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By Otto H. Henry, J. J. Cordiano

Though powder metallurgy is one of the oldest of metallurgical processes, it is in its infancy as a branch of the modern field of metallurgy. As early as 3000 B.C., the ancients produced implements and weapons of iron by reducing iron ore in a charcoal fire and forging the resultant sponge while still hot into the desired shape. Melting and casting soon superseded the powder metallurgy practice, and it was not until the turn of the twentieth century that powder metallurgy was applied to the commercial manufacture of tungsten lamp filaments. From this beginning, the use of the powder metallurgy process spread" to the manufacture of porous oil-impregnated bearings, cemented carbide tools, contact points, diamond grinding wheels, clutch facings, starter brushes, resin-bonded radio cores, magnets, metallic filters, and other parts. More recently, and particularly during the present war, there has been a definite trend toward the manufacture of machine parts by this process in competition with the more conventional fabricating methods. Notwithstanding these many developments, scientific progress has been slow. While many articles on powder metallurgy have appeared in the press and various publications, there have been too few contributions to our technical societies, with the result that there is a dearth of the type of information necessary for the full development of this science. It is of primary importance that the results of research and development be freely disseminated if powder metallurgy is to realize its fullest potentialities. Purpose of Investigation Excluding the refractory metal field, all commercial powder metallurgical products made by a cold compressing and sintering operation contain voids. These voids can be substantially eliminated by subsequent severe working and heat-treatment, but ' such practice has not proved economical in most applications. By applying heat during the compressing operation, plastic deformation of the powder particles is substantially increased, and the resultant compact or briquette is practically free of voids. However, hot-pressing has many attendant problems, which at present preclude its use in the manufacture of precision parts. The more important of these are to devise: 1. A die steel capable of withstanding high pressures at elevated temperatures. 2. A high-temperature lubricant capable of preventing excessive die wear and welding of the compact to the die walls. 3. A method of retaining a neutral or reducing atmosphere in the compacting zone of the die assembly. 4. A means of ejecting the hot-pressed specimen into a neutral atmosphere until reasonably cold.

Jan 1, 1946

By N. J. Grant, C. F. Flo, F. B. Cuff

An investigation of the Ti-Cr system has shown the presence of a complete series of solid solutions in the ß phase, with a minimum in the solid us near 50 pct Cr. An intermetallic compound, TiCr2, forms during cooling from the ß solid solution. There is a eutectoid reaction at the low chromium side of the system. IN view of the recognition of the potentialities of titanium and its alloys as important structural materials there has arisen a need for a systematic investigation of various titanium binary diagrams. Of these, the Ti-Cr system was of particular interest because of the improved properties imparted to titanium by small chromium additions. A literature survey disclosed a partial constitution diagram that had been suggested by Vogel and Wenderott,¹ This diagram, shown in Fig. 1, is the result of work done on the Fe-Cr-Ti ternary system. Alloys containing less than 43 pct Cr were not investigated. In addition to this, a study of alloys in the Ti-Cr system was reported by McPherson and Fontana.² Adenstedt3 furnished a number of dilato-meter curves obtained from specimens containing 5, 10, and 15 pct Cr. Very recently, McQuillan4 submitted a proposed diagram covering the complete range of compositions between titanium and chromium. This diagram was in general agreement with the one described here. Experimental Procedure Sponge titanium (99.7 pct) and electrolytic chromium (Bureau of Mines analysis: 99.03 pct Cr, 0.40 pct Fe, 0.53 pct O2) were used for the initial determination of the system. Subsequent refinement of points, particularly for the low chromium side of the diagram, was accomplished by using iodide titanium and high purity chromium (National Research Corp. analysis: 0.050 pct C, 0.045 pct O2). The Ti-Cr alloys were prepared by melting 25 g charges in an enclosed, water-cooled copper crucible using a movable tungsten electrode. The atmosphere was helium, prepurified by first passing it through a drying agent and then through sponge titanium heated to 850°C. Particular care was taken to avoid any large scale segregation resulting from insufficient mixing of the components. Each charge was melted and kept molten for approximately 1 min, allowed to cool, inverted, remelted, and kept molten for an additional minute. In order to ascertain the amount of impurity pickup (oxygen and nitrogen) during the melting procedure, three iodide titanium specimens were melted in sequence and then checked for increased hardness. The specimens in the as-cast condition showed a relatively large hardness increase over the hardness of the as-received iodide titanium. The scattering on these readings was large. However, after the specimens had been homogenized at 840 °C for 2 hr and slowly cooled, the measured increase in hardness over the as-received titanium was small with a corresponding decrease in scatter. The reason for this hardness change is that in heat treating these specimens below the transformation temperature (885°C) there was a conversion from the as-cast structure to the more nearly equilibrium structure of the annealed specimen. The alloys were analyzed for chromium content and although the actual percentage was invariably low the deviation from the nominal composition never exceeded 0.6 pct Cr.

