellows, Howard S. Avery

Heat-resistant alloys of the higher nickel and chromium ranges have been empirically developed through the practical experience of the past two decades to a position of significant industrial importance. Few quantitative metallurgical data have been published, however, that could be related directly to engineering specifications and design. Inadequate information led to the inauguration, in 1934, of a research directed toward a better understanding of the properties and performance of these steels at temperatures between 1400° and 2000°F. Investigation was concentrated upon the 26 per cent Cr:I2 per cent Ni type, which owes its extensive application to an excellent combination of economy, strength and surface stability at elevated temperatures. The requirements of the oil-refining industry have been prominent in influencing the scope and evolution of the program. Creep strength at the maximum allowable deformation ranges has been employed to evaluate the merits of these alloys. Permissible elongation rates vary from 1.0 per cent in one million hours for turbine blades to I .o per cent in ten thousand hours for tube supports and heating-furnace equipment. Throughout this paper, the limiting creep stress (L.C.S.) for a given temperature has been defined as the stress that will produce a uniform elongation rate (stage 11) of one per cent in ten thousand hours as extrapolated from rates per hour for tests of less than two thousand hours duration. Certain current specifications require a minimum L.C.S. of 1600 lb. per sq. in. at I800°F., of 3600 lb. per sq. in. at 1400°F., and include a tensile acceptance test at room temperature, after aging for 24 hr. at 1400°F., as a measure of embrittlement in service. An elongation of 4 per cent is considered satisfactory for general use, while a minimum of 9 per cent is prescribed on occasion. Many differences of opinion concerning the properties and performance of these alloys appear to result from their sensitivity to unappreciated variations of chemical analyses which, for example, may change their L.C.S. from 3500 to 350 lb. per sq. in. at 1800°F. Within broad chemical limits, it is possible to balance the six or more important elements present to attain desirable properties with a number of different combinations. One method for adjusting the composition has been described under the sponsorship of the Alloy Casting Institute,' in which, by maintaining a substantially austenitic* alloy, optimum creep strength and minimum tendency to embrittle after aging are expected. Unless, however, the limits are undesirably close, it has not been possible to ensure a moderate range of strength and of ductility by any chemical specification alone. There is a definite need

Jan 1, 1942

By J. A. Fellows, Earnshaw Cook, Howard S. Avery

Heat-resistant alloys of the higher nickel and chromium ranges have been empirically developed through the practical experience of the past two decades to a position of significant industrial importance. Few quantitative metallurgical data have been published, however, that could be related directly to engineering specifications and design. Inadequate information led to the inauguration, in 1934, of a research directed toward a better understanding of the properties and performance of these steels at temperatures between 1400° and 2000°F. Investigation was concentrated upon the 26 per cent Cr:I2 per cent Ni type, which owes its extensive application to an excellent combination of economy, strength and surface stability at elevated temperatures. The requirements of the oil-refining industry have been prominent in influencing the scope and evolution of the program. Creep strength at the maximum allowable deformation ranges has been employed to evaluate the merits of these alloys. Permissible elongation rates vary from 1.0 per cent in one million hours for turbine blades to I .o per cent in ten thousand hours for tube supports and heating-furnace equipment. Throughout this paper, the limiting creep stress (L.C.S.) for a given temperature has been defined as the stress that will produce a uniform elongation rate (stage 11) of one per cent in ten thousand hours as extrapolated from rates per hour for tests of less than two thousand hours duration. Certain current specifications require a minimum L.C.S. of 1600 lb. per sq. in. at I800°F., of 3600 lb. per sq. in. at 1400°F., and include a tensile acceptance test at room temperature, after aging for 24 hr. at 1400°F., as a measure of embrittlement in service. An elongation of 4 per cent is considered satisfactory for general use, while a minimum of 9 per cent is prescribed on occasion. Many differences of opinion concerning the properties and performance of these alloys appear to result from their sensitivity to unappreciated variations of chemical analyses which, for example, may change their L.C.S. from 3500 to 350 lb. per sq. in. at 1800°F. Within broad chemical limits, it is possible to balance the six or more important elements present to attain desirable properties with a number of different combinations. One method for adjusting the composition has been described under the sponsorship of the Alloy Casting Institute,' in which, by maintaining a substantially austenitic* alloy, optimum creep strength and minimum tendency to embrittle after aging are expected. Unless, however, the limits are undesirably close, it has not been possible to ensure a moderate range of strength and of ductility by any chemical specification alone. There is a definite need

Jan 1, 1942

By Nicholas J. Grant, A. W. Mullendore, Paul Predecki

The Duwez technique of splat cooling in which a molten droplet of metal is accelerated and made to impact on a cold, highly conducting substrate was investigated.- An apparatus for producing "splat" was constructed employing an explosive powder charge to accelerate the molten metal. Transport of the molten metal just prior to impact with the substrate was studied by means of high-speed photography. The molten particles are small spherical droplets from about 1 to 50 µ diameter. The average cooling rates for aluminum, silver, and a gold alloy splatted on nickel substrates at room temperature were determined experimentally and were found to vary from 18 to 5 x 10' C per sec. The heat-transfer coeficients for pure aluminum and pure silver splats cooled on nickel substrates at room temperature were found to be 2.7 to 6.8 and 13.6 to 54.2 cal per sec sq cm 'C (2 to 5 x 104 and 1 to 4 x 18 Btu per hr sq ft OF), respectively. Solidification rates in pure aluminum and silver splats were calculated. "SPLAT cooling" is a term describing a technique for extremely rapid freezing and cooling of molten metals and alloys to room temperature or below. The technique was developed by Duwez et al.' in 1960, and after refinement2 consisted in transferring a few tenths of a gram of molten metal to near sonic velocity to strike a suitably placed cold copper surface. Upon impact, the metal spread into a thin nonuniform film called a splat, about 10-4 cm thick. The splat particles produced in this manner were thin enough in some areas to be suitable for transmission electron microscopy, without further treatment, and, together with X-ray data, yielded a variety of structures which would be classified as follows: 1) supersaturated solid solutions (increase in solubility limit), 2) metastable crystalline stoichiometric and non-stoichiometric intermediate phases ( in Au-20.5 at. pct Si and in Au-14 at. pct Sb, respectively), 3) amorphous alloys, 4) retained high-temperature phases, 5) alloys having equilibrium phases present, but with unusual, markedly altered structures. The crystalline phases present in splats were usually extremely fine-grained and had low dislocation densities. Present interest in splat cooling centers around a study of the unusual structures produced by the technique, their contribution to alloy theory, and the possibility of the development of new or unusual properties. In addition, the technique is being examined for the production of bulk quantities of new high-strength alloys for low- and high-temperature use. Although a number of alloy systems have already been investigated by the technique, relatively little is known about the mechanism of splat formation and the physical conditions encountered by the metal during solidification and cooling. The purpose of this work was: 1) to determine the velocity, shape, and size of the molten metal droplets just prior to impact with the cold substrate; 2) to try to estimate or measure solidification rates, heat-transfer coefficients, and cooling rates in splats. EXPERIMENTAL TECHNIQUE Splatting Apparatus and Procedure. An apparatus similar to the type originated by Duwez et a1.,2 but employing an explosive charge to accelerate the molten metal, was used. A "Ramset" fastening tool (a gun normally used for driving studs into concrete, steel, and so forth) was mounted vertically above a resistance-heated graphite crucible as shown schematically in Fig. 1. A splat product was produced by melting a few tenths of a gram of metal at the position shown in Fig. 1, and then exploding a powder charge in the breach. The shock wave thus generated traveled down the barrel into the furnace, ejecting the molten metal through the 0.06-in.-diameter hole at the bottom of the crucible. The ejected metal was impacted on a high-conductivity metallic substrate where it formed a splat. High-Speed silhouette Photography. In order to investigate the size and shape of the ejected, molten metal just prior to impact with the cold substrate, high-speed silhouette photographs were taken, and are shown in Figs. 2 and 3. These photographs were obtained with the aid of an Edger-ton, Germeshausen, and Grier microflash unit and a submicrosecond flash drive equipped with a variable time delay. Two types of triggers were

