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By P. D. Zemany, G. W. Sear, B. W. Roberts

Vanadium metal is embrittled by hydrogen at a temperature as low as 250°C when held in the presence of manganese metal and water vapor in a rough vacuum. It is established that the property changes are caused by the catalytic decomposition of water vapor at the vanadium surface and the diffusion into and solution in the vanadium of the resultant hydrogen. It is found that manganese is a necessary component of the catalyst. The manganese is transported in the vapor phase by an unknown molecule. A deuterium tracer experiment demonstates the role of water vapor in the embrittle-ment process. VANADIUM metal foils were observed to become embrittled' at a temperature of about 300 °C when held in the presence of manganese metal and a small amount of moist air, This paper describes the investigation to find the embrittling agent and an understanding of the relatively low temperature reactions that are involved. Experimental The vanadium metal foil used was prepared by cold-rolling and pack-rolling 32 mil sheet" in a series of steps down to 1 mil foil. The original observation was confirmed by sealing vanadium foils of 3 x 10 sq cm into individual Pyrex tubes with manganese powder† and a con- trol tube containing only the vanadium foil. These tubes were evacuated to 10 -5 mm Hg without baking and sealed. After heat treatment for 200 hr at 300°C, the control foil showed no change in duetility, whereas the foil contained in the manganese— containing tube was embrittled. The visual appearance of each was unchanged. A series of Pyrex sample tubes, about 2.5 cm diam and 25 cm long, were prepared, each containing a 3 x 10 sq cm piece of foil and 5 g manganese powder at the lower end of the tube. By reducing the time of anneal and the temperature of these samples, it was found that embrittlement could be created at 250°C in a time as short as 1 hr. Since the vanadium metal used here has been drastically cold-worked by rolling, it is assumed that it contains a maximum number of dislocations. To check the possible necessity of dislocations in this low temperature reaction, a vanadium foil sample was annealed in Vycor for 2 hr at 800°C to re crystallize and reduce the dislocation concentration. Metallographic examination showed grains which were not visible before annealing. The embrittlement procedure was carried out at 300°C and 3 hr. Upon checking the foil no embrittlement was observed. Further experiments demonstrated that about 6 hr at 300°C are required to create embrittlement in the foil. This delay in the onset of embrittlement in the vanadium foil suggests but does not prove that dislocation channels play a role in the embrittlement phenomena. If manganese metal is necessary for this low temperature embrittlement, do other elements in the transition metals group yield the same result? To check this qualitatively, a group of elements of similar atomic radii were obtained and sealed as before into Pyrex tubes with a sheet of vanadium foil. These tubes were annealed at 250°C for 6 hr and included (with radii)-2 A1 (1.4A), As (1.25A), Be (1.2A), Co (1.25A), Cr (1.45A), Cu (1.25A), Fe (1.25A), Ga (1.2A), Ge (1.25L%), Mn (1.3A), Ni (1.25A), Si (1.2A), Ti (1.45A), Zn (1.3A), air, H,O, 10 cm Hg of dry hydrogen, and MnO, powder. Upon testing the above sample foils for brittleness, only the manganese-containing tube yielded a brittle foil. Manganese Transport—To eliminate contact of manganese metal powder and vanadium foil, sample tubes were prepared with fritted glass barriers. The embrittlement reaction was still found to occur. Thus, the mode of transfer of manganese is certainly vapor transport. A vanadium foil was embrittled by this mechanism in an evacuated Pyrex tube for 8 hr at 300°C. By means of X-ray fluorescence analysis,' the amount of manganese added to the surface was established at 5 ±2 x 10 -6 g per sq cm. Since the average rate of manganese deposition is known, an effective average pressure of an assumed carrier compound can be computed. ___ P = M/T v2p mkT

Jan 1, 1959

By C. S. Smith

The solubility of tantalum at 1100°C is 0.025 pct in pure copper, 1.2 pct with 20 pct Ni, and 2.7 pct with -30 pct Ni. The solubility decreases with temperature, and the alloys are precipitation hardenable. A 79/20/1 Cu-Ni-Ta alloy reaches maximum hardness after aging at about 750°C and, if cold worked, does not recrystal-lize below that temperature. The alloys have good tensile properties at moderately elevated temperatures and, since they can be hot and cold worked nearly as easily as cupronickel, they are suggested for service at temperatures above the usual limits for copper alloys. MASSIVE tantalum dissolves extremely slowly in copper-rich alloys, and tantalum powder forms refractory surface layers which prevent its solution unless special precautions are taken. After many trials, an 80 pit recovery was consistently obtained by using a mixture of a finely divided tantalum powder with two or three times its volume of potassium tantalum fluoride (K2TaF1,-melting point about 750°C) poured in a slow stream directly on to the surface of the molten copper alloy, so that each particle would remain coated with flux and be immediately and individually wetted by the molten metal. A 50-50 nickel-tantalum alloy has a melting point of about 1400°C and small lumps of it slowly but satisfactorily dissolve in molten copper-nickel alloys as does commerical "ferro-tantalum." The high affinity of tantalum for carbon makes it necessary to avoid carbonaceous crucibles. COPPER-TANTALUM ALLOYS Castings with good shrinkage resulted when 0.1 pct Ta or more was added to copper melted under charcoal and cast in air. Copper in two-pound melts to which 0.1, 0.15, and 0.2 pct Ta had been added retained residual amounts, by analysis, of 0.025, 0.023 and 0.026 pct, respectively. The solubility of tantalum in molten copper at about 1200 °C is therefore about 0.025 pct. The electrical conductivity of these three samples was 100.48, 100.00 and 100.20 pct IACS in the annealed condition, and 98.24, 97.98 and 98.08 in the cold-drawn condition. Wires 0.080 in. in diam withstood from thirteen to nineteen reversed bends after annealing for 30 min in hydrogen at 850°C and the metal was therefore completely deoxidized. The annealing temperature corresponding to 50 pct loss of work hardness in 1 hr in a strip cold rolled 77 pct reduction was 175C, compared to 230°C for equivalent undeoxidized material. It is unusual for a decrease of recrystallization temperature to result from an alloying addition, and the tantalum probably combines with and removes some minor impurity in the original copper that restrained its recrystalliza- tion. No difficulty whatever was encountered in hot or cold rolling or drawing these ingots, and were tantalum not so expensive it would make an excellent deoxidizer for copper. COPPER-NIcKEL-TANTALUM ALLOYS Studies were made of the solubility of tantalum in a large number of copper-rich alloys, but of these significant solubilities were found only in the case of an iron alloy (which, however, separated into two immiscible liquid layers) and alloys of copper and nickel, in which tantalum is freely soluble and which were found to have interesting properties. The latter alloys are the subject of U. S. Patent No. 2,430,306 (1947). They are not at present available commercially. The Ternary Constitution Diagram—The copper-rich corner of the constitution diagram was constructed on the basis of microscopic examination of samples of various composition and treatments. The alloys were hot rolled to 0.25 in. from castings 0.625 in. thick, then annealed for 1 hr at 1000°C, quenched, cold rolled to 0.10 in. and finally annealed for various periods of time at temperatures between 800" and 1100°C in a nonoxidizing atmosphere and quenched. When the solid-solubility limit had been transgressed, particles of a compound believed to be Ni2Ta*were

