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By A. E. Bellis

THERE is a wide difference of opinion among rifle experts in the matter of barrel steel, and the relative importance to the life of the barrel of the steel's composition, heat-treatment, structure, and physical properties. In order to promote discussion, which it is hoped may lead to the discovery of the source of 'some of the traditional beliefs on the subject, the results of some tests are presented here, with data on heat-treating, physical tests, and erosion. The detailed metallographic study of this steel is given in another paper.1 The conclusions from these tests should be capable of wider application than their direct bearing on the barrel-steel problem. The Government specifications for barrel steel for model 1903 rifle (Springfield) call for carbon 0.50 to 0.60 per cent., manganese 1.00 to 1.29 per cent., silicon under 0.25 per cent., sulfur under 0.06 per cent., phosphorus under 0.08. per cent. The minimum physical requirements are: Tensile strength,- 110,000 lb. per sq. in. (7733 kg. per sq. cm.); yield point, 75,000 lb. per sq. in (5472 kg. per sq. cm.); elongation, 20 per cent.; reduction of area, 45 per cent. Material fulfilling these requirement's gives practically no failures when proof-fired, with charge of 70,000 lb. per sq. in. pressure, which is 40 per cent. in excess of the service charge. The turning and drilling operations show up any seams or laps such as accidentally occur even in- the best heats of open-hearth steel, so that' faulty stock is practically all eliminated before proof-firing. These physical requirements are met either by rolling billets at, a proper temperature clown to the form of tapered barrel blanks or by heat-treating bars that have been upset to enlarge the breech section. The transformation temperature of this steel, on heating, is around 1350° F. (732° C.), varying somewhat with the composition and the rate of heating. The actual treatment of the barrel blanks, as carried out by different

Jan 7, 1919

Jan 1, 1935

Jan 1, 1935

Jan 1, 1935

Jan 1, 1935

Jan 1, 1935

By M. J. Taylor

A common method of classifying non-water- impounding mine waste structures is necessary to facilitate professional comnunications within the industry and between representatives of industry and government. Camnon classifications can a1 so he1 p quickly create and maintain a mental picture of the structures being discussed. Three categories for classification presented in this paper include: (1) configuration and location with respect to existing topography. (2) material characteristics of waste, and (3) environmental sensitivity or characteristic; of the disposal area. Configuration and location classification involve terminologies typically utilized for waste dumps over the past several years. General classification categories for non-impounding structures include: (a) valley-fill, (b) cross-valley, (c) side-hill, (d) ridge, and (e) heaped. Material characteristics can govern the location, engineering details, and environmental sensitivity of the waste structure. Material classification includes: (a) hazardous, (b) non-hazardous, (c) inert, (d) fugitive dust susceptible and (e) leachate susceptible. A1 though many characteristics are potentially applicable, this paper will focus on the categories mentioned above. Environmental sensitivity of a waste structure encompasses many objective and subjective evaluations plus many multi-disciplinary considerations. Environmental sensitivity incorporates numerous factors in addition to those highlighted in configuration and location, and material characteristic classifications. A few examples of how this categorization can govern the engineering design and location of waste dumps are provided in this paper. In addition to the verbal descriptions of classifications, a simplified version of the combined classifications with an alphanumeric format and its methodology is discussed in this paper. Surface water controls often govern the classification or vice versa. In order to maintain a non- impounding status, water must be prevented from collecting behind the structures. Depending on material characteristics, water inflow to or over the site can cause significant environmental or human consequences. Finally, the access of water to the waste dumps can play a large role in the degree of environmental sensitivity to the site. This paper provides a framework in which to discuss the other aspects of this conference and also provides insights as to how the framework is based on actual practice.

