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By C. A. Stickels

The preferred crystallographic orientations developed during recrystallization of polycrystalline electrolytic iron sheet, cold-rolled 90 pct, were ivestigated. Recrystallization at 500° or 565° C for relatively short limes produces a texture which is similar to the rolling texture. A duplex partial fiber texture, significantly different from the rolling texture, is found when the annealing time is increased. Recrystallization at 700°C also produces a sequence of textures with increasing annealing time. In order of their appearance, these textures are: 1) the duplex partial fiber texture found at lower temperatures, 2) a four-component texture "near {322}(296)", 3) a {112)(110) +(111)(110) texture, and 4) a two-component "near {554)(225)" texture. Secondary recrystallization, or discontinuous grain growth, accompanies the development of the latter texture. However, the "near {554} (225)" texture was nol observed when secondary recrystallization occurred under other conditions. All of the ideal orientations found in recrystallization textures can be accounted for by the growth of minor components of the deformation texture. IN recent years there has been increased interest in improving the drawability of metals by controlling the preferred crystallographic orientation of the sheet.1-l3 Since more low-carbon steel is drawn than any other material, considerable attention has been focused on the properties of sheet steel. Efforts to improve drawability through texture control have been hampered by the lack of any published systematic study of the recrystallization textures developed in iron annealed below Acl. The purpose of the present work was to supply some of this missing information, specifically, the recrystallization textures obtained by isothermal anneals of polycrystalline iron, cold-rolled 90 pct. LITERATURE REVIEW—DEFORMATION TEXTURES Barrett14 reviewed and summarized the investigations of rolling textures in iron and low-carbon steel prior to 1952. The texture of heavily deformed iron (rolled 90 pct or more) is described as consisting of two partial fiber textures. The dominant fiber texture (designated here as fiber texture A) has a (110) fiber axis in the rolling direction and includes the orientations (001)[110], {112}( 110), and {111}(110). The secondary texture (designated here as fiber texture B) is described as having a (111) fiber axis in the sheet normal direction, and includes the orientations {111)( 110) and (111)(112). Since the publication of Barrett's book, there have been two detailed studies of the deformation texture in polycrystalline iron. In both instances, the more sensitive diffractometer methods of pole-figure determination were used. Bennewitz15 studied the cold-rolling textures developed in polycrystalline low-carbon steel and a 3 pct Si steel. He determined (110), (2OO), and (222) pole figures for specimens reduced 30, 50, 60, and 90 pct and analyzed his results in terms of partial fiber textures. He distinguished three stages in the development of the final deformation texture. 1) Grains rotate to form two incomplete fiber textures, with (110) fiber axes inclined 30 deg to the sheet normal toward the rolling direction (fiber texture B). After 50 pct reduction, the highest density of poles is near an ideal orientation {554}(225), the two components of which are members of this duplex fiber texture. 2) With increasing amounts of reduction, grains rotate about the former fiber axes toward ideal orientations of the type {112)( 110) (common to fiber textures A and B). 3) Finally, grains rotate about their ( 110) axis in the rolling direction, clockwise and counterclockwise from the orientations {112}(110). This produces the range of orientations from (111)(110) to (001)(110) commonly found in the rolling texture of heavily deformed iron (fiber texture A). No significant difference was found between the deformation textures of low-carbon and silicon steels. A few years earlier, Haessner and weik16 had determined (110) pole figures for carbonyl iron rolled to 30, 60, 80, and 90 pct reduction in thickness. heir data agree quite well with the more complete data of Bennewitz, but a somewhat different description of the evolution of the deformation texture was given. The appearance of (110) poles in the transverse direction is ascribed to a (100)[011] component rather than "near {554}( 225) " components, and the (110) fiber axes of fiber texture B are described as located 35 deg rather than 30 deg from the sheet normal. Both of these studies agree that the secondary texture present in heavily rolled iron, the duplex fiber texture B, has (110) fiber axes and not a single ( 111) fiber axis normal to the sheet, as

Jan 1, 1965

By Frank Harris

THE strength of hard drawn copper wire is a question of considerable importance to both manufacturer and consumer. Unlike steel and alloy wires, in which strength is governed by both chemical and physical considerations, the high conductivity required for copper wire in the electrical industry generally precludes the possibility of obtaining tensile strength through the use of alloys. Consequently, this strength is largely a matter of the physical structure, and all factors known to affect this structure, such as rolling temperature, die contours, lubricants, drawing speeds, etc., are matters of prime importance. The experiments described here have been carried out during the past year at the Baltimore plant of the American Smelting & Refining Co., in connection with a general investigation of some of the questions mentioned above, with the object of studying the physical structure of hard drawn copper, wire and the influence of certain factors on that structure. The work has been of an exploratory nature, rather than comprehensive, but it was thought that the results might prove of sufficient general interest to warrant their publication. The bulk of the hard drawn copper wire used in places where strength is an important factor is manufactured to strict specification, based usually in the United States on that of the American Society for Testing Materials (B-1-23). This specification calls for almost as high a value of tensile strength; consistent with elongation, as it is possible practically to obtain, and only physically perfect material can be relied on to meet such a specification. Obvious defects, such as laminations produced through fins in the original rods or brittleness due to overdrawing, are usually recognized and need only passing reference. Laminations can be detected by twisting and untwisting the wire several times; they will show up on the surface in the form of fins. Fig. 1 shows a very bad example of this defect. The wire may be perfectly smooth before twisting but when such laminations are present they will readily show up under this treatment. This trouble is caused by similar defects in the original rod produced by overfills in rolling. Fig. 2 shows the cross-section of the rod from which the wire in Fig.1 was drawn.

Jan 1, 1928

By C. C. Harris

The modern flotation machine is the result of decades of adaptation to a changing environment. In response to falling ore grades, and rising demands, tonnages, and costs, attempts are now underway to expand machine capacity by about an order of magnitude, while at the same time maintaining metallurgical performance and if possible reducing power requirements per unit throughput. Empiricism and rule-of-thumb have guided flotation machine design and scale-up, and a current need is for the development of a fundamental understanding of the subject. Design features of current and some obsolete machines are discussed in an attempt to outline those principles of flotation machine design which are known, and to delineate topics needing investigation. Attempts are made to distinguish between essential machine features required by the physics of the process, those features resulting from speculative theories on the flotation mechanisms, and features which are essentially styling variations. Topics discussed include: flotation mechanisms; tank and impeller/ stator design and dimensions; impeller speed; aeration; power consumption; particle suspension and dispersion; capacity; correlation of design parameters and operating variables; scale-up and scale-down; materials flow; problems in machine design. Design details and operating data for several current machine families are provided in Appendices.

