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By Charles Hawes

WHEN sulfur is encountered in objectionable amounts, it is regarded as the most trouble-some element for the mine operator to control. It exists in two conditions in iron ores, as sulfide in iron pyrites and as sulfate in gypsum and barite. The sulfate radical (SO,) may also be linked to one or more of the several other elements found in the iron ore compound. The only dependable method for the de-termination of sulfur, particularly when found in the insoluble residue, has been the so-called fusion method. This is a long and tedious process involving several hours for its completion. In looking for a more rapid method which would be .applicable to all iron ores, the writer investigated the work of M. Groger1 who in 1881 showed that by heating pyrite with iron out of contact with air, the pyrite is reduced to iron sulfide and from the latter all of the sulfur will be given off as hydrogen sulfide gas on treat-ment with hydrochloric acid. FeS2 + Fe = 2 FeS FeS + 2 HCl = FeC12 + H2S In 1891 F. P. Treadwell2 not only confirmed the results of M. Groger but also found that iron powder would reduce to sulfide the sulfate radical of any substance, including barite, if not found in excessive quantities. The writer found this to be true and developed a method, in cooperation with F. J. Baker, chief chemist, mining department, of the Cleveland-Cliffs Iron Co., Ishpeming, Mich., for the determination of sulfur in iron ores which is not only very rapid but also accurate. Pure iron ore powder, reduced by hydrogen, gave good results but was found to be very slow in reacting. A high silicon iron with a low sulfur content of 0.012 per cent was found to be. much more active than the pure iron and gave excellent results.

Jan 11, 1927

By P. Lelong, P. Lacombe, J. Herenguel

IN previous studies on the effect of orientation on the rate of anodic oxidation of A1-3 pct Mg mono-crystals of high purity, certain anomalies were observed on the cold-worked metal. These anomalies were shown to correspond to regions of changed orientation within the grains resulting from plastic deformation.' In fact, they are deformation bands revealed by differences in rate of oxidation and produced by more or less continuous bending of the lattice about a [Ill] axis common to the original crystal and the band.2,3 Preliminary observations made after relatively large degrees of cold working revealed an approximate twin relationship between the bands and the surrounding crystal. Fig. 1 shows the etch figures formed on such a band and on the original monocrystal after a 67 pct rolling reduction. It is evidently not a true twinning, since there is a broad transition zone between the two systems, however the two extreme orientations are nearly in a twin relationship. It was decided to investigate these preliminary observations by progressive rolling-reductions on monocrystals of varied initial orientation. From these studies, which form the subject of the present article, it became apparent that the initial observation of an approximate twin relation between the bands and the monocrystal was simply a special case. Our experiments show that deformation in rolling proceeds by a mechanism analogous to those which have already been observed in the simpler cases of tension and compression. Micrography—in particular, the formation of etch figures—has been applied to the quantitative study of these orientations. In addition, observations have been made of slip lines, of thin oxide films yielding interference colors, of thick films in polarized light, and of etchants giving "engraving." These micrographic methods for studying orientation have established the orientation relationship without any ambiguity, whereas X-ray methods would have given only a statistical result because we were often dealing with thin bands (2 to 100 microns) and even narrower connecting zones covering a maximum of 20 microns. Specimen Preparation and Methods The Al-3 pct Mg alloy was prepared from 99.95 pct A1 and is of the type used industrially in France for brilliant anodized finishes under the name of Brillalumag. This material combines the normal commercial properties with an enduring brilliance due to the protection of a thick, but perfectly transparent, film of alumina.' The monocrystals were prepared by the method of critical deformation followed by prolonged annealing. Great care was taken to insure perfect homogeneity of composition, to eliminate included crystals,6,7 and to avoid the formation of undulations during oxidation. The critical deformation for this alloy is about 1.2 pct (1-l0/10x100) and the subsequent annealing was for one week at 500 °C after gradual increase of temperature (35°C per hr). The crystals thus prepared had a section of 10x5 mm and a length of 20 to 40 mm. Each of these crystals was polished electrolytically and etched for the formation of etch-figures with the classical reagent.' The orientation was then calculated from their angles, as measured with a goniometer microscope.

Jan 1, 1954

YEAR OF ELECTION 1913. FRANK DAWSON ADAMS, Montreal, Canada. 1938. HENRY CORT HAROLD CARPENTER, London England 1933. KARL EIKERS New York, N. Y. 1922. FEDERICO GIOLITTI Torino, Italy. 1906. SIR ROBERT A. HADFIELD London, England. 1921. FRANK WILLIAM HARBORD London, England. 1917. HERBERT HOOVER Washington, D. C. 1931. WALDEMAR LINDGREN Cambridge, Mass. 1921. C. McDERMID London, England. 1913. EZEQUIEL ORDONEZ Mexico City, Mexico. 1935. ROBERT PEELE New York, N. Y. 1938. A.M. PORTEVIN Paris France 1937. JOHN V.W. REYNDERS, New York, N. Y. 1937. GEORGE S. RICE Washington, D. C. 1911. ROBERT H. RICHARDS Cambridge, Mass. 1934. T.A. RICKARD Victoria, B.C., Canada