Jan 1, 1965

By A. R. Elsea, E. E. Fletcher

Cast X-40 alloy with the lowest nitrogen content studied had a 100-hr rupture stress at 1500°F about equal to the reported value for the commercial alloy. Increases in nitrogen content progressively decreased the rupture strength, and with about 0.1 to 0.2 pct N a minimum strength was reached, but there was no accompanying change in microstructure. Further increases in nitrogen resulted in the fomzation of progressively increasing amounts of a lamellar, pearlitic-type structure at the grain boundaries, with accompanying increases in rupture strength, and above about 0.25 to 0.33 pct N, the rupture strength exceeded the normally reported values. Two series of melts representing different melting stocks exhibited minimum rupture strengths at different nitrogen contents. This suggests that some unidentwed trace element not removed by vacuum melting has a strong influence on the response of this alloy to variations in nitrogen content. In a series of melts with manganese contents ranging from 0.6 to 3.2 pct, the material with the lowest manganese content had the highest 100-hr rupture strength at 1500°F. The manganese variations had no appreciable effect on the microstructure. CAST cobalt-base alloys such as X-40* exhibit high strengths at elevated temperatures (1500°F and above), but they also show considerable scatter in high-temperature properties. Since the design of parts for critical applications must be based on the minimum expected properties, this broad scatter appreciably lowers the design strength of the alloy. If the relatively poor properties exhibited by a few test parts from any given group could be improved, the spread in properties would be narrowed, and the design stress could more closely approach the average expected properties. The purpose of this research investigation was to learn what were the causes of these erratic properties and, if possible, to find methods for eliminating them. Earlier unpublished work showed that vacuum melting markedly influenced the high-temperature strength of cast X-40 alloy. There were indications that this behavior may have been the result of changes in the nitrogen or manganese content of the castings. Consequently, in the current investigation, the influence of variations in nitrogen and manganese content on the structure and high-temperature properties of cast X-40 alloy was studied. In this program, X-40 melting stock was vacuum melted to produce a base material that was very low in nitrogen content and in volatile trace elements. Additions were then made to this base material to produce alloys of the desired composition. The metal was cast into investment molds producing coupons which were machined into stress-rupture bars. These bars were subjected to stress-rupture tests and me-tallographic study to determine the influence of changes in analyses (nitrogen and manganese content) on the structure and properties. Material—The melting stock used for these experi-ments was in the form of shot from two master heats of commercial X-40 alloy. Only one master heat was used in a given series of melts. Thus, any differences in structure or properties that showed up in the specimens from a given series would be the result only of the intended variations in composition. The certified analyses of the master heats were as follows: C Mn Si Cr Ni Lot No. 2 0.50 0.71 0.66 25.44 10.53 Lot No. 3 0.54 0.73 0.70 25.95 10.00 W Fe Co P S Lot No. 2 7.45 1.34 Bal. 0.011 0.013 Lot No. 3 7.58 1.53 Bal. 0.005 0.016

Jan 1, 1960

By Peter A. Rona

Investigations of the Gorda Ridge within the proclaimed U.S. Exclusive Economic Zone off northern California and southern Oregon, and the Mid-Atlantic Ridge extending along the center of the Atlantic Ocean, have received new impetus from the discovery of hot black smokers, sizeable polymetallic massive sulfide deposits, and vent animals at the Mid-Atlantic Ridge by the NOAA VENTS Program in 1985. Prior to that time, the oceanographic community focused investigations of seafloor venting at the faster-spreading oceanic ridges of the Pacific Ocean including the East Pacific Rise, the Galapagos spreading center and the Juan de Fuca Ridge because a consensus held that the faster spreading rates at these ridges were necessary to supply the heat to drive black smoker-type venting. As the only oceanic ridge in the Pacific with slow-spreading characteristics, the Gorda Ridge is an outlier of the slow-spreading ridge system that extends through the Atlantic Ocean and western Indian Ocean comprising more than half the 55,000-kilometer-length of the globe-encircling oceanic ridge system. In response to recent discoveries at the Mid-Atlantic Ridge the oceanographic community is shifting more of its activity to investigation of hydrothermal venting and related processes at slow-spreading oceanic ridges. The discovery cruise by a team of NOAA and university scientists followed by

Jan 1, 1988

By J. P. Nielsen, H. Margolin, E. Ence

The Ti-Ni phase diagram has been investigated up to 68 pct Ni with iodide titanium base alloys by metallographic, X-ray, and melting point methods, and from 68 to 90 pct Ni by examination of as-cast structures of sponge titanium base alloys. NVESTIGATION of the nickel-rich portion of the I Ti-Ni phase diagram was first reported by Vogel and Wallbaum in 1938.' This work was subsequently extended to lower nickel contents by Wallbaum' who indicated the possibility of a eutectic reaction for nickel contents below 38 pct. Long et al.3 studied the titanium-rich portion of the phase diagram and found eutectic and eutectoid reactions below 38 pct Ni. However, the temperature of the eutectic indicated by Long et al. was considerably lower than that suggested by Wallbaum. Long and his coworkers synthesized their alloys by powder metallurgical techniques and encountered oxygen and/or nitrogen contamination. Thus the diagram which was obtained did not represent binary alloying conditions. However from these results the features of the binary diagram were predicted. At Battelle Memorial Institute4 the Ti-Ni diagram was investigated up to approximately 11.5 pct Ni with sponge titanium alloys. The range of temperatures used was not sufficient to define the eutectoid temperature or composition. The data of Wallbaum2 and Long et al.8 were of particular interest for the present study, and although the work was originally concerned with the region below 40 pct Ni, the investigation was extended to higher nickel contents in an attempt to resolve the differences between these workers. Experimental Procedure Preliminary work on the Ti-Ni system was carried out with duPont Process A sponge titanium alloys to reduce the amount of subsequent work to be done with iodide titanium base alloys. The sponge titanium used contained 99.71 to 99.77 pct Ti, 0.1 pct Fe and 0.005 to 0.009 pct Ni. The iodide titanium obtained from the New Jersey Zinc Co. contained 99.9 to 99.95 pct Ti. Nickel used with sponge titanium was 98.9 pct pure. The high-purity nickel alloyed with iodide titanium was cobalt-free with approximately 0.05 pct C and was obtained through the courtesy of the International Nickel Co. The 15 g sponge titanium charges for melting were prepared by compacting in a die or by placing the weighed portions of nickel and titanium directly into the furnace. Iodide titanium charges were made by drilling holes in the as-received rod and inserting the nickel or by wrapping the nickel in sheet. Sponge titanium alloys containing from 0.2 to 90 pct Ni and iodide titanium alloys containing 0.2 to 68 pct Ni were prepared by these methods. In addition to these alloys several 1/2 1b sponge titanium alloys were supplied by the Allegheny Ludlum Co. The charges were melted in an arc furnace under an argon atmosphere. The procedures used were similar to those reported in the literature5,' and the furnace has been described.' Except for iodide titanium alloys with 40 to 68 pct Ni (see section on copper contamination), each alloy was melted for 1 min, then either turned over or broken before re-melting for an additional minute. Currents of 200 to 400 amp were used depending on the melting point of the alloy. Prior to heat treatment, alloys containing less than 14.5 pct Ni were hot-forged at 750°C. With the exception of alloys in the homogeneity range of the compound TiNi, alloys of higher nickel contents could not be hot-forged. Heat treatment of iodide titanium base alloys was carried out in argon-filled quartz capsules which were broken under water at the conclusion of heat treatment to quench the specimens. Temperatures were controlled to ±5oC and annealing times up to 48 hr were used. For melting point determination, specimens were placed in carbon crucibles which were in turn en-capsuled in argon-filled quartz capsules. The start of melting was determined by rounding of corners and by metallographic examination. Complete melting was considered to have occurred at that temperature at which the specimen assumed the shape of the crucible. Specimens were prepared for metallographic examination by mechanical polishing or by an electrolytic procedure." For alloys containing up to 80 pct Ni Remington A etch7 50 pct glycerine, 25 pct HNO,, 25 8 HF) was used. For higher nickel alloys aqua regia and Carapella's etch (5 g FeCl,, 2 ml HNO,, and 99 ml methyl alcohol) were employed. Specimens to be exposed for powder patterns were prepared by filing, by breaking specimens in a

