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By L. Graydon Carpenter, Donald E. Garrett

Probably the largest single tungsten deposit in the U. S. is one that has yet to produce any tungsten; it is not even listed in tables showing U. S. reserves. This deposit is at Searles Lake, Calif., where the brine contains about 70 ppm WO3. Small though the concentration is, the amount of brine in the lake is so great that it is estimated at 170 million lb of WO3 (8.5 million units), equaling the total of all other known reserves in the U. S. This dry lake is located in the basin formed by the Argus and Slate ranges in the northeastern corner of the Mojave Desert (Fig. 1). It consists of a massive salt body some 35 sq miles in extent, varying in thickness from 0 to 120 ft and containing more than 3 billion tons of salt. A central section of about L. G. CARPENTER and D. E. GARRETT are with the Research Dept., American Potash & Chemical Corp., Trona, Calif. TP 47881. Manuscript, Nov. 4, 1958. AlME Trans., Vol. 214, 1959. 8 sq miles is exposed salt; the remaining area is overlain by mud. Except for the exposed sodium chloride (10 to 15 ft thick) the body consists of a very complex and heterogenous mixture of soluble crystals, containing at least eight major phases. It is only about 53 pct solid, the voids being occupied by a dense brine which is the raw material for two chemical companies operating in the valley.

Jan 3, 1959

By W. R. Ingalls

Jan 1, 1922

By W. R. Ingalls

Jan 1, 1922

By H. J. McQueen, G. C. Kuczynski

ThE useful properties of k6-Alfenol and the thermenol alloys (summarized by J. E. Nachman and W. J. Buehler') have encouraged research into those alloys which, having compositions near 25 at. pct Al, 75 at. pct Fe, undergo an ordering transformation between 500° and 550°C. The equilibrium diagram of the iron-rich iron-aluminum alloys was first determined by Bradley and Jay.2,3 They distinguished three structures: the disordered, which is bcc; FeAl, which is simple cubic with one Fe atom and one Al atom per lattice site; and Fe3Al, which is fcc with three Fe atoms and one Al atom per lattice site. This last structure may be conceived as eight body-centered cubes with Al atoms at the centers of alternate ones, so that Al atoms have only Fe as first and second nearest neighbors. Later, Osawa and Murata4 determined the boundaries of the different superlattice regions in the iron-rich Fe-Al solid solution by dilatometric and X-ray methods. Two of the six phases, which they found between 0 and 50 at. pct Al, were never identified by other workers. A. Taylor and R. M. ones,' concurrently with and independently of this research, redetermined this diagram by measuring lattice parameters at temperature. Further reference will be made to their results. EQUILIBRIUM DIAGRAM A) Experimental Procedure—The alloys were prepared from NRC gas-free iron (99.9 pct Fe) and 99.99 pct Al. They were melted in alumina' crucubles under vacuum in an induction furnace. If the ingots were unsound or showed more than a 3 pct weight loss, they were rejected. The ingots were homogenized at 1100°C for 24 hr before hot-rolling at 700" to 900°C from 1 in. to 0.25 in. thick. It is interesting to note that alloys containing more than 28 at. pct Al were able to take an 80 pct reduction of area in the ordered FeAl region. The dilatometer specimens and specimens for chemical analysis were cut from these plates. A Leitz dilatometer was used to measure the coefficient of linear expansion as a function of temperature. The specimens were slowly cooled from 900" to 300°C where they were annealed for 10 hr before slowly cooling to room temperature. The expansion was recorded photographically as the alloys were heated and cooled at 1.5oC per min between 100" and 1000°C. The magnetic measurements determining the Curie points were performed on the dilatometer specimens in the dilatometer, by using direct current in the furnace windings both to heat the specimen at 1.5°C per min and to serve as the primary coil. The regular specimen holder was replaced by a thin alundum tube on which was wound, as a secondary, a closely spaced coil of fine copper wire. A reversal of current in the furnace windings produced, in the secondary, an emf which was proportional to the induction in the specimen, and which could be detected on a galvanometer. In this way, it was possible to determine the Curie points of alloys to within * 5°C. B) Results and Discussion—The coefficients of linear expansion, as obtained from the dilatometer curves, are plotted as a function of temperature and composition in Fig. 1. Since in many cases the heating and cooling curves were superimposed, it was decided to measure slopes of heating curves only. Definite high peaks, due to the transformations, were observed on each curve between 400" and 520°C, and, for alloys containing more than 23 at. pct Al, a second set of smaller peaks at temperatures above 660°C. The sharpness of the peaks indicates that ordering to Fe3Al is a first-order reaction. The compositions and the temperatures of these peaks are listed in Table I. When these temperatures are plotted as a function of the composition, the result is the constitution diagram presented in

Jan 1, 1960

By Y. L. Yao

The solubility limit of oxygeu in a titanionn at 850°C has been determined by magnetic measurements as 12.5 + 0.5 pct (29.0—30,9 at. pct). Also in the susceptibility-co~centmtion curve, there is n distinct but not sharp maximum at 19.0 * 0.5 z~t pet (40.5--42.1 at. pct), which indicates the presence of the 6 phase. The magnetic susceptibilities of the alloys of the titanium-oxygen system were determined byEhrlich in 1939.' is alloys were prepared by sintiring mixtures of titanium and its oxide at high temperatures. On the titanium-rich side, despite the fact that the alloy compositions he measured crossed two phase boundaries, the limited results obtained

