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By E. T. Turkdogan, W. L. Worrell

The rate of reaction of hydrogen sulfide with ferrous sulfide was studied by measuring the initial rates of sulfidation of iron strips in hydrogen sulfide-hydrogen-argon mixtures at 670°, 800°, and 900" C. The time-dependent surface sulfur activity is derived from the instantaneous rate of sulfidation with the assumption that diffusion in the sulfide layer is in a pseudo-steady state with the gas-sulfide chemical reaction. The rate of sulfur transfer from hydrogen sulfide to the surface of iron sulfide is proportional to the partial pressure of hydrogen sulfide and inversely proportional to the activity of sulfur at the surface of the sulfide layer. The derived rate equation is based on the assumption that most of the surface sites on the chemisorbed layer of iron sulfide are occupied by sulfur atoms and that the slow rate-controlling reaction is the dissociation of hydrogen sulfide on the chemisorbed layer. The experimental results are in reasonable accord with this reaction model. 1 HE slow approach to parabolic growth rate in the sulfidation of iron in hydrogen sulfide-hydrogen mixtures is a manifestation of slow approach to surface equilibrium between the gas and the surface of the sulfide layer. For example, at 800°C and Fig. 7 in Part I, the parabolic growth rate begins after approximately 1 day of sulfidation time. With decreasing temperature and decreasing partial pressure of hydrogen sulfide, the time necessary to reach gas-sulfide surface equilibrium is much longer. These observations are similar to those reported previously by Turkdogan et al.' on the oxidation of iron to wustite in hydrogen-water vapor mixtures. The slow approach to gas-sulfide equilibrium is well-demonstrated by the results in Fig. 1 where the square of the weight gain per unit area, ( g S per sq cmI2, is plotted against time for gas mixtures having PH g//>H = 1.0 with 0, 25, 50, and 67 pct Ar at 800°C. The value of 1.5 X lo-' (g S per sq cm)' corresponds to almost complete sulfidation of the iron strip (- 0.05 cm thick) to iron sulfide. The points fall on S-shaped curves. If the inflection parts of the curves are considered to be linear, thus indicating parabolic growth rate and the establishment of gas-sulfide equilibrium, their slopes would have to be the same for a fixed ratio in the gas mixture. Such is not the case, and the slopes for the "linear" parts of the S-shaped curves are lower than the value when gas-sulfide surface equilibrium is established. The subject matter of this paper is the kinetics of the surface reaction of hydrogen sulfide with iron sulfide during the early stages of sulfidation of an iron strip. EXPERIMENTAL RESULTS The apparatus and materials used were the same as those discussed in Part I. In the present experiments of relatively short duration, a purified iron strip, 5 by 2 by 0.05 cm, which was suspended in the uniform hot zone of a vertical zircon tube, reacted with a flowing gas mixture of hydrogen sulfide-hydrogen-argon. The sample was suspended from a gold chain attached to a calibrated silica spring. The amount of sulfur picked up by the iron, which formed a layer of iron sulfide, was determined by measuring the displacement of a reference point on the silica spring, using a cathetom-eter. In all cases, the degreased sample was first annealed in a stream of oxygen-free dry hydrogen for several hours to remove any impurities such as oxygen, nitrogen, or carbon which might have been present on the surface of the sample.' The rate measurements were carried out at 670°, 800°, and 900° C in hydrogen sulfide-hydrogen-argon mixtures with pH2 s/Ph2 ratios from 4 to 0.1. In a few experiments of short duration no sulfur deposition was observed in gas mixtures with (comments concerning sulfur deposition at high pH s/ph ratios were made in Part I.) Typical examples of the results obtained are shown in Fig. 2 for two temperatures and two gas compositions. With the present experimental technique, reproducible rate data could be obtained only after a uniform

Jan 1, 1969

By J. P. Morgan, J. R. Dawson

A review of Jones & Laughlin's Vesta coal preparation plant flowsheet in 1959 indicated that, of the three sizes making up the company's metallurgical product, the -48-mesh fraction offered the greatest potential for improvement. Efforts to improve quality and yield were directed toward the use of froth flotation. After laboratory and plant tests were completed, it was decided that treatment of the entire -@-mesh fraction of the coal mix by flotation was a significant step forward in providing better coal quality and recovery. Reduction of ash in the metallurgical coal from 6.75 to 6.25% provided material savings at the coal washer, coke plant and blast furnace. The Vesta preparation plant of the Jones & Laughlin Steel Corp. is located on the Monongahela River, eight miles south of Brownsville, Pa. This plant is capable of cleaning 2400 tph of a mixture of Pittsburgh seam high volatile coals from the Vesta and Shannopin mines to produce a metallurgical coal of 6.75% ash and a steam coal containing 14% ash. Heavy media vessels are utilized to upgrade the coarse coal while launders and tables handle the -1/4-in. fraction of the raw coal feed. Cleaned -1/4-in. coal goes to boot classifiers where the 48 mesh material settles and is removed by bucket elevators. This 1/4x48 mesh material is then sent to centrifugal dryers for dewatering. As designed, the -48 mesh fraction overflows the boot as a 7 to 8% solids slurry and is recovered by cyclones and a thickener and then dewatered by filtration. The products of the heavy media circuit, the launder-table-boot-dryer circuit and the cyclone-thickener-filter circuit, are combined to constitute the shipped metallurgical coal. The Vesta-Shannopin mix is combined at Jones & Laughlin's Pittsburgh and Aliquippa Works with purchased low and medium volatile coals for the production of metallurgical coke. A review of the preparation plant flowsheet in 1959 indicated that of the three sizes making up the metallurgical product, the -48-mesh fraction, containing 16.7% ash, offered the greatest potential for improvement of coal quality. This fraction was considered low, averaging 55%, while the rejects averaged only 33% ash. It was evident that while the boot overflow fraction represented only 6% of the product mix, a reduction in ash to 6 or 7% would reduce the ash of the metallurgical coal by 0.5%. The original plant flowsheet handling the boot overflow (Fig. 1) consisted of ten 14-in. cyclones treating 70% of the overflow and producing a 48x150-mesh underflow and a -150-mesh overflow. The cyclone overflow was combined with the remaining 30% of the boot overflow to constitute the feed to a 150-ft thickener. The cyclone underflow and the 48x325-mesh thickener underflow were fed to vacuum disk filters for dewatering prior to combination with the coarse metallurgical coal. The thickener overflow stream containing the -325-mesh material was split, 40% being recycled directly to the plant water system and 60% being flocculated and clarified in an 85-ft thickener. The underflow of this thickener was pumped to waste. Efforts to improve quality and yield on the -48-mesh coal were directed toward the use of froth flotation early in the investigation. A review of the fine coal flowsheet indicated that three streams should be considered for treatment by froth flotation. These were the -325-mesh thickener overflow, the -150-mesh cyclone overflow and the -48-mesh boot overflow. Flotation of the thickener overflow was considered basically to improve recovery by upgrading the reject stream so that it could be fed to the filters along with the thickener and cyclone underflows. The overflow from the thickener contained 3.7% solids and had an ash analysis of 33%. Flotation of the cyclone overflow and boot overflow were considered to upgrade coal quality as well as yield. Treatment of the cyclone overflow, a 5.4% solids slurry averaging 27% ash, would eliminate the coal thickener and provide a feed of flotation concentrate and cyclone underflow to the filters. Treatment of the boot overflow would eliminate both the cyclones and thickener and feed only flotation concentrates to the filters. Flotation feed in this