Jan 1, 1954

By M. T. Simnad, G. Derge, I. George

Measurements of current efficiency on iron-silicate slags in iron crucibles showed that conduction is about 10 pct ionic in slags with less than 10 pct silica and about 90 pct ionic in slags with more than 34 pct silica, increasing linearly in the intermediate range. The balance of the conduction is electronic in character. Silicate ions are discharged at the anode with the evolution of gaseous oxygen. Transport experiments show that the ionic current is carried almost entirely by ferrous ions, which may be assigned a transport number of one. THERE has been increased evidence in recent years that the constitution of liquid-oxide systems (slags) is ionic.1-3 The principal studies designed to establish the structure of liquid slags have been by electrochemical methods', " and conductivity measurements1,6,7 which also have indicated the presence of semiconduction in several silicate systems1,4-0 and in pure iron oxide.' It is well known that many slag-forming metallic oxides have an ionic lattice type in the solid state, and their properties are determined to a large extent by the lattice defects and ion sizes. As Richardson8 as pointed out, the detailed models of liquid slags cannot be found on thermodynamic data only but "must rest on a proper foundation of compatible structural and thermodynamic knowledge, combined by statistical mechanics." A careful thermodynamic study of the iron-silicate slags has been carried out by Schuhmann with Ensio9 and with Michal.10 They obtained experimental data relating equilibrium CO2: CO ratios to slag composition and made thermodynamic calculations of the activities of FeO and SiO, and of the partial molal heats of solution of FeO and SiO2 in the slags. It was found that the activity-composition relationships deviate considerably from those to be expected from an ideal binary solution of FeO and SiO2. However, the partial molal heat of solution of FeO into the slags was estimated to be zero. Their experimental results were correlated with the constitution diagram for FeO-SiO2 of Bowen and Schairer,11 with the results of Darken and Gurry" on the Fe-O system, and with the work of Darken"' on the Fe-Si-O system. All these studies were found to be consistent with one another. The variation of the mechanism of conduction with composition in the liquid iron-oxide-silica system in the range from pure iron oxide to silica saturation (42 pct SiO2) in iron crucibles was reported in a preliminary note." The current efficiency, or conformance to Faraday's law, showed some ionic conductance at all compositions, the proportion increasing with the concentration of silica. The current-efficiency experiments since have been extended. Furthermore, transport-number measurements have been completed in silica-saturated iron silicates to determine the nature of the conducting ions. Experimental Current Efficiency in Liquid Iron Oxide and Iron Silicates using Iron Anodes: This study was carried out by passing direct current through slags in the range from pure iron oxide to iron oxide saturated with silica (42 pct silica), using pure iron rods as anodes and the iron container as the cathode. A copper coulometer was included in the circuit to indicate the quantity of current passed during electrolysis. Assuming that the cation involved is Fe-+, the theoretical quantity of iron lost from the anode according to Faraday's law may be calculated and when compared with the actual loss observed, gives an indication of the extent to which Faraday's law has been obeyed. It also gives an indication of the presence and extent of ionic conduction in the melt. Preparation of the Slags: About 100 g of chemically pure Fe,O, powder is placed in an iron pot which is heated by induction until the contents liquefy. In this way, FeO is produced according to the reaction Fe2O3 + Fe = 3 FeO. More Fe2O3 or SiO, powder is added and, when a sufficient quantity of molten slag is obtained, the induction unit is turned off, the pot withdrawn, and the molten slag poured on to an iron plate. Homogenization and Electrolysis of the Slag: Apparatus—After considerable development, the setup illustrated in Fig. 1 proved to be quite satisfactory. A is an Armco iron cylinder, 1 in. ID and 1/8 in. wall, consisting of three sections placed one on top of the other. The bottom section is a pot about 5 in. long with a small hole drilled in its bottom to allow withdrawal of gases during evacuation of the apparatus. The middle section is 6 in. long and consists of a pot which serves as the slag container, while the top section is a hollow-cylinder continuation of the slag-container pot. The height of this latter section is about 5 in., giving an overall length of approximately 16 in. The iron cylinder is constructed in this way for ease of fabrication, the individual sections becoming welded together after the

Jan 1, 1955

By S. G. Cupschalk, J. J. Wert, R. A. Buchanan

The uniaxial deformation of thermally ordered and disordered polycrystalline Cu3Pt was studied by means of the X-ray line - broadening analysis according to Warren and Averbach and the extension of this analysis to ordered fcc materials by Mikkola and Cohen. Because of the heat treatment history, extinction had a pronounced effect on the X-ray spectra of ordered and disordered C%Pt at small plastic strains. After an appropriate correction for extinction, the long-range order in thermally ordered ChPt was found to decrease at a slow constant rate with plastic strain. Furthermore, the antiphase domain probability increased at a constant rate to 17.5 pct strain. The effective particle size behavior indicated that the stacking fault energy is lower in ordered than in disordered Cu3Pt. Analysis of the stress-strain curves shouled that ordered Cuzt has a slightly lower yield Point but a much higher work-hardening rate than disordered Cu3Pt. THE presence of long-range order in a solid-solution alloy has a marked effect on its mechanical properties. While this effect has been known qualitatively for many years, it was not until recently that detailed investigations have been performed to determine the exact role long-range order plays in this strengthening mechanism. The development of an advanced, quantitative. X-ray diffraction analysis by Warren and Averbachl and the extension of this analysis to the L1, type super lattice by Mikkola and cohen2 have greatly accelerated research in this field. The research reported in this paper consisted of two primary phases. The first phase was to determine the effect of long-range order on the tensile properties of polycrystalline Cu3Pt. This objective was accomplished by comparing the stress-strain behavior of thermally ordered CusPt to that of thermally disordered CusPt. The second phase of the research was to determine the difference between the atomic arrangements in thermally ordered and thermally disordered Cu3Pt as a function of uniaxial deformation and thereby gain a deeper insight into the mechanism by which long-range order affects the tensile properties. This second objective was accomplished by applying the Warren-Averbach method of peak profile analysis to the X-ray diffraction patterns obtained from ordered and disordered Cu3Pt after given amounts of uniaxial deformation. EXPERIMENTAL PROCEDURE The Cu3Pt was prepared by vacuum melting and casting. After a homogenization anneal, the ingot was cold-rolled to sheet form. Two tensile specimens with gage sections of 2.50 by 0.500 by 0.115 in. were carefully machined from the sheet. The specimens were polished with a final step of 600-grit paper to insure smooth diffracting surfaces. Finally, one specimen was heat-treated to yield an average grain diameter of 0.016 mm and an initial degree of long-range order, S, of 0.825. The other specimen was water-quenched from above the critical temperature, 645"C, to yield an average grain diameter of 0.017 mm and zero long-range order. The heat treatment history of each specimen is shown in Table I. The tensile tests were performed utilizing a Research Incorporated Model 900.95 Materials Testing System. This unit employs a closed-loop feedback system which maintains a constant strain rate through an extensometer clipped to the gage section of the tensile specimen. A strain rate of 1.32 i0.02 x 10"4 sec-' was employed in testing both specimens. In the X-ray diffraction analysis, a General Electric XRD-5 diffractometer equipped with a pulse-height analyzer set for 90 pct efficiency was employed. The goniometer speed selected was 0.2 deg, 20, per min. Filtered Cu radiation was used for all peaks and all peaks were chart-recorded. Because of nonuni-form grain size. it was necessary to spin the specimens during X-ray analysis in order to obtain reproducible integrated intensities. The spinning rate was 2000 i100 rpm. The application of the Warren-Averbach method of peak broadening analysis to a diffraction pattern is very time consuming if done manually. In this research, the calculations involved were performed with the aid of a computer program by wagner.3 As reported by Wagner, the program is written in Fortran TV computer language. It was modified to Fortran I1 so as to be handled by the IBM 7072 computer at Van-derbilt University. In the X-ray diffraction analysis of uniaxially deformed Cu3Pt, the 100, 200. 400. 111, and 222 reflections were recorded from the initially ordered sample after 'plastic strains of 3.0, 6.0, 9.0, 12.0,