Jan 1, 1939

By A. D. Mutch

This contribution is prompted by the fact that the writer has recently published an article" which has in part the same general conclusions as B. S. Butler's, namely, the form of the standard mineral paragenesis in hypogene deposits and the parallel transition from early high temperature acid to late low temperature alkaline or neutral solutions. Since the writer also recognizes that he is not fully qualified to discuss the formation of ore deposits as a chemical problem, this discussion will be limited to summarizing other observations and will suggest how Butler's Table II might be made more effective. The form of Butler's table is not unlike that widely used by European geologists30 in which the various minerals are related to the types of ore fluids-pegmatitic, pneumatolic and hydrothermal. Thus the decrease in temperature and change in pH of the ore- bearing fluids as recognized by Butler may also be paralleled by perhaps even greater changes in composition, density, and character.

Jan 5, 1958

By Gordon G. Bentle, W. P. Fishel

GREINER, Marsh, and Stoughton1 have reviewed the literature in a monograph on the iron-silicon system. The lack of agreement among the various studies may be due to the difference in the purity of materials in some cases. The results of the X-ray method indicate only a general rather than a specific limit of the y loop. The iron-titanium system was investigated by Roe and Fishel,' whose results are used in determining the iron-silicon-titanium y loop. The literature indicates that no study has been made thus far of the iron-silicon-titanium system in the y loop region. Experimental Procedure The alloys were made from Armco iron, 99.86 pct Si, and 98.5 pct Ti. The titanium contains 1.07 pct Fe. The Armco iron contains 0.01 pct C, 0.005 pct Si, 0.02 pct Mn, 0.006 pct P, 0.01 pct S, and 0.04 pct Cu. A dilatation sample of Armco iron was prepared in the same manner as were the other samples. Several dilatation runs indicated that Ac3 begins at 905°C and ends at 910°C, which compares favorably with the study of Wells, Ackley, and Mehl' employing specially purified iron. All alloys were melted in an induction furnace in Alundum crucibles and were top poured into baked core sand molds. The ingots weighed about 600 g as cast. Dilatation samples of 4 in. by % in. diam were taken from the bottom end of the ingot, from which turnings for analysis were taken. Turnings from each end of two of the samples indicated that the alloys possessed uniform silicon content. A hole was drilled halfway through each end of a 4 in, dilatation sample. A chromel-alumel thermocouple was inserted in each hole and the surrounding metal peened in to hold the thermocouple in place. Fig. 1 represents the dilatation apparatus. The heating unit is an Alundum tube wound with resistance wire; the dilatation tube is a McDanel combustion tube; the push rod is zircon; the dial gage is graduated in 0.0001 in. and was vibrated to prevent sticking. A blank run over the temperature range of interest gave no significant dilatation. The thermocouples were attached to a multiple recording micromax which indicated the temperature at each end of the sample at 35 sec intervals. A Lindberg control unit governed the rate of heating of the furnace, and lowering or raising the

Jan 1, 1957

By J. E. Cannaday, R. J. Austin, R. K. Saxer

PREVIOUS investigators have determined activation energies and have postulated various controlling mechanisms for creep.'-' Recently Barrett et 01.I5 have suggested, as the result of their work with cadmium, that the apparent activation energies for creep (AHc) above 0.5 T, could be correlated by means of activation energies for self-diffusion (AHSD) and modulus of elasticity functions: AHC = AHSD - (5RT/E) dE/dT. The wide disparity of apparent activation energies previously obtained with nickel compared to those reported for other fcc metals and the desire to test the postulate of Barrett el al. prompted this study. The as-received 17.5-mil nickel wire used during this work assayed 99.87 pct pure, Table I. X-ray analysis indicated that the wire was in a stress-re- lieved state and metallographic examination revealed an equiaxed grain size of approximately 0.01 mm. Specimens degreased in acetone and cut to 1-1/2 in. lengths were tested in a cam-mechanized constant-stress creep machine. A vacuum of at least lo-5 mm Hg was maintained during all tests. Stresses were applied when thermal equilibrium was established. When a measurable (at least 0.5 pct) secondary creep was observed, the temperature was pulsed to a higher temperature. Pulsing time varied from 30 sec at 480°C to 2 min at 1300'C and above. Slopes of creep curves were determined after the initial acceleration in the strain rate had subsided and steady-state creep was achieved. To confirm the results obtained, additional tests were performed incorporating a temperature pulsing method suggested by sherby. In this method the load was removed during temperature pulsing and reapplied after thermal expansion had occurred. No appreciable differences in results were obtained. Tests were performed over a wide range of temperature (470" to 1310°C) and stresses (250 to 17,000 psi). The expression was used to calculate activation energies. In this expression T1 and T2 are the initial and final temperatures, and el and i2 are the creep rates at T, and T2. The variation of the apparent activation energies with the mean test temperature calculated by means of Eq. [I] are shown in Fig. 1. The mean test temperature (Tmean) was calculated from the relationship