Jan 1, 1960

By R. L. Orr, O. D. Sherby, J. E. Dorn

Creep data for pure metals at temperatures above those at which rapid recovery occurs (above about 0.45 the melting temperature) are correlatable by means of the equations and These correlations were applied successfully to data for aluminum, iron, nickel, copper, zinc, platinum, gold, and lead as well as for simple alloys. For a given metal, AH is a constant about equal to the activation energy for self-diffusion. THE early recognition that creep is stimulated by thermal activation prompted numerous investigators', ' to apply the same laws that are valid for the viscous behavior of liquids to analyses of creep data. To a good first approximation these laws are summarized by the equation were i is the creep rate; T, the absolute temperature; R, the gas constant in cal per degree per mol; T, the applied stress; AH, the activation energy in cal per mol; S', a constant dependent on the entropy of activation, the frequency of activation, and the contribution of each activation to the strain; and B', a constant dependent upon the size of the flow unit that is activated. Although Eq. 1 describes the flow of viscous fluids very accurately, its application to the creep of solids has been disappointing. Two reasons for the failure of Eq. 1 for creep are now known: First, the decelerating creep rates during primary creep demand that the internal structures of the metals are changing and that these changes might be reflected by changes in one or more of the three creep parameters S', AH, and B'. Consequently when the conventional methods of evaluating these parameters are employed by comparing the secondary creep rates at a series of temperatures and at a series of stresses, the true effects of temperature and stress are masked because of simultaneous changes in one or more of the creep parameters. Second, as will be illustrated more fully later, three distinctly different types of investigations have shown that the stress term for high temperature creep enters the analysis as and not as t/T. This rather unexpected result points sharply to a significant fundamental difference between the mechanisms for the viscous behavior of fluids and the high temperature creep of metals. More recent extensive investigations",'4 on the creep of high purity aluminum and several of its dilute a solid solutions have shown that where E is the total creep strain; temperature compensated time = for a constant temperature test; t, is the initial stress in a constant load test or the constant true stress in a constant stress test;" and t is the duration of test. • Eq. 2 is valid for either constant stress or constant load tests; but the total creep curve is difierent (i.e. the function, f, is differentl for each test. The fundamental origin of the validity of Eq. 2 is thought to arise from some equivalence of substructures in constant load creep tests at equal values of 0 or E. Both X-ray and metallographic examination revealed that practically identical lattice distortions and subgrain sizes are obtained for the same creep strain under the same load over wide ranges of test temperature.' Furthermore the room temperature tensile properties following equal creep strains at different temperatures under constant loads were also identical, illustrating that identical substructures were indeed obtained.' " A second type of correlation was obtained by differentiating Eq. 2 with respect to time, giving For the secondary creep rate, denoted by i,, But Eq. 2 reveals that 0, is a function of a, alone because the secondary creep rate in a given constant load test is always reached at a fixed value E,". Therefore where is the function that was introduced a number of years ago by Zener and Hollomon." Eq. 4 is not only valid for creep but will permit correlations between constant load creep and tensile data where e, is the rate of tensile straining and (TC is the engineering ultimate tensile strength.

Jan 1, 1955

By James P. Graham

In this paper the authors present the background information to preheating and its development and describe the facilities available at the British Carbonization Research Association. The results obtained at BCRA including those collected during commissioning of the new plant at BCRA are summarized and reference is made to the research work planned for the next stage. The enlarged test-plant facilities at the British Carbonization Research Association embrace both 10-t (4.2 m high) and 17-t (6.5 m high) test ovens and two types of preheater, the preheated coal being charged by either gravity or pipeline. During the commissioning of the new equipment, an extensive programme on preheating was initiated and the paper includes the results obtained so far. These cover tests on wet and preheated coals and commercial blends, the use of Rosin- and Cerchar-type preheaters and charging by both gravity and pipeline systems to the 10-t and 17-t ovens. Accordingly, for the first time a direct comparison is possible of both the performances of the plants and the qualities of the respective products. The paper also outlines the findings from a study of the carryover of solids during the charging of preheated coal and of the treatment of those solids.