Jan 1, 1976

By E. I. Larsen, E. F. Swazy, F. R. Hensel

HIGH strength, high-conductivity copper-base alloys have found considerable use in the resistance welding and electrical industry in the form of castings, forgings, or wrought products. There are a number of parts utilized by these industries which would lend themselves to manufacture by powder metallurgical methods provided the required physical and electrical properties are obtainable with such methods. The present paper describes the results of a study of certain copper-base powder metallurgical products susceptible to precipitation hardening. COPPER POWDERS The general types of commercially available copper powders are flake, spherical and electrolytic. Micrographs of these three types are shown in Figs. I, 2 and 3. The difference in particle shape is pronounced and obvious. The type of powder used is determined, or at least indicated, by the final properties desired in the finished article. The type of powder employed during the present investigation was electrolytic powder shown in Fig. 3, and was chosen because of its particular adaptability to the work described in this paper. The screen and chemical analyses of this powder are shown in Table I. In the preparation of compacts of high electrical conductivity, the purity of the copper powder is of importance in addition to other factors, such as density and the state of internal stresses. These variables were studied briefly by pressing a number of compacts to various pressed densities, sintering them for one hour in hydrogen at 975°C. and repressing. The size of the finished compacts was ¼ by ¾ by 8 in. The time interval from the application to the termination of pressure was approximately 20 sec. Some of the compacts were annealed before testing. The results are listed in Figs. 4 to 8, inclusive. [ ] In Fig. 4 the curves of pressed density vs. compacting pressure, and sintered density versus compacting pressure show the typical direct relationship of pressure and density. Compacting pressures exceeding about 30 tons per sq. in. result frequently in expansion during sintering for the type of powder used. Fig 5. indicates that variations in the initial pressed density within the range of about 5 to 6 grams per c. c., affect only slightly the pressure required to obtain a

Jan 1, 1945

By Henry D. Hibbard

The finishing and casting temperature of steel is a subject concerning which there is so much to learn that some temerity is required to write about it. This paper is, therefore, suggestive rather than dogmatic, though the views expressed seem now to the writer to be correct; nevertheless, they may be modified by future developments or discoveries but until then may be of some utility, particularly to beginners in the making of steel. Consideration is limited to the temperatures of simple steels, also called carbon steels, which are those in which carbon is the only element introduced for the purpose of conferring desired physical properties. Alloy steels will receive attention later. Three things largely determine what a batch of steel really will be; via., composition, casting temperature, and rate of teeming. Each has its advocates as being the most important; but they are so interrelated that each must conform to or be right for the other two. When one is claimed to be the most essential to have right, that, probably, was the one that it was found had to be of a certain description, within narrow limits, to suit the already adopted variations of the other two; that is, it was the last of the three to be brought into line. The term "casting temperature" is rather indefinite as there ale several temperatures to which the term may be applied. Most writers who have mentioned casting temperature of steel have alluded to it in relative terms. The few who have given actual temperatures, including Langmuir, Burgess, Styri, and Bash, do not give full information about them. Their results demonstrate the efficiency of pyrometers used rather than aid the steelmaker. In order to understand and use the temperatures, one should know in each case what the steel was, the stage at which the temperature was taken, whether it was usual or not, and whether or not it was right for that steel at that stage. When these

Jan 1, 1925

By S. H. Ash

Diesel locomotives, trucks, bulldozers, and other diesel-powered equipment are fast proving their superiority for mine transportation purposes. MORE than 100 minerals are mined and processed in the U. S. Management and labor have negotiated wage-scale agreements that have balanced wages and affected cost of labor in such a manner that comparisons can be made of transportation costs. Transportation of materials underground has become of vital importance.1"0 This report describes underground haulage methods and equipment, presents cost analyses, and discusses future plans for underground transportation in metal mines. It does not cover hoisting in shafts or slopes by cage, skip, or cars. Advancements have been made by the metal mining industry in lowering transportation costs in several mines, making it possible to mine orebodies or portions of orebodies not considered economically recoverable a few years ago.7,8 Wumerous reports have been written on transportation at metal mines; this report covers recent haulage at some representative mines throughout the U. S. The term metal mine is used to include all noncoal mines from which up-to-date information was obtainable. The best haulage equipment is subject to break- downs and accidents; but unless the rolling equipment, track, and signal systems are kept in good repair and are operated systematically, these breakdowns and accidents are likely to be frequent, destructive, and more costly than are necessary.10 Rail haulage crews using dependable rolling stock can move trips safely over grades low enough to be traveled without loss of control; this applies particularly if the track is straight and well maintained. Few accidents occur on main haulage lines if visibility and clearance are good and clear track is assured. One-way traffic underground is a recent development contributing to reduction of accidents and increased efficiency in mine haulage (Kelly mine, Anaconda Co., Butte, Mont., and Climax mine, Climax Molybdenum Co., Climax, Colo.)." Locomotives: Six types of mine locomotive used in recent years are trolley, storage battery, combination trolley and storage battery, diesel, compressed air, and gasoline. Electric haulage has been displaced by compressed air locomotives, which have been used successfully in certain instances, especially on dirty and wet tracks and steep grades." Maintenance cost is high, and operating efficiency, from a power standpoint, is low. Compressed air locomotives operate under self-contained power, and the absence of bare elec-