Jan 1, 1967

YEAR or ELECTION 1913. FRANK DAWSON ADAMS Montreal, Canada. 1933. KARL EILERS New York, N. Y. 1922. FEDERICO GIOLITTI Q. Torino,Italy. 1906. SIR ROBERT A. HADFIELD London, England. 1921. FRANK WILLIAM HARBORD London, England. 1917. HERBERT HOOVER New York, N. Y. 1931. WALDEMAR LINDGREN Cambridge, Mass. 1921. C. MCDERMID London, England. 1913. EZEQUIEL ORDOAEZ Mexico City, Mexico 1935. ROBERT PEELE New York, N. Y. 1938. A. M. PORTEVIN Paris, France. 1937. Joan V. W. REYNDERS New York, N. Y. 1937. GEORGE S. RICE Washington, D. C. 1911. ROBERT H. RICHARDS Cambridge, Mass. 1934. T. A. RICKARD Victoria, B.C., Canada. 1933 Cambridge, Mass.

Jan 1, 1938

By A. D. LeClaire

This paper claims to provide a description of diffusion processes in solids which is said to be a reformulation of Fick's Laws and therefore, presumably, is supposed to be as general as those laws. In fact, it appears to provide no more than an atomic model of a rather restricted type leading even to results in conflict with experiment. The restriction lies in the assumption, carried throughout the paper, that even in a chemical gradient the jump frequency i of an atom of species i is a function only of the composition on the plane from which it jumps, i.e., ?i = ?i(cl, c2 ...). This means that atom jumps in opposite directions from a site would occur with equal probability. There would then be no tendency for atoms to migrate preferentially in one direction along the gradient, so the model is incapable of accommodating the well established phenomenon of atom drift in a chemical gradient.10-" A more general model that does not suffer this difficulty was discussed in a paper13 I wrote in 1958, and the results therein already contain as a special case those given by Dr. DeHoff. The simplest sufficiently general assumption is to suppose that ?i is a function ?i(a1 a2,...; wl, w2, ..-) of the concentration aj and ? ?j at the site from which and to which, respectively, a jump is made. Any such function can always be expressed as a function of the sum and difference of its arguments. Thus d is the jump distance and g a geometrical constant. Obviously this is zero, i.e., no drift, if ??ri/??j = 0. X is written in this form, revealing two components of the total drift, because it can be shown that only the second component, depending on ??'i/???(?j -aj), actually contributes to the net diffusion flux Ji. This component can be shown,13 by a straightforward ther-modynamic transcription, to be responsible for the so called "thermodynamic factor" in the Darken approximation for a chemical diffusion coefficient is the activity coefficient of i. Because the jump rate ?i in DeHoff's approximation has no dependence on the concentration gradient it does not seem possible to derive the Darken equation from his Eq. [6] without the further rather arbitrary assumption that ?i = const ri. It is to be noted too that the author's derivation of Eq. [12] is only valid when ?i is in fact a function only of the cj at the plane from which the atom jumps—for only then is ci?i a unique function of position and therefore of A. Author's Reply R. T. DeHoff I am indebted to Dr. LeClaire for his penetrating discussion of my paper, "A Reformulation of Fick's First Law". I concede that it does appear initially that the relationship presented in my paper could be derived from the detail atomic model of random walk and atom drift presented in Dr. LeClaire's approach. However, there are some important differences in concept and viewpoint which deserve further amplification. Dr. LeClaire's description is penetrating, detailed, and fundamental. His approach describes an extremely complex phenomenon, the displacement of atoms in a multicomponent gradient, at a sophisticated level. For precisely these reasons, his analysis has not proved to be useful for experimental studies that have practical objectives. For example, in order to obtain the information required to complete the analysis of the jump frequency distribution in a ternary system, the following series of experiments would, I believe, be required: 1) chemical (gradient) experiments, covering the composition range;

