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By W. Philippoff

ALTHOUGH Gaudin1 and more recently Sutherland2 have calculated the probability of collision of a falling mineral particle with a rising bubble, there is no published information concerning the details of the mechanism of attachment of a collector-coated particle to a bubble. During the past year the writer has developed a theory for the mechanism of attachment, which has been substantiated experimentally. Funds for the investigation and for some of the equipment used have been supplied by the Mines Experiment Station of the University of Minnesota. Motion picture studies of the phenomena involved in the collision between mineral particles and bubbles, such as those of Spedden and Hannan," show that the contact can be completed within 0.3 millisec. Formulas developed for rigid bodies have hitherto been used' for the calculation of the motion of bubble and particle, but it is obvious that a bubble cannot be regarded as a rigid body. On the contrary, Spedden and Hannan's pictures show a great degree of deformation during the collision. The time of attachment was calculated as the time the particle drifted past the bubble. Time of Collision The theory presented in this paper enables calculation of the time of collision, using the concept that the bubble, or more generally, a liquid-air interface, acts as an elastic body. The elasticity, defined as the restoring force on a mechanical deformation, is caused by the surface tension and is the result of the principle of the minimum of free surface energy. It is well known that an elasticity together with a mass determines a frequency of vibration. The vibrations of jets and drops caused by the elasticity of the interface are known to comply exactly with the classical theory of capillarity.' However, the vibrations of isolated bubbles, as distinct from foams, have not been investigated previously. The following equation, presented elsewhere,' has been deduced for these frequencies: 3_____________________ fB = 9.20.vV.vn. (n-1).(n+2)/8 [I] in which fB is the frequency of a harmonic of the bubble in cycles per second, V the volume of the bubble in cc, n a number determining the order of the harmonic, and n = 2 the basic vibration. The first (basic) harmonic describes a change of the spherical bubble to an ellipsoidal bubble. The higher harmonics are more complicated, for the circumference of the bubble is divided approximately into as many parts as the order of the harmonic. As an example, Spedden and Hannan's published motion picture of a vibrating bubble corresponds to the sixth harmonic. Eq 1 shows that only the first and third harmonics are simple multiples (1 and 3), all the others being irrational fractions of the basic frequency. This means that the shape of the vibration can change with time and is in general unsym-metric in respect to the time axis. Such conditions prevail when there is a distributed elasticity or mass, as in the case of vibrating membranes or rods. The constant 9.20 is valid for water at room temperature, but a general solution involving the physical constants of the liquid has not been found. The case of the floating particle is much easier to treat than that of the bubble. It can be assumed that the elasticity is caused exclusively by the interface and that the mass is concentrated in the particle together with some adhering water. The following expression for the frequency of a system of one degree of freedom can be applied: fP = -1/2-vE/m [2] Here fp is the frequency of the particle vibration in cycles per second, E the elasticity in dynes per cm, and m the mass in grams. The classical theory of impact phenomena gives the time of collision during the striking of a spring (in this case the surface of the bubble) by a mass, as: t, = 2/f = nvm/E [3] It is now possible to develop an expression for the elasticity of a floating cylindrical particle. The force equilibrium of a cylinder floating end on at the air-liquid interface is given by the well-known equation (Poisson7 1831) P = /4 D2-pL-g-h + D.y sin a [4] which accounts for the buoyancy and the action of the surface tension where P is the force acting on the particle in dynes (weight-buoyancy), D the diameter of the cylinder in cm, PL the density of the liquid in grams per cc, g the acceleration of gravity = 981 cm per sec2, h the depression of the cylinder below the surface of the liquid in cm, y the surface tension in dynes per cm and a the supporting angles or the one required to insure equilibrium, a being smaller than the contact angle 0. Although demonstrated by Poisson, it has not

Jan 1, 1953

By W. Philippoff

ALTHOUGH Gaudin1 and more recently Sutherland2 have calculated the probability of collision of a falling mineral particle with a rising bubble, there is no published information concerning the details of the mechanism of attachment of a collector-coated particle to a bubble. During the past year the writer has developed a theory for the mechanism of attachment, which has been substantiated experimentally.' Funds for the investigation and for some of the equipment used have been supplied by the Mines Experiment Station of the University of Minnesota. Motion picture studies of the phenomena involved in the collision between mineral particles and bubbles, such as those of Spedden and Hannan,3 show that the contact can be completed within 0.3 millisec. Formulas developed for rigid bodies have hitherto been used' for the calculation of the motion of bubble and particle, but it is obvious that a bubble cannot be regarded as a rigid body. On the contrary, Spedden and Hannan's pictures show a great degree of deformation during the collision. The time of attachment was calculated as the time the particle drifted past the bubble. Time of Collision The theory presented in this paper enables calculation of the time of collision; using the concept that the bubble, or more generally, a liquid-air interface, acts as an elastic body. The elasticity, defined as the restoring force on a mechanical deformation, is caused by, the surface tension and is the result of the principle of the minimum of free surface energy. It is well known that an elasticity together with a mass determines a frequency of vibration. The vibrations of jets and drops caused by the elasticity of the interface are known to comply exactly with the classical theory of capillarity.5 However, the vibrations of isolated bubbles, as distinct from foams, have not been investigated previously. The following equation, presented elsewhere,' has been deduced for these frequencies: [3fB = 9.20•'./V•Vn- (n-1) • (n+2) /8[1]] in which fB is the frequency of a harmonic of the bubble in cycles per second, V the volume of the bubble in cc, n a number determining the order of the harmonic, and n = 2 the basic vibration. The first (basic) harmonic describes a change of the spherical bubble to an ellipsoidal bubble. The higher harmonics are more complicated, for the circumference of the bubble is divided approximately into as many parts' as the order of the harmonic. As an example, Spedden and Hannan's published motion picture of, a vibrating bubble corresponds to the sixth harmonic. Eq 1 shows that only the first and third harmonics are simple multiples (1 and 3), all the others being irrational fractions of the basic frequency. This means that the shape of the vibration can change with time and is in general unsymmetric in respect to the time axis. Such conditions prevail when there is a distributed elasticity or mass, as in the case of vibrating membranes or rods. The constant 9.20 is valid for water at room temperature, but a general solution involving the physical constants of the liquid has not been found. The case of the floating particle is much easier to treat I than that of the bubble. It can be assumed that the elasticity is caused exclusively by the interface and that the mass is concentrated in the particle together with some adhering water. The following expression for the frequency of a system, of one degree of freedom can be applied: [1E/m[2] fP = 27] Here f, is the frequency of the particle vibration in cycles per second, E the elasticity in dynes per cm, and m the mass in grams. The classical theory of impact phenomena gives the time of collision during the striking of a spring (in this case the surface of the bubble) by a mass, as: [t~ = 2/f = 7r\/m/E[3]] It is now possible to develop an expression for the elasticity of a floating cylindrical particle. The force equilibrium of a cylinder floating end on at the air-liquid interface is given by the well-known equation (Poisson' 1831) [aP = 4 D2.pL•g•h +7rD•y sin a[4]] which accounts for the buoyancy and the action of the surface tension where P is the force acting on the particle in dynes (weight-buoyancy), D the diameter of the cylinder in cm, pL the density of the liquid in grams per cc, g the acceleration of gravity = 981 cm per sec2, h the depression of the cylinder below the surface of-the liquid in cm, y the surface tension in dynes per cm and a the supporting angle' or the one required to insure equilibrium, a being smaller than the contact angle ?. Although demonstrated by Poisson, it has not

Jan 1, 1952

By N. S. Stoloff, R. G. Davies

A study has been made of the efject oj different dislocation-precipitate interactions upon the temperature dependence of the flow stress of aged Ni-14 at. pct A1 alloy. It is observed that when the dislocations bow between widely spaced (-20004 coherent Ni3Al particles the flow stress decreases with increasing temperature in the normal way. However, when the dislocations cut closely spaced (-5004 particles the flow stress is independent of temperature from -100 to 600°C, due to a balance between softening of the matrix and an increase in strength of the particles with increasing temperature. The retention of strength at high tempera-tures of commercial nickel-base alloys, which are strengthened by the precipitation of a phase based upon Ni3Al, is thought to be due to the unusual strength properties of Ni3Al. The flow stress of Ni3Al increases continuous1y from -196"C to a maximum at -600"C. It is concluded from a series of thermal-mechanical tests that the sevenfold increase in flow stress over this temperature interval is due to a lattice effect and is not diffusion-controlled. The flow stress of precipitation- or dispersion-hardened materials depends on the resistance to dislocation motion within the matrix and the extra energy required for dislocations to bow between or to cut particles. If the dislocations bow between the particles or if the strength of the cut particles is constant with temperature, then the flow stress of the precipitation-hardened alloy must decrease with increasing temperature due at least to the decrease in elastic modulus of the material. There will be softening also from thermally activated cross-slip or climb, offering an additional degree of freedom for dislocations to avoid particles. For example, in the case of nickel containing a dispersion of thoria,' which most probably deforms by dislocations bowing between particles, the flow stress decreases by about 50 pct between 25" and 650°C. In A1-Cu alloys2 aged to produce the 8" precipitate, dislocations cut the particles, and the flow stress decreases by about 20 pct between -269" and 25°C. However, many commercial high-temperature nickel-base alloys, for example Inconel-X and Udimet-700, exhibit little or no decrease in flow stress with increasing temperature up to about 700°C. A characteristic feature of these alloys is that they are strengthened by the precipitation of a phase based upon Ni3A1. Guard and westbrook4 and flinn' have shown that Ni3Al (and alloys in which a third element such as molybdenum or iron is substituted for part of the aluminum) is unusual in that the hardness and flow stress increase with temperature to a maximum at about 600°C. For the flow stress of a precipitation-hardened alloy to be independent of temperature we propose that the particles must be cut by dislocations moving through the matrix and that the strength of the particle must increase with increasing temperature. Theories of precipitation hardening do not take into account the flow stress of the dispersed particles that are cut during deformation; the only dissipative process usually considered7 is the creation of interface within the particle and between the precipitate and matrix. The purpose of the present investigation has been to study in detail the temperature dependence of the flow stress of a nickel-base alloy strengthened by the precipitation of Ni3Al in two structural conditions such that when deformation occurs it does so by dislocations a) bowing between the particles and b) cutting the particles, respectively. A simple binary Ni-14 at. pct A1 alloy was chosen because considerable information is already available for this system concerning phase equilibria and precipitation reactions and rates.' Dislocation-precipitate interactions in the binary alloy should be similar to those in the more complex commercial alloys. In addition, the mechanical and physical properties of NisAl were studied in detail in the hope of elucidating the mechanism by which the strength increases with increasing temperature up to 600°C. EXPERIMENTAL PROCEDURE For the study of the effect of precipitation of Ni3A1 upon the temperature dependence of the flow stress, an alloy containing 14 at. pct A1 was utilized; a Ni-8 at. pct A1 solid-solution alloy was employed as a comparison material. Vacuum-cast ingots were hot-rolled at 1000°C and cylindrical compression samples, 0.20 in. diam by 0.40 in. high, were prepared from the 1/4-in.-diam rod. Specimens were recrystallized and solution-treated at 1000°C for 1/2 hr and then water-quenched. A preliminary study revealed that, when the Ni-14 at. pct A1 alloy was aged for 1 hr at 700°C, significant precipitation hardening was obtained, and that the structure was free from grain boundary discontinuous precipitation; an overaged condition was produced by annealing the aged specimens at 850°C for 1 hr. To circumvent the difficulties involved in the hot rolling and swaging of Ni3A1, compression samples,