Jan 1, 1985

By Edwin Ludlow

IN CONSIDERING the waste in the coal industry, which is the title of this discussion, we must entirely eliminate the anthracite region. The demand for anthracite has been constantly increasing and the mines have been able to run during this last year, while the bituminous field was idle for two-thirds of the time. It is on a 100 per cent. basis, which is the best answer to any waste in industry investigations. The anthracite region is also in a very much better position on account of its control by very large companies which realize that their resources are limited; they must husband these resources in every way and they are making the most careful study of methods of mining and preparation to avoid any possibility of waste. This is well known and realized by the managers of those companies. Even the fine slush which was formerly thrown away is now being stored in the hope that methods will be found by which it can be used economically, either in briquet or pulverized fuel in coking. This has been tried. The anthracite factor may be dismissed from the waste in industry discussion, except as it applies to the general public that burns anthracite; there is a tremen-dous waste of anthracite in the burning of it. Looking at the ash piles where anthracite is burned and seeing the amount of carbon that comes away with the ashes, one realizes that proper efficiency is not obtained in the burning of that coal. Even its higher cost does not seem to bring people to a realization of the necessity of conserving and burning properly the entire coal carbon content of what has cost so large a price. Many people feel that they want to have the coal as large as they can shovel into the furnace; they do not realize that the small coal makes just as many heat units and can be bought at half the price. Consequently, the demand is continually in excess of the probable pro-duction for the larger sizes of anthracite, while the smaller sizes, such as large buckwheat and pea, are a drug on the market and can be bought at nearly half the figure; and yet, if properly burned, these produce just as much heat and can be used with a much greater economy than the large egg and stove.

Jan 3, 1922

By C. S. Sivil

To modern civilization the platinum metals are of inestimable value. Their distinctive properties, both physical and chemical, render them indispensable in an age in which the processes of the laboratory are becoming an essential part of the daily life of the world. Generally, their function, though vital to particular processes, is not spectacular, and so remains unknown to all except the operators. EARLY METHODS OF MANUFACTURE The properties of platinum were recognized as of prime importance for scientific work at. a very early date, probably between 1750 and 1760. The difficulty of making articles from it seemed almost insurmountable, because no means of attaining the temperature necessary to melt it were known at that time. It is true that solid platinum could be obtained by fusing the metal with a more volatile metal such as arsenic, which was subsequently driven off by baking, but such a method, at its best, would give only a very brittle product, which was totally unsuitable for the purposes for which it was intended. Little, if any, progress could be made until it was discovered that platinum sponge, heated to white heat and hammered in an iron or clay mold, would metallize and retain the shape of the mold. This was a process that could not be wholly satisfactory but from it developed the method of pressing the sponge in a mold and forging it at white heat into wire or sheet. At present a somewhat similar process must be used for fabricating articles of tungsten, molybdenum and other metals of high fusing point, as with these metals we face today the same problems that the scientists of the eighteenth century encountered with platinum, but possibly the next significant advance in the metallurgy of these metals will be the devising of means to melt and cast them.