Jan 1, 1953

By H. Warlimont

Precipitation of cementite plates from carbon-rich austenite gives rise to localized formation of martensite in the matrix next to the precipitate upon quenching. The growth of this martensite is urmsually restricted. A distinct lattice relationship is found between this martensite and the cementite phase. It is concluded that the martensite is nacleated in the region of lowest carbon concentration at the austenite-cementite integace, with low-energy matching to the cementite lattice. The restriction of the growth of the martensite phase to the vicinity of the cementite plates can be explained by misfit in the austenite -martensite glissile interface and the uphill carbon gradient in the austenite. The investigation was carried out by light- and electron-microscopy and by electron-diffraction of thin foils. THE precipitation of cementite from carbon-rich austenite gives rise to a localized formation of martensite in the matrix next to the precipitate upon quenching. This behavior has been described previously by A. B. Greninger and A. R. Troiano1 for a 1.78 pct C steel ("pseudomorphic martensite"), and by A. Hultgren2 for a Fe-0.99 pct C—5.20 pct Mn steel. By measuring the axial ratio of the martensite phase,' it was found that the carbon content of this martensite was lower than that of the martensite which formed without prior partial transformation to cementite. Therefore, it was concluded that lowering of the carbon concentration in the matrix next to the precipitate raised the M, temperature locally and thus led to nucleation of martensite at the ce-mentite-austenite interface. However, the restriction of the growth of the martensite phase to the vicinity of the cementite plates cannot be explained by localized carbon impoverishment alone. Under ordinary conditions a martensite plate which has been nucleated in a favorable site can continue to grow until it encounters a physical barrier, or until it encounters a compositional barrier in the matrix, where the total free-energy change for the reaction is not negative, since it has been shown by numerous experiments that the activation energy for growth of ordinary martensite is virtually zero.' An investigation was conducted by transmission electron microscopy and electron diffraction to study this extraordinary martensite. It was expected that crystallographic in addition to compo- sitional conditions caused this abnormal mode of martensite formation. MATERIAL AND PROCEDURE An Fe-C-Mn alloy containing 1.15 pct C and 5.1 pct Mn was used for most of the experiments. For metallographic work, specimens 15 by 15 by 3 mm were used. For preparation of thin films, the initial specimen size was 15 by 15 by 0.7 mm. The heat treatment of each specimen consisted of austenitizing for 1/2hr at 1100'~ in an argon atmosphere, subsequent isothermal holding in a lead bath at one of a series of different temperatures, and, finally, quenching into brine at room temperature. For metallographic examination, the specimens were mounted in bakelite (slight tempering effects resulted from the heat in the mounting press) ground, polished, and etched in nital. Thin films for transmission electron microscopy and electron diffraction were obtained by grinding the specimens to approximately 0.2 mm thickness and subsequent electrolytical thinning. The orientation relationships were evaluated from sets of electron micrographs and selected-area electron diffraction patterns of single units of the transformation products and the adjacent matrix. RESULTS In the alloy which was used for the experiments, proeutectoid carbide precipitates over a wide tem-

Jan 1, 1962

By R. E. Holmes, T. W. Leland

The solution to the unsteady-state mass- and heat-transfer equations describing the adsorption of a dilute component from a gas stream flowing through a packed bed is readily applicable to the design of hydrocarbon recovery units. The solution to these equations depends on the assumption of a linear equilibrium equation to relate concentrations on the surface to concentrations in the gas. This assumption is shown to be justifiable for components which are very dilute in the gas stream. This approach permits the prediction of the effects of bed shape and size, cycle time, and gas mass velocity on the recovery of each component. It also allows the time and gas flow rate needed for regeneration to be determined. The equations involve a constant for each component and another constant characterized by the type of adsorbent. These constants must be evaluated from operating data. Values of these constants determined for silica gel are presented. This paper points out the type of experitrrental data needed for a thorough design of hydrocarbon adsorption units. INTRODUCTION The growing importance of adsorption processes to recover hydrocarbons from natural-gas streams has created an urgent need for better methods to design and predict the performance of these processes. Since the original use of solid adsorbents in natural-gas processing was for dehydration, most of the current design methods for hydrocarbon recovery re extensions of those developed earlier for dehydration. These methods are based on an adsorptivity or bed loading factor defined as the amount of a particular component adsorbed per unit weight or per unit volume of the packed bed. The recovery of a given component is then expressed as an empirical function of its adsorptivity on a given adsorbent. In spite of the simplicity and initial utility of these methods, they are not completely satisfactory in that they do not predict the effect of the dimensions of the bed in relation to flow rate and cycle time, and do not predict the distribution of various components along the length of the bed. It is possible, however, to make a considerable improvement in the design of these processes by applying the basic equations of mass and heat transfer applicable to gases containing small concentrations of adsorbing components flowing through granular beds. These equations have been developed and solved by Anzelius,' Furnas,% nd Hougen and Marshall." where G = mass velocity of the inert constituents (assumed to be methane), lb/hr-sq ft, y = weight of an adsorbing component in the gas per unit weight of inert, lb of a given constituent per lb of methane, x = bed length, ft, p, = packed bed density, lb/cu ft, P, = gas density, Ib/cu ft, f = void fraction in the bed, t = time, hours, and w = weight of a given component adsorbed per unit weight of adsorbate free solid, Ib/lb pure adsorbent. The term on the left side of Eq. 1 represents the flow rate per unit bed volume of a component entering less the flow rate per unit bed volume of that component leaving the differential bed length. The terms on the right side represent, respectively, the accumulation rates per unit bed volume on the solid and in the gas phase within the differential bed section. The last term is negligible for this application since pgfv is quite small in comparison with p, and G. The partial differential is determined by the mass-transfer rate between gas and solid in the differential section.

By Harold P. Rue

The problem of removing hydrogen sulphide from natural gas has become within the last few years one of major importance to the natural gas and petroleum industry. Several gas fields in the United States produce natural gas containing hydrogen sulphide, and a number of oil fields produce gas which contains hydrogen sulphide. 3 The percentage of hydrogen sulphide in natural gas produced in these fields varies from a fraction of 1 per cent t6 as high as 20.5 per cent by volume. 4 Hydrogen sulphide is probably the most toxic gas associated with natural gas,5 and from the standpoint of health and safety it is important that the hydrogen sulphide be removed if the gas is to be used for domestic purposes. Hydrogen sulphide, besides being toxic and therefore a hazard to animal life, is under certain conditions an active corrosive agent and very destructive to equipment. 6 The annual loss caused by the corrosive action of hydrogen sulphide on equipment amounts to hundreds of thousands, perhaps millions, of dollars. 7 The largest percentage of natural gas produced in the United States comes from wells which also produce oil. This associated gas invariably carries with it light fractions of petroleum, which when separated from the gas, condense to the liquid form at ordinary temperatures and pressure. Commercial application is made of this natural condition, and "wet gas" is processed to obtain so-called natural gasoline by employing the principles of absorption, adsorption, or compression and subsequent expansion accompanied by cooling. In the manufacture of natural gasoline from wet gas which contains hydrogen sulphide, if this contaminating gas is not first removed, some of the hydrogen sulphide may oxidize to form free