Jan 1, 1977

By W. V. Green

Creep of tantalum was measured at temperatures from 0.6 to 0.89 of the absolute melting temperature. The creep curves include first, second, and third stages. Steady-state creep rate depends on the fourth power of stress. The activation energy for creep throughout this temperature range is approximately 114 kcal per mole, measured by the aT technique. Subgrain formation occurs as a result of creep strain, and pile-up dislocation arrays are observed in etch-pit patterns. BECAUSE of its high melting point-which is exceeded only by those of rhenium and tungsten—and its high room-temperature ductility compared to most of the other high-melting-point metals, tantalum will undoubtedly be utilized in an increasing number of high-temperature applications. Alloying studies directed toward increased high-temperature strength must use data on tantalum itself as a base line in order to evaluate the effectiveness of the alloying additions. However, to date, no systematic study of creep of tantalum at temperatures above one-half of its melting point has been reported in the literature. Conway, Salyards, McCullough, and Flagella1 have measured linear creep rate of tantalum sheet as a function of stress, but at only one temperature, 2600°C. This paper describes a relatively thorough study of the high-temperature creep of tantalum. METHOD Material Tested. The commercially supplied, l/2-innch-diameter tantalum rod used for this work was electron-beam-melted, cold-forged, rolled, swaged, cleaned chemically, and vacuum-annealed for 1 hr at 1000°C, all by its manufacturer. The vendor's analysis included 60 to 170 ppm C, 3.4 to 4.2 ppm H, 60 to 80 ppm 0, 15 ppm N, and a hardness ranging from 66 to 81 Bhn and averaging 76 Bhn. Creep eimens Used. Two creep-tested specimens are shown in Fig. 1. The 1/4 in.-diameter gage section was 3/4 to 1 in. long, and terminated either at shoulders 5 mils high or at 20-mil-diameter tantalum wires spot-welded to the circumference of the gage section. Both kinds of shoulders served equally well as fiducial marks for optical strain measurements. The spot welding did not alter the creep behavior in any detectable way; the 5-mil- high sharp shoulders did not result in any detectable localized effect on the strain. Before testing, each tensile bar was first mechanically polished -id then electrochemically polished according to the method referred to by Forgeng2 as the "Thompson Ramo Woolridge" method, which was suitable for tantalum after small adjustments of technique were made. Two tensile bars tested at low stresses had 1/8-in.-diameter gage sections and utilized only the weight of the bottom grip for the applied load. Although these diameters were smaller than were desired for other reasons, applied loads were known with high precision in the tests in which they were used. Testing Procedure. Two different constant-load creep-testing machines were employed, one of which has been described by Smith, Olson, and Brown.3 In both, the tensile bar is held vertically on the axis of a cylindrical tungsten tube or screen heater by threaded tungsten grips. The tensile bars and associated grips are heated by radiation from the incandescent heaters, which are heated by their own electrical resistance. Both testing machines use pins to hold the bottom grips in place. The load is applied to a tensile bar through hanging weights, a constant force-multiplication lever, a pull rod sealed to the chamber lid, and a top grip threaded to the pull rod at one end and to the tensile bar at the other. In one machine, the vacuum seal is a bellows with a low spring constant; in the other, the seal involves a rotating "0 ring". With the latter, rotation is converted to translation with a crank shaft, so that elongation of the tensile bar is accommodated with no change of tensile load. The incandescent tensile bar is viewed by an external optical system through slots in the radiation shields and heater, and an enlarged image is projected on a ground-glass screen. Gage-length measurements are made on this image with cathetometers on traveling microscopes. With regard to creep-test results, the two machines were identical. Thorium oxide coatings were applied to the threaded ends of the tensile bars, to prevent diffusion welding of the tensile bars to the grips during testing. Specimen temperatures were measured with an L. & N. optical pyrometer which had been calibrated against a standard carbon arc, and were corrected fir window absorption by calculation from the measured spectral transmittance of the quartz observation windows. Longitudinal temperature gradients in the tensile-bar gage length and temperature drifts during testing were detectable but small, and were estimated to be 10°C or less. Accuracy of temperature measurement was confirmed by comparing the temperature measured on the surface of a special

Jan 1, 1965

By Clyde E. Williams

MODERN steelmaking has gradually evolved from an inefficient small-scale operation, utilizing tiny units, to a highly efficient one utilizing large units almost completely mechanized. The leading position of the United States in the steel industry is due to the possession of easily assembled high-grade raw materials and large-scale highly developed plants. Modern integrated steel operation is largely one of material handling. Mining and transportation of iron ore, coal, and fluxes, coking the coal in by-product ovens, smelting the ore with coke and limestone in blast furnaces, converting the resulting pig iron into steel in the Bessemer, or in the open hearth when supplemented with scrap, and fabricating the resulting steel ingots in the shapes of commerce all are done on a huge scale, mostly by machinery, most of the labor being employed either in operating the machine or keeping it in repair. At the same time the quality and uniformity of the product have been greatly increased by the development of scientific knowledge and operating skills. Thus, the steel industry in this country is highly developed both mechanically and technically.

Jan 1, 1942

By M. E. Grimes

A method of predicting uranium extraction rate in underground bacterial leaching of as-blasted ore has been developed. The method is based on the hypothesis that extraction is directly proportional to the volume of a piece of ore that is penetrated by the leaching agent, and that each piece of ore, regardless of size, has been penetrated radially from the external surface to the same depth at a given leaching time. Consequently, given an initial size and assay distribution, extraction should be predictable for a given penetration. Two laboratory bacterial leaching tests under simulated stope leaching conditions were carried out for a periods of 89 and 96 weeks. On completion, the residual ore was screened and analyzed. The information so obtained was employed to provide an elapsed time base for the penetration model, from which predictions of future extraction were made. The finding of solubilized uranium trapped within apparently sol id ore pieces was unexpected. However, it is not recoverable except after crushing, and hence does not affect calculations of underground leaching .

