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By O. N. Carlson, H. A. Wilhelm, J. M. Dickinson

On the basis of microscopic studies, melting-point observations, and X-ray analyses, a phase diagram is proposed for the Cb-V system. A complete series of solid solutions is formed with a minimum in the solidus at 1810°C near 35 wt pct Cb. No compounds or intermediate phases were found in the system above 650°C. THERE is an ever increasing need for better structural metals and alloys for use in nuclear reactors. In addition to the normal properties of engineering structural materials, such as high temperature strength, resistance to corrosion, and fabric-abil~ty, the nuclear properties of the material must be considered. In a nuclear reactor it is important to conserve neutrons, so a material which removes these neutrons from the reaction excessively is considered to have unfavorable nuclear properties. In nuclear-reactor design the engineer must have nuclear as well as other data available on alloys in order to make a wise selection of materials. Due to the fact that many of the common structural materials have undesirable nuclear properties, it is vital that new alloys of metals having more favorable nuclear properties be investigated. Columbium and vanadium are both high melting metals, both exhibit resistance to chemical attack, and no great difficulty is encountered in fabricating them into desired shapes. With proper treatment both metals can be cold rolled extensively without failure. In addition they have desirable nuclear properties for certain types of reactors. Therefore, the alloys of columbium and vanadium should be of interest in the atomic energy program. Since an alloy-development program is enhanced by a knowledge of the phase equilibria of the components, this investigation was undertaken to establish the phase diagram for the Cb-V system. According to the Hume-Rothery rules for alloying,' the chemical similarity, crystal structure, and atomic-size factor are favorable for a complete series of solid solutions for this system. Both elements are in the same family of group V of the periodic table and thus are quite similar chemically. The crystal structures of columbium and vanadium are compatible for extensive solid solubility, since both have body-centered-cubic structures. The atomic diameters of columbium and vanadium are 2.85 and 2.62Å, respectively. This difference of slightly more than 8 pct is well within the 15 pct maximum difference allowed for extensive solid solubility. Experimental Procedures Source of Materials: Columbium powder and sheet trimmings were obtained from the Fansteel Metallurgical Corp. According to the manufacturer the metal contains less than 1 pct impurity. An analysis of the metal showed approximately 1800 ppm C in the powder while the sheet trimmings contained less than 500 ppm C. Spectrographic analysis showed minor amounts of Ca, Cr, Fe, Mn, Si, Ti, V, and Zr in both forms of the columbium. No commercial source of vanadium having the ductility and purity desired was available to the authors at the beginning of this investigation. As a result, all of the vanadium used in this study was prepared by the bomb reduction of vanadium pen-toxide with calcium employing the method reported by Long.' Yields of massive vanadium normally were about 80 pct. Chemical analysis of the vanadium prepared in this manner showed the presence of 200 to 500 ppm N and 800 to 1000 ppm C. Minor amounts of Ca, Fe, Mn, Si, Zr, Cr, and Cb were detected by spectrographic analysis. This vanadium metal was ductile and was cold rolled into 5 mil sheet. Annealing was not necessary during this rolling and the metal retained its cold-rolling characteristic after are-melting. Preparation of Alloys: The Cb-V alloys were prepared by melting pieces of vanadium sheet togethel-with columbium in the form of sheet or pellets of powder. The melting was carried out under argon in conventional arc-melting equipment employing a tungsten electrode and a water-cooled copper crucible. Each alloy was remelted three or four times, inverting the alloy after each melting in order to assure complete mixing. Alloys normally were obtained as round flat disks, weighed approximately 70 grams, and had roughly the shape of a disk 1 1/2 in. in diameter and 1/4 in. thick. Half of each alloy

Jan 1, 1955

By J. H. Scaff, W. G. Pfann

The effects of impurities on the electrical properties of silicon are discussed in a companion paper by Messrs. Scaff, Theuerer, and Schumacher.' It was shown that an ingot of silicon which contained boron and phosphorus in certain concentrations consisted partly of p-silicon and partly of n-silicon2 and that the common boundary between these regions was the source of a photo-voltage. These features were ascribed to the segregation of boron and phosphorus during solidification of the ingot. Because of different rates of segregation the first portion of the ingot to freeze contained a molar excess of boron over phosphorus and hence was p-silicon. The last-solidified region contained a molar excess of phosphorus and was n-silicon. This paper is concerned with the microscopic examination of such ingots. The microstructures are rather unusual in certain respects and their study has helped to show how the segregation of minor elements modifies the electrical properties of silicon. The microstructures of ingots prepared from two lots of silicon, designated 1 and 2, will be described. These lots were obtained from the Electro Metallurgical Co. Both contain 99.8+ pct silicon, but they differ in impurity analysis. For each lot will be shown the microstruc-ture of a slowly cooled ingot, for which the time of solidification is 6 min., and that of a rapidly cooled ingot, for which the time of solidification is 2 min. The ingots were prepared from 45 g charges of silicon. When an alloy freezes a cored structure is generally produced because of insufficient diffusion to remove concentration differences established during freezing. As a result of coring the first portion of a crystal to freeze is usually poorer in solute than equilibrium requires, while the last material to solidify is richer. In a columnar grain, growth is principally unidirectional and coring can produce a concentration gradient along the entire grain. By carrying the picture of the coring process from a single grain to an entire ingot in which columnar grains are aligned more or less radially, one can easily see that a composition gradient could be established between the outside and center of the ingot. This, to a considerable extent, occurs during the solidification of a slowly cooled ingot of silicon. Slowly Cooled Ingot In slowly cooled ingots of both lots of silicon, columnar grains which are approximately radial in direction and longest at the top of the ingot enclose a region in which the shapes and align- ment of the grains are considerably less regular. The central region is last to solidify and contains cavities and inclusions. The sketch of Fig 1 illustrates some of these features, as well as others which will be described below. If a polished section of an ingot is etched in a mixture of 2 parts of 20 pct HF and 98 parts of HNO3 the p-n barrier becomes visible. The overall shape of the barrier in a slowly-cooled ingot is roughly oval, as may be seen in Fig 1. The proportion of n-silicon, that is, the region enclosed by the barrier, is considerably greater in ingots of lot 1 than in those of lot 2. A sample about 34 x 1/8 X 1/8 in. was cut from a slowly cooled ingot of lot 1 with its long direction extending from the top of the ingot to the porous region. Thus its upper half was psilicon, its lower half n-silicon. A longitudinal face was polished and etched in the acid mixture. The areas of interest are identified in Fig 2. A band of parallel markings extends across the sample in the columnar region. These striae are roughly parallel to the upper surface of the ingot and cross grain boundaries with only slight deviations. They are quite distinct in the n-silicon,

Jan 1, 1950

By W. F. Kieschnick, Eugene Harrison, W. J. McGuire

W. J. McGuire, et al, are to be commended for their undertaking of a mathematical solution of a very difficult problem. Unfortunately, however, a mathematical approach requires the application of several assumptions. These assumptions appear to be unrealistic and lead to answers which do not describe what actually happens when hydraulically fracturing oil and gas wells. Considering laboratory confirmation of breakdown phenomena, the authors appear to have tested their theories only on cement specimens and on samples of Austin limestone, much too small to provide any fracture system. This work resulted in the formation of vertical fractures. If the authors had tried similar experiments on thick walled cylinders made from almost any sandstone cores, they would have found that, using crude oil as the breakdown fluid, horizontal fractures would almost always occur, and at pressures much lower than any calculated. They would also find that by confining the fluid to within the bore (using oil base mud for example) on similar samples, the pressures required to burst the cylinders would be considerably higher and most of the fractures would be vertical. This breakdown pressure behavior has been duplicated in wells in Texas, Oklahoma, Kansas and in Wyoming. Considering field data the phenomena of different breakdown pressures for different breakdown techniques can be further illustrated. Most production and service personnel will agree that a breakdown can be more easily obtained if injection into a formation can be established prior to the occurrence of the breakdown. This is true whether the formation being treated is completed as open hole or as a perforated interval. This is clearly illustrated by a Lakota well in Wyoming, completed open hole at a total depth of 7,358 ft. An attempt to vertically fracture this well failed when a bottom hole pressure of 10,326 psi was insufficient to break down the formation. A non-penetrating type fluid (oil base mud) was in the well at the time the breakdown was tried." The oil base mud was then cleaned out of the well and replaced by a 30" API gravity crude oil. With this oil in the hole the formation breakdown was easily accomplished at a bottom hole pressure of 3,607 psi. This large difference in fracture pressures would be impossible according to the theories presented by McGuire, et al. The authors have used as an example the breakdown pressures experienced when acidizing Permian Basin wells. During acid treatments of limestone and dolomite the "breakdown" (drop-off in pressure) seldom occurs until some injection of acid has been accomplished. In these cases the breakdown is most likely to result from the chemical reaction of acid and rock in already existing vugs and fractures rather than from making a new fracture by hydraulic pressure. If this is true, then results in the Permian basin should not be used to validate the authors' calculations. *** AUTHORS' REPLY to ROSCOE C. CLARK and HENRY F. COFFER The purpose of our laboratory experiments in which thick-walled rock cylinders were hydraulically fractured was to determine the validity of the "thick pipe" formula for brittle materials, and not to predict nor demonstrate directly the orientation of field fractures. Our conclusions concerning field results resulted from calculations involving the "thick pipe" relationship as well as considerations of overburden stresses, rock strengths, and the geometry and dimensions of the field system. Clark suggests that had the models been more porous or contained weak bedding planes, horizontal fracturing would have occurred. This is undoubtedly true provided external stresses similar to those in the earth's crust are nor imposed. However, if we were going to design experiments to represent directly the field case we would impose the proper stresses on the models. It is generally recognized that the vertical compressive stress in the earth's crust arising from the weight of the overburden is approximately 1 psi/ft of depth. Then, as an example, even though a horizontal bedding plane has zero strength, the formation cannot be separated to form a horizontal fracture unless the hydraulic pressure exceeds the stress due to overburden. And in those cases in which the stress resisting vertical fracturing is significantly less than that resisting horizontal fracturing, vertical fractures should result, notwithstanding horizontal plane weaknesses. We agree that breakdown pressure will be less if the fracturing fluid penetrates the formation. In Appendix HI of our paper it is shown that leak-off reduces the pressure necessary to initiate either a horizontal or vertical fracture. It would be difficult to attempt to