Jan 1, 1932

By R. L. Stephenson

The effect of fabrication variables on the creep-rupture properties of TZM (Mo-0.5 pct Ti-0.08 pct Zr-0.03 pct C), Cb-TZM [columbium-modified TZM (7140-1.5 pct Nb-0.5 pct Ti-0.8 pct Zr-0.05 pct C)], and TZC (Mo-1.2 pct Ti-0.25 pct Zr-0.15 pct C) has been studied at 982º. 1093º, 1204º, and 1315°C for rupture times up to 1000 hr. The creep tests were performed in vacuum creep furnaces, employing cold-trapped oil-diffusion pumps, operating at pressures 2 X 10-7 tory. After-test chemical analyses showed no significant changes in interstitial content. The creep-rupture properties of TZM, fabricated according to current mill practice, were determined in the stress-relieved and in the re crystallized conditions to provide control data. Specimens of TZM, Cb-TZM, and TZC fabricated according to each of two experimental fabrication schedules were then tested. The schedules differed in terms of in-process annealing temperatures and warm -working temperatures. The data make several conclusions evident. Fabrication variables do indeed have a profound effect on the long-time creep-rupture properties of molybdenum alloys. Strengths of these alloys compare very favorably with those of other refractory metal alloys in the temperature range studied. FOR several years we have been engaged in the evaluation of the creep-rupture properties of promising refractory metal alloys. We consider the determination of these properties to 1000 hr at several temperatures necessary for a valid comparison. Such an evaluation indicated that the creep strengths of Mo-0.5 pct Ti and TZM (Mo-0.5 pct Ti-0.08 pct Zr-0.03 pct C) compare favorably with those of other refractory metal alloys. Tensile and recrystallization data of other investigators1,2 show an effect of fabrication variables. More specifically, those fabrication schedules which produce high strength through work hardening exhibit the least thermal stability. A similar conclusion was reported by perkins3 for Cb-TZM. This would lead one to expect that, in terms of creep, those materials showing higher strength at short times and low temperatures would exhibit lower strength at long times and high temperatures. If these expectations were fact, a substantial improvement could be made in these alloys by optimizing their fabrication schedule in terms of their anticipated service temperature. We therefore undertook a study of the effect of fabrication variables on the long-time creep properties of TZM, Cb-TZM (Mo-1.5 pct Cb-0.5 pct Ti-0.08 pct Zr-0.05pct C) and TZC (Mo-1.2 pct Ti-0.25 pct Zr-0.15 pct C). MATERIALS All materials used were supplied by the Climax Molybdenum Co. of Mich. Our original intention was to test the same heat of each alloy in different conditions. Fabrication difficulties prevented this; however, heats with nearly identical analysis were selected. The compositions of the materials used are listed in Table I. All materials were supplied in the

Jan 1, 1970

By George C. Fryburg

The oxidation of platinum at high temperatures (above 800°C) is controlled by boundary-layer diffusion, except at the lowest pressures. The rate of oxidation is determined by the rate of diffusion of oxide through the boundary layer away from the platinum. This is in contrast to the situation in tungsten and molybdenum oxidation, where the rate is controlled largely by the diffusion of oxygen through the boundary layer towards the metal. The difference in mechanism is a consequence of the much lower rate of oxidation of platinum. Evidence is presented that supports this mechanism for the oxidation of platinum. SOME years ago we reported1 a study of the oxidation of platinum in ordinary oxygen in the temperature region above 800°C, where the oxide is volatile. All the data available up to that time had been obtained with platinum crucibles or plates heated in tube furnaces. The results were widely variant and highly dependent on the flow rate of gas through the furnace: the faster the gas flow rate, the higher the measured oxidation rate. These results indicated to us2 that reaction equilibrium was being approached in the furnace and that the rate was diffusion limited. To minimize these effects, we employed1 electrically heated ribbons placed in a cooled bulb. We were primarily interested in the low-pressure region, but some experiments were performed at pressures up to 1 atm. We found a peculiar pressure dependency which was explained on the basis that the rate of oxidation was proportional to the pressure at all pressures, but that with increasing pressure an increasing fraction of the volatilizing oxide was back-reflected to the ribbon by the surrounding gas molecules. The oxide molecules are not stable at the temperature of the platinum and are decomposed on impact, the platinum being re-deposited on the specimen. Because of the limitations of space, this explanation was offered summarily without detailed support. Recently, Jaffee and associates3, 4 have suggested that the pressure dependency of our data up to 0.5 torr could be explained on the basis of a Langmuir-type adsorption of molecular oxygen followed by dissociation. Since the correct explanation of the effect of pressure on the oxidation rate is a necessary requirement to the development of the basic mechanism of the oxidation reaction, we felt that the back-reflection theory should be presented in more detail. This is the object of the present paper. RESULTS AND DISCUSSION A typical set of results showing the pressure dependency of the oxidation reaction is presented in Fig. 1, where we have plotted, on logarithmic scales, the rate of oxidation, R, vs the pressure, p. For comparison some results of Krier and Jaffee4 and of Raub and plate5 are also shown. These data were taken at atmospheric pressure using furnaces. The effect of reaction equilibrium within the furnace is apparent from the much smaller values that these observers obtained: Raub and Plate In O2 found a rate of 3.1 µg cm-2 hr-1 compared to our value of

Jan 1, 1965

By N. P. Goss

Asterism appearing in X-ray Laue diagrams is an extremely sensitive index of changes in the internal structures of individual grains of polycrystalline metals. It indicates the existence of various forms of subboundary structure or subdivision within the grain. The present viewpoint is that the grain is made up of lattice blocks having substantially the same size and shape.l-4 When these blocks all assume exactly the same orientation, the crystal is perfect and asterism is absent. Cleaves and Thompson5 observed that the hardness of iron is virtually constant over a range of grain sizes. It is quite possible that the subboundary structure caused this, and that the latter is just as important as grain size itself. While the experiments reported here are not conclusive, they do support this viewpoint. Metals can, under certain conditions, be prepared entirely free from asterism, but this preparation does not necessarily require long heating or slow cooling, as will be shown here. In this paper a distinction is made between mosaic and distorted grain structures but the difference is only a matter of degree. In a previous paper9 it was shown that the grains can become distorted even during heat-treatment. The block structure exists inherently, even in a perfect crystal; the size and shape of these crystallites, or lattice blocks, varies, depending upon a number of factors. Goetz, for example: has shown that a metal before crystallizing from the molten state must go through a paracrystalline phase with oriented groups of atoms about 10-l4 to 10-l5 c.c., or 10-15 cm. in size. These crystallites, or blocks, appear to survive the most drastic cold-working, for it has been shown, and more elaborately by Wood's recent work," that cold-working cannot reduce the size of these crystallites smaller than of the order 10-5 cm. The interesting correlation brought out by the experiments of Goetz and Wood is that the block size after severe cold-working is of the same order of magnitude as the groups that formed in the metal just before crystallizing from the molten state. The work of Ewald and Rennenger17 and the now classical experiments of Darwin and James12,13 all strongly suggest the existence of a periodic structure of crystal. lites smaller than the grain itself. Related to this effect is the appearance of veinlike markings in the microstructures. Ammer-man and Kornfeld,16 in their experiments, did not observe any evidence of veining in specimens of alpha iron annealed below A3 and believed that veining was not due to impurities in the iron, as Northcott16 suggested; veining was observed by them only in specimens annealed above A3. Hanemann,16 Schrader and Tangerding found that annealing in hydrogen did not diminish veining and so it could not be caused by oxides. They concluded that veining was due to a slight difference in the orientation of ferrite lattice blocks. This viewpoint is well supported by the evidence presented here.