Jan 1, 1962

By J. W. Snavely

A practical mathematical treatment is presented for the determination and control of conveyor belt acceleration, particularly for conditions of starting where vertical curves are involved. A typical sample problem is analyzed, with required calculations, to clarify the procedure. THE part that acceleration plays in starting a belt conveyor and its effect on belt conveyor design are well understood in a general way. Its practical importance is easily overlooked, however, and under some conditions, it is absolutely necessary to give the problem of acceleration detailed study. Most handbooks on conveyor belting design adequately present basic data for the determination of acceleration values. This paper will only attempt to present practical thinking and a convenient method of treatment of acceleration in belt conveyor design. Mathematical Analysis In working out the various problems of conveyor belt acceleration, the starting point, as presented by the handbooks, is the familiar formula of "force of acceleration is equal to the mass times acceleration." By expressing these fundamental quantities in terms of belt conveyor design, it is possible to arrive at the unsuspected conclusion that the acceleration time for horizontal belt conveyors is independent of the load, and instead, dependent upon the belt speed, the type of drive arrangement and drive pulley, and the idler coefficient of friction. The mathematics leading to this conclusion are shown in Table I, which has been prepared to show this derivation. While at first the conclusion just given may not seem to be reasonable, further reflection indicates that obviously the type of drive pulley and the type of drive do affect materially the tension in the conveyor belt, and thus, as clearly shown, the time of acceleration is dependent upon the factors mentioned. Inasmuch as all of the factors except time are predetermined by the belt conveyor design, it becomes relatively easy to establish the accelerating time and to reduce further this time determination to a simple graph from which the time in seconds can be read directly. Such a graph is given in Fig. 1. The table appearing on Fig. 1 should be explained further. For a given belt speed, the time of acceleration can be expressed as a percentage of the belt speed. The time of acceleration is also dependent on the drive arrangement, and changes in the drive arrangement consequently change the time of acceleration. It further follows that for a given belt speed, the time expressed as a percentage of that belt speed also changes with the type of drive. Obviously then, it becomes possible to graph the percentage of speed for .each type of drive against the belt speed and accelerating time, after which, for a given belt speed and type of drive, the time can be read directly in seconds. Two constants were established for Fig. 1, the first one being the limiting of the maximum acceleration tension to 35 pct of the full load operating tension in the belt. The purpose of this is to limit the total tension imposed upon the belt during the acceleration period to 135 pct of the full load operating tension, 'which is the amount required to start or breakaway the fully loaded belt conveyor from rest. The other constant is the friction factor used for the idler equipment, which has been established as 0.022. For installations where it is necessary to establish the values of acceleration, invariably high grade idler equipment is used, and it has been established from field experience that 0.022 for the idler friction factor is amply conservative. The use of this friction factor for idlers must be tempered with judgment, of course, for occasions will arise where more power than indicated is required to start, even with the very best of equipment, such as low temperature operations that tend to congeal the grease in the bearings and thus produce additional friction drag. An inspection of the table in Fig. 1 affords a convenient rule of thumb method for determining the acceleration time, which conveniently can be 5 pct of the belt speed in seconds. The 5 pct of belt speed figure is close to the average for most types of drives. In using Fig. 1 it must be emphasized that it applies accurately to horizontal belt conveyors only.

Jan 1, 1953

By K. E. Manchester, A. H. Clark

Hall measurements have been made on three groups of silicon samples, which were implanted with boron, aluininunz, and phosphorus ions. Boron and phosphorus implants show essentially bulk properties when annealed under previously determined conditions, i.e., 15 min at 80O°C for boron and 10 mirz at 600°C for phosphorus. Bulk properties were also observed in alu?rrilzum-implanted samples when ion concentrations were below 1015 cm-'; however, strange low- terrzperalure behavior and a falloff in cavrier concenlralion was observed for sarnples with ion concenlration abozle loL5 ern-'. This can be attributed to solid-solubility effects since the average bulk concentration f0.r the sample exhibiting bulk properties was below the reported solubility limit for aluminum, while the other samples were above the limit. A detailed study of the annealirzg process indicates that mobilities reach bulk c~alues in phosphorus implants at 250DC, in alutninurn ivnplanls a1 500°C, and in boron implants at 700°C. A simple rnodel has been proposed to fit the annealing data. THE technique of semiconductor doping by direct injection of energetic ions is being actively investigated at this Center. Previously published results from this laboratory have been concerned with the investigation of implanted profiles in silicon and the electrical properties of the resultant p-n structures.1'2 Diode properties have been reported for junctions produced by implantation of boron ions and aluminum ions in n-type silicon and phosphorus ions in p-type silicon. Reverse current characteristics, as well as electron diffraction data, have indicated that the damage to the structure produced by the energetic ions during the stopping process can be minimized by a gentle anneal. The time-temperature conditions for this process have been determined by sheet resistivity studies.3 Electrical profiles of implanted phosphorus ions which have been determined by a differential sheet conductivity technique are based on the assumption that the carrier mobilities and activation energies for ionization are the same as in bulk silicon. These assumptions have been substantiated experimentally by measurement of Hall constants and resistivities of these implanted structures and the results are presented herein. After establishing the bulk property equivalence of the annealed layers, a more detailed examination of the annealing process was undertaken. The sheet resistivity technique initially used to determine anneal conditions has been broadened to include Hall constants and mobility behavior with anneal conditions. PROCEDURE The three groups of samples prepared for Hall measurements were implanted with mass-separated beams of boron, aluminum, and phosphorus ions. Ion concentrations are equivalent surface concentration and they have been determined from beam current measurements during implantation. Implant conditions for the samples are shown in Table I. In all cases, the samples had previously been angle-lapped on one edge and the junction delineated with an HF-copper sulfate solution. The surfaces of the implanted regions of the samples were exposed to the delineating solution and in the case of the "n"-type phosphorus-implanted areas it was later discovered that a finite amount of surface could be etched by this solution. This is discussed in the section on results. The implanted samples were ultrasonically cut into the conventional six-probe configuration shown in Fig. l.* In the annealing studies, for which the meas- urements were taken at room temperature, the sample holder was a Bakelite jig incorporating tungsten mechanical probes. The other measurements were taken using a conventional liquid-nitrogen cryostat, in which the temperature could be varied from 77" to about 400°K. For these measurements, 10-mil aluminum dots were evaporated onto the samples and alloyed with the silicon at 530°C for 15 min. The samples were then mounted in flat-packs, and gold wires were bonded to the aluminum dots. The electrical measurements were made with either a Leeds & Northrup K-5 Potentiometer, a Keithley 147 Nanovolt Null Detector, or a Keithley 601 Electrometer. No external bias was applied to the substrate. There is an inherent self-biasing arrangement at one or the other of the current contacts (depending upon the polarity), and this suffices to limit the current flow to the implanted layer. The magnetic field was generated by an A. D. Little electromagnet. Except for occasional checks of linearity of the Hall voltage with

Jan 1, 1969

By P. M. Robinson, M. B. Bever

THE heats of formation at 273°K of the compounds InBi, In2Bi, and TIBi2 have been determined by metal solution calorimetry with bismuth as solvent. The published information on the thermody-namic properties of these compounds is limited. The structure of the compound InBi (mp, 383°K) is tetragonal of the PbO (B10) type.1 The compounds In2Bi (mp, 362°K) and TIBi2 (mp of highest-melting composition, 486°K) are now believed to have structures of the NizIn (B8) type273 and not the A1B2 (C32) structures reported earlier.4, 5 The compounds In2Bi and TIBi2 have limited homogeneity ranges at room temperature, while the compound InBi has no detectable homogeneity range.' Samples of the compounds were prepared from 99.99 pet Bi (Mallinckrodt Chemical Works), 99.999 pet in (American Smelting and Refining Co.), and 99.99 pet Tl (Amend Drug and Chemical Co.) by melting the components in sealed, evacuated Vycor tubes. The compounds InBi and in2Bi were of the stoichiometric compositions. The compound TIBi2 contained 40 pet T1; this composition was chosen because at room temperature the alloy of stoichiometric composition does not lie within the homogeneity range.' The melts were held approximately 50°C above the liquidus for 8 hr and then quenched into iced brine. The samples were homogenized for 48 hr at temperatures 20°C below the temperature of complete solidification. Metallographic examination of sections taken from the center and the ends of the ingots did not reveal second phases or segregation . The heats of formation were determined as the difference in the heat effects on solution of a compound and of a mechanical mixture of its components added from the reference temperature of 273°K to liquid bismuth at 623°K. The details of the calori-metric procedure and method of calculation have been described elsewhere? The calorimeter was calibrated by adding bismuth at 273°K to the bismuth bath at 623°K. The calculated values of the heats of formation of the compounds, listed in Table I, are based on a value of 4.96 keal per g-atom for the difference between the heat contents of bismuth at 273" and 623oK7 No published values are available for the heats of formation of the compounds InBi and In2Bi. Hultgren et a1.7 have calculated a value of -0.50 i 0.10 keal per g-atom for the heat of formation at 423°K of the compound TIBi2 containing 40 at. pet T1 from heat-content measurements in the range from 398" to 700°K and the heats of formation of the liquid alloy. The difference between this value and the value of -0.66 * 0.01 keal per g-atom at 273°K is too large to be attributed to a difference between the heat capacities of the compound and the components over the temperature range 273o to 423oK7 However, the direct determination of the heat of formation by a calorimetric method should give a more accurate value. The heats of formation at 273°K of the compounds InBi and In2Bi lie on a straight line when plotted as a function of composition, Fig. 1. The compound In2Bi, therefore, appears to be barely stable with respect to its neighboring phases. In order for the free energy vs composition curve to be concave towards the composition axis, the entropy change on formation of In2Bi must be more positive than that on formation of InBi. As the compounds are ordered,2, 3, 8 there is no configurational entropy change on formation. The difference between the entropy changes on formation of the two compounds, therefore, is probably associated with the vibrational entropies. In view of the low heats of formation, the changes in vibrational entropy, due to changes in bond strength on formation of the two compounds, are likely to be small. Owing to the volume contractions on formation of the two compounds,5 the vibrational entropies probably decrease slightly but the decrease of the vibrational entropy of InBi is expected to be larger than that of In2Bi.

