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By Robert Baschwitz, John Chipman

The distribution of silicon between liquid silver and Fe-Si-C alloys has been studied at 1420oand 1530°C. The data are consistent with earlier studies. New data of Hager on the liquidus lines of the system Ag-Si and the distribution data are used to obtain the activity coefficient of silicon in both liquid phases. Data on the heat of mixing in iron permit accurate extension to 1600°C. Equilibrium data involving SiO2 and silicon in liquid iron together with revised data on the free energy of SiO2 are used to fix the activity of silicon in the infinitely dilute solution. The binary system exhibits strong negative deviation from ideality. At infinite dilution ? Si at 1600" is 1.25 x 10'3, and at concentrations up to NSi = 0.4 the slope d InySi/dNSi has a constant value of r; = 13. It is found that logysi in the ternary solutzon is approximately but not exactly the same function of Nsi + NC as of NSi in the binary. The results are consistent with currently available data on the free energy of Sic and its solubility in molten iron. LIQUID solutions of the system Fe-Si-C have acquired considerable importance as the laboratory prototypes of blast furnace hot metal. Equilibrium studies involving such solutions and slags approximating those of the blast furnace have yielded useful information concerning the thermodynamic properties of blast furnace slags. In studies of this kind great importance attaches to a knowledge of the thermodynamic activity of silicon in the solution as a function of temperature and composition. An attempt was made by Chipman, Fulton, Gokcen, and askey' to evaluate all of the pertinent data on this system and to deduce the desired relation between activity, composition, and temperature. These authors published data on the solubility of graphite and Sic in molten Fe-C-Si solutions and on the distribution of silicon between liquid iron and liquid silver. They showed further how the activity of silicon in very dilute solutions in liquid iron could be calculated from equilibrium data involving the molten alloy and solid SiO,. These calculations rested on the published thermodynamic properties of SiO, in- cluding its heat of formation which at that time was recorded as -209.8 kcal. This value has been under suspicion for some time and has recently been replaced by the concordant results from two independent laboratories2,3 which place the heat of formation of a-quartz at -217.6 kcal. This revision necessitates a re-evaluation not only of the activity of SiO2 in slag but also of silicon in molten iron. It is the purpose of this paper, therefore, to recalculate the activity of silicon, and in furtherance of this objective to present new data on its distribution between liquid Fe-Si-C alloys and liquid silver. HEAT OF SOLUTION OF SILICON IN IRON In order to determine the effect of temperature upon the activity coefficient it is necessary to know the heat of solution of silicon in iron as a function of composition. This is found in the data of Korber and Oelsen4 shown in Fig. 1. The curve corresponds to the following equation, which is of a form suggested by Wagner:5 Here AH is the heat absorbed in kilocalories in forming one gram atom of molten alloy from its molten elements and the N's are atom fractions. The relative partial molal enthalpies of the components, each referred to its pureliquid state and defined as zFe = aFe - PFe and zsi = HSi — -psi, are shown graphically. At low concentrations zSi = -28.5 kcal, in agreement with Kijrber and Oelsen's computation. This is in good agreement with the value of -29.3 kcal obtained by Chipman and Grant6 using an entirely different method. ACTIVITY AT INFINITE DILUTION From the known free energy of SiO, it is possible to obtain the activity of silicon in dilute solution in liquid iron from equilibrium studies. The heat of formation of a-quartz is —217.6 kcal and the heat capacity and entropy data are given by Kelley and ~ing.' The free energy of formation of ß-cristo-balite at temperatures above the melting point of silicon is expressed by the following equation: Si(Z) + O2(g) = SiO2 (crist); AF" =-226,500 + 47.50T [I] The value of the deoxidation product for silicon [%Si] x [%O]2 at 1600°C according to Gokcen and chipmans is 2.8 x 10"5, in agreement with results of Hilty and Crafts.9 More recent works of Matoba,

Jan 1, 1963

By J. P. Zannaras

Referring to the article by R. J. Charles and P. L. de Bruyn, let us assume that W = weight of glass bar; P = weight of hammer; e = total deformation; K = unit of deformation; K = potential stress energy; E = modulus of elasticity; L = length of the bar; 7 = coefficient of inertia; h = height, ft; V = velocity, fps; and a = cross section area. The portion of kinetic energy which is effective in producing stress energy in a fixed bar struck horizontally is given by the formula" 1 + 1/3 W/P K =1+1/3w/p/(1+1/2w/p)2 P.V2/2g = Ph where 1 + 1/3 W/P ? (1 + 3W/P/(1=1/2w/p)2 [8] Putting e e = W/P =------------ From the above equation it can be seen that the maximum transfer of kinetic energy to stress energy is when e = 0 or W/P = 0 which indicates that the weight of the hammer must be very large as compared with the weight of the impacted rod. Eq. 8 diametrically opposes the conclusions reached by the authors of this article. In fact, if their suggestions were followed to the extreme when e = co when P = 0, there would be no transfer of kinetic energy to stress energy at all, as 7, becomes zero. Eq. 8 presumes that the velocity with which the stress is propagated through the bar is infinite, whereas the authors claim that the compression waves reflected are reaching the struck end of the bar prior to the complete transfer of the kinetic energy to cause such modification of the conditions there as to make them reach the reverse conclusions demonstrated by the above formula. That such interference exists is unquestionably demonstrated by the authors and others. However, if my observations are correct, such interference for this specific experiment and also for practical comminu- tion is insignificant, and the conclusions of the authors are in error and must be reversed to comply with Eq. 8. Eq. 5, w. = AE 2/2, given by the authors on page 51, is derived from the following equation (Eq. 9): K = 1/2Pe, where P = Sa, S = ?E, e = EL, and L = 1. The above formula, Eq. 9, cannot be applied in this case. This formula is applicable for static loads where the load increases from zero up to its final value, P, in such a way that the deformation at different instants is proportional to the loads acting at those instants and actually represents the area of a right triangle in the strain load diagram of base e and height P. The typical photographs shown in Figs. 3 and 4 represent the familiar strain load diagrams, and since the line of the wave marks the existence and intensity of the strain with the unquestionable conclusion that such strain has been caused by the action of a load acting continuously all along the wave until it reached the horizontal axis, the work stored at this point is represented by the area under the wave line and the horizontal axis and not by the area of the fictitious triangle given by the authors. Then if this is correct, even visual estimation of these areas at gage stations given in the typical photographs of Figs. 3 and 4 suffice to contradict the authors' calculation given in Figs. 6a and 6b and Figs. 7a and 7b. The typical photographs presented by Charles and de Bruyn show a considerable variation of the intensity of the strain at different stations but very small variation of areas which actually represent the stress energy at the corresponding stations. And, apparently, by squaring the small quantities, the authors magnified their error tenfold. J. M. Frankland's paperV iscusses the relative strain intensity and not the total energy for different types of impact loading. He states in his paper, "The reader is explicitly warned not to confuse the results in this report with those obtained when the load is applied by a blow as from a hammer. In this case the peak load rises to very large but mostly unknown values. The accompanying large deflections and stresses are the result of high values of P, not of the dynamic load factor n. According to Frankland "the dynamic load factor" is the numerical maximum of the response factor. It therefore appears that the authors followed the same procedure in obtaining the relative strain energy ab-

Jan 1, 1957

By H. C. Gatos, A. D. Kurtz

WITH the growing interest in the mechanism of self-diffusion of metals, the study of accurate and convenient methods for determining self-diffu-sion coefficients appears highly desirable. It was with this objective in mind that the present investigation was undertaken. Gatos and Azzam1 employed an autoradiographic technique for measuring self-diffusion coefficients of gold. This method involved sectioning of the specimen through the diffusion zone and recording the radioactivity directly on a photographic film. Because of the very short range of the emitted ß rays in gold, the activity recorded on the film was essentially the true surface activity. With proper choice of the sectioning angle, sufficient resolution could be obtained and the entire concentration-distance curve recorded in one measurement. For the boundary conditions of the experiment, where an infinitesimally thin layer of radioactive material diffuses in positive and negative directions into the end faces of a rod of infinite length, the solution of the diffusion equation is C/Cn = 1/v4pDt exp (-x2/4Dt) where C is the concentration of diffusing element (photographic density in this case), C,, is the constant (depending upon amount of radioactive material), x is the diffusion distance, D is the diffusion coefficient, and t is the time. Thus, by plotting the logarithm of the concentration vs the square of the diffusion distance, a straight line results and the slope contains the diffusion coefficient. In this manner, the self-diffusion coefficient of gold can be obtained as a function of temperature. In the present investigation the results reported by Gatos and Azzam1 have been verified, and the autoradiographic technique has been further developed and applied for the determination of the self-diffusion coefficient of gold at a number of temperatures. Furthermore, the energy of activation for the self-diffusion of gold has been conveniently determined. . Experimental Techniques Preparation of Specimens: The inert gold of high purity was received in the form of a rod from which cylinders were cut and machined to a diameter of 0.500 in. The specimens were annealed to a suitably large grain size and the faces were surface ground prior to the deposition of the radioactive layer. The radioactive isotope Au198 was chosen. It was produced in the Brookhaven pile by means of the reaction Au197 + n ? Au108. It decays by ß emission (0.96 mev) to Hg108 with the subsequent emission of a y ray (0.41 mev). 70Au 108 ? 80Hg 108 + -1e°. The half life of the Au108 is 2.7 days so that a strict time schedule had to be maintained in order to secure sufficient activity until the end of the experiments. For this reason, initial activities as high as 10,000 millicuries per gram were used. The gold arrived in the form of foil and was evaporated onto one face of each gold specimen cylinder to a thickness of about 100A. A sandwich-type specimen was formed by welding two such cylinders together. Evaporation of Gold: The gold was evaporated under vacuum from heated tantalum strips which were bent in such a way as to limit the solid angle through which the gold was allowed to vaporize, thus insuring a more efficient utilization of the gold. The specimens rested on flat brass rings which had an inner diameter of 0.475 in. The entire specimen-holding assembly could be manipulated from outside the vacuum system by means of a magnet which attracted a slug of soft iron attached to the assembly. By evaporating inert gold on glass slides under conditions identical to those employed for the radioactive gold, it was found that the thickness of the films was about 100A. Welding: The welding was performed by hot pressing in a stainless steel cylinder. The inside of the cylinder was threaded and fitted for two plugs. The specimens to be welded were placed in the middle of the cylinder and two pressing disks, one at each end, were inserted to avoid shearing stresses as the plugs were tightened. Mica disks were placed between the pressing disks and the specimens to prevent them from welding. The plugs were then tightened with a hand wrench and the entire unit was placed in an argon stream for about an hour to remove the oxygen. The unit was then inserted in the center of an argon atmosphere furnace maintained at about 700°C and left there for about an hour. Because of the difference in the temperature coefficient of expansion of the two metals, as the temperature rose. the pressure on the specimen-rollple increased and a weld resulted Welding was generally satisfactory under the conditions described.