Jan 1, 1941

By N. P. Goss

Asterism appearing in X-ray Laue diagrams is an extremely sensitive index of changes in the internal structures of individual grains of polycrystalline metals. It indicates the existence of various forms of subboundary structure or subdivision within the grain. The present viewpoint is that the grain is made up of lattice blocks having substantially the same size and shape.l-4 When these blocks all assume exactly the same orientation, the crystal is perfect and asterism is absent. Cleaves and Thompson5 observed that the hardness of iron is virtually constant over a range of grain sizes. It is quite possible that the subboundary structure caused this, and that the latter is just as important as grain size itself. While the experiments reported here are not conclusive, they do support this viewpoint. Metals can, under certain conditions, be prepared entirely free from asterism, but this preparation does not necessarily require long heating or slow cooling, as will be shown here. In this paper a distinction is made between mosaic and distorted grain structures but the difference is only a matter of degree. In a previous paper9 it was shown that the grains can become distorted even during heat-treatment. The block structure exists inherently, even in a perfect crystal; the size and shape of these crystallites, or lattice blocks, varies, depending upon a number of factors. Goetz, for example: has shown that a metal before crystallizing from the molten state must go through a paracrystalline phase with oriented groups of atoms about 10-l4 to 10-l5 c.c., or 10-15 cm. in size. These crystallites, or blocks, appear to survive the most drastic cold-working, for it has been shown, and more elaborately by Wood's recent work," that cold-working cannot reduce the size of these crystallites smaller than of the order 10-5 cm. The interesting correlation brought out by the experiments of Goetz and Wood is that the block size after severe cold-working is of the same order of magnitude as the groups that formed in the metal just before crystallizing from the molten state. The work of Ewald and Rennenger17 and the now classical experiments of Darwin and James12,13 all strongly suggest the existence of a periodic structure of crystal. lites smaller than the grain itself. Related to this effect is the appearance of veinlike markings in the microstructures. Ammer-man and Kornfeld,16 in their experiments, did not observe any evidence of veining in specimens of alpha iron annealed below A3 and believed that veining was not due to impurities in the iron, as Northcott16 suggested; veining was observed by them only in specimens annealed above A3. Hanemann,16 Schrader and Tangerding found that annealing in hydrogen did not diminish veining and so it could not be caused by oxides. They concluded that veining was due to a slight difference in the orientation of ferrite lattice blocks. This viewpoint is well supported by the evidence presented here.

Jan 1, 1941

By J. Doig, A. A. Bauer, F. A. Rough

AN investigation of the phase relationships between the 6 phases of the uranium-molybdenum and uranium-titanium systems was conducted by studying alloys ranging in composition from uranium-31.5 at. pct Mo to uranium-34 at. pct Ti. A ternary cut showing these phase relationships is presented upon the basis of thermal-analysis data obtained and upon the basis of metallographic and X-ray examination of heat-treated and quenched specimens. Isothermal ternary sections, inferred from phase relationships determined between the two 6 phases, are presented for the uranium-molybdenum-titanium system. In the compositional range investigated, a y-ura-nium solid solution exists at elevated temperatures. At about 900°C the 6-U2Ti phase, with limited molybdenum solubility, forms. Epsilon, a molybdenum-titanium-rich solid solution appears at the same temperature and appears in all phase regions that exist below this temperature except in the terminal 6-phase regions. The y phase is stable to approximately 660°C in titanium-rich alloys; the y and 6-U,Ti phases decompose at this temperature to a uranium in a peritectoid reaction. The y phase in molybdenum-rich alloys of the ternary cut studied decomposes in an apparent eutectoid reaction, at about 593°C, to the 6-uranium-molyb-denum and a-uranium phases. The e phase while present, presumably plays a negligible role in these reactions. Titanium was found to markedly decrease the stability of retained y in these alloys. Age-hardening characteristics developed in titanium-contain ing uranium-molybdenum alloys quenched from 900°C and aged at 500°C are ascribed to decomposition of Y to the a and E phases. INTRODUCTION Alloys of uranium with molybdenum, titanium, and zirconium, respectively, are of great interest because of their potential use as reactor fuels. As a result, this interest extends to the constitution of these alloys, particularly in view of their uniqueness of constitution. In these systems, alone, among known uranium-alloy systems, an intermediate phase forms by direct decomposition of the high-temperature y-uranium phase. As a prelude to possible future alloy development, an investigation has been undertaken of the phase relationships that occur between the intermediate phases of these systems. These phase relationships have already been reported for the intermediate 6 phases of the uranium-zirconium and uranium-titanium systems.' The results of an investigation of the phase relationships between the 6 phases of the uranium-molybdenum and uranium-titanium systems are reported here. The phase relationships that occur in the binary systems pertinent to this investigation are shown in Fig. 1. Phase nomenclature appearing throughout this report is as shown in these diagrams; the crystal structure of the phase is given in Table I. EXPERIMENTAL PROCEDURE Alloys were cast and fabricated to flat plate. Specimens were heat treated and examined metal-