Jan 1, 1965

By R. W. Snyder, H. J. Ramey

This paper presents the results of applying the Buckley-Leverett' displacement theory to petroleum reservoirs consisting of a finite number of layers. The layers are assumed to communicate only in the wellbores, and the reservoir may be represented as a linear system. Most previous investigations of this nature were limited by assumptions and by inconsistent calculation techniques. This study improves on previous work by applying the Buckley-Lev-erett displacement theory to a noncommunicating layered system where permeability, porosity, initial saturation, residual saturation and relative permeability vary from layer to layer in a logical and consistent manner. Gravity and capillary-pressure effects are neglected. A modification of the Higgins-Leighton calculation method was used in this study. Waterflood predictions were made with all properties varying, and then with only permeability varying using several inability ratios. These results were compared with the Stiles and Dykstra-Parsons predictions. It is shown that the latter methods generally give poor values for the breakthrough recovery and pessimistic predictions for the performance after breakthrough. Similar results were obtained for a gas-displacement case. lNTRODUCTION Field experience with immiscible displacement usually shows constant producing conditions until breakthrough of the displacing fluid. Then oil production continues at increasing displacing-to-displaced fluid ratios until the economic limit is reached. Three different ideal mechanisms are known that will produce this behavior: (1) relative permeability effects as described by Buckley-Leverett frontal advance theory,' (2) vertical stratification as considered by Stiles,2 Dykstra and Parsons5 and others and (3) different path lengths involved in areal (two-dimensional) flow between wells as described by Dyes et al.4 Without question, a combination of these factors modified by formation heterogeneity and other known and unknown factors actually does control the behavior of real systems. This paper presents results of an investigation of certain factors that should affect performance but which have received little attention to date. In 1944, Law5 demonstrated that porosity and perme- ability are often found to have normal and logarithmic-normal distributions, respectively. throughout cored intervals in natural formations. This led to the concept of the noncommunicating, multilayered reservoir model for immiscible displacement. This model assumes that the reservoir is composed of a number of layers that communicate only at the wellbores. Each layer is individually homogeneous, but may be different from every other layer. Stiles' presented one of the earliest applications of this model to waterflood performance. In addition, Stiles assumed that the initial saturations and relative permeabilities were the same for each layer, porosity was the same. displacement was piston-like, fluids were incompressible and injection into each layer was proportional to that layer's permeability capacity (permeability-thickness product). The last assumption would be true if the mobility ratio for the displacement were unity.21 Dykstra and Parsons" used the same model as Stiles, but rigorously included mobility ratios other than unity for piston-like displacement. Dykstra and Parsons used their general result to produce charts for log-normal permeability distributions between layers. Similarly, Muskat6 Pub1ished analytical solutions for linear and exponential permeability distributions. In 1959, Roberts' described a scheme for calculating water-drive performance for the noncommunicating, layered reservoir model which considered two-phase flow in the displaced region. Roberts used the same model and assumed that the injection rate into a layer was proportional to that layer's permeability capacity, but that flood front locations could be evaluated from the Dykstra-Parsons results. These assumptions are inconsistent, and a material balance cannot be maintained except for a mobility ratio of unity. At the same time, Kufus and Lynch8 coupled Buckley-Leverett displacement theory with the layered model to provide an improvement of the Dykstra-Parsons method that was consistent. In 1960, Higgins and Leighton9 resented a numerical method for calculating waterflood performance also considering two-phase flow in the displaced region. The result was used to investigate variation in absolute permeability and oil viscosity. An excellent, detailed history of using the noncomrnunicating, layered reservoir model was presented by Nielsen.'" The preceding techniques (and many related ones) were similar in that differences in initial saturations, residual saturations and relative permeabilities from layer to layer were neglected. It is well known that the irreducible water saturation is an important function of absolute permeability. Calhoun11 showed that the irreducible water saturation

By B. L. Averbach

Recovery and primary recrystalliza-tion in cold worked metals are usually considered as two competing processes. Some of the effects which usually accompany recovery are: alleviation of stress corrosion tendencies, changes in thermal emf,1 damping capacity,2 electrical resistivity,2 and magnetic properties,3 and only minor changes in hardness or the related strength properties. During primary recrystalliza-tion new unstrained grains are formed at the expense of the strained matrix. These new grains eventually become visible metallographically, and nucle-ation and growth kinetics have been indicated for this process.4,5 Frequent attempts have been made to study the cold-working phenomenon by observations on the line broadening by X ray diffraction patterns. Relatively few measurements of line intensities have been made, although Brind-ley and his collaborators, 6,7,8 by means of film techniques, compared the intensities of cold worked Cu, Ni, and Rh patterns with those from chemically precipitated powders. These precipitated powders were presumed to be strain free, and it was found that the intensities for the cold-worked materials progressively decreased as the Bragg angle increased except for the first line, where there was an increase due to reduction in extinction. This was interpreted as a randomness in atomic position induced by cold work. Such randomness is similar to that caused by thermal agitation and has been described as "frozen heat" displacement of 0.08-0.10 A from the mean atomic position. In a recent study9 on the effect of cold work in metals on their powder pattern intensities, the changes in integrated intensity for heavily cold worked alpha brass were observed as a function of the annealing temperature. These measurements were made with a manually operated Geiger-counter spectrometer using CuKa radiation monochromated with a rock salt crystal. Intensity measurements were made with a scaling meter over small intervals of angle, and the equipment was so arranged that the diffracted and incident beams made equal angles with the specimen. Intensities could be compared directly by simply interchanging specimens, and comparisons from day to day were made with a standard whose line intensities did not change on aging. It was shown that a cold worked alpha brass standard was stable for at least a year. Table 1 indicates the results of the integrated intensity measurements on a 70 Cu-30 Zn brass. In the sample preparation, a brass plate was first cold rolled 50 pct and then filed, screened to —325 mesh, compacted into briquettes at a pressure of 60,000 psi and finally annealed for one hour at various temperatures up to 400°C. The briquetting pressure did not seem to influence the integrated intensities, and most of the cold work was introduced by the filing. Although this method of cold work is not quantitative, it was used to obtain random orientation (and thus uniform diffraction lines) in order to make accurate measurements of integrated intensity. Back reflection patterns were taken in each case to check the uniformity of the lines, and from the observed line broadening it was apparent that this type of plastic deformation was quite severe. Care was taken to traverse the entire background of the pattern and to assign to each peak the total intensity above this background. The bases of the diffraction lines were quite broad and spread out over several degrees, even for the narrow peaks. The theoretical intensities were calculated to include a temperature correction, a dispersion correction, and a Lorenz- polarization factor corrected for the crystal monochromated beam. In Table 1 it was necessary to match the calculated and observed values at only one point, and the rest of the experimental values were converted directly to this arbitrary scale. The integrated intensities in Table 1 are listed in arbitrary units, and the accuracy was sufficient to reproduce any of the measured line intensities to within + 1.5 units. It is evident that the percentage error on the strongest line (111) was quite low. The calculated values and the observed intensities for the cold worked material matched reasonably well. As the annealing temperature was raised, however, the intensity of the strongest reflections, particularly the (lll), decreased measurably. Since the background intensities of all of these patterns were identical, such behavior could be interpreted as a primary