Jan 1, 1960

By R. C. Hassinger, D. V. von Rosenberg

Transverse dispersion has received considerably less treatment in the literature than has longitudinal dispersion. Different methods for determining transverse dispersion coefficients have been used in different investigations, and the results obtained have not been consistent enough to permit accurate generalizations as to the effect of various physical parameters on the magnitude of these coefficients. A numerical solution to the differential equation describing transverse dispersion in the absence of longitudinal dispersion was obtained to enable one to calculate the dispersion coefficient from experimental results. The more general dispersion equation including longitudinal dispersion also was solved numerically to give quantitative limits of a dimensionless group within which the assumption of negligible longitudinal dispersion is justified. Possible experimental procedures were examined, and one utilizing a cylindrical packed column was chosen for the determination of transverse dispersion coefficients. Values of these coefficients were determined for a system of two miscible organic fluids of equal density and viscosity, for two sizes of packing material over a wide range of flow rates in the laminar regime. The dispersion coefficient was found to decrease, for a constant value of the product of packing size and interstitial velocity, as the size of the packing material particles increased. INTRODUCTION Longitudinal dispersion has received extensive treatment in the literature, and consequently is better understood than its orthogonal counterpart, transverse dispersion. Many mathematical models of dispersion processes assume that transverse dispersion is rapid enough to damp out any radial** concentration gradients and therefore may be neglected. Laboratory and production results, however, indicate that this is a poor assumption.l,2 Various experimental procedures for determining transverse dispersion coefficients have been used in previous investigations, but the results have generally been expressed by similar correlations. The transverse dispersion coefficients obtained, however, have often varied considerably for given values of the correlation parameters. We feel that further experimental determinations of transverse dispersion coefficients will help alleviate some of the inconsistencies in these empirical correlations. One assumption implicit in all previous investigations is that of negligible longitudinal dispersion in the experimental system. An attempt to justify this assumption often is made using intuitive reasoning, but it is apparent that this reasoning must break down as the condition of zero flow rate is approached. A mathematical examination of the equations describing the system yields physical limits outside of which the assumption of negligible longitudinal dispersion is invalid. BACKGROUND In a porous medium, the "effective molecular diffusivity" De is less than the molecular diffusivity D measured in the absence of a porous medium, due to the tortuous path which a diffusing molecule must travel. Various authors3-5 have reported values of the ratio De/D in the range of 0.6 to 0.7. When there is fluid flow within the porous medium, mass transfer occurs by convective dispersion as well as by molecular diffusion. These are separate phenomena and can be treated as such on a microscopic scale. However, the mathematical complexity is such that only extremely simple geometries could be considered, and the results hardly would be applicable to the complex geometries existent in actual porous media. A more

Jan 1, 1969

By E. J. Rapperport, J. P. Pemsler

The thermodynamic activities of zinc in six solid solution Cu-Zn alloys ranging from 5 to 35 at. pct Zn were determined experimentally in the temperature range 400° to 600°C. This low temperature investigation was canducted in order to evaluate techniques developed to utilize the inherently high sensitivity of atomic absorption flocesses in the measurement of thermodynamic activities. Analytical expressions ,for the activity and actizlity coeflcient are given for the six alloys in the temperature ranges investigated. RELATIVELY few experimental methods are available for investigation of thermodynamic activities of alloys, especially in the solid state. The techniques most frequently used have been the electrochemical potential and the effusion methods, both of which have severe limitations in many instances. It is therefore desirable to expand the ability to perform such measurements in order to obtain new information as well as to provide an additional independent verification capability. In this work, we present a significant improvement in the spectrophotometric method for sensing small vapor pressures in static absorption cells. Similar techniques have been used previously;1"5 however, applications had been limited to relatively high pressures, often greater than 1 torr. Prior investigators have, for the most part, used broad spectral sources such as xenon or mercury lamps, and high intensity arcs. Hollow cathode sources were first suggested in 1956 6 and were used soon afterwards.4'5 These sources offer significant improvements in sensitivity and freedom from interfering spectral lines.'-' EXPERIMENTAL High purity zinc was obtained from Cominco Products, Inc., and copper from American Smelting and Refining Co. Both elements were of 99.999 pct purity. Copper-zinc alloys were vacuum melted in a high fired carbon crucible with each alloy pulled from the melt as a 4 -in. diam bar. The bars were swaged to -1/4 in. rods and vacuum annealed for 160 hr at 800° + 1°C. Samples for gross chemical analysis were taken at intervals along the length of the rods to ascertain the axial zinc gradient. Electron microprobe analysis of homogenized specimens indicated that the alloys had uniform compositions over their cross sections on a macro (200 p) and micro (1 u) scale to better than *1 pct (20) of the gross composition. This tolerance was determined by counting statistics, rather than assured composition fluctuations. All SiO 2 windows were high-ultraviolet-transmission grade to minimize intensity losses. Silica absorption cells were scrupulously cleaned consecutively in organic solvents, dilute HF, and distilled water before use. The empty cells were then flamed while under a dynamic vacuum, cooled, and removed to an argon-filled glove bag. Alloy pieces were cut and filed in the glove bag to produce fresh surfaces, and then loaded into the cells. The loaded cells were temporarily sealed, removed from the glove bag, reevacuated to 10-5 torr or better, and permanently sealed. The instrument used is schematically shown in Fig. 1. The spectral emission from a commercially made hollow cathode lamp (A) of a selected element is focused through an absorption cell (B) inside a well-controlled furnace (C). The intensity of the transmitted beam is measured using the spectrometer* (D) 'Techtron model AA4 atomic absorption spectrometer. which contains a grating (E) that disperses the light prior to impingement on the photomultiplier (F). The monochromator grating is adjusted so that only the wavelength of interest is measured. The power supply delivered an interrupted voltage to the lamp, causing a chopped radiation output to be transmitted. The detector read only the intermittent component of radiation incident upon it, so that all continuous noise signals (furnace radiation, and so forth) were eliminated. Three recording thermocouples contained in the muffle furnace were positioned along the length of the absorption cell: one at each end and one at the center. An effort was made to keep the ends of the cell several degrees hotter than the center to avoid window condensate. Appropriate thermal corrections were then necessary to relate cell pressure to radiation attenuation. Water-cooled heat shields, as shown in Fig. 1, were found to aid signal stability by protecting the hollow cathode and the photomultiplier from furnace radiation. The furnace had a 2-in. diam muffle, Kan-thal wound, with SiO 2 windows at its ends to minimize convective effects. The hollow cathode radiation was masked and focused to form a conic beam that was a maximum of { in. diam within the furnace. Thus, the 1.5 in. diam absorption cell easily contained the entire beam. The furnace was mounted on ball-bearing slides with positive positioning detents. This arrangement allowed the removal of the entire furnace assembly from the radiation path, position [I], Fig. 1, so that frequent sampling of the unattenuated beam intensity could be obtained. In all cases the beam intensity was kept constant to 0.1 pct as judged by readings taken immediately before and immediately after data collection. Only data for absorptions of less than 80 pct were utilized, as systematic deviations from linearity were found for greater absorptions.

Jan 1, 1970

By A. M. Skov, L. F. Elkins

First response to large-scale water flooding in the fractured very low permeability Spraberry sand has led to a new unique cyclic operation. Capacity water injection is used to restore reservoir pressure. This is followed by marly months production without water irzjection and the cycle repeated. Expansion of the oil, rock and water during pressure decline expels part of the fluids but capillary forces hold much of the injected water in the rock. At least with reservoir pressure restored and with partial water flood development, field performance has proved this cyclic operation is capable of producing oil from the nzatrix rock at least 50 per cent faster and with lower water percentage than is imbibition of water at stable reservoir pressure. INTRODUCTION The Spraberry Field of West Texas presents unusual problems for both primary production and water flooding. Extensive interconnected vertical fractures in the fractional-md sandstone permitted recovery of oil on 160-acre well spacing, but they made capillary end effects dominant. Primary recovery by solution gas drive is less than 10 per cent of oil in place. The concept of displacement of oil from the sand matrix by capillary irnbibition of water has led to field techniques which promise greatly increased oil recovery. Free exchange of laboratory research, reservoir information and results of field pilot tests among the various companies has been very important in development of this technology. Five units covering a total of 170,000 acres have been formed for water flooding, and 10 other areas covering an additional 175,500 acres are in various stages of unitization. Part of the Driver Unit reaching fillup first has demonstrated very unusual waterflood behavior and indicated numerous operating problems that will develop within and among the various units. SPRABERRY ROCK AND PRIMARY PERFORMANCE The Spraberry, discovered in February, 1949, is a 1,000-ft section of sandstones, shales and limestones with two main oil productive members: a 10-15 ft sand near the top and a 10-15 ft sand near the base. In part of the field some thinner intermediate sands are oil productive, and others are water bearing. All sands have permeabilities of 1 md or less and porosities of 8-15 per cent. Ordinary core analysis and electric and radiation logs are ineffective in differentiating between oil productive and nonprcductive sands. Sands capable of containing producible oil are best identified by mercury injection capillary pressure measurement and, in some cases, by core water saturation. About 3,500 wells have been drilled in the 500,000-acre trend. Vertical fractures were observed in practically all Spraberry cores. Continuity and interconnection of fractures were confirmed by pressure interference among wells during early development.' Major fractures trend northeast-southwest as indicated by oriented cores and confirmed by five fluid injection tests, by analysis of the pressure transients observed during development,''' and by three interference tests in the Driver Unit Water Flood reported herein. Fracture spac- ing probably averages inches to a few feet. Spraberry wells typically produced 100-400 BOPD initially after hydrauLic fracture treatments. By 1962 oil production had declined to an average of 12 bbl/well/day, near the economic Limits of operation. Reservoir pressure had declined from 2,300 psi initially in the Upper Spraberry and 2,500 psi in the Lower Spraberry to 500-1,000 psi. Partial closing of the fractures with declining reservoir pressure is believed to be the cause of such low oil production rates at these relatively high reservoir pressures. Cumulative recovery of 208 million bbl of oil is 80 to 90 per cent of that recoverable by primary means. Performance of the entire reservoir is summarized in Fig. 1. IMBIBITION WATER FLOODING By 1952 reservoir performance indicated low primary recoveries. Most engineers, expecting serious channeling of injected fluids through the fractures, held little hope for secondary recovery. With its extensive background of research on the fundamentals of fluid flow within reservoir rocks, Atlantic's Research and Development Division on short notice in 1952 conceived that displacement of oil by capillary imbibition of water into the rock might significantly increase Spraberry recovery. Laboratory data reported by Brownscombe and Dyes scaled to probable reservoir conditions showed potential waterflood recovery equal to or greater than primary recovery with a 10-15 year flood life.= A pilot test using three 40-acre injection wells, one central producing well and 18 surrounding observation wells demonstrated technical feasibility of the process. Injection of 1.5 million bbl of water from November 1952 to August 1955 proved water entered the rock and displaced oil