Jan 1, 1931

By Barry D. Lichter, Pierre Sommelet

High-temperature heat contents of InSb, GaSb, and AlSb were measured over the temperature range 400" to 1450°K using a diphenyl ether drop calorimeter. Smoothed ualues of the thermal properties, H$ - H:9s, have been derived and are tabulated at even temperature intervals. The heats of fusion of the three compounds were determined as, respectively, 5707 *100, 7780 * 100, and 9800 5 3011 cal per g-atorn at the determined melting points of 797" l°, 985" * , and 1330" * 5°K. The calculated entropies of fusion are, respectively, 7.16 0.12, 7.90 * 0.11, and 7.37 * 0.23 cal per deg per g-atotn. The heat capacities increase substantially on tnelting in contrast to the behauior of structurally related germanium and silicon. Deriations from the Kopp-Neumann rule are negative for solid compounds and positive for the liquid phases. Previously obsevrrd "post melting" in InSb is confirmed. The high-temperature thermal properties of 111-V compounds are presently not well-established, despite the technical importance of these semiconducting cbmpounds. Uncertainties in available heats of fusion and heat contents have seriously hampered thermo-chemical evaluations' and thermodynamic analyses of phase equilibria2"" in these systems. This paper reports results of high-temperature heat content investigations of InSb, GaSb, and AlSb measured in the range 400" to 1450°K employing a diphenyl ether drop calorimeter. Similar measurements for InAs and GaAs will be reported in a following publication.~ EXPERIMENTAL PROCEDURES Samples. High-purity samples of the compounds InSb and GaSb were supplied in the form of crushed crystal fragments by Dr. Carl Thurmond of the Bell Telephone Laboratories and in the form of single crystals by Dr. A. Strauss of the M.I.T. Lincoln Laboratory. Single-crystal samples of semiconductor-grade AlSb were supplied by Dr. W. P. Allred of the Bell and Howell Research Center. Chemical analyses indicated that all compounds were stoichiometric to *0.1 at. pct, which is within the experimental uncertainties of the analyses. Samples were crushed, weighed, and encapsulated in evacuated, thin-walled, fused silica capsules. The capsules were nearly identical in external shape, 2 cm by 2 cm diam, but varied in weight due to differences in wall thickness. One sample of AlSb was contained in a thin-walled, high-purity alumina cup encapsulated in silica and used for heat content determinations of liquid AlSb. The capsule materials showed no visual evidence of reaction with any of the compounds. The sample and capsule weights are given in Table I. Calorimeter. A Bunsen-type calorimeter, similar in design to a previously described instrument6 but employing diphenyl ether (C6H5)'0 as the calorimetric substance, was used for measurements of heat contents above 300.0°K, the melting point of diphenyl ether. Heat input to the calorimeter caused isothermal melting of diphenyl ether, and the resulting increase in volume was measured by displacement of mercury from the calorimeter into a 200 cm horizontal, calibrated capillary, 1.25 * 0.01 mm diam, or into a weighed beaker. The advantages of diphenyl ether over water have been previously pointed out7 and include: i) an increase by a factor of 3.5 in sensitivity as measured by the ratio of the volume change to the enthalpy change on fusion, ii) the smaller required extrapolation from the melting point to the standard temperature of 298.17"K, and iii) the positive volume change on fusion of diphenyl ether in contrast to the contraction which occurs on fusion of ice. Diphenyl ether was purified by fractional crystallization to 99.95 mol pct as determined from the melting point depression with fraction crystallized. During assembly of the calorimeter, the ether was repeatedly outgassed under high vacuum to remove dissolved air. The calorimeter receiving vessel consisted of an 8-in.-long by 1:-in.-diam copper tube with twelve horizontal 3-in.-diam radiator "fins" for dissipating heat to the surrounding mantle of diphenyl ether. Before forming the mantle the chamber surrounding the receiving vessel contained 3300 cu cm of liquid diphenyl ether above 250 cu cm of mercury which was

Jan 1, 1970

By D. A. Munro, G. P. Contractor

Fourteen 50-lb laboratory melts were investigated to determine the effect of uranium on the tenpering characteristics of loo-carbon (0.06 to 0.1 pct C) steels. It was found that uranium additions, particularly in the range 0.30 to 0.45 pct, enhanced the hardness and both ultimate and yield strength of the experivzental steels in the quenched and tempered condition. The structural and morphological chazges indicated that uranium retarded tempering of the tnartensite, thereby hindering the normal formation of polygonal ferrite formed in the late stages of tempering. The effect of this was to make possible the re-tension of the acicilar ferritic structure in the uranium-bearing' steels. The iraniuin-bearing steels also showed IVidnzanstatten-type growth of ferrite plates and had large prior austenite grains containing assenzblies of fine ferrite grains, mainly acicular in geometry. The fine-grained ferrite structure and the presence of more numerous and apparently smaller precipitates in the uranium-bearing steels are thought to he principally responsible for the itnproved tensile strength and hardness of the experinzental uranium-bearing steels. At ternperirzg temperatures above 455% (850'F) the ferrite in the higher-uraniun steels nzaintained acicularity and, hence, its strength and resistance to tempering. Uranium did not produce a secondary hardening peak. However, it retarded softening during the third stage of tempering because of its effect of inhibiting the grouth of cementite particles and of retaining the acicularity of ferrite plates. The resistance to coalescence accounted for the slow grocth of the ferrite grains in the uranium-modified steels and, hence, fov the persistence of the acicular ferrite structure. IT had been found previously1 that uranium additions up to about 0.45 pct had no significant effect on the tensile properties of low-carbon steel (0.06 to 0.10 pct C) in the as-rolled and normalized conditions, Fig. 1. On the other hand, it was observed that uranium in excess of about 0.30 pct had an embrittling effect as revealed by Charpy V-notch impact results. It was also noted that, as the uranium content increased, the morphology of pearlite changed from lamellar to feathery and the ferrite grains showed an etching effect resembling striated or dashed markings, suggestive of precipitation. The sharp drop in the impact properties shown in Fig. 2 warranted an assumption that the uranium content of about 0.30 to 0.45 pct may produce some secondary hardening reaction on tempering, analogous to that associated with a Cr-Mo-V steel, which shows very poor CVN toughness at the secondary hardness peak in the tempering curve.1' With this background and the reported findings of Hasegawa and noda that low-carbon uranium-treated steel showed signs of secondary hardening, the present investigation was undertaken to determine the effect of uranium additions on the mechanical properties of 0.10 pct C steels. No attempts were made to investigate in detail the mechanisms of hardening, although some suggestions based on the experiments are made. MATERIALS AND PROCEDURES A series of 50-lb induction-furnace melts was made using AISI 1008 rimming steel billets as the melting stock. The melting, forging, and rolling techniques proven satisfactory in previous projects'-3 were employed as a guide for this investigation. The steel was deoxidized with aluminum (2 lb per ton) prior to the addition of high-purity uranium. The analysis of each melt is given in Table I. Properties were evaluated as a function of heat treatment and are presented in terms of hardness and tensile strength vs tempering temperatures. The variation of hardness with the tempering temperature was studied on the quenched and tempered specimens, some of which measured 0.50 by 0.25 in. diam and the others 0.40-in. cubes. Before quenching, the specimens were vacuum-sealed in glass tubes and normalized at 900°C (1650°F) for 20 min. Following this treatment, the sealed specimens were hardened by austenitizing at 955°C (1750°F) for 20 min and water quenching, and then tempered for 1 hr in the range 150 to 730°C