Jan 1, 1950

By R. W. Hertzberg, F. D. Lemkey, J. A. Ford

The effect oj unidirectional solidification upon the microstructural, crystallographic and mechanical characteristics of high -purity A1-A13Ni eutectic a1loy specimens has been investigated. varying the growth rate caused the morphology of the Al3Ni phase to gradually change from platelets to rods. In addition, the effect of increasing the rate of planar liquid-solid interlace movement is to decrease the Al3Ni rod diameter and inter-rod spacing. The platelets and rods grow approxinlately parallel to each other and are aligned in a faulted substructure. A single unique crystallographic relationship hetween the A1 and ASNi phases was found and may he descrihed as : interface {001} Al3Ni {331} Al and growth direction <010>Al3Ni || <110> Al The platelet interlace and the aligned rows of rods are both uniquely defined by the above statements. Alignment of the Al3Ni phase by unidirectional solidification has given rise to a threefold increase in strength over that exhibited by specimens with an as-cast microstructure. These results illustrate the possible use of this eutectic alloy as a zohisker -reinforced structure. It has been demonstrated by Winegard et al.,1 Kraft and Albrighht,2 Chilton and winegard,3 Chad-wick,4 Yue,5 Tiller,6 and others that a planar liquid-solid interface may be established in binary eutectic alloys by proper control of heat flow during the solidification process. The unidirectional movement of such an interface results in a eutectic mi-crostructure consisting of an essentially parallel array of the two phases over an entire ingot. Two dominant phase micromorphologies have been produced using this technique: that of substantially parallel alternating lamellae of each phase or long thin parallel rods of one phase imbedded in a continuous matrix of the other phase. While not all eutectics can be controlled in this manner, it has been shown that a "normal" eutectic may, under properly controlled conditions of purity, liquid and solid thermal gradients, and solidification rate, be forced to solidify with essentially parallel phase particles. Kraft and Albright2 have demonstrated that, when the growth rate is too rapid or when the thermal gradient in the liquid at the growing interface is too low, a layer of constitutionally supercooled liquid, formed by impurity build-up, will stabilize a cellular rather than a planar interface. The unidirectional movement of the cellular interface forms the macrostructure of eutectic colonies.7 Chilton,3 Chadwick,4 and Tiller,&apos; studying eutectics of zone-refined Pb-Sn, A1-CuAl2, Zn-Sn, and Cd-Zn systems, noted that in the absence of impurities a planar interface is stabilized even when rapidly solidified. Although it is generally agreed that impurities break down a planar interface to form the eutectic-colony macrostructure, at present the origin of the varying micromorphologies (i.c., plates or rods) in a given normal eutectic alloy is not completely understood. Tiller8 has predicted that the micromor-phology produced may be dependent on solidification rate (i.e., at fast rates a rodlike structure is preferred whereas at slow rates a lamellar structure should form), and also suggested that rod formation may be favored at large phase-volume ratios. yue5 has experimentally verified Tiller&apos;s prediction in the eutectic Mg-Mg17Al12 by observing a lamellae-to-rod transition with increasing growth rate where the phase-volume ratio was approximately 2.3:1. However, Hunt and chilton9 more recently unidirec-tionally solidified over a wide range of growth velocities six different eutectic systems with phase-volume ratios between 12:l and 2.7:1 and observed no lamellae-to-rod transition. chadwickl0 has proposed that the change in micromorphology in some eutectic alloys is due entirely to the presence of impurities and further states that the lamellar structure is the characteristic structure of pure eutectic alloys even when the phase-volume ratio is as great as 12:l. Kraft11,12 has shown that the parallel lamellae in certain eutectic alloys assume a unique crystallographic relationship during unidirectional growth. This preferred crystallography may be developed

Jan 1, 1965

By R. C. Himes, K. Zanio, J. B. Wagner, J. W. Moody, J. F. Miller, W. Johnson

The recent research with which this paper deals has been concerned with the preparation of pure PbTe crystals suitable for control of radiative processes and othe.v electronic applications. The research has been directed toward the developrnent of methods for the preparation of puve, homogeneous, loul-ca,wier-density PbTe crystals suitable for use in the fabrication of optically pumped and electrically pumped diode lasers. Steps in the preparation of the crystals include 1) purification of the starting materials, 2) synthesis and purification of the PbTe, 3) growoth of PbTe crystals, and 4) adjustt~zertt of composition of the crystals This Paper describes the general preparation procedure employed and discusses some investigntiotzs of purification processes. THE as-received lead and tellurium are nominally 99.999 to 99.9999 pct pure. Assays furnished by the suppliers, obtained by emission spectrographic analysis, show only the presence of about 2 ppm Mg in the tellurium. Such analyses are not very informative, however, since they give no information on the concentrations of normally gaseous impurity elements such as oxygen. Both the lead and the tellurium in the as-received state usually contain appreciable concentrations of oxygen—primarily in the form of oxides on the surfaces of the material. To reduce the oxygen content, the lead is held in the molten state in hydrogen at -700°C for several hours. For this treatment, the lead is contained in a graphite boat which is rf-heated, and the hydrogen is passed through a liquid-nitrogen trap prior to use. Tellurium is subjected to multiple distillation in hydrogen to remove oxygen and other foreign impurities. If, in conducting the distillation, approximately the last 10 pct of the charge at each "stage" is left as the discard cut, purification of the tellurium proceeds as indicated in Fig. 1.l Tellurium used in this work has been distilled two or three times to reduce carrier concentration to about 10" per cu cm. The PbTe is prepared from the purified elements by reacting them in the molten state in sealed, evacuated, or hydrogen-filled silica ampoules. The charge is usually held above the melting point of PbTe (923°C) at 1000 C for 24 hr to ensure that it has been homogenized. Two methods have been employed for the growth of the required single crystals of PbTe: growth from the melt, which will be discussed first, and growth from the vapor phase. Large single crystals of p-type PbTe can be grown readily from the melt by the Bridgman method. For this method of crystal growth, the PbTe is synthesized in situ from proper proportions of the purified elements. Growth (dropping) rates in the range 1.3 to 6.5 mm per hr are employed. Material near the maximum melting composition (about 50.012 at. pct Te) is utilized to avoid composition change on freezing, and thus to obtain crystals of nearly uniform composition. Crystals prepared by use of procedures described up to this point were sufficiently pure and perfect that they could be used for the preparation of laser crystals in studies at Lincoln Laboratory.2-5 ADJUSTMENT OF COMPOSITION Good n-type crystals of PbTe cannot readily be grown from the melt. Further, growth from the melt of crystals of any compositions (n or p type) other than those near the maximum melting composition is difficult because of the large composition changes that must occur on freezing.&apos;" However, melt-grown crystals can rapidly be brought to other desired compositions by isothermal equilibration through the vapor phase with a two-phase Pb-Te stock material, essentially as described by Brebrick and Allgaier.1 The work to be discussed has been concerned with the preparation of n-type crystals; thus, the two-phase stock which was employed contained excess lead, and all ensuing discussion is of the preparation and properties of n-type crystals. The equilibration was designed to the shift composition to the lead-rich limit of stability at the selected temperature. Initial treatment of the crystals at high temperature was employed to minimize the total time required for the equilibration. Final adjustments of the composition and state of the crystals were made by use of low-temperature heat treatment at temperatures in the range 200" to

Jan 1, 1968

By Richard J. Fruehan

ThE thermodynamic activities of nickel and chromium in liquid Ni-Cr alloys have not been previously reported. There have been several investigations on the solid alloys.1-3 In the present investigation the activity of chromium in liquid Ni-Cr alloys at 1600° C was determined by the use of the following galvanic cell, The use of galvanic cells employing solid oxide electrolytes for determining the activity of metals in solid alloys was first demonstrated by Rapp and Maak4 on Cu-Ni alloys between 700° and 1000°C. Recently Schwerdtfeger5 tested the performance of the ZrO2(CaO) electrolyte at 1600°C and found it satisfactory for oxygen pressures as low as 10"12 atm. He also demonstrated the use of solid oxide electrolytes for determining the activities in liquid metal systems in measuring the activity of manganese in Fe-Mn alloys at 1500°C. Extensive investigations have been carried out in this laboratory, exploring the oxygen partial pressure lower limit at which ZrO2(CaO) electrolyte can be used satisfactorily in oxygen galvanic cells at 1600°C. In a paper to be published at a later date,9 it is demonstrated conclusively that ZrO2(CaO) exhibits insignificant electronic conductivity at oxygen pressures down to about 3 x 10"13 atm which is lower than the oxygen partial pressure in equilibrium with Cr-Cr2O3 at 1600°C. Therefore, the activity of chromium can be related to the measured reversible electromotive force of the cell in the present investigation since no electronic short circulating will occur. Fig. 1 is a schematic diagram of the cell used. The Cr-Cr2O3 electrode was made of equal weights of chromium and Cr2O3 powders which were mixed and packed into a recrystallized alumina crucible. A Pt-wire lead was embedded into the Cr-Cr2O3 compact to insure good contact. A ZrO2(CaO) crucible, supplied by the Zirconium Corp. of America, contained the Ni-Cr melt equilibrated with solid Cr2O3. The ZrO2(CaO) crucible was placed on the Cr-Cr2O3 electrode and additional Cr-Cr2O3 mixture was packed around it. Contact with the Ni-Cr alloy was made with a Cr2O3 tube, which is a weak electronic conductor. A platinum wire lead was attached to the Cr2O3 tube. The galvanic cell was lowered slowly into the hot zone of the furnace tube in order to avoid too large a thermal shock which could break the ZrO2 (CaO) crucible. The furnace was con- tinuously flushed with purified argon during the experiment. The temperature was measured with a Pt-Pt/10 pet Rh thermocouple and maintained at 1600° ± 5°C. The electromotive force was measured with a recording potentiometer. The furnace was turned off when measurements were being made in order to avoid induction effects and the reversibility of the cell was demonstrated in the usual manner. The electromotive force was measured several times over a period of approximately 1 hr. During that period the electromotive force remained constant within ± 3 pet and did not continuously drift in any one direction. After the experiment the alloys were analyzed for chromium in the usual manner. The measured electromotive forces of the cell at 1600°C, (2912°F), are listed in Table I. The activity of chromium, with solid pure chromium as the reference state, (a), was calculated using Eq. [l], , n F E I, n lna RT- U] where E is the measured electromotive force, F is the Faraday constant, and n is equal to 3, the number of equivalents per mole, (nF/RT = 18.59 volts&apos;1 at 1600°C). The resulting activities are listed in Table I. The activity of chromium relative to pure liquid chromium was also calculated. The free energy of fusion of chromium at 1600° C was calculated using