Jan 1, 1965

By J. L. Huitt, B. B. McGlothlin, D. K. Lowe

Published correlations for the principal aspects of hydraulic fracturing were combined into a digital computer program to facilitate the study of interrelated variables. The computer program includes individual relationships for fracture width during pumping, fracture area generated, propping agent embedment, flow capacities of propped fractures and transport of propping agents in horizontal fractures. The effects of more than 20 treatment and formation parameters on the predicted results of hydraulic fracturing treatments were studied. The effects of these parameters were determined for (I) fracture width during injection, (2) fracture width after the overburden comes to rest on the propping agents, assumed not to be crushed, (3) generated and propped fracture area, (4) location and concentration of propping agents in the fracture when injection ceases, (5) flow capacities of the various propped sections of the fracture and (6) expected increase in the well productivity. The effects of propping agent, formation and fracturing fluid parameters on well productivity are discussed. The parameters that were found to have the most pronounced effects on hydraulic fracturing treat~nents are injection rate, treatment volume, fracturing fluid coefficient, size and amount of propping agent, spearhead volume, well drainage radius and formation capacity. INTRODUCTION Many correlations have been published for predicting effects of various parameters that are considered in the design of hydraulic fracturing treatments. The Carter equation' can be used to predict generated fracture radius as a function of fracture width, fracturing fluid leakoff and other parameters. Fracture width can be determined by use of the Perkins and Kern correlation' in which the fracture width is related to the fracture radius, fluid injection rate and certain formation and fracturing fluid parameters. Wahl and Lowe et aL4 have reported methods of predicting the location of propping agents in fractures when pumping ceases. The former study is applicable to the case where the ratio of propping agent diameter to fracture width is less than 0.1. The latter is applicable when this ratio is greater than 0.1. These studies showed that the propping agent placement in horizontal-radial fractures depends principally on how the individual particles are transported in the fracture by the carrying fluid. Particle transport in fractures is determined by local fluid velocity in the fracture, fluid and particle properties, and the size of the particle relative to the fracture width. The distribution of propping agents, effective overburden pressure and formation rock strength control the propped fracture width6 by controlling the extent to which the propping agent particles embed into the fracture faces. From the distribution of propping agents and the propped fracture width, fracture flow capacities can be calculated or the various regions of the fracture. The flow capacities and the radial extent of these regions can be combined with reservoir information to predict the productivity increases for fractured wells. In all these studies, the effects of certain treatment and/ or reservoir parameters on one facet of fracturing can be predicted only if other facets which the parameters affect are fixed. For instance, fracture width and radius are interrelated; that is, to calculate the value of one, the value of the other must be known. Also, some parameters influence more than one aspect of fracturing. For example, prop-pant transport is a function of both fracture width and fluid viscosity. but fracture width is itself a function of fluid viscosity. Since these calculations are complex and the parameters interrelated, it is not possible to write an equation with which the over-all effects of treatment parameters can be solved explicitly. For these reasons, the correlations for determining the effects of the parameters which are most significant in hydraulic fracturing treatments have been incorporated into a digital computer program. COMPUTER PROGRAM The program, which was written for an IBM 7094 computer, can be used to predict results of most of the combinations and values for the treatment parameters that are ordinarily considered for fracturing treatments. A spearhead of fracturing fluid and a propping agent-carrying fluid with different fluid properties can be taken into account. Also, the total volumes and relative amounts of the spearhead and carrying fluids can be varied. Two different propping agents (as used in tail-in operations) and a wide range of formation properties and injection rates are considered. The computer program (Fig. 12) consists of several sets of calculations. First, the final flooded fracture radius and average fracture width at the cessation of pumping are calculated. This is done by simultaneous solution of the Perkins and Kern fracture width equation and the Carter equation for flooded fracture radius (equations used in the computer program appear in the Appendix). The next step is to determine local fluid velocity in the fracture as a function of time and radius. Since it is not possible to write this function in closed form expression, a table of velocity values is generated by the program and stored for subsequent use. The time span from the beginning of

By R. L. Bell, T. G. Langdon

A technique was developed- for determining the grain strain, and hence the grain boundary sliding contribution, occurring during the high- temperature creep of a magnesium alloy, from the distortion of a grid photographically printed on the specimen surface. The results were compared with those obtained from measurements of grain shape, both at the surface and interrwlly, and it was concluded that the grain shape technique may substantially underestimate the grain strain and overestimate the sliding contribution due to the tendency for migration to spheroidize the grains. ALTHOUGH a considerable volume of work has been published on the role of grain boundary sliding in high-temperature creep, many of the estimates of Egb (the contribution of grain boundary sliding to the total strain) have been in error due to the use of incorrect formulas or inadequate averaging procedures.' One of the most easy and convenient measurements from which to compute Egb is that of v, the step normal to the surface where a grain boundary is incident. Unfortunately, this parameter is also the one associated with the treatest number of pitfalls. Values of v have been used to calculate Egb from the equation: egb =knrVr [1] where k is a geometrical averaging factor, n is the number of grains per unit length before deformation, v is the average value of v, and the subscript ,r denotes the procedure of averaging along a number of randomly directed lines. If the dependence of sliding on stress were assumed, it would be possible, in principle, to calculate k from the known distribution of angles between boundaries and the surface. This in itself is difficult because the distribution depends on the history of the surface,' but the problem is even further complicated by the fact that v depends on other factors such as the unbalanced pressure from subsurface grains.3 However, the great simplicity of the measurement procedure for v makes it highly desirable that this problem of k determination should be overcome. In the present experiments, this was achieved by the use of an indirect empirical method in which the grain strain, eg, at the surface was determined by the use of a photographically printed grid. The assumption here is that the total strain, et, is simply the sum of that due to grain boundary sliding, egb, and that due to slip or other processes within the grains, eg. SO that: Et = Eg + Egb [2] Thus k is given by: In practice, it is customary to indicate the importance of sliding by expressing it as a percentage of the total creep strain; this quantity is termed y (= 100Egb/Et). The determination of Eg from a printed grid within the grains avoids the difficulties due to boundary migration which should be considered when the grain strain is calculated from measurements of the average grain shape before and after deformation. As first pointed out by Rachinger,4,5 however, this latter technique has the particular advantage that it can also be applied in the interior of a polycrystal. Recently, several workers have produced evidence on a variety of materials6-'' to support the observation, first made by Rachinger on aluminum,4,5 that 7 can be very high, 70 to 100 pct, in the interior, even when the surface value, determined from boundary offsets, is very much lower.10'11 Although there have been criticisms both of the shortcomings of the grain shape technique'' and of the different procedures used to determine y at the surface,' it seemed important to check whether measurements of sliding by grain shape gave values of y which were truly representative of the material. In the present experiments, grain shape measurements were therefore made both at the surface and in the interior for comparison with one another and with the independent measurements of grain strain using the surface grid technique. EXPERIMENTAL TECHNIQUES The material used in this investigation was Magnox AL80, a Mg-0.78 wt pct A1 alloy supplied by Magnesium Elektron Ltd., Manchester. Tensile specimens, about 7 cm in length, were prepared from a 1.27-cm-diam rod, with two parallel longitudinal flat faces each approximately 3 cm in length. The specimens were annealed for 2 hr in an oxygen-free capsule, at temperatures in the range 430° to 540°C, to give varying grain sizes, and, prior to testing, the grain size of each was carefully determined using the linear intercept method. This revealed that the grains were elongated -0.5 to 5 pct in the longitudinal direction. Testing was carried out in Dennison Model T47E machines under constant load at temperatures in the range 150" to 300°C. At temperatures of 200°C and below, tests were conducted in air with the polished flat faces coated with a thin film of silicone oil to prevent oxidation; at higher temperatures, an argon atmosphere was used. To determine v,, each test was interrupted at regular increments of strain and the specimen removed from the machine. At the lower strains, when v, was less than about 1 pm, measurements were taken on a Zeiss Linnik interference microscope;