Jan 1, 1974

By H. B. Aaron

DIFFRACTION micrography provides a useful tool for studying complex strain fields. Newkirk1 observed an X-ray diffraction effect due to strain interactions at the intersection of slip lines in LiF and attributed this effect to a reduction of primary extinction due to strain-gradients. It is the object of the current study to correlate the observations and rationalize the appearance of the strain fields around intersecting slip planes in x-ray extinction contrast studies by a semiempirical method. Fig. 1 shows typical strain patterns observed by Newkirk on (200) reflections in LiF using chromium radiation and a variation of the Berg-Barrett (B-B) technique. In other reflections (e.g., {220}) the strain patterns are absent. If one defines a vector 3 along the striations from the more to the less dense portion, and a vector v in the direction of the X-ray beam parallel to the reflecting plane, such that both 3 and v lie in the reflecting plane of the crystal, Fig. 1, then the positions of the striations on (200) reflections appear to be confined such that where is the angle between v and 3, Fig. 1. It will be seen later that 3 is the projection of Sb on the reflecting planes. In some instances the strain effects are observed only above the slip plane, Fig. l(C), in others only below, Fig. l(D). Where the strains appear to be both above and below the slip line, Fig. l(E), one sees that there are actually two slip lines present, Fig. 1(F) and (G). Furthermore, the strain patterns are observed in the vicinity of the intersection of emergent edge dislocations (diagonal lines on the micrograph) with emergent screw dislocations (horizontal lines on the micrograph), e.g., Fig. l(HJ) and (MN), but not in the vicinity of the intersection of two emergent edge dislocation arrays, Fig. l(HJ) and (KL). However, in the latter case the contrast of the dislocation array itself is reduced along (KL) and increased along (HJ). Line (AB) in Fig. 1 is the edge of the crystal, and the material to the left of (AB) is the side of the sample. Any empirical correlation will have to include most, if not all of the aforementioned observations. Lang2 and Newkirk1 determined that the condition for contrast is given by g-b#0 [2] where b is the Burgers' vector and g the diffraction vector (normal to the reflecting planes). If one assumes that the direction of the observed strain fields is given by the vector sum of the slip vectors (Sb1) of the intersecting slip planes, then Eq. [2] may be generalized to g-S#0 [3]

Jan 1, 1963

By Roscoe C. Clark

A study of the effect of various corrosive waters on five different types of cements indicated that those cements containing less than 5 per cent tricalcium aluminate were the most resistant to corrosion. INTRODUCTION Oil field experience indicates that some formation waters are excessively corrosive to some cements. A study has been made to determine the most corrosive waters and the relative resistance of various types of cement to these corrosive waters. PROCEDURE This study of the corrosion of cement was limited to an evaluation of the effect of various corrosive waters on five different types of cements. A determination was made of the types of cements which offer the greatest resistance to corrosion and of the chemical contents of waters found to be the most destructive to cements. The laboratory tests were conducted by immersing small cylinders of 20 samples of set cement in various corrosive waters. The initial mixing of most of the cement slurries was done using Tulsa tap water as the mixing water; however, a few of the slurries were mixed using a sulfate water as the mixing water. When the 20 samples had been immersed for a period of nine months, they were removed from their corrosive baths, inspected and photographed. The field tests were conducted by placing ten cement specimens, each contained in a screen wire thimble, in the gas anchors of insert pumps in wells in six different fields. The specimens were observed whenever it was necessary to pull the pumps and were returned to the 1aboratory when deterioration had reached a serious state. EXPERIMENTAL RESULTS In both the field and laboratory tests the resistance of the different cements to corrosion was determined by the condition of the samples and the number of samples of each that were in good condition at the end of the immersion period. Table I lists the various cements tested in the order of their resistance to corrosion with the most resistant at the top and the least resistance at the bottom. An analysis of Table I reveals that the upper half of the table predominates in slow set and sulphate resisting cement.: (low 3CAO.A12O3 content) while in the lower half the port-land and high early strength cements are in the majority. Of the last two the high early strength cements appear to offer less resistance to sulfate water corrosion. An investigation of the composition of the cements listed in Table I indicated that as a general rule the cements containing less than five per cent tricalcium aluminate will exhibit the greatest resistance to corrosion. The destructive action of various corrosive waters was compared by listing the waters in the order of those containing the largest number of samples that failed completely. These ratings are presented in Table II with the most corrosive waters listed first and the least corrosive last. Figs. 1 and 2 show the end views of cement samples which had been immersed in the baths of five per cent MgSO, water and the five per cent Na2SO4 water respectively in the laboratory. These .pictures are representative of the results obtained and present a pictorial evaluation of the corrosive resistance of various types of cements.