Jan 1, 1959

By G. Derge, R. E. Grace

Diffusivity of bismuth in liquid Pb-Bi alloys has been measured by the capillary reservoir method as a function of temperature and composition. Fair agreement between theory and experiment is found for the measured diffusion coefficients and activation energies of diffusivity and viscosity in dilute lead alloys. The measured diffusion coefficients in the high bismuth alloys describe mass transport in special concentration gradients. THE study of diffusion in liquid metals has been restricted largely to the amalgams over small ranges of temperature and composition. In contrast to this, diffusion and transport of ions and molecules in aqueous and organic solvents have been studied extensively for the past 50 years. The general difficulty of measuring diffusion coefficients in both gases and liquids at high temperatures has been that of eliminating convection currents. Measuring techniques based on freezing and sectioning after the diffusion period1 are not particularly satisfactory because of the volume changes and segregation which occur during freezing. The most promising technique which is currently being developed for liquid metals is the capillary reservoir method. This eliminates convection currents from the diffusion zone and requires measurements only of the length and initial and final compositions of the alloy within the capillary. Anderson and Saddington' have measured the diffusion coefficient of tungstate anions in aqueous Na,W04; Wane4' and Hoffman- ave studied structure and self-diffusion in water and self-diffusion in mercury, respectively, by similar techniques. Selection of the Pb-Bi system for initial study was made for several reasons: I—It afforded a low temperature system for development of techniques which would be suitable for high temperature measurements in liquid iron. This research on diffusion in liquid iron is now in progress as a part of a general program on the rates and mechanisms of slag-metal reactions. 2—Activity data on liquid Pb-Bi alloys of all compositions are available from the emf meas- urements of Strickler and Seltz.9 3—Some data are available on the diffusion of bismuth into solid lead. The measurements on solid alloys were made by von Hevesy and Seith7 for the concentration range of 0 to 2 pct Bi. An appropriate solution to Fick's law (unsteady state) for the capillary reservoir has been found by considering the mass flow which passes through a unit area of surface on a slab of infinite thickness as a function of time. This is the first application of this solution to diffusion measurements, although a similar solution is available for identical boundary conditions.' By analogy with heat conduction, the problem reduces to a perfectly insulated infinite slab at temperature C,, one surface of which is quenched to C, for all times, t. Heat is permitted to flow from quenched surface of the slab to an infinite supply of coolant at C.' (see Appendix). For the corresponding solution for diffusion in liquids, these terms become Where Cu is the initial concentration of component A" in entire capillary at t - 0, C, is the final con- In this case. component A is bismuth, and C, = 0. centration of component A in entire capillary at t = t, C, is the concentration of component A in bath for all t, 1 is the length of the capillary in centimeters, t is the time of diffusion period in seconds, and D is the diffusion coefficient in square centimeters per second. By comparison with solutions to Fick's law used in the solid state, this solution for liquids is similar to the Grube solution for solids, where the change in D with chemical composition is not taken into consideration. Difficulty is to be expected if large concentration gradients are employed and if the concentration gradients extend over an appreciable length of the capillary.

Jan 1, 1956

By Paul L. Russell

Introduction As the US and world population increases, so does the problem of supplying expanding energy needs. Early requirements worldwide were mostly met by abundant forest vegetation. Man's large scale use of minerals-energy resources is a fairly recent development. It occurred only during the past 200 years. Coal was the first mineral fuel to be used extensively. It was not a large factor in US energy production until the beginning of the 19th century. Petroleum and natural gas use is about 125-years-old. Electricity generation by falling water is about a century old. Atomic energy is only about 30-years-old. Extensive use of it is still in the future. Solar energy is in its infancy. Other energy sources such as wind, peat, and geothermal energy are minor and primarily of local importance. There were three main features of the fuel supply picture during the past 125 years. By 1900 coal replaced wood as the principal energy source. Since 1900, oil and gas have captured many coal markets. And after the mid-1970s, there was a resurgence of coal usage. (Figs. 1 and 2). Wood Until the advent of mineral fuels for energy, the major global energy source was wood. But populations increased and became centered in towns and cities. The problem of a sufficient energy supply first developed as nearby, easily obtained wood was exhausted. This problem became more acute as industries such as shipbuilding, smelting, and pottery making, became more widespread. Results of wood use as an energy source and building material are starkly evident in areas of the world where civilization is oldest. Entire forests were consumed and have never been replaced. In the US, wood supplied 90% of all fuel requirements until about 1850; and more than 50% until after 1880. By 1900, wood still supplied 20% of all energy uses. In the mid-1970s, energy costs increased, due largely to the oil embargo. This increase resulted in an increased wood demand for domestic heating. While large quantities of wood were being used, its share of the total energy market was still less than 3%. By 1983, wood was becoming increasingly scarce and more costly near most large US cities. Wind Wind powered ocean vessels for many centuries. Sailing ships were still common in 1900. Windmills provided an early, cheap energy source and continue to be used today on farms. Design and development of windmills for electricity generation is being conducted. Interest in wind power is growing worldwide where availability of wind currents warrants the investment. Also, US regulation requires power companies to "buy back" electricity generated by windpower in excess of that used by the windmill owner. Coal By the 1600s, England was producing, 900 kt/a (1 million stpy) of coal for heating and other uses. The first coal-fired blast furnace for iron smelting was developed in 1621. Coke-making was developed in 1755, and Watt invented the steam engine in 1769. Illuminating gas made from coal was introduced in London in 1800. During this same 200-year period, France and Germany developed coal production, but on a smaller scale than England. US, bituminous coal was mined near Richmond, VA, as early as 1748. Major bituminous coal-producing states in the early period were Pennsylvania, Virginia, Ohio, Kentucky, Illinois, and Maryland. West Virginia replaced Virginia as a major producer, when that state was formed in 1863. As the US expanded west, new coal deposits were found and exploited. Indiana, Alabama, Tennessee, Iowa, Kansas, Missouri, Washington, Colorado, Wyoming, Arkansas, Oklahoma (Indian Territory), Montana, New Mexico, and Texas had significant coal mining operations by 1900. Today the eight major coal producing states are Kentucky, West Virginia, Pennsylvania, Illinois, Ohio, Virginia, Wyoming, and Montana. These states accounted for about 80% of all coal (lignite, bituminous, and anthracite) produced in the US during the past several years.