Jan 1, 1967

By W. A. Jr. Slippy, E. P. Dahlberg, R. B. Reed-Hill

A large room-temperature mechanical-hysteresis effect under cyclic tensile loading was observed in zivconium specimens prestrained at 77°K so as to form large numbers of (1121) twins. The observed hysteresis loss was a maximum when the prestrain was just under 1 pct It also depended strongly on the maximum applied cyclic stress, but only moderately on loading rate and temperature (-72°to 100°C). The phenotnena have also been studied in terms of the elastic aftereffect where the data has been found to cotlforln closely to the functional relationship between strain and time, c = -RT/a In tank ft. t to)/Zr, for strains cts large as about 10-4 This eyication can be derived directly from the strain-rate equation There may be important commercial implications to the present discovery since it appears that within rather wide limits it may be possible to obtain any desired magnitude of damping capacity in zirconium melal. The damping propevties may also be given directional characteristics. It is shown that the observed anelaslic phenomena can he explained on the assumiplion that they are the resuilt of stress-induced twin-boundary movements in which the average twin increases its thickness by only a few percent. In polycrystalline zirconium, a grain is least favorably oriented for slip when its basal plane, containing the (1130) slip directions, lies nearly perpendicular to the principle normal stress. Mechanical twinning is favored by this orientation and under simple tensile deformation at room temperature the primary twinning mode is {1012).' The (1012) twinning shear is small (0.17) so that, during deformation, twin growth does not greatly enhance ductility. At 77°K, on the other hand, the predominant twinning mode shifts to {1121), and many thin (1121) twins can form at small strains.' Twins thus nucleated can grow readily during subsequent room-temperature deformation and permit easy deformation in unfavorably oriented grains. In addition, the (1131) twinning shear is large: 0.63. These facts have lead to a process' for greatly improving the room-temperature ductility of zirconium when it contains a sizable fraction of grains with basal planes nearly normal to the stress axis. Optimum results are obtained by slightly less than 1 pct prestrain. The effect is large so that a 50 pct elongation increase is readily attained. A large room-temperature mechanical-hysteresis effect has also been observed in zirconium specimens prestrained at 77°K. The present paper is concerned with some aspects of this effect which the experimental observations strongly indicate is primarily the result of strain-induced {1121) twin-boundary movements. This clearly suggests that mechanical twins formed by plastic deformation can cause anelastic effects similar to those previously observed2 for annealing twins and twins associated with phase transformations. EXPERIMENTAL PROCEDURES All specimens were prepared from arc-melted sponge zirconium plate stock as previously described,3 which has a preferred orientation with basal planes aligned parallel to and uniformly distributed around the rolling direction. Both lonti-tudinal and transverse tensile specimens, with axes parallel and transverse, respectively, to the plate rolling direction, were cut from this plate. In a longitudinal specimen most grains are favorably oriented for slip, while in a transverse specimen a large fraction are unfavorably oriented. Prestraining at 77°K was performed in two ways. In one, annealed zirconium plate stock was cooled in liquid nitrogen and then rolled in the transverse plate direction to strains varying from 0.5 to 12 pct. Cylindrical 1/8-in.-diameter by 1-1/4-in.-gage-length tensile specimens were machined from this material. In the other, previously machined tensile specimens were prestrained in tension at 77°K between 0.2 and 3.8 pct. Type FA 25-12 SR4 strain gages were cemented to all specimen gage sections. All tests were performed on an Instron testing machine of 10,000-lb capacity. EXPERIMENTAL RESULTS Fig. 1 shows a typical two-cycle room-temperature stress-strain diagram for a specimen prestrained 0.65 pct by rolling at 77°K. The maximum stress in both cycles was 14,500 psi. Note that in the first cycle the loading and unloading curves are unsymmetrical so that a residual strain, er, remains after unloading. In the second cycle the stress-strain curves form a nearly symmetrical hystere-