Jan 1, 1950

By David W. James

DURING a recent investigation of the transition from a planar to a nonplanar solid/liquid interface for the systems H20-NH4F and 0-AIR,' an apparatus was developed which enabled the morphology of the ice-water interface to be examined during the transition. The slightly curved interface was viewed at a glancing angle with a Bausch and Lomb X30 stereomicro-scope. The interface was obliquely illuminated by a polarized light source. It was thus possible to differentiate between small projections or depressions at the interface, both by a shadowing effect and by the way in which they distorted the colored interference fringes across a particular crystal surface. Many attempts were made to photograph the various stages of interface breakdown. Unfortunately, the limited depth of focus of the microscope rendered the system unsuitable for photographic recording. Visual examination, however, was facilitated by continuous adjustment of the microscope focus control. The following observations of the interface transition were recorded for dilute solutions of ammonium fluoride in water. They are compared with iller's postulated morphologies shown in Fig. 1. i) The initial departure from planarity was characterized by the appearance of a large number of small depressions or "pores" in the interface. ii) The pores were arranged in wavy rows all lying roughly in the same direction within a particular grain. iii) The pores did not appear to form near to grain boundaries. iv) The pores appeared in certain grains before others. v) The transition from the pore structure to the elongated cell structure, Fig. l(d), was difficult to follow. It occurred quite rapidly and no truly irregular cell structure, Fig. l(c), was noticed. vi) The elongated cells gradually became more uniform in size and shape, with straight cell boundaries. They remained elongated, however, and no regular hexagonal cells were observed, cf. Fig. l(e). Several conclusions may be drawn from these observations. Firstly, the appearance of a poxed or, more accurately, a pore structure can be used as a criterion for the onset of constitutional supercooling. Care should, of course, be taken to distinguish between the true pore structure and artifacts caused by splashed-back liquid, particles of oxide, or inefficient decantation of liquid retained between cell boundaries, as described by Spittle et al? Secondly, and contrary to the suggestions of Tiller el l., it seems that interfacial perturbations are characterized not by curved projections but by small depressions or pores. A recent independent investigation by Biloni, Bolling, and cole5 reported in this issue indicates that the depressions might be explained in terms of dislocations intersecting the solid/liquid interface. It is possible that the dislocation concentration is reduced in the vicinity of grain boundaries. This would then explain the grain boundary depleted areas noted in (iii) above. This work was partially supported by the Air Force Office of Scientific Research, Contract No. AF 49-(638)-1029.

Jan 1, 1967

By O. W. Simmons, L. W. Eastwood, C. M. Craighead

Eighty-four quaternary titanium-base alloys from the following systems were investigated: 1. Titanium-chromium-carbon-nitrogen. 2. Titanium-chromium-carbon with copper, vanadium, molybdenum, manganese, iron, or nickel. 3. Titanium-chromium-nitrogen with vanadium, molybdenum, or nickel. The tensile properties, minimum bend radii, hardnesses, response to heat treatment and aging treatment, and phase relationships of these alloys are reported. HIS paper is the third in a series describing J- titanium-base alloys. The first paper describingdescribed titanium binary alloys, and the second contained data on a large number of titanium ternary alloys. Eighty-four heats, comprising three groups of quaternary titanium-base alloys, are described in this paper as follows: 1. Titanium-chromium-carbon-nitrogen alloys. 2. Titanium-chromium-carbon alloys containing copper, vanadium, molybdenum, manganese, iron, or nickel. 3. Titanium-chromium-nitrogen alloys containing vanadium, molybdenum, or nickel. The methods of preparing, fabricating, and testing the alloys are described in the paper Binary Titanium Alloys. Quaternary Titanium - Chromiurn - Carbon -Nitrogen Alloys The properties of titanium-chromium-nitrogen and titanium-chromium-carbon ternary alloys have been described in the paper Ternary Alloys of Titanium. Because small amounts of nitrogen and carbon both appear to be somewhat beneficial, these additions have been investigated fairly completely in quaternary alloys containing chromium at the following levels of composition: 2.5, 3.5, and 5.0 pct chromium 0.25 and 0.5 pct carbon 0.1 and 0.2 pct nitrogen Mechanical Properties: The mechanical properties of these alloys are listed in table I. Fig. 1 and 2 show graphically a comparison of the tensile properties and hardnesses of quaternary alloys with those of the titanium-chromium binary alloys. Fig. 1 shows the comparison at 0.25 pct carbon, whereas fig. 2 shows the comparison at 0.5 pct carbon. It is concluded from these data that 0.25 pct carbon and 0.1 pct nitrogen are desirable additions to the 5 pct chromium-titanium base. It is further concluded that most of this improvement is produced by the nitrogen and that carbon is either harmless or mildly beneficial, at least up to 0.25 pct. Structure and Transformation Range: The effects of carbon and nitrogen on the transformation range of the titanium-chromium alloys were determined as part of the routine investigation. Additions of 0.25 pct carbon and 0.1 pct nitrogen had no appreciable effect upon the transformation range. However, the addition of 0.25 pct carbon and 0.2 pct nitrogen definitely raised the beta solvus temperature, as shown by fig. 3.