Jan 1, 1970

By C. M. Schwartz, A. E. Austin, W. B. Wilson

High-temperature X-ray diffraction studies were conducted with Th-U alloys with up to 10 wt pet U. The solid solubility of uranium in thorium as a function of temperature was determined by the method of lattice parameters. Thorium will dissolve up to 2.5 wt pet U at 950°, 4.5 wt pet U at 1150°, and 7.5 wt pet U at 1250°C. Determinations were made of the temperature of the transition of thorium and of the effect of uranium upon the transition. The a to ß transition for thorium was observed to occur at 1330' ±20°C. Mean coefficients of expansion were calculated for thorium and two alloys, and for ThO2 in contact with the thorium, using X-ray lattice-parameter data. Values obtained at 950° C for thorium and Tho2 were 12.1 and 9.40 X 10-6 per OC, respectively. Impurities obtained during the X-ray exposure were identified by diffraction and were essentially Tho2 and ThC, with two additional unknown phases being detected. The effect of the impurities upon the results is discussed. DIRECT investigation (i.e., high-temperature X-ray diffraction studies) of the phase diagram of thorium-rich uranium alloys has been shown to be necessary since recent work1 - has disclosed the presence of an allotropic transformation near 1400°C in pure thorium with the room temperature face-centered-cubic phase transforming to a body-cen-tered-cubic structure at the elevated temperature. The effect upon the transition of the addition of uranium to thorium and of the solubility of uranium in thorium at high temperatures remained unknown, yet was of interest in understanding fabrication procedures and elevated temperature use. The present work was undertaken to provide information in this area by determining the transition temperature, the effect of uranium on the transition temperature, and the solubility of uranium in thorium as a function of temperature. Experimental Work The high-temperature diffraction data were obtained using a camera especially designed for the purpose3 and capable of reaching temperatures in excess of 2000°C at pressures as low as 1 x 10." mm Hg. Temperature regulation was provided by regulating the power input to ±0.1 pet variation, and by regulating the water flow through the camera jacket to provide a constant thermal load. The X-ray sample was a rod nominally 80 mil diam, which was further turned down to 20 mil and then etched to 18 mil over 1/2 in. of one end. This was placed in the sample holder and mounted on the camera so that the smaller part was surrounded by a cylindrical tantalum-sheet radiation-type heating element. Diffraction from the sample was recorded on film after passing through a slot in the heating element and radiation-baffle shield and through beryllium vacuum windows. The X-ray film mounting was of the Straumanis type4 with a camera diameter of 114.59 mm. Since previous work of Chiotti' indicated that impurities considerably alter the transition temperature, chemical analysis of the arc melted iodide crystal-bar thorium samples was obtained prior to testing. The analysis disclosed the material to contain as low as 0.001 ±0,0002 wt pet H and 0.007 ±0.001 wt pet 0. Carbon was 0.003 wt pet and nitrogen less than 0.002 wt pet. This material was sealed in mild steel in an inert atmosphere and subsequently hot rolled to 3/8-in. diam rods at a temperature of 732°C. Following removal of the jacket, the material was pickled and cold swaged to 1/8-in. rods, from which the diffraction samples were prepared. The alloys were similarly prepared, with the uranium being added during arc melting. The uranium analyses of the alloys prepared appear in Table I. The experimental procedure for diffraction examination of the three samples of high-purity thorium differed from those of the Th-U alloys. The original practice, later modified, consisted of pumping down the camera with the diffusion pump on and then admitting liquid nitrogen to the cold trap of the system. This was modified for the Th-U alloys by maintaining liquid nitrogen in the cold trap at all times before and while the diffusion pump was heated. This minor change produced a reduction in the amount of carbon pickup by the sample during exposure to the diffusion-pump vapors. The sample was brought to the desired test temperature and exposed for 21/2 hr at pressures which were usually 2 x 10-6 mm Hg, or lower. Exposures were made at

Jan 1, 1959

By C. K. Gupta, P. K. Jena

Niobium (columbium) metal in the form of a button has been produced by calciothermic reduction of niobium pentoxide using sulfur as the heat booster. In these experiments with 50 g of niobium pentoxide per batch, the influence of percentage of calcium in excess of the stoichiometric, the percentage of sulfur by weight of the niobium oxide, and the outer-wall temperature of the bomb on the quality and yield of the metal has been studied. The maximum yield on this scale was 78 pct using 50 pct excess Ca and 20 pct of S and at a bomb-wall temperature of 600°C. In experiments conducted on 500 g of niobium pentoxide, reduction with 40 pct excess Ca and 20 pct of S, and at a bomb-wall temperature of 500°C, the.yield of the metal improved to 82 to 84 pct. Some of the reduced-metal samples have been electron-beam melted and the button melts thus obtained have been observed to be extremely ductile and capable of cold reduction to 98 pct without intermediate annealing. NIOBIUM, because of its high melting point (2468oC), its ductility, its resistance to corrosion especially in liquid metals, its high-temperature mechanical properties, and its relatively low capture cross section (1.1 barns) for thermal neutrons, has potentialities as a structural material in the field of nuclear engineering. Niobium and its alloys are also gaining importance in high-temperature technology as components in aircrafts, jet engines, and missiles, electronics, chemical engineering, and surgery. In 1904 Weiss and Aichel 1 produced niobium by reduction of its pentoxide by misch-metal. Since then a large number of investigations have been carried out for the extraction of niobium from its compounds. The various routes through which the metal niobium has been extracted from its compounds can broadely be classified into the following groups: a) reduction of its halides, oxides, and fluo-salts with metals like calcium, magnesium, sodium, and aluminum;2-10 b) reduction of its halides with hydrogen and oxides with carbon or carbides;11-18 c) electrolytic winning from its halides and fluo salts;19-25 and d) thermal decomposition and disproportion ion of its halides26-30 Among the processes in commercial use at present are sodium reduction or fused-salt electrolysis of potassium niobium fluoride, carbide or carbon reduction of the oxide, and hydrogen or metallothermic reduction of the chlorides. In 1907 Von Bolton2 prepared niobium by reducing niobium pentoxide with aluminum. Dennis and Adam-sona attempted reducing the oxide with magnesium powder in argon atmosphere and the resulting metal contained as much as 5 pct O. Mondolfo8 claimed a method for the reduction of the oxides with aluminum. In the case of both magnesium and aluminum reduction, it has been stated31 that by-product oxides such as MgO or A12O3 would involve a major separation problem. Further, in the case of alumino-thermic reduction, the solubility of aluminum in niobium and intermetallic formation have to be reckoned with. In 1922 Bridge3 produced niobium by calcium reduction of niobium pentoxide. Dennis and pdamson' obtained the metal by a similar method with an oxygen content of 0.2 pct. In most of the cases of metallothermic reduction of niobium oxides mentioned above, the metal niobium is obtained in the form of powder. In an attempt to produce massive niobium metal, lock' used iodine as the thermal booster in the bomb reduction of niobium pentoxide by calcium. Massive niobium metal was obtained with a maximum yield of 75 pct containing 0.1 pct 0. Joly32 carried out experiments on calcium reduction of niobium pentoxide on a 2-kg scale using sulfur as the heat booster in a sealed bomb filled with argon. In his experiments, the reaction was initiated by passing a current through niobium spiral wire embedded in the charge. In a typical run consisting of 2000 g of niobium pentoxide, 3100 g of calcium, and 800 g of sulfur (S/Nb2O5 = 3.3), he obtained massive metal in the form of a very uneven mass with a yield of -57 pct. With lower amounts of sulfur and calcium, powder niobium mixed with small metal globules was obtained and the yield was also poor. From these experiments he concluded that this method is not a suitable one for industrial development considering the quality and yield of metal obtainable and the possible hazards with higher ratios of sulfur to niobium pentoxide in the -~charere. In the present work, a detailed investigation on the calciothermic reduction of niobium pentoxide in a stainless-steel bomb in presence of sulfur as the heat booster was undertaken on a laboratory scale.

Jan 1, 1964

By D. L. Albright, R. W. Kraft

High-purity materials have been used in producing as-cast, controlled, colony, and degenerate solidification structures in the Fe-FeS eutectic. Experiments disclosed that this eutectic can be classified as normal and has a natural morphology composed of rodlike iron particles dispersed in a matrix of iron sulfide. The metallography of the various structures was studied, and a preferred crystallography was revealed in the controlled specimens produced by unidirectional solidification. The orientation effects found in these latter specimens are an [001] fiber texture in the -mowth direction of the bcc iron bhase and a texture corresponding to bicrystalline behavior in the hexagonal iron sulfide, with the growth direction near to (2111) poles. The observed texture of the iron phase is considered as indirect evidence that the alloy un-dercooled by at least 75°C before solidification. The unidirectional solidification of binary eutectic alloys has produced materials which exhibit a structure and properties markedly dependent upon the solidification process. In many cases a controlled microstructure with pronounced metallographic and crystallographic anisotropy can be experimentally achieved by proper regulation and balance of the growth rate of the alloy, the chemical purity of the starting materials, and the thermal gradient in the liquid at the liquid-solid interface. The purposes of this investigation were to produce various micro-structures in the Fe-FeS eutectic for subsequent study of their magnetic properties and to correlate the different structures with the solidification conditions in order to obtain a better understanding of the structure of eutectics. The Fe-S equilibrium diagram exhibits a eutectic composed of nearly pure iron and stoichiometric iron sulfide (FeS1.00), with the eutectic reaction occurring at 988°C and 31.0 wt pct S.1 Calculations indicate that this eutectic should solidify with about 9.5 vol pct Fe and 90.5 vol pct FeS, which in turn suggests2 that the micros tructure will consist of a rodlike iron constituent dispersed in a matrix of FeS. This characteristic has in fact been revealed some years ago.3 Thus, controlled solidification of this alloy might yield a material whose micromorphology would consist of very small ferromagnetic iron particles, rod-like in shape and aligned parallel to one another, supported in a matrix of antiferromagnetic FeS. Such specimens, because of the magnetic characteristics of the two phases, would be interesting subjects of study as magnetic materials. Hence the magnetic properties were considered in detail and are reported elsewhere.4 EXPERIMENTAL PROCEDURE The specimens of Fe-FeS eutectic were prepared from ultrapure iron (99.99+ pct) and high-purity sulfur (99.999+ pct). The iron was estimated to contain 60 ppm impurities (99.994 pct Fe) after zone purification.5 The ingots of iron were cut into chips, and the lumps of sulfur were ground into powder. In order to redice any nometallic impurities which might have accumulated during handling, the iron chips were annealed for 5 hr at 750° ± 10°C in a dry hydrogen atmosphere. Immediately after this treatment the chips were blended with the sulfur powder in eutectic proportions; the mixture was tamped into transparent fused quartz tubing and then vacuum-encapsulated under a pressure of 40 to 60µ of Hg. Because FeS expands upon solidification it was necessary to re-encapsulate the initial capsules so that oxidation reactions would be avoided when the inner tube cracked during solidification. For purposes of homogenizing the blended mixtures before solidification, the double capsules were heated to 750° ± 20°C and held for 20 hr; after this treatment the reacted product was weakly agglomerated. Each sample was then loaded into an apparatus for very rapid melting and freezing; this was accomplished by passing a molten zone through the specimen, using induction heating and a traverse mechanism. The resulting specimens solidified in the shape of the quartz tubing. Two sizes of specimens were used in this work, 18 mm diam by 100 mm long and 5 mm diam by 30 mm long. Metallographic examination of several ingots of both sizes after the above consolidation indicated no lack of compositional homogeneity and a random "as-cast" structure, because the travel rate was so rapid that unidirectional solidification was not achieved. Unidirectionally solidified specimens were resolidified in the apparatus shown schematically in Fig. 1, This equipment consisted of a kanthal resistance furnace mounted on the carriage of a zone-melting unit so that the heating element could traverse the length of the sample at a selected rate of speed. Large specimens were solidified with the mechanism tilted at ap-