Jan 1, 1969

By J. J. Gilman, T. A. Read

FOLLOWING the deformation 01 zinc monocrys-tals, sharply bent basal planes are observed near several types of inhomogeneities. Three of these in-homogeneities have characteristics which are quite regular so that they can be studied and analyzed. These are compressive kink bands, "deformation bands," and the inhomogeneities near end restraints. The present paper describes experiments in which "deformation bands" were artificially produced, and bend plane phenomena are discussed in terms of dislocation theory. Also, two new bend plane phenomena are described. The importance of bend plane phenomena in the deformation of crystals is not widely recognized. Many phenomena may be explained in a manner similar to the discussion in this paper. Jillson1 has pointed out that the "punching effect" in zinc is a bend plane phenomenon and is not caused by prismatic slip.' Bowles3 has suggested that they may be involved in diffusionless phase changes. Cahn4 has discussed the role of bend plane formation in the polygonization of zinc. Experimental Work Tensile Kink Bands: Because of the geometrical similarity between "deformation bands" and "kink bands" (compare Fig. 1 of this paper with Fig. 1 of the paper by Hess and Barrett"), the band shown in Fig. 1 of this paper will be called a "tensile kink band," and that shown by Hess and Barrett will be called a "compressive kink band." It is felt that the term "deformation band" should be reserved for banded structures in polycrystalline materials such as iron." Tensile kink bands seem to form spontaneously in aluminum crystals deformed by tensile loading.7-10 In zinc and cadmium crystals they do not form in good, carefully loaded specimens.'." However, tensile kink bands can be produced artificially in zinc crystals. The present authors did this by scratching one of the flat surfaces of triangular crystals transversely with a sharp needle. Natural tensile kink bands caused by inhomogeneities sometimes appeared in deformed crystals which were identical in appearance with the artificially produced ones. Zinc monocrystals were grown by the Bridgman method in graphite molds. Chemically pure zinc (99.999+ pct Zn) was used and the molds were sealed inside evacuated pyrex tubes during growth. The crystal cross sections were equilateral triangles with a typical base of 0.210 in. The artificial kink band shown in Fig. 1 is typical of tensile kink bands in zinc. The band lies between two bend planes which run from upper right to lower left and is inclined oppositely to the slip bands which are sharply bent at the two bend planes. The general form of the artificial tensile kink bands was independent of the scratch depth (1 to 5 mils deep) and also independent of which side was scratched. These variables did cause variations, however. Deep scratches produced more localized kink bands than light scratches. Also, if the angle between the slip plane traces and a transverse scratch varied appreciably among the three sides, then localization of the resulting kink bands also varied. Furthermore, if the slip direction lay nearly parallel to the scratched side, the band was more developed near the scratched side than at the opposite edge. Scratches produced tensile kink bands for crystal orientations from xo = 15" to x, = 75". Fig. 2 shows a scratched crystal after deformation. One triangular side lies in the plane of the photograph. The right hand tensile kink band was produced by a transverse scratch on the upper right side. The next two kink bands were the result of scratches on the front surface. The kink band at the left was caused by a scratch on the lower back side. All four bands have the same general form. A longitudinal scratch was also made on the crystal shown in Fig. 2 to determine the effect of a scratch on the critical shear stress. The critical shear stress of the scratched region was 33.9 g per sq mm compared to 24.4 g per sq mm for the un-scratched region above it. Fig. 3 shows Laue patterns of the crystal shown in Fig. 2. Fig. 3a shows the pattern of the undeformed crystal. The orientation was x, = 21°, A, = 31". After deformation, Fig. 3b was made of the homogeneously deformed portion of the crystal. The spots are compact but split into two halves. This region was elongated 45 pct and its orientation was x = 14", X = 20"; the sine law predicts x = 14", A = 20.5". Fig. 3c was taken near the center of the middle tensile kink band of Fig. 2. The pattern shows a range of orientations and polygonization in this region. The spread in orientation was due to the fact that the basal planes were curved (see Fig. 1) rather than flat as in the ideal case. Some may also have been the result of elastic distortions and "local curvatures." The orientation range was x = 23" to 32", A = 30" to 42". It is apparent from Fig. 3 that the material inside and outside the kink band rotated in opposite directions with respect to the tension axis during deformation. The orientation calculated from the ideal configuration of Fig. 9,

Jan 1, 1954

By J. D. Miller, J. E. Sepulveda

Tar sand deposits in the state of Utah contain more than 25 billion bbl of in-place bitumen. Although 30 times smaller than the well-known Athabasca tar sands, Utah tar sands do represent a significant domestic energy resource comparable to the national crude oil reserves (31.3 billion bbl). Based upon a detailed analysis of the physical and chemical properties of both the bitumen and the sand, a hot-water separation process for Utah tar sands is currently being developed in our laboratories at the University of Utah. This process involves intense agitation of the tar sand in a hot caustic solution and subsequent separation of the bitumen by a modified froth flotation technique. Experimental results with an Asphalt Ridge, Utah, tar sand sample indicated that percent solids and caustic concentration were the two most important variables controlling the performance of the digestion stage. These variables were identified by means of an experimental factorial design, in which coefficients of separation greater than 0.90 were realized. Although preliminary in nature, the experimental evidence' gathered in this investigation seems to indicate that a hot-water separation process for Utah tar sands would allow for the efficient utilization of this important energy resource. The projected increase in the ever-widening gap between the domestic energy demand and the domestic energy supply for the next few years has motivated renewed interest in energy sources other than petroleum, such as tar sands, oil shale and coal. Although a number of research programs on the exploitation of national coal and oil shale resources have already been completed, very few programs have been initiated on the processing of tar sand resources in the United States. In recognition of their significance as a domestic energy resource, investigators at the University of Utah have designed an extensive research program on Utah tar sands. An important phase of this program, and the main subject of this publication, is the development of a hot-water process for the recovery of bitumen from Utah tar sands, as a preliminary step toward the production of synthetic fuels and petrochemicals. The term "tar sand" refers to a consolidated mixture of bitumen (tar) and sand. The sand in tar sand is mostly a-quartz as determined from X-ray diffraction patterns. Alternate names for "tar sands" are "oil sands" and "bituminous sands." The latter is technically correct and in that sense provides an adequate description. Tar sand deposits occur throughout the world, often in the same geographical areas as petroleum deposits. Significantly large tar sand deposits have been identified and mapped in Canada, Venezuela and, the United States. By far, the largest deposit is the Athabasca tar sands in the Province of Alberta, Canada. According to the Alberta Energy Resources Conservation Board (AERCB),2,3 proved reserves of crude in-place bitumen in the Athabasca region amount to almost 900 billion bbl. To date, this is the only tar sand deposit in the world being mined and processed for the recovery of petroleum products. Great Canadian Oil Sands, Ltd. (GCOS) produces 20 million bbl of synthetic crude oil per year. Another plant being constructed by Syncrude Canada, Ltd. is expected to produce in excess of 40 million bbl of synthetic crude oil per year. According to the Utah Geological and Mineral Survey (UGMS), tar sand deposits in the state of Utah contain more than 25 billion bbl of bitumen in place, which represent almost 95% of the total mapped resources in the United States.4 The extent of Utah tar sand reserves seems small compared to the enormous potential of Canadian tar sands. Nevertheless, Utah tar sand reserves do represent a significant energy resource comparable to the United States crude oil proved reserves of 31.3 billion bbl in 1976.5 Tar sands in Utah occur in 51 deposits along the eastern side of the state.4 However, only six out of these 51 deposits are worthy of any practical consideration (Fig. 1). As indicated in Table 1, Tar Sand Triangle is the largest deposit in the state and contains about half of the total mapped resources. Information regarding the grade or bitumen content of Utah deposits is still very limited. The bitumen content varies significantly from deposit to deposit, as well as within a given deposit. In any event, the information available6-8 seems to indicate that Utah deposits are not as rich in bitumen as the vast Canadian deposits which average 12 to 13% by weight.9 Although many occurrences of bitumen saturation up to 17% by weight have been detected in the northeastern part of the state (Asphalt Ridge and P. R. Spring), the average for reserves in Utah may well be less than 10% by weight. Separation Technology As in any other mining problem, there are two basic approaches to the recovery of bitumen from tar sands. In one