Jan 1, 1965

By C. W. Arnold, H. L. Stone, D. L. Luffel

The purpose of this research is to determine (I) the efficiency of small banks of enriched gar driven by methane in displacing oil from a porous medium and (2) the effects of variation in bank size and composition of that efficiency. Most of the experiments were conducted in a sand-packed tube 20-ft long and 1/2-in. in diameter. The hydrocarbon system generally used was methane, butane and decane at 2,500 psia and 160°F. The results of these experiments indicate that, in the regions contacted by the gas, a small bank of an oil-miscible gas driven by methane can displace all of the oil in a piston-like manner. If the enriched gas is of such composition as to remain immiscible with the oil, displacement of oil is less efficient than for the miscible case, and the gas bank travels through the sand with a velocity less than that of the driving gas. These data along with theories discussed imply that smaller banks and less total gas are required when the enriched gas and oil are miscible. INTRODUCTION Widespread application of enriched-gas drive to the recovery of oil rests upon a key factor — the use of limited quantities, or "banks", of enriched gas. At the present time, the value of liquefied petroleum gas or other enriching agents discourages their use in a continuous injection technique, or even in a large bank, except in a few isolated reservoirs. If small banks of enriched gas driven by methane were as effective in displacing oil as is continuous injection, the enriched-gas drive process might be applied to a larger number of reservoirs. Previous research on the mechanics of the enriched-gas drive process reported by Stone and Crurnpl and by Kehn, Pyndus and Gaskell has utilized continuous injection of enriched gas. This work has shown that two types of displacements occur. With gases containing sufficient intermediates. the oil is displaced misciblv and complete recovery is obtained from the regions swept. When gases are used which contain insufficient intermediate hydrocarbon for miscible displacement, oil is displaced immiscibly. In the latter type, selective solution of the intermediate hydrocarbons causes a swelling and reduction in viscosity of the oil and leads to an increased recovery over that obtained by dry-gas (methane) drive. The size of the enriched-gas bank necessary for efficient displacement of oil is determined by those factors which cause deterioration of the bank. A differentiation may be made between those factors which operate on a microscopic scale and those which act on a macroscopic scale. On the smaller scale, the enriched gas mixes in the direction of flow by diffusion and convection with the fluids immediately preceding and following it. On the larger scale, the gas may by-pass the oil by flowing through permeable streaks, by overriding the oil because of density difference, or by fingering because of unfavorable viscosity ratios. In such cases, the enriching material tends to mix with the oil both laterally and in the direction of flow. The increase in effective area available for diffusion and dispersion of the enriching components leads to a faster degradation of the bank and a need for a larger bank than is necessary for those cases in which no by-passing occurs. The effects of such macroscopic factors in the deterioration of enriched-gas banks have been reported in a separate paper by Blackwell, Terry and Rayne. The present study was confined to the factors which operate on the smaller scale, in particular to the behavior of banks of enriched gas in sands uniformly swept by the gas. Experiments were designed to answer the following questions. 1. Can small banks of enriched gas driven by methane be used to secure oil recoveries comparable to those obtained by continuous injection of enriched gas? 2. What is the optimum bank size (the minimum bank size necessary to obtain a recovery comparable to that obtained by continuous injection of the same enriched gas)? 3. How many total pore volumes of gas must be injected to obtain the maximum recovery when the optimum bank size is used? 4. What is the effect of varying the number of enriching components in a gas bank? This report describes the experimental investigations and discusses the results in terms of their significance to reservoir behavior.