Jan 1, 1970

By P. Gordon

An isothermal jacket microcalorimeter, supplemented by metallographic, microhardness, and X-ray measurements has been used to study the isothermal annealing of high purity copper after room temperature tensile deformation. The amount of stored energy released during annealing has been measured as a function of deformation in the range 10.8 to 39.5 pct elongation. The data have shown the major heat effect to be associated with recrystallization and have allowed an analysis of the recrystal-lization kinetics and the calculation of activation energies of recrystallization. WHEN a metal is deformed plastically, some of the energy expended is dissipated as heat during the working process, while the remainder is stored within the metal in the form of lattice distortions and imperfections. During subsequent heating of the metal, the distortions and imperfections can be largely annealed out and the associated stored energy released as heat. It is apparent that measurements of the evolution of stored energy during such annealing may produce important information concerning the nature of the annealing mechanisms and the imperfections involved. Some excellent studies of this type have been made in the past, notably those of Taylor and Quinney,' Suzuki,2 Bever and Ticknor,3 Borelius, Berglund, and Sjöberg,4 and Clarebrough et al.5,6 None of this work, however, employed isothermal techniques, with the exception of the Borelius studies' in which only the early annealing stages were investigated. Since isothermal measurements, as compared with heating or cooling curve, have the merits that 1—they reveal the kinetics of a process more clearly, 2—the results obtained are more easily applied to theory, and 3—most fundamental investigations of annealing using techniques other than calorimetry have been carried out isothermally, it was considered important to apply calorimetry to the study of the isothermal annealing of metals. Accordingly, an isothermal jacket calorimeter of the Borelius type,' supplemented by metallographic, hardness, and X-ray measurements, has been used to study the annealing of high purity copper after room temperature tensile deformation. Experimental The microcalorimeter has been described fully elsewhere." Briefly, the specimen to be studied is placed in a constant temperature environment of virtually infinite heat capacity achieved, as shown in the drawing of Fig. 1, by means of a vapor thermostat. A high thermal resistance is provided between the sample and the environment and a sensitive differential thermopile (see Figs. 2 and 3) arranged with half its junctions in contact with, and thus at the constant temperature of, the environment, and the other half in contact with the sample. A reaction in the sample develops a small difference in temperature, AT, across the thermopile, which is followed by a recorder-galvanometer set-up as a function of time, t, and is converted to reaction heat per unit time, P, by the use of the equation AT P=a?T + b AT dt The constants, a and b, in Eq. 1 are determined by a simple calibration, making use of the Peltier heat developed by a small current run through the junction of a thermocouple located in an axial hole in the specimen (Fig. 2). In its present form, the limit of sensitivity of the calorimeter is a heat flow of 0.003 cal per hr. The copper used was the spectroscopically pure metal supplied by the American Smelting and Refining Co. in the form of 3/8 in. diam continuously cast rod, reported to be 99.999+ pct Cu. A small amount of the copper was available at the start of this work and is referred to hereafter as lot A. A second batch, lot B, was obtained later, most of the results described subsequently being for this lot. As will be seen, there is some indication that lot A was somewhat purer than lot B, but it is not known whether this difference was present in the as-received metal or arose during subsequent handling. The two lots of copper were remelted and cast into two 1½ in. diam ingots in vacuo, using high purity graphite crucibles and molds. The ingots were upset several times to break up the large cast grains, and then rolled and swaged to rods 0.391 in. in diameter, using several intermediate anneals with about 40 pct reduction in area between anneals. The penultimate anneal was 2 hr at 350°C. X-ray examination showed no marked general preferred orientation in the resulting rods. The grain structure typical of the two rods is shown in the micrograph of Fig. 4." It was found to be virtually im- possible to get an unambiguous measure of the absolute grain size in the two annealed rods because of the profusion of annealing twins and the lack of regularity of the grain boundaries. However, counts of the number of boundaries intersected per unit length along a random line on a polished section, making a correction for the proportion of boundaries (about half) estimated to be twin boundaries, gave a figure of about 0.015 mm for the average grain diameter. The grain size of the rod from lot A was about 5 pct smaller than that from lot B. The rods were cut into 1 ft long bars and these deformed in tension at room temperature to various total elongations in the range 10.8 to 39.5 pct. A strain rate of 1 pct per min was used. The deformed bars were then stored in a dry ice chest until such time as samples were to be cut from them. Five bars deformed as indicated in Table I were used for the subsequent tests. In all cases, all the calorimeter.

Jan 1, 1956

By Eric R. Jette, Earl S. Greiner

The constitution of the iron-rich iron-nickel-silicon alloys at 600ºC. has been determined by X-ray crystal structure methods. If the results are represented isothermally by a conventional equilateral triangular plot, the single-phase solid solution (body-centered cubic structure) of the iron-rich alloys at 600°C. is bounded by four two-phase areas. Three three-phase areas separate the respective two-phase areas. Three of the phases in the polyphase areas have the same structures as the corresponding phases in the binary iron-nickel or iron-silicon alloys. There are two phases, however, that have not been reported previously and whose structures have been tentatively identified in the present investigation. The solid solubility limit of the iron-rich solid solution at 600°C. was determined by comparing the lattice constants of this body-centered cubic phase in the single and polyphase areas. It has been shown that the lattice constants of ternary alloys may be represented by surfaces whose distances above the ternary plot (representing atomic composition) are proportional to the lattice constants of the respective alloys. Then the intersections of the surface for the single-phase alloys with those for the corresponding phase in the polyphase alloys, at the respective temperatures, are the solid solubility limits of the single-phase alloys at these temperatures. The data of the present investigation were evaluated by this method, although by a system of orthographic projection it was possible to represent the three-dimensional model in two-dimensional diagrams. This investigation has shown that the solid solubility of nickel in the body-centered cubic phase of the iron-rich iron-nickel-silicon alloys at 600°C. changes from 3.7 weight per cent (Owen and Sulley) in the binary iron-nickel alloys to 5.0 per cent in the ternary alloys containing 4.8 per cent silicon. Further additions of silicon increase the solid solubility limit of nickel to 16.0 per cent at 12.6 per cent silicon. From this point the limit decreases to 9.5 per cent nickel at 15.5 per cent silicon and is, of course, zero at 15.3 per cent silicon in the binary iron-silicon alloys.