Jan 1, 1979

By A. J. Shaler, H. Udin, J. Wulff

In the study of the sintering of meta powders, we have come to the conclusion in this laboratory that further progress requires a more basic understanding of the operating mechanisms. This is emphasized in detail by Shaler. He has shown that a knowledge of the exact value of the surface tension is imperative for a solution of the kinetics of sintering. This force plays a principal role in causing the density of compacts to increase.2 Furthermore, a knowledge of the surface tension of solids is also applicable to other aspects of physical metallurgy. C. S. Smith3 points out the relation between surface and interfacial tension and their function in determining the microstructure and resulting properties of polycrystal-line and polyphase alloys. This paper describes one group of results of an experimental program designed for the study of the surface tension in solid metals. As a by-product of this work, considerable information has been obtained on the rate and nature of the flow of a metal at temperatures approaching the melting point and under extremely low stresses, a field of mechanical behavior heretofore scarcely touched by metallurgists. The importance of this additional information to students of powder metallurgy need not be stressed. Theoretical Considerations Interfacial tension arises from the condition that an excess of energy exists at the interface between two phases. Gibbs proves that this energy is a partial function of the interfacial area; thus: ?F/?s = ? where ?F/?s is the rate of change of free energy of the system with changing surface area, at constant temperature, pressure and composition, and ? is the interfacial tension, or interfacial free energy per unit area. If one of the phases is the pure liquid or solid, and the other the vapor of the substance, ? may properly be termed "surface tension," and is a characteristic of the solid or liquid. The attempt of a body to lower its free energy by decreasing its surface gives rise to a force in the surface which is numerically equal in terms of unit length to the free energy per unit area of the surface. Thus ? may be expressed either in erg-cm-² or in dyne-cm-1. Similarly, surface tension may be determined either by a thermo-dynamic measurement of the surface energy or by a mechanical measurement of the surface force. We have chosen the latter approach. Tammann and Boehme4 determined the surface tension of gold by measuring the amount of shrinkage or extension of thin weighted foil at various temperatures and interpolating to zero strain. The method is of questionable accuracy because of the tendency of foil to form minute tears when heated under tension. Their assumption of F = 2W?, where W is the width of the foil, is unsound, as the foil can decrease its surface area by transverse as well as by longitudinal shrinkage. Although their experimentation was meticulous, the paper does not include details of the sample configuration required for recalculating ? on a correct basis, even if such a calculation were possible. In the experimental procedure chosen here, a series of small weights of increasing magnitude are suspended from a series of line copper wires of uniform cross-section. This array is brought to a temperature at which creep is appreciable under extremely small stress. If the weight overbalances the contracting force of surface tension, the wire stretches; otherwise, it shrinks. The magnitude of the strain is determined by the amount of unbalance, so a plot of strain vs. load should cross the zero strain axis at w = F?. If balance is visualized as a thermodynamic equilibrium, the critical load is readily calculated. At constant temperature, an infinitesimal change in surface energy should be equal to the work done on or by the weight: ds = wdl [A] For a cylinder, s = 2pr2 + 2prl [2] If the volume remains constant, r = vV/pl [31 s = 2vpl+2V/l [4] ds = vpv/l - 2V/l²) dl [5] Substituting [5] into [I] gives for the equilibrium load, w = ?(z/rV- 2V/12) [6] and, again expressing V in terms of r and l, w = pr?(1 - 2r/l [7] Here the end-effect term, 2r/l, is neglected for thin wires in subsequent work. Eq 7 can be confirmed by means of a stress analysis. If the x-axis is chosen along the wire, then the stress is 2pr? - w pr² pr2 [8] A cylinder of diameter dis equivalent to a sphere of radius r, insofar as radial surface tension effects are concerned.³ Thus xv = 2?/d = ?/r = sz [9] For the case of zero strain in the x direction, the strain will also be zero in the y and z directions. Since the wire is under hydrostatic stress, Eq 8 and 9 are

Jan 1, 1950

By Paul R. Bluekamp

The Mather Mine property is composed of a 5.2 sq km (2 sq mile) area within the Cities of Ishpeming and Negaunee which are located in the Upper Peninsula of Michigan. Production in this mine started in 1943; it ran continuously until its closing in 1979, having produced slightly over 55 million tons of iron ore. This mine was a joint venture between several steel companies and The Cleveland- Cliffs Iron Company (CCI), with CCI being the fee holder and operator. For the first 17 years of operation, the ore was shipped in its natural state directly from the mine to the steel mill. By 1960, iron ore pellets were on the market and they proved so superior to the natural soft Mather ore that the latter became difficult to sell except for the coarser portions. It was decided to develop a pelletizing process for the Mather ore, and this was accomplished by 1965 when the first Mather pellets were produced. From this date, all but the coarse fraction of the Mather ore was shipped as pellets. The geological setting is that of a large east-west trending synclinorium which plunges to the west. The lowest member of this trough-like structure is, for the most part, a quartzite-like graywacke, the upper 20 m of which grades into a softer, fine grained slate. Lying conformably on this graywacke footwall member is an iron- formation member which is over 1,300 m in thickness. This iron-formation is cow posed of thin alternate bands of iron oxides and chert and is intruded by a number of diorite sills - some up to 122 rn (400 ft) in thickness. The north limb of this synclinorium dips at approximately 45' and bottoms out at about 1,000 m from surface in the central part of the mine. There are two sets of faults, many of which are intruded by diorite dikes, which trend east-west and southeast-northwest. Displacements are varied, reaching a maximum displacement of 243 m (800 ft). The ore is found lying directly on the slatey footwall and its position is largely controlled by the faults and dikes, with the bulk of the ore being on the upper side of these structures. The ore is composed of soft earthy hematite and martite with vertical thicknesses up to 122 m (400 ft) although the average thickness would be closer to 46 m (150 ft). The ore averaged 60% iron and 7% silica on a dried analysis. The Mather Mine is located on the north limb of the syncline and was worked from two shafts, the deepest of which. was 1,09 7 m. These two shafts are about 2 km apart and serviced a total of 8 working levels between them during the life of the mine. The level spacing was about 61 m (200 ft). The main haulageways were driven parallel to the ore/footwall contact in the hard cow petent graywacke wherever possible. On the lower levels, deeper into the footwall, naked development was common. This material graded into roof bolting ground towards the upper stratigraphic portion. As drifting progressed further into the upper stratigraphic portion of the footwall, progressively stronger steel sets had to be used. From the main haulageway, cross-cuts were turned into the orebody on 61 m (200 ft) centers and extended as far as needed to recover the ore available to that particular cross-cut. The main haulageways and cross-cuts were driven 3 m (10 ft) high, 3 m (10 ft) wide at the top and 4.6 m (15 ft) wide at the bottom. This configuration would accomodate a set composed of a 2.7 m (9 ft) cap on top of two 2.7 m (9 ft) legs angled out at 18'. A sill plate was used under each leg to prevent its sinking into the ore below. The type of steel used was dependent on the expected weight to be experienced. Sets were placed on 1.63 m (5 ft 4 in.) centers; however, in extremely heavy areas, it was sometimes necessary to install sets on 0.8 m (2 ft 8 in.) centers. On the top two levels (5th and 6th), the ore was considerably hard-er and was mined by sub-level stoping and long hole drilling. Very light steel sets were used in the main haulageways and cross-cuts and timber was used in the production drifts. Some con- crete production drifts were installed on 6th level, but proved to be uneconomical. However, as mining reached greater depths, the ore became softer and more massive, reaching its maximum vertical heights on 11th and 12th levels. On levels 7 through 10, yielding steel sets were used extensively in the slusher drifts. While they were satisfactory on 7th and 8th levels, their success diminished with depth and they were

Jan 1, 1981

By V. Pasupathi, R. E. Hummel, J. W. Koger

Specimens of P1 brass were plastically deformed at room temperature to various degrees of deformation and subsequently cooled in order to transform them to low-temperature martensite. Deformation shifts Ms. A, , and the temperature of minimum resistivity to lower temperatures, and also decreases the temperature coefficient of electrical resistivity. These properties change rapidly up to about 15 pct reduction but vary very little with higher deformation. The possible relationships between martensite formed by deformation and the M, temperature of low-temperature martensite are discussed. Evidence is given that deformation martensite delays the formation of low-temperature martensite. BETA' brass undergoes at least two different types of martensitic transformations. One of these transformations (B1- B2) was first observed by Kaminski and ~urdjumov' and occurs when 81 brass with a zinc content between 38 and 42 wt pct (quenched from the single-phase region) is cooled below room temperature. Jollev and Hull' determined the structure of 0" from X-ray and electron-diffraction data as ortho-rhombic. Kunze came to the conclusion that the super-lattice cell of 0" is one-sided face-centered triclinic (pseudomonoclinic). The second martensitic transformation (B1-A1) occurs when the specimens are deformed at or somewhat above room temperature. This type of martensite will be called deformation martensite. Horn-bogen, Segmuller, and Wassermann4 determined the structure of deformation martensite to be bct. (An intermediate phase, az, occurs before the final phase appears.) At deformations higher than 70 pct, a, transforms into a.4 A critical temperature Md exists above which no transformation occurs during deformation and is estimated to be around 400°C in P1 brass.5 This martensite has elastic properties.6 When the sample is stressed, martensitic plates appear; when the stress is released, the plates disappear. The present paper studies the effect of deformation martensite on the formation of low-temperature martensite. The experiments involved samples of 8, brass which were plastically deformed by various amounts and were subsequently cooled below the transformation temperature. EXPERIMENTAL PROCEDURE The 13 brass investigated was made from 99.999 pct pure copper and 99.9999 pct pure zinc and contained 38.8 wt pct Zn. The specimens, consisting of foils 0.1 mm in thickness, were heat-treated at 8'70°C for 15 min in an argon atmosphere and then quenched into ice water. They were then deformed by cold rolling and subsequently cooled at a rate of 1°C per min. The martensitic transformation that occurred during cooling was followed by electrical resistivity measurements. The resistance measurement technique and its accuracy have been described in a previous paper. Because the transformation 81 —-8" occurs below room temperature, the samples were placed in a cryo-stat which contained isopentane as a cooling medium. The isopentane was cooled by liquid nitrogen pumped under pressure through a 15-ft coil of copper tubing which was immersed in the isopentane. The nitrogen flow was regulated by a temperature controller using two thermistors in the cooling medium. The cryogenic liquid could be heated with an immersion heater. The useful temperature range with this device was from +25° to approximately -155~C. EXPERIMENTAL RESULTS Resistivity Measurements. The following abbreviations are used in this paper to label the characteristic temperatures during the martensitic transformation. M, is the starting point of the martensitic transformation and is defined as that temperature where the resistivity vs temperature curve on cooling first deviates from a straight line. Mf is the temperature at which the martensitic transformation is completed. On reheating, the transformation from martensite to the parent phase starts at a temperature A, and ceases at a temperature Af. Fig. 1 presents five different resistivity vs temperature curves corresponding to the transformation of brass from Dl to 8" after different degrees of reduction in thickness. The following observations can be made from these curves. 1) With increasing degree of deformation the Ms temperature is shifted to lower temperatures. This shift ranges up to 35°C compared to the undeformed state. This is also indicated in Fig. 2, where AM, (the shift of Ms, compared to the undeformed state) is plotted vs the degree of deformation. AM, increases rapidly until a reduction of about 15 pct is reached. With higher deformations, no additional increase in AM, was found. 2) With increasing degree of deformation the temperature of minimum resistivity (M) is also shifted to lower temperatures. The shift, attains a maximum of about 61°C compared to the undeformed state. In Fig. 3, AM is plotted as a function of deformation. It can be seen that, as in 1 above, AM increases rapidly and no further shift of M occurs for deformations greater than 15 pct. 3) The temperature coefficient of resistivity, is given by the slopes (dp/dT) of the linear portions of