Jan 1, 1957

By B. L. Landrum, P. B. Crawford

The rise of a new laboratory model for studying tran-sient fluid flow problems, is described. The theory of he model is based on the analogy between the equa-ions which describe the flow of compressible fluids in porous media and the conduction of heal m solids. The "thermal model" is used to study the Iransient Pressure distributions during fluid displacement in five Standard well patterns. The studies indicate quantita-ively how die pressures change with lane at any poit in the reservoir. It was found that, in general, the equipressure contours are more radial about the wells early in the program than after steady-state conditions are established. In addition two studies of fractured five-spot patterns are presented to illustrate the opplicability of the model for a wide range of reservoir conditions. INTRODUCTION In reviewing the oil recovery techniques involving fluid displacement it was found that a very considerable quantity, if not all of the fluid within the pattern, mav be displaced (during a time-dependent or transient period rather than the usually applied steady-state conditions. Since the oil is displaced during a transient period, it was believed desirable to study the nature of the transients for several of the commonly used fluid displacement patterns. A literature review indicated fairly complete studies of the streamlines and isopotentials for steady-state: yet few data were available on the transient displacement process. Several simulated steady-state multiwell displacement programs were carefully mapped by Foster and Lodge on an electrical conducting sheet as early as 1875, hut the more recent work of Muskat provides the best source for steady-state phenomena having direct petroleum production applications. In order to study the transient phenomena a new lab-oratory model was developed which is based on the similarity of the equations that describe the flow of fluids and the conduction of heal. The equations describing heat flow were developed early in the nineteenth century. Regorous solutions tor most transient problems are in gcneral complex and usually limited to fairly restricted houndary conditions.2 the thermal model may be used to obtain approximate solutions for many transient problems without some of the restrictions required for analytical solution. The application of the analogy between heat conduc-. tion in solids. electrical flow and fluid flow, in a porous media has long been recognised. Techniques originally developed for the solution of heat conduction prol?lems have in many instances proved useful in ohtaining solutions to fluid flow prohlcms. Ir is not surpri5ing then that a heat conduction or thermal model might bf usefull in sludying transient fluid flow. As an example ol the problems suitable for the study with the thermal model, transient pressure distributions during fluid displacement programs are studied for the five-spot, line-drive, staggered line-diibe. seven-spot 211 two fractured patterns. A model was constructed of an iron plate of geometry similar to the reservoir patterns under study. The principle of mirror images as suggested by Bryan was employed in shaping the model. A temperature differential was obtained by soldering two sections of copper tubing to the plate at points analogous to the well location in the field. Hot and cold water was passed through the tubing. The temperature distribution, analogous to the pressure distribution in the field. was obtained a a function of time from thermocouplcs distributed over the plate.

By P. Greenfield, C. C. Smith, A. M. Taylor

The Mg-Zr alloy ZA and Mg-Mn alloy AM503(S) are shown to have a markedly improved resistance to creep deformation after suitable heat treatments. This improvement makes them suitable for certain stress-bearing fuel element components in nuclear reactors. The extent of strengthening is described and an explanation of the behavior of both materials is given, based on a combination of strain-aging and grain growth. The increase in operating temperatures of fuel element components in Calder Hall type nuclear reactors has necessitated the development of magnesium base alloys with a very high resistance to creep at temperatures up to 500°C. Such alloys are not required for fuel element cans, which require high-creep ductility rather than strength, but for can supporting and stabilizing components, which are needed to support the imposed loads without deforming more than about 1 pct in times of up to 40,000 hr. The amount and type of alloying addition made to magnesium for these parts is limited by the necessity of keeping the cross-section to thermal neutrons as low as possible. The alloys must also possess a high resistance to oxidation in CO2. Alloys which have been developed for this application include ZA, an alloy of magnesium with 0.5 to 0.7 pct Zr and AM503(S), an alloy of magnesium with 0.5 to 0.75 pct Mn. In the as-extruded condition these alloys are very weak and ductile in creep but it has been found that they can be strengthened to a significant extent by heat treatment. This paper describes the method of developing a high-creep resistance in ZA and AM503(S), the extent of the strengthening produced and discusses the probable mechanisms of strengthening. TEST MATERIALS Specimens were taken from typical casts of ZA and AM503(S) alloys extruded into 2 1/4-in.-diam bars, supplied by Magnesium Elektron Ltd. Typical analyses of the bars were as follows: The as-extruded mean grain diameter was 0.001 to 0.002 in. for the ZA alloy and 0.003 in. for the AM503(S) alloy. EXPERIMENTAL METHODS Extruded bars of ZA alloy, 2 1/4 in. in diameter and 9 in. long, were heat treated in electrical resistance furnaces in an atmosphere of flowing CO2 containing 50 to 300 ppm water, thereby reducing the extent of oxidation compared with that which would have occurred in air. Heat treatments were carried out at 600oc for times of 8, 24, 48, 72, and 96 hr and material was subsequently both furnace cooled and water quenched. In order to measure the effect of time of heat treatment, specimens were creep tested at 400°C and 336 psi for about 1000 hr. Subsequently, the behavior of material heat treated for 96 hr at 600°C and furnace cooled was tested at a variety of stresses from 200° to 500°C. Tests were also conducted at 200o and 400°C on material in the as-extruded condition for comparative purposes. With the AM503(S) alloy, only the effect of heat treatment at 565°C for 4 hr was examined. It has been shown1 that such a heat treatment produces marked strengthening in this alloy. Tests on this material were conducted at a variety of stresses at 200°, 300°, and 400oc with comparative tests on as-extruded material at 200o and 400°C. The creep tests were carried out on machines using dead-weight loading and direct micrometer strain measurements on specimens 5 in. long and 0.357 in. diameter. At temperatures of 400° C and below, the creep tests were conducted in air, but at higher temperatures an atmosphere of CO2 was used. Grain size measurements were made on ZA in the extruded and heat treated states and on each specimen after creep testing. This was done by a line count of a minimum of 20 grains in two or three random fields in the longitudinal and transverse directions. The same method was used for measuring the grain size of as-extruded AM503(S), but the grain size of the heat treated material was so large that this method could not be employed. For heat-treated AM503(S) the large grained characteristics (between 0.1 and 1 in.) were confirmed by the measurement of individual grains. In the case of the ZA alloy, specimens taken from various stages in the program were analysed for hydrogen by a combustion method. Material in various states was also analysed for the soluble and insoluble zirconium content by dissolving in dilute hydrochloric acid. This technique has been useda for the determination of amounts of zirconium present