Jan 1, 1967

By C. H. Samans, G. F. Tisinai, J. K. Stanley

The addition of nitrogen (0.10 to 0.20 pct) to Fe-Cr-Ni alloys of simulated commercial purity results in a real displacement of the u phase boundaries to higher chromium contents. The effect is small for the (Y + s)? boundary, but is pronounced for the (y + a +s)/(y + a) boundary. Although there is an indication of an exceptionally large shift of the n boundaries to higher chromium contents, especially in steels with nitrogen over 0.2 pct, the major portion of this apparent shift results from the fact that carbide and nitride precipitation cause "chromium impoverishment" of the matrices. The effect of combined additions of nitrogen and silicon to the Fe-Cr-Ni phase diagram is demonstrated also. Nitrogen can nullify the effect of about 1 pct Si in shifting the (y + o)/? phase boundary to lower values of chromium at all nickel levels from 8 to 20 pct. NItrogen can nullify this U-forming effect of about 2 pct Si at the 8 pct Ni level, but not at the 20 pct Ni level. The alloys studied were in both the cast and the wrought conditions. There are indications that the u phase forms more slowly in the cast alloys than in the wrought alloys if both are in the completely austenitic state. The presence of 6 ferrite in the cast alloys accelerates the formation of U. Cold working increases the rate of o formation in both cast and wrought alloys. THE major improvement in Fe-Cr-Ni austenitic alloys in recent years has been in the addition or removal of minor alloying elements to facilitate better control of corrosion resistance, sensitization, and heat resistance. One shortcoming of the austenitic Fe-Cr-Ni alloys, which never has been completely circumvented, is their propensity toward u formation. In the AISI-type 310 (25 pct Cr-20 pct Ni) and type 309 (25 pct Cr-12 pct Ni) steels, sufficient amounts of u phase can form, if service or treatment is in a suitable temperature range, to cause severe embrittlement. Also, there is a growing conviction that this phase may be contributory to some unexpected decreases in the corrosion resistance of certain of the 18 pct Cr-8 pct Ni-type steels. The present paper discusses the effect of nitrogen additions on the location of the (r+u)/d and the (y+a+u)/(y+a) phase boundaries in the ternary Fe-Cr-Ni system, for cast and wrought alloys of simulated commercial purity, and in similar alloys containing up to about 2.5 pct Si. The objective is to define compositional limits for alloys which will not be susceptible to u formation when used near 1200°F (650°C). An excellent review of the early studies of the u phase in the Fe-Cr-Ni system has been compiled by Foley.1 Rees, Burns, and Cook2 have determined a high purity phase diagram for the ternary system, whereas Nicholson, Samans, and Shortsleeve3 are- stricted themselves to a portion of the simulated commercial-purity phase diagram. Both groups of investigators show almost an identical position for the commercially significant (y+u)/y phase boundary. Further comparison of the work of the two groups indicates that, below the 8 pct Ni level, the commercial alloys have a decidedly greater propensity toward u formation than the high purity alloys. The two groups of workers agreed that both the AISI-type 310 (25 pct Cr-20 pct Ni) and the type 309 (25 pct Cr-12 pct Ni) steels are well within the (y+~) region and that the 18 pct Cr-8 pct Ni-type alloys straddle the U-forming phase boundaries. Nicholson et al.3 showed, in addition, that these boundaries shift toward lower chromium contents if greater than nominal amounts of silicon or molybdenum are added. The effect of nitrogen on the location of the s phase boundaries in the Fe-Cr-Ni system has not been known with any certainty. In 1942, an approach to this problem was made by Krainer and Leoville-Nowak,' but at that time they apparently were unaware of the slow rate of s formation in strain-free samples and aged their samples for insufficient times, e.g., 100 hr at 650°C (1200°F) and 800°C (1470°F). For this reason, it would be expected that their (y+ u) /y boundary would be shifted toward lower chromium contents (restricted ?-field) when equilibrium conditions were approximated more closely. Procedure for Studying the Alloys The alloys used were prepared in the following way: Heats of 200 lb each were melted in an induction furnace. A 5 lb portion of each heat was poured into a ladle containing an aluminum slug for de-

Jan 1, 1955

By P. S. Kotval, O. H. Nestor

Stacking fault energy measurements of nickel-base alloys have been mainly confined to binary and ternary systems. In this paper, the stacking fault energy has been measured by the rolling texture method in a series of ten alloys which comprise successive additions of Cr, Mo, Fe, and C to pure nickel, eventually resulting in an alloy of the composition of Hastelloy alloy X. The alloys studied here are single-phase, solid solutions with the exception of two alloys in which some undissolved particles of "primary" carbide have been retained. It is found that successive additions of chromium, molybdenum, and iron all lower the stacking fault energy, with iron having only a minor effect. The stacking fault energy is found to increase when carbon is added in solid solution. The results from the rolling texture measurements are correlated with thin foil observations of dislocation substructures in these alloys. In a recent paper' it was pointed out that the dislocation substructure of various superalloy matrices could be classified into three broad categories based on 'high', 'medium', and 'low' stacking fault energy. It has also been demonstrated2 that the dislocation substructure in each of these categories has a well defined role in the nucleation of strengthening precipitates which is different from the role played by the dislocation substructure in other categories. Thus, it becomes desirable to understand the influence of various solute elements on the stacking fault energy and hence on the dislocation substructure of the matrix, before any further development of superalloys by mi-crostructural predesign can be undertaken. Recently, Beeston and France have studied the influence of increasing solute additions on the stacking fault energy of a series of binary nickel-base alloys relevant to the Nimonic series using the rolling texture method, and have then estimated the effect of a given alloy addition in five commercial Nimonic alloys. However, comparison with stacking fault energy data from other investigations''5 suggests that the influence of a given solute element in a nickel-base binary system is not necessarily the same in a ternary or more complex superalloy system. Accordingly, the present work was undertaken to study the effect of successive addition of solute elements to pure nickel, the final composition being the nominal composition of Hastelloy X. The rolling texture method of stacking fault energy measurement was used since it can be used for the whole range of stacking fault energy values and does not have the disadvantage of, say, the Node method which is only applicable to low values of stacking fault energy. In addition, the rolling texture method provides a means of determining the stacking fault energy which is statistically more significant than that provided by other methods. EXPERIMENTAL TECHNIQUES Button heats of alloys of the compositions shown in Table I were prepared. Each button was remelted not less than four times. After a slight deformation (approximately 5 pct) all alloys were homogenized at 2200°F except alloys, H . I, and J. Alloys H and I were solution heat treated at 2150°F and alloy J at 2282OF. The buttons were cold worked by rolling, using "end-to-end" passes and intermediate anneals at the homogenization temperatures mentioned above. After each annealing treatment the samples were rapidly water quenched to avoid any precipitation. In alloys F and I, however, a few particles of "primary" carbides were retained even after the homogenization treatments at the temperatures mentioned above. Part of the solution heat treated material was cold worked to 0.04-in.-thick sheet and the penultimate reduction was -50 pct of deformation as recommended by Dillamore et al. All annealing was carried out in vacuo within sealed quartz capsules. Some of the material from each alloy was rolled down further to 0.004 in. strip for thin foil transmission electron microscopy specimens. Specimens of this strip were annealed at the homogenization temperature for 1 hr and then strained 7 pct by rolling at room temperature. Thin foils were prepared from the strip specimens by the 'window" technique using an Ethanol-Perchloric acid electrolyte at 32°F and a voltage of 22 v. Stainless steel cathodes were employed. All transmission electron microscopy was performed in a JEM-7 electron microscope using an accelerating voltage of 100 kv. Specimens from the 0.04 in. sheet which had been rolled -60 pct in the final pass were electropolished to remove the surface layers to a depth of approximately 0.002 in. Rolling texture pole figures for all the alloys were determined using a Schulz ring and nickel filtered CuKa radiation at 50 kv and 20 ma. The texture parameter Io/(lo + I,,) (where Io is the