Jan 1, 1955

By F. R. Jones

AT Tahawus, N. Y., National Lead Co. operates the MacIntyre development. Here the world's largest titanium mine produces 5200 long tons of ore per day and pours 8000 long tons of waste rock over its dumps. Concentrated ilmenite is sent by rail to National Lead Co. pigment plants, and a second product, magnetite, is sold to steel producers in raw form or is agglomerated and shipped as sinter. Several earlier attempts had been made to produce iron from the deposits, which have been known since 1826. These attempts failed, chiefly because of titanium impurity. In 1941 the present owners reestablished the operation for production of war-scarce ilmenite, and the impurity became the main product. The Ore: The MacIntyre ore zone is about 2400 ft long and 800 ft wide in horizontal measurements. Ore outcrops were found on the northwest side of Sanford Hill, 450 ft above Sanford Lake and 2500 ft southeast. The zone dips at about 45" toward the lake and plunges to the southwest. The ore minerals, ilmenite and magnetite, are unevenly distributed in bands roughly parallel to the long axis of the ore zone and are interspersed with bands and horses of waste. Hanging wall ores are fine grained and grade from rich ore to waste rock or gabbro. Footwall ores are coarse grained and are almost entirely ilmenite and magnetite. The foot-wall waste rock, anorthosite, is the common country rock. Several faults cut the ore zone. These faults have no great displacement but do contribute to the great physical variations in ore rock and surrounding waste. The Mine: The MacIntyre mine is an open pit operation, with benches at 35-ft intervals. The lowest bench is now 54 ft below lake level. Loading equipment consists of three electric-powered shovels (a P & H model 1400 with 4-yd dipper and two Bucyrus-Erie models 85-B with 2%-yd dippers) and one diesel-powered shovel (a Northwest model 80D with 2%-yd dipper). Ore and waste are transported to a 48x60-in. jaw crusher in ten 22-ton Euclid trucks with 300-hp diesel engines. Ordinarily the two Bucyrus-Erie 2 % -yd shovels load ore into a fleet of three or four trucks. This combination works two 8-hr shifts per day, moving 5200 long tons of ore to the crusher and removing a small portion of the waste rock. The P & H model 1400 shovel, with a fleet of four trucks, loads waste on three shifts per day. The mine operates on a 5-day week, with a small maintenance crew working Saturday. Oversize rock is broken by a dropball handled by an Osgood model 825 rubber-mounted crane.' Ore and waste are broken by drilling and blasting 9-in. diam vertical holes behind the benches. Bucyrus-Erie 42-T churn drills are used to drill the holes, which are extended 4 ft below the bench level on which the broken rock will fall. Drilling and Blasting History: In its early years the mine was equipped with Bucyrus-Erie 29-T churn drills, which drilled 6-in. holes. To keep up with production requirements the hole diameter was soon increased to 9 in., and by 1950 the three 42-T drills now in use had been acquired. Early blasting experiments with different kinds and grades of explosive led to adoption of 90 pct straight gelatin dynamite as standard. It was recognized that this explosive was expensive, and from the start of operations until 1950 extensive experiments were made using blasting agents of the ammonium nitrate family. Results were recorded as uniformly poor, with great build-up of oversize rock. The expense of these experiments, and the discouraging results, caused the abandonment of any expectation of breaking MacIntyre rock with anything but 90 pct straight gelatin dynamite. Further standardization led to 9-in. well drillhole spacings set at 16 ft in ore and 18 ft in waste, exceptions being permitted only for unusual conditions. The hole burdens were theoretically about 22 ft. Due to the extreme back-slope of bench faces, caused by blasting with heavy charges of dynamite, actual burdens were commonly well over 30 ft. Lack of precise control resulted in many holes having a burden as light as 15 ft. General practice was to stem 6 or 7 ft of hole with magnetite concentrate, the amount of stemming being left to the discretion of the pit foreman. Usually all holes in a row were fired instantaneously with Primacord detonating fuse. Millisecond delays were

Jan 1, 1957

By E. W. Stewart, G. P. Conard, J. F. Libsch

The effects of several additives upon the reduction characteristics of hydrogen-reduced ferrous formate are described. The various additives inhibit sintering of the reduced iron particles by apparently different mechanisms. The magnetic properties of the low density compacts produced from the resulting ultra-fine iron powders were improved markedly. THE permanent magnetic characteristics of ultra-fine iron powder prepared by various means have been a subject of considerable interest and experimentation in the past few years. When such particles are small enough to show single domain behavior, they possess' 1—permanent saturation magnetization, and 2—high coercive force. In the absence of domain boundaries, the only magnetization changes in a particle occur through spin rotation which is opposed by relatively large anisotropy forces. With decreasing particle size, the coercive force tends to increase to a maximum and then decrease because of the instability in magnetization associated with thermal fluctuations. Kittel' has calculated the critical diameter at which a spherical particle of iron can no longer sustain domain boundaries or walls to be approximately 1.5x10-' cm. Stoner and Wohlfarthr in England and Neel4,6 in France have shown from purely theoretical calculations that the high coercive force expected from single domain particles is dependent upon crystal anisotropy, shape anisotropy, or strain anisotropy contributions. Further work by Weil, Bertaut,' and many others has contributed much to the understanding of fine particle theory. Neel and Meikeljohn" have demonstrated that a decrease in particle size below a critical value of approximately 160A leads to a quite rapid decrease in coercive force because of the prevention of stable magnetization by thermal agitation. Lih1, working with powders prepared by the reduction of formate and oxalate salts of iron, has shown the marked influence of powder purity upon magnetic properties. Maximum coercive force was obtained in powders of approximately 65 pct metallic iron content while the maximum energy product, (BxH) occurred in powders of 85 pct metallic iron content. Careful consideration of the preceding theoretical considerations and experimental results has led to the manufacture of permanent magnets from ultra-fine ferromagnetic powders by powder metallurgy techniques. Such work has been done by Dean and Davis," the Ugine Co. of France, and Kopelman." The aforementioned work of Kopelman and the Ugine Co. was concerned somewhat with the effect of various additives upon the properties of hydrogen-reduced ferrous formate. Virtually no work, however, has been published on the effects of additives on the reduction rates of metal formates, although unpublished work by Ananthanarayanan16 howed promise of improved energy product in ultra-fine iron compacts prepared by the hydrogen reduction of a coprecipitated mixture of magnesium and ferrous formate. After consideration of the preceding information, it was hoped that a better balance between the metallic iron content and particle size of the reduced iron powder could be accomplished by a prevention of the attendant sintering of the partially reduced iron powder during the reduction reaction. It appeared possible that magnesium oxide might interpose a mechanical barrier between adjacent iron particles and prevent their sintering together, while metallic cadmium and metallic tin would interpose a liquid barrier which might accomplish the same purpose. The degree to which these materials were effective in accomplishing the foregoing objective and the experimental details associated with the work are reported in the following sections of this paper. Experimental Procedure Preparation of Formate and Oxide Mixtures: To obtain ferrous formate of reproducible reduction characteristics, a slight modification' was made in the technique of Fraioli and Rhoda." A supersaturated solution of ferrous formate was mixed with an equal volume of 95 pct ethyl alcohol and the formate crystals precipitated by stirring and screened to —325 mesh. These crystals were in the shape of elongated hexagons, approximately 4x10 micron in dimension. Various preparations of such ferrous formate, designated as lot 111, were reduced for 2 hr, yielding ultra-fine iron particles of exceedingly reproducible size, metallic iron content, and magnetic properties. The magnesium and cadmium formates were prepared by the reaction of dilute formic acid with their respective carbonates, while the tin formate was prepared by the reaction of dilute formic acid with stannous hydroxide. To evaluate the effect of metallic formate additives in intimate mixture with the ferrous formate, varying amounts of magnesium, cadmium, and tin formates were coprecipitated with the latter. The designations of these materials and their chemical compositions are given in Table I. Due to the differing solubilities of the various formates in aqueous media,