Jan 1, 1943

By Earl S. Greiner, Eric R. Jette

The constitution of the iron-rich iron-nickel-silicon alloys at 600ºC. has been determined by X-ray crystal structure methods. If the results are represented isothermally by a conventional equilateral triangular plot, the single-phase solid solution (body-centered cubic structure) of the iron-rich alloys at 600°C. is bounded by four two-phase areas. Three three-phase areas separate the respective two-phase areas. Three of the phases in the polyphase areas have the same structures as the corresponding phases in the binary iron-nickel or iron-silicon alloys. There are two phases, however, that have not been reported previously and whose structures have been tentatively identified in the present investigation. The solid solubility limit of the iron-rich solid solution at 600°C. was determined by comparing the lattice constants of this body-centered cubic phase in the single and polyphase areas. It has been shown that the lattice constants of ternary alloys may be represented by surfaces whose distances above the ternary plot (representing atomic composition) are proportional to the lattice constants of the respective alloys. Then the intersections of the surface for the single-phase alloys with those for the corresponding phase in the polyphase alloys, at the respective temperatures, are the solid solubility limits of the single-phase alloys at these temperatures. The data of the present investigation were evaluated by this method, although by a system of orthographic projection it was possible to represent the three-dimensional model in two-dimensional diagrams. This investigation has shown that the solid solubility of nickel in the body-centered cubic phase of the iron-rich iron-nickel-silicon alloys at 600°C. changes from 3.7 weight per cent (Owen and Sulley) in the binary iron-nickel alloys to 5.0 per cent in the ternary alloys containing 4.8 per cent silicon. Further additions of silicon increase the solid solubility limit of nickel to 16.0 per cent at 12.6 per cent silicon. From this point the limit decreases to 9.5 per cent nickel at 15.5 per cent silicon and is, of course, zero at 15.3 per cent silicon in the binary iron-silicon alloys.

Jan 1, 1943

YEAR OF ELECTION 1913. FRANK DAWSON ADAMS Montreal, Canada. 1935. GUILLAUME DANIEL DELPRAT Melbourne, Vic., Aust. 1920. HENRY STURGIS DRINKER Merion Station, Pa. 1933. KARL EILERS New York, N. Y. 1922. FEDERICO GIOLITTI Trino,Italy. 1906. SIR ROBERT A. HADFIELD London, England. 1921. FRANK WILLIAM HARBORD London, England. 1917. HERBERT HOOVER New York, N. Y. 1931. WALDEMAR LINDGREN Cambridge, Mass. 1921. C. MCDERMID London, England. 1913. EZEQUIEL ORDOAEZ Mexico City, Mexico. 1935. ROBERT PEELE New York, N. Y. 1937. GEORGE S. RICE Washington, D. C. 1911. ROBERT H. RICHARDS Cambridge, Mass. 1934. T. A. RICKARD..... Victoria, B.C., Canada. 1933. ALBERT SAIVEUR CAMBRIDGE, MASS.

Jan 1, 1937

By Patricia J. Rand, Kailish Narayan

Atlantic Richfield Co. began operating a commercial uranium solution mining project in April 1975, in Live Oak County, Tex., 10 miles southwest of the town of George West. The operation is designed to produce 250,00 lb of U308 per year and is jointly owned by Atlantic Richfield Co., Dalco, and U.S. Steel. The proprietary process employed in this operation consists of injecting an alkaline leaching solution into the ore-bearing formation through a pattern of injection wells. The uranium-bearing solution is then removed through producing wells using submersible pumps in central production wells. The produced solution is pumped through a pipeline to the plant area where the uranium is extracted. The barren solution is refurbished and returned to the pattern area for reinjection. The injection field and plant site are located in the drainage of Spring Creek, a small intermittent stream tributary to the Nueces River. Land use in the area is primarily agricultural cattle ranching and farming---but several producing gas fields are present in the areas as well. A nine month enviromental baseline study was conducted on the property beginning in Feb. 1975. Plants and animals of the Clay West area were identified and representative terrestrial and aquatic communities were analyzed. Physical and chemical parameters of the soils were described and analyzed. The surface and underground waters were analyzed chemically; radionuclides, heavy metals, nutrients, and pesticide concentrations were measured in the creek sediments; bacteriological measurements were made in creek water and sediments; radionuclide and chemical levels in plants and animals, both terrestial and aquatic, were measured. These studies, though not required by existing, state or federal laws, were considered necessary to assess the existing conditions before operations began on site. The in-place leaching technique for uranium mining eliminates the enviromental disturbances associated with open pit or underground mining. No grinding and crushing of ore is needed, nor are large tonnages of waste tailings produced. Surface subsidence is not expected since less than 5 lb of uranium is present per ton of rock and even complete removal of such small amounts should not cause subsidence. Solution mining has fewer environmental problems than conventional mining. Members of the Texas Water Quality Board recognized the differences which exist and have recently formulated new environmental control procedures which apply specifically to solution mining. Some of the solution mining activities which are of particular concern are discussed in the following pages. Process Description The produced uranium-bearing solution is pumped through a pipeline from the pattern area to the plant area. The first pattern area consists of a nest of 125 fluid injection, production, and monitor wells [(Fig. 1)]. Injection wells have 4-in. casings and are spaced 50 ft apart on a square grid. Production wells have a 6-in. casing and are located in the center of each square. A block flow sheet of the plant is given in [Fig. 2]. The produced solution goes through parallel trains, each consisting of a carbon column and an ion exchange column in series. The carbon column removes any sand present in the produced solution and the ion exchange column extracts the uranium. The depleted solution is treated with chemicals and returned to the field for reinjection. The uranium retained in the ion exchange column is removed periodically with a sodium chloride solution. Uranium-bearing solution from the ion exchange column is passed through a charcoal column for removal of impurities and then to a precipitation tank where the uranium is precipitated as yellowcake slurry. The yellowcake slurry is clarified, filtered and dried to produce a granular yellowcake product. Environmental Considerations Sources of waste solutions to storage reservoirs, the backwash reservoir and atmospheric discharges are shown in the block flow sheet [(Fig. 2)]. These waste solutions pose a relatively small problem and are adequately controlled by conventional means. There is one area, however,