By E. Douglas Sethness

The uranium industry is booming. In Texas alone, there are about 22 different companies with active exploration programs. Twelve solution mines have been permitted; three surface mines have been authorized; and two mills are currently in operation. However, the industry also has a problem, and that is the disposal of radioactive wastes. Over the past several years, stories concerning nuclear wastes have appeared frequently in the news. One of the most frequently cited cases occurred in Grand Junction, Colorado. In 1966, after ten years of investigations, the U. S. Public Health Service (PHS) discovered that tailings from a uranium mill were being used as fill material and aggregate for local construction purposes. It was estimated that between 150,000 and 200,000 tons of material had been removed and used under streets, driveways, swimming pools, and sewer lines. In addition, tailings had been used under concrete slabs and around foundations of occupiable structures. Further studies prompted the Surgeon General to warn that the risk of leukemia and lung cancer could be doubled at the measured radiation levels. More recently, the L. B. Foster Company discovered that its building site in Washington, West Virginia, was radioactive. While digging a foundation, the ground erupted and a ball of fire 30 feet high shot out. Evidently, the dirt was laced with radioactive thorium and zirconium, a potentially explosive mixture contained in a Nigerian sand which had been used by the previous site owners in the manufacture of nuclear fuel rods. Just this month we have read about legal suits to stop exploration for a nuclear waste disposal site in Randall County, Texas. The U. S. Department of Energy is trying to locate a deep underground nuclear waste depository for final burial of over 76 million gallons of high-level wastes. The problem is acute, the wastes are accumulating at a rate of about 300,000 gallons per year. Nor do these numbers include the spent fuel elements from nuclear power plants that are in temporary storage facilities. Fortunately, public awareness of these and other related issues is high. Unfortunately, the differences in the waste products from the nuclear fuel cycle are not always apparent to the general public. There are two distinct types of radioactive wastes: "high-level", which consist of spent fuel or wastes from the reprocessing of spent fuel; and "low-level", which, in general, are by-product wastes. There are numerous non-technical definitions that can be applied to help the layman differentiate between high-level and low-level wastes. For this latter purpose, it is best to think of them in terms of what we can see and feel. In general, high-level wastes are physically hot and can cause acute radiation sickness in a short period of time. Low-level wastes are not hot, but may cause chronic health effects after long exposure. The wastes which we are concerned with in the uranium mining and milling industry are low-level wastes. As recently as ten years ago, there were very few controls or regulations governing tailings disposal methods. At the same time, mine reclamation was not enforced through either state or Federal laws and the long-term viability of abandoned tailings ponds was not assured. The regulatory climate has changed significantly in the last decade, however. The low-level radioactive wastes generated by uranium mining and milling are generally contained in a tailings pond. Approximately 85-97% of the total radioactivity contained in uranium ore is present in the mill waste that goes to such tailings ponds. The isotope Radium-226 is probably the most potentially harmful radioactive parameter in the ponds. Radium emits gamma radiation and is also an alpha particle emitter. Because gamma radiation is very penetrating, it presents a potential health problem when a source is located external to the body. Gamma radiation will go through the body, causing damage to each cell encountered on the way. Although alpha particles have very little penetration capability, they can cause extensive cell damage. For this reason, alpha particles are a problem after inhalation or ingestion. Radium creates a health hazard by both of these mechanisms. Radium decays to radon gas which can be inhaled and serve as an alpha particle emitter. Additionally, radium is very soluble and readily enters the natural hydrologic cycle if allowed to leach from a tailings pond. With a half-life of 1620 years, radium has plenty of time to be taken into the food chain and end up in our bodies, emitting alpha particles. Because the potential health problems are better understood today than ten years ago, and because the Nuclear Regulatory Commission (NRC) has developed increasingly stringent government regulations, the uranium mining industry applies a high level of technology to the disposal of nuclear wastes. In most cases, low-level radioactive wastes are disposed of at or near the site where they are produced. There are six commercial burial grounds for low-level wastes, but it would not be economical to ship all mine or milling wastes to these sites. The on-site disposal methods most often used are ponding

Jan 1, 1979

By John W. Weeton, Max Quatinetz, Robert J. Schafer

Nickel powders of 1-, 0.4-, and 0.2- average particle size were combined with 0.05- p MgO powder. Oxide was added in quantities of 4, 12, and 20 vol pct. The mixtures were vacuum hot pressed and extruded at 2100°F. In spite of a general tendency to growth of oxide particles during processing, dispersions of fineness comparable to those of Sintered Aluminum Powder, SAP, were achieved. A small amount of oxide considerably improved 1800°F stress-rupture strength, but further oxide addition resulted in a decrease of strength. RECENTLY there has been much speculation concerning the requisites for strong and stable dispersion-strengthened alloys. The nature of the dispersed phase, the stability of the alloys, and variables influencing the mechanical properties of alloys have been given consideration. Such topics have been reviewed by Bunshah and Goetzel. Of the many parameters related to the mechanical properties of dispersion alloys, mean free path between dispersed particles (a measure of the fineness of the dispersion), particle size, and vol pct of the dispersed phase have been shown to be very important. Generally, strength increases with decreasing mean free path, and there is an optimum vol pct of dispersed phase. Just how fine a dispersion ought to be is somewhat open to debate, but Irmann, Lenel, Backensto, and Rose, Gregory and Grant,4 and others have shown in the original SAP-type alloys, aluminum plus aluminum oxide alloys, that better strength is generally obtained with finer dispersions and the best strengths were obtained when the mean free path was less than . Further, Cremens, and rant' showed that in nickel plus aluminum oxide alloys, stress-rupture strength at 1500°F was definitely increasing with decreasing mean free path. In their investigation, however, the finest dispersions that were produced were about 2.0 mean free path. To determine the potential of the nickel plus oxide type of product, it would be neces- sary to obtain dispersions analogous to those of SAP alloys; namely, ones where the mean free path between oxides is less than . Unfortunately, however, it is very difficult to produce extremely fine dispersions of oxides in metals other than those which naturally form refractory oxide skins. There are several approaches to this problem including such procedures as mechanical mixing, internal oxidation, reduction of mixed oxides, and metal deposition from solution. Work at the Lewis Research Center has shown that metal powders 0.1p or finer could be produced by ball milling, when selected grinding aids were utilized.' Because these unusually fine powders were available, it was an initial objective of this study to determine if by mixing of these fine Ni powders with various amounts of MgO (0.05) a fine dispersion of MgO in Ni comparable to those in Al-Al203 might be achieved. A second objective was then to attempt to relate the parameters of structure to 1800°F stress-rupture strength of the alloys. Nickel powders of 1-, 0.4-, and 0.2-p average particle size were combined with 0.05- average particle size MgO powder. This oxide was added in quantities of 4, 12, and 20 vol pct. The mixtures were hot pressed and extruded at 2100°F.