Jan 1, 1956

By Robert E. Green, Kenneth Reifsnider

Several experiments have been performed in order to illustrate the application of a recently developed X-ray image intensifier system to metallurgical investigations. In the present work the system has been used to study the instantaneous alterations in Laue transmission X-ray diffraction patterns during tensile deformation of aluminum single crystals. Expem'mental results are presented which demonstrate the capability of the system for crystal orientation, for following orientation changes due to lattice rotation during tensile deformation, and for showing changes in the homogeneity of the lattice planes along the specimen length as a function of strain rate. RECENTLY, a new X-ray system has been developed which incorporates a cascaded image intensifier and permits direct viewing and recording of X-ray diffraction patterns produced on a fluorescent screen.1"3 In the present work the results of several experiments are presented which demonstrate the usefulness of this system for metallurgical applications. EXPERIMENTAL PROCEDURE A schematic diagram of the experimental arrangement is shown in Fig. 1. In this system a Machlett AEG-50-S tungsten target X-ray tube, normally operated at 50 kv and 40 ma, serves as the X-ray source. The X-ray tube is placed in direct contact with a 10-in.-long collimator, which transforms the X-ray beam from one with a circular cross section to one with a rectangular cross section 3 in. high and 1/6in. wide. By blocking off all but a small portion of the rectangular slit, it is possible to work with the more conventional "pinhole" collimated X-ray beam commonly used for obtaining Laue diffraction patterns. In the present work the test specimens were 99.99+ pct aluminum single crystal wires & in. in diam and 3 in. long. For the deformation tests the wire crystals were mounted in a special set of grips in a table model Instron machine so that diffraction patterns could be recorded during specimen deformation. For the orientation tests the wire crystals were mounted in a rotating goniometer so that diffraction patterns could be recorded during specimen rotation. At a distance of 3 cm from the specimen axis, a 6 in. diam DuPont CB-2 fluorescent screen is positioned to transform the X-ray image to a visible one. A Super Farron f/0.87 72 mm coupling lens, corrected for 4 to 1 demagnification, transmits the visible image to the image tube. The image intensifier used is a three-stage magnetically focused RCA type C70021A with an S-20 input photocathode and a P-20 output phosphor. The tube has unity magnification and useful input and output screen diameters of 1.5 in. The image on the output phosphor is of sufficient intensity to be viewed directly, to be recorded cine-matographically, or to be displayed by vidicon pick-up on a television monitor. The recording device most commonly used is a 16 mm Bolex motion picture camera fitted with a Canon f/0.95, 50 mm lens. The overall gain of the system is 16,000 for direct viewing and 2240 for recording on 16 mm movie film. The resolution of the system is limited to 1 line pair per mm which is approximately that of the fluorescent screen. This system has been used for cine recording of transmission Laue X-ray diffraction patterns with exposure times as short as 1/220 sec and for vidicon television pick-up and display at a scan time of 1/30 sec. Quantitative information may be obtained from each frame of the movie film, by either stopping the vertical slit down to a point source in order to obtain a conventional Laue photograph or else by retaining the linear beam and introducing fiducial marks as described in a previous paper.4 In either case, each frame may be enlarged to appropriate size for analysis by either using a photographic enlarger and making prints of the desired frames, or, more conveniently, by using a microfilm reader. EXPERIMENTAL RESULTS The first series of photographs which are presented in Fig. 2 serves to demonstrate the usefulness of the system for crystallographic orientation determination. This series of prints, made from enlargements of a 16 mm movie film, shows the dynamic Laue transmission patterns produced by an aluminum single crystal wire which was rotating about the wire axis when the patterns were recorded. The movie films were taken at 16 frames per sec and the crystal was rotated at a rate of 15 rpm.

Jan 1, 1970

By ED. E. Furman, R. H. Atkinson

HITHERTO the practical creep testing of precious metals has received little or no attention. The only previous creep tests of precious metals have been made with wires under conditions such as to yield much more rapid rates of creep than in engineering tests.', ' Up to the present time the value of creep bars of adequate size, in the absence of real need for engineering data, has deterred investigators. However, the increasing use of platinum at high temperatures has demonstrated the need for reliable creep data for the guidance of engineers, especially those engaged in designing certain specialized chemical plant equipment. In order to supply this need, creep tests were conducted at 1382°F (750°C) on 0.290 in. diam specimens of platinum, 90 pct Pt, 10 pct Rh and palladium. The platinum was high purity, nominally 99.95 pct Pt. The 90 pct Pt, 10 pct Rh was of the same high quality as is used for making gauzes for the catalytic oxidation of ammonia. The palladium was also of high purity; two batches of palladium bars were tested, one deoxidized with calcium boride and the other with aluminum. Spectrographic examination of the palladium confirmed its good quality; the only significant impurities apart from the residual deoxidizers were traces of silicon and lead. Procedure The creep bars, which were furnished by Baker and Co. to our specification, were 6 ¾ in. in overall length with a 4½ in. (4 in. gage length) reduced section 0.290 in. in diam and had the ends threaded (?-NC16). It may be of interest that the bars were valued at up to $600 each. The specimens were supplied in a 50 pct cold-worked condition to facilitate attachment of the creep extensometer, which was of the push rod type. Because of the softness of the platinum and palladium, the extensometer rings were secured to the test section by means of circular knife edges instead of the usual pointed set screws. The extensometer rods extended through the bottom of the furnace and readings were taken with a 0.0001 in. "Last Word" dial gage fastened to the rods for the duration of the test. The bars were directly loaded by hanging weights from the lower specimen grip. All tests were conducted at 1382°F ± 2°F, and an effort was made to maintain the temperature gradient over the test section within 2°F. The ends of the furnace tube were packed with asbestos wool, which allowed a very slow circulation of air through the tube. Annealing was accomplished in the creep furnace before the load was applied. The platinum and palladium specimens were annealed at the test tem- perature for about 17 and 24 hr respectively; in the case of the rhodioplatinum it was found expedient to anneal for 1 hr at 1922°F (1050°C). Pilot samples cut from the same stock as the bars were used to check annealing procedures. Pertinent measurements of grain size and hardness were recorded. Results and Discussion The creep data obtained are given in Table I and the creep curves are plotted in Figs. 1, 2, and 3. Two platinum specimens, tested under a stress of 250 psi, had almost identical creep rates at 2000 hr, namely 0.000008 and 0.000009 pct per hr. A third platinum specimen, stressed at 400 psi, had a creep rate at 2000 hr of 0.000026 pct per hr; the reason for a rather sharp decrease in creep rate during the period from 1200 to 1600 hr is unknown. As it was thought that 90 pct Pt, 10 pct Rh would have a lower creep rate than platinum, the first sample was tested at 400 psi; however, the creep rate was approximately 50 pct greater. Microex-amination revealed that differences in grain size might be responsible for the unexpected result, as annealing at 1382°F developed an average grain diameter of 0.0021 in. in the rhodioplatinum specimen compared with 0.004 in. in the platinum bar. Annealing the alloy for 1 hr at 1922°F (1050°C) increased the average grain diameter to 0.0032 in. and materially improved the creep resistance, making it much better than platinum. A second specimen annealed at 1922°F (1050°C) and tested under a stress of 550 psi had a creep rate of 0.000022 pct per hr at 2000 hr, which was still substantially lower than that shown by the specimen annealed at 1382°F (750°C) and stressed at only 400 psi. In contrast to the creep behavior of the platinum and rhodioplatinum specimens, the palladium bars, whether deoxidized with calcium boride or aluminum, were characterized by high first stages of creep. However, after about 1200 hr of test, the creep

Jan 1, 1952

By Victor deWolfe

INTRODUCTION The Henderson Mine, located near Empire, Colorado, utilizes a continuous panel caving system to extract ore as one of the world's major producers of molybdenum. Any mine using a caving-by-gravity technique of mining must rely on closely controlled draw of the caved ore. This control is essential to insure proper caving action, to avoid damaging load concentrations of weight and to minimize the dilution of ore with waste material. Henderson is no exception. Draw control is a major factor in all production planning, from long- range plans to short-range and day-to-day ore scheduling. Draw control is reviewed constantly and administered daily in an effort to optimize production efficiency, ore recovery, and cave management. MINING METHOD The cave at Henderson is massive, moving slowly through large panels that are 244 m (800 ft.) wide by 610 m (2,000 ft.) long. Generally two cave areas are drawn at one time. The areas under active draw vary in size but can be as large as 244 m (800 ft.) by 244 m (800 ft. ) containing 400 draw points. Each draw point contains 45,360 mt (50,000 st) on the average and takes about two and one half years to exhaust. A complete panel is worked for seven to ten years. No pillar exists between panels, but rather a buffer zone of broken ore, or "static face," is left in each panel to be drawn with the adjacent, yet-to-be-caved panel in efforts of minimizing dilution of a working area from an exhausted one. (Figure 1) Production drifts are driven on 24.4 m (80 ft.) centers through the ore body. Between the production drifts are funnel-shaped draw bells on 12.2 m (40 ft.) x 24.4 m (80 ft.) centers to receive ore from the cave. Each bell is accessed by two draw points, one from the production drift on either side, thus forming a 12.2 m (40 ft.) x 12.2 m (40 ft.) draw pattern. Extraction of the ore is by rubber-tired, 3.8 m3 (5 yd3) load-haul-dump equipment. The LHDs then tram the ore a maximum of 49 m (160 ft.) to ore passes. Cave initiation and bell development are done from the undercut drifts which are parallel to and 17 m (55 ft.) directly above the production drifts. Longhole rings are drilled and blasted from the undercut drifts to define the bells and establish the undercut for caving. (Figure 2) DRAW CONTROL Since the cave line at Henderson is constantly advancing, it is necessary to be continually initiating new cave at one end while exhausting it at the opposite end. There must exist, therefore, an angle on the ore-waste contact in the broken rock from initiation to exhaustion. The basic concept of draw control is to keep this angle as smooth and even as possible, particularly at the time of exhaustion. If this is achieved, draw points are exhausted more or less in a line, avoiding pockets of remaining ore surrounded by exhausted areas. These pockets would cause spotty ore extraction at the time of exhaustion, increasing the amount of dilution occurring while introducing the potential for significant weight problems in the production area. To arrive at the desired angle on the ore- waste contact, maximum tonnage percentages are assigned to each row of draw points increasing at 10% or 15% increments (depending on cave size and velocity of draw) working away from the cave line. The available tonnage indicated by these percentages is the maximum allowable tonnage to be extracted from each draw point until the available tonnage percent- age is increased. As the cave moves, these percentages increase for each draw point regularly. However, in general the tonnage drawn from each draw point is kept at about 50% of this allowable maximum in order to maintain adequate available tonnage in the cave to sustain production for seven months if cave initiation were to cease. This available tonnage cushion is a safeguard built into the draw control program at Henderson to accommodate fluctuations in the rate of cave advance. When draw points move past the row of 100% tonnage availability, they are drawn past the desired 50% at the same increments per row until exhausted. (Figure 3) To achieve proper draw control, the number of LHD buckets to be taken from each draw point is assigned daily. The actual buckets taken, which may at times deviate from the