Jan 1, 1962

By P. Douglas, G. M. Leak, B. Mills

The sinusoidal surface relaxation technique has been used to measure the surface self-diffusion coefficient of spectroscopically pure nickel over a wide temperature range under a hydrogen atmosphere. A kink in the Arrhenius plot has been observed. In the temperature range T/T 0.98 to 0.80 (T in O K and T, is the melting temperature) the average self diffusion coefficient is given by Below the temperature T/T,- 0.80a decrease in the slope of the log Ds us 1/T plot is observed. This is associated with a diffusion process characterized by a lower activation energy (-20,000 cal mole'') and smaller preexponential term (-10- sq cm sec"). A series of experiments were carried out at T/Tm = 0.61 under a hydrogen atmosphere of higher oxygen partial pressure than for the rest of the experiments. It was found that Ds was significantly depressed due to oxygen adsorption. This evidence supports the opinion that the low temperature process (activation energy -20,000 cal mole-') is unlikely to be due to oxygen adsorption. An interesting feature of the present data is that the transition temperature (T/Tm - 0.80) is a function of orientation. For a small number of crystals of measured orientation the transition temperature was observed to be higher towards the low index (100) pole. Theories of surface diffusion are briefly reviewed and it is concluded that the present reszuts are best explained by invoking a surface roughening process. GJOSTEIN has recently analyzed available surface diffusion data for a wide range of metals. He suggested that two mechanisms were operative for fcc metals, an adatom process at high temperatures and a vacancy process at low temperatures. Results for nickel can be summarized as follows. At low temperatures (T/T, - 0.3 to 0.44) under ultra high vacuum conditions, Melmed2 measured an activation energy Q of 21 kcal mole-' using field electron emission microscopy. At higher temperatures (T/T - 0.7 to 0.9) under a vacuum of 10- ' torr, Maiya and lakel measured y as 39 kcal mole-' using the multiple scratch smoothing technique. The present work was undertaken to try to find out if two distinct processes could be observed. High temperature results give Q about 47 kcal mole-': there is evidence also for a low temperature value of about 20 kcal mole-'. These measurements were all made under a hydrogen atmosphere, in the temperature range 860" to 1412°C. Concurrent with the present study Bonze1 and jostein> have also observed a break in the Arrhenius plot for the (110) surface of nickel. These measure- ments under ultrahigh vacuum conditions using the laser diffraction technique are in excellent agreement with the work reported here under hydrogen annealing conditions. THEORY The available surface relaxation techniques include single and multiple scratch smoothing and grain boundary grooving. The processes have been compared in detail by Gjostein for conditions where surface diffusion dominates6 and Mills et al? where volume diffusion dominates. In summary the relevant points are as follows. Grain boundary grooving gives an average Ds for the two surfaces adjacent to the boundary and this can, to some extent, be simplified by using symmetrical bicrystals. This technique has been used to study the effect of environment on Ds for silver and copper.'-'' Scratch techniques yield Ds values for the small orientation range exposed by the scratches (-2 deg). The multiple scratch process is preferable because the profile rapidly becomes sinusoidal and can then be interpreted theoretically in a relatively simple way. Also corrections for mass transport processes other than surface diffusion can be introduced easily. Mullins" considered a sinusoidal profile described the wavelength of the profile. After time t the profile can be described by the equation The terms A, A', C, and B which account, respectively, for contributions due to evaporation-condensation, diffusion through the gas phase, volume diffusion through the lattice, and surface diffusion are defined as: where Ds = the surface self diffusion coefficient ys = the surface energy per unit area p = the equilibrium vapor pressure over a flat surface pa = the equilibrium vapor density over a flat surface DG= the diffusion coefficient of vapor molecules in the inert gas DM = the mass transfer diffusion coefficient which for a pure cubic metal is Dv/f where Dv is the radiotracer diffusion coefficient and f is the correlation factor H = the molecular volume V = the surface density of atoms, il2'3 M = mass of an evaporating molecule

Jan 1, 1970

Uranium was the prima donna of 1978, and the exploration pace was fast and furious in traditional uranium producing areas such as the Powder River Basin, Grants mineral belt, San Juan Basin, and Uravan mineral belt. Adding fuel to the fire were several significant discoveries: The Date Creek Basin of western Arizona, the McDermitt caldera of northern Nevada, and Tallahassee Creek in Colorado's intermontane basin. Advances in solution mining haw resulted in deeper drilling on the Texas Coastal Plain, and the price of uranium has encouraged widespread exploration in frontier areas and non-sandstone environments.

Jan 5, 1979

By Alan J. Strauss

The Group IV elements germanium, tin, and lead form nine 1:1 compounds with the Group VI elements sulfur, selenium, and tellurium. This paper reviews the properties of the pseudobinary solid solutions formed by these compounds, including the extent of mutual solid solubility, temperature-composition phase diagrams, transport properties, deviations from stoichiometry, optical properties, and energy band structure. Particular emphasis is placed on the Pbl-xSnxTe and Pb1-xSn,Se alloys with rocksalt structure, because of current interest in these malerials for generating and detecting infrared ,radiation. ThE Group IV elements germanium, tin, and lead form nine 1:1 compounds with the Group VI elements sulfur, selenium, and tellurium. Some of the physical and electronic properties of these compounds, including their melting points1-8 and energy gaps,4'9-12 are listed in Table I. Four compounds (SnTe, PbS, PbSe, and PbTe) have the cubic rocksalt (Bl) structure. At room temperature GeTe has a rhombohedra1 structure closely related to the B1 structure, into which it is transformed at about 400°C. Four compounds (GeS, GeSe, SnS, and SnSe) have the orthorhombic B29 structure. In samples which have not been intentionally doped with impurities, the electrical conductivity is due primarily to electrons or holes produced by the ionization of donor or acceptor lattice defects associated with deviations from stoichiometry. Undoped samples of PbS, PbSe, and PbTe may be either n type or p type, depending on whether they contain excess lead or an excess of the Group VI element, respectively, but only p-type samples of the other compounds have been reported. This paper will review the properties of the pseudo-binary solid solutions formed by the nine 1:l compounds. The topics to be considered include the extent of mutual solid solubility, temperature-composition phase diagrams, transport properties, deviations from stoichiometry, optical properties, and band structure. Particular emphasis will be placed on the Pb]-xSnxTe and Pbl-xSnxSe alloys with B1 structure, which are promising materials for generating and detecting infrared radiation in the 8 to 14 µm atmospheric window and beyond. MUTUAL SOLID SOLUBILITY The extent of mutual solid solubility in the pseudo-binary systems has been investigated for fourteen of the eighteen ternary systems (in which the two terminal compounds have a common element) and for nine of the eighteen quaternary systems. In most cases, X-ray diffraction measurements made at room temperature were used to determine the structure and lattice parameter(s) of the phase(s) present in samples of various compositions prepared by freezing from the melt. In many investigations, including the extensive studies of Krebs and co-workers,13'14 the samples were annealed at elevated temperatures before X-ray measurements were made. In some cases, metallographic examination and thermal analysis have also been employed. The results for the ternary and quaternary systems are summarized in Tables II and III, respectively. (The original references should be consulted for the annealing temperatures.) The solubility of one compound in the other is at least 5 mol pct in all cases, and is often much larger. Complete solid solubility has been observed in all systems so far investigated in which the terminal compounds have the same structure, although in the PbS-PbTe system complete solubility is limited to elevated temperatures. Thus complete solid solubility occurs in five systems where both compounds have the B1 structure (SnTe-PbTe, PbS-PbSe, PbS-PbTe at temperatures above 805°C, PbSe-PbTe, and SnTe-PbSe) and in three where both compounds have the B29 structure (GeS-SnS. GeSe-~n~e-, and SnS-SnSe), as well as in two involving GeTe and a compound with B1 structure (GeTe-SnTe and GeTe-PbTe). On the basis of thermal analysis data, complete solid solubility at sufficiently high temperatures has also been reported for the GeSe-GeTe2 and SnS- pbS22 systems. This seems unlikely, however,