Jan 1, 1962

By E. Martinez

Samples of chrysocolla, a hydrated copper silicate, from several sources were submitted to differential thermal analysis (DTA) and thermal gravimetric analysis (TGA). Pure samples of chrysocolla are difficult to obtain, but comparison of the thermograms revealed that certain reactions occurred in all the samples and it is believed that these are due to chrysocolla. Endothermic reactions occurred at 50' to 175°C, 350° to 650°C, and 980° to 1100°C each of which caused a weight loss. The first was due to loss of absorbed water, the second to a dehydroxyla-tion, and the last to conversion of cupric oxide to cuprous oxide. In addition, two exotherms were detected at 685° to 695°C and 935° to 950°C. Infrared spectra of heated samples indicated the exotherms were due to changes in the Si-O bonds leading to the formation of cristobalite. The results suggest that thermoanalysis may be a more sensitive and reliable method for detection and identification of chrysocolla than X-ray diffraction, infrared spectroscopy, or optical techniques. The recovery of copper from oxide ores has been the subject of much research in recent years. One of the principal oxide copper minerals is chrysocolla, a hydrated copper silicate. The information available indicates that it has widely varying properties depending on the origin of the sample. Chrysocolla usually is found associated with cuprite, malachite, azurite, and native copper. Most of these other minerals are recoverable by present day flotation practice, but chrysocolla is lost in the tailings.1,2 Reagent combinations that are successful in laboratory chrysocolla flotation with ore from one source have been found to be ineffective with another. 3 The copper segregation process offers an altemative to beneficiation of oxide copper ores.4'5 The process is of particular interest in processing ores that cannot be treated by conventional hydrometal-lurgical methods because of high acid consumption.6 Fuller understanding of the properties of chryso- colla may prove of value in the search for improvements in both the flotation and segregation processes. Chrysocolla samples from Arizona, Nevada, Chile, Australia, and South Africa were studied by differential thermal analysis (DTA) and thermal gravimetric analysis (TGA). In addition, a sample from Miami, Ariz., was submitted to X-ray diffraction, infrared spectrophotometry, and surface area determination. EQUlPMENT Differential thermal analysis (DTA) detects, amplifies, and records exothermic-endothermic reactions occurring in a sample as its temperature is raised at a constant rate, usually 10°C per minute. A unit made by the Robert L. Stone Co. was used in this study. The tests can be run if desired with a gas streaming through the sample. A thermobalance records weight losses or gains in a sample as its temperature is raised at a uniform rate. The gravimetric analyses (TGA) were run on a Chevenard thermobalance converted electronically for graphic recording. A heating rate of 5°C per minute up to 1100°C was used in these tests. A Perkin Elmer Infrared Spectrophotometer, Model 221, was used with sodium chloride optics obtaining IR spectra in the 2 to 15 region. A pressed pellet technique with potassium bromide was used in preparing the samples. DESCRIPTION OF SAMPLES Chrysocolla is found in the oxidation zones of copper deposits commonly associated with malachite, azurite, and limonite. It is described as varying in composition and in color from bluish green to brown or black. Inclusions of other minerals are usual so chemical analysis of chrysocolla samples may be suspect. However, it is considered by many to be a solid solution of CuO, SiO2 and H2O with a general formula of CuO.SiO2.2H2O or 2CuO.2SiO2. 3H20. Chukhrov7 has described a sample from the Urals as CU3,5(OH)2 (A1Si3) O10.nH2O and stated that chrysocolla is a montmorillonite-type mineral. On the basis of X-ray analyses, DTA, and infrared spectra of chrysocolla from the Inspiration Mine, Ariz., Sun8 concluded that there was no substantial evidence to classify it as a montmorillonite mineral. Chrysocolla samples from many sources throughout the world were used in this investigation. The following is a brief description of each:

Jan 1, 1963

By Robin O. Williams

AN investigation has been made of the Ni-Cr system for the purpose of elucidating certain points, namely the nature of aging in both terminal solid solutions and the nature of the phase diagram. Information pertaining to solubilities and precipitation has been obtained. Experimentation Five alloys, Table I, were arc melted in a cold copper crucible using electrolytic chromium and car-bony1 nickel, both dry hydrogen treated. These 100 g buttons were homogenized 24 hr at 1300°C in dry hydrogen and air cooled. Powders were prepared by filing or pulverizing and subsequent heat treatment was done in vacuum or helium using titanium chips as a getter. Powder of —80 mesh was filed from the 60 pct Ni alloy quenched from 1000°C and was sealed in silica under vacuum using a 250 °C outgassing. After aging as indicated in Table I1 the lattice parameters were measured on the quenched samples using the standard cos" 6 extrapolation. These parameters are considered accurate to roughly 0.0001A. In all cases chromium lines of 2.8812 ± 0.0005Å at 30°C were found. Drastic quenching from sufficiently high temperatures produced very sharp body-centered-cubic lines in the first four alloys without indications of transformations. Temperatures to 1250°C were used. Solid samples less than 1/16 in. thick were quenched in water without transformation and the powders could be adequately quenched in small helium filled thin wall silica tubing using a water quench. For those powder samples which were quenched from the two phase field the relative intensity of the body-centered-cubic lines and the face-centered-cubic lines were estimated and extrapolated to give the indicated solubility data in Fig. 1. The data for the two higher alloys were somewhat limited, the plotted points being the lowest temperature where no nickel phase was found. Neither filing, abrading, pulverizing, nor cooling to —190°C produced any new diffraction lines for solid samples quenched from the single phase region, nor did the character of the body-centered-cubic lines change Single phase body-centered-cubic powders likewise did not change on cooling to —190°C. Also, samples which had some precipitation due to inadequate quenching showed no additional changes under these conditions. The first change apparent by X-ray diffraction form samples quenched almost fast enough to prevent precipitation was the diffuseness of the body-centered-cubic lines, particularly on the low angle side, For slower cooling rates the diffuse face-centered-cubic lines appeared. Work on the large grained castings showed profuse streaking through some of the Laue spots while oscillating patterns showed broad body-centered-cubic and face-cen-tered-cubic lines as well as some new lines. For the 23.6 pct Ni alloy the new lines corresponded to 2.16, 1.96, and 1.86A and were more similar in character to the face-centered-cubic lines than the body-centered-cubic lines. Samples which were air cooled gave only face-centered-cubic and body-centered-cubic lines which were still broad. One pattern indicated that face-centered-cubic (111) plane was parallel to a body-centered-cubic (110) plane. For those samples which were examined by light microscopy there were details which were not resolved. However, varied and beautiful structures were obtained. Fig. 2 is of an alloy quenched in a helium filled silica tube from the single phase region and shows particles associated apparently with dislocations which are arranged in low angle boundaries. Finer, general precipitation has also taken place within the grains. Figs. 3 to 5 show the variety of structures produced in these alloys on continuous cooling. It appears that there are four distinct modes of precipitation as evidenced by these, figures. Annealing these structures at higher tem-peratures in the two phase field gives structures as shown in Fig. 6, which shows nickel plates in the chromium matrix which reprecipitated nickel on a much finer scale of the final quench. Lower annealing temperatures and shorter times naturally give finer plates of the nickel-rich phase. Samples of the first four alloys were annealed for appreciable times between 900º and 1250°C and gave structures like Fig. 6. The relative amounts of the two phases were measured and extrapolated to give solubility data as indicated in Fig. 1. The point at 1250°C was deduced from data of Oxx.1 These and most of the other samples were checked for ferromagnetism but none was apparent. In fact, it appeared that the magnetic susceptibilities were not more than three times that for paramagnetic chromium. Discussion In Fig. 1 it is seen that the solubility of nickel in chromium can be represented by a slightly curved line on the usual log X vs 1/T plot, These data are believed to be accurate to roughly 5 to 10º. There is only fair agreement with the data of Taylor and