Jan 1, 1955

By H. E. Cline, E. F. Koch, J. L. Walter

It has been postulated and, in a few instances shown, that some kind of dislocation structure will be present at semicoherent interfaces to accommodate small lattice mismatches. In the present study of the NiAl-Cr eutectic, regular arrays of interface dislocations are observed at the boundary between the chromium-rich rods and the NiAl-rich matrix. The networks were examined by transmission electron microscopy and selected area diffraction. The rods and the matrix have a crystallographic relationship in which all directions and planes of the two phases are parallel. The dislocation networks are cmposed of a<100> dislocations lying on the intersections of the cylinders with (100) planes. Dislocations forming hexagonal rather than square arrays are also observed at certain areas of the network. The morphology of the network is consistent with the interpretation of mismatch being accommodated by interface dislocations in the cylindrical geometry. The measured spacing between dislocations was used to calculate an apparent lattice mismatch between the phases (˜0.35 PCt)interface network energy (-140 ergs per sq cm), and network strengthening (-10,000 psi). It has been proposed by Frank and Van der Merwe1 that dislocations should be present at the boundary between two semicoherent crystallographically related phases. The role of the interface dislocations would be to reduce the internal stresses, caused by the mismatch in atomic spacing across the interface. Such dislocations have been observed at the interface between expitaxially grown films and Substrates.2-4 Interfacial dislocations have also been observed at precipitate-matrix interfaces.&apos;-&apos; Directionally solidified eutectics have been shown to have semicoherent phases1&apos; and would, therefore, be expected to have interfacial dislocations as found by Weatherly at a lamellar fault in A1-A12Cu.11 The NiAl-Cr eutectic appeared to be a promising system to examine because both phases are cubic, the lattice mismatch is small, and the phases are crys-tallographically related. Furthermore, the eutectic is easily thinned for transmission microscopy. Indeed, interfacial dislocations were observed and this report describes the nature of the dislocation networks in the boundary between the NiAl-rich ma-trix-phase and the fine chromium-rich rods.21 I) EXPERIMENTAL PROCEDURE Ingots, 3/4 in. in diam and 6 in. long were made by melting in vacuum and casting under argon using 99.9 pct pure material. The composition, in at. pct, was 33 pct Ni, 33 pct Al, and 34 pct Cr. The ingots were then placed in A1203 crucibles on a water-cooled base, melted by means of induction, and withdrawn from the hot zone at the rate of 1 in per hr under argon.* * T his material was first directionally solidified in this laboratory by E. R. Stover Slices were taken perpendicular to the growth direction of the directionally solidified ingot for metallography and for transmission electron microscopy. The electron transmission samples were thinned mechanically, then thinned electrolytically in A-2 electrolyte* *A-2 electrolyte: 62 ml perchloric acid, 700 ml ethanol, 100 ml butylcellosolve, 137 ml distilled H20. until a hole appeared in the foil. 11) EXPERIMENTAL RESULTS A) Optical Microscopy. The microstructure, viewed on a plane perpendicular to the growth direction, is shown in Fig. 1. The structure consists of cells or colonies of parallel chromium-rich rods in the NiAl matrix. The cells occur when there are impurities present12 or, in a ternary eutectic, if the composition is slightly off the eutectic composition. The axis of the chromium-rich rods is parallel to the growth direction except near the cell boundaries. Here the rods may assume angles to the growth direction; however, examination shows that the crystallographic relationship between the rod and the matrix remains the same. Fig. 1 includes cell boundaries where the rods formed at a large angle to the growth direction. The variation of rod position across the cells made it possible to Fig. 1-Structure on plane perpendicular to growth direction. Rods near cell walls are at large angle to growth direction. Magnification 315 times.

Jan 1, 1970

By R. O. Williams

Analytical representation of the thermodynamics of solutions is highly desirable from the standpoint of accuracy, compactness, and numerical manipulations. In particular, computer calculations are greatly implemented. Mathematical considerations show that previous expressions have one or more serious defects. This investigation shows a Fourier series to be satisfactory but that it is also possible to derive a new series which fits certain additional conditions. Included examples show the value of analytical expressions in giving a simple characterization of each system using some two to five parameters, the elimination of the Gibbs-Duhem integration, and the es timation of the error for the experimental function as well as derived functions. It is further shown that the present characterization provides easy comparison between systems. IN the past, thermodynamic calculations have depended to a considerable extent on tabular and graphical methods. As the volume and precision of such data increase such methods become less satisfactory. Specifically, the selection of the optimum representation and the estimation of errors require statistical methods which in turn require analytical representation. The utilization of such data require further manipulations which are best done analytically for maximum precision. For example, phase equilibria are determined by common tangents to free-energy curves: a graphical determination is normally of low accuracy. As computers are increasingly used analytical representations become almost mandatory. Insufficient mathematical consideration has been given previously to the selection of empirical expressions. Those expressions having some theoretical justification are generally too inflexible and mathematically unattractive. We consider the problem in some detail and show that a Fourier series can be effectively used. Also a new series is defined which has certain advantages. ANALYSIS We wish to consider the analytical representation of the heat of mixing, AH, the excess free energy, ?Gxs, and the excess entropy, ?sXS, as a function of composition, X, for binary solutions relative to the pure components in the same state. When a distinction is not required, we use W to denote any one of the above functions. One may use a Taylor expansion around X = 0 to generate a power series. As the derivatives are un- known we represent the series as W = A + BX + CX2 + DX3 + EX4 + ... [l] where the constants A , B, C , ..- are to be selected to provide some optimum fit. For the extremes of composition W is necessarily zero so it follows that A = 0 [2a] B +C + D + E +••• = 0 [2b] Nonelectrolytes, which we are considering, appear to satisfy the condition that d3W/dx3 = 0 [3] in the terminal regions. This is the basis of the a, ß, and Q functions used by Hultgren et al.&apos; and others. While this condition does not have a strong theoretical basis it does appear desirable that any analytical relation should satisfy this condition. Darken2 and Turk-dogan and Darken3 have shown that many systems exhibit this behavior over an extended range from each terminal region, departure being restricted to a limited intermediate region. Since we have no a priori knowledge as to where this transition occurs we can require that this condition be satisfied only as a limit at the extreme compositions as a general condition. We will show later how more restricted conditions can be included in specific solutions. Darken2 has called this behavior the quadratic formalism; we call our application the limiting quadratic formalism, LQF. This condition applied to the above power series requires that D = 0 [4a] 4-3-2E +5-4-3_F + 6 • 5 . 4G + ••• =0 [4b] The form of the power series normally used, due to Margules,4 is W=X(1-X)(A + BX + CX2 + DX3 + EX4 + •••) [5] where A, B, C, --. are a new set of constants. (Guggenheim5 has given a variation of this expression in a more desirable form. Since, however, it is contained in the above expression it does not require separate consideration.) This form is precisely what results by incorporating the conditions in Eq. [2] into the power series and regrouping the constants. The LQF requires that B =C [6a] and 4.3.2(D-C) +5-4-3(E-D) + ••• =0 [6b] Thus, the correct form of the Margules expression with two adjustable parameters is w =X(1-X)[A + B +X2-2/3x3)] 171 and the EX4 term must be included before three adjustable parameters are permitted.