Jan 1, 1981

By E. Teghtsoonian, E. M. Schulson

The tensile behavior of bcc ordered P' AuZn single crystals (CsCl structure) has been investigated under varying conditions of temperature, composition, and orientation. Between -0.2 and 0.4 T, multi-stage hardening occurs fm stoichiometric and nonstoichio-metric crystals oriented near the middle of the primary stereographic triangle. At higher and lower temperatures, parabolic type hardening occurs, followed by work - softening at the higher temperatwes. Deviations from stoichiometry give rise to increased flow stresses. Multi-stage hardening was observed for most orientations, except along the [loll-[lll] boundary and near the [001] corner of the stereo -graphic triangle, where parabolic type hardening occurs. Along two slip systems, (hk0)[001] and (, operate simultaneously while in the [001] comer, slip occurs mainly on the system. Electron microscopy of deformed crystals revealed bundles of edge dislocations forming walls approximately Perpendicular to the glide plane. In general the plasticity of 4' AuZn closely resembles the plasticity of bcc crystals. In recent years, considerable interest has arisen concerning the mechanical properties of the CsCl type intermetallic compounds Ag Mg,'- Fe co,' and Ni Al.'-' The compound P'AuZn is structurally similar. It has a low and congruent melting point of 725"~,'" remains ordered up to the melting point,16 and pos-esses a range of solid solubility from 47.5 to 52.0 at. pct Au at room temperature.15 The present paper reports the results of an investigation on the general tensile behavior of material in single crystal form. Some dislocation configurations characteristic of the deformed state are also reported. The results of a detailed study of the slip geometry in AuZn are presented in a separate paper.17 PROCEDURE Alloy preparation, crystal growing techniques, and the procedure followed in selecting specimens of minimum composition variation are reported elsewhere.17 Dumb-bell shaped tensile specimens were prepared by carefully machining single crystals in a jewellers' lathe to a gage length of 0.80 in. and diam of 0.090 in. Back-reflection Laue X-ray patterns and room temperature tensile tests revealed that machining damage could be eliminated by electrochemically polishing 0.005 in. from the machined surface followed by annealing at 300°C for 1 hr. Specimens were polished in fresh 5 pct KCN solution (40°C, 12 v). Experiments were performed by gripping specimens in a self-aligning pin-chuck and threaded collet system, then straining in a floor model Instron tensile machine. All tests were performed in duplicate. Experimental variables included temperature, composition, and orientation. Unless otherwise stated the strain rate was 2.5 x 10"3 per sec. Liquid testing environments included nitrogen (WOK), nitrogen cooled petroleum ether (133" to 293"K), and silicone oil (293" to 488°K). Resolved shear stress-shear strain curves were electronically computed from autographically recorded load-elongation curves. Stress and strain were resolved on the macroscopic noncrystallographic (hkO) [001] system operative under the specific test conditions of temperature, strain rate, and orientation reported earlier.17 RESULTS The temperature dependence of the work-hardening curves is shown in Fig. 1 for gold-rich crystals of 51.0 at. pct Au oriented near the center of the stereo-graphic triangle. Over the range of intermediate temperatures from -200" to 400°K, they are very similar to those classically observed for fcc metals (reviewed by Nabarro et al.).'' The beginning of deformation is characterized by a region of decreasing hardening rate, stage 0, which is followed by a region of low linear hardening, stage I, and then a region of higher linear hardening, stage 11. At the higher temperatures, stage 111 is observed, a region of decreasing hardening rate. Over the intermediate temperature range, the extent of stage 0 and of the slow transition between stages I and I1 decreases with increasing temperature. Total ductility is large, often greater than 300 pct shear. As the temperature is either increased or decreased, the extent of stage I is decreased, giving rise to parabolic type flow and reduced ductility. Similar temperature effects have been reported for bcc ~r~stals.~~-~~ Below -14O°K, hardening is terminated in brittle fracture while above -400°K. initial hardening is followed first by work-softening and then by chisel-edge type ductile fracture. Stoichiometric (50.0 at. pct Au) and Zn-rich (51.0 at. pct Zn) crystals were also tested from 77" to -500°K. The effect of composition on the flow behavior is illus-

Jan 1, 1970

By H. T. Green, R. M. Brick

True stress-strain data on alloys of pure iron with up to 2.4 pct Al were obtained in the temperature range +100° to —185°C. Alumi-num was found to reduce yield and flow stresses of iron at low temperatures but to have little or no effect on ductility. The effects of temperature and composition on strain hardening are discussed. SEVERAL independent studies of the behavior of high purity iron binary alloys at low temperatures are now in progress in attempts to evaluate systematically the variables affecting the low temperature brittleness of ferritic steels. This paper reports the results of one such investigation in which the tensile properties of aluminum and aluminum plus silicon ferrites were measured from 100" to —192°C. True stress-natural strain data have been obtained in order to evaluate as many as possible of the parameters which describe the behavior of the materials involved. In comparable studies at the National Physical Laboratory in England, iron and iron alloys of high purity have been produced' and tested at subat-mospheric temperatures.' True stress-natural strain curves were obtained there also. The purest iron contained 0.0025 pct C and 0.001 pct O and N. Even this, as normalized at 950°C following hot rolling, showed little ductility at -196°C. The grain size was ASTM No. 3, and the room-temperature yield strength was 17,800 psi (which seems too high for pure iron). Some of the NPL irons contained considerably more oxygen and demonstrated intergran-ular fracture at —196°C. The authors2 carefully differentiated between intergranular fractures associated with excessive oxygen content and transcrys-talline cleavage with little ductility encountered at —196°C in the purer material. The cleavage stress was half again as great as that associated with inter-granular fracture. Test Material, Preparation, and Procedures Of a number of Fe-A1 alloys produced, eight were considered to be sufficiently pure for testing. Partial chemical analyses (Table I), low observed yield points, and high ductilities indicate these alloys to be comparatively pure for vacuum-melted irons of sizable ingots, 5 Ib or more. To produce the binary Fe-A1 alloys, electrolytic iron was melted in air, cast into slabs, and rolled to strips 0.010 in. thick. These strips, joined into a continuous ribbon and wound into 2 1/2 in. diameter spools, were subjected for four weeks to a moving atmosphere of purified dry hydrogen in a stainless-steel tube at 1050" to 1150°C. Charges of these spools were melted in beryllia crucibles under good vacuums (1 micron), and aluminum (99.97 pct Al) was added to the melts. Compositions of these alloys are recorded in Table I. The ingots were hot forged and then cold rolled at least 65 pct to 3/8 in. rods which were vacuum annealed to the desired grain size, approximately ASTM No. 4, prior to machining into tensile test bars. All tensile specimens had gage sections 1 in. long, with a fillet of 1.5 in. radius to the shoulder. Gage diameters were 0.250 in, except for a few rods where additional cold work required use of a 0.200 in. gage section. After machining, 0.002 in. was removed from the gage diameter using 240, 400, and 600-grit metallo-graphic papers. The final polish with 600 grit left the fine scratches running in the longitudinal direction. By this means, surface metal strained during machining was removed. A few specimens heat treated after machining were similarly reduced 0.004 in. to remove any material affected chemically by the atmosphere during heat treatments, as is discussed in a later section. Tensile tests of the eight alloys at constant temperatures from +100° to —185°C were performed in apparatus which has been described." The essentials include a double-walled insulated metal vessel which contained the liquid heat-transfer medium surrounding the test specimen. A constant temperature was maintained by means of a pyrometer which regulated the pressure of dry air driving liquid air through a copper coil. Temperature variation was less than ±2°C during a specific test. For axial straining, two lengths of case-hardened chain, terminating in simple shackles, loaded the specimen through threaded grips. The lower grip bar passed through a hole in the bottom of the test vessel to which it was joined by a thin-walled