Jan 1, 1951

Jan 1, 1951

Jan 1, 1970

By F. W. McQuiston

INTRODUCTION IDARADO Mining Company's mill and surface plant are at the portal of the Treasury tunnel at elevation 10,625 ft, 12 miles south of Ouray, Colo. In 1943 and 1944 this tunnel was extended in a westerly direction to tap the Black Bear vein, and, together with the recently completed 1100-ft inclined Treasury raise, provides access to the Black Bear mine. In 1945 a 250-ton mill was placed in operation to process development ores containing much needed copper, lead and zinc for war requirements, and to serve as a pilot testing unit for larger scale milling operations. In the years 1902 to 1927 some 329,000 tons of ore were produced from mine workings on the Black Bear vein that extended downward from elevation 12,325 ft at the Black Bear tunnel portal to elevation 11,618 ft, the lowest level of the mine. The ores were taken to the Cascade and Smuggler-Union mills near Telluride over two aerial tramways. Gold bullion, a copper-lead concentrate, and a zinc concentrate were produced. ORES The Black Bear is a compound fissure vein that has undergone several stages of mineralization. The first important stage is represented by complex ores of copper, lead and zinc in a quartz gangue. Precious metal content in this stage is low. The galena is argentiferous and varying amounts of tetrahedrite, freibergite and bornite tend to increase the normal copper content represented in chalcopyrite. The zinc occurs as a true sphalerite with little or no iron. Later re-openings of the fissure took place more or less co-extensively with the successive introduction of (1) abundant quartz gangue containing minor amounts of chalcopyrite, galena and sphalerite; (2) free gold in a fluoritic quartz gangue; and (3) barren, low-grade quartz with minor amounts of pyrite as the only important sulphide. The sulphides show practically no oxidation and little mineralogic intergrowth. Because of the common overlapping of these successive stages of mineralization within a given mining area, considerable dilution of the base-metal content takes place, thus lowering the average mill-head grade. PLANT PRACTICE Metallurgical testing was started shortly after the Black Bear vein was encountered on Treasury tunnel level in 1944. Laboratory tests showed good flotability of the minerals in the ore and indicated an economic separation could be made of the copper, lead and zinc minerals into their respective concentrates. The success of the separation is close reagent control to maintain a balance between promotion and depression of the sulphides. To maintain this balance reagent consumption for the three metal separation is considerably higher than in ordinary lead-zinc flotation.

Jan 1, 1947

By H. F. Yancey

PHYSICAL properties have been used for a long time in characterizing different kinds of coal, and physical properties, such as friability and slacking, have been included with chemical properties in grouping coals according to the widely used system of classification adopted by the U. S. Geological Survey. Nevertheless, the usefulness of physical properties for classifying the different ranks or kinds of coal has been limited because it has not been possible to express these properties in definite form as numerical values. Because of the present wide interest in the formulation and adoption of a satisfactory system of coal classification, the Northwest Experiment Station of the U. S. Bureau of Mines, in cooperation with the College of Mines of the University of Washington, has recently completed a study of the physical and chemical properties of approximately 100 different coals ranging from lignite to, anthracite. PURPOSE AND SCOPE For this investigation samples were collected in approximately 45 mines and prospects in Washington and 23 mines in other states. The following studies, in addition to proximate and ultimate analyses, were made of the samples: Determination of (1) friability, (2) slacking or weathering characteristics, (3) yield of products by low-temperature carbonization in a laboratory retort, and (4) agglutinating value. Most of the coals examined were subjected to each of the tests enumerated, but with coal of some ranks this was unnecessary in some instances, and in others it was not feasible to do the work involved.