Jan 1, 1969

By J. T. Norton, H. Blumenthal, S. J. Sindeband

The interstitial phases formed by the transition elements with carbon, nitrogen and boron constitute a unique class of substances which are of considerable technical interest because of their well developed metallic properties and their high hardness and melting points. They are also of interest from the point of view of structure. It was pointed out by Hagg1 some years ago that the principal factor determining these structures was the relative sizes of metal and metalloid atoms. He predicted that if the ratio of the radii of metalloid to metal atoms was less than 0.59, the metal atoms would be arranged in a close-packed fashion of relatively simple type. This prediction has been amply confirmed. It appears, however, that this specification applies most generally to those structures in which the metalloid atoms occupy isolated positions in the lattice and there is no tendency for strong binding between the metalloid atoms themselves. The series of borides of the type MeB2 formed from the transition elements of the fourth (Ti, Zr) and fifth (V, Cb, Ta) groups of the periodic table provide an opportunity to investigate this situation further. If one takes as a basis the radius of 0.87 A for the boron radius and the Goldschmidt values for the metal atoms in 12 coordination, then the radius ratio varies from 0.54 for zirconium to 0.64 for vanadium, which latter is considerably greater than Hagg&apos;s limit of 0.59. It was of interest, therefore, to see if these five elements would form borides of the MeB2 type and have simple isomorphous structures. When this investigation was initiated, relatively little information had been published on the structure of this group of borides. McKenna2 reported the preparation of ZrB2 by a carbon reduction process. Its structure was hexagonal and the approximate lattice constants were given as a = 3.15 A, c = 3.53 A. Ehrlich3 reported the structure of TiB2 as being of the C 32 type with a = 3.02, c = 3.21 and c/a = 1.06 A. Recently, Kiessling4 published the results on ZrB2, showing it to be one in which the metal atoms are arranged in a simple hexagonal lattice with the boron atoms in flat sheets, midway between the layers of metal atoms. The structure is C 32 type with a = 3.169, c = 3.530 and an axial ratio =1.11 A. He also states that CbB2 and TaB2 are isomorphous with ZrB2. Preparation of Borides The borides of titanium, zirconium, columbium and tantalum were made by the electrolysis of fused salt baths after the method described by Andrieux.5 The bath, totalling about 6500 g in each case had the following general composition: Metal Oxide + B2O3 + CaO + CaF2. The actual proportions of the ingredients were determined for each individual case. Although Andrieux reports the preparation of vanadium boride by the same electrolytic method, several attempts proved unsuccessful. However, it was possible to prepare VB2 by the carbon reduction method used by McKenna to produce ZrB2. The mixture consisted of V2O5 + 6B2O3 + 11C and the reduction was carried out in an induction heated graphite crucible. This method proved satisfactory for ZrB2 and TiB2 and undoubtedly the others could be made in the same way. The borides, which were all in the form of fine gray metallic crystals from the electrolytic bath or gray powders from the reduction process, were carefully separated from other constituents and analyzed chemically for metal and boron content. The densities of the boride powders were measured using a standard 10 cc pyenometer at 18°C and ethyl benzene proved to be the most satisfactory liquid for this purpose. X Ray Technique The boride crystals, as prepared, were too small for single crystal examination so that powder techniques were employed. Patterns of the finely ground borides were prepared, using the Norelco recording X ray spectrometer. The lines on these patterns were indexed with the aid of Hull-Davy charts and the approximate dimensions

Jan 1, 1950

By Marston G. Fleming

E. C. Peterson (Anaconda Copper Mining CO., Darwin, Calif.)—A study of this quite comprehensible and interesting paper by Dr. Fleming brings to mind several observations in the practical application of alkalinity and related factors to the actual practice of lead mineral flotation. Soda ash has long been a widely used and a very helpful alkaline conditioning agent in the flotation of galena from the usual run of lead-zinc ores. Soda ash is one of the common "standard" conditioning agents tried in any laboratory investigation of lead-zinc ores because it has so often proved helpful to galena flotation. However, the use of soda ash in the actual flotation of oxidized lead ores is certainly not widespread. In the flotation of certain lead-zinc ores from the Park City district of Utah, it was found in the usual cyanide-zinc sulphate-xanthate circuit that soda ash had an effect of producing a very condensed and flat froth regardless of the many frothers tried and that substitution of lime for soda ash to produce the same pH (8.0 to 8.4) improved the froth condition. However, the flotation of coarse particles of galena became so critical that some would pass through into the tailing, but caustic soda as a substitute for either soda ash or lime produced a very desirable froth condition in the same pH range and greatly improved the metallurgical outcome. Milling was carried out in typically "hard water" from watersheds of limestone and other calcareous rocks and the ore also apparently contributed magnesium and calcium salts to the pulp. Certain lead-zinc ores of Mexico have shown their greatest flotation response in circuits conditioned with sodium bicarbonate, and such actual mill use was also believed to be related to the water necessarily used in milling. On other lead-zinc ores the optimum has been obtained in actual milling treatment by use of caustic soda in both circuits of the operation. In flotation of oxidized lead ores in which sulphidi-zation is employed by addition of sodium sulphide, very high pH&apos;s exist in the flotation circuit. With the usual oxide lead ores, demanding 5 to 15 lb of sodium sulphide per ton of ore, pH&apos;s in the circuit (closed and open water-circuits) may range from 9.5 to 12, and for ores that contain much anglesite or a high percentage of iron oxide minerals, sodium sulphide consumption may reach as much as 30 lb per ton of ore and the resultant pH will be correspondingly higher. Yet in such flotation treatment employing either xanthate or oil (high-sulphur crude oil or diesel oils) as the collector, satisfactory to excellent grades of concentrate and lead recovery are obtained. Although there have been great studies and accomplishments in the investigations of flotation fundamentals and flotation theory, and although an investigation to determine the effect of all factors entering into the actual plant flotation of sulphide and oxide ores must become very complicated, as Mr. Fleming has mentioned in his caution regarding the drawing of general conclusions for other ores, it seems that it would be of greater interest if the techniques and investigations could, with time, approach the conditions of actual practice and lead the way to improved and more efficient flotation plant performance. Marston G. Fleming (author&apos;s reply)—The gulf between fundamental flotation research and plant operation is, perhaps, less wide than Mr. Peterson suggests. For example, the work described in my paper developed from an extensive investigation of a complex lead-vanadium ore from southwest Africa. This investigation started as a standard ore-testing problem but, at almost every stage, results were obtained which could not be interpreted in terms of previous experience. A program of fundamental research was therefore undertaken and was closely interlocked with the ore testing. This coordinated investigation resulted in a flotation process which has now been proved by two years of successful plant operation, and although the grade and mineralogical constitution of ore as well as smelter requirements have altered more than could have been anticipated, our knowledge of the fundamental character of the problem has made it possible to meet each change in conditions with much more confidence than would have been the case had we neglected this aspect of the investigation.