Jan 1, 1970

By E. V. Potter

For a nurnber of years, it has been known that manganese made by electro-deposition under certain conditions is ductile while under other conditions it is very brittle. The ductile metal is gamma manganese normally stable only between 1100 and 1138°C1; the brittle metal is alpha manganese, stable up to 727OC. The ductile metal is not stable, but gradually changes to the brittle form; the time required to complete the transfornlation is about 20 days at room temperature. Other observations have indicated that the transformation is completed in 10 to 15 min. at about 125°C, while at — 10°C, no appreciable change occurs in 9 months. The properties of gainma and alpha Illanganese in the pure state are ordinarilj difficult to determine because the gamma structure cannot be retained by normal quenching procedures and alpha manganese is so brittle, it is difficult to obtain specimens free from flaws. In a recent investigation2 some properties of gamma and alpha manganese were determined by studying the ductile electrolytic metal and determining the changes in its properties as it transformed to the brittle alpha form. These investigations provided an excellent opportunity for following the progress of the transition and studying its mechanism. The results of a series of such investigations are reported in this paper. Procedure Various properties of manganese were determined starting with the metal in the original ductile gamma form and following the subsequent changes in its properties as the metal transformed to the brittle alpha form. These observations were made at various temperatures, the data providing information regartling the mechanism of the transformation as well as the effect of temperature 011 the transition rate. Structure and resistivity values gave the most significant results, so this paper is concerned primarily with them. The structure was studied microscopically as well as by X ray diffraction. The resistivity was determined on strips of the metal by measuring the potential drop across a given length of the specimen. Current was passed through the specimen by wires soldered to its ends, and the potential connections were made by wires looped around the specimen near its center. The current was determined by the potential drop across a standard resistor connected in series with the specimen, the potential drop being measured on a potentiometer. In the temperature range from room temperature to 100°C an ordinary drying oven was used to heat the specimen. This was entirely satisfactory except at 100°C, where the time required to heat the specimen was long compared to the transition time, making the initial section of the resistivity curve unsatisfactory. To overcome this limitation, at 100°C and higher a thermostatically controlled oil bath was used to heat the specimens. The block on which the specimen was mountetl was plunged into the hot oil at the start of each test. The heating time was thereby reduced from 5 min. to about 6 sec, and dependable resistivity values could be obtained through 160°C. At this point the whole transition, including the warm-up time for the specimen, required only about 20 sec and it was not considered worth while trying to extend the temperature range further. Aside from the heating problem, the problem of making a sufficient number of accurate resistivity determinations became more and more difficult as the temperature was raised. Using the manually operated potentiometer, 100°C was about as far as it was possible to go. At this temperature and above, a self-balancing photoelectric recording potentiometer was used. Its response was quite rapid, and it proved to be entirely satisfactory all the way through 160°C, where the tests were stopped because of the specimen heating problem rather than any limitation of the potentiometer recorder. The metal used in these tests was prepared at the Salt Lake City laboratory of the Bureau of Mines. The method of preparation is discussed in a paper by Schlain and Prater.3 The sheets were about 2 3/8 by 5 3/16 in. and varied from 10 to 16 mils in thickness. They could be cut readily into pieces suitable for the various tests. X ray and microstructure determinations were made on pieces about 1/8 to 1/4 in. wide and about 1 in. long, while resistivity measurements were made on strips as long as possible and about 55 in. wide. The thickness of each sheet was not uniform over all its surface. This had no bearing on the X ray and microstructure determinations, but sections as nearly uniform and free from flaws as possible were chosen for the resistivity determinations. The gamma manganese was electro-deposited at 30°C, the time of deposition ranging from 5 to 12 hr for each sheet. Whenever possible, the tests were started directly after the metal was stripped from the cathode; otherwise the sheet was placed immediately

Jan 1, 1950

By Robin O. Williams

AN investigation has been made of the Ni-Cr system for the purpose of elucidating certain points, namely the nature of aging in both terminal solid solutions and the nature of the phase diagram. Information pertaining to solubilities and precipitation has been obtained. Experimentation Five alloys, Table I, were arc melted in a cold copper crucible using electrolytic chromium and car-bony1 nickel, both dry hydrogen treated. These 100 g buttons were homogenized 24 hr at 1300°C in dry hydrogen and air cooled. Powders were prepared by filing or pulverizing and subsequent heat treatment was done in vacuum or helium using titanium chips as a getter. Powder of —80 mesh was filed from the 60 pct Ni alloy quenched from 1000°C and was sealed in silica under vacuum using a 250 °C outgassing. After aging as indicated in Table I1 the lattice parameters were measured on the quenched samples using the standard cos" 6 extrapolation. These parameters are considered accurate to roughly 0.0001A. In all cases chromium lines of 2.8812 ± 0.0005Å at 30°C were found. Drastic quenching from sufficiently high temperatures produced very sharp body-centered-cubic lines in the first four alloys without indications of transformations. Temperatures to 1250°C were used. Solid samples less than 1/16 in. thick were quenched in water without transformation and the powders could be adequately quenched in small helium filled thin wall silica tubing using a water quench. For those powder samples which were quenched from the two phase field the relative intensity of the body-centered-cubic lines and the face-centered-cubic lines were estimated and extrapolated to give the indicated solubility data in Fig. 1. The data for the two higher alloys were somewhat limited, the plotted points being the lowest temperature where no nickel phase was found. Neither filing, abrading, pulverizing, nor cooling to —190°C produced any new diffraction lines for solid samples quenched from the single phase region, nor did the character of the body-centered-cubic lines change Single phase body-centered-cubic powders likewise did not change on cooling to —190°C. Also, samples which had some precipitation due to inadequate quenching showed no additional changes under these conditions. The first change apparent by X-ray diffraction form samples quenched almost fast enough to prevent precipitation was the diffuseness of the body-centered-cubic lines, particularly on the low angle side, For slower cooling rates the diffuse face-centered-cubic lines appeared. Work on the large grained castings showed profuse streaking through some of the Laue spots while oscillating patterns showed broad body-centered-cubic and face-cen-tered-cubic lines as well as some new lines. For the 23.6 pct Ni alloy the new lines corresponded to 2.16, 1.96, and 1.86A and were more similar in character to the face-centered-cubic lines than the body-centered-cubic lines. Samples which were air cooled gave only face-centered-cubic and body-centered-cubic lines which were still broad. One pattern indicated that face-centered-cubic (111) plane was parallel to a body-centered-cubic (110) plane. For those samples which were examined by light microscopy there were details which were not resolved. However, varied and beautiful structures were obtained. Fig. 2 is of an alloy quenched in a helium filled silica tube from the single phase region and shows particles associated apparently with dislocations which are arranged in low angle boundaries. Finer, general precipitation has also taken place within the grains. Figs. 3 to 5 show the variety of structures produced in these alloys on continuous cooling. It appears that there are four distinct modes of precipitation as evidenced by these, figures. Annealing these structures at higher tem-peratures in the two phase field gives structures as shown in Fig. 6, which shows nickel plates in the chromium matrix which reprecipitated nickel on a much finer scale of the final quench. Lower annealing temperatures and shorter times naturally give finer plates of the nickel-rich phase. Samples of the first four alloys were annealed for appreciable times between 900º and 1250°C and gave structures like Fig. 6. The relative amounts of the two phases were measured and extrapolated to give solubility data as indicated in Fig. 1. The point at 1250°C was deduced from data of Oxx.1 These and most of the other samples were checked for ferromagnetism but none was apparent. In fact, it appeared that the magnetic susceptibilities were not more than three times that for paramagnetic chromium. Discussion In Fig. 1 it is seen that the solubility of nickel in chromium can be represented by a slightly curved line on the usual log X vs 1/T plot, These data are believed to be accurate to roughly 5 to 10º. There is only fair agreement with the data of Taylor and