Jan 1, 1969

By N. Plint, E. T. Pinkney

Copper segregation developments within the Anglo American Group are described with particular reference to the work of the TORCO Project Team. Some radical differences have been adopted in the mechan-ical approach as compared with methods used by previous workers. Most important among these are the following: (1) The heating unit is a fluidized-bed roaster. (2) A moving bed segregation chamber is used in which a packed bed of heated ore, to which reagents have been added, moves slowly downwards under gravity and is discharged at a controlled rate by means of a suitable valve mechanism. (3) The quantities of reagents used (sodium chloride and some form of carbon) are only a small fraction of those recommended by other investigators. (4) Retention times of only 2 min have been satisfactory for the complete segregation of many types of ore. Project TORCO is the code name given by Anglo American Corp. of South Africa, Ltd., to the investigation undertaken in Zambia to develop a plant capable of treating large tonnages of refractory copper-bearing ores by the segregation process. The process has now been under investigation for six years (1960-66) and during that period work has progressed from bench-scale laboratory tests through small-scale pilot-plant work to a 500-tpd commercial pilot unit. Test work has been carried out mainly on the ores of the Zambian Copperbelt and results have been encouraging. Recoveries on the order of 90% have been obtained on all Copperbelt ores tested with high concentrate grades. PRELIMINARY INVESTIGATION Laboratory scale segregation test work on Nchanga banded sandstone ore and oxidized ore from Kansan-shi was originally conducted by Professor Maurice Rey. SociCtC Miniere et Mktallurgique de Pefiarroya, in 1957. Both ores responded favorably to the process. In November, 1960, the consulting engineers of Anglo American Corp. of South Africa, Ltd., initiated preliminary investigations on the application of the segregation process to Copperbelt ores. In early laboratory work, samples were heated in a closed, externally heated tube furnace for one hour at temperatures ranging from 800 to 830°C. The free moisture content of the ore feed was 4% and reagent additions were 0.5% salt and 1.0% coal. Both salt and coal were pre-crushed to -48 mesh Tyler.* All samples except an ore from the Bancroft mine responded well to the process and recoveries of over 90% were obtained. The presence of dolomite was responsible for the relatively poor recoveries from the Bancroft ore, as tests on one of the other ores with additions of calcite or dolomite confirmed. Recoveries dropped from 94 to 45% Grades of all the rougher flotation concentrates produced were low at 10-15% copper. The laboratory investigation was then followed up on a pilot-plant scale. An attempt was made to carry out the segregation process in a coal-fired 5-ft-diam fluidized-bed roaster. The feed consisted of a banded sandstone from Nchanga to which 0.5% sodium chloride had been added. Reducing conditions were maintained in the roaster by the addition of coal in excess of that required for heating the feed. The excess coal and added salt constituted the segregation reagents and the temperature of the bed was held at approximately 750°C. Recoveries of copper and grades of concentrates were low. The reasons for lack of success were: (1) The reaction is a gaseous one and a fluidized-bed roaster is therefore unsuitable as a segregation unit. (2) Coal in excess of that required for heating was introduced into the fluidized bed and under these conditions some of the copper would be reduced in situ. The term "reduced in situ" is used here to describe any copper mineral reduced to the metallic state in its original site within the ore particles. If this copper is to be recovered by subsequent froth flotation it must first be liberated from the gangue by grinding or some other method which, in practice, is usually not feasible on ores of this type.

Jan 1, 1969

THIS chapter and Chapter 19 deal with the heat quantities involved in open-hearth steelmaking, including the thermal efficiency of the furnace as a generator of high-temperature heat, the heat storage in the steel and slag produced and in the furnace walls, and the heats of the various chemical reactions involved in combustion, in the formation of slag, and the refining of steel. The fundamental laws to be considered are the first two laws of thermodynamics as applied to problems involving sensible and potential (or latent) heat only: 1. Heat input equals heat outflow plus storage. 2. Heat may be transferred only from a hotter body to a colder one. The first of these laws relates to the various heat quantities or the heat balance of the process. No one has as yet made an open-hearth heat balance that came out almost exactly balanced. The cause probably lies in the difficulties of accurate temperature and flow-rate measurements at the high temperatures involved and the inherent complexities in flow distribution and air leakage present in any real open-hearth furnace system. Some idea of the magnitude of the heat quantities involved can be obtained by a brief, crude outline of the heat balance of a 90-ton heat of steel. We will assume a 60 to 70 per cent hot- metal charge in a fairly fast working and only moderately efficient furnace, burning fuel oil, with a fuel consumption of around 4 million Btu per net ton of ingots produced. The time from start of charge to tap is 9 hr. During this period, the oil burned represents 360 million Btu and the sensible heat in the preheated air at the port ends is 117 million Btu, so that the total heat input at the ports is 477 million Btu. Heat stored in the steel and slag at the end of the heat, plus the various amounts absorbed by certain bath reactions, amounts to a total of 195 million Btu, but 120 million of this is supplied by the heat in the hot metal and by various heat-evolving reactions in the bath,