Jan 1, 1955

By C. W. Beattie, R. L. Solter

ALUMINUM-KILLED, low carbon steel sheets are used extensively for severe deep drawing and other difficult forming operations. They usually, but not always, have a characteristic grain structure in which the grains are elongated both in the lengthwise and in the transverse direction. As described by Burns and McCabe,&apos; a typical grain in the plane of the sheet has its two axes in that plane from 1 Y2 to 4 times as long as the axis normal to the plane of the sheet. Rickett, Kalin, and MacKenzieZ have also reported on the recrystallization behavior of such steel. The contrast in grain structures of fully processed sheets of aluminum-killed and rimmed steel is illustrated by Figs. 1 and 2. The elongated grain structure of the aluminum-killed sheet is not developed on all heats or lots of this metal, and studies of the factors controlling and influencing its formation are reported in this paper. Jeffries and Archerb tate that unstrained grains are normally equiaxed, but exceptions are common. For example, if a metal containing a material mechanically obstructing grain growth is subjected to considerable working followed by thorough annealing, it may exhibit grains consistently elongated in the direction of working. Our experiments demonstrate that aluminum-killed, low carbon steel is such a metal, and that the substance mechanically obstructing grain growth is aluminum nitride. The effectiveness of aluminum nitride in inhibiting grain growth has been found to be influenced by the degree of cold reduction, the rate of heating in annealing, the thermal history of the sample before cold reduction, and the residual aluminum content. A correlation between grain shape and austenitic grain coarsening temperature also was indicated and additional experiments demonstrated that aluminum nitride is also the principal cause for the fine grain characteristic of aluminum-killed steels. Manufacture In conventional practice, aluminum-killed sheet steel is manufactured from a low carbon steel containing approximately 0.02 to 0.07 pct residual (HC1 soluble) Al. With the exception of certain samples containing greater or lesser amounts of aluminum, the steels used in these investigations were within the following composition range: C, 0.03 to 0.06 pct; Mn, 0.28 to 0.38; S, 0.017 to 0.032; Al, 0.03 to 0.06; P, <0.01; and Si, <0.01. Properly heated ingots are rolled to slabs about 4 in. thick. After surface conditioning, the slabs are reheated to about 2300°F and hot rolled continuouslv to strip about 1/10 in. thick. The strip rolling is completed at a temperature of 1550°F or higher, and the strip is coiled, usually at a temperature near the lower critical transformation. After cooling, the strip is pickled to remove oxide, cold reduced 40 to 70 pet to final thickness, then annealed to 1250° to 1350°F in 20 to 80 ton charges, the size of which results in slow heating and cooling rates. Effect of Cold Reduction According to Sachs and Van Horn,&apos; the deformations of the individual grains in rolling are similar to those of the total volume. Thus individual grains would elongate in rolling according to the amount of cold reduction imposed. This is true theoretically, but as cold reduction increases the individual grains tend to fragment, and measured grain elongations become less than theoretical. The amount of grain elongation may be described by a numerical rating based on grain counts made by the intercept method. Specimens are polished normal to the plane of the sheet, with the polished surface extending parallel to the rolling direction. After etching, grain intercepts are counted along a 50 mm line on a micrograph of suitable magnification. In random locations parallel to the plane of the sample 20 counts are made and 20 are made in the thickness direction of the sample the average count in the thickness direction divided by the average count parallel to the plane of the sample gives a numerical rating of the grain shape called grain elongation. For example, a grain elongation of 2.00 means that the average grain is twice as long as it is thick. The average of both counts may be converted to grains per sq mm by a nomograph relating intercept counts and grain count. By the same procedure the grain elongation in the plane of the sheet but transverse to the rolling direction may be determined, using transverse metallographic samples. A comparison of theoretical and measured grain elongation was obtained on an aluminum-killed

Jan 1, 1952

By A. Lubinski

In drilling operations, attention generally is given to hole angles rather than to changes of angle, in spite of the fact that the latter are responsible for drilling and production troubles. The paper presents means for specifying maximum permissible changes of hole angle to insure a trouble-free hole, using a minimum amount of surveys. It is expected that the paper will result in a decrease of drilling costs, not only by avoiding troubles, but also by removing the fear of such troubles. SUMMARY, CONCLUSIONS AND RECOMMENDATIONS Excessive dog-legs result in such troubles as fatigue failures of drill pipe, fatigue failures of drill-collar connections, worn tool joints and drill pipe, key seats, grooved casing, etc. Most of these detrimental effects greatly increase with the amount of tension to which drill pipe is subjected in the dog-leg. Therefore, the closer a dog-leg is to the total anticipated depth, the greater becomes its acceptable severity. Very large collar-to-hole clearances will cause fatigue of drill-collar connections and shorten their life, even in very mild dog-legs. Another finding regarding fatiguing of collar connections in dog-legs is that rotating with the bit off bottom sometimes may be worse than drilling with the full weight of drill collars on the bit, mainly in highly inclined holes when the inclination decreases with depth in the dog-leg. Means are given for specifying maximum dog-legs compatible with trouble-free holes. An inexpensive technique proposed is to take inclinometer or directional surveys far apart; then, if an excessive dog-leg is detected in some interval, intermediate close-spaced surveys are run in this interval. The application of the findings should result in a decrease of drilling costs, not only by avoiding troubles, but mainly by removing the fear of such troubles. The result would be much more frequent drilling with heavy weights on bit, regardless of hole deviation. Because of errors inherent to their use, presently available surveys are not very suitable for detecting dog-legs. There is a need for instruments especially adapted to dog-leg surveys. Crooked hole drilling rules should fall into two distinct categories—(1) those whose purpose is to bottom the hole as desired, and (2) those whose purpose is to insure a trouble-free hole. Three kinds of first-category rules in usage today are as follows. 1. A means to bottom the hole as desired is to prevent the bottom of the hole from being horizontally too far from the surface location; this may be achieved by keeping the hole inclination below some maximum permissible value such as, for instance, 5. 2. Another means to achieve the same goal is to limit the rate at which the inclination is allowed to increase with depth. A frequently used rate is 1/1,000 ft. In other words, a maximum deviation of l° is allowed at 1,000 ft, 2 at 2,000 ft, 3 at 3,000 ft, etc. 3. Whenever application of the first two means precludes carrying the full weight on bit required for most economical drilling, then the best course is to take advantage of the natural tendency of the hole to drift updip, displace the surface location accordingly and impose a target area within which the hole should be bottomed. This method has already been successfully applied,&apos;.&apos; and its usage probably will become more frequent in the future. Means for calculating the amount of necessary surface location displacement are avail-able.3&apos;5&apos;6 If in high-dip formations the full weight on bit should result in unreasonably great deviations, the situation could be remedied by increasing the size of collars and (if needed) the size of both hole and collars,351 or in some cases by using several stabilizers. Rules which would fall into the second category (i.e., rules whose purpose is to insure a trouble-free hole) are seldom specified today. It is vaguely believed that following Rules 1 and 2 of the first category will automatically prevent troubles. Actually, this is not true. If at some depth the only specified rule is that the hole inclination must be less than 4", the hole may be lost if the deviation suddenly drops from 4 to 2, or if the direction of the drift changes, etc. Rule 3 of the first category is generally used in conjunction with a rule belonging to the second category, namely, that the hole curvature&apos; (dog-leg severity) must not exceed the arbitrarily chosen value of 1½ /100 ft. Moreover, when using this rule, the industry is not clear over what depth intervals the hole curvature should be measured. All this results in a frequent fear

By R. R. Vandervoort

The creep behavior of Nb(Cb), Ta, Mo, and W was determined under conditions of constant atomic dif-fzisivity, constant stress to elastic modulus ratio, and nearly equivalent grain size, and the steady-state creep rates obtained from these tests were correlated with calculated stacking fault energies for the metals. These results, in conjunction with similar data for several fccMetals,13 suggest that stacking fault energy may influence the creep strength ofbcc metals. The interrelationship between steady-state creep rate, subgrain size, and stacking fault energy was examined. It was found that the subgrain size for a given creep stress, increased as stacking fault energy increased, but that this relationship did not cormpletely account for the effect of stacking fault energy on creep rate. The crystallography and energetics of stacking fault formation in bcc metals has been discussed by a num-ber of authors,1-5 and impurity stabilized stacking faults on (112) planes have been observed in Nb,6,7 w,8,9 Fe,] and V" by transmission electron microscopy. However, a crucial question is whether or not stack-ing faults influence the mechanical strength of bcc metals. Potentially, stacking faults could increase strength by reducing the mobility of the partial dis-locations bounding the fault, by acting as barriers to slip dislocations, and by retarding the climb of dislo-cations during high-temperature deformation. The objective of this study was to seek a correlation be-tween creep strength and stacking fault energy for several bcc metals; namely, Nb, Ta, Mo, and W. The creep behavior of most polycrystalline metals and alloys at high temperatures and moderate stresses can be described by the following relation:11,12 im=Af(s) where i, = minimum creep rate, A = constant, j(s) = a function involving metallurgical structure, a = applied stress, E = average elastic modulus at the test tempera-ture, w = constant (equal to 5 for most pure metals), D = diffusion coefficient. One factor in the structure function F(s) which sig- R. R. VANDERVOORT, Member AlME is Research Metallurgist, Process and Materials Development Division, Chemistry Department, Lawrence Radiation Laboratory, University of California, Livermore, Calif. Manuscript submitted February 28, 1969. IMD nificantly affects the creep resistance of fcc metals is stacking fault energy, and creep rate has been shown to vary directly with stacking fault energy to the 3.5 power." In the latter investigation, four fcc metals of widely different stacking fault energies (Ag, Cu, Ni, and Al) were creep tested at a constant stress to modulus ratio of 1.21 x 10-4, at a constant diffusivity of 2.7 x 10-12 sq cm per sec, and at nearly equivalent grain sizes of about 0.7 mm. The creep data were then correlated with stacking fault energies. In the present study, a similar procedure was followed. All materials used in this work were consolidated by powder metallurgy techniques. Impurity contents in the as-received materials are listed in Table I. Chemical analyses showed that no measurable contamination of the test specimens occurred during pretest annealing treatments or creep testing. Specimens with a gage section 0.75 by 0.125 by 0.050 in. were creep tested in tension in a vacuum of less than 10-9 torr. Deformation at temperature was measured by tracking fiducial marks on the gage section of the specimen with an optical comparator. Optical deformation measurements also permitted observation of the macroscopic characteristics of the deformation Table I. Typical Specimen Impurity Content, ppm Nb Ta Mo W C 45 10 155 6 O 185 30 4 10 N 30 6 3 2 H 5 I 1 <1 als 3 10 2 15 Ca <5 I3 5 Cr 5 <3 10 <5 Cu 10 50 2 15 Fc 10 10 150 35 Ni 2 150 20 <5 Si <I0 1 3 <10 Ta 100 Ti 10 8 1 Zi 15 50 1 3 Table II. Test Conditions for Constant Stress-Modulus Ratio of 6 X 10.&apos; and Constant Diffusivity of 2.7 X 10-12 sq cm per see, and Grain Size Values for the Given Pretest Annealing Treatments Literature references Pretest Annealing for E and D Treatment Stress, Temperature, ___"&apos;Values__ Grain Tempera-Metal psi "C E D Size, mm ture, .C Time hr Nb 745 1525 14 15 to 17 0.83 1650 I Ta 1220 1770 18 19.20 O.91 1800 I Mo 1975 1630 18 21 0.77 2200 I W 2140 2265 18 22 040 2400 5