Jan 1, 1958

By James W. Shaffer

Aluminum is the most abundant metallic element and forms 8% of the earth&apos;s crust. Because of its chemical activity it does not occur in nature in a metallic form but principally in the silicates, feldspar, mica, shale, and clay. Bauxitic raw materials-bauxitic clay, aluminous laterite, and bauxite-develop from these aluminum silicate sources. Bauxite and aluminous laterite are naturally occurring mixtures of aluminum hydroxides and mineral impurities, and they are the major sources of aluminum. End Uses Bauxite is the basic ore of the aluminum industry. It is refined into aluminum hydroxide and calcined for feed stock to electrolytic smelters. Over 90% of the tonnage mined is refined and smelted into aluminum. The total tonnage consumed by other industries is small in comparison, but bauxite and alumina are important for the manufacture of refractories, ceramics, chemicals, abrasives, pharmaceuticals, catalysts, cements, dessicants, and flame retardants. Some aluminum-containing industrial products are made from clays and other materials but the majority originate from bauxite directly or from aluminum hydroxide and alumina refined from bauxite. Lateritic type bauxite is also used as building and decorative stone. Geology Mineralogy and Properties Bauxite is a mineral raw material composed principally of one or more aluminum hydroxide minerals and impurities of silica, iron oxide, titania, and other elements in minor or trace amounts. The important aluminum hydroxide minerals are gibbsite, boehmite, and diaspore. Bayerite, nordstrandite, and tucanite are naturally occurring aluminum hydroxides but are rare and not commercially important. Mixtures of gibbsite and boehmite are common, boehmite and diaspore less common, and gibbsite and diaspore rare. Gibbsite is aluminum trihydroxide. Boehmite and diaspore are aluminum oxide hydroxides and have essentially the same composition, but diaspore has a greater hardness and a more tightly bonded and denser atomic structure. Although not technically correct, the commercial terms trihydrate and monohydrate are frequently used to distinguish types of bauxite. If bauxite is all or nearly all gibbsite, it is referred to as a trihydrate ore; if boehmite or diaspore or a mixture of both are the dominant minerals, the bauxite is called a monohydrate ore. If substantial amounts of both gibbsite and boehmite are present, the bauxite is frequently referred to as a mixed ore. Properties of major bauxite minerals are shown in Table 1. Silica, iron oxides, and titania are major impurities always present in deposits of bauxitic raw materials.. Silica impurities usually occur as clay minerals, micaceous minerals, or quartz, but cristobalite and feldspar are found in some deposits. The most common clay minerals are kaolins; however, illites and chlorites are often found in small quantities. Varying proportions of aluminum, iron, and clay minerals make bauxite, aluminous laterite, ferruginous laterite, and bauxitic clay deposits gradational with each other. Hematite and goethite are the common iron mineral impurities in many bauxites. Magnetite and ilmenite occur in deposits derived from basic igneous rocks. Other iron minerals associated with bauxite are maghemite, pyrite, siderite, and marcasite. The titanium minerals found in bauxites are anatase, leucoxene, rutile, ilmenite, titaniferous

Jan 1, 1983

By AIME AIME

MESSRS. BURT and Shockley and others have been for three years urging upon the Institute the matter of simplified spelling. The Institute endeavors to be progressive in the matter of spelling and. is in cooperation with the national body which has this in hand and which has its principal office in New York. Our spellings are changed as fast as-good usage will permit (so it seems to the Board of Directors) and the national body, which is the real representative of the simplified spelling interests, deems that the Institute is making real progress, although not quite as fast as extremists would wish. In January, 1919, Mr. Shockley sent out letter ballots which created a misconception in the minds of many members, who received the impression that the communications came from the Institute instead of from a small group of members. As a result of this propaganda 163 proposal-ballots were sent in to the Institute divided as follows: Yes ?114 No ?20 Negative letters ?9 Split ballots ?20 Total ?163 The total membership at this time was about 7200. The disgusted members would not vote.

Jan 1, 1920

By O. D. Gonzalez

Existing measurements for the steady-state permeation of hydrogen in a iron above 100°C have been examined for contribution of determinate errors. The analysis leads to a recommended equation for the permeability of hydrogen in a iron: o= (2.9 ±0.5) x 10-3 exp - (8400 ± 400)/RT cu cm (ntp H2) cm-1 sec-1 atm-1/2 THE permeability of a iron to hydrogen has been the subject of numerous investigations over the past 40 years, and at present there are thirteen sets of published results for the rate of steady-state permeation of hydrogen in a iron above 100°C. The numerical values in each set of results are entirely self-consis-tent, but the spread among the sets is too large to be attributed solely to experimental error, i.e., to error other than in the specimen itself. Several reasons have been advanced to explain the disparities, but to date the relative importance of experimental inaccuracy to the spread remains uncertain. The purpose of this report is to examine in detail the sources of determinate errors inherent in the experiments and to assess as far as possible the contribution of the errors to the results. The ultimate goal is the selection of values for the permeability and heat of permeation most nearly representative of hydrogen in a iron. The analysis is limited to those experiments in which the permeation rate was observed at steady state—a condition in which traps for hydrogen within the metal are filled to a fixed level15 so that the trapping mechanism is not reflected in the rate of passage of the gas. Furthermore only data are examined in which surface processes are judged to have little or no influence on the flow. It is hoped with these restrictions to obtain values of the permeability and the heat of permeation which will be as closely related as possible to the mechanism of lattice diffusion. I) DEFINITION OF TERMS; UNITS In this report the data for permeation are given in terms of a coefficient oj permeability, ?, which is defined by the equation: jt=?A/?x{p1/2-po1/2) [1] where jt is the total flow of gas normal to the surface of a membrane of planar geometry, e.g., a disc, of area A and thickness ?x; pi and po are the pressures in the input and output sides, respectively. For flow radial to the walls of a membrane of cylindrical geometry, e.g., a tube, the corresponding equation is: where 1 is the length of the cylinder, and ri and ro are the inner and outer radii, respectively. The flux normal to the surface is given by Fick&apos;s law: j= -D(dc/dx) [3] At steady state the concentration gradient will be constant, and integration of Eq. [3] gives for the total flow through a disc of area A and thickness Ax: h =-DA(co - ci) [4] where c, and ci are the concentrations of solute at the output and input surfaces, respectively. When surface control is absent, co and ci are given by Sievert&apos;s law c = Kp1/2, and substitution therewith into Eq. [4] gives directly Eq. [I] where ? = DK. Integration of Fick&apos;s Eq. [3] in cylindrical coordinates will give Eq. [2] where again ? = DK and is thus shown to be independent of geometry (provided that surface control is negligible). The coefficient of permeability, or simply the permeability,* must be expressed in proper units. In *The term permeability will refer in this report always to the coefficient defined above; permeation will be used to specify the general phenomenon of gas passage through a membrane. this report ? will be expressed in the units of cu cm (ntp H2) cm-1 sec-1 atm-1/2. The variations of D and K with temperature are given by D = Do exp(-Ea/RT) and K = KO exp(-?Hs/RT) where E, is the activation energy for diffusion and AH, the heat of solution, each usually expressed in calories per mole of solute. The variation of permeability with temperature will thus be given (for conditions where surface control is negligible) by ? = ?o exp(-?Hp/RT) where ?0 = DoKo and ?Hp = Ea + ?Hs. The units of ?0 are the same as those of 6, and??Hp will be expressed in calories per mole H. 11) SUMMARY OF PERMEABILITY RESULTS Table I gives the values reported to date for the permeability of H2 in a iron in terms of ?o and ?Hp. Except where noted the parameters listed were taken directly from the numbers reported by the various investigators with only a change in units. The temperature limits within which the listed ?o and ?Hp hold are given in column 7; the limits marked in parentheses in this column indicate the entire temperature range covered in each investigation. The listed values of ?o and ?Hp are those giving a linear plot of ln? against T-1 at the higher temperatures in each set of measurements, and thus presumably represent the case for which surface control was negligible. Column 6 gives values of 9 at a representative tem-