Jan 1, 1978

By B. S. Tani, R. V. Schablaske, M. G. Chasanov

MEASUREMENTS of the solubility of scandium in liquid cadmium from 349" to 606°C by Chasanov, Hunt, Johnson, and eder' led to the isolation and characterization of the intermetallic compound ScCQ. The solid phase found in equilibrium with a saturated solution of scandium in liquid cadmium above 422°C is ScCQ. The equilibrium solid phase below this temperature has not as yet been characterized. Hexagonal needles of ScCG were isolated from a 5 wt pct Sc alloy which had been ice-quenched from 500°C. Free cadmium was leached away from the intermetallic with saturated ammonium nitrate solution. The concentration of cadmium in successive leaches was followed polarographically and the leaching process was terminated when the rate of cadmium dissolution decreased markedly. Chemical analysis of the intermetallic thus isolated is as follows: 11.5 wt pct Sc and 88.3 wt pct Cd. The composition of ScCQ is 11.8 wt pct Sc and 88.2 wt pct Cd. X-ray diffraction powder photographs of ScCds taken in a 114.6-mm Debye-Scherrer camera with CuKa, radiation can be indexed on the basis of a hexagonal nit cell. Lattice coonstants are a. = 6.330 ± 0.001A to 6.324 i 0.001A and co = 4.853 i 0.001A. The a0 has a slight range with the lower value seemingly associated with the cadmium-rich terminus of the phase. With eight atoms in the unit cell the calculated intermetallic density is 7.54 g per cu cm; the density as determined by a liquid-displacement method is 7.7 g per cu cm. Powder-diffraction data for ScCQ are recorded in Table I. Stoichiometry, symmetry, unit-cell dimensions, and the number of atoms per unit cell indicate the MgsCd (D019) type structure for ScCQ; space group is D,&-P6,/mmc. Atomic coordinates are:

Jan 1, 1964

John Langton, New York, N. Y. (communication to the Secretary*):—In Mr. Firmstone's discussion (Trans., xli., 912) the formula he gives for the periodical payment—his equation (2)—may be simplified in form. The expression given by Mr. Firmstone is : q=S(l+ r) 1-(1+r)-(1+r) r which is a simpler form to use in calculations. Where S= principal to bc extinguished. Q = periodical payment or charge for sinking-fund, r = rate of interest per period (i. e., interest on $1 per period). n = number of payments or periods. Kent's Mechanical Enyineer's Pocket Book, under the heading Annuities, equation No. 5, gives : " The annuity which $1 will purchase for any number of ytlars n is * Received Jiily 16, 1911.

Jan 1, 1912

By H. H. Reamer, B. H. Sage

The formation volume and the relative volume of the liquid Phase of mixtures of Condensate and gas from five different parts of a field in the Louisiana area have been established experimentally. These studies included gas-condensate ratios as high as 25.000 cu ft/bbl at pressures up to 5,000 psi. The gravities of the tank condensates were between 53.9° and 63.9° API. The compositions of the condensate and gas samples inves-ligated are presented in tabular form. The formation volume. the relative volume of the liquid phase. and the specific volume of the mixtures are available. A rough correlation of the formation volume as a function of pressure. temperature, and gas-condensate ratio has been presented. This correlation i-considered applicable only to the range of conditions covered in this investigation for systems of similar nature and com-position. INTRODUCTION Often a knowledge of the influence of temperature. Pres sure, and relative quantities of condensate and gas upon the volumetric behavior of fluid*. from a given field has been employed in estimating the Conditions existing in the reservoir. The early work of Beecher and Parkhurst' and the more rrcent Studies of the Dominguez' and San Joaqnin Valley Field;: are examples of experimental studies of this character. Correlation of the volumetric behavior of naturally occurring hydrocarbon mixtures has Proved To be feasible' . at low gas-conden-sate ratios-; for fluids*. from a single geographic area. Progres 5 is being made in the Prediction of the volumetric behavior of distillate systems at high gas-condensate ratios and pressure.. However. the estimation of the retrograde". ' dew-point pressure still remains uncertain and probably can best be established by experiment. For this reason it appears of interest to make available experimental information about the volumetric behavior of mixtures of condensate and gas from five different part.; of a single field. This material shows a rough similarity to data for mixtures of condensate and gas from San Joaqnin ' Valley Fields. In the present paper. 110 attempt has been made to coordinate the available information on tbe volumetric behavior of fluids from condensate fields. It appears hopeful that progress can be realized by the utilization of recently developed equa-tions of state.!' However. the primary problem in any such correlation is the proper identification of the components and constituents of naturally occuring systems. Additional experimental information may be required to permit the determination of the values of the constants of the equation of state for References given at end of paper. Manuscriptt received in the office of the Petroleum Branch June 11. 1951. *For present purposes low gas- condensate ratios will he considered as those below 3,000 cu ft/bbl, and the region of high gas-condensate ratio will be limited to compositions containing more than 5,000 cu ft of pas per bbl of condensate. Mixtures Containing between 3,000 and 5.000 CII ft OF gas per bbl will be considered of intermediate gas-condensate ratio.