Jan 1, 1970

By R. T. Hukki

This paper presents a series of equations for mechanical ball wear, relating parameters of ball size, mill speed, and mill diameter. The fundamental equation, Eq. 12, presented here is introduced to correlate these basic parameters and thus define and clarify the concept of ball wear. This equation is offered as a general rule, which may be modified to apply to individual problems of grinding. BALL wear as observed in grinding installations is the combined result of mechanical wear and corrosion. Corrosion should be a linear function of the ball surface available. Ball corrosion, however, has been studied so little that its effect, although of great importance, cannot be included in the analyses given here. In a separate paper&apos; it is shown that 1 n = 0.7663 np----=— rpm [1] vD P = c, np D kw [2] T=c²(np)n De tph [3] In these equations n — actual mill speed, rpm np = calculated percentage critical speed D = ID of mill in feet P = power required to operate a mill, kw T = capacity of a mill, tph C¹ and c² - appropriate constants in = exponent of numerical value of 1 5 m 1.5 Exponent m is the slope of a straight line on logarithmic paper relating mill speed (on the abscissa) and mill capacity (on the ordinate). It is generally accepted, although not sharply defined, that ball wear in mills running at low (cascading) speeds is a function of the ball surface available. Accordingly, the wear of a single ball may be considered to be a homogeneous, linear function of its surface and of the distance traveled. Thus dw = f¹(d2) . f2(ds) [4] where dw is the wear of a single ball in time dt, d the diameter of the average ball in ball charge, and ds the distance traveled by the ball in time dt. Indicating that ds - a D n dt, the wear of the average ball in time dt becomes dw = f¹(d2) . f2(Dn dt) 1 --- f¹ (d1) f² (D c3 np-----— dt) \/D = c,d² n, D dt The rate of wear of the average ball is given by dw/dt. dw/dt = c, d² np D lb per hr [5] The weight of the ball charge per unit of mill length is a function of D The number of balls of size d in the ball charge is = f³(D2)/f4(d³). The rate of wear of the total ball charge equals the number of balls times rate of wear of the average ball. Thus rate of total ball wear = — . (dw/dt) w. c, . (l/d) . n,, D lb per hr [6] which is the equation of ball wear in low speed mills. In a mill running at a low speed, grinding is the result of rubbing action within the ball mass and between the ball mass and mill liners. When the speed of the mill is gradually increased toward the critical, the impacting effect of freely falling balls becomes increasingly prominent in comparison with the rubbing action. Reduction of ore takes place partly by rubbing, partly by impact. The share of the freely falling balls in the reduction of ore reaches its practical maximum at a speed somewhat less than the critical; at that speed grinding by rubbing has decreased to a low value. It may be reasonable to think that size reduction by freely falling balls should reach its theoretical maximum at the critical speed, if the fall of the balls were not hindered by the shell of the mill beyond the top point; grinding by rubbing would cease at the critical speed. As a first approximation, wear of freely falling balls may be considered to be a homogeneous, linear function of the force at which they strike pieces of rock and other balls at the toe of the ball charge. The force equals mass times acceleration. The mass of a ball is a function of d3 and its acceleration is a function of the peripheral speed of the mill. The wear of a single ball of size d representing the average ball in a ball charge will therefore be w¹ = f3(F) = f (d3) f7 (v). [7] Indicating that v = D n, and n = c³ np 1/vD, Eq. 7 becomes W1 = cn d3 np Do.5 lb per hr. [8] Total wear of the ball charge equals number of balls times the wear of the average ball. Number of

Jan 1, 1955

By W. M. McKewan

Magnetite pellets were reduced in flowing hydrogen at pressures up to 40 atm over a temperature range of 350° to 500°C. The rate of weight loss of oxygen per unit area of the reaction surface was found to be constant with time at each temperature and pressure. The reaction rate was found to be directly proportional to hydrogen pressure up to 1 atm and to approach a maximum rate at high pressures. The results can be explained by considering the reaction surface to be sparsely occupied by adsorbed hydrogen at low pressures and saturated at high pressures. PREVIOUS investigation1,2 have shown that the reduction of iron oxides in hydrogen is controlled at the reaction interface. Under fixed conditions of temperature, hydrogen pressure, and gas composition, the reduction rate is constant with time, per unit surface area of residual oxide, and is directly proportional to the hydrogen pressure up to one atmosphere. The reduction rate of a sphere of iron oxide can be described3 by the following equation which takes into account the changing reaction surface area: where ro and do are the initial radius and density of the sphere; t is time; R is the fractional reduction; and R, is the reduction rate constant with units mass per area per time. The quantityis actually the fractional thickness of the reduced layer in terms of fractional reduction R. It was found in a previous investigation2 of the reduction of magnetite pellets in H2-H,O-N, mixtures, that the reaction rate was directly proportional to the hydrogen partial pressure up to 1 atm at a constant ratio of water vapor to hydrogen. Water vapor poisoned the oxide surface by an oxidizing reaction and markedly slowed the reduction. The enthalpy of activation was found to be + 13,600 cal per mole. It was also found that the magnetite reduced to meta-stable wüstite before proceeding to iron metal. The following equation was derived from absolute reaction-rate theory4,8 to expfain the experimental data: where Ro is the reduction rate in mg cm-2 min-&apos;; KO contains the conversion units; Ph2 and PH2O are the hydrogen and water vapor partial pressures in atmospheres; Ke is the equilibrium constant for the Fe,O,/FeO equilibrium; Kp is the equilibrium constant for the poisoning reaction of water vapor; L is the total number of active sites; k and h are Boltzmann&apos;s and Planck&apos;s constants; and AF is the free energy of activation. Tenenbaum zind Joseph5 studied the reduction of iron ore by hydrogen at pressures over 1 atm. They showed that increasing the hydrogen pressure materially increased the rate of reduction. This is in accordance with the work of Diepschlag,6 who found that the rate of reduction of iron ores by either carbon monoxide or hydrogen was much greater at higher pressures. He used pressures as high as 7 atm. In order to further understand the mechanism of the reduction of iron oxide by hydrogen it was decided to study the effect of increasing the hydrogen pressure on rebduction rates of magnetite pellets. EXPERIMENTAL PROCEDURE The dense magnetite pellets used in these experiments were made in the following manner. Reagent-grade ferric oxide was moistened with water and hand-rolled into spherical pellets. The pellets were heated slowly to 550°C in an atmosphere of 10 pct H2-90 pct CO, and held for 1 hr. They were then heated slowly to 1370°C in an atmosphere of 2 pct H2-98 pct CO, then cooled slowly in the same atmosphere. The sintered pellets were crystalline magnetite with an apparent density of about 4.9 gm per cm3. They were about 0.9 cm in diam. The porosity of the pellets, which was discontinuous in nature, was akrout 6 pct. The pellets were suspended from a quartz spring balance in a vertical tube furnace. The equipment is shown in Fig. 1. Essentially the furnace consists of a 12-in. OD stainless steel outer shell and a 3-in. ID inconel inner shell. The kanthal wound 22 in. long, 1 1/2, in. ID alumina reaction tube is inside the inconel inner shell. Prepurified hydrogen sweeps the reaction tube to remove the water vapor formed during the reaction. The hydrogen is static in the rest of the furnace. The sample is placed at the bottom of the furnace in a nickel wire mesh basket suspended by nickel wire from the quartz spring. The furnace is then sealed, evacuated, and refilled with argon several times to remove all traces of oxygen. It is then evacuated, filled with