Jan 1, 1958

By Nicholas J. Grant, Ulf Kalling, John Chipman

THE operation of a blast furnace is dependent to an important extent upon the sulphur content of materials charged and the desired limit of sulphur in the product. It has long been known that the blast furnace is the most efficient tool for desulphurization in common use and that this efficiency is associated with the strongly reducing conditions of the hearth and is enhanced by increased basicity and fluidity of the slag. The chemical reactions of desulphurization may be studied from the viewpoint of the ratio of the process or of the final equilibrium conditions. Both kinds of studies contribute to an understanding of the process and both are included here. A simple measure of the desulphurization power of a slag is given by the ratio: Pct sulphur in slag (Pet S) Pct sulphur in metal [Pct S] This ratio was used by Holbrook and Joseph',' to measure relative desulphurizing powers under controlled laboratory conditions. It was also used by Hatch and Chipman3 as a measure of the equilibrium distribution. For the latter purpose it would be preferable to employ thermodynamic activities rather than percentages, but until very recently this has been impossible for lack of data. Now, thanks to the work of Morris and Williams and Morris and Buehl," the effects of carbon and silicon upon the activity of sulphur in the metal are known. The confirmation of this work and its extension to include the effects of other elements by Sherman and Chipman and by Rosenqvist and Cox' make it possible to calculate the activity of sulphur in pig iron of any composition. Hence it is now possible to use data on the equilibrium distribution of sulphur to find its activity in the liquid slag and to approach an ultimate solution of the thermodynamic aspects of the problem. The rate of transfer of sulphur from metal to slag is the problem of major industrial importance and indeed the principal need for equilibrium data has been as a necessary adjunct to the kinetic studies. The rate of approach to equilibrium under laboratory conditions seems slow compared to the requirements of industrial practice, and it is clear that further laboratory studies of rates are needed. In the research reported below, the items which were investigated were the following: I—The role of mechanical stirring on the approach to equilibrium. 2—The role of MgO in desulphurization as compared to CaO. 3—The role of MnO in desulphurization. 4— The limiting reactions which constitute the slow steps in desulphurization. Experimental Procedure The experimental set-up and procedure previously described by Hatch and Chipman" were essentially followed with several small modifications. The graphite crucible containing the slag and metal charge was altered to provide considerably more active stirring and mixing of the slag and metal in the carbon monoxide atmosphere. For this purpose the crucible was machined to provide two deep cylindrical wells which were interconnected at top and bottom as shown in Fig. 1. A graphite screw with a flat thread and of shallow pitch (4 threads per in.) spinning at 600 to 800 rpm was used to lift the slag and metal over the partition between the two wells and throw them over into the second well, where the metal settled through the slag into the reservoir at the bottom. It was possible to see actual particles of slag and metal being thrown over the partition. In this respect, the stirring was more vigorous than used in the work of Hatch and Chipman. A charge of 400 g of wash metal was first melted, and 20 g of FeS was then added to yield a bath containing 1.65 pct S. Immediately 400 g of slag (as pure mixed oxides) was added and fused. The slag was generally fused in 1 hr * 10 min. Within 30 to 45 min after melting, the temperature was adjusted to 1525"C, and the first slag and metal samples were taken. The slag was picked up on the end of a cold Armco iron rod, whereas the metal was sucked into a silica tube. The wash metal composition was (in percent): 4.29 C; 0.022 S; 0.021 P; 0.38 Si. The slags used were of four fixed starting compositions covering a wide range of acid-base ratios shown in Table I. Deliberate variations in MgO were made in these slags to check the role of MgO in blast-furnace desulphurization. Changes due to additions and reactions were followed by analysis of samples. Additions of Mn and MnO were made to most of the heats to note the role of Mn and MnO on desulphurization. Three heats (62 through 64) were made in an open pot induction crucible (graphite) using a

Jan 1, 1952

By D. Harker, M. J. Murphy

The mechanism by which age-hardening takes .place is still not completely understood. The principal theories range from the extreme of "precipitaiion-hardening" to that of "order-hardening," with many intermediate gradations. In the hope of obtaining new data on which to base a choice among the many theories, the authors have made electron micrographs of formvar replicas taken from metallic specimens at various stages of age-hardening. The results so obtained on an alloy of 2 per cent beryllium in copper and one of 20 per cent molybdenum in iron will be described in the following pages, as well as the technique used in preparing the replicas. Preparation of Formvae. Replicas of Metallic Specimens This section is concerned with the techniques that have roved successful in preparing formvar replicas for studying the microstructure of metals with the electron microscope. Because of its simplicity and accuracy, the formvar replica1 method developed in the General Electric Research Laboratory has been used exclusively, since it is much simpler than the silica replica method of Heidenreich and Peck2 and produces as good results. The formvar method to be described requires a technique acquired only by practice, and no foolproof set of rules can be given. it is a rapid method and a replica can be obtained in less than five minutes under favorable conditions. This does not mean that perfect replicas can be produced every five minutes, but rather that one can well afford to take several strippings* in order to produce a replica of high quality. The metallographic technique required for replica work—very fine polishing and etching with very dilute solutions—is described here. It can be said, in general, that almost any surface that has been carefully prepared for a photomicrograph at 1000 or more diameters can be used to obtain a replica for the electron microscope. Preparation of Metallographic Specimens The preparation of the metallographic specimens is somewhat more exacting than that usually given to samples to be used for inspection under the light microscope. After much experimentation, it was concluded that the best method consists in polishing the metal through to the coarse cloth, then etching and repolishing on final cloth until the layer of distorted metal is removed. After this procedure has been completed, a dilute etch—to approximately the depth necessary for a micrograph at 1000 diameters—usually produces a sur-

Jan 1, 1945

By R. I. Jaffee, E. S. Bartlett, D. N. Williams

The oxidation resistance and fabricability of molybdenum powder-metallurgy alloys containing up to 20 wt pct nickel in the form of a fine dispersion were studied. The effects of dispersion distribution and of amount and type (pure nickel or nickel compounds) of additive were considered. A definite improvement in oxidation resistance was observed, but the amount of improvement fell far short of requirements for an oxidation-resistant alloy. Fabricability by hot-rolling was limited. INVESTIGATIONS at Ohio State university1 and Battelle Memorial institute2 have indicated that the attainment of a truly oxidation-resistant molybdenum-base alloy for high-temperature service, with the coincident retention of high-temperature strength and toughness, will be a difficult, if not impossible, feat. Among the most promising alloy systems studied was the molybdenum-nickel system shown in Fig. 1. Alloys in this system show good oxidation resistance at temperatures well above the normal melting point of molybdenum trioxide because of the regenerative formation of protective nickel-molyb-date (NiMoO4) on the surface of the alloy during exposure at high temperatures. Unfortunately, the required alloy content to obtain oxidation resistant molybdenum-nickel alloys far exceeds the limits of the narrow terminal solid-solution boundary, and closely approaches the peri-tectic composition involving the brittle intermetal-lic compound, MoNi. Because of the peritectic reaction, fusion alloys are characteristically brittle, since the MoNi phase, being the last to form, tends to occur as a continuous grain-boundary network around each molybdenum grain. An added problem with this system is that the protective nickel-molybdate layer formed at high temperatures is unstable, undergoing one or more allotropic modifications upon cooling. The volume changes associated with this modification during cooling are such as to cause violent expulsion of the protective molybdate coating from the surface of the piece. This is, of course, quite detrimental in material exposed to cyclic temperature during service. The research reported here resulted from an attempt to overcome the aforementioned problems associated with the molybdenum-nickel alloys primarily by controlling the amount, size, and location of the MoNi phase. By using powder-metallurgy techniques of alloy preparation and sintering the blended molybdenum-nickel powder mixtures at a temperature below the peritectic reaction temperature, molybdenum-nickel alloys with a randomly dispersed MoNi phase were produced. MATERIALS The molybdenum powder used in the preparation of alloys for this study was commercial hydrogen reduced, minus 325-mesh (Tyler) material. Master alloys, by which nickel was added, consisted of elemental carbonyl nickel powder (12 average particle size), an alloy of 50 wt pct Ni-50 wt pct Cr, and compounds of MoNi, Ni,Si, and Ni2B. The nickel-chromium powder was prepared by filing a vacuum-arc-melted ingot, separating impurities magnetically, and screening to minus 200-mesh. The MoNi powder was prepared by pulverizing a hydrogen-sintered powder-metallurgy compact to minus 325-mesh material. Further sizing of this powder for study of effects of dispersion distribution was accomplished by air classification of the minus 325-mesh powder. The Ni2Si and Ni2B powders were prepared by crushing vacuum-sintered powder-metallurgy compacts to minus 200-mesh material. EXPERIMENTAL PROCEDURE Specimen Preparation—The general procedure followed in the preparation of alloys for testing is described as follows: 1) Blending. Either hand-blending of the powders in a ceramic mortar and pestle for X 1/2 or mechanical blending was used. Hand-blending seemed preferable for optimum homogeneity. 2) Compacting. Most of the blended powders were pressed to green compacts at pressures of 15 tons per sq in. in a single-action die, although some were pressed at 80 tons per sq in. A 1 X 11/4 X 1/4-in. green compact was the standard size for oxidation test material, and 7 X 1/2 X 1/4-in. compacts were used for fabricability test specimens. 3) Sintering. All green compacts were sinteredat