Jan 1, 1951

By F. P. Haver, M. M. Wong

When a mixture of chalcopyrite concentrate and lime is heated in air, the following reaction takes place: 2CuFeS2 + 4Ca (OH)2 + 8 ½ 02 ; 2Cu0 + Fe203 + 4CaS04 + 4H20. Ca (OH)2 is the only common reagent that will adequately control SO2 evolution in roasting. Other additives tried, including CaO, CaCO3, Na2CO3, NaOH, MgO, Mg (OH)2, and MgCO3, were not effective. All reacted with sulfur, but not to the extent that will be required to meet the new control regulations for sulfur oxide emissions. For example: when the concentrate was heated with the theoretical amount of limestone (CaCO3), about 55% of the sulfur was retained in the calcine; whereas use of the stoichiometric amount of Ca (OH), converted approximately 99% of the sulfur to CaSO1. In most of the experiments, Ca(OH)2 was formed by adding CaO to concentrates containing enough water to simulate filter cake from flotation circuit. This eliminated the need for drying the concentrate and utilized the heat of hydration of CaO.

Jan 6, 1972

By Maurice L. Pinel, Thomas T. Read, Thomas A. Wright

THE erroneous, but until recently widely prevalent, belief that iron castings were first made in Europe in the fourteenth century has been adequately refuted in a number of earlier papers;1,11,12 but except for an unpublished metallographic study by the late William Campbell, and a phosphorus analysis by T. A. Wright of the metal of the Han dynasty cast-iron stove described by Laufer,2 nothing was known of either the chemical composition or the metallographic structure of the metal of any ancient iron casting, although data on castings of recent manufacture are abundantly available. The metal of the stove, which had been buried in a grave for at least 15 centuries, was so corroded that it seemed inadvisable to publish the results on it until more and better specimens from other ancient castings could be obtained. During a journey through China, extending from Peking in the north to Canton in the south and Cheng-tu, Ssu-chuan, in the west, in the late summer and autumn of 1936, one of us (T. T. Read) was fortunate enough to obtain, and bring home for study, nine castings, all more than 1000 years old, of which the date of manufacture is precisely known because of inscriptions cast in them, another that lacks a dated inscription, but of which the date, as explained below, can be otherwise established, and a specimen from the largest iron casting ever made," which is also dated. The period covered by this suite of specimens extends from 502 to 1093 A.D. In addition, a sample was obtained from a casting of date unknown, but which was certainly cast before 1000 A.D. and possibly as early as the third century A.D. A detailed description of these specimens is as follows: 1. 502 A.D. A pair (right and left) of recumbent lions (Fig. 1). Base 7 ¼ by 14 3/8 in., height over-all 5 ½ in., weight 26 lb. each. The translation of the inscription on the base is: "Made on the twenty-fourth day of the seventh month of the third year of Ching Ming of Great Wei" (Sept. 11, 502 A.D.).

Jan 1, 1938

By J. H. Hamilton, J. R. Lewis, J. H. Dismant, W. M. Fassell

RECENTLY this laboratory undertook the task of preparing some low carbon metallic cobalt by the carbon reduction of precipitated cobalt oxides. The oxides came in two lots and had the compositions given in Table I. Industrially, cobalt rondels' are made by the carbon reduction of cobalt oxide in small reverberatory furnaces at temperatures near but not above 1050 °C. This is a slow and tedious small batch process but does give cobalt metal which is low in carbon (1 pct C). Inasmuch as these cobalt oxides were relatively low in copper and iron it was decided to reduce them with petroleum carbon in a small three-phase electric arc furnace of about 3 to 5 Ib of metal capacity. It was hoped such a procedure would save considerable time and perhaps also indicate an improved commercially feasible process. The furnace was preheated by using small pieces of carbon electrodes to maintain the arc until a temperature of 1500°C was reached. The broken pieces of electrodes were then removed and 500 g of the cobalt oxides and 60 to 100 g of petroleum coke added. At the end of the reduction reaction, a high lime flux was added to the furnace to facilitate the removal of sulphur. Due to the fineness of the cobalt oxide powder used and also to the presence of chlorides there was considerable dust lost in charging the furnace. This was overcome by preparing briquettes composed of the cobalt oxides, charcoal, and a little bentonite to serve as a binder. The results obtained are summarized in Table 11, which shows that the cobalt metal produced in the arc furnace was high in carbon varying from 2.5 to 3.5 pct. This, of course, is much too high. For certain uses, such as in the preparation of permanent magnets, the carbon content of cobalt must be low. It was necessary therefore to find an effective method of reducing the carbon content of the cobalt that had been prepared in the arc furnace. Van Arkel3 recommends the use of Co2O3 as a means of removing carbon from cobalt. Experiments by the authors indicate that carbon is removed with difficulty by this method. Other methods used, including lancing of the molten high carbon cobalt with a jet of oxygen, proved to be ineffective. The method finally used makes use of the fact that cobalt metal is decarburized by carbon dioxide at the temperature of the molten metal. Two methods were used: 1—carbon dioxide was bubbled up through the molten cobalt, and 2— molten cobalt was maintained in a carbon dioxide atmosphere. In the first method the gas was introduced below the surface of the molten metal by means of a quartz tube. A high frequency induction furnace was used to keep the cobalt molten in both of these experiments. This kept the metal somewhat agitated and no doubt was an important factor in hastening the decarburization 'process by the second method. Both methods gave satisfactory results as shown in Table 111. An interesting observation in connection with this 'work was that a small sample of carburized cobalt metal could be decarburized by heating it with an acetylene torch adjusted to give an oxidizing flame. This may have been due to the presence of carbon dioxide in the products of combustion. Samples of molten high carbon steel were decarburized by the same method and it is suggested that nickel and perhaps manganese can be decarburized in a similar manner. Additional work, including rate studies and other theoretical aspects of this problem, is in progress in this laboratory.