Jan 1, 1962

By J. F. Martin, L. M. Melnick, R. Rapp, R. C. Takacs

Recent studies on the determination of hydrogen in steel have shown that the hot-extraction method for removing hydrogen from a solid sample is preferable to its removal from a molten sample by vacuum fusion or by fusion in vacuum with tin. A number of techniques are available, however, for determining the hydrogen so extracted. They include: thermal conductivity, gas chromatography, pressure measurement before and after catalytic oxidation of the hydrogen to water and removal of the water, and pressure measurement before and after diffusion of the hydrogen through a palladium membrane. These techniques have been evaluated on the basis of initial cost, maintenance, speed and accuracy of analysis, and applicable concentration range. The results of this study showed that the palladium-membrane technique is best suited for routine use. FOR some time investigators have been concerned with the origin, form, and effect of hydrogen in steel. In such stdies&apos;, the analysis for hydrogen constitutes one of the most important phases. It is quite apparent that the results for hydrogen concentrations in a given steel are dependent on the method of obtaining the sample, storage of the sample until analysis, preparation of the sample, and analysis of the sample, including all the facets inherent in the calibration and operation of an apparatus for gas analysis. There are a number of means available for determining hydrogen. This is a critical study of some of the more common techniques in use today. In most conventional melting and casting methods, hydrogen concentrations of 4 to 6 parts per million (ppm) in steel are quite common. Because of the undesirable effects of hydrogen on steel there has been increased use of techniques such as vacuum melting,&apos; vacuum casting, and ladle-to-ladle stream degassing, which lower the hydrogen content to levels on the order of 1 to 2 ppm. Therefore, the method used for determining hydrogen in steel must be sensitive and precise. In any analytical procedure for gases in metals there are two distinct operations—the extraction of the gas from the metal and the analysis of the extracted gas. To extract the gas from the steel, three methods have been employed: 1) fusion of the sample with graphite at high temperature; 2) fusion with a flux, such as tin, at a lower temperature; and 3) extraction of the hydrogen from the solid sample at a temperature below the melting point of the steel. Fusion with graphite is the least-acceptable method. The blank in this method is higher and more variable than in either of the other two methods. The hydrogen fraction of the total gas composition usually is between 10 and 50 pct; thus, a larger analytical error is possible. The vacuum-tin fusion4 extraction of hydrogen is probably the most rapid method in use today; the extraction time is usually about 10 min. However, with this system a bake-out of the freshly charged tin for 2 hr is necessary and a change of crucible and a charge of fresh tin are required after each day of operation whether one or thirty samples have been analyzed. In addition, frequent checks of blank rates are required since CO and Na are continually being given up by the steel samples dissolved in the tin bath. The composition of the gas in this method lends itself readily to analysis; although the hydroge concentration may fall to as low as 50 pct, more often it is above 90 pct, thus allowing a more precise analysis (because of less interference from other gases). In 1940 ewell&apos; published the hot-extraction method for extracting hydrogen from the solid sample, comparing analysis for hydrogen extracted at 600°C with similar analysis for the gas extracted at 1700°C by fusion with graphite. Good agreement for hydrogen was obtained between these two methods, provided sufficient time was allowed for extraction at the lower temperature. carsone obtained good results in his comparison of this hot-extraction method with vacuum-tin fusion. Subsequent work by Geller and sun7 and Hill and ohnson&apos; has shown that steel samples should be heated to at least 800°C to effect the release not only of the diffusible hydrogen but also of the "residual" hydrogen that may be present as methane. Since the rate of evolution of hydrogene9l0 depends on such factors as sample size and composition, thermal history, and extent of cold work, a fixed extraction time is not possible. Extraction times of 30 min are normal, but 2 hr are not unusual. Induction or resistance heating may be used in the hot-extraction method. With resistance heating the

Jan 1, 1964

By E. M. Cox, A. S. Skapski, N. H. Nachtrieb, M. C. Bachelder

As a result of using copper-containing scrap in the steelmaking process, the copper content of steels has been steadily increasing for years. Consequently the possible role copper may play in the steelmaking process and in the finished product begins to attract the metallurgists&apos; attention. Some time ago one of the present authors forwarded the idea—based on the results of the analysis of nonmetallic inclusions extracted electrolytically from steels—that sulphur in plain carbon steels is distributed mainly between copper and manganese, the amount of iron sulphide being very small; and that, consequently, the problem of copper and that of sulphur in steel cannot be treated separately.&apos; At the time of the publication of the quoted paper little was known about the relative affinities of copper and manganese for sulphur at high temperatures except that at moderate temperatures (below 1000°C) the affinity of manganese for sulphur is much greater. To gather more experimental data on this subject, the present authors undertook the investigation of the equilibrium constants of the reactions: 2Mn(8 or 1) + S2(g) = 2MnS(s) 4Cu(s or 1) + S2(g) = 2Cu2S (S or I)* 2Fe(s) + S2(g) = 2FeS (s or 1) over a range of temperatures wide enough to establish the dependence of these equilibrium constants on temperature. From the equilibrium constants (K = l/Ps2) the free energy of formation (affinity) can be calculated from F° = -RTln 1/PSt (1) where the standard conditions chosen are: 1 atm of sulphur pressure and the activities of condensed components equal one. The decomposition pressure, Ps2, of sulphur over the respective sulphides is too small to be measured directly, but there is a way of eliminating this difficulty by measuring the equilibrium constant of the reaction between the sulphide and hydrogen. From the latter and from the equilibrium constant of the thermal dissociation of H2S we then calculate Ps2 for the respective sulphide. 2Mn + 2H2S = 2MnS + 2H, 2H2 + S2 = 2H2S_________ 2Mn + S2 = 2MnS The numerical values of the equilibrium constant of the thermal dissociation of H2S at different temperatures were taken from Kelley&apos;s paper, "The Thermodynamic Properties of Sulfur and its Inorganic Compounds."² In previous experimental work published by Jellinek and Zakowski3 and by Britzke and Kapustinsky4 the equilibrium constants of the reactions Metal sulphide + H2 = H2S + metal were determined by passing hydrogen, at different rates of flow, over the sulphide, analyzing the resulting H2S + H2 mixture and then extrapolating the H2S/H2 ratio (which is a function of the rate of flow) to the zero speed of flow, a method necessarily involving considerable uncertainty. In the present work the equilibrium ratio was actually measured instead of being extrapolated. The apparatus is shown in Fig 1. Experimental Procedure The sulphides were prepared by the following methods: FeS Powdered iron which had been reduced with hydrogen (ferrum reduc-tum) was mixed in stoichiometric ratio with sublimed sulphur and carefully ground. The mixture was put into an alundum crucible, covered with pure sulphur, and the reaction started by touching the mixture with a glowing iron rod. After the reaction was completed the product (still containing some metallic iron) was again ground with sulphur, put into a Rose crucible, covered with sulphur, and heated in a strong current of pure hydrogen. Analysis of the final product showed 62.46 pct Fe and 36.59 pct S. Theoretical for FeS: 63.53 pct Fe and 36.47 pct S. MnS Manganese sulphide (precipitated and carefully washed with distilled water containing H2S) was dried in a Rose crucible in an atmosphere of H2S and heated in a current of hydrogen for 2 hr at red heat. The product was then ground and ignited for several hours at 1000°C in a current of hydrogen sulphide. Analysis showed 64.53 pct Mn and 36.63 pct S. Theoretical: 63.15 pct Mn and 36.85 pct S. Some MnS samples were prepared from metallic manganese and sublimed sulphur by mixing and grinding them and then heating in a current of hydrogen sulphide in an alundum tube.

Jan 1, 1950

By J. B. Wagner, p. Hembree

The iron tracer diffusion coefficient of umstite has been measured at 110(fC across the phase field and at a single composition at 800°C. Assuming a simple cation vacancy model the tracer diffusion coefficient was found to be a linear function of the cation vacancy concentration at 1100°C. The equation is D = 3 x 20 29 where denotes the concentration of vacancies in numbers per cc. The tracer work at 800°C was carried out to investigate the reported "pinning" of tracer to the wustite surface at low temperatures. No evidence for the "pinning" of the tracer was found at 800°C in COz-CO gas mixtures. HIMMEL, Mehl, and Birchenall,&apos; Carter and Richardson,2 and Desmarescaux and La combe3 have measured the diffusion of iron tracer in wustite at several temperatures and compositions. The present work was undertaken to extend the measurements over a large composition range at 1100°C and to resolve certain apparent discrepancies in the data, expecially at lower temperatures. EXPERIMENTAL Wustite was prepared by oxidizing rectangular iron plates* in C02-CO mixtures. The samples were •The iron was supplied by the Battelle Memorial Institute courtesy of the American Iron and Steel Institute. The analysis is presented in Table I. quenched. Due to the inward flow of cation vacancies during oxidation, the center of the sample contained a thin void. The edges of the wustite slab were sanded until the sample could be split into two parts. Each part was then sanded on the front and back flat area until a smooth surface was obtained. The specimens were then replaced in the furnace and equilibrated at llOO°C in a predetermined COa-CO mixture by methods described elsewhere.4"6 The specimens were again quenched and the surfaces were lightly sanded to remove any roughness following the first equilibration. The specimens were then reequi lib rated in the same C02-CO mixture for thirty minutes in order to relieve any mechanical damage on the surface due to the polishing. The specimens were then quenched and the tracer was applied by an electroplating technique. The work of Carter and ~ichardson&apos; demonstrated that there was no systematic difference in the iron tracer diffusion coefficient in wustite if the tracer was plated, dried, or evaporated on the specimen. In the present study a piece of filter paper was saturated with an iron chloride solution of pH <* 3 that contained the tracer FeS5. The wustite was placed on the filter paper and made the cathode. A current density of 0.4 to 0.6 ma per sq cm was passed for about five to ten minutes. The thickness of the tracer layer was estimated to be about 7 x lom6 cm. This estimate was made by considering the area plated, the current flow, and time for plating and the activity of the iron in the plating solution. Different areas of the specimen were counted using a collimator to determine the uniformity of the tracer. Any specimen which exhibited a variation from the initial count rate (about 1500 cpm) by more than 15 pct was rejected. An estimate of the time necessary to convert the thin layer of iron tracer to wustite was made using the data of Pettit and wagner." he estimated time was 1 sec at 1100°C assuming linear oxidation kinetics. The shortest diffusion anneals were 1800 sec. The samples were suspended in the hot zone of a furnace by two platinum wires. Two separate specimens were run at the same time. Only the edges of each sample were in contact with the wires. The C02-CO gas of the same composition as that used in the pre-diffusion anneals flowed freely around the samples at a linear velocity of 0.9 cm per sec. To initiate a run, the specimens were lowered from the cold zone of a furnace to the hot zone by a magnetic lowering device." bout 60 sec were required for lowering. To terminate a run, the sample was withdrawn from the hot zone to the cold zone. Time zero for the beginning of the experiment was taken when the sample blended into the red glow of the furnace and conversely for the end of the experiment. The surface decrease method of measuring the tracer diffusion coefficient was used to collect the data. This method requires that counting geometry be reproducible because the specimen is counted before the diffusion anneal and after the anneal. A special jig was constructed for each specimen so the specimen could be removed from the jig and returned to the jig such that the well geometry was reproducible.