Jan 1, 1959

By M. Kaufman, A. E. Palty

The phase structure and aging characteristics of two nickel-base alloys, Inconel 718 and 702, were investigated. Wrought and cast Inconel 718 showed ArisCb as the major hardening phase, as well as ture was positively identified. Inconel 702 had Ni3Al as the major hardening phase, as well as Cr7C3-of the phase amounts with temperature was determined. 1 HE value of phase structure and aging rate information in the utilization of alloys has been well established. As part of a continuing program, the results obtained on two nickel-base alloys, hconel 718 and Inconel 702, are presented here. Inconel 718 is a relatively new alloy considered for application in both sheet and cast form. Its novelty lies in the use of high columbium tantalum levels in a Ni-Cr-Fe base. No previous work on this type of alloy has been reported in the literature. Inconel 702 is being used now as an oxidation-resistant sheet alloy. Occasional fabrication problems spurred work on its basic structural behavior. The composition of 702 is essentially nichrome plus aluminum and titanium for strengthening and added oxidation resistance. This puts it in the basic class of other nickel-base alloys which have already been studied. The two alloys will be treated separately. INCONEL 718 Phase Study Methods. The phases present after the various conditions investigated were determined by X-ray diffraction analysis of two different elec-trolytically extracted residues and microscopic examination. The extraction methods, using 10 pct HC1 in alcohol and 10 pct H3PO, in water and the X-ray technique have been described previously.' The effect of temperature on the phase behavior was studied on sheet material using the following heat treatments: 2250°F, 2 hr, ice brine quench 1300 100 hr, water quench 2250°F, 2 hr, ice brine quench 1400°, 100 hr, water quench 2250°F, 2 hr, ice brine quench 1500°F, 100 hr, water quench 2250o 2 hr, ice brine quench 1700°F, 48 hr, water quench 2250°F, 2 hr, ice brine quench 1850°, 24 hr, water quench 2250°, 2 hr, ice brine quench 2000°F, 6 hr, water quench The high solution temperature was an attempt to dissolve as many phases as possible without causing melting. Aging exposures were made longer than normal heat treating times in order to approach equilibrium conditions more closely. Cast material was examined only in the as-cast condition and with the normal heat treatment (1700°F, 1 hr, air cool 1325oF, 16 hr,air cool). Materials. The full phase study was performed on sheet specimens approximately 1/2 in. by 2 in. by 1/16 in. The cast pieces were cut from the gating system of a vacuum-cast test piece. An all weld-metal specimen was made by running many weld beads on the sheet using filler made from the sheared edges of the sheet. All external contamination was removed by belt-sanding followed by an electro-polishing treatment. The typical composition of wrought Inco 718 is as follows: The alloy is air-melted. Cast hco 718 may have higher molybdenum. Due to the expectation of finding a Ni-Cb phase (Ni, Cb) for which no X-ray diffraction pattern was available in the ASTM Index, a special 5 Ib heat to the stoichiometric composition of this intermetallic

Jan 1, 1962

By J. H. Swisher, E. O. Fuchs

Rate equations were derived to describe the kinetics of internal oxidation of cylinders and spheres. The derived equations for cylinders were checked experimentally by means of sub scale thickness and electrical conductivity measurements on Cu-Cr alloy wires. The properties of the internally oxidized samples were examined with conductivity applications in mind. It was possible to produce uniform dispersions of Cr2O3 in copper with an initial chromium content as high as 3 wt pct. While electrical conductivities only a few pct less than that of OFHC copper were obtained, the Cr2Os particle size and spacing were too large for effective dispersion hardening. T.HE process of internal oxidation has been used widely in basic studies of the permeability of gases in metals. In a review article, Rapp1 has discussed the principles of internal oxidation in considerable detail. From a technological standpoint, internal oxidation is often considered undesirable, since it is a means by which inclusions can be introduced into an otherwise clean material. Another important aspect of internal oxidation is its use as a means of dispersion hardening a material. Broutman and Krock2 discuss this and other methods for making dispersion hardened alloys. The only internally oxidized material known to the authors which is commercially available is a Cu-BeO alloy.3'4 This alloy is made from Cu-Be alloy powder, using a so-called Rhines pack. It has a tensile strength of 80,000 psi and retains its strength at relatively high temperatures. The objectives of the present study were to derive rate equations for the internal oxidation of cylinders and spheres, to check the derived equations for cylinders experimentally, and to examine the structure and properties of internally oxidized Cu-Cr alloys. The Cu-Cr system was chosen for this study because uniform dispersions are obtainable at high alloy contents, which is a desirable characteristic in dispersion hardened materials. RATE EQUATIONS FOR VARIOUS GEOMETRIES A number of authors5--9 have derived equations to describe internal oxidation kinetics. These derivations differ somewhat in mathematical assumptions and approximations, and all except one of the derivations deal exclusively with the internal oxidation of plates. The exception is a brief treatment of cylindrical and spherical geometries given by Meijering and Druy-vesteyn9 as a part of a comprehensive paper on the general subject of internal oxidation. These authors did not obtain rate data to check their derivations, although they did show that the hardness profile across an internally oxidized sample is directly related to the rate of interface movement. For cylindrical and spherical geometries, a quasi-steady-state approximation is needed to circumvent mathematical complications in obtaining a solution to the basic differential equations. In using this approximation, we consider the concentration gradient of dissolved oxygen in the internally oxidized zone or sub-scale to be the same as the gradient which would be present if there were no movement of the subscale interface. The steady-state approximation introduces an error of about 1 pct in computing the rate of internal oxidation of an Fe-1.0 pct Mn alloy plate, if the present method is compared to the more exact method of Wagner.7'10 The details of the derivations of the rate equations for cylinders and spheres are given in the Appendix, and only the results of these derivations are given below. The final equations obtained by Meijering and Druyvesteyn9 can be shown to be equivalent to our Eqs. [1] and [2], although the two approaches are somewhat different. Cylindrical Geometry. [2] where r1 is the outer radius of the cylinder or sphere, cm, r2 is the radius of the unreacted core, cm, see Fig. l(a), D is the diffusion coefficient of oxygen in copper, cm2 per sec, %O is the concentration of dissolved oxygen at the surface of the specimen, wt pct, %Cr is the initial chromium concentration in the alloy, wt pct, and t is the reaction time, sec. Plate Geometry. The analogous rate equation for a plate has been derived previously for internal oxidation of Fe-Al alloys.8'11 For Cu-Cr alloys, we may write the same equation as follows: [3] where r1 is the half-thickness of the plate, cm, and r2 is the distance from the mid-plane to the subscale intherate is An analysis of Eqs. [1], [2], and [3] shows that for a plate the rate is completely parabolic. The initial