Jan 1, 1970

By F. H. Wilson, E. W. Palmer

Brass mills are familiar with a recurring problem which reveals itself during deformation of annealed metal as an opening up of cracks which are suggestive of a grain boundary pattern. A typical example is seen in Fig 1, which shows part of the convex surface of a cartridge brass disc which was slightly dished by the punching operation. Another illustration of the same type of defect was found in the surface of a finished cartridge case which split open on firing. These cracks, shown in Fig 2, were away from the split but indicate the presence of the type of weakness which permitted the splitting. Usually the weakness consists of separate cracks, the lengths of which are of the same order of magnitude as the grain size prior to the final anneal. Their typical appearance in a polished section is shown in Fig 3. This structure was found below the surface of a fractured tensile-specimen that had been cut from a large cartridge blank showing the defects on its convex surface. While the pattern formed by these cracks suggests a grain boundary origin, the cracks bear no relation to the currently existing grain structure. Thus Fig 3 shows that the grains have grown up to cracks already present. Since it is doubtful that the cracks were present in the metal before the anneal, they must have formed during the anneal but prior to recrystallization. The term "fire-cracking" has been generally applied to cracking that occurs during an anneal under the influence of internal stress, but refers more specifically to obvious macroscopic cracking attributable to the presence of a low melting phase, usually lead. Since the type of cracking described above may occur when the lead content is very low, we have considered it a somewhat different phenomenon, and have called it "inter-granular parting." While most examples involve cartridge brass, intergranular parting has also been observed in the fabrication of large seamless tubes from discs of both 85/15 red brass and 70/30 cupro-nickel. Similar cracking in nickel silver was investigated by Jones and Whitehead,1 who showed that it could occur during healing or cooling. That which occurred on heating they called "fire cracking," and they suggested that, a transformation at about 320°C might account for its occurrence. It was felt in this laboratory that the observed parting was probably one aspect of the general observation, first made by Rosenhain and Archbutt,2 that a tendency toward intergranular fracture under tensile stress increases with increasing temperature and decreasing rate of strain. In this case the stress involved would be internal. The rate of strain would be exceedingly slow, a localized internal creep. Reference to the literature uncovered no efforts to extend the observations of Rosenhain to conditions involving only internal stress. Accordingly, an exploratory research, sufficient in detail to satisfy us as to the probable truth of this explanation of the observed intergranular parting, was undertaken. The internal stresses present during the early stages of an anneal (prior to recrystallization) would be from two sources: residual stresses developed during deformation, and thermal stresses caused by uneven heating and expansion. Thermal stresses would vary widely according to shape, size and manner of heating, and can be considered as supplementary to residual stresses. It seemed necessary to determine the stress and temperature conditions under which parting would occur in a relatively short time, and then to establish whether or not in- ternal stresses may persist to an extent adequate to cause parting at such temperatures. (Jones and Whitehead1 and Moore and Beckinsale3 made tests which indicated that stress relief required an appreciable time.) During the course of the research the desirability of studying the effect of grain size became apparent, and this factor is one of the major variables of our investigation. Intergranular Parting under Tension at Elevated Temperatures In early experiments, hard 70/30 brass tensile specimens, with unknown residual stresses, held for 10 min. at applied stresses from 30,000 to 40,000 psi at temperatures from 300 to 350°C, showed no macroscopically visible cracks when unloaded and cooled, but cracks very similar in appearance to those observed in commercial practice were revealed in these specimens by pulling them to fracture. Annealed specimens, on the other hand, showed no cracks in such experiments, apparently because they deformed plastically at stresses below those necessary to cause parting in a reasonable time. Hence, the specimens used for this study were first strengthened by cold tensile elongation, giving them an unavoidable residual stress pattern (determined largely by grain size and orientation) which would, however, correspond in direction, at least, to the stress pattern obtained under stress at temperature. A variation in these residual stresses might be expected with variations in grain size, but following any given anneal the reproducibility of stress pattern should be as good as that of the grain size measurement. With no information as to the effect of the amount of prior deformation, tests were conducted on specimens given the same cold elongation, using a stress applied at temperature which was a constant proportion of the stress required to produce this elongation. In

Jan 1, 1950

By K. Tangri, M. Chaturvedi

On water-quenching from within the (a + ß) phase region Zr-2.5 Nb and Zr-2.5 Nb-0.5 Cu alloys can undergo w transfirmation. This transformation has been attributed to the enrichment of ß Zr phase, at the solution-treatment temperature, by the solute atoms. The ß—w transformation is accompanied by a hardening 01 the alloys and in the optimum condition w-bearing specimens of the ternary alloy are about 14 pct harder than those bearing only a' phase. On aging crl reverls thermally to more stable phases, which results in a considerable softening of the specirnens, as follows: w quenched -waged + ß Zr (enriched) waged - a + ß Nb The high-temperature bcc ß Zr in certain zirconium alloys can transform to a metastable w phase which has a primitive hexagonal structure. In the Zr-Nb alloy system w phase has been observed in the composition range of 7 to 15 wt pct Nb by a number of investigators.'-' In this composition range w can form by water-quenching ß Zr or by aging. known as athermal and thermal w, respectively. Hatt and Roberts' have found that thermal w has a constant c/a ratio of 0.622 +0.002 regardless of the composition and tha: the lattice parameters are of the order of a = 5.02A and c = 3.0 ?. The actual lattice parameters depend upon composition and the temperature of aging. The lattice constants of athermal w. however, depend upon alloy composition.3 The presence of w phase in zirconium alloys can cause substantial strengthening which has been attributed as probably due to the coherency between w and parent ß Zr phase. Robinson et al.' found that the presence of w can raise the strength of Zr-5.0 Nb-2.0 Sn and Zr-5.0 Mo-2.0 Sn alloys to about 170.000 psi: however. the ductility is considerably reduced. With the growing use of Zr-2.5 Nb and Zr-2.5 Nb-0.5 Cu alloys in the nuclear reactors, and the significant effect of w transformation on the mechanical properties of zirconium alloys studies were undertaken to study the w transformation in Zr-2.5 Nb and Zr-2.5 Nb-0.5 Cu alloys. This paper is concerned with the formation and reversion of w phase in these alloys. 1) MATERIALS AND EXPERIMENTAL TECHNIQUES The following two alloys were studied: 1) Zr-2.5 Nb-0.5 Cu (referred to as the ternary alloy): 2) Zr-2.5 Nb (referred to as the binary alloy). The chemical analysis of the alloys is given in Table I. The alloys, in the form of 3/8-in.-diam rods and 1/8-in.-thick sheets, were supplied by the Chalk River Nuclear Laboratories of the A.E.C.L. Initial fabrication for preparation of various specimens was carried out by cold rolling and swaging with intermediate anneals at 1000°C. All the heat treatments were carried out after the specimens were wrapped in zirconium foils and encapsulated in silica tubes under a vacuum of 5 x 10-% m of Hg. For optical metallography and hardness measurements disc specimens + in. diam by + in. thick were used. After the required heat treatments the specimens were mechanically and then chemically polished in a 45 pct HNO3, 45 pct H20. and 10 pct HF solution. The hardness measurements on the chemically polished specimens were carried out on a Vickers hardness tester using a 10-kg load. For each specimen at least fifteen indentations were made in order to obtain a representative hardness value. The phase identification and structural analysis were carried out using X-rays and electron diffraction techniques. Wires of 1.5 mm diam reduced to 0.12 mm diam by chemical etching were used for making Debye-Scherrer powder patterns using copper K-a radiations in 114.6-mm-diam camera. Thin films for transmission electron microscopy were prepared by electropolishing heat-treated a by 7 by 0.005-in.-thick strips using a modified Bollmann-Window technique. The 10 pct perchloric acid-90 pct methyl alcohol electropolishing bath was kepl at -5OCC and polishing was done at 5 to 10 v. The thinned specimens were washed in ethyl alcohol at -30" to -40°C and dried between filter papers. The thin films