Jan 1, 1952

By J. L. Taylor, Pol Duwez

THE present knowledge of the Ti-Al system is limited to the portion of the diagram extending from pure aluminum to the intermetallic compound TiAl3' A preliminary investigation of the titanium-rich Ti-A1 alloys revealed that the solubility of aluminum in titanium is quite extensive and that the transformation temperature of titanium is raised by alloying with aluminum. From an X-ray diffraction study of alloys containing up to 75 atomic pct Al (TiAl3), the phase boundaries at 750°C were located as follows: the a titanium solid solution extends from 0 to about 36 atomic pct Al, a two-phase region exists from 36 to 46 atomic pct Al, and an intermediate phase of varying concentration extends from 46 to 62 atomic pct Al. The purpose of the present paper is to describe the crystal structure of this new phase. The alloys were prepared by melting in a helium arc furnace on a water-cooled copper plate, using a furnace essentially the same as that described in ref. 2. The titanium metal was of the iodide type furnished by the New Jersey Zinc Co., and a piece of pure aluminum (99.99 pct) was obtained through the courtesy of the Aluminum Co. of America. After melting, the samples, weighing approximately 5 g, were sealed in evacuated fused silica tubes and homogenized for 4 hr at 1000°C. The specimens were then held for 10 days at 750°C and rapidly cooled to room temperature. Filings were then taken from each sample, sealed in evacuated vials, annealed at 750°C for 4 hr, and rapidly cooled to room temperature by quenching the vials in water. Powder diffraction patterns were obtained with a 14.32 cm diam camera, using K copper radiation filtered through a nickel foil. Structure Determination The X-ray diffraction patterns of the alloys containing 46, 50, 55. 60, and 62 atomic pct Al were identical, except for a slight shift in the positions of the reflections. The alloy containing 55 atomic pct Al was chosen far structure determination. The powder pattern contained 37 reflections and the a, doublets were well resolved in the back-reflection range. The relatively small number of reflections and the sharpness of the pattern were considered as an indication of a simple atomic arrangement in a relatively small unit cell of high symmetry, and an attempt was therefore made to solve the structure without single crystal work. A satisfactory fit was found on a large scale Hull-Davey tetragonal chart for an axial ratio of approximately 1.45. All the reflections were readily indexed and the tentative unit cell dimensions were computed to be a - 2.81 kX, c - 4.07 kX, and c/a = 1.448. The indices (hkl) and the values of d and sin2 A are given in Table I. The agreement between observed and computed sin2 is quite satisfactory. Once the unit cell was known, the details of the structure were readily found. First, it is obvious that the number of atoms per unit cell cannot exceed two, since more than two atoms would lead to a density greater than that of pure titanium. The next logical step is to assume that the two atoms occupy the corner and center of the cell. If this assumption is correct, the structure may be described as follows: space group D1, - P 4/mmm; one titanium atom in a: 000; one aluminum atom in d: 1/2 1/2 1/2: and no extinctions. Assuming this structure to be the correct one. intensities were computed by means of the usual equation: sin2 0 cos 0 where F is the structure factor; 8, the Bragg angle. and p, the multiplicity factor. The calculated values of intensities are compared in Table I with the visually estimated intensities. The agreement is quite satisfactory and the structure is therefore confirmed. The crystal structure of TiAl just described is the same as that of AuCu ordered. The AuCu ordered