Jan 1, 1970

By J. H. Swisher

The solubility of sulfur and rate of solution of sulfur in pure Lron were measured in H2S + H2 and H2S + H2 H2O gas mixtures. The solubility and diffusivity of sulfur at 1410°Care 0.13 pet S and 1.0 x 10-5 sq cm per sec, respectively. The solubility iS the same, but the rate of sulfurization is slower in the presence of H2O in the reacting gas. Under these conditions, the over-all rate is controlled jointly by a slow surface reaction and by solid-state diffusion; the mechanism for the surface reaction has not been identified. KNOWLEDGE of the behavior of sulfur in solid iron is desirable for the metallurgy of such products as free machining steel, where a high sulfur level is required, and inclusion-free high-strength steels, where the sulfur specifications are very low. The present investigation was undertaken to check previously reported values for sulfur solubility and diffusivity in 6 iron, and to study the poisoning effect of chemisorbed oxygen on sulfurization kinetics in H2-H2S-H2O gas mixtures. All of the experiments were performed at 1410°C. The thermodynamic behavior of sulfur in 6 iron was the subject of a paper by Rosenqvist and Dunicz.' The sulfur solubility at 1400" and 1500°C was determined by equilibrating pure iron specimens with H2-H2S gas mixtures. The maximum solubility of sulfur in 6 iron was alsc determined by Barloga, Bock, and parlee2 by reacting iron wires with sulfur in sealed capsules. In another investigation, the diffusion coefficient of sulfur in 6 iron at temperatures up to 1450°C was measured by Seibel.3 The method used was to measure sulfur concentration profiles in diffusion couples containing radioactive sulfur EXPERIMENTAL Apparatus. A vertical resistance furnace wound with molybdenum wire and containing a recrystallized alumina reaction rube was used for the experiments. The hot zone in the furnace was approximately 2 in. long with a temperature variation of ±3oC. The hot zone temperature was automatically controlled to within ±2°C, and the test temperature was measured with a pt/Pt-10 pet Rh thermocouple before and after each experiment. Flow rates of the reacting gases were obtained using capillary flow meters. Materials. The source of H2S in the gas train was a premixed cylinder containing 5 pet H2S in H2. This mixture then was diluted with additional hydrogen and argon. In some experiments, water vapor was introduced by passing hydrogen and argon through a column containing 10 pet anhydrous oxalic acid and 90 pet oxalic acid dihydrate. The vapor pressure of water above this mixture is well-known.4 Argon was used as a diluent to minimize thermal segregation of H2S in the furnace5 and to reach higher H2O:H2 ratios than could be obtained in mixtures of H2 and H2S alone. Argon was purified by passage over copper chips at 350°C and subsequently over anhydrone. Hydrogen was purified by passage over platinized asbestos at 450°C and then over anhydrone. The H2-H2S mixture was purified by passage over platinized asbestos and then over P2O5. The specimen stock was made by melting and vacuum-carbon deoxidizing electrolytic "Plastiron" in a zirconia crucible. The principal impurities are listed in Table I. In some of the equilibrium experiments, six-pass zone-refined iron was used to minimize impurity side effects. This zone-refined iron had a total impurity level of about 25 ppm. Procedure. Specimens were annealed in hydrogen for a period of at least 2 hr at the beginning of each experiment. The specimens were held in the reacting gas for times varying between 10 min and 17 hr, and cooled to room temperature in a water-cooled stainless-steel block at the bottom of the furnace. The pH2S/pH2 ratios reported are those for gas equilibrium at 1410°C. Calculations based on available thermodynamic data8 showed that the only other gaseous8 species that formed in significant amounts in the furnace were S2 and S. Even when water vapor was introduced into the gas mixture, the concentrations of SO2, SO, and so forth, were negligible. The initial partial pressure of H2S was therefore corrected for its partial dissociation to S2 and S in determining the equi-

Jan 1, 1968

By R. G. Wheeler, J. W. Goffard

The Larson-Miller parameter was used to correlate time, temperature, and indentation creep of magnesium, aluminum, and some of their alloys. In the temperature range 300" to 450°C, the short-time Meyer hardness of pure magnesium was less than that of the magnesium alloys tested, but for long times the pure magnesium has greater indentation creep resistance. Aluminum (1100 alloy) had 1.5 to 2.5 times more indentation creep resistance than magnesium at 300" and 450oC, respectively. Hardening of aluminum with a dispersion of Al2O3 was effective in the time and temperature ranges studied. New technologies have required the development of new materials and the utilization of the more familiar materials for new and unusual applications. The use of magnesium and aluminum and some of their alloys, because of their desirable nuclear characteristics, light weight, low cost, and ready availability, has been extended to the 300" to 450°C temperature range. In this temperature range the basic consideration of these materials must be their rate of plastic flow rather than offset yield strengths. The indentation testing reported here arose from a need for design data for the load-holding ability of supports made of these materials. Test Procedure—Hardness indents were made with a 0.275-in.-diam quartz indentor and a 10.65-lb load. The indentor was made by fire-polishing a spherical surface on the end of a fused quartz rod. The samples were held at temperature in a graphite crucible controlled to ±2°C. A thermocouple was attached to the sample and test temperatures were recorded. The diameter of the spherical indentation was measured at the end of a test period and the compression stress (Meyer Hardness) was determined by: H___________load__________ m = projected area of indent Samples were 1 in. in diam and at least 1/4 in. thick. It was observed that at the higher temperatures and longer times, the quartz indentor would stick to the magnesium sample. The quartz indentor was, therefore, frequently inspected and fire-polishing repeated when necessary. The area of sticking was always a small fraction of the area of indent and was therefore considered to have an insignificant effect on results. Correlation of Hot-Indentation Test Data with Time-Temperature Parameter—Sherby and Dorn' have correlated creep or tensile data of a' solid solutions of aluminum with a temperature and strain-rate parameter suggested by Zener and Holloman. underwood2 used this parameter to correlate creep properties of some steels with hot hardness, and upon the basis of this correlation a means of obtaining creep properties from short-time (and inexpensive) hot hardness tests has been demonstrated. Since the validity of the correlation of creep properties with a time-temperature parameter and the correlation of creep properties with hot hardness have been shown, it follows that hot hardness may correlate with the time-temperature parameter. The hot-indentation data obtained was expressed as Meyer hardness, and was shown to be time and temperature dependent. Correlation of Meyer hardness, time, and temperature with the parameter was made using the relationship: Hm = Meyer hardness t = time, hours T = absolute temperature, OK K = constant A value for the constant K was calculated by equating In l/t + K/T at different temperatures and times but at the same hardness. The correlation was tested by plotting Hm vs the parameter, In 1/t +K/T. Since materials are being sought which have high hardness at low indentation creep, i.e., a high Meyer hardness for long time at high temperatures, low values of the parameter are ofthe most interest. TEST RESULTS Magnesium—Pure magnesium (99.98 pct) cut from extruded rod was indentation tested perpendicular to the rod axis at temperatures of 300°, 350°, 400°, and 450°C for times ranging from 6 sec to 112 hr. Fig. 1 shows the time dependency of Meyer hardness at the four constant temperatures. Fig. 2 shows the correlation of the Meyer hardness of pure magnesium with the time-temperature parameter using a K of 22,720 in Eq. [I]. At the bottom of Fig. 2, the effect of doubling the time of indentation t2 = 2(t1), on the abscissa for any time is shown graphically. This effect is of constant magnitude. Also shown graphically are the magnitudes of the effects on the

Jan 1, 1960

By G. O. Hiers

IN 1948 in the United States, 184,300 tons of lead was fabricated as coverings for electric power and communication cables. Such covering generally is called "sheathing" for the principal lengths of the cables, while the short lead encasements of the joints are called "sleeves." For the purpose of this paper, both lead sheathing and lead sleeving can fittingly be called lead tubing. In the United States in the same year, 34,262 tons of lead was consumed in making extruded lead pipes. Such pipe is employed frequently for water or chemical solution conveyance with gage pressures up to 100 psi. Consequently, the walls are thicker in ratio to diameter than is the case with the cable sheaths. For this discussion, lead pipes can be thought of as tubes or tubing. About 19 centuries ago, the Romans produced metallic lead of about 99.9 pct purity. Evidently this was cast in flat sheets of Y8-in. thickness, for instance. Such sheet of 10-ft lengths was bent around a removable core producing a tube somewhat oval in cross-section and with a typical internal diameter of, say, 3 in. This method of production necessitated a seam or longitudinal joint evidently made by casting the seam metal adherently. The seam metal, as analyzed recently, contains about 1 ½ pct Sn. Blow pipes were said to have been employed when the joint was made. Specimens of such pipe were found both in Italy and in the British Isles. It is believed the pipes conveyed water for drinking purposes and for baths. Numerous specimens are still in existence in a good state of preservation, due to lead's usual inertness in air or water. According to Claude E. Pearson,1 there were various stages in the evolution of lead extrusion presses for making tubes: I. First press by Joseph Bramah, 1797. 2-First successful press by Thomas Burr, 1820. 3-First bridge die press, J. and C. Hanson, 1837. In 1845, Wheatston and Crooke made the first lead sheathed electric cable. Flat lead strip was wound spirally around the insulated conductor. Overlapping edges of lead were soldered. Early cables were also made by pulling insulated conductors into 50-ft lengths of lead tube. Borel in France, and Wesselau are credited with devising the first presses for extruding lead sheaths directly on cables in 1879. In the United States, Eaton, in 188o, and Roberston, in 1885, introduced new presses of improved design. Following Cunningham's original idea of making lead traps and. bends in 1873, the presses apparently were improved and made successfully soon thereafter. Between the years 1930 and 1940, two continuous extrusion presses for lead were developed abroad. One apparently was originated by Dunsheath and made by the Henley Telegraph Works Co. The other was developed by the Pirelli-General Cable Co. CURRENT PRODUCTION In addition to the several descriptive chapters on lead extrusion in Pearson's book, the Metals Handbook2 recently described pipe, sheath, and collapsible tube production. The latter is by impact extrusion.