Jan 1, 1932

By R. W. Bartlett

The rates of oxidation of tungsten have been determined at temperatures between 1320" and 3170°C and oxygen pressures to 1 amn using a surface -recession measurement technique. Above approximately 2000°C and 10-6 atm the rate is independent of temperature and can be calculated from gas collision theory assuming a constant reaction probability, e, of 0.06. Oxygen molecules react at surface sites where oxygen atoms have previously chemisorbed. This provides a direct pressure dependence at low pressures but at high pressures tungsten oxide molecule s form an adjacent gas boundary layer which lowers the PO2 at the tungsten surface. A correction for this effect using free-convection theory fits the rate data over the entire oxygen-pressure range from 10-8 to 1 atrn as well as data using O2-A mixtures. Below 10-6 atrn and above 2000°C, e decreases with increasing temperature because of desorption of oxygen atoms. Below 2000°C the rate decreases with decreasing temperature at all oxygen pressures following an apparent activation energy of 42 kcal per mole and depending on (Po2)n with n varying between 0.55 and 0.80. MOST of the previous tungsten oxidation studies have employed gravimetric methods and have been limited to temperatures below 1000°C where the weight loss associated with evaporation of tungsten oxides is negligible compared with the weight gain from oxidation.' At higher temperatures, oxygen-consumption rates have been determined from pressure measurements, usually at constant flow rates, by Langmuir,2 Eisinger,3 Becker, Becker, and Brandes,4 and Anderson.5 The sensitivity of this method decreases with increasing pressure and, with the exception of Langmuir's work, these investigations were confined to pressures below 10-6 atm. Above approximately 1300°C, depending on the oxygen pressure, the rate of oxide evaporation is greater than the oxide-formation rate and the recession of the tungsten surface can be measured optically without interference from an oxide layer. This was first done by Perkins and crooks6 who heated tungsten rods in air pressures from 1 to 40 torr at temperatures between 1300" and 3000°C. The present investigation of the oxidation kinetics of tungsten at high temperatures emphasizes oxygen pressures from 10-6 to 1 atm. This is the range of interest for earth atmosphere re-entry applications of tungsten for which little data were previously available. APPARATUS The apparatus is a modification of the type used by Perkins and crooks.' Ground tungsten seal rods, 6 in. long by 0.125 in. diam, were mounted vertically between two water-cooled electrodes, one fixed and the other having free vertical travel. The movable counter-weighted electrode is prevented from undergoing horizontal displacement by three sets of runners mounted at 120-deg intervals. Electrical contact is made by means of a water-cooled mercury pool. A 24-in. vacuum bell jar having a volume of approximately 267 liters was used as the reaction chamber with the sample holder mounted in the middle of the chamber. Power was supplied from an 800-amp dc variable power supply. Temperature readings were made by means of a Latronics automatic two-color recording pyrometer. With this instrument, corrections for emissivity are not necessary provided the spectral emissivi-ties at two closely spaced wavelengths are equal. Supporting measurements were made with a micro-optical pyrometer corrected for emissivity of bare tungsten and window absorptivity. The micro-optical pyrometer was calibrated against a National Bureau of Standards calibrated tungsten lamp and both pyrometers were periodically checked against the melting points of tungsten and molybdenum using the oxidation apparatus. Above 10-6 atm, pressures were measured with an Alphatron gage calibrated against a McCleod gage. At 10-6 atm, a hot-filament ionization gage was employed. A magnified image of the self-illuminated tungsten rod was formed using a 360-mm objective lens mounted outside the bell jar. When the experiment exceeded 1 hr, the image was focused on a ground-glass plate about 10 ft from the tungsten rod at about X8 and the recession of the thickness of this image was monitored with a Gaertner cathe-tometer. When faster rates were encountered, a 35-mm time-lapse cinecamera with a telephoto lens and bellows extension was substituted for the ground-glass plate and cathetometer. Diameter recession rates were determined from the photograph image projected on the screen of an analytical film reader. EXPERIMENTAL PROCEDURE After installing the rod in the apparatus and cleaning it with acetone, the system was evacuated to 5 1 x 10-5 torr. Before oxygen was introduced,