Jan 8, 1984

By D. B. Smith, John Chipman

THIS investigation was undertaken as a prerequisite to the study of sulphur activities in the liquid system Fe-Cu-Ni, a continuation of the work of Sherman, Elvander, and Chipman,¹ using the same equipment and the same experimental technique. Temperature measurements are made with a disappearing filament optical pyrometer, sighting through a prism onto the surface of the melt which is heated by induction. To obtain true temperatures by this means, a calibration is necessary, since, in the absence of true black-body conditions, the observed temperature is always lower than the true temperature. This calibration is actually an emissivity correction, since the emissivity is a function of both temperature and melt composition. The original furnace was used with a deeper metal bath to provide sufficient depth of immersion for the thermocouple which was inserted through the sampling hole. The atmosphere was a 50-50 mixture or argon and hydrogen passing through the furnace at a rate of 1000 ml per min, and preheated by means of a molybdenum resistance coil. It was assumed that the rapid gas flow effectively eliminated any error which might have resulted from the presence of metallic vapor. The charge in each case consisted of about 300 g of metal, contained in a high purity, dense alumina crucible. At each melt composition, readings of the optical pyrometer and of a thermocouple immersed in the molten metal were taken simultaneously. Calibration curves of optical vs. true temperature were then made for each melt composition investigated. Thermocouples were made of 0.010 in. platinum and platinum-10 pct rhodium. They were checked against the melting point of chemically pure copper. Silica protection tubes with an outer diameter of approximately 1/8 in. were used. The depth of immersion of the tip in the molten metal was approximately 1 14 in. The life of these small thermocouples was relatively short, varying from approximately 1 min at 1550°C to 10 min or longer at 1300°C. Melts containing iron corroded the protection tubes more than those without iron present. Measurements taken while heating checked those taken while cooling, indicating the absence of a temperature gradient between the molten metal and the thermocouple. Results The temperature range of the. measurements on copper-free metals extended from the melting point to 1600°C. For copper the upper limit was 1380°C, and the data were extrapolated linearly for comparison with observations on the other metals and alloys at 1535°C. Examples of typical calibration curves for various melt compositions are shown in Fig. 1. These curves, besides providing a calibration system for the given experimental conditions, also give a basis for the calculation of true emissivities at the effective wavelength of the pyrometer, 0.65 micron. The starting point of the calculations is the value of the emissivity of iron at its melting point. Taking the melting point of electrolytic iron under hydrogen as 1535°C, and using the corresponding emissivity value of 0.43, as determined by Dastur and Gokcen,² the emissivities can be determined as functions of composition (and temperature, if desired). The relationship used is the Wien equation: ln(Ea)=C2/?(1/T1-1/Ta where E is the emissivity; a, the transmissivity of the optical system; C,, the fundamental constant, 14,380 micron-degrees; A, the wave length of light used (0.65 micron); Tt, the true temperature, OK; and T., the apparent temperature, OK. At the melting point of iron the observed temperature was 1358"; this gives 0.62 for the transmissivity, which compares with 0.65 in the similar system used by Dastur and Gokcen. The emissivities of all the compositions investigated were determined at 1535°C, after obtaining from each calibration curve the corresponding temperature. These results

Jan 1, 1953

By S. S. Lisnyak, B. D. Lichter

An important question in the study of the "hot corrosion" phenomenon (i.e., accelerated oxidation of Ni-Cr-base and Co-Cr-base "superalloys" in the presence of sulfur) is the condition of formation and physico-chemical properties of different sulfides in the system Ni-Co-Cr-S. seyboltl has observed the presence of ((chromium-rich" sulfides of compositions M3S4, and M2S3 in the scales of superalloys after hot corrosion; however, as he points out, much confusion exists as to the identity of sulfides found on Ni-Cr and Co-Cr alloys. According to the literature,&apos; ternary sulfides, e.g., NiCr2S4, (monoclinic structure) and CoCr2S4 (spinel structure) exist in these systems. Therefore, it is likely that these compounds are found in the scale of alloys undergoing hot corrosion. At present this possibility has not been discussed in the literature. Accordingly, we have begun investigation on the properties of ternary sulfides and their mutual interaction. The magnetic and electrical properties of pure NiCr2S4 and pure CoCr2S2 have been previously determined."-&apos; In the present work we have studied the interaction between these sulfides at high temperatures by means of X-ray diffraction and magnetic susceptibility. The measurements were carried out at room temperature after appropriate thermai treatment of samples at high temperatures. For alloys containing the three elements nickel, chromium, and cobalt this interaction may be of importance in the study of the hot corrosion phenomenon. The compounds NiCr2S4 and CoCr2S4 were prepared from stoichiometric amounts of the elements and CoS in the case of CoCr2S4. Samples were mixed and sealed in evacuated silica tubes, slowly heated to 900&apos; to 100O°C, and held at temperature for a total of 120 hr (NiCr2S4) or 80 hr (CoCr2S4) with intermediate re-grinding~. X-ray diffraction showed a pure spinel structure for CoCr2S4 and a monoclinic structure for NiCr2S4. Samples of mixtures of the two ternary sulfides were prepared with different compositions, see Table I. Samples were heated for a total of 80 hr at 900&apos;C in evacuated silica tubes, with one intermediate regrind-ing. The capsules were slow-cooled to room temperature in the furnace. Samples for X-ray diffraction were prepared by placing powder on a glass slide with vaseline. Dif- fractometer traces were obtained on a G.E. XRD-5 diffractometer using a chromium tube and vanadium filter (40 kv and 31 ma). The size of the spinel unit cell was determined from the eight highest angle peaks (80 deg < 2b < 160 deg). Calculated lattice parameters were plotted vs cos B . ctn 0 and extrapolated to 6 = 90 deg. Measurements of the magnetic susceptibility were carried out at room temperature using the Faraday method in a magnetic field of 3000 G and using a Cahn RG automatic microbalance. The system was calibrated from the well-established susceptibilities of platinum and palladium: (Xpt = 0.968 x 10-6 g-1 and Xpd 7 5.04 x 10~6 g-1). The weight of samples were approximately 0.200 g. X-ray diffraction patterns for unsintered mixtures showed Bragg peaks from both the spinel (CoCr2S4) and monoclinic (NiCr2S4) structures. After sintering no monoclinic peaks were observed in samples containing up to 40 mol pct NiCr2S4. However, the spinel peaks were displaced to higher angles with increasing