Jan 1, 1954

By W. W. Stephen, J. P. Walsted, W. J. Kroll

FOR the production of carbides, various furnace types are available, especially those using arc, resistance, and high-frequency heating. Selection of a specific means of heating depends primarily on the material to be treated and the physical properties of the carbide produced. In the present case, zirconium carbide had to be prepared on an industrial scale as a raw material for the production of anhydrous zirconium chloride. Considering that a rather expensive pure oxide was to be used, the arc-furnace treatment recommended for zircon sands in a previous publication&apos; was ruled out because of the considerable volatilization and dust losses caused by the blast of the arc. For this reason, either high-frequency or resistance heating seemed to offer more promise. Since there was not enough capacity of the former available, resistance heating was chosen. It was first thought that the Acheson silicon carbide furnace would be suitable for the present purpose, but the voltage in such a furnace, in which the current passes through the batch, varies from 220 to 75 v from the start to the end of a run. This variation is so great that a special tap transformer would have been required. Trouble was also expected by local melting of the carbide. Pure zirconium carbide melts at about 3527°C, but the melting point is brought down to 2427°C, according to Agte,2 when an excess of 6 pct C is present in the carbide. This we found confirmed by experiments in a high-frequency furnace. Excess carbon is needed in the batch to obtain a complete reduction. Fusion of the charge would cause great difficulties in an Acheson-type furnace because of the good electrical conductivity of the carbide as compared with that of the loose batch. Also, fused carbide is much more difficult to chlorinate than the spongy product that can be made in the radiation furnace described below. It was apparent that, to obtain a good-qual-ity zirconium carbide, the heat input would have to be well-controlled. The hairpin-resistor principle seemed to offer possibilities in this regard, and a furnace of this type was therefore developed. The advantages of the hairpin-resistor radiation principle have been discussed in previous publications, and a split-tube graphite-resistor furnace," now increasingly used in various laboratories, as well as a centrifugal quartz melting furnace4 of this type, has demonstrated the usefulness of this heating method. The hairpin-heater element has the following definite advantages over a straight resistor of the type used, for instance, by Georges:5 Its resistance is four times greater; it can expand freely; it is sturdier because of the larger diameter, and it has a larger radiation surface; there are no hot contacts that might wear out or overheat; only one clamp is used which permits assembling all electrical leads at one side of the furnace, making the other sides easily accessible to the operator. The shorter element and its larger diameter permit greater concentration of heat. The furnace developed is shown in fig. 1. The box (I), made of 2 1/2-in. graphite plates, has inside dimensions of 23x17x16 in. It contains the briquet-ted batch (2). The box is embedded in lampblack (3) up to the cover plate. The cover plate contains an opening for the gas escape (4) and for the observation hole (5), which permits measuring the temperature with an optical pyrometer. The cover plate is embedded in charcoal (6). The lampblack is contained in the insulating brick lining (7), held in the 1/4-in. sheet steel box (8). The graphite box is set on two rows of triangular graphite bars (9). The hairpin-heater element (10), the dimensions of which are given below the main drawing, extends horizontally in the graphite chamber and radiates freely on the batch. A graphite tube (11) keeps the lampblack from falling into the slot. The split electrode, which in reality is turned 90" against the drawing, is so arranged that the slot is vertical. The water-cooled packing gland (12) is insulated by an airgap from the heater element. A thin pipe (13)

Jan 1, 1951

By A. E. Back, S. F. Ravitz

When manganese ores are leached with sulphur dioxide, a large part of any zinc in the ore usually is extracted with the manganese.1 In the dithionate process,2,3 in which the manganese ore is leached with dilute sulphur dioxide gas and the manganese is subsequently precipitated by adding lime to the pregnant solution, the zinc is precipitated with the manganese. These facts suggested the possibility of applying the dithionate process to the recovery of zinc from oxidized ores or calcined sulphide concentrates. Accordingly, an investigation was begun, and preliminary tests were made on several ores and concentrates. The results were encouraging, but the need to devote available funds and personnel to other projects made it necessary to drop the work before the various factors involved in the process could be studied in detail. A brief mention of the process,4 however, elicited a number of inquiries, so that it appeared desirable to report the results of the preliminary work. The dithionate process as applied to zinc may be summarized briefly as follows: The ore or calcined concentrate is leached with dilute SO2 gas to dissolve the zinc, largely as zinc sulphate, in accordance with the following equation in which [ZnO] represents the "oxide zinc" in the ore or calcine: [ZnO] + SO2(g) + 1/2 O2 = Zn++SO4— [1] If, in the leaching step, the ore or calcine is suspended in calcium dithionate solution, the sulphate ion, as it is formed, is precipitated immediately as calcium sulphate, which is subsequently filtered off with the insoluble ore residue: Zn++SO4— + Ca++S2O6— = CaSO4(s) + Zn++S2O6- [2] An alternative procedure is to suspend the ore in water, filter off the insoluble ore residue after leaching, add calcium dithionate to the resulting zinc sulphate solution, and then filter off the calcium sulphate precipitate. In either event the zinc is obtained in the form of a pregnant solution of zinc dithionate, from which it can be precipitated as relatively pure zinc hydroxide by adding milk of lime: Zn++S2O6— + Ca(OH), = Zn(OH)2(s) + Ca++S2O6- [3] The zinc hydroxide precipitate is calcined to form the final zinc oxide product: Zn(OH)2 = ZnO(s) + H2O(g) [4] Theoretically the process is completely cyclic with respect to dithionate ion, the calcium dithionate required for reaction 2 being replaced by that formed in reaction 3. In practice, of course, small amounts of dithionate would be lost in the various filter cakes, depending on the economic washing limits. In the leaching step, however, some of the SO2 is oxidized to dithionale instead of sulphate, accord- ing to the following overall reaction: ZnO + 2SO2(g) + 1/2 O2(g) = Zn++S2O6-, [5] so that the dithionate losses are made up, at least in part. Ores and Concentrates Tested The analyses of the ores and concentrates tested are summarized in Table 1, which also includes the analyses of the calcines produced from the concentrates. The concentrates were roasted in a laboratory muffle furnace at the temperatures indicated until evolution of SO2 could no longer be detected; shallow clay dishes ordinarily were used; but calcines 2a and 2c were prepared in iron dishes, with the result that these two calcines were badly contaminated with iron scale. The "complex" concentrate, sample 1, consisted essentially of sphalerite grains containing inclusions of pyrite and other sulphide minerals too minute to be liberated by grinding within economic limits. In the " marmatitic" concentrate, sample 2, the iron was present largely as marmatite (that is, as iron sulphide in solid solution in zinc sulphide) rather than as pyrite inclusions. The zinc was present largely as the carbonate in the two "oxidized" ores, samples 3 and 4, and as the silicate in the "willemite" ore, sample 5. Leaching Tests The leaching cell was a 3-liter glass beaker placed in an electrically heated water bath manually controlled to maintain a temperature of 35°C in the leach slurry. The slurry was agitated by means of a flat paddle-type stirrer

Jan 1, 1950

By Robert D. Pehlke, Robert G. Blossey

The solubility of nitrogen in liquid iron alloys containing chromium and nickel has been measured in the temperature range 1550° to 1700°C at nitrogen pressures to 1 ah. The solubility surface has been determined for the iron corner of the liquid metallic ternary to 21 pct Cr and 11 pct Ni. The nitrogen solubility increases markedly with increasing chromium concentration. The effect of nickel and the bilinear effect of chromium and nickel are less pronounced. The solubility data are presented, and interaction coefficients are determined from them. THE solubility of nitrogen has been determined in liquid Fe-Cr-Ni binary and ternary alloys. The work in this laboratory has been directed toward a relatively limited composition range, that is to 20 pct Cr and 10 pct Ni. Humbert and Elliott1 studied the entire ternary; but their investigation was, understandably, limited in this region. The region considered here is of the greatest commercial interest. Knowledge of the solubility in these alloys is important for accurate control of the nitrogen levels in various stainless alloys. Also, the solution behavior is of theoretical interest because of its deviations from ideal dilute solution rules at moderate chromium levels. The nitrogen solubility in the iron binaries has been reviewed extensively by Pehlke and Elliott,2 particularly with respect to the dilute solution interaction parameters. More recently, there has been some interest in accurate descriptions of solute interactions in non-dilute solutions. Most of the investigators have adopted a power series fit to accommodate any non-linearity of the data. The Taylor series expansion is adopted here, using Scimar&apos;s notation3 for the interaction coefficients. An alternate method, instead of requiring higher-order expansions, is to restrict the range of the composition variables. A planar solubility surface may be fitted to the data and its range of validity determined. If the axes of this region are displaced from the usual origin—the solvent pure iron—a change in standard state is required for the solution reaction. The standard state for nitrogen in iron is commonly taken as the hypothetical 1 wt pct solution of nitrogen in pure iron. However, for an alloy of appreciable solute concentration, e.g., 13 pct Cr or 18 pct Cr-8 pct Ni, it would be convenient to use that base composition as the solvent and redefine the solution reaction in terms of that particular solvent. A brief derivation follows, defining the free energy of solution in an alloy solvent. Following the treatment by Darken and Gurry,4 for a nitrogen pressure of 1 atm: N (pure component) = N (1 wt pct in solvent) [l] AG=RT ln100/ wt pct N in solvent Taking this reaction for an alloy solvent less the same reaction for a pure iron solvent the relation between nitrogen in pure iron and an alloy is: N (1 wt pct in Fe) = N (1 wt pct in alloy) [2] AG = RT[ln(wt pct N in Fe) - ln(wt pct N in alloy)] Adding ½ N2(g) to each side and subtracting N (1 wt pct in Fe) = ½N2(g), the total free energy for solution of nitrogen in an alloy, referred to nitrogen at infinite dilution, is given by: ½N2 (g) = N (wt pct in alloy) L3] AG°3 = -RT ln(wt pct N in alloy )PN =1 atm Interaction parameters have their usual meaning when ?G°/RT is differentiated with respect to the proper composition coordinate. If desired, the dilute alloy treatment may be utilized with the new reference composition as the solvent. This has. indeed, been done by Small and pehlke5 in their analysis of nitrogen solution in liquid 18-8 alloys. The Sieverts&apos; technique was used for this investigation on an apparatus previously described.6 The materials used in this program were: Ferrovac-E iron, 99.95 pct, Crucible Steel Co.: chromium, 99.95 pct, Union Carbide Metals Co.; nickel, 99.9 pct, International Nickel Co. Recrystallized alumina crucibles were used with no evidence of crucible-melt reaction. RESULTS The locations of the ternary compositions investigated are shown in Fig. 1. The data of Turnock7 and small8 have been alluded to in other papers.5&apos;6 There were no departures from Sieverts&apos; law observed for