Jan 1, 1970

By R. F. Hehemann, S. T. Zegler

The w phase in zirconium alloys containing more than G pct Nb can form in a difjUsionless manner during quenching or with composition change during aging at temperatures below 550°C. The latter treatment establishes a metastable equilibrium between an w phase containing 5 pct Nb and a ß phase containing 45 pct Nb. The superconducting transition temperat~ire of quenched Z?-10 pct Nb and Zr-15 pct Nb alloys is increased by aging at 375°C and both alloys exhibit the same Tc after prolonged aging. The superconducting transition both in the quenched and in the aged alloys is relatively sharp (occurring ovev a temperature range of <0.5°K), and is interpreted in terms of the microstructure that characterizes the ß-w) mixture and proximity effects postulated for thin-film superconductors. At elevated temperatures alloys of the transition elements from groups 4 to 6 in the periodic table exhibit wide ranges of mutual solubility in the bcc structure. Except for compositions that are high in a group 4 element? the bcc structure can be retained at room temperature by quenching and several systematic studies of the influence of composition on superconductivity in bcc alloys have been reported.&apos;&apos;2 The high-temperature bcc form (ß) of the group 4 elements—titanium, zirconium, hafnium—is not retained upon quenching but instead transforms marten-sitically to the hcp structure (a). The Ms temperature can be lowered by alloying the group 4 metals with a solute from a higher group and with a critical amount of solute the martensitic transformation can be suppressed. The 3 phase, however, is not retained without modification. Rather, a metastable phase, w (in particular in titanium- and zirconium-base alloys), occurs over a fairly wide range of compositions above that required for suppression of the martensitic transformation.3,4 As in the formation of a, the transformation of ß to w occurs in a diffusionless manner during quenching and the temperature for its initiation (ws) is lowered by increasing solute content. The ws temperature can be well below room temperature and reversibility of the transformation allows w that forms on cooling to low temperatures to disappear again when the sample is returned to room temperature.5 In Zr-Nb alloys the w phase can also be formed during aging at temperatures below 550°c.5 In this treatment, a metastable equilibrium is established between a niobium-rich ß phase and a zirconium-rich w phase. The present study of Zr-Nb alloys was undertaken to examine the influence of microstructure: as controlled by composition and heat treatment, on the superconducting transition temperature, Tc. MATERIALS AND PROCEDURE The alloys were prepared from zirconium crystal bar that was spectrochemically pure except for 30 ppm of Al, 40 ppm of Cu, 100 ppm of Fe, and 1 ppm of each B and Pb. The niobium used was 99.5+ pct pure containing according to the suppliers&apos; (Stauffer Chemical Co.) analysis a maximum of 50 ppmO, 60 ppm of N, 30 ppm of C, 5 ppm of H, 200 ppm of Ta, 50 ppm of Mo, and 15 ppm each of Fe, Co, and Ni. The alloys were prepared in the form of buttons weighing approximately 4 g by arc melting on a water-cooled copper hearth in an inert-gas atmosphere. All alloys were heat-treated in the ß-phase region for 3 days at 1000°C and water-quenched. For heat treatment, the specimens, 0.5 cm in diam by 1.0 cm long, were wrapped in zirconium foil and sealed in an inert-gas atmosphere inside quartz capsules. Quenching was done by breaking the capsules under water. Subsequent aging treatments of Zr-10 pct Nb and Zr-15 pct Nb alloys were done by placing the specimens in a lead bath at 375°C and quenching in chilled brine. The phases present in the quenched and aged alloys were identified by optical and X-ray metallography. X-ray powder patterns were taken at room temperature with a 114.6-mm-diam Debye-Scherrer camera and CuK-a radiation. Superconducting transition temperatures were determined from magnetic-permeability measurements6 in a 10 oe field. The superconducting transition temperature was determined from a plot of galvanometric deflection as a function of temperature by extrapolation of the linear portion of the curve to zero deflection. In Figs. 1 and 3 the arrows indicate the temperature range of the transitions. RESULTS AND DISCUSSION Superconductivity in Quenched Alloys. In Fig. 1 Tc for alloys quenched from 1000°C is plotted against nominal composition. Indicated also in Fig. 1 are results of metallographic studies. Alloys containing less than 5 pct Nb transform martensitically to a during quenching from 1000°C and Fig. 1 demonstrates that Tc increases significantly as the niobium concentration in a is increased. On the basis of the Bardeen-Cooper-Schreiffer theory of superconductivity7 it seems most likely that this increase in Tc results from an increase in the density of electronic states as zirconium is alloyed with niobium. This increased density of states has been revealed clearly in low-temperature specific-heat measurements.8 The martensitic transformation to a is suppressed when the alloy content exceeds 5 pct Nb. In the range from approximately 6 to 30 pct Nb, quenched alloys consist of a mixture of B and w phases and the amount

Jan 1, 1967

By G. R. Leverant, M. Gell

The high- and low-cycle fatigue poperties of MAR-M200 directionally -solidified into columnar-grained and single crystal forms were determined at 1400" and 1700°F. These results were compared with the corresponding properties of conventionally -cast MAR-M200. The low-cycle fatigue lives of the columnar-grained and single crystal materials were similar at both temperatures and were one to two orders of magnitude greater than those of conventionally-cast material. The variations in the fatigue lives among the three forms of MAR-M200 were related to the more rapid rate of intergranular muck propagation compared to that of transgranular propagation. In conventionally-cast MAR-M200, cracks were initiated in grain boundaries and crack popagation occurred rapidly along an almost continuous grain boundary path. In the columnar-grained material, crack initiation occurred on short transverse segments of grain boundaries, but crack propagation was transgranular. The fatigue lives of columnar-grained and single crystal materials were approximately the same because most of the life in both materials was spent in trans-granular propagation. For the directionally -solidified materials, the number of cycles to failure, Nf, can be related to the total strain range, , by: heT = K where n and K are 0.16 and 0.044 at 1400&apos;F and 0.29and 0.098 at 1700, respectively. In addition to intergranular crack initiation in the columnar-grained material, initiation also occurred at we-cracked MC carbides and micropores in both directionally-solidified materials. At 1400°F, fatigue life was reduced with increasing MC carbide size, but at 1700 there was no effect of carbide size. THE creep and stress-rupture properties of conventionally-cast nickel-base superalloys can be greatly improved by directional solidification into either single crystal or columnar-grained forms. The improvement in properties results from a reduction in intergranular cavitation and crack growth in the columnar-grained materials and the complete absence of this fracture mode in the single crystals. This paper describes the effect of grain boundaries on the elevated temperature fatigue properties of the nickel-base superalloy MAR-M200. The effect of cycling frequency on cracking arid fatigue life and the role of MC carbides and micropores on crack initiation are also emphasized. The fatigue properties of columnar-grained and single crystal MAR-M200 at room tem- perature,4,5 and the change in the mode of fatigue crack propagation with temperature6 have recently been described. I) EXPERIMENTAL PROCEDURE The material used in this study was the nickel-base superalloy MAR-M200, directionally-solidified into columnar-grained and single crystal forms. The columnar grains were approximately 0.5 mm in diam. The nominal composition of these materials in wt pct was 0.15C, 9Cr, 12.5W, loco, 5A1. 2Ti, lCb, 0.05Zr. 0.015B, bal Ni. They were solutionized for 1 to 4 hr at 2250°F followed by aging at 1600°F for 32 hr. which resulted in 0.2 pct offset yield stresses of 150,000? 144,000, and 95,000 psi at room temperature? 1400°, and 1700°F, respectively. The corresponding elastic moduli parallel to the testing direction were 19.2. 15.0, and 12.5 x 106 psi, respectively. Specimen design, testing procedures and alloy mi-crostructure have been described previously and will only be summarized here. Following the 1600°F aging treatment. MAR-M200 contains an ordered, cuboidal, y&apos; precipitate 0.3 1 on edge, which is coherent with the 1 matrix. The y&apos; precipitate is quite stable; even after testing at 1700°F. the precipitate is only slightly enlarged and its corners somewhat rounded. The alloy also contains a small volume fraction of micropores, and MC carbides. some of which contain preexisting cracks5 formed during casting. These cracks are always parallel to the long dimension of the carbide. Measurement of MC carbide size has been described previously.5 Axial fatigue tests were conducted in air over a wide range of strain amplitudes in both the high - and low-cycle fatigue regions, with specimen lives varying from about 10&apos; to 10&apos; cycles. Low-cycle fatigue (LCF) tests were strain-controlled with strain varied between zero and a maximum tensile value at a frequency of about 2 cpm. High-cycle fatigue (HCF) tests were stress-controlled with the stress varied between 5000 psi and a maximum tensile value less than the yield stress at a frequency of either 10 cps or 0.033 cps (2 cpm). The temperature in the gage section was controlled to 52°F. Specimen axes were within 5 deg of the [001] growth axis of the single crystals and the common [001 ] growth axis of the columnar-grained material. Specimen gage sections were electro polished prior to testing. After the standard heat treatment, three specimens were coated with a typical aluminide coating applied as a slurry. An additional specimen was given the coating heat treatment without actually being coated. In all cases: specimens were reaged at 1600°F for 32 hr after receiving the coating heat treatment at 1975&apos;F for 4 hr.