Jan 1, 1951

By George S. Ansell, John B. Hudson, Stephen A. Koffler

The permeability of hydrogen through the a phase of palladium has been measured by a low pressure permeation technique under conditions such that bulk diffusion was the rate-controlling process. The observed permeability is described by the equation: J = 1.80 x 1O-3 P½exp(-3745)/RT) cc(stp)/sec/cm2/en over a range of hydrogen pressure, P, from 2.9 x 10-5 m Hg to 5.0 x 10-3 cm Hg, and over a temperature range, T, from 300" to 709°K. The fact that the permeability shows a square root dependence on pressure and a reciprocal dependence on thickness was taken as evidence that bulk diffusion, rather than surface reactions, was the rate-controlling process. The permeability data were used in conjunction with the solubility data of Salmon et al., to determine the diffusivity of hydrogen through palludium as: D = 4.94 x 10-3 exp(-5745/RT) cm2/sec There was no influence of sub structural defects observed over the temperature range employed. From the permeability data obtained, coupled with grain size measurements, it was concluded that the ratio of grain boundary diffusivity to bulk diffusivity was less than 105 over the range of temperature investigated. THE diffusive mass transport of hydrogen in the a phase of palladium has been studied previously by numerous investigators.' In spite of the large amount of attention this system has received, there is not good agreement between the results obtained in different investigations.' This is due in part to the fact that the mass transport was surface-limited during some of these studies, rather than being diffusion-controlled3-5 The reason for the disagreement in other cases is not clear. These studies made use of such techniques as rate of absorption from solutions6 and gases,' electrochemical potential,8 time lag,' and permeation10,11 to determine the mass transport behavior. Of these, the gas permeability technique is the only method which allows an easy test to determine if diffusion is the rate-controlling mechanism, thus eliminating the uncertainty regarding the limiting transport processes inherent in the other techniques. The two most recent permeability studies are those of Toda10 and Davis." Toda determined the permeability of hydrogen in the a phase of palladium over the temperature range from 170" to 290°C, and over the pressure range from 36 to 630 mm Hg utilizing a steady-state gas-permeability technique. Toda's result was: J = 1.41 x 10-3 P½ exp(-3220/RT) where J = specific permeability in cc(stp)/sec/cm2/cm, P½ = square root of the inlet pressure in (cm Hg)½, R = Universal gas constant, and T = temperature in deg Kelvin. Davis11 also employed a steady-state gas-permeability technique over the temperature range from 200" to 700°C and over the pressure range from 0.02 to 760 mm Hg. His result for the permeability of hydrogen in the a phase of palladium was: J = 3.15 x 10-3 P½ exp(-A440/RT) In the range of overlapping temperature for these two investigations, the values of the specific permeability calculated from the above two equations differ by a factor of about 1.8. In the present investigation, the permeation of hydrogen in the a phase of palladium was determined over a wide temperature range, 27" to 436oC, and over the pressure range from 2.9 x 10-5 to 5.0 x 10-3 cm Hg. This temperature range overlaps that of the previous investigations of Toda and Davis, but also covers the lower temperature range which has never before been investigated. The lower pressure range used here avoided the interaction between the dissolved hydrogen atoms observed at higher hydrogen concentrations.' MATERIALS The as-received palladium specimens were cold rolled from a casting and were supplied as 5.08 cm discs of 0.508 and 0.762 mm thickness. According to J. Bishop and Company specifications, the composition of the discs was 99.95 pct Pd, the balance being Cu, Ag, Au, and Ir. In order to obtain samples of varying grain size, the as-received discs were then heat treated. Sample 1 was treated for eight minutes in a nitrogen atmosphere at 810°C and then air-cooled. Sample D was heated first to 550°C in helium. The helium atmosphere was then immediately replaced by hydrogen and the temperature was slowly raised to 1220°C and held for 22 hr. The sample was then cooled to 550°C where the hydrogen was replaced by helium and the disc was further furnace-cooled to room temperature. Sample H was given the same heat treatment, only with hydrogen substituted for helium below 550°C. The reason that hydrogen was not used throughout the entire annealing cycle for samples D and H was to prevent the distortion encountered by low-temperature cycling in hydrogen observed by Darling." After these heat treatments were completed, the

Jan 1, 1970

By Edwin B. Wilson, Joseph W. Leonard, Richard W. Borio

preparation of specification coals for the utility industry is approached from the standpoint of a cooperative effort with the power company to assure that the shipped product will be a noncorrosive coal. Selective mining, coal preparation, the use of additives, and coal blending are mine-site techniques proposed as means for upgrading corrosion-retarding constituents of utility coals (CaO and MgO) while, at the same time, reducing the concentration of those constituents of coal (Na20 and K,O) which cause corrosion. Data obtained from large-scale boiler testing in combination with bench-scale physical and chemical tests are used as a basis for the formulation of mine-site procedures to improve coal combustion utilization. Electric power generation remains the greatest market for coal in the United States. A major reason for this strong market has been the ability of the coal industry to supply a superior product for combustion while maintaining competitive prices. Now that the coal industry is engaged in intensive competition with nuclear power for future energy demand, it becomes ever more imperative that superior coal products are produced. One important problem when burning some coals is a liquid phase corrosive attack on boiler tube metal in the superheater and reheater areas of the furnace where steam temperatures may exceed 1050°F. In addition to causing power outages and high maintenance costs, corrosion is a continuing barrier to the future development of more efficient steam cycles which permit lower generation costs when using coal. Research methods of the past which have been directed at finding solutions which could be applied at the power plant site have been only partially successful and are generally costly. This paper contains a suggested --method of corrosion control through investigation of the problem at the source of the trouble—the coal supplier facility. It is proposed that mine-site pre-ventative procedures be considered as an effective means of cooperative corrosion control between the power company and the coal supplier. Economic justification must be applied to the individual situation. This paper is based on recent work by Combustion Engineering, Inc. and the Coal Research Bureau of West Virginia University under a contract from the Department of the Interior Office of Coal Research.',' Results which are of importance to this discussion may be summarized as follows: 1) The alkali minerals of coal (Na2O and K2O), as measured by an acid leaching process, can be related to corrosion rate in utility boilers. As these constituents increase in a coal product, its corrosiveness can be expected to worsen. 2) The alkaline-earth minerals contained in coal, reported as total CaO and MgO reduce corrosiveness 2s their concentration increases. 3) Naturally occuring, alkaline-earth-bearing materials, such as limestone and dolomite, may be added to a coal to reduce its corrosiveness. 4) Coals with very low iron contents tend to be less corrosive. These results provide the groundwork for examination of the system of operations which occur at the supplier facility from the seam face where mining is begun to the point of loading for shipment. The Mining System Fig. 1 is a cross section of one of the typical coal seams (designated coal A) that was boiler tested during this research program. This seam is overlain by material which consists primarily of a gray-black shale. In the lower third of this seam, there exists a thin rock parting of less than 4 in. in thickness which is consistent throughout the seam. The floor material is fireclay. In order to examine mineral distribution throughout this particular seam, the Coal Research Bureau collected a

Jan 1, 1971

By James R. Cost

The increasing emphasis upon the use of refractory metals has resulted in a closer look at some of the properties of these materials. In the course of a study of the Nb-N system it became necessary to re-examine one such property, the spectral emissivity, e. Experiments were performed in order to determine e for solid niobium as a function of both temperature and nitrogen content. The results of these experiments, which indicate a small but significant correction to the spectral emissivity of pure niobium, are reported herein. In addition, data on the total emissivity, et, of niobium and the spectral emissivity of the nitride Nb2N are given. To this author's knowledge, neither of these emissivities are as yet reported in the literature. Previous studies of the spectral emissivity of niobium have been few. In 1915, Burgess and Waltenburgl obtained a value of 0.49. Whitney? in 1935, used a Mendenhall V to obtain the commonly accepted value of 0.374. A recent investigation of sintering of niobium graphically reported a value of about 0.36= using a solid block of the metal with a blackbody hole drilled into it. In the present experiment a method similar to that of Worthing4 was used. An 8-in. long and 3/16 in. diam hollow tube was formed from smooth surfaced 0.0013-

Jan 1, 1962