Jan 1, 1964

By E. J. Rapperport, E. D. Levine

The boundaries of the (a +ß) field in the Al-Li system were determined between 150°and 550°C utilizing quantitative metallography and lattice-parameter measurements. The solubility of lithium in aluminum decreases from 12.0at. pct Li at 550°C to 5.5 at. pct Li at 150°C. P Al-Li is saturated with aluminum at 45.8 at. pct Li and has this boundary value constant over the temperature range 150°to 550°C. THE solid solubility of lithium in aluminum has been determined by several investigators, 1-6 but, as shown in Fig. 1, there is little agreement among the various determinations. The earliest investiga-tions'-' are suspect because of the use of impure materials. Although high-purity materials were employed in more recent work,4'5 the experimental techniques may have led to contamination of the specimens. Probably the best work has been that of Costas and Marshall,6 who obtained close agreement between results obtained by two independent phase-boundary techniques: electrical resistivity and mi-crohardness. No detailed studies of the solubility of aluminum in the bcc ß phase, Al-Li, have been reported. Cursory investigations1,2,6 have indicated only that the (a+ß) -p boundary lies between 40 and 50 at. pct Li and is relatively independent of temperature. The present work was undertaken in order to provide an independent check on Costas and Marshall's determination of the solubility of lithium in aluminum, to extend knowledge of this solubility limit to temperatures below 225°C, and to make an accurate determination of the solubility of aluminum in Al-Li. EXPEFUMENTAL Alloy Preparation. In view of the difficulties encountered in previous investigations of the A1-Li system, close attention was paid to the use of methods of alloy preparation and treatment that would minimize contamination. Aluminum sheet (99.99 + pct Al) was vacuum-induction melted in a beryllia crucible to remove hydrogen. Lithium (99.9 pct Li) was charged with pre-melted aluminum into a beryllia crucible, in a helium-filled drybox. The crucible was sealed in a Vycor tube and transferred from the drybox to an induction furnace. Melting of alloys was performed by induction heating in a helium atmosphere. Solidification was accomplished by means of a suction apparatus, shown in Fig. 2, in which the alloy was forced by changes of pressure into a 3/16-in. inside diam closed-end beryllia tube. This technique produced rapid solidification of a small portion of the melt, resulting in alloys with a high degree of homogeneity. Typical lithium distributions are presented in Table I. Transverse sections 1/8 in. long were cut from the alloy rods, and each section was split in half longitudinally. One half of each section was analyzed for lithium, and the opposing halves were employed for phase-boundary determinations. Lithium contents were determined by flame photometry with an accuracy of 1 pct of the amount of lithium present. Thermal Treatments. Homogenization and equilibration heat treatments were performed in electrical-resistance furnaces with temperatures controlled to ± 2OC. Calibrated chromel-alumel thermocouples were employed to measure temperature. Homogenization was performed in helium-filled l?yrex tubes for 1 hr at 565°C. The encapsulated specimens were then transferred directly to furnaces maintained at lower temperatures for equilibration. Equilibration times were 2 hr at 550°C, 8 hr at 450°C, 27 hr at 350°c, 90 hr at 250°c, and 285 hr at 150"~. These times were chosen on the basis of conditions employed by previous investigators. Alloys were quenched from the equilibration temperatures by breaking the capsules into a silicone oil bath. By performing all possible operations either in sealed capsules or in a helium-filled drybox, the specimens were given minimum exposure to the atmosphere. Quantitative Metallography. Metallography of Al-Li alloys is difficult because of the atmospheric reactivity of the ß phase. It was found possible, however, to prepare surfaces of good metallographic quality by preventing contact with moisture during preparation. Grinding through 4/0 paper was performed in the drybox. The specimens were then transferred under kerosene to the polishing wheel. Three polishing stages were employed: 25-p alundum with kerosene lubricant on billiard cloth, 1-µ diamond paste on Microcloth, and 1/4-p diamond paste on Microcloth. Between stages the samples were cleaned by rinsing in trichloroethylene and buffing

Jan 1, 1963

By Ralph Hultgren, Prodyot Roy

Manganese vapor pressures from 1250° to 1500°K were measured by conventional Knudsen and torsion-effusion methods in twelve Fe-Mn alloys with compositions from 9 to 80 at. pct Mn. The Knudsen re-sults agveed approximately with previous measurements found in the literature which indicated the solutions were nearly ideal. However, except for the higher manganese compositions, the initial torsion readings indicated much (up to 50 pct) higher vapor presszcres than the Knudsen method. These high readings decreased steadily with time. The results are interpreted as due to depletion of surface concentration of manganese during evaporation. Thus the initial torsion readings are the most nearly correct and Knudsen methods on alloys are to be regarded with suspicion unless it can be demonstrated that diffusion rates are rapid enough to replenish surface concentrations depleted by evaporation. For the lowest manganese contents and highest temperatures, depletion causes initial torsion readings to be too low. FEW thermodynamic data are available for alloys of the high-melting transition metals due primarily to experimental difficulties at the high temperatures involved. Measurement of vapor pressure is one of the most promising techniques to be applied to this problem. From the vapor pressure of a component of an alloy phase, its activity or its partial molar Gibbs energy can be directly calculated. From measurements over a sufficient range of temperatures and compositions the partial and integral Gibbs energies, enthalpies, and entropies may be determined through Gibbs-Duhem integration. For solid phases of variable composition, which commonly are found in alloys, measured vapor pressures are low because of depletion of the surface concentration of the more volatile component as it is selectively vaporized. Diffusion from the interior tends to restore the depleted concentration. The seriousness of the effect depends on the relative rates of diffusion and vaporization. Although depletion has been recognized as a factor,' direct measurements of its rate are lacking and quantitative estimates of its effect on measurements are uncertain. In the present work the vapor pressures above a series of Fe-Mn alloys have been measured by the torsion-effusion method. This method permits a measurement of the vapor pressure of manganese as soon as the sample comes to temperature and continuously throughout the run. The rate of decline of the apparent vapor pressure measures the rate of depletion and the torsion reading at zero time should be correct. At high manganese concentrations (70 and 80 at. pct Mn), the torsion reading remained nearly constant; surface depletion was negligible. There was a slow drop in measured pressure due to bulk loss of manganese from the sample. However, for lower manganese contents, especially at high temperatures, the depletion effect was considerable and the apparent vapor pressure decreased steadily with time. For these alloys the Knudsen method should give pressures which are too low. The measurement of pressure in the Knudsen method cannot be made until enough manganese has been vaporized to be weighable. The Knudsen result is therefore an integrated average between the initial and final vapor pressures. To verify this, conventional Knudsen measurements were also made with results in approximate agreement with previously published Knudsen measurements for these alloys. As expected, for the lower manganese contents the initial torsion readings of pressure were as much as 50 pct higher than the Knudsen. THE APPARATUS The apparatus, shown in Fig. 1, is capable of operating to temperatures up to 2000°K. A vacuum of 5 x 10"6 mm Hg can be maintained. Temperatures were measured to +3° by a W-Re thermocouple placed in the dummy cell, N. The torsion cell consisted of two alumina crucibles with holes in their covers. They were placed in the molybdenum holder shown suspended from the torsion wire. Rectangular torsion wires (1 by 4 mm) were found to have superior sensitivity and less residual distortion than circular (2-mm-diameter) wires. The precision of angular measurement was found to be *0.025 deg; the variation is no doubt due to temperature fluctuation and vibration. For the Knudsen experiments the alumina cell was set on the support, N. SAMPLE PREPARATION 800-g samples were prepared by melting together electrolytic iron and electrolytic manganese and pouring the melt into a water-cooled copper mold. Melting and pouring were done under a helium at-