Jan 1, 1970

By Y. C. Liu, G. A. Alers

measurements of the anisotropy in Young&apos;s modulus produced in copper by rolling 95 pct reduction in thickness below room temperature have been carried out in order to study the dependence of the texture on rolling temperature. The results clearly show the transition from a copper-type texture to a brass-type texture as the temperature of rolling is lowered. The intermediate textures observed can be described very well as a simple mixture of the two terminal textures. These results cormbined with other texture measurements make possible afresh review of the experimental facts velating to rolling textures in fee metals and, as a consequence, a critical examination of the current theories is presented. PREVIOUS experiments have shown that the transition from the copper- to the brass-type rolling texture is clearly displayed and can be quantitatively analyzed by measurements of the anisotropy of Young&apos;s modulus.&apos; Application of this method to the Cu-Zn alloy system showed that the description of the texture transition as a gradual rotation of the grains from the orientation characteristic of the copper texture to the {110}(112) texture of brass2 was inconsistent with the data. Instead, the data suggested that the texture within the transition region could be described as a simple mixture of the two terminal textures.5 Unfortunately, it was difficult to establish this point conclusively because of the inadequacy of corrections for the composition dependence of the single-crystal elastic constants. Since a rigorous establishment of the nature of this texture transition is essential to our understanding of the formation of rolling textures in fee metals, it is clearly important to undertake an investigation in which the composition dependence of the elastic constants would not enter. A suitable composition-independent texture transition is provided by the well-established variation in the rolling texture of copper with rolling temperature. This temperature-dependent texture transformation has been studied by smallman&apos; in several fee alloys and by Müller5 and others"&apos; in copper. They observed that the texture characteristic of copper rolled at room temperature changed to a brass-type texture when the rolling temperature was lowered to 77°K. Although it is not possible to decide unequivocally from the published pole figures whether or not the 77°K rolling texture of copper is entirely of the brass type,&apos; this complication does not affect the main purpose of the present investigation. In addition to establishing the nature of the texture within the transition region, the modulus data should also provide a determination of the temperature at which the transition occurs as well as the temperature range over which the transition extends. This information when combined with the modulus data on Cu-Zn alloys would then provide a considerable body of new information on textures in fee metals. Since these modulus results and the data obtained from pole-figure studies must be internally consistent, it is appropriate to compile a brief summary of the experimental observations based on all available methods rather than on the pole-figure data alone as has been done in the past. The primary purpose of such a summary would be to yield a more precise definition of the experimental facts on the rolling textures of fee metals, and thus greatly facilitate our evaluation of various proposed theories in this field. The final section of this paper is devoted to this compilation of consistent, experimental facts and their application to the various theories. EXPERIMENTAL PROCEDURE Two 18-lb ingots of cathode copper of 99.99 pct purity were induction-melted under a nitrogen atmosphere in a graphite crucible and chill-cast into a steel mold. The ingots were repeatedly cold-rolled and annealed (I hr at 500°C) into slabs about 1 1/8 in. thick. Blocks 3 1/4 in. wide, 2 1/4 in. long, and 1.000 in. thick were machined from each slab. The rolling schedule used was the same as in the previous investigation1 and the final thickness of the sheet was 0.050 in. with a rolling reduction of thickness of 95 pct instead of 97.5 pct as in the previous work.&apos; The compositions and temperatures of the cold baths used for the low-temperature rolling were as shown in Table I. After each pass the rolled strip was immediately immersed in the cold bath for about 1 min or until the bubbling of the bath had subsided. The modulus data were taken within 2 hr after the rolled strip was warmed to room temperature for the first time, so that effects due to recrystallization were minimized. The modulus specimens were in the shape of flat bars, 3 in. long, 4 in. wide, and 0.050 in. thick, cut with their long dimensions oriented at 15-deg intervals between the rolling direction and the transverse direction. The values of Young&apos;s modulus were deduced from measurements of the frequency at which these long narrow bars were set into longitudinal, resonant vibration as previously described.9 To excite the mechanical vibrations in the specimen, an electromagnetic drive similar to that employed by Thompson and lass" was used. The maximum in the amplitude of

Jan 1, 1968

By C. L. Sollenbeger, R. B. Greenwalt

PERHAPS the most widely used dispersants or gangue depressants in nonmetallic flotation are sodium silicates, which vary in silica-to-soda ratio from 1 to 3.75. Typical manufactured silicates in order of decreasing solubility and increasing amounts of silica are Metso, silica-to-soda ratio of 1.00; D, 2.00; RU, 2.40; K, 2.90; N, 3.22; and S-35, 3.75.* References in flotation literature1,2 to the use of sodium silicates are often weak because they fail to mention the type of silicate used. Metso and silicate N have occasionally been mentioned, but when the type of silicate is not mentioned, it is usually assumed to be N, the cheapest of the soluble silicates and the one recommended by sodium silicate manufacturers as a flotation agent. In the All is-Chalmers Research Laboratories a systematic study was made of the effect of different alkali-silica ratios on the concentration by flotation of two scheelite ores. One of these was a high grade ore from the Sang Dong mine in Korea. The effect of such factors as pH; addition agents; and conditioning time, temperature, and pulp density on the flotation efficiency of this ore have been described previously. The other ore was a low grade ore from Getchell Mines Inc., Nevada. The mineralogy and techniques of concentrating this ore have been described by Kunze. Hereafter these ores will be referred to as the Korean and Nevada ores. Experiments were made with both to determine the effect of three factors—-type of silicate, concentration of silicate, and pH of the pulp—on recovery and grade of tungsten in a rougher concentrate. Average WO, content of the Korean ore was 1.50 pct and of the Nevada ore 0.27 pct. The predominant tungsten mineral in both ores was scheelite, which was accompanied by a small amount of powellite. The powellite and scheelite were finely disseminated through both ores and required a —200 mesh grind for liberation. Major gangue minerals in the Korean ore, in decreasing order of abundance, were amphi-boles, quartz, biotite, garnet, fluorite, and calcite. Bulk sulfides composed about 3 pct of the total weight. Gangue in the Nevada ore, in descending order of abundance, was garnet, alpha quartz, calcite, phlogopite, wollastonite, and amphiboles. Sulfide minerals were 3 to 4 pct of total weight. Batch flotation experiments were made with 500-g samples of ore, each sample wet-ground to 90 pct passing 200 mesh. The finely ground ore was floated in a Fagergren batch cell at 25 pct solids. The natural pH of the Nevada ore was 8.9 and of the Korean ore, 8.5. The D, RU, K, N, and S-35 sodium silicates were obtained in colloidal dispersions with varying amounts of water. The most alkaline, Metso, was in dry powdered form. For convenience in addition, 5 pct solutions by weight were prepared from each of the silicates, on the basis of dry sodium silicate dissolved in the correct amount of distilled water. Chemical analyses of the various silicates are given in Table I, together with the pH of the 5 pct solutions. A preliminary bulk sulfide float was made with secondary butyl xanthate as the collector and pine oil as the frother. The WO] analysis of the sulfide concentrate was nearly 1 pct for the Korean ore and about 0.1 pct for the Nevada ore. The tungsten contained in the sulfide concentrate constituted about 3 pct of the total tungsten in each ore. No effort was made to recover these tungsten values. The scheelite was floated with oleic acid. Adjustments in pH were made with sulfuric acid or sodium carbonate. A 1 pct solution of 85 pct Aerosol OT was sprayed on the froth and sides of the cell during the scheelite float to aid in dispersing the minerals and to decrease the entrapment of gangue particles. Six tests were planned for each of the six types of silicate in which concentrations of 1, 2, and 4 1b of silicate per ton of dry ore were investigated at both 6.5 and 10 pH. All tests were made at room temperature. The performance of each silicate was judged from the grade and recovery of WO, in the scheelite rougher concentrate. Tungsten recovery was calculated on the basis of the scheelite remaining in the ore after the preliminary sulfide float. Testing of each silicate at three levels of concentration and two levels of pH required 36 tests with each scheelite ore. Variance analyses were performed on the concentrate grades and recoveries to determine whether or not the type of sodium silicate, the concentration of sodium silicate, or the pH significantly affected recovery or grade. Results Concentrate Grade: A variance analysis of the concentrate grades for the Korean ore showed that concentration of the silicate and pH of the ore pulp were major factors in producing a high grade concentrate. Also, the silica- to-so da ratio was important as an interaction with pH. The concentrate grade vs silica-to-soda ratio is plotted in Fig. 1. The curves show that the concentrate grade improved with an increase in concentration of sodium silicate and also

Jan 1, 1959

By E. W. Davis

THIS nation has been depending upon the Lake Superior iron ranges for most of its iron ore requirements for over half a century. Furthermore, it can continue to draw the major portion of its ore requirements from the same source for another century or more, but to accomplish this, great changes must take place in the present ore mining industry. Preliminary preparations for these changes are now underway and in the next few years decisions of far-reaching importance must be made. But to understand the situation, it is first necessary to know something of the nature and occurrence of the Lake Superior iron formation. So far as is known, the merchantable grades of iron ore remaining in the various Lake Superior ranges are as shown in Table 1, which shows that