Jan 1, 1969

By J. C. Purcupile, R. L. Morris

Present methods of roof support require that first a hole must be drilled, and then a toggle belt or grouted resin bolt be set. Automating this system is very difficult. Therefore, we decided to develop a system which would insert pins into the roof in one step. We think that this type of system can be automated much easier. It would then speed up the roof bolting operation and eliminate the need for a man to work under unsupported roof, making roof bolting a much safer job. There are two ways that a pin can be inserted in a mine roof without drilling a hole. It can either be pushed in, using a very high, constant axial force; or it can be hammered in like a nail. The former method involves very high compressive stresses in the pin, and requires complicated machinery to prevent buckling. Since hammering the pin does not require high axial forces buckling is not as likely. Also, experiments have shown that there are no problems associated with getting the pin to go in straight. We chose to develop the second method. While investigating methods for hammering a pin into a mine roof, we learned that piles were being driven 10 times faster than by conventional methods with a method called "orboresonance." We decided to turn the pile driver upside down, and drive pins up into the mine roof. Sonic Pin Setting Machine The machine consists of a steel bar supported at the nodal points of its first mode of transverse vibration (Fig. 1). A mechanical oscillator, driven by a hydraulic motor, is attached to one end of the bar. This provides the forcing function for the vibrations in the bar. The hydraulic power is supplied by a portable unit. When the forcing function is near the natural frequency of the bar, large amplitudes of vibration occur and require little power input. An anvil is attached to the bar at a point of maximum deflection, providing a means for striking the pin. Since the bar is being driven near its natural frequency, it takes little energy to accelerate the bar back and forth between each hit. This means that most of the power being put into the bar is going into driving the pin. This is unlike conventional pile drivers which use most of their power to accelerate the hammer back and forth between hits. The resonant frequency of the bar we used is about 250 Hz, which is in the audible range, hence the term "sonic pin setting machine." The system benefits are: (1) elimination of the two-step drilling and pin-setting roof bolt operation, (2) increased speed, (3) more intimate contact between the formation and bolts, (4) less expensive roof bolts, (5) higher integrity, (6) potential for automatic operation. Modifications During the months from Oct. to Dec. 12, 1974, under U.S. Bureau of Mines grant No. 00144104, "Projects on Coal Extraction," six members of the senior class at Carnegie Mellon University designed and built the experimental machine. They drove three pins into the Safety Research coal mine at Bruceton, Pa., and determined that the project was worth continuing. In January 1975 we began making modifications and repairs in preparation for a more extensive program of testing at Bruceton. The major items were: 1) Design and build a collapsible pin guide to keep the pin parallel with the line of travel of the lift table. 2) Move the anvil to the center of the bar. 3) Refinish the nodal holes and install new node pins and bushings. 4) Realign the bar in its frame. 5) Install a relief valve on the lift unit to control the upward force. The upward force that we exert on the pin is relatively small, so there is no tendency for it to buckle. This means that all we need to do is hold the pin straight during its initial penetration. This allows the bushing to travel up and down in a straight line. At the beginning of penetration, the bushing

Jan 1, 1977

By James L. Poole

INTRODUCTION Companies in the mining industry are subjected at times to currency exchange rate risks and interest rate risks. The former occurs any time a firm deals in more than one currency. The latter occurs anytime a firm is financing itself with borrowings that have a floating interest rate (although a fixed rate can be a problem at times). FOREIGN EXCHANGE RISKS There are two general reasons why exchange rate risks (FX risk) can occur: First, a firm may be building a mine or plant in another country where financing for the plan is denominated in funds other than the country hosting the plant. For example, if a Canadian firm is building a smelter in another country which is being financed with Canadian dollars, the C$100 million for labor, which must be paid in the local currency, has subjected the firm to a C$100 million FX exposure risk. If after the cost estimates were made the local currency appreciates 30% against the Canadian dollar, the cost of labor has increased to C$130 million. (Of course, the local currency could also devalue during this period making the labor cost component less than estimated.) The other reason exchange rate risk arises is due to the fact that the costs of production in one country producing minerals may be incurred in one currency while the production is sold in another currency, such as coal being exported from Canada to Japan under contracts denominated in Yen. In this example, the Canadian firm is exporting coal to Japan and being paid in yen in 30 days at a contract price of 100 Y per ton. Furthermore, assume that the mining firm anticipates a depreciation of the yen against the Canadian dollar of 5% in the next 30 days. Therefore, the mining firm stands to lose 5% of the value of its shipment in a month&apos;s time not including carrying costs. In both cases, the mining company risks an adverse change in the exchange rate, that is, the currency the company holds may decrease in value relative to the other currency. There are a number of ways to mitigate the exposure to foreign exchange which can be utilized by the mining company. Optimally, a company should try to match all their inflows and all their outflows in the same currency. For example, if a Canadian firm has mining costs and debt service costs both denominated in Canadian dollars then 100% of their receipts would need to be in Canadian dollars. If a Canadian firm has 70% of its cash outflows in Canadian dollars and 30% of its cash outflow in American dollars in the form of a loan amortization then their receipts would need to be 70% in Canadian dollars and 30% in U.S. dollars. This general principle of matching cash receipts and cash costs in the same currency works fine in a perfect world but what of the real world when these cannot be matched? In actual practice, however, the financial markets have developed a number of ways in which foreign exchange risks can he managed or reduced. FORWARD MARKETS One method is to use the forward markets which are made on a world-wide basis by the commerical banks for foreign exchange. If a firm wanted to lock in the current exchange rate for the receipt of a foreign currency at some point in the future, they could do so by contracting on the forward market with a bank selling that foreign currency at the current rate at the future date. For example, assume that a Canadian firm is exporting coal to Japan which would pay yen in 90 days. If a firm wanted to lock in today&apos;s exchange rate they could contract with a bank to sell on the forward market yen in 90 days. When the yen was received the contract would be executed by selling the yen to the bank and receiving the previously agreed upon number of Canadian dollars. Generally speaking, the forward market can be used to sell forward about 12 months. The costs of selling (or buying) forward