Jan 1, 1970

By Paul Billingsley, Augustus Locke

Certain adequately developed mining districts give complete three-dimensional patterns of ore bodies as clusters rising from roots in basement rocks with details controlled by structure of cover rocks. Examples are Goodsprings, Tintic, Bisbee, Leadville, Hedley, Eureka, Ely and Goldfield. In these and others the controlling structures in both basement and cover rock are elements of the neighboring regional structure and have been produced by rock movements in harmony with continental adjustments, which are usually steep tear faults in basement and thrust crumples in cover. Invariably the structures and movements in the districts have long geologic history and result from intersecting or superimposed deformations of different ages. Salt Lake districts show intersecting Uinta and Cordilleran elements; Colorado districts show intersecting pre-Cambrian and Tertiary elements, and Bisbee shows Tertiary superimposed on fina1 Paleozoic folding. Thus the position of districts has been determined by the patterns of North American orogenic belts, which make crossroads in Rocky Mountains and blocks of superimposed deformation in Sierra Nevada, northwest Coastand Canadian Coast Ranges. Essential function of repeated deformation is to strengthen rock, as from shale to slate to schist to gneiss, and thus make it competent to carry channels to depth with or without addition of intrusives. The essence of a mining district is the presence of such competent rocks with long-lived, deep, penetrating breaks reopened to permit passage of heat and associated products from depths to surface. It is suggested that melts such as magmas; metamorphism up to and including granitization; alteration, such as orthoclase-quartz or quartz-sericite or garnet, and also mineralization, may all be by-products of such heat escape along channels made by tectonic forces at specific foci. Magnitudes of Districts (Fig. 1) Any synthesis of ore deposits must deal with comparable units. Ore deposits are innumerable, but great ore deposits, commercially important for a generation or more, are few. The bulk of metals has come from districts of the first and second magnitudes (over $1,000,000,000 and over $250,000,000, respectively), and of these the United States and adjacent provinces of Canada and Mexico contain only 28, of which 9 are first magnitude: Butte, Coeur d&apos;Alene, Bingham, Bisbee, Sierra Gold, in the West; Tri-State, Southeast Missouri, Michigan copper, Sud-bury, in the East. Another 28 or 30 districts are of third magnitude (over $50,000,000).

Jan 1, 1941

By Augustus Locke, Paul Billingsley

Certain adequately developed mining districts give complete three-dimensional patterns of ore bodies as clusters rising from roots in basement rocks with details controlled by structure of cover rocks. Examples are Goodsprings, Tintic, Bisbee, Leadville, Hedley, Eureka, Ely and Goldfield. In these and others the controlling structures in both basement and cover rock are elements of the neighboring regional structure and have been produced by rock movements in harmony with continental adjustments, which are usually steep tear faults in basement and thrust crumples in cover. Invariably the structures and movements in the districts have long geologic history and result from intersecting or superimposed deformations of different ages. Salt Lake districts show intersecting Uinta and Cordilleran elements; Colorado districts show intersecting pre-Cambrian and Tertiary elements, and Bisbee shows Tertiary superimposed on fina1 Paleozoic folding. Thus the position of districts has been determined by the patterns of North American orogenic belts, which make crossroads in Rocky Mountains and blocks of superimposed deformation in Sierra Nevada, northwest Coastand Canadian Coast Ranges. Essential function of repeated deformation is to strengthen rock, as from shale to slate to schist to gneiss, and thus make it competent to carry channels to depth with or without addition of intrusives. The essence of a mining district is the presence of such competent rocks with long-lived, deep, penetrating breaks reopened to permit passage of heat and associated products from depths to surface. It is suggested that melts such as magmas; metamorphism up to and including granitization; alteration, such as orthoclase-quartz or quartz-sericite or garnet, and also mineralization, may all be by-products of such heat escape along channels made by tectonic forces at specific foci. Magnitudes of Districts (Fig. 1) Any synthesis of ore deposits must deal with comparable units. Ore deposits are innumerable, but great ore deposits, commercially important for a generation or more, are few. The bulk of metals has come from districts of the first and second magnitudes (over $1,000,000,000 and over $250,000,000, respectively), and of these the United States and adjacent provinces of Canada and Mexico contain only 28, of which 9 are first magnitude: Butte, Coeur d&apos;Alene, Bingham, Bisbee, Sierra Gold, in the West; Tri-State, Southeast Missouri, Michigan copper, Sud-bury, in the East. Another 28 or 30 districts are of third magnitude (over $50,000,000).

Jan 1, 1941

By N. Tiner

The mechanical properties of magnesium-alloy castings are greatly improved by grain refinement, and at present considerable attention is being paid to methods of obtaining fine-grained castings. One method of achieving this end is the use of proper melting and superheating techniques. These have been employed com-rnercially to a great extent, but no comprehensive study of the subject appears to have been published at present. Numerous investigations found in the literature include some reference to superheating techniques,1"4 and Achenbach, Nippcr and Piwowarsky,% ased principally on thermal analysis, ascribe the grain-refilling action of superheating to the influence of solid foreign particles or dissolved gases present in the melt upon the course of crystallization. The primary purpose of this paper is to present the principal facts concerning the cffect of superheating on grain size of magilesium alloys. .ittempts are also made to explain the experimental results in terms of a general hypothesis. The investigations are limited chiefly to common commercial casting alloys. In order to carry on this work, it was necessary to determine the factors influencing grain size and to develop a proper technique for the preparation of samples. Thc preliminary tests made by the author indicated that the grain size of magnesium alloys resulting from the solidification of a melt is determined by alloy composition, presence of impurities, rate of cooling during solidification, thermal history of the melt and to some extent the structure of the charge used. The rate of cooling is in turn dependent upon the mass of the melt, pouring temperature, and dimensions, properties and temperature of the mold into which the melt is poured. The final grain size of the alloys is still further modified by heat-treatment subsequent to casting. It is not within the scope of the present paper to discuss all of these factors, and it would suffice to note here that when equal amounts of magnesium alloys are prepared from the same alloying materials, melted in graphite crucibles under the same thermal conditions, then cast into graphite molds at a constant temperature, the resultant solids have the same grain size. By maintaining all other factors constant, it is thus possible to vary thermal conditions of a melt and to examine the grain size of resultant castings. Preparation of Samples and Grain-size TestiNg The alloys investigated in this paper were taken from ingots made by Permanente the Or or prepared by the author with pure carbothermic magnesium and necessary alloying elements. Melting was carried out under a flux in a graphite crucible heated uniformly by an