Jan 1, 1970

By H. J. Levinstein, H. J. Guggenheim, C. D. Capio

An optical incestigation of the phase transformations in RbFeF, has been conducted. Details of the ferromagwetic phase transition and the metamagnetic state are disczrssed. The three-dimensional magnetic domain structure existing in bulk crystals of RbFeF3 is described, as well as the effect of the magnitude and direction of the applied magnetic field on the domain struclure. The strong magnetoelastic interaction in RbFeF, is demonstrated. The results of the direct observation of the role of dislocation subboundaries and slip bands as nucleation sites and as pinning sites for magnetic domain walls are reported. THE perovskite compound RbFeF, has recently been shown to exhibit three magnetic states.''' Above 102°K it is paramagnetic, between 102" and 87°K it is antiferromagnetic. and below 87°K it is ferromagnetic (i.e., exhibiting remanence) with a modification of the magnetization occurring at 40°K. The changes in magnetic structure at 102o, 87°. and 40°K are accompanied by shear transformations to successively lower crystal symmetry classes.3 At 102°K a second-order phase transformation to a tetragonal crystal structure occurs. The c/o ratio increases with decreasing temperature to a value of 1.0034 at 87°K where the crystal undergoes another shear transformation to an ortho-rhombic crystal structure. The magnetic modification at 40 K is also accompanied by a shear transformation to a lower crystal symmetry. RbFeF3 is unique in that in the ferromagnetic state it is transparent in the bulk to visible light, has a low saturation magnetization, a large magnetic rotation. and has good optical properties.' All these features make it an ideal material for the investigation of magnetic domain structures in bulk crystals. In addition a dislocation etch has been developed which reveals the point of emergence of dislocations with the (100) surfaces of RbFeF3,5 making it possible to determine the dislocation arrays in the material. As a result domain wall dislocation interactions can be observed in the bulk crystal. In this paper we report on 1) the crystallography of the phase transformations in RbFeF3. 2) the domain configurations as a function of magnetic field and crystal orientation. 3) the interaction of dislocations and magnetic domain walls in RbFeF,. EXPERIMENTAL PROCEDURE Since the temperature range of interest is below room temperature the dewar shown in Fig. 1 was employed. It was designed to fit on the stage of a Leitz panphot metallograph, and permitted examination of the sample at a magnification up to 150 times. The dewar consists of a double-walled glass cylinder bent into an L shape. The space between the walls is evacuated. The viewing wirldows were made of four optically flat quartz discs aligned parallel to each other and sealed to the sides of the concentric cylinders. The specimens were mounted in a holder attached to a flexible plastic shaft. Various holder designs were employed depending upon the type of observations to be made. For studies of the phase transformation between the antiferromagnetic and ferromagnetic state the holder shown in Fig. 2(d) was employed. This holder permitted accurate temperature control at 82" to 88°K by balancing the heat input from the carbon resistance heater against the heat loss to a heat sink immersed in liquid N2. Magnetic field studies were conducted by employing the holders shown in Figs. 2(a) and (b). The holder shown in Fig. 2(n) has a solenoid imbedded in it. such that the magnetic field direction is in the same direction as the incident light. The magnetic field provided by the holder in Fig. 2(b) is perpendicular to the incident light direction. The holder in Fig. 2(r) was employed when observations of the specimen were desired while an elastic bending stress was applied to the sample. The stress was applied to the sample by pulling a wire from outside of the dewar. The wire was attached to a lever pivoted on the holder. which caused the knife edge of the lever to push against the sample. The crystal growth and the sample preparation were described previously.' PHASE TRANSFORMATIONS The first phase transformation in this system is from the cubic perovskite structure to n tetragonal

Jan 1, 1970

By H. F. Ramstad, F. D. Richardson

The equilibria (a) SiOz +Hz =SiO +H20 and (b) Si + SiO, = 2Si0 have beet1 studied at temperatures of 1425"to 1600°C ad 1310°to 1485°C respectively. The stattdard free energy changes for the tzrro reactions are given by the equatiotts Combination of the results for both equilibria leads to tiotz removes certain anomalies in existing high-terlzperature data for equilibria involving silica and silicon in iron. In many metallurgical processes and in many laboratory investigations silicon monoxide undoubtedly plays an important role. It is unfortunate therefore that wide differences exist between the results obtained by different investigators1-7 in their studies of such equilibria as In an attempt to put our knowledge of SiO on a surer basis, an exhaustive study has been made of equilibria [I] and [2] at temperatures ranging from 1300" to 1600°C. Reaction [I] was studied by measuring the amounts of silica which could be condensed from streams of Hz or Hz + HzO which had previously been brought into equilibrium with silica at temperatures ranging from 1425" to 1600°C. Reaction [2] was studied by measuring the material that could be condensed from streams of Hz or argon which had been brought into equilibrium with mixtures of silicon and silica at temperatures ranging from 1310" to 1485°C. EXPERIMENTAL Materials. The silicon was "superpure" grade and contained less than 0.1 pct impurities. The silica was prepared from pure mineral quartz; this was crushed and treated with concentrated hydrochloric acid to remove particles of iron, washed with water, and finally dried at 120°C. For the hydrogen + silica reaction, the silica was sized to —20+100 mesh. For the silicon + silica reaction, the two materials were ground to a fine powder in an agate mortar. The hydrogen and argon were commercial oxygen-free gases. The gas streams were controlled with capillary flow meters and the volumes were measured by wet gas meters. After passing through the meters, the gases were partially dried by silica gel. The hydrogen for the HZ + SiOz reaction was then passed through palladised asbestos at 300°C and dried with magnesium perchlorate. The efficiency of oxygen removal was checked throughout the experiments by passing the gas over an electrically heated strip of nichrome, used as an indicator as described by Rathman and de itt.' When mixtures of HZ + Hz0 were required, the partial pressures of water vapor (1.8 to 22 mm) were obtained by passing the hydrogen through oxalic acid dihydrateg' lo held at various controlled temperatures, O.l°C, by means of a water bath. The hydrogen for the Si + Si02 reaction was purified by passing it over a mixture of 3 parts of magnesium to 5 of lime heated to 600°."l1 u The argon for this reaction was passed through titanium powder (-3/16 in. + 100 mesh) heated to 900°C. The nitrogen used to prevent the reaction products escaping from the condenser (see later), was deoxidized by copper or iron at 600°C. All these gases were finally dried with magnesium perchlorate. Furnace, Temperature Contr01, and Measurement A molybdenum resistance furnace was used for both sets of experiments. The reactions were conducted inside a high-grade alumina tube, 36 in. long and 1 in. in diam as indicated in Fig. 1. With this arrangement an even temperature zone (2"C) 4 cm long was satisfactorily obtained. The temperatures were kept constant by means of a proportional controller actuated by a Pt-Pt 13 pct Rh thermocouple. This was placed between the two alumina tubes, so that the temperature at the junction was 1400" to 1450°C. Up to 1485"C, the temperatures were measured with Pt-Pt 13 pct Rh thermocouples. For higher temperatures an optical pyrometer was used, this being sighted (through the glass window 1 in Fig. 1) on the end of the alumina tube, that held the SiOz or Si +SiOz mixture, 10 in Fig. 1. The optical pyrometer was recalibrated whenever a change was made in any part of the apparatus situated in the hot zone. Successive readings with the optical pyrometer were reproducible to within 1"C. Equilibrium Apparatus and Procedure. Hydvogen and Silica Reaction. The apparatus is shown in Fig.

Jan 1, 1962

By I. Javandel, P. A. Witherspoon

The finite element method was originally developed in the aircraft industry to handle problems of stress distribution in complex airframe configurations. This paper describes how the method can be extended to problems of transient flow in porous media. In this approach, the continuum is replaced by a system of finite elements. By employing the variational principle, one can obtain time dependent solutions for the potential at every point in the system by minimizing a potential energy functional. The theory of the method is reviewed. To demonstrate its validity, nonsteady-state results obtained by the finite clement method are compared with those of typical boundary value problems for which rigorous analytical solutions are available. To demonstrate the power of this approach, solutions for the more complex problem of transient flow in layered systems with crossflow are also presented. The generality of this approach with respect to arbitrary boundary conditions and changes in rock properties provides a new method of handling problems of fluid flow in complex systems. INTRODUCTION Problems of transient flow in porous media often can be handled by the methods of analytical mathematics as long as the geometry or properties of the flow system do not become too complex. When the analytical approach becomes intractable, it is customary to resort to numerical methods, and a great variety of problems have been handled in this manner. One such method relies on the finite difference approach wherein the system is divided into a network of elements, and a finite difference equation for the flow into and out of each element is developed. The solution of the resulting set of equations usually requires a high speed computer. When heterogeneous systems of arbitrary geometry must be considered, however, this approach is sometimes difficult to apply and may require large amounts of computer time. The finite element method is a new approach that avoids these difficulties. It was developed originally in the aircraft industry to provide a refined solution for stress distributions in extremely complex airframe configurations. 27 Clough has recently reviewed the application of the finite element method in the field of structural mechanics.6 The technique has been applied successfully in the stress analysis of many complex structures.l, 27, 28 Recognition that this procedure can be interpreted in terms of variational procedures involving minimizing a potential energy functional7 leads naturally to its extension to other boundary value problems. In the field of heat flow, there recently have been introduced several approximate methods of solution that are based on variational principles.2-4, 17 By employing the variational principle in conjunction with the finite element idealization, a powerful solution technique is now available for determining the potential distribution within complex bodies of arbitrary geometry. In the finite element approximation of solids, the continuum is replaced by a system of elements. An approximate solution for the potential field within each element is assumed, and flux equilibrium equations are developed at a discrete number of points within the network of finite elements. For the case of steady-state heat flow, the technique is completely described by Zienkiewicz and Cheung.33 Since the flow of fluids in porous media is analogous to the flow of heat, Zienkiewicz et al. have employed the finite element method in obtaining steady-state solutions to heterogeneous and anisotropic seepage problems. 34 Taylor and Brown have used this method to investigate steady-state flow problems involving a free surface.25 The work of Gurtin has been instrumental in laying the groundwork for the application of finite element methods to linear initial-value problems.12 As a result, Wilson and Nickell have recently