Jan 1, 1985

By J. Washburn, R. Drouard, E. R. Parker

Temperature dependence of the rate of recovery in zinc single crystals after a simple shear deformation at low temperature was investigated. Some tentative suggestions regarding the annealed and strain-hardened states of a crystal are discussed. RECOVERY may be defined as the gradual return of the mechanical and physical properties of strain-hardened metal to those characteristic of the annealed material; an increase in temperature increases the rate of recovery. The annealing process in strain-hardened polycrystalline metals is complicated by the inhomogeneity of strain which always exists in aggregates. Polygonization in bent regions of the crystals and growth of new almost strain-free grains starting at points of severe local distortion1-:&apos; make it almost impossible to isolate and study the recovery process. Homogeneously strained single crystals, however, do not polygonize or re-crystallize and hence they can be used advantageously to study recovery. In such crystals strain hardening is completely removed by recovery alone. Since recovery is a process whereby certain lattice disturbances introduced by plastic flow are gradually reduced, a knowledge of the rate and temperature dependence of this process for various conditions of prestrain might be helpful in formulating a model of the strain-hardened state. For simplicity it seemed desirable to limit the type of prestrain to the simplest obtainable, i.e., simple shear strain. In the experiments to be described, recovery was studied by observing changes in the stress-strain curve of prestrained zinc single crystals held for various times at temperatures above that employed for straining. Single crystals were grown from the melt by a modified Bridgeman technique from Horse Head Special zinc 99.99 pct pure, and from spectrographically pure zinc 99.999 pct pure. They were grown as 1 in. diameter spheres and acid-machined&apos; to the final specimen contour. The test section was a cylinder about 1/8 in. high and 3/4 in. in diameter. The conical sections adjacent to the test section were cemented into the grips so the load could be transmitted to the crystal as uniformly as possible. The specimens were oriented so that in testing the maximum shear stress was applied along one of the slip directions, [2110], in the (0001) plane. Details of the production and testing of such specimens have been presented.&apos; Each test was carried out according to the following schedule: 1—The crystal was strained at — 50°C until it reached a maximum shear stress, ,,,. The strain rate was approximately 5 pct per min in all cases. 2—After straining, the crystal was unloaded before the temperature was changed. Unloading required about 3 min. 3—The temperature of the specimen was then increased from — 50°C to the temperature, T, of recovery. This change in temperature was completed in a time of less than 2 min. The specimen remained at temperature, T, for a time, t, which differed for the various specimens. 4—Thereafter the temperature was again reduced to — 50 °C in approximately 3 min. 5—While at —50°C, the stress-strain curve after recovery was obtained. 6—The specimen was then unloaded and annealed for 1 hr at 375 °C in a helium atmosphere to bring about complete recovery. Cooling to room temperature after anneal required 90 min. 7—The same crystal could be re-used for another test because the plastic properties after annealing closely duplicated those of the original crystal. The specimen was immersed during the test in a bath of methyl alcohol which, through a system of tubes, could be pumped through either of two heat exchangers to regulate the temperature; this was accomplished by circulating the liquid through coils immersed in a bath of acetone and dry ice for cooling or in a bath of warm water for heating. Test temperatures were thus maintained constant within ±1°C. The — 50°C temperature was low enough so that no measurable recovery occurred during unloading and reloading. The stress-strain curve continued after recovery along a path below, but approximately parallel to, the path of a curve obtained in an uninterrupted test. Fig. 1 shows some of the results from a specimen of 99.999 pct Zn. The amount of downward displacement of the curve due to recovery was a

Jan 1, 1954

By D. S. Bloom

RECENTLY a report has been published on an investigation of the MO-RU system by E. Raub.&apos; In this report it is stated that below approximately 1200°C the system consists of two terminal solid-solution phases and the intervening two-phase field; at 1200°C and above, an intermetallic compound labeled Mo,Ru, makes an appearance. Along with some other little known systems in which it was considered possible for a u phase to form, the Mo-Ru system had been under investigation in this laboratory. Since molybdenum and ruthenium have very high melting points and also since ruthenium is expensive, the investigation was confined to X-ray diffraction by sintered powder compacts. The compositions which were studied were (in atomic percentages): 75 Mo-25 Ru, 67 Mo-33 Ru, 50 Mo-50 Ru, and 25 Mo-75 Ru. All compacts were made from the pure powders, and no chemical analyses were attempted. Each compact weighed about 5 g. All annealing was done in evacuated Vycor capsules, and even though the capsules completely collapsed at 1250°C, they still admitted no air. The results of this investigation in general corroborate the work of Raub. Up to 1150°C the two terminal solutions and the intervening two-phase field were found, and at 1200°C and above an intermetallic compound was found. However, the composition limits of the compound were estimated to lie close to 70 atomic pct Mo rather than 62.5 atomic pct Mo as Raub had reported. At least the X-ray diffraction results indicated that the 75 atomic pct Mo specimen consisted of the molybdenum solid solution plus the compound, while the 67 atomic pct Mo specimen consisted of the ruthenium solid solution plus the compound and the equiatomic composition contained only very little of the compound. The most interesting observation, however, was that the diffraction pattern of the compound was similar to that of the well known u phases, except that unusually large lattice constants were indicated. The d-values, or interplanar spacings, of the diffraction lines (as produced by filtered copper KCX radiation) of the compound are given in Table I, as determined from specimens annealed at 1250°C for 9 hr and then quenched in water. Also shown are the lines of the Mn-Mo u phase as given by Decker, Waterstrat, and Kasper,&apos; the lines of the Fe-Mo a as given by Goldschmidt," and the lines of the Mo-Ru phase multiplied by constants in order to make comparison with the Mn-Mo and Fe-Mo patterns easier. These results indicate that the compound in the Mo-Ru system is apparently isomorphous with the other o- compounds. The rather large lattice constants of the Mo-Ru a phase are acceptable in view of the comparatively large size of the constituent atoms. The data in Table I for the Mo-Ru compound check quite well with those given by Raub, not shown here, though there are some discrepancies between the two sets of data in the estimated intensities of a few lines. Although very little is known at present about this u phase, there are some things which may be pointed out, mainly in comparison with other better known u phases. Even though molybdenum and ruthenium are in the same columns in the Table of Elements as chromium and iron, respectively, the Mo-Ru u phase contains much more molybdenum than ruthenium, while the Cr-Fe u phase centers around or very nearly around the equiatomic composition; furthermore, the Mo-Ru u is stable only above approximately 1200°C, but the Cr-Fe u is stable only below about 825 °C. With respect to its being unstable below 1200°C, the Mo-Ru u is similar to the other binary molybdenum u phases in the Co-Mo and Fe-Mo systems. Similar to the latter two systems, the Mo-Ru u apparently exists over a rather narrow composition range, whereas the Cr-Fe u at lower temperatures is stable over a much wider composition range. The Mo-Ru u thus is closer to the other molybdenum u phases than to the Cr-Fe cr phase. The most striking feature of the Mo-Ru u phase, however, is the absence of any of the transition elements of the First Long Period.

Jan 1, 1956

By Gearge C. Wallick

This paper describes experimental procedures for the calibration of capillary tubes to be employed as comparison standards in gas flow-rate measurements and considers several types of flow which were observed in the calibration of seven capillaries of varying diameter. It is shown that under certain experimental conditions deviations from Poiseuille flow are observed which may not be attributed to turbulence, and that this flow behavior may be described empirically by introducing the classical kinetic energy correction into the Poiseuille equation. INTRODUCTION Techniques have been developed for the measurement of the volume rate of flow of nitrogen through glass capillary tubes as a function both of the differential pressure applied and of the absolute gas pressure. Capillary tubes thus calibrated are frequently used to determine the rate of gas flow through a porous sample in the measurement of slip-corrected permeability. The purpose of this note is to describe the calibration procedure and to discuss some observed deviations from the ideal flow behavior predicted by the Poiseuille equation. EXPERIMENTAL PROCEDURE The volume flow rate of nitrogen as a function of the differential pressure was determined experimentally for a series of seven glass capillary tubes of varying diameter. As is shown in Table 1, the capillary radii in the series ranged from 0.004 cm for K1 to 0.07 cm for K7. The capillary tubes were cut from capillary tubing purchased on special order from several manufacturers and the bore radii were determined by microscopic measurement. The gas volume flow rate was measured for each capillary tube at a series of differential pressures in the range from 0 to 3 cm of mercury and at average absolute pressures of 1.0, 2.5, 4.0, and 6.0 atm. Experimental Procedure for Small Capillaries A schematic diagram of the apparatus used to calibrate the three smallest capillaries, Kl, K2, and K3, is shown in Fig. 1. By means of a standard Ruska Instrument Corp. volumetric pump (equipped with a synchronous motor drive, two stages of gear reduction, and a quick-change lathe gear box), nitrogen was displaced by mercury from a metal cell immersed in a constant-temperature bath. The pump, the line to the cell, and a part of the cell were filled with mercury. As mercury was forced into the cell, it displaced nitrogen which passed through the capillary tube being calibrated and into a large container with a volume of approximately 30 liters. Since the volume of this container was very large compared to the total volume change of the gas resulting from the action of the displacement pump, the absolute pressure of the system was essentially constant throughout any one run. Prior to any determination this pressure could be set at any desired level by means of an auxiliary nitrogen cylinder. The differential pressure across the capillary was measured with a Meriam Red Oil manometer and recorded together with the corresponding pump displacement rate which was equal to the gas flow rate. By repeating this process for a series of different displacement rates at each of the absolute pressures of interest, the data necessary for plotting the calibration curves for capillaries K1, K2, and K3 were obtained. All calibration measurements were made in a constant temperature room to minimize errors due to temperature fluctuations. In each run. the flow was maintained until equilibrium was established. Experimental Procedures for Large Capillaries The flow rates through Capillaries K4, K5, K6, and K7 required to obtain the pressure drops desired were larger than could be obtained from available constant displacement pumps, and it was thus necessary to employ an alternative calibration procedure. The arrangement of apparatus used in calibrating these capillaries at an average pressure of one atmosphere is shown in Fig. 2. Each capillary was calibrated by flowing nitrogen through the capillary and recording the differential pressure as a function of the flow rate. &apos;The differential pressure was measured with a Meriam Red Oil manometer and the flow rate was determined by collecting the effluent gas over water at constant pressure and recording the increase in volume per unit time after an equilibrium had been established. Appropriate corrections were made for the

Jan 1, 1953