Jan 1, 1953

By Kehsin Kuo

DIE steel with 1.47 pct C, 0.42 pct Mn, and 8.22 pct W contains Fe,C and WC in the annealed state and WC embedded in a martensitic matrix in the hardened state (quenched from 800°C). The presence of a large amount of undissolved angular WC carbide particles makes this steel extremely wear-resistant. However, the WC carbide begins to dissolve in the austenite when the heating temperature is raised above 1000°C, but this solution process is still not complete at 1250°C although partial melting has already taken place at the austenite grain boundaries. At 1325 C, this steel contains austenite and liquid metal but no WC carbide, Fig. 1. On quenching into brine, the liquid metal solidifies into a constituent which etches differently from the austenite-martensite matrix. This constituent appears lighter than the matrix at the 1250°C stage when etched with 2 pct nital but darker in Fig. 1 when etched with Murakami's reagent. Such a constituent has also been found in burnt tungsten high speed steels. This constituent has been isolated from the matrix by anodic dissolution in a 5 pct HC1 electrolyte and was found by the X-ray powder method to be iso-morphous with &manganese, Table I. The agreement in line intensity between these two phases indicates that there is no ordering in this new phase. The calculated cube edge of this phase is 6.395A, somewhat larger than that of ß-manganese (6.30A). Besides this phase, the anodic residue also contains free carbon from the matrix and amorphous WO,, although the X-ray powder pattern of the residue shows the presence of only this phase. This makes it impossible to determine directly the composition of this phase from the residue. Attempts to prepare an alloy containing only this phase have hitherto been fruitless. The following experiments serve to show that this phase can only be obtained by drastic quenching from the liquid state and that its tungsten and carbon contents must be lower than those in the Fe,W,C carbide. On air cooling from 1325oC, the liquid metal solidifies into a eutectic consisting of austenite and the Fe3W3C carbide. The nature of this dark-etching Fe3W3C has also been ascertained by X-rays. This eutectic becomes coarser on furnace cooling. The solidification of a specimen cooled for 5 min from 1325" to 1200°C and water quenched is still not complete and the fine eutectic is formed during the quenching. Since this new phase can be replaced by a eutectic of austenite and the Fe,W,C carbide, its composition must lie somewhere between the latter two phases. The special features of this phase are: 1—it can only form directly from the liquid state by drastic quenching and 2—it is a ternary phase of iron, tungsten, and carbon. The ß-manganese structure of this phase is also interesting. In the Periodic Table, iron and tungsten lie on opposite sides of manganese; a certain combination of these two elements, in the presence of carbon, may have average properties similar to those of manganese. It has already been established that the x phase, a ternary one of iron, chromium, and molybdenum, has the a-manganese structure."' A similar phase has also been found in the ternary system Fe-Cr-W and in the following binary systems: Re-Cb, Re-Ta, Re-Mo, and Re-W.6 References 1 G. D. Preston: Philosovhical Magazine 11928) 5. 7th Series, p. 1207. - K. W. Andres: Nature (1949) 164, p. 1015. 35. G. McMullin, S. F. Reiter, and D. G. Eberling: Trans. ASM 11954) 46, p. 799. ' J. S. Kasuer: Acta Metallurgica (19541 2, p. 456. • H. J. Goldschmidt: Trans. ASM (19541 46, p. 807. P. Greenfield and P. A. Beck: Private communication.