Jan 1, 1947

By G. W. Brownell, H. T. F. Lundberg, R. W. Pringle, K. I. Roulston

The rapid improvement of the airborne scintillometer and the perfection of its efficiency for counting low energy gamma radiation has made it possible to work out a technique to map in great detail the radiation pattern at the earth&apos;s surface. On such maps low radiation over certain areas appears to indicate the existence of oil accumulations, forming a pattern similar to that obtained by the geo-chemists. RADIOACTIVE analyses of samples from the surface of oil fields were carried out more than 10 years ago in Alberta by the alpha particle ioniza-tion chamber technique,&apos; but large enough tracts could not be covered in these investigations to make possible any evaluation of the method as a means of oil exploration. Considerable interest has recently been revived, however, as a result of certain striking advances which have been made in the instrumentation available -for the measurement of radioactivity. It is the object of this paper to indicate the nature of these improvements in radiation technology and then to describe the attempts that have been made to interpret the radioactive patterns obtained in the course of airborne recordings with the new instruments. Since the survey can be carried out from the air and records can be accumulated over vast areas in a short time, the result may easily lend itself to statistical treatment. Areas have been surveyed in Alberta, British Columbia, Saskatchewan, Quebec, Texas, New Mexico, Nebraska, Colorado, Utah, and Montana. Producing fields in Alberta and West Texas have been flown over several times in different directions, Fig. 1. The operations were then extended into unknown territory and drill holes were put down on the anomalies which looked promising. The results from these drillings were encouraging and have given hopes for the development of an entirely new method of oil exploration. Any large scale method for the survey of radioactive anomalies must be based on the measurement of gamma rays, as beta and alpha rays have much too short a range to be of any significance. Thus the essential improvement which has made the present stage of this work attainable is the development of new highly sensitive detectors for gamma radiation. In the past the only detectors of any consequence that were available were the ionization chamber and the geiger counter, but both of these suffer from the defect that only a small proportion of the gamma rays passing through the counter are detected, possibly 0.1 to 0.2 pct. The recent development of the scintillation counter2,3 has completely transformed the situation and has had a considerable impact on many branches of nuclear technology. The detection of alpha particles in zinc sulphide screens by visual observation of the individual scintillations which these particles produce dates back to the early spinthariscope of Rutherford and Crookes, but the combined use of an appropriate scintillating phosphor and photomultiplier tube had to await the technical development of the latter many years later. With this development came the modern era of the scintillation counter and a knowledge of phosphors which have a large light output under the bombarding action of gamma radiation. Some of these phosphors are relatively dense and are capable of stopping a large proportion of the incident gamma radiation. As the sensitive region is the whole volume of the crystal, a very high detection efficiency, 50 pct or more, can be obtained for medium energy gamma rays. Scintillation counters for geological purposes were first developed in 19494-6 in an attempt to utilize this remarkable improvement in efficiency, which has the attractive consequence that only a small portion of the normal background of the counter is due to cosmic radiation. In 1949 tests were made in northern Saskatchewan by Lundberg Explorations Ltd. with portable scintillation counters which gave excellent results in the search for uranium and served to indicate unknown uranium deposits in areas previously closely surveyed with geiger counters. Portable scintillometers (registered in Canada) are now commercially available and in regular use,&apos; and the adaptation of the instrument to radioactivity oil well logging has also been very successful.8 Initial attempts to measure radioactivity from aircraft with scintillation counters were made during this period in the same area and yielded most encouraging results. It would be appropriate to consider some specific requirements for airborne investigations. The essential problem to be met in the detection of any radioactive source is the necessity of obtaining a signal greater than the statistical fluctuations of the background counting rate for the instrument. It is possible to show that Nt>2k2 is the condition for detectability of a signal where N = average background counting rate for the detection. t = time constant of the counting rate meter, used to determine the average number of counts arriving in a certain predetermined time interval. N&apos; = average source counting rate at the detector. k = N/N&apos;, and N>>N&apos;. Sample values are given in Table I. Assume that the aircraft carrying the equipment is travelling at 120 mph, in which case it will cover 176 ft in 1 sec. Assume also, as a first approximation, that a point source target is in range when the air-

Jan 1, 1954

By A. Raman, P. P. Singh

The constitution of the La-Rh system was studied by powder X-ray diffraction, metallopaphic, and differential thermal analysis techniques and an equilibrium diagram is presented. Eleven intermediate phases occur in the system and the crystal structural data for nine of them were determined. La3Rh crystallizes in an orthorhombic structure of undetermined type, whose unit cell is obtained by doubling the &apos;a; and &apos;c,,&apos; edges of an FesC type unit cell. The other intermediate phases of the system are LarRh-3( undetermined structures also occur in the system. LaRh, undergoes a polymorphic phase transformation at 1240°C. LaRh3 and La2Rh7 also exhibit polymorphisnz. The phases Laah and LazRh7 melt congruently. The latter undergoes a eutectoid transformation into LaRh, and Rh at 1205°C. Laah3 is formed by a peritectoid reaction between Laah and La,Rh,,. The other Phases result from peritectic reactions between the liquid and the adjacent rhodium-rich phases. The intermediate Phases of the La-Rh system are compared with those of the La-Co and La-Ni systems. DURING the course of a detailed investigation to study the occurrence of CrB, FeB, A1B2, and related structures in the rare earth alloys it was found that much information is lacking for the rare earth noble metal systems. Although the structures of several rare earth alloys containing the noble metals at the AB and AB2 stoichiometries have been reported, the occurrence of related structures at other stoichiometries has not been studied. We have initiated a project to study the crystal structural features of selected rare earth-rhodium alloys and to map the equilibrium diagrams of representative systems with conventional methods. The results of our investigation in the La-Rh system are presented in this paper. Two phases were known in the La-Rh system. LaRh has the CrB-type structure.&apos; LaRhz is a MgCu2-type Laves phase.z EXPERIMENTAL PROCEDURE Alloys weighing less than 1 g were prepared from commercially pure lanthanum (99.9 pct +), supplied by Lunex Company, Pleasant Valley, Iowa, and rhodium (99.92 pct +), supplied by Engelhardt Industries, Newark, N.J., in a conventional arc melting furnace under argon atmosphere. The buttons were turned upside down and remelted three times to insure homogeneity in the samples. Since negligible loss of material was encountered during melting, a chemical analysis of the alloy buttons was not undertaken. Powder specimens for X-ray diffraction studies in the as cast state were then prepared. The buttons were wrapped in thin molybdenum foils and homogenized by heating in vacuum at suitable high temperatures for more than 1 week. They were then broken into three or four pieces for annealing experiments. The pieces were wrapped in molybdenum foils and annealed at various temperatures in evacuated quartz capsules. The annealing was carried out for 2 hr at or above 1200°C, 1 day at temperatures close to llOO°C, 2 days at 1000°C, and for 1 week at temperatures below 1000°C. After annealing the alloy pieces were again broken and powder specimens for X-ray diffraction were prepared. The powders of the lanthanum rich alloys with more than 80 at. pct La were prepared by filing. The filings were sealed in molybdenum tubings and stress-relieved at 600°C in vacuum. It was not deemed necessary to stress-relieve the powders of the other alloys, since the alloys were very brittle and were ground easily. POWDER X-RAY DIFFRACTION X-ray diffraction photographs of powders (-325 mesh size) of the alloys in the as cast and annealed states were prepared in a Guinier-de Wolff focussing camera with copper K, X radiations. These patterns were studied to identify the stoichiometries and the crystal structures of the intermediate phases. The lattice parameters of the phases were calculated after minimizing the differences between the observed sin2 6 values, calculated from the diffraction angles 8, and the sin2 8 values, calculated using approximate lattice constants obtained from a few lines. These differences were minimized manually to less than 0.0005. The latLice constants are judged to be accurate to *0.005A for values less thp about 10A and to k0.01~ for values greater than 10A. The relative intensities of the lines were calculated using a computer program written by Jeitschko and Parthk.~ No attempt was made to refine the atomic positional parameters in the phases. METALLOGRAPHY The phase equilibria in the investigated alloys in the as cast and annealed states were also studied by metallographic examination. The polished specimen surfaces were etched with 10 pct picric acid in alcohol (alloys up to 25 pct Rh), concentrated picric acid (from 25 to 37.5 pct Rh), 2 pct nital (40 to 50 pct Rh), 10 pct nital (from 50 to 66.7 pct Rh) and with concentrated 48 pct HF for the other rhodium-rich alloys. Selected microstruture~ were then photographed using a Po-laroid Land camera. THERMAL ANALYSIS Differential thermal analysis of the alloys was carried out in DTA-668 Stone differential thermal ana-

Jan 1, 1970

By Mara Kosovac, Sujan K. Kundu

The Iron Ore Co. of Canada (IOC) has developed a computerized plan analysis method for its open-pit iron mining operations which will eliminate much of the tedious manual drafting of pit design plans performed by engineers and geologists at its two major operational areas near Schefferville, Que., and Labrador City, Nfld. The new technique permits analysis of section or bench plans for a specific ore body, based on drillhole information; preparation of a three-dimensional unit profit matrix for pit design purposes, and storage of extensive ore-block information for other applications. Other advantages include enhanced design accuracy through reduction of human errors normally associated with preparation of large data masses, and by elimination of errors inherent in planimeter-prepared ore block diagrams.