Jan 1, 1970

By A. S. Grove, C. T. Sah, E. H. Snow, B. E. Deal

A summary of- several recent investigations in to the properties of the metal-oxide-silicon system is presented. A major portion of these studies makes use of the MOS capacitance-z)oltage method of&apos; analysis. The particular areas of investigation which are reported include: 1) a general survey of the electvical properties of thermally oxidized silicon surjbccs; 2) a study of ion migration through silicon dioxide films ; 3) measurements of electron and hole mobilities in surface inversion layers; 4) a study of impurity redistribution due to thermal o.ridatiotz; and 5) measurements of the rates of oxidation oj-heavily doper7. silicon. THE importance of the metal-oxide-semiconductor (MOS) system in the semiconductor industry is well-known. In addition to its importance in the "planar" device technology,&apos; the MOS structure is now also used in the fabrication of active solid-state devices. Consequently, extensive efforts have been made recently to obtain a better understanding of the characteristics of this system. A summary of some studies of the MOS system conducted in our laboratories during the past year is presented. For the most part these studies used silicon as the semiconductor, along with silicon dioxide and aluminum as the other two components of the system. Since the MOS capacitance-voltage method of analysis was used extensively in these studies, we will first briefly describe its nature and consider some of the possible causes of deviation of experimental observations from the simple theory. We will then outline the various related areas of investigation carried out in our laboratories and will briefly indicate some of the results. It should be noted that the purpose of this paper is merely to provide a brief summary of MOS studies. More detailed discussions of the various areas of investigation are given in the references cited. PRINCIPLES OF THE MOS C-V METHOD OF ANALYSIS&apos; A sketch of the MOS structure is shown in the upper portion of Fig. 1. In this case the insulating film is Si02 and the semiconductor p-type silicon. If a large negative bias is applied to the metal field plate, holes are attracted to the silicon surface. The silicon then behaves much like a metal and the capacitance measured is that of the oxide layer alone, Co. If a small positive bias is applied to the aluminum, holes are repelled and a region depleted of majority carriers is formed at the silicon surface. This depletion I-egion adds to the width of the dielectric and the measured capacitance begins to drop. With increasing positive bias, the width of the electrical depletion region increases. At some large positive bias an inzevsion regiotr is formed at the surface and additional charges induced in the silicon appear in the form of electrons in this narrow inversion region. Thus the depletion-region width approaches a maximum value and, consequently, the capacitance reaches a minimum value and then either levels off or rises again depending on the measurement frequency and the rate of equilibration of the minority carriers in the inversion layer.3 Band diagrams, along with the corresponding charge distributions, are shown in Fig. 1 for the above bias conditions. If minority carriers cannot accumulate at the surface to form an inversion region, the depletion-region width continues to increase with increased positive bias and the capacitance drops toward zero as in a reverse biased p-n junction. The effect of a work-function difference $hs between the metal and the silicon, and of surface charges per unit area Qss located at the oxide-silicon interface, is simply to attract charges in the silicon much like the applied bias. It can be shown that this results in a parallel shift of the capacitance-voltage characteristic along the voltage axis by an amount corresponding to AV = -$bIs + Qss/Co. Theoretical curves have been calculated4 giving the capacitance of the MOS structure C normalized to the oxide capacitance Co vs the quantity VG here VG is the voltage applied to the metal field plate. In Fig. 2 such calculations are shown as points for a particular oxide thickness and bulk impurity concentration for a p-type semiconductor. (For an n-type semiconductor the curves would be mirror images of these.) All three cases, i.e., low frequency. high frequency, and depletion, are indicated. Also shown in the figure are recorder tracings of the characteristics of actual devices. These characteristics have been shifted along the voltage axis to compensate the effect of surface charges and work-function difference. It is evident that agreement between experiment and theory is good. The nature of this shift along the voltage axis is

Jan 1, 1965

By Pol Duwez

IN the course of an investigation of the V-CO system, two intermediate phases were found. One of these phases corresponds approximately to the stoichiometric composition VCo and is isomorphous with the sigma phase in the Fe-Cr system.&apos; The second phase has the composition V3Co; its crystal structure is described in the present note. The alloys were prepared by mixing the two metals in the powder form, pressing a small disk weighing about 5 g at 80,000 psi, and arc melting this disk on a water-cooled copper plate in an atmosphere of pure helium. The details of this technique have been described.&apos; The vanadium powder was obtained from Westinghouse Electric Corp., Bloomfield, N. J. This powder is probably of very high purity, since when it is properly sintered or melted in the above-mentioned arc furnace, ductile specimens are obtained. The cobalt powder, from Charles Hardy, Inc., New York, contained 0.5 pct Ni, 0.1 pct Cr, and traces of Si and Fe. After melting, the V,Co samples were sealed in evacuated quartz tubes and homogenized for ten days at 800°C. Powder diffraction patterns were obtained with a 14.32 cm diam camera, using Ka copper radiation. The patterns were readily indexed on the basis of a primitive cubic lattice with a parameter equal to 4.675A. The density, determined by the immersion method, was 6.71 g per cu cm; hence the number of molecules per unit cell is approximately 1.95; i.e., 2. At this point, the possibility that the structure might be that of beta tungstena became apparent. The beta tungsten structure is described as follows: Space group 03,, — Pm3n 2 Co in (a) : 000; ?4lhYZ (hhl) reflection present only if 1 = 2n. Assuming this structure to be the correct one, intensities were computed by means of the usual eauation: 1 + cos220 I oc p F sin 0 cos 6 where F is the structure factor, 0 the Bragg angle, and p the multiplicity factor. The observed and calculated values of sin 0 and the intensities are given in Table I. The agreement between the observed and the calculated sin 0 is good and there are no flagrant discrepancies between the calculated intensities and those estimated visually. The (hhl) reflections for which 1 is odd are not observed, as required by the space group. In addition, the (410), (430), and (531) reflections are missing as expected, because of the special (a) and (c) positions in0%. However, six reflections—(llo), (220), (310), (411), (422), and (510)—which have very weak computed intensities were not observed. For these reflections, the structure factor is proportional to the difference between the scattering factors of the two atoms in the structure. Since the scattering factors of vanadium and cobalt are not very different, these reflections are weak. However, by using Ka chromium radiation, whose wavelength is just above the absorption edge of vanadium, the effective scattering factor of vanadium may be decreased by one or two units; consequently the difference between the cobalt and vanadium scattering factors is increased. It was, indeed, found that in a powder pattern taken with chromium Ka radiation, the three reflections (110), (220), and (310) were actually present. The three other reflections (411), (422), and (510), with spacings smaller than half the wavelength of chromium Ka, were obviously not obtainable with chromium radiation. All the experimental results appear to confirm the beta tungsten structure for V,Co. In this structure, each cobalt atom is surrounded by twelve vanadium atoms at 2.61A; each vanadium atom is surrounded by two vanadium atoms at 2.34A, four cobalt atoms at 2.61 A, and eight vanadium atoms at 2.86A. Acknowledgment This work was done at the Jet Propulsion Laboratory, California Institute of Technology, under contract number W-04-200-ORD-455 with the Army Ordnance Department, Washington, D. C. The author wishes to thank this agency for the permission to publish the results of this investigation. References &apos;P. Duwez and S. R. Baen: X-Ray Study of the Sigma Phase in Various Alloy Systems. Symposium on the Nature, Occurrence, and Effect of Sigma Phase. ASTM Special Tech. Pub. No. 110, pp. 48-54. Philadelphia, 1951. 2 C. H. Schramm, P. Gordon, and A. R. Kaufmann: The Alloy Systems Uranium-Tungsten, Uranium-Tantalum, and Tungsten-Tantalum. Trans. AIME (1950) 188, pp. 195-204; Journal of Metals (January 1950). 3 M. C. Neuburger: The Crystal Structure and Lattice Constants of Alpha and Beta Tungsten. Ztsch. fiir Krist. (1933) 85, pp. 232-237.

Jan 1, 1952