Jan 1, 1969

By David P. Shoemaker, Clara B. Shoemaker

It has been shown for an important class of complex transition intermetallic compounds (a, P, R, 6, and p phases) characterized by "normal" coordination [CN12 (icosahedral), CN14, CN15, CN16/ that interatomic distances nay be calculated to a good approximation as the sum of characteristic atomic radii. Two radii, one for major ligands and one for minor ligmds, are specified for each atom, except in the case of CN12 where only a miaaor-ligand radius is specified. The same appears to be true of transition-metal phases of simpler struc-ture: Laves phases (CN12, CN16), and p-tungsten phases (CN12, CN14). In the case of known examples of the more complex phases, a simple rule is given which specifies these radii. However, only a fraction of the known examples of the simpler phases obey this rule closely. To include the latter phases the rule may be modified by considering the radii as linear functions of the weighted average of the Pauling CN12 radii of the two kinds of atonzs, with the radii weighted according to the over-all chemical composition of the alloy. With very few exceptions interatomic distances for both tlze complex and the simpler transition phases can b$ predicted with this modified rule to within 0.06A. ManY intermetallic compounds are known of composition A,By, in which A is a transition element to the left of the manganese column in the periodic table and B is a transition element in or to the right of it. Frequently the coordination numbers (CN) found in these compounds are CN12 (icosahedral), CN14, CN15, and CN16 (called "normal" coordinations by Frank and Kasperl). Well-known examples are the cubic and hexagonal Laves phases which have CN12 and CN16, and the 0-tungsten (CrsO) phases which have CN12 and CN14. In the more complicated (often ternary) phases, such as the a phase,2 the Beck phases p3 and R~, the 6 phase,5 and the p p atoms occur with CN12, CN14, CN15, and (except for a) CN16; in many cases several crystallographically independent atoms of one particular CN occur in the asymmetric unit. A large number of independent interatomic distances are found in these complicated phases, varying from 20 in the a phase to 94 in the 6 phase. These distances show a large spread; they vary, for example, from 2.358 to 3.278A in the 6 phase. In our analysis of these distances we found that in each of these compounds every atomic position can be characterized by either one or two radii. The CN12 positions are characterized by a single radius, The higher coordinated positions are characterized by two radii, namely: the CN14 positions by 4 in the direction of the twelve "5-coordinated" ligands3 (called 'minor" by Frank and Kasperl) and by r:, in the direction of the two "6-coordi-nated" ligands (called "major" by Frank and Kasper); the CN15 positions by r15 for the twelve minor and r:, for the three major ligands; the CN16 positions by rlE for the twelve minor and r:, for the four major ligands. We have expressed the experimentally determined interatomic distances in observational equations as the sums of the appropriate pairs of these characteristic radii and the value of these radii have been determined by the method of least Squares. Despite their wide range, the interatomic distances could then be predicted by the sums of these atomic radii with an average deviation in any one compound of 0.06A or less. The results are summarized in Table I. Inspection of the radii thus obtained shows that in the structures in Table I the radii (in A) are given to a first approximation by the simple relationship: Where CN is the coordination number (12, 14, 15, or 16), and A = 1 for major ligands and = 0 for minor ligands. The interaLomic distances can be predicted within about 0.1A by sums of these atomic radii. Another phase belonging in this group with CN12, 14, 15, and 16 is the y phase & B7, in which A is molybdenum or tungsten and B is iron or cobalt. Recently the M%C phase has been refinedE and the observed distances also agree well with those calculated with Eq. [I]. (In the original determination of the structure of W6FeV7 the F$(II)-W(II1) distance was erroneously given as 2.84A, but we have recalculated it fro? the published parameters and found it to b? 2.57i4, in good agreement with the value of 2.6A predicted with Eq. [I.].) Many binary transition alloys are known to crystallize with the simpler structures having "nor-

Jan 1, 1964

By M. O. Abo-el Fotoh, J. B. Mitchell, J. E. Dorn

The effect of strain rate and temperature on the tensile flow stress of a polycrystalline bcc alloy of magnesium containing 14 wt pct Li and 1.5 wt pct Al was investigated for strain rates of 3.13 x lom5 to 3.13 x 10-3 per sec over the range from 20° to 300°K. From about 180° to 300°K the alloy exhibited an ather-ma1 deformation behavior where the flow stress was independent of strain rate and increased only slightly with decreasing temperature. At lower temperatures the flow stress was strongly strain-rate- and temperature-dependent, characteristic of deformations controlled by thermally activated mechmzisms. The activation volume for thermally activated plastic defornzation was between 5 and 30 cu Burgers vectors, independent of plastic strain. This low-temperature thermally activated deformation behavior was found lo be in satisfactory agreement with the theoretical dictates of the Dorn-Rajnak1 formulation of the Peierls mechanism where deformation is controlled by the rate of nucleation of pairs of dislocation kinks over the Peierls energy barriers. SEVERAL studies of the low-temperature thermally activated deformation of bcc metals and alloys (molybdenum,1 tantalum,1 Fe-2 pct Mn,2 Fe-11 pct MO,3 and AgMg4) have revealed that the strain rate is controlled by the activation of dislocations over the Peierls-Nabarro energy hills. Although there is some uncertainty as to the nature and effect of solute atom-dislocation interactions during low-temperature deformation of bcc metals, it has been concluded by Dorn and Rajnak,1 Conrad,1 and Christian and Masters6 among others that overcoming the Peierls-Nabarro stress which arises from the variations in bond energies of atoms in the dislocation core as it is displaced is the probable mechanism controlling low-temperature deformation. The purpose of this research was to investigate the low-temperature plastic deformation of the bcc alloy Mg-14 wt pct Li-1.5 wt pct A1 to determine if the behavior of this alkali metal alloy might be analogous to that for other bcc metals. This alloy was selected because of its availability and its current industrial importance as a lightweight material for aircraft and aerospace applications. I) EXPERIMENTAL PROCEDURE Polycrystalline tensile specimens having cylindrical gage sections 2 in. long by 0.2 in. in diam were machined from as-received alloy sheet stock of Mg-14 wt pct Li-1.5 wt pct Al. Specimens were annealed in an argon atmosphere at 423°K for 4 hr and maintained in a kerosene bath together with the sheet stock to prevent corrosion. The resulting specimen microstructure consisted of a coarse uniform dispersion of incoherent precipitate MgLi2Al particles7 in a bcc 0 phase matrix having an average grain size of 150 p. Prior to testing the specimens were chemically polished in dilute hydrochloric acid. Comparison of tensile properties and microstructures of specimens cut from center and edge sections of the sheet stock revealed no effects of inhomogeneities in the sheet material. Tensile tests were performed on an Instron machine at crosshead speeds corresponding to tensile strain rates of 1.56 x 10-5 and 1.56 x 10"3 per sec. Stresses were determined to ±2 x 106 dynes per sq cm and strains to within ±0.0001. Average values of shear stress t and shear strain y reported were taken as one half the tensile stress and three halves the tensile strain, respectively. Flow stresses were taken at 0.05 pct strain offset. Test temperatures down to 77°K were obtained by immersing the specimens in constant-temperature baths. Lower-temperature tests were performed in a liquid helium cryostat to within ±2°K of the reported values. Prior to testing at the various temperatures and strain rates all specimens were prestrained at 2 35°K at a shear strain rate of 3.13 x 10-5 per sec to a stress level of 0.606 x 10' dynes per sq cm to obtain a uniform initial state. Additional tests were made to determine the effect of changes in strain rate and strain on the flow stress by rapidly changing the crosshead motion during testing. Shear moduli of elasticity, needed for analyses of the data, were obtained at several temperatures by a common technique of determining the resonant frequencies of vibrations of rectangular test specimens. 11) EXPERIMENTAL RESULTS Fig. 1 shows the experimentally determined flow stress vs temperature for two strain rates. Two distinct regions of behavior are evident. Below about 180°K the strong increase in flow stress with increased strain rate and decreasing temperature indicates that deformation is controlled by a thermally activated dislocation mechanism. At higher temperatures an athermal region is evident where the flow stress is independent of strain rate and only slightly dependent on temperature. The applied stress t to cause plastic flow was separated into two components:

Jan 1, 1969

By Martin F. Littmann

Inclusions of MnS and Fe3C have been introduced into single crystals of 3 pct Si-Fe to study their effects on recrystallization behavior and textures after cold rolling and annealing. The presence of MnS in (110) [001] and (111)[112] crystals inhibited primary grain growth and promoted secondary recrystallization but did not alter the texture significantly after annealing at 1200°C. The presence of Fe3C in (llO)[OOl] and (100)[001] crystals caused a refinement of the primary re crystallized grain size but did not promote secondary recrystallization. THE texture behavior of single crystals of 3 pct Si-Fe during deformation and recrystallization has been studied by numerous investigators. The early work of Dunn&apos; followed by Decker and Harker2 involved relatively small cold reductions. More detailed studies of Dunn3&apos;4 and of Dunn and Koh5&apos;6 involved a reduction of 70 pct and recrystallization at 980°C for several crystals. Walter and Hibbard7 studied a greater variety of initial orientations and sought to relate the textures to those of polycrystalline material. Attention was focused on the nucleation process during early stages of annealing and on surface energy effects in studies by Walter and Dunn8 and by HU.9&apos;10 One of the most extensive investigations has been reported by T. Taoka, E. Furubayashi, and S. Takeuchi.11 Most of this work has been conducted using relatively pure crystals with minimal amounts of precipi-tate-forming elements such as carbon, oxygen, sulfur, and nitrogen. Recently, however, S. Taguchi and A. Sakakura have observed that AIN precipitates can alter the recrystallization textures of rolled (100)[001] crystals.12 The present studies were initiated to determine effects of MnS and Fe3C precipitates on recrystalli-zation and grain growth behavior of rolled single-crystals of 3 pct Si-Fe. Both of these types of inclusions play significant roles in the recrystallization behavior leading to the formation of the (110)[001] or cube-on-edge texture in commercial grain-oriented silicon iron. It is well known that (110)[001] primary grains are formed by recrystallization of (110)[001] or (11 l)[ 112] crystals after cold reduction of about 60 pct or more. Crystals of these orientations, therefore, were selected for study of the effect of MnS in-clusions on grain growth. On the other hand, a major component of the texture of cold-rolled, polycrystal-line 3 pct Si-Fe is the (100)[011] orientation. The function of Fe3,C inclusions is of interest for this orientation. EXPERIMENTAL PROCEDURE The single crystals used are listed in Table I and were obtained from commercial Si-Fe alloy processed to produce (110)[001] and (100)[001] texture by secondary growth. The cube-on-edge material was 0.59 mm thick. Suitably large (110)[001] crystals 25 mm wide were selected and their orientations were determined using an optical goniometer. Etch pits for texture determination were formed by a ferric sulfate solution. The other crystals used in the study with (100)[001], (100)[011], and (111)[112] orientations were obtained from sheet which contained large grains developed from secondary recrystallization by a surface-energy driving force.13 Most crystals had a (100) plane very nearly parallel to the sheet surface and the rolling direction could be selected readily. The same sheet also contained a few crystals with (111) planes parallel to the sheet surface, these also being a result of growth by surface energy. The crystals selected from the sheet were about 25 mm wide and 0.25 to 0.28 mm thick. As shown in Table 11, the crystals already contained about 0.070 to 0.10 pct Mn. Inclusions of MnS were incorporated into crystal 36 in the following manner. The crystals were first sulfurized by holding them Table I. Initial Orientations of Crystals Crystal No. Initial Orientation Thickness, mm Special Treatment 34 (I10) [00l]* 0.59 None 36s (110) [001] 0.59 Sulfide precipitates added 30,40 (111)[Ti21 0.28 None 43s (III) [Ti21 0.28 Sulfide precipitates added 37 (100) [Oll] 0.30 None 37C (100) [01I] 0.27 Carbon added 41 (100) (01I] 0.25 None 41C (100) [OI11 025 Carbide precipitates added 42 (100) [OOl] 0.25 None 42C (100) [001] 0.25 Carbide precipitates added *Tilted 4 deg to r~ght about R.D. Table II. Compositions of Crystals Special Treatments Base Analysis ~ ______________________£________________Crys- Crystals Pct Si Pct C Pct Mn Pct S Pct N Pct Al tal Pct C Pct S 34.36 2.93 • 0.099 <0.005 - 0.0014 36S 0.011 30.37 to 42 2.78 0.0057 0.070 0.001 0.0008 0.0011 43S 0.022 37C 0.029 -41C 0.028 -42C 0.026 *Estimate 0.004 pct. Oxygen estimated <0.003 pct on all samples

Jan 1, 1970

By Amitava Gangulee, Michael B. Bever

The order-disorder transition in Ag-Mg alloys in the range 17 to 26 at. pct Mg was investigated and some thermodynamic, electrical and mechanical properties of ordered Ag-Mg alloys were measured. A modification of the phase diagram is proposed on the basis of measured transition temperatures. The stability of the ordered structure is analyzed in terms of the quasichemical theory. Kinetic aspects of the order-disorder transition were also investigated. The energies of ordering at 273°K of Ag-Mg alloys were measured by liquid metal solution calorimetry. The electrical resistivities and the tensile Properties of the ordered alloys were measured. The exothermic energy of ordering increased with magnesium concentration up to the composition Ag3Mg; it is discussed in terms of the quasichemical theory. Pronounced order hardening was observed and can be explained by several strengthening mechanisms, which are particularly effective because of a change in crystal structure associated with the ordering transition. IN silver-rich solid solutions containing approximately 17 to 26 at. pct Mg, long- range ordering may occur during appropriate therma1 treatments. Clare-brough and Nicholas1 first detected such long-range order in the solid solution of composition Ag3Mg and suggested that the unit cell of the superlattice contained 256 atoms. x-ray2 and electron diffraction3 investigations confirmed the occurrence of long-range order, but indicated a shifted Ll2-type structure. On the basis of recent X-ray measurements the crystal structure of ordered Ag3Mg has been confirmed as type D023.4 The present paper is concerned with the order-disorder transition in silver-rich Ag-Mg solid solutions and some thermodynamic, electrical, and mechanical properties of the ordered alloys. The effects of short-range order on the properties of silver-rich solid solutions containing up to 26 at. pct Mg are discussed in a concurrent paper.5 1) EXPERIMENTAL PROCEDURES 1.1) Preparation of Specimens. Specimens of Ag-Mg alloys containing from 18 to 26 at. pct Mg were prepared in the form of 1.0-mm-diam wires as described in Section 1.1 of Ref. 5. Disordered specimens were prepared by annealing at 773°K for 1 hr and quenching into iced brine. Ordered specimens were prepared by annealing at 773°K for 1 hr and slowly cooling to room temperature over a period of 15 days. The specimens were stored at 78°K. 1.2) Resistivity Measurements. Electrical resistivi- ties were measured by a potentiometric method.&apos; Equilibrium values were obtained at several temperatures. The kinetics of ordering were investigated by following the time-dependent changes of the resistivity of initially disordered specimens during annealing. The specimens were enclosed in copper capsules and immersed in salt pots; the heating up was accelerated by injecting preheated helium into the capsules. 1.3) Calorimetry. The energies of ordering of the alloys were measured in a tin solution calorimeter as the difference between the heat effects of additions of completely ordered and disordered specimens of the same composition. The procedure and the method of calculation have been described.6 Magnesium was used for thermal compensation in most calorimetric runs in order to improve the accuracy.677 1.4) Mechanical Tests. Tensile tests were carried out with wire specimens at room temperature as described in Section 1.5 of Ref. 5. Microhardness measurements were also made.? 2) RESULTS AND DISCUSSION The characteristics of the order-disorder transition of silver-rich Ag-Mg solid solutions will be discussed first. Some thermodynamic, electrical, and mechanical properties of the ordered alloys will then be considered and compared with the corresponding properties of dis-disordered alloys. 2.1) The Order-Disorder Transition. 2.1.1) Thermodynamic and Structural Aspects. The transition temperatures Tt were determined from discontinuities in the slope of the equilibrium resistivity vs temperature curves. Normalized curves for three compositions are shown in Fig. 1. The transition temperature increases with the magnesium concentration and reaches a maximum at the stoichiometric composition Ag3Mg. The transition temperature of the alloy Ag3Mg was measured as 665° ± 2°K and compares with published values of 660°,1 663°,8 and 66°k.3 In the slopes of the resistivity vs temperature curves of alloys containing 22.2 and 22.5 at. pct Mg, discontinuities were observed at two temperatures. Such upper and lower discontinuities indicate a two-phase field. A modified form of the published phase diagram9 is shown in Fig. 2. A two-phase field was found only on the low-magnesium side of the composition Ag3Mg. On the high-magnesium side, the existence of a two-phase field could not be established because of insufficient resolution, but such a field must be present. This part of the phase diagram can be made complete by a eutec-toid-type reaction a = (a&apos; + ß). The existence of a boundary between the two-phase fields (a + ß) and (a&apos; + ß) is also in accord with published lattice parameters.3 The crystal structure of ordered Ag3Mg (type Do23)

Jan 1, 1969

By E. Miller, K. Komarek, W. Johnson

The activity of aluminum in solid Cr-A1 alloys has been measured by an isopiestic technique between Cr-A1890&apos; and 1126" and 13 and 80 at. pct Al. The integral free energy of mixing has a minimum value of —5600 cal per g-atom at 59 at. pct Al. The maximum solid solubility of aluminum in chromium was determined to be 43 at. pct Al, and the composition limits of the compounds CrA14, Cr4A19, and Cr5Al, at 1000"~ were found to be 79 to 80, 66 to 70, and 59 to 63 at. pct Al, respectively. The thermodynamic properties of the Cr-A1 system have been investigated as part of a thermodynamic study of aluminum-transition metal systems.172 Little information is available on the equilibrium properties of the Cr-A1 system. The heats of formation of solid Cr-A1 alloys have been determined by Kubaschewski and Haymer at 600" and low-temperature specific heat data have also been obtained.~ More extensive work has been performed on the phase diagram, and a compilation has been provided by Hansen and Anderko,~ their phase diagram at elevated temperatures being essentially based on the work of Bradley and LU.~ The high-temperature portion of the phase diagram shows an intermediate phase CrA14 decomposing peritectically at 1018°C and existing at 82 at. pct A1 at 1000°C. They also identified the compounds with solubility limits of 72 to 75 at. pct A1 at 1000°C, and Cr5A1,, existing at 61 at. pct A1 at 1000°C. The maximum solid solubility of aluminum in chromium at 1000°C was found to be 46 at. pct Al. These elevated-temperature data were obtained by examination of quenched samples and were considered as less precise than the lower-temperature data. Koester, Wachtel, and Grube7 have revised the phase diagram as a result of their magnetic susceptibility and X-ray study. The results of this work differ appreciably from those of Bradley and Lu at temperatures above 800°C. The CrA1, compound is given as existing between 79 and 81 at. pct A1 at 1000°C, and they do not indicate the presence of a CrA13 phase reported by Bradley and Lu. They also report the compound Cr4Alg as having solubility limits of 66 to 70 at. pct A1 at 1000°C, while Bradley and Lu show this compound stable only up to 870°C. Koester et al. state that the high-temperature modification of the compound Cr5A18 is stable down to 1125"C, and not 980°C as stated by Bradley and Lu, and that the low-temperature modification of Cr5Al, has a range of homogeneity of 58 to 63 at. pct A1 at 1000°C. They also report that the maximum solid solubility of aluminum in chromium is 43 at. pct A1 at 1000°C. APPARATUS AND EXPERIMENTAL PROCEDURE An isopiestic method was employed which has been successfully applied to the determination of aluminum activities in solid ~e-All and Ni-Al alloys. Alloy specimens were held at different positions in a temperature gradient and were equilibrated with aluminum vapor from an aluminum reservoir kept at the temperature minimum of an impressed thermal gradient in a closed alumina system. Diffusion of aluminum into the specimens occurred until equilibrium was reached, at which the partial pressure of aluminum in each of the specimens was given by the vapor pressure of the pure aluminum reservoir. The activity of aluminum referred to liquid aluminum as the standard state in a given equilibrated sample at temperature T could therefore be expressed by: vapor pressure of pure aluminum at _ the temperature of the reservoir Vapor pressure of pure liquid aluminum, at specimen temperature T Since both the temperature of the aluminum reservoir and the specimen temperatures were determined experimentally, and the vapor pressure of pure aluminum is known as a function of temperature,&apos; the activity of aluminum in a given aluminum alloy of known composition could be calculated. Initial runs were made with samples consisting of pure chromium chips placed in alumina crucibles. These runs exhibited large inconsistencies, indicating that equilibrium was not attained. High aluminum content Cr-A1 alloy powders were therefore substituted for the pure chromium specimens. The starting composition of the alloys was adjusted through experimentation until the concentration change necessary to attain equilibrium was small. In this manner, consistent results were obtained in reasonable times. SPECIMEN PREPARATION Alloy specimens were prepared from chromium of 99.997 pct minimum metallic purity: with 0.028 to 0.038 pct H, 0.0002 pct N, and 0.27 to 0.46 pct 0 (Aviquipo, Inc.). The aluminum had a purity of 99.99+ pct and the following impurities: 0.003 pct Cu; 0.002 pct S; 0.002 pct Fe; 0.001 pct Pb; 0.001 pct Ga (Aluminum Corp. of America). Alloy powders were prepared from weighed mixtures of chromium and aluminum by double-arc melt-