Jan 1, 1965

By L. J. Buck

Lead smelting in the ore hearth, first practiced in Europe, whence the traditional name "Scotch hearth, and later developed in America to its present efficiency, is based upon the roast reduction process and is applicable to relatively high-grade galena concentrates. The roast reduction process embraces not only the roast reactions: 2PbS + 70 = PbO + PbSO4, + SO2 PbS + 2PbO = 3Pb + SO2 PbS + 2PbSO4 = Pb + 2PbO + 3SO2 but also, by reason of the addition of extraneous fuel, the reduction reaction: 2PbO + C = 2Pb + CO2 Thus in the ore hearth the roasting of lead sulphide, forming lead oxide and lead sulphate, the reaction of lead oxide and lead sulphate with lead sulphide producing lead, and the reduction of lead oxide by carbon producing lead, go on simultaneously. The Ore Hearth Function of the Ore Hearth.—The ore hearth, through the roast reduction process, provides for the extraction of lead and the immediate disposal of the smelting products—fume, hearth slag and lead, all of which are subsequently treated as follows: the fume is recovered in the flue system and baghouses, mixed with concentrates, and returned to the ore hearth in the form of ore mix; the hearth slag is smelted in a blast furnacc for the further recovery of lead; and the lead is refined and cast into a marketable product. Concentrates Amenable to Ore-hearth Smelting.—Only concentrates that are high in lead content, 70 per cent or more, are ordinarily considered amenable to ore-hearth smelting. Silica should not exceed 2 per cent and iron is allowable only to 4 per cent. These slag-forming constituents occlude both sulphide and mctallic lead, prevent the elimination of sulphur and decrease the lead extraction. Lime and magnesia are beneficial in that both favor a dry slag, promote elimination of sulphur, and prevent occlusions of metallic lead.

Jan 1, 1937

By M. J. Richards, A. Szirmae, G. R. Speich

The formation of austenite from ferrite, ferrite plus retastable carbide, spheroidite, and pearlite has been studied in a series of irons, Fe-C alloys, and plain-carbon steels using fast heating techniques. In the absence of carbide, austenite nucleates at ferrite/ferrite grain boundaries; nucleation is followed by the rapid growth characteristic of a massive transfornation. The trarnsformation occurs at 950°C at heating rates of 106º C per sec and cannot be suppressed. Metastable carbide dissolves before austenite forms and does not influence the transformation kinetics. For spheroidite structures, austenite nucleates preferentially at the jinction between carbides and ferrite grain boundaries. Growth from these centers proceeds until the carbide is completely enveloped; subsequent growth occurs by carbon diffusion through the austenite envelope. For pearlite structures, austenite nucleates preferentially at pearlite colony intersections. Carbide la)?zellae dissolve at the advancing austenite interface but complete solution of carbide does not occur; the residtial carbide is eventually dissolvled or spheroid-ized depending on the carbon cuntent. The magnitude and temperature dependence of the austenite growth rate into Fe-C pearlite when incomplete carbide dissolution is assumed are satisfactorily explained by an approximate colume diffusion model. The impurities present in plain-carbon steel reduce the growth rate of austenite in comparison to that jound in an Fe-C alloy. The formation of austenite has been studied in much less detail than the decomposition of austenite. This is primarily a result of the importance of harden-ability in determining the mechanical properties of steel. Recently, more interest in the kinetics of austenite formation has resulted from the discovery by Grange1 that rapid heating techniques strengthen steel by refining the austenite grain size. Although the strengthening effect is not large, it is accompanied by no loss in ductility. In addition, interest continues in rapid heat treatment of low-carbon steel sheet for tin plate applications.2,3 Among the few systematic studies of austenite formation are the early work of Roberts and Mehl4 on formation of austenite from pearlite and recent work of Molinder5 and of Judd and paxton6 on formation of austenite from spheroidite. Also, Boedtker and Duwez7 and Haworth and paar8 have recently studied the formation of austenite from ferrite in relatively pure iron, Kidin et al.9,10 have studied the formation of austenite in 8 pct Cr steels, and Paxton has recently discussed various aspects of austenite formation in steels." The present work was undertaken to determine the kinetics of austenite formation for a variety of starting structures including ferrite, ferrite plus metastable carbide, ferrite plus spheroidal cementite, and ferrite plus pearlitic cementite. Emphasis was placed on determining the active sites for austenite nucleation, determining the temperature and time range of austenite formation, and in the case of pearlite a careful study of the growth rate of austenite was made in the absence and presence of impurities. By using a variety of heating techniques including laser-pulse heating, it has been possible to study austenite formation in an isothermal fashion over a wide range of temperatures. EXPERIMENTAL PROCEDURE The alloys studied in the present work are a zone-refined iron with 4 pprn C, an Fe-C alloy with 130 pprn C, 2 Fe-C alloys with 0.77 and 0.96 wt pct C, and a plain carbon steel with 0.96 wt pct C. The zone-refined iron and Fe-C alloys contained 60 pprn and 200 pprn total substitutional impurities, respectively. The plain carbon steel contained 2400 pprn Si, 2000 pprn Mn, and 900 pprn Cr. Various heat treatments were given to these alloys to produce different starting structures of equiaxed ferrite, ferrite plus metastable carbide, fine pearlite, and spheroidite. These heat treatments are given in Table I. A wide range of heating rates were employed in this work because many of the reactions occur so quickly at temperatures in the austenite range that they are completed during the initial heating cycle unless very fast heating rates are used. Essentially the same heating techniques employed by Speich et a1.12 and Speich and Fisher13 were used in this work. For time intervals of 2 sec to 20 hr, simple hand immersion of 0.010-in. thick specimens in a Pb-Bi bath was employed. These specimens were quenched in a 10 pct NaC1, 2 pct NaOH aqueous bath. For time intervals of 100 m-sec to 2 sec, an automatic dunking and quenching device was employed with 0.002-in. thick specimens. Again, liquid Pb-Bi baths were used for a heating medium but now helium gas quenching was employed. For time intervals of 2 to 100 m-sec a laser heating device was employed with 0.002-in. thick specimens; a helium plus fine water-droplet spray was now used for quenching. Additional information on heating times shorter than 2 m-sec was obtained by study of the zones around the centrally heated laser spot. Here diffusion of heat from the centrally heated zone raises the temperature of the specimen locally to all temperatures between ambient and the peak temperature, but for times of the order of microseconds. All the heat-treated specimens were examined by