Jan 7, 1978

By Charles E. Lundin

The character of the Er-D system was established by determining pressure-temperature-composition relationships. A Sieuerts&apos; apparatus was employed to make measurements in the temperature range, 473" to 1223"K, the composition range of erbium to ErD3, and the pressure range of 10~s to 760 Torr. The system is characterized by three homogeneous phase regions: the nzetal-rich, the dideuteride, and the trideuteride phases. These phases and their solubility boundaries were deduced from the family of isotherms of the system zchich relate the pressure-temperature-composition variables. The equilibrium plateau decomposition relationships in the two-phase regions were determined from can&apos;t Hoff plots to be: The differential heats of reaction in these two regions are AH = - 53.0 * 0.2 and -20.0 *0.1 kcal per mole of D2, respecticely. The differential entropies of reaction are AS = - 36.3 * 0.2 and - 31.0 * 0.2 cal per mole D2. deg, respectively. Relative partial molal and intepal thermodynamic quantities were calculated from the pure metal to the dideuteride phase. The study of the Er-D system was undertaken as a logical complement to an earlier study of the Er-H system.&apos; The primary interest was to compare the characteristics of the two systems and relate the difference to the isotopic effect. Studies of rare earth-deuterium systems by other investigators have been very limited in number and scope. Furthermore, there is even less information available wherein an investigator has systematically compared a binary rare earth-hydrogen system with the corresponding rare earth-deuterium system. The available information consists primarily of dissociation pressure measurements in the plateau pressure region of a few rare earths. Warf and Korst&apos; determined dissociation pressure relationships for the La- and Ce-D systems in the plateau region and several isotherms for each system in the dideuteride region. They compared these data with those of the corresponding hydrided systems. The study of these systems as a whole was very cursory and did not give sufficient data for a thorough comparison of the effect of the hydrogen vs the deuterium in the respective rare earths. The heat capacities and related thermodynamic functions of the intermediate phases, YH, and YD2, were determined by Flotow, Osborne, and Otto,~ and the investigation was again repeated for YH3 and YD3 by Flotow, Osborne, Otto, and Abraham.4 This investigation studied only these specific phases. Jones, Southall, and Goodhead5 surveyed the hydrides and deu-terides of a series of rare earths for thermal stability including erbium. They experimentally determined isotherms of selected hydrides and plateau dissociation pressures for deuterides. These data allowed comparison of the enthalpy and entropies of formation of the dihydrides and dideuterides. To date, no one rare earth has been selected to thoroughly establish the complete pressure-temperature-composition (PTC) relationships of binary solute additions of hydrogen and deuterium, respectively. The objective in this investigation was to provide the first comparison of a complete family of isotherms of a rare earth-deuterium system with those of a rare earth-hydrogen system. This would allow one to determine what differences exist, if any, in the various phase boundaries and the thermodynamic relationships in various regions of the systems. I) EXPERIMENTAL PROCEDURE A Sieverts&apos; apparatus was employed to conduct the experimental measurements. Briefly, it consisted of a source of pure deuterium, a precision gas-measuring buret, a heated reaction chamber, a mercury manometer, and two McLeod gages (a CVC, GMl00A and a CVC, GM110). Pure deuterium was obtained by passing deuterium through a heated Pd-Ag thimble. A 100-ml precision gas buret graduated to 0.1-ml divisions was used to measure and admit deuterium to the reaction chamber. The reaction unit consisted of a quartz tube surrounded by a nichrome-wound furnace. The furnace temperature was controlled by a recorder-controller to . An independent measurement of the sample temperature in the quartz tube was made by means of a chromel-alumel thermocouple situated outside, but adjacent to, the quartz tube near the specimen. Pressure in the manometer range was measured to k0.5 Torr and in the McLeod range (10~4 to 10 Torr) to *3 pct. The deuterium compositions in erbium were calculated in terms of deuterium-to-erbium atomic ratio. These compositions were estimated to be *0.01 D/Er ratio. The erbium metal was obtained from the Lunex Co. in the form of sponge. The metal was nuclear grade with a purity of 99.9+ pct. The oxygen content was reported to be 340 ppm and the nitrogen not detectable. Metallographically the structure was almost free of second phase (<i vol pct). A quantity of sponge was arc-melted for use as charge material. The solid material was compared with the sponge in the PTC relationships. They were found to be identical. Therefore, sponge material was used henceforth, so that equilibrium could be attained more rapidly. The specimen size was about 0.2 gr for each loading of the reaction chamber. The procedure employed to obtain the PTC data was to develop experimentally a family of isothermal curves of composition vs pressure. First, a specimen

Jan 1, 1969

By E. S. Fisher, L. C. R. Alfred

The elastic anisotropy factors, c4,/c6,, c3,/cll, and c12/cl,, for hcp metal crystals vary significantly among the dgferent unalloyed metals. Significant variations with temperature are also found. The effects of elastic anisotropy on the dislocation in an elastic continuum with hexagonal symmetry have been investigated by computing the elasticity factors for the self-energies of dislocations in fourteen different metals at various temperatures where the elastic moduli have been reported. For most of the metals the effects of the orientation of the Burgers vector, dislocation line, and glide plane are small and isotropic conditions can be assumed without significant error. Significant effects of anisotropy are, however, found in Cd, Zn, Co, Tl, Ti, and Zr. The elasticity factors have been applied in the calculations of dislocation line tensions, the repulsive forces between partial dislocations, and the Peierls-Nabarro dislocation widths. It is predicted that the increase in elastic anisotropy with temperature in titanium and zirconium makes edge dislocations with (a), (a + c), and (c) Burgers vectors unstable in basal, pyramidal, and prism planes, respectively. The probability of stacking faults forming by dissociation of Shockley partials in basal planes also decreases with increasing c4,/c6, ratio, when the stacking fault energy is greater than 50 ergs per sq cm. The widths of screw dislocations with b = (a) in titanium and zirconium increase very significantly in prism planes and decrease in basal planes as c4,/c6, increases. The effects of elastic anisotropy on various dislocation properties in cubic crystals have received considerable attention during the past few years. In the case of cubic symmetry the departure from isotropic elasticity depends entirely on the shear modulus ratio, A = 2c4,/(cl, —c12); i.e., the medium is elastically isotropic when A = 1. Foreman1 showed that an increase in the ratio A produces a systematic lowering of the dislocation self-energy for a given orientation and Poisson&apos;s ratio. ~eutonico~, has shown that large anisotropy can have a marked effect on the formation of stacking faults by the splitting of glissile dislocations in (111) planes of fcc and (112) planes of bcc crystals. ~iteK&apos; made similar calculations for (110) planes of bcc metals. Both studies of bcc metals showed that the large A values encountered in the alkali metals tend to reduce the repulsive forces between Shockley partial dislocations. In fcc metals, however, A does not vary over the large range encountered in bcc metals; consequently, the effect of A on the forces between Shockley partials is masked somewhat by the differences in Poisson&apos;s ratio between metals. The effect of A on the line tension of a bowed out pinned dislocation has also been investigated for cubic crystals, first by dewit and Koehler5 and more recent- ly by Head.6 In both cases the line energy model is applied and the core energy is not taken into account, thus making the conclusions somewhat tenuous with regard to the physical interpretation. Nevertheless, the fact that a large A decreases the effective line tension is clearly evident and the tendency for large A to produce conditions that make a straight dislocation unstable (negative line tensions) also seem evident. Head, in fact, shows visual microscopic evidence that stable V-shaped dislocations occur in 0 brasse6 For hcp metals the definition of elastic anisotropy is more complex and, furthermore, significant deviations from an isotropic continuum are found among a number of real hcp metals, especially at higher temperatures. The present work was carried out to survey the effects of elastic anisotropy on the elasticity factors, K, that enter into the calculations of the stress fields around a dislocation core. Some isolated analytical calculations have previously been carried out for several hcp metals but they are restricted in the dislocation orientations and temperature.8&apos;9 The present computations are based on single-crystal elastic moduli that have appeared in the literature and consider various orientations requiring numerical computations. The results are then applied to survey the effects of temperature on the dislocation line tension and dislocation splitting in hcp metals. PROCEDURE Anisotropy Factors. The degree of elastic anisotropy in hcp crystals cannot be described by a single parameter, such as the A ratio in cubic crystals. The following three ratios must be simultaneously equal to unity in order to have an elastically isotropic hexagonal crystal: The magnitudes of these ratios at several temperatures, as computed from the existing data for the elastic moduli of unalloyed hcp metals, are given in Table I. There are no cases of complete elastic isotropy, but the large anisotropy ratios encountered in the cubic alkali metals are also missing. There are, however, several significant differences among the hcp metals, the most notable being the relatively small A and B ratios in zinc and cadmium and the differences in the magnitudes and temperature dependences of A. It has been noted that the temperature dependence of A has a consistent relationship to the occurrence of the hcp — bcc tran~formation. For cadmium, zinc, magnesium, rhenium, and ruthenium, A is less than unity at 4&apos;~ and, with exception for rhenium, decreases with increasing temperature. In the case of rhenium, A has essentially no temperature dependence between 923&apos; and 1123"~, so that it is clear that A does not approach unity at higher temperatures. Cobalt is similar to the above-mentioned group of metals in that it also does

Jan 1, 1969