Jan 1, 1969

By F. H. Froes, D. H. Warrington

Electron microscope evidence which indicates that TaC may precipitate at random sites in the matrix is presented. Initially the particles are almost spherical and coherent with the matrix. However, as they grow in conditions in which there are insufficient vacancies to relieve lattice strain, the particles rapidly lose coherency in two directions and continue to grow as plates with approximately the full lattice mismatch strain present perpendicular to the plane of the plate. The necessary relief of strain comes from dislocations loops which do not become visible until the later stages of aging. The rapid decrease of apparent strain to low values of appoximately 1 pct at small particle sizes arises not from a complete incoherency but from applying a model wrong for the particle shape and strain distribution. PREVIOUS work has shown that MC-type carbides may precipitate intragranularly in austenitic stainless steel on dislocations,1&apos;2 in association with stacking faults,3&apos;4 and randomly through the matrix,5-7 In investigations of the matrix precipitate by thin-foil electron microscopy, considerable lattice strain has been found to occur around the precipitating phase.7&apos;8 Attempts have been made to evaluate the amount of lattice strain by using the methods developed by Ashby and brown.9,10 Values of the linear strain, much less than the 17 pct theoretical mismatch (for TaC), have been reported; it has been suggested that this is due to either a loss of coherency1&apos; or vacancy absorption which occurs during either the initial nucleation or growth of the precipitate." This report is an extension of earlier work7 that dealt with the precipitation of TaC from an 18Cr/12Ni/ 2Ta/O.lC alloy after it had been quenched from 1300°C and aged between 600" and 840°C. In particular, the shape of the precipitate particles and the amount of strain in the matrix, due to the precipitate, have been studied. The work described here is part of a wider investigation of factors that affect carbide precipitation in austenitic stainless steel," details of which are to appear elsewhere. RESULTS The present investigation can be conveniently split into two aspects of the strain-fields surrounding the matrix particles: 1) information derived from the strain-field which indicates the shape and habit plane of the precipitate particles and 2) the magnitude and sign of the strain-field. The Shape and Habit Plane of the TaC Precipitate. In the early stages of aging twin lobes (normally black F. H. FROES, formerly at the University of Sheffield, Sheffield, England, is Staff Scientist, Colt Industries, Crucible Materials Research Center, Pittsburgh, Pa. D. H. WARRINGTON is Lecturer, Department of Metallurgy, University of Sheffield. Manuscript submitted November 1, 1968. IMD on white background, i.e., for the deviation parameter, S > 0) that indicate the strained region of the matrix define the position of the particles by bright field transmission electron microscopy. The actual particles were not detected until they were approximately 120Å diam; below this size they were too small to be imaged in the electron microscope. This meant that particle growth that had occurred before this stage had to be inferred from the matrix strain-field contrast. In all cases when diffraction effects were observed from the precipitate particles, a cube-cube orientation relationship (i.e., (llO)ppt Il<llO>matrix and {1ll }ppt {III} matrix) existed between the precipitate and the matrix. From the matrix precipitate particles lying along edge-on {111} planes (e.g., at A, Fig. I), the precipitates are seen to be plate-like with their diameter being roughly 18 times their thickness after 5000 hr at 650°C. However, the exact shape of the particles cannot be determined because of the masking effect of the strain-field contrast. If a dark-field micrograph, using a precipitate reflection, is studied, Fig. 2, a number of the projected images of the TaC particles [on the (110) foil surface] apear to have straight edges parallel to projected f111) planes. Thus, it appears that in the later stages of aging the TaC particles are plate-like with some tendency for the edges of the plate to be bounded by the matrix close-packed {ill} planes (though the general shape of the particles in the plane of the plate is circular and thus the "diameter" of the particles has a real physical significance). It should be noted that the bands of fine discrete particles observed in Figs. 1 and 2 are not the matrix precipitate discussed in this paper but are precipitates associated with extrinsic stacking faults3j4 occurring on (111) matrix planes. **£** ****** \ *x 23 Fig. 1—18/12/2~a/0.1~ alloy. Solution treated at 1300°C for 1 hr, water quenched, and aged 5000 hr at 650°C. The (112) directions shown are the traces of the e&e-on (111) planes. Foil normal [110]; operating reflection (331); bright field micrograph.

Jan 1, 1970

By M. Cohen, D. J. Blickwede

Quantitative observations concerning the carbide phases in high speed steel are of importance for two general reasons: (1) the carbides, being inevitable constituents of the final structure, exert a direct influence on the properties of the steel; and (2) a substantial proportion of the total alloy content is tied-up in the carbides, and hence the extent of their solution on austenitizing governs the composition of the steel matrix. The latter relationship has a vital bearing on the response of the steel to tempering as well as on its performance in subsequent service. Accordingly, in the course of a long-term study of the behavior of high speed steels, the authors were confronted with the problem of securing quantitative data on the carbide phases. The obvious method for acquiring such information is to isolate the car-bides from the steel and subject them to chemical, X ray diffraction and other measurements. There are well-known extraction techniques which involve the chemical or electrolytic solution of the less noble matrix (ferrite, marten-site or austenite), thus leaving a residue of the carbide phases. However, the results obtained must be scrutinized carefully1,2 since the carbides may be affected by the chemical or electrolytic action. It is the purpose of the present paper to describe the experiments leading to an electrolytic-extraction technique for quantitatively isolating the carbides from both I and hardened high speed steel. Particular attention is paid to the amount, as well as the composition, of the carbides so that the matrix analysis becomes ascertainable by subtraction from the overall steel composition. Illustrative results are given for the M-2 grade of tungsten-molybdenum steel. Review of the Literature The chemical method of dissolving the matrix selectively with respect to the carbides makes use of dilute non- oxidizing reagents such as hydrochloric or sulphuric acid. Although this simple procedure has led to the determination of the cementite composition,3,4 it achieved only limited success because of the interaction between the acid and the carbide residue. Some of the carbides may not only be destroyed in this way, but the hydrogen released is likely to remove part of their carbon as hydrocarbon gases. The electrolytic technique of isolating carbides has the advantage of rapidly dissolving the specimen (anode) in the presence of less reactive solutions than are practicable with the chemical method. This reduces the possibility of chemical attack on the carbides, and furthermore, the hydrogen evolved during the electrolysis is released at the cathode which is not in close proximity to the carbides. The common electrolytes adopted for this purpose are hydrochloric and sulphuric acids.5-l1 Aqueous solutions of ferrous salts have also been used.12,13 A considerable advance in experimental technique was introduced by Treje and Benedicks14 who developed a double-compartment cell for electrolytic extraction, the anode and cathode chambers being separated by a porous diaphragm. A solution of 15 pct sodium citrate, 2 pct potassium bromide and 1 pct potassium iodide was selected for the anolyte, while the catholyte consisted of a 10 pct solution of copper sulphate, with copper serving as the cathode. This type of cell has a number of desirable characteristics: 1. The anolyte has a pH value close to 7, at least at the beginning of the run. 2. The iron that dissolves from the anode-specimen forms a water-soluble complex ion with the citrate, thereby preventing the precipitation of iron hydroxide (which would contaminate the carbide residue) despite the neutrality of the solution. 3. Copper deposition instead of hydrogen evolution occurs at the cathode, and this avoids an increasing concentration of hy-droxyl ions which (in an otherwise neutral solution) might cause the precipitation of insoluble hydroxides. 4. Contamination of the anode chamber by copper sulphate is inhibited by the porous diaphragm. Houdremont and coworkers15 applied the above method (with the further refinement of excluding oxygen during the electrolysis, washing and drying) to the extraction of carbides from a series of plain carbon steels after various heat treatments. They had quantitative success only with specimens in the annealed condition, and concluded that the size and shape of the carbide particles play an important role in the isolation process, with large spheroids exhibiting the least tendency to decompose during the electrolysis. Up to the present time, the citrate double-cell has not been used to any extent for isolating the carbides of high alloy steels, apparently on the grounds that the complex carbides are more resistant than cementite to attack in the simpler acid electrolytes. In particular, Bain and Grossmann7 and Gulyaev10.10a have employed the hydrochloric acid cell for their investigations of the carbides in high speed steel.* It will be demonstrated here that this type of cell is capable of yielding quantitative results in the case of high speed steel, and actually has certain advantages over the more complicated double cell. However, in order to provide a rigorous test of the quanti-tativeness of electrolytic procedures for the problem at hand, both methods were studied in considerable detail.