Jan 1, 1970

By Arthur F. Taggart

The major portion of the work described in this paper was performed by R. W. Young, † a graduate student in the department of Mining and Metallurgy, Sheffield Scientific School, Yale University, working under a cooperative agreement between the Hardinge Conical Mill Co., the Sheffield Scientific School, and himself. Since this cooperative scheme is at present in effect in the case of a considerable number of other students in the department and since it is the hope of the school that the privilege thus extended may be utilized even more freely in the future by mining and manufacturing companies, it may not be amiss at this point to give a summary of the general plan. Briefly it is as follows: A graduate student, whose undergraduate work in this or other universities shows promise of ability to handle research work, is chosen by conference between the company and instructor involved. The aim of the company in the agreement then entered into is to obtain the solution of one or more of the technical problems with which it may be confronted, or, at the end of 1 or 2 years, to obtain as an employee a man especially trained in its work. As a means to accomplish one or both of these ends, the company furnishes the machine, apparatus, or material to be tested and pays the student during his graduate work a small salary, usually just sufficient to cover his living expenses, tuition and fees. The aim of the student is special training along a line in which he is particularly interested, the attainment of his advanced degree, and the chance to show to his future employer ability to handle such problems as may be presented to him. In return for the financial aid which he receives he agrees to devote at least half of his working time to the special problem submitted by his company. The other half is devoted to study of the collateral subjects required by the department for the granting of the degree which the student seeks. The student further agrees to enter the employ of the company in question at a wage not greater than that paid in like positions to recent graduates not specially trained and to remain with his employer at such wage for at east 1 year. If the student is to obtain a degree, the special work forming the basis of his investigation must be such as will involve real research and not mere routine manipulation. The subject is chosen by conference between the three parties to the agreement. The work is carried on under the direct supervision of the instructor involved. The school furnishes the general laboratory and library equipment essential to the pursuit of any extended investigation. In

Jan 1, 1918

By T. A. Rickard

From time immemorial the fissure-vein has been held the simplest type of ore-deposit. The prominence given to it by Cotta and his disciples, from their study of the mines of the Erzgebirge, is impressed upon technical literature; and, in consequence, the ores which carry the valuable metals have been supposed to occur mainly in fissures, cleaving the rocks in diverse directions, and the noblest type of vein has been deemed that which cut across the country independent of its structure, whether evidenced as bedding, foliation or cleavage, and which was identified with rents produced in the rocky crust of the earth. As so conceived, the vein was a fissure filled with ore, extending through the country for a varying distance, and continued downward to a depth more or less proportionate to its longitudinal extent. The vein-material was bounded by an encasement of rock, and those immediate surfaces which limited it on either aide were called " walls." These primary conceptions have become modified by the experience of modern mining in widely separated regions. The study of lode-formations has led to the recognition of notable departures from the supposed normal structure of the veins of Saxony and Cornwall, the two classic homes of early economic geology. Typically the walls of a vein are conceived as parallel rockplanes enclosing the ore; the upper one being called the " hanging-," and the lower the "foot-wall."* Walls are rarely alike. Even where a vein traverses a homogeneons formation, such as a massive crystalline rock, it is usually found that the surface which bounds it underneath

Jan 1, 1897

By D. Cox, M. J. Duggin, L. Zwell

The carbides of approximate composition and Mn have been studied using X-ray diffraction techniques. Those carbides of the type (Fe,Aln)zC ave isostructural with cementite. The cell pararmeters a and c have minimum values at approximately 10 at. pd substitution of manganese for iron; no satisfactory explanation has yet been found for this phenomenon. The carbide fFeMn4)C has a monoclinic unit cell whose dimensions are close to those of ,11,15Cz A neu-troip-dij~ractiot~ study of (F'eAlrz4)C~ reveals that, like MnsCZ, it is isostructural with Pd5Bz. The iron and manganese atoms occupy the palladium atom sites, while the carbon atoms were found to have the same atomic coordinates as the hovon atoms. A neutrorr-diffraction study of indicates that the carbon-atom positions are very close to those occupied in (Fez.,ll/lr~,.3)C. In both carbides studied, tlre iron and manganese atomzs were found to be essentially randomly distributed, although, in the case of (Fe,.811fn1.2)C, it is possible that there may be a slight preference of manganese atoms for- the general (d) positions and a corresponding slight preference of iron atoms for the special (c) positions. It has been found that a complete range of solid solution exists between Fe3C and Mn3C at 1050°C,I although Mn3C becomes unstable when the temperature is reduced to 95O0C,' and can only be retained by rapid quenching. It is also known that a complete range of solid solution exists from Fe5Cz to M~SC~,~ although the stability range of carbides of the type (Fe,Mn)sCz as a function of the relative proportions of iron and manganese is not known. X-ray examinations of Oh-man's carbide3 and Spiegeleisenkristall,~ which have the approximate compositions (Fe3.67Mnl.33)C2 and (Fe3-,Mn,)C, where x lies between 0.4 and 1, respectively, have been made. The following carbides have also been studied: ] The lattice parameters determined during these investigations are listed in Table I. It is seen that carbides of the type (Fe,Mn)sCz have a monoclinic unit cell while carbides of the type (Fe,Mn)3C have an orthorhombic unit cell. It is evident that the variation of lattice parameters with manganese content is not linear for carbides of the type (Fe,Mn)3C. The coordinates of the atoms in (Fe2.7Mno.3)C have recently been determined by single-crystal analysis., The fractional atomic coordinates have been shown by Fasiska and jeffrey to be in good agreement withj those deduced from an earlier analysis of Fe3C by Lipson and etch.' However, it was impossible to determine whether iron and manganese atoms occupied ordered positions because of the small difference between the atomic scattering factors of iron and manganese. The atomic positions in Mn5Cz (Refs. 8 and 9) and Fe5C2 (Refs. 7 and 8) have been obtained only by comparisons made with the isostructural compounds P~SB~.' Since X-ray diffraction techniques were used in these investigations, accurate positioning of the carbon atoms, which have a low atomic scattering factor, was difficult. No attempt has been made to determine the atomic positions in the other carbides previously studied. It was felt that an investigation of the lattice parameters of a number of intermediate carbides of the types (Fe,Mn)sCZ and (Fe,Mn)& would be of interest. It seemed likely that a neutron-diffract ion study of such carbides would indicate whether ordering occurred between the iron and manganese atoms because of the large difference between the neutron-scattering cross sections of iron and manganese. It also seemed probable that such an investigation would provide a determination of the atomic coordinates of the carbon atoms. I) EXPERIMENTAL DETAILS Specimens, each weighing approximately 20 g, were carefully prepared according to the following proportions: The components were 500-mesh powders of 99.995 pct purity iron and spectroscopically pure carbon and a 200-mesh powder of 99.995 pct purity manganese. The component powders were intimately mixed by prolonged shaking, then each specimen was inserted into a spot-welded cylindrical container of tantalum foil, whose end was closed but not sealed. Each specimen in its envelope was then sintered at 1050° C for 24 hr in a thin-walled evacuated quartz capsule, such a time having been previously found sufficient for equilibrium to be attained.' Each specimen was then quenched in order to attempt to retain the high-temperature phase, as the literature indicates that transformations may occur on cooling. Debye-Scherrer X-ray photographs were taken of each specimen using a 114.6-mm-diam camera, Fig. 1, patterns 2 to 6. The exposure time was 6 hr using filtered iron radiation at a tube voltage of 40 kv and a tube current of 12 ma.

Jan 1, 1967