Jan 1, 1969

By H. H. Rachford

Numerical solutions of immiscible flow problems in which dispersive effects of capillarity are dominated by convection require excessively fine grid spacing with attendant high computing costs. The use of coarser spacing reduces cost but often produces oscillation or undue dispersion associated with displacement fronts, A numerical formulation is proposed here which should be applicable to two-dimensional flow problems. it is in part analogous to an approach previously tested for miscible systems. The convective transport is approximated using a change of variables to yield a coordinate system moving approximately with the local characteristic velocity. The capillarity-induced dispersive terms in the differential system describing the process are approximated with respect to a fixed coordinate system by the usual implicit formulation. One-dimensional tests of the procedure yielded results in which the saturation profiles tended smoothly to the zero-capillary pressure solution as the ratio of viscous to capillary forces was successively increased in a sequence of calculations. This contrasted favorably with solutions by other numerical procedures which would require attendant grid refinements to approach the zero capillary pressure results. INTRODUCTION Numerical solution of displacement problems has until recently relied on applying methods developed primarily for transient heat-flow problems. Such problems are classified as parabolic in type, and where the heat transport is purely by diffusion their solutions are characterized by a high degree of smoothness. It is not surprising, therefore, that for approximating these solutions available finite difference methods are quite adequate. In flow problems the transport is partly by diffusion, partly by convection or flow. Although the problem remains of parabolic type because the dispersive effects of capillary forces or diffusion play some role in every displacement, at high flow rates the problem is dominated by convection, and solutions tend toward those of equations of the hyperbolic type. Solutions of hyperbolic problems are characterized by the translation of fronts, or discontinuities, that may progressively increase in sharpness. Numerical methods for treating parabolic problems become less and less satisfactory as displacement rates increase and the role of dispersion due to concentration or capillary pressure gradients becomes small relative to transport due to flow. In computation the difficulty manifests itself as an error associated with the grid size chosen. 1-6 In summary, if the heat-flow type approximations are to include the terms arising due to convection, one of several choices may be made: (1) an upstream (to the direction of flow) approximation for the convection terms may be used; (2) a centered-in-distance (CID) approximation may be used; or (3) a recently developed approximation based on the theory of oscillation matrices may be chosen.6 The last appears to have significant promise for one-dimensional flow problems; its extendibility to two or three dimensions is an open question. In either of the first two approaches, a suitably small ratio of v&/D must be maintained, where v is the velocity, & is the grid spacing and D the effective dispersivity in the direction of flow. In the first choice, the approximation of the convective part is only first-order correct and errors introduced appear as a numerically induced dispersivity of magnitude proportional to v?x. In the CID choice, the approximation can be second-order correct, but the difference formulation fails to satisfy the maximum principle unless a condition on v?x/D is met. Practically, this means that for high flow rates oscillatory solutions may result in the neighborhood of a front unless exceedingly small grid intervals are taken. While the procedure proposed by Stone and Brian4 permits a less severe limitation to be placed on this ratio, ultimately the flow rates increase relative to the dispersivity the oscillation obtains. Further, extensions of their approach to higher dimensional systems may be attended by considerable

Jan 1, 1967

By Che-Yu Li, J. L. Gregg, W. F. Brehm

Oriented, oxygen-doped niobium bicrystals were tested in liquid lithium. The grain boundaries were attacked preferentially. The depth of the penetrated zone varies as (time)2. The penetration was aniso-tropic, had a high activation energy, and increased with the increased oxygen doping level. A possible model was proposed to account for the experimental observations. 1 HE grain boundary penetration of a metallic system by liquid metal has been studied by several investigators. Their results are summarized by Bishop.' Most of these works show that the penetration by liquid metal corresponds to the phenomenon of liquid metal wetting. In the case of a grain boundary, wetting will occur when twice the solid-liquid interfacial tension is smaller than the grain boundary tension resulting in the replacement of the grain boundary by two new solid-liquid interfaces. Other possibilities exist; for example, the atoms of the liquid metal may diffuse into the grain boundary region due to chemical potential gradient. The gradient can be produced by impurity segregation or simply be due to the increase in solubility in the grain boundary region. The penetrated grain boundary in these cases may remain solid at the test temperature. The Nb-Li system has been of considerable interest because of its possible technological applications. For fundamental interest it provides a possibility of studying the grain boundary penetration process which is not controlled by the wetting mechanism. The pure niobium is not attacked by the liquid lithium, but if niobium containing more than 300 to 500 ppm oxygen by weight is exposed to liquid lithium, corrosion occurs at the solid-liquid interface and preferentially at grain boundaries. Previous investigators2-' have proposed that this preferential corrosion at grain boundaries is caused by oxygen segregation there, with subsequent inward diffusion of lithium to form a Li-Nb-0 compound. These investigators also found that the corrosion could be retarded by adding 1 pct Zr to the niobium to precipitate the oxygen as ZrO2 upon proper heat treatment. However, there are no quantitative data on the kinetics of the grain boundary penetration process to test the validity of the proposed corrosion mechanism. In this work an investigation of this penetration process in oriented bicrystals was made as a function of the oxygen doping level in the bulk niobium and the grain boundary orientation. A possible model for the penetration process based on the experimental results was proposed. EXPERIMENTS Oriented niobium bicrystals were grown by arc-zone melting oriented single-crystal seeds.7 These bicrystals contained simple tilt boundary. The [001] directions in the two grains were tilted about a common [110]. The bicrystals were 31/2 in. long and 5 by 4 in. in cross section with the straight, symmetric, planar grain boundary longitudinally bisecting the crystal rod. The bicrystals were doped with oxygen by anodically depositing a layer of Nb2O on the surface in a 70 pct HNO solution at 100 v, using a stainless-steel cathode. The specimens were homogenized by annealing in evacuated quartz tubes at 127 5°C. Oxygen content of the niobium was measured from microhardness values, after DiStefano and Litmman.' Supplementary checks were made with vacuum-fusion analysis.7 Individual test specimens cut from the doped bi-crystal rods, about by by % in. in size, were tested inside double jacket sealed capsules. The inner jacket was niobium, the outer was stainless steel. The niobium inner jacket eliminated the problem of dissimilar-metal mass transfer.' The lithium (99.8 pct pure, obtained from Lithium Corp. of America) was handled only in a purified argon atmosphere in a Blickman stainless-steel glove box. After introduction of lithium, the capsules were sealed by welding. Further detailed experimental procedures are given in Ref. 7. The capsules were heat-treated in vertical Marshall resistance furnaces. Temperatures were controlled to When heating above 1100°C, it was necessary to seal the furnace work tube and flow argon through to prevent failure of the stainless-steel outer jacket of the capsule. Tests were made on 6" 2", 16" 2, and 33" i2" bicrystals at oxygen levels up to 2600 ppm by weight in the 6' and 16" crystals and with 1300 ppm oxygen in the 33' crystals. The oxygen levels were controlled to 100 ppm. Most of the quantitative data were obtained from 16" bicrystals between 800" and 1050°C. The capsules were quenched into water after the test and cut open with a water-cooled abrasive wheel. The capsules were then submerged in water, which dissolved the lithium and freed the specimen. Measurement of the depth of the penetrated zone in the grain boundary was done either on metallographically prepared surfaces or directly on the grain boundary plane after the specimen was fractured in tension in the grain boundary plane. The depth of penetration measured by both methods agreed well. Further details describing these techniques have been reported elsewhere.'p7

Jan 1, 1969