Jan 1, 1957

By R. J. Sherman, M. R. Achter

Creep rates have been measured for nickel at 1200°F in air and in vacuum, and related to the depth of surface cracks in the specimens tested in the two environments. The surface cracks were observed to grow faster in air than in vacuum in both short-time and long-time tests. Crack depths, however, could be correlated with creep rates only in the short-time tests. Although no cracks were observed in a relatively brittle Ni-Cr-A1 alloy creep tested at 1300°F, the appearance of alloy specimens fractured after interruption in creep indicated that there was a decrease of intergranular cohesion resulting from the diffusion of atmospheric gases into grain boundaries. PREVIOUS investigations1-3 have demonstrated that the effect of atmosphere on the high-temperature strength of nickel and nickel-base alloys may be reversed by changes in the stress or temperature of the test. At high stress and low temperature the creep-rupture life is longer in vacuum than in air, while at high temperature and low stress the relative strengths are reversed. A mechanism to explain this behavior was proposed based on two competing processes. Oxidation hardens and strengthens a metal, but adsorption of gas, by lowering surface energy, reduces the work required to propagate an intergranular crack. The controlling process is determined by temperature and stress. Basic to this mechanism is the concept that environment determines the ease of propagation of intergranular cracks and thus affects creep rate and rupture life. To explore the mechanism in greater detail, this investigation was undertaken to determine the effect of atmosphere on crack initiation and propagation rates in nickel1 and a relatively brittle Ni-Cr-A1 alloy.3 EXPERIMENTAL PROCEDURE The materials were the same as used previously. Nickel of 99.8 pct purity had the following impurity content: 0.02 pct Co, 0.03 pct Fe, 0.02 pct Mg, 0.03 pct Si, and 0.008 pct C. Rods were cold swaged from 1/2 in. to 3/8 in. in diam and machined into speci- mens with gage sections 1/4 in. in diam by 1-1/4 in. long. They were annealed in vacuum at 1600" F for 1 hr to achieve a grain size of 0.07 mm. The alloy was composed of 19.1 pct Cr, 3.6 pct Al, 1.7 pct Si, 0.57 pct Fe, 0.10 pct Mg, 0.027 pct C, 0.004 pct S, and the balance Ni. Sheet specimens of 0.500-in. width and 1.25-in. gage length were machined from 0.033-in.-thick strip and annealed in vacuum at 2000° F for 2 hr, resulting in a grain size of 0.036 mm. Large amounts of an intergranular phase were introduced as a result of the anneal. Creep tests on both materials, using equipment previously described by Shahinian; were conducted in air and in a vacuum of 1 x 10-5 mm Hg. They were continued to rupture and interrupted at various stages of the test. Although extensions were measured from dial gage readings of movement of the lever arm, it is considered that valid comparisons of creep rate may be made in this manner.l Two parameters were measured metallograph-ically* on longitudinal sections of nickel specimens interrupted during creep—the mean depth of surface cracks 2, and the density n as the number of cracks per unit length of surface. They were combined into one parameter by the relation p = n-d [1] As used here, p is equal to the total of crack depths per unit length of surface. RESULTS Nickel. It is well known that the incidence of intergranular cracking is increased with increasing rupture life. Indeed, it is evident in Tables I and I1

Jan 1, 1962

By H. H. Klepfer, M. E. Snyder

UNTIL very recently, the commercial availability of the rare earths as metals has been very limited. Fabrication of mill products from these metals has not been studied in most cases. This note reports the results of the development of fabrication techniques for thulium. Thulium has a melting point of about 1550°C and a hexagonal close-packed crystal structure. It oxidizes in air to give a black oxide (TmzO,). The procedures for producing thin thulium foil were developed on one ingot weighing about 220 g and were subsequently applied in processing about a pound of metal to foil. Several alternates for the various fabrication steps were investigated and will be discussed. The metal fabricated was in the form of commercial chill-cast ingots 1 in. in diam and 2 in. long weighing approximately 220 g. Impurities in the ingots were reported by the vendor to be 4000 ppni tantalum, 2000 ppm calcium, 200 ppm nickel, 100 to 200 ppm iron, 100 ppm europium, and less than 100 ppm copper, lutetium, and ytterbium. In addition to these impurities, several salt-like inclusions as large as l/8 in. in diam were revealed along the center line of the one ingot sectioned. Preliminary tests indicated that small wafers cut from the as-cast ingot would not fabricate readily by rolling. Forging of copper-jacketed wafers was therefore attempted. At 1550"F forging was satisfactory but an apparent reaction of copper with thulium demanded investigation of lower temperatures. Therefore, the remainder of the test ingot was forged at 1450°F—with only minor cracking. All ingots forged were inserted into copper tubes of 1 in. 1D and 0.125-in. wall thickness. The jackets were sealed by flattening the ends of the tubes and welding under helium. Heating time at 1450°F was 30 min. Press forgings of 1/8 in. per pass were used, followed by 10 min reheats. When the ingots had been squared and reduced to 0.250 in. in thickness, the original copper jacket was stripped off and replaced by a new jacket in preparation for hot rolling. Hot rolling at 1450" F without edge cracking was readily accomplished after forging. Excellent results were obtained with 10 pct reductions of thickness followed by 5 to 10 min reheats. After reduction of thickness from 0.250 to 0.100 in. the copper jacket was removed. It was found, in fact, that hot rolling in air was possible. A tenacious black oxide similar to that seen on zirconium was formed during 3 min reheats at 1450°F. Reduction in air to 0.010 in. foil was possible taking 10 pct reductions per pass. The oxide coat formed during hot rolling in air could best be removed by sand blasting and pickling. Common pickling solutions containing polar solvents were found to attack the metal too rapidly and a concentrated nitric-hydrofluoric acid mixture attacked neither the oxide nor the metal. The most satisfactory pickling solution was 52 vol pct concentrated nitric acid-48 vol pct glacial acetic acid. After forging to 0.250 in., vacuum annealing and cold rolling was found to be another satisfactory alternate to hot rolling in copper jackets. After forging. a hardness of Rockwell B76 was found. Annealing in vacuum (2 X l0-5 mm of Hg) for 1 hr at 1200"F did not alter this value. Annealing at 1470"F for 1 hr brought the hardness down to R;]63. With cold rolling (5 pct per pass) the hardness returned to about R,1:76 after 20 pct reduction and edge cracking became noticeable. However, cold rolling to a total of 40 pct reduction in thickness (RI,,83) was possible before edge crack propagation became serious. Good surface finish was obtained, and the metal loss due to oxidation was minimized by cold rolling and vacuum annealing. Using this procedure the yield of 1.25 in. wide by 0.010 in. thick foil from a 1-in. diam ingot was about 40 pct.

Jan 1, 1961