Jan 1, 1953

By B. F. Decker, J. H. Keeler, W. R. Hibbard

Textures of hot-rolled, of cold-rolled, and of cold-rolled and annealed zirconium sheets were determined by use of an X-ray spectrogoniometer. All textures showed a tilt of the basal planes 240" from the rolling plane about the rolling direction. Deformation textures showed the [1010] direction to be approximately in the rolling direction. After recrystalliza-tion, the 111201 direction was approximately in the rolling direction. This texture was sharpened by annealing above the allotropic transformation. Textures of sheet produced by alternately cold rolling and then annealing were less strongly oriented than sheets that underwent severe deformation without intermediate anneals. THIS study of the deformation and annealing textures of high purity zirconium was carried out to supplement an investigation of zirconium and zirconium-base alloys. The spectrogoniometer method of pole figure determination was used and yielded more detailed and precise information than the photographic techniques used by previous investigators. Earlier investigations of the rolling textures of hexagonal metals titanium,1,2 beryllium," and zirconiumL having a c/a ratio slightly less than 1.6 disclosed a rotation of the basal planes from the rolling plane about the rolling direction so that the basal poles were inclined toward the transverse direction, and with one exception,& the [1010] direction was found parallel to the rolling direction. Similarly, hot rolling produced a spread of the basal poles in the transverse direction,2,3,5 With the [1010] direction parallel to the rolling direction. The textures of these metals after recrystallization differ somewhat from each other although still exhibiting the transverse rotation of the basal poles about the rolling direction. Zirconium" was found to have the [1120] direction in the rolling direction, beryllium" continued to have the [1010] direction in the rolling direction, and titanium' was shown to have the [1010] direction in the rolling direction (the cold-rolling texture) when annealed at 1000°F, changing to the [1120] direction in the rolling direction when annealed at 1500°F. It was suggested2 that the titanium retains the deformation texture after re-crystallization, and preferred grain growth at higher temperatures produces the second texture. Annealing above the allotropic transformation temperature did not appreciably change the high temperature recrystallization texture for titanium.' Coarsegrained zirconium showed more spread" in the transformation texture than in the recrystallization texture. The material used in this investigation was crystal-bar zirconium prepared by the iodide process. The major impurity was hafnium which was less than 0.3 pct. The composition obtained by spectro-graphic analysis is given in Table I. The hot-rolled specimen was prepared from a 50-g button of zirconium arc-melted in a water-cooled crucible under an argon atmosphere. This button was enclosed in a welded sheath of soft iron and the ensemble was rolled at 825°C. After each pass the sheath was put back into the furnace to allow it to return to the desired temperature. The zirconium ingot was reduced in this manner to a thickness of about 0.035 in., a reduction of approximately 90 pct. The sheet after removal from the sheath was alternately polished and etched to a thickness of 0.001 in. Cold-rolled specimens were prepared from crystal bar by a series of reductions of approximately 0.010 in. until the sheet reached a thickness of 0.010 in. after which it was cold-rolled to 0.001 in. thickness between sheets of alloy steel. The total cold reduction amounted to about 99.8 pct. Specimens were annealed in an evacuated Vycor or quartz tube. The transmission pole figures were obtained by the quantitative method of Decker, Asp, and Harker6 modified by Geisler7 using an X-ray spectrogoniometer with Cu Ka radiation. Unless otherwise specified all four quadrants were examined with 5" increments along the radius of the pole figure. Absorption corrections were made for changes in length of the X-ray path through the sample as inclination of the

Jan 1, 1954

By Alfred R. Powell

LATE in 1944 a group of petroleum and coal technologists was organized in Wash¬ington under the sponsorship of the Petroleum Administration for War and the U. S. Bureau of Mines. This group, known as the Technical Oil Mission, was placed under the authority of the Technical Industrial Intelligence Committee of the Joint Chiefs of Staff, which was then promoting the organization of other missions to be sent to Germany to obtain technical information regarding various industries in that country. The objective of the Technical Oil Mission was to obtain technical data on German production of synthetic oil from coal and all other commercial practices and research projects connected with the oil industry. The primary purpose was to obtain such data as were immediately essential to assist the war against Germany and Japan, with secondary emphasis on technical information that would assist American industry after the end of the war. Since oil-production facilities in Ger¬many and the occupied countries con¬trolled by Germany were the top targets for strategic bombing, the Technical Oil Mission was the first group under the Technical Industrial Intelligence Com¬mittee to enter Germany. A few members of the group flew to London in February 1945 and the majority of the remainder followed in March. In London they were joined by an equal number of British investigators and the combined, mission was then sent into Germany under the authority of the Combined Intelligence Objectives Subcommittee (CIOS). A sum¬mary of the more important information gained by this combined British and American mission is presented below. GERMAN OIL PRODUCTION The oil available to Germany, both natural and synthetic, and including pro¬duction from occupied countries, reached a peak in late 1943 and early 1944. Pro¬duction at this time was at the rate of 1,200,000 metric tons per month, equiva¬lent to i00,000,000 bbl. per year, or about 7 per cent of the wartime production in the United States. Of this production, 50 per cent was crude oil from Romania and Hungary, 12 per cent was crude oil from Germany itself, including Austria, 9 per cent com¬prised coal-carbonization products, that is, tar and light oil, and only the remaining 29 per cent was synthetic liquid fuel made from coal or tar. Synthetic liquid fuel made in Germany was equivalent therefore to less than 2 per cent of war¬time production of crude oil in the United States. The success of the strategic bombing campaign against German oil-production facilities is demonstrated by the fact that by September 1944 the Ploesti fields of Romania and much of the syn¬thetic oil production had been eliminated, so that the amount of oil available to the German war machine was only 25 per cent of the peak quantity mentioned above. By early 1945, when the combined British and American oil mission was entering

Jan 1, 1946

By R. B. Roof, J. J. Lombardo

DURING the course of investigations on U-Cb alloys by X-ray diffraction techniques, an impurity or an additional phase was observed in severa1 alloys of various compositions. The X-ray dif-

Jan 1, 1959