Jan 1, 1964

By Nicholas J. Grant, John S. Benjamin

A series of composite alloys containing a high volume of NiAl, Ni3Ah or CoAl, bonded with 0 to 40 vol pct of a ductile metal phase, were prepared by powder blending and hot extrusion. The binder metals were of four types: pure nickel or cobalt, near saturated solid solutions of aluminum in nickel and cobalt, type 316 stainless steel, and niobium. Sound extrusions were obtained in almost all instances. Studied or measured were the following: interaction between the alunzinides and the binders, room-temperature modulus of rupture values, 1500° and 1800°F stress rupture properties, hardness, structure, and oxidation resistance. Stable structures can be produced for 1800°F exposure, with interesting high-temperature strength and good high-temperature ductility. Oxidation resistance was excellent. A large number of experimental investigations have been made of the role of structure on the properties of cermets and composite materials. Gurland,1 Kreimer et al.,2 and Gurland and Bardzil3 have indicated the preferred particle size in carbide base cermets to be about 1 µ, with a hard phase content of 60 to 80 vol pct. The optimum ductile binder thickness was noted to be 0.3 to 0.6 µ.1 Complete separation of the hard phase particles by the binder is important in reducing the severity of brittle fracture.' The purpose of the present study was to produce structures comparable to the conventional cermets, using a series of relatively close-packed intermetal-lic compounds rather than carbides as the refractory hard phase, and to study the effects of binder content and composition on both high- and low-temperature properties. The selected intermetallic compounds were particularly of interest because of the potential they offered in yielding room-temperature ductility. The highly symmetrical structures are known to possess high-temperature ductility and room-temperature toughness. Based on a ductile binder, the alloys were prepared by the powder-metallurgy route to avoid melting and subsequent alloying of the matrix, and were extruded at relatively low temperatures. It was expected that the composite alloy would retain useful ductility. In contrast, infiltration and high-temperature sintering led to alloying of the matrix and to decreased ductility. The systems Ni-A1 and Co-A1 were selected for this study. In the Ni-A1 system the compounds NiA1, having an ordered bcc B2 structure, and Ni3Al(?1), having an ordered fcc L12 structure, were chosen. In the system Co-A1 the intermetallic compound CoAl with an ordered bcc B2 structure was used. ALLOY PREPARATION The intermetallic compounds, see Table I, were prepared by using master alloys of Ni-A1 and CO-A1, with additions of either cobalt or nickel to achieve the desired compositions. The master alloy in crushed, homogenized form, was melted with pure nickel or cobalt in an inert atmosphere, cold copper crucible, nonconsumable tungsten arc furnace. The resultant intermetallic compounds were homogenized at 2192°C in argon, crushed, and dry ball-milled in a stainless mill to -100 and -325 mesh for the Ni-A1 compounds and to -325 mesh for the CoAl compound. Finer fractions were separated for some of the composite alloys. Several ductile binders were utilized. These included Inco B nickel, 5µ ; pure cobalt, 5 µ, from Sher-ritt Gordon Mines, Ltd.; fine (-325 mesh) niobium hydride powder; fine (15 µ) type 316 stainless-steel powder; and near-saturated Ni-A1 and Co-A1 solid-solution alloys, also in fine powder form. The niobium hydride was decomposed above about 700°C in processing of the compacts in vacuum to produce niobium powder. The Ni-7.1 pct A1 and the Co-5.3 pct A1 solid-solution alloys were prepared from pure nickel or cobalt and pure aluminum by nonconsumable tungsten arc melting under an inert atmosphere. The ingots were homogenized, lathe-turned to fine chips, and dry ball-milled in air to -325 mesh powder. These solid-solution alloys are designated NiSS and CoSS; see Table I. Subsequently the hard and ductile phases were dry ball-milled as a blend. Experiments clearly established the need to coat the hard particles with the ductile binder to optimize subsequent hot compaction by extrusion. Ordinary dry mixing usually resulted in nonhomogeneous alloys which were quite brittle. Conventional cermets are consolidated by liquid phase sinteiing or infiltration, which resulis in undesirable and uncontrolled alloying of the binder phase. For this study, a loose (unsintered) powder-extrusion process was emploved, minimizing reactions between binder and hard particle, thereby permitting much greater control of composition and structure. The constituent powders were first mixed in the desired

Jan 1, 1967