Jan 1, 1946

By N. Tiner

The mechanical properties of magnesium-alloy castings are greatly improved by grain refinement, and at present considerable attention is being paid to methods of obtaining fine-grained castings. One method of achieving this end is the use of proper melting and superheating techniques. These have been employed com-rnercially to a great extent, but no comprehensive study of the subject appears to have been published at present. Numerous investigations found in the literature include some reference to superheating techniques,1"4 and Achenbach, Nippcr and Piwowarsky,% ased principally on thermal analysis, ascribe the grain-refilling action of superheating to the influence of solid foreign particles or dissolved gases present in the melt upon the course of crystallization. The primary purpose of this paper is to present the principal facts concerning the cffect of superheating on grain size of magilesium alloys. .ittempts are also made to explain the experimental results in terms of a general hypothesis. The investigations are limited chiefly to common commercial casting alloys. In order to carry on this work, it was necessary to determine the factors influencing grain size and to develop a proper technique for the preparation of samples. Thc preliminary tests made by the author indicated that the grain size of magnesium alloys resulting from the solidification of a melt is determined by alloy composition, presence of impurities, rate of cooling during solidification, thermal history of the melt and to some extent the structure of the charge used. The rate of cooling is in turn dependent upon the mass of the melt, pouring temperature, and dimensions, properties and temperature of the mold into which the melt is poured. The final grain size of the alloys is still further modified by heat-treatment subsequent to casting. It is not within the scope of the present paper to discuss all of these factors, and it would suffice to note here that when equal amounts of magnesium alloys are prepared from the same alloying materials, melted in graphite crucibles under the same thermal conditions, then cast into graphite molds at a constant temperature, the resultant solids have the same grain size. By maintaining all other factors constant, it is thus possible to vary thermal conditions of a melt and to examine the grain size of resultant castings. Preparation of Samples and Grain-size TestiNg The alloys investigated in this paper were taken from ingots made by Permanente the Or or prepared by the author with pure carbothermic magnesium and necessary alloying elements. Melting was carried out under a flux in a graphite crucible heated uniformly by an

Jan 1, 1946

By Russell Franks

IN recent years a considerable tonnage of the 18 per cent chromium and 8 per cent nickel steels has been used in lightweight high-strength structures that must be resistant to deterioration under atmospheric conditions. These steels are austenitic in character and their strength can best be increased by cold-working. Although a marked increase in tensile strength is obtained, claims have been made that the cold-worked steels do not possess fully satisfactory elastic properties. Under certain conditions, however, this has caused no great concern, because, as explained by Ragsdale,1,2 the cold-rolled steels have been widely applied for construction of lightweight, high-strength trains using spot welding for joining purposes. Ffield3 agrees that the ease with which thin sections of the cold-rolled 18-8 steels can be joined by spot welding is greatly in their favor, but that they would be more satisfactory if higher elastic properties and a higher modulus of elasticity could be obtained. He states that, as far as the mechanical properties are concerned, the straight chromium stainless steel is superior to the cold-worked 18-8 steel, and in order to illustrate, he presented the data of Fig. 1, which also give a stress-strain curve for a low-carbon 13 per cent Cr steel. Ffield? pro-posed that columbium-bearing or titanium-bearing cold-rolled 18-8 steels be given a heat-treatment (preferably for 1 hr.) in the range 400° to 600°C. to precipitate carbides in order to obtain better mechani-cal properties. Discussions by Sutton,4,5 de Ganahl,6,7 8 Krivobok and Lincoln,9 Knerr,10 and Thaden11 have given further information on the properties of the cold-rolled 18-8 steels compared with other materials used in light-weight, high-strength structures. These investigators and others point out that the ultimate tensile strength of the cold-rolled 18-8 steel can be increased to 200,000 lb, per sq. in. or higher; Krivobok et al. indicate that

Jan 1, 1940

By E. F. Ketterer, D. B. Collyer, A. B. Wilder

The microstructural stability of 59 carbon and low alloy steels after 34,000 hr exposure at 900&apos; and 1050°F, including the weld heat-affected zone, is discussed. The tensile and creep rupture properties of the parent metal of many of the steels before and after 10,000 hr exposure are presented. SEAMLESS pipe has been used for many years at elevated temperatures. When seamless first began to be used, operating temperatures and pressures were relatively low and the steel produced adequately met the existing requirements. The need for knowledge of creep properties and graphitiza-tion was nonexistent. Today the requirements for tubular products have changed considerably. Operating temperatures and pressures are increasing continually and corrosion, particularly at high temperatures, presents a problem. Carbon steels have, under these conditions, certain economic advantages particularly in oil refinery use. However, the use of low alloy steels in the generation of steam power has increased appreciably due to the influence of molybdenum on creep strength and the control of graphitization with chromium. The behavior of carbon and certain low alloy steels after long periods of exposure at 900° and 1050°F (480° and 565°C) has been investigated. The steels were examined before and after 10,000 to 34,000 hr exposure. Results obtained with material exposed at 1200°F (650°C) are not presented due to the severe oxidation and resulting decarburiza-tion after long periods of exposure. The graphitization characteristics of carbon and low alloy steels are of fundamental concern to the users of these materials. This is particularly true where welded joints are employed. In order to evaluate properly the behavior of steels at elevated temperatures, it is desirable to employ not only long periods of exposure under controlled conditions, but to use material with known properties and steelmaking practices. The purpose of this investigation has been to study the microstructure of the weld and the tensile and creep rupture properties of the parent metal after long periods of exposure at elevated temperatures. In the report&apos;.&apos; of a previous investigation data were presented on the microstructure, impact strength, hardness, and oxidation characteristics after 10,000 hr exposure at elevated temperatures. When this investigation was planned, methods for controlling the graphitization characteristics of steel were not established and, therefore, a number of special alloy steels were included in the study. During the exposure of these steels, the influence of aluminum in accelerating the rate of graphitization and the inhibiting power of chromium were established by various investigators. Also, the vanadium-bearing steels, because of their outstanding creep properties, received attention. As a result of these developments, considerable progress has been made in the field of high temperature tubular products, and the necessity for special alloys to prevent graph-itization has been restricted largely to the use of chromium. Material Chemical analysis, deoxidation treatment, and austenitic grain size of the 59 steels discussed in this paper are shown in Table I. The steels were forged to 1x1 in. bars with the exception of steels 12B, 12DX, 12DY, 12DZ, 70, 70A, 72 to 78, 82, and 83 which were machined from heavy wall pipe. All bars were surface ground before exposure. Heat treatment before exposure is shown in Table 11. The majority of the steels included were melted in either a basic open hearth or basic electric furnace; how-

Jan 1, 1955