Jan 1, 1950

By Pradeep K. Rohatgi, Clyde M. Adams

Structures produced upon solidification of the eu-tectic composition (33 wt pct Cu) aluminum copper alloy have been examined as a function of freezing rate dfs /d? , the rate of change of fraction solid (fs) with time (8). Slow (dfs/d? = 0.0016 sec-1), intermediate (dfs/d? = 0.02 sec-1) and rapid (dfs/d? = 0.4 to 7.30 sec-1) freezing rates were used. The lamellar Al-Cual2 eutectic is arranged in the form of rod-shaped colonies at rapid freezing rates. The colonies are aligned parallel to the direction of heat flow, whereas the lamellae within the colonies are aligned at various angles, as high as 90 deg, to the direction of heat flow. The colony spacing (C) is proportional to the square root of inverse freezihg rate. The relationship is C = 15.5(dfs/d?)-1/2 where C is in µ and 8 is in sec. The ratio of colony spacing to lamellar spacing is greater than 20.0 and increases with a decrease in the freezing rate. A duplex dendritic structure is produced at intermediate freezing rates. A fine lamellar eutectic is arranged within the dendrites (exhibiting side branches at an angle close to 60 deg from the main stem) and a coarse irregular eutectic appears in the interdendritic regions. The duplex eutectic structure is also produced at slow freezing rates. However, at slow freezing rates there is a Platelat of CuAl2, along the center of the main stem of each dendrite and the other lamellae are arranged perpendicular to the central platelet. THE eutectic between CuA12 and a! aluminum has been reported to freeze in a lamellar form by several workers.&apos;-3 chadwick4 has measured the interlamel-lar spacing as a function of growth rate. Kraft and Albright2 have reported on irregularities in the lamellar structures, and have proposed growth models which account for the formation of faults during solidification. In certain instances the lamellar eutectic has been found to exist in colonies. The colony formation315 has been attributed to the breakdown of a planar liquid-solid interface due to rejection of impurities. The aim of the present work is to study the structures produced from the eutectic aluminum-copper alloy under relatively fast solidification rates, such as encountered in casting and welding operations. The solid-liquid interface presumably remains planar under conditions of slow unidirectional freezing which produce lamellae aligned parallel to the direction of heat flow. The local growth velocities are the same over the entire interface and are equal to the rate of growth of the all-solid region. The spacing between the eutectic lamellae is inversely proportional to the square root of the growth rate of the all-solid region. Under the freezing conditions used in the present study, the solid-liquid interface is cellular or dendritic and the local growth velocities are different in the different regions of the interface. The relationship between the growth rate of the all solid region and the local growth velocities varies with the location and the shape of the interface. The growth rate of the all-solid region is, therefore, an inadequate parameter to describe the eutectic micro-structures which depend upon the local growth velocities. For this reason the structures have been examined as a function of freezing rate, dfs/d?, where fs is the fraction solidified at time 0. The freezing rate was varied by a factor of 4000. The relationship between the freezing rate, dfs/d?, and the growth velocit of the all solid region depends upon the specimen geometry and the shape of the interface. EXPERIMENTAL PROCEDURES The A1-33 pct Cu alloy used throughout this study was made in an induction furnace, using electrolytic copper and aluminum of commercial purity (99.7 pct), the primary impurities being silicon (0.12 pct), iron (0.14 pct), and zinc (0.02 pct). Three ranges of freezing rates were investigated: 1) A spectrum of rapid freezing rates (ranging from 0.40 to 7.30 sec-1) was obtained in arc deposits made on 2-in. thick cast plates of the eutectic alloy. The arc was operated at constant power and was made to travel at constant velocity on the surface of the plate that was in contact with the chill surface during solidification. The pool of liquid metal formed under the moving tungsten arc solidified rapidly by heat extraction through the unmelted plate. Conditions of unidirectional heat flow were achieved near the fusion zone interface, especially in the center of the arc deposits. The great advantage of the arc technique is that rapid cooling and freezing rates can be varied in a qualitative way. The correlation between the arc parameters and the solidification rate is given by the following relationship:6-8

Jan 1, 1970

By D. L. Harris, A. M. Gaudin

Observations were made of the distribution of mercaptan containing S35 between aqueous solution and mineral and between aqueous solution and the gaseous phase. Although equilibrium may not have been attained, adsorption of the reagent was shown to occur reasily from air or aqueous solution on sphalerite, zincite, and willem-ite and to correspond to flotation. Adsorption on quartz did not similarly occur. THE following results, presented here in condensed form,&apos; were obtained in a preliminary study of the adsorption of n-hexane thiol, hexyl mercaptan, on sphalerite, zincite, willemite, and quartz, from aqueous solution and from a gas. Interest in this subject was aroused by a Belgian report&apos; of effective use of hexyl mercaptan for flotation collection of oxidized zinc minerals. The relatively low boiling point, 149°C, of the mercaptan3 suggested the desirability of extending the usual measurements of partition of collector between aqueous solution and gas and between gas and mineral. It is believed that this paper presents the first measurements of this type on a flotation system. Attempts were made to carry out the measurements at equilibrium, but as the work progressed it became increasingly doubtful that this desirable condition had been achieved. To control composition and extent of the gas phase, the apparatus was a wholly-enclosed thermally-controlled glass system. Because of these constraints and the desirability of dealing with pure minerals, a scale of operations was chosen in which a few grams of deslimed mineral were used in each test. It was also necessary to choose a particularly sensitive method for mercaptan analysis, and in fact a method that would permit the experimenters to follow the approach to equilibrium. For these reasons mercaptan marked by radiosulphur 35 was used. An analysis was made for the radiosulphur by a modification of the method of Gaudin and Carr. Coarsely-crystallized sphalerite was handpicked, stage-crushed in the dry state, wet-screened on a 200-mesh sieve, and deslimed in water at about 5 microns. Further treatment consisted of a wash in dilute aqueous hydrogen peroxide, drying, removal of the dark-colored fraction in a Frantz magnetic separator, washing in very dilute hydrochloric acid, repeated washing in distilled and conductivity water, and drying. The last washings showed a conductivity equivalent to a few ppm NaC1, that is, much more than would be provided, theoretically, by a saturated ZnS solution. The material was stored dry in sealed bottles. Analyses were as follows: Zn, 62.3 pct; Fe, 0.43 pct; Cd, 0.44 pct; S, 31.2 pct; Mn, 0.001 pct. The specific surface (BET method) was 2000 cm2/g. Zincite from Franklin furnace of the New Jersey Zinc Co. was hand-picked, dry-crushed, wet-screened at 100 mesh, and deslimed at about 10 microns. After drying, the associated zinc, manganese, calcium, and silicate minerals were removed in a Frantz magnetic separator. The purified zincite was washed in distilled water and conductivity water to a conductance of less than 2 ppm equivalent NaC1, dried, and stored. Analyses were as follows: Zn, 75.1 pct; Fe, 0.9 pct; Mn, 2.78 pct. The specific surface (BET method) was 1740 cm 2/g. Willemite, also from Franklin furnace, was purified similarly. Analyses were as follows: Zn, 52.5 pct; Fe, 0.12 pct; SiO², 27.3 pct; loss on ignition, 0.13 pct. The specific surface was 1760 cm 2/g. Conductivity water (double-distilled) and demin-eralized-distilled water were used in most of the tests. The specific resistance was not less than 600, 000 ohms, and usually above 1,000,000. Radiosulphur-marked hexyl mercaptan (1-hexane thiol) was synthesized by Tracerlab, Inc., Boston. Two lots were secured several months apart. The last lot, consisting of about 0.5 g of the mercaptan, had a total activity of about 10 millicuries. Tracerlab Co. guaranteed only the activity; hence a quasi -vapor pressure determination (based upon an S analysis) of the mercaptan was made. The calculated value, 4.2 mm of mercury at 25.5&apos; C, has been compared with that of a sample of Highest Purity 1-hexane thiol from Fisher Scientific Co. The latter had a vapor pressure of 4.5 mm of mercury at 2.5 C. Analytical Procedures The sample containing radiosulphur-marked mercaptan was oxidized to convert the mercaptan sulphur to sulphate, carrier barium sulphate being added to provide a suitable quantity of total barium sulphate in a filter cake. The precipitate was filtered and dried, and counting was carried out either in a streaming-gas (Q-gas) counter for high sensitivity or with an end-window G-M counter for convenience. The oxidized and precipitated mercaptan gave a radioactive count of 65 counts per minute per microgram in the end-window Geiger-Mueller counter and 1100 counts per minute per microgram in a Q-gas counter. For standardization of the mercaptan solution, 15 replicate analyses were made. The average deviation per measurement was about 1600 cpm in 65,000 cpm, the probable error in the mean being 275 cpm. It

Jan 1, 1955

By W. C. Ellis, E. S. Greiner, M. E. Fine

Discontinuous changes of Young&apos;s modulus, internal friction, coefficient of expansion, electrical resistivity, and thermoelectric power are evidence for a transition in chromium near 37OC. Although the X-ray diffraction pattern gives no clue, a difference between the thermal expansivity and the temperature dependence of the lattice parameter suggests a crystal-lographic change. Young&apos;s modulus data disclosed another transition near THE thermal dependence of a number of proper- ties of chromium indicates a transition occurring over a temperature range near room temperature. Bridgman&apos; noted this first from a minimum in the electrical resistivity near 12°C. In a sample of greater purity, Sochtig&apos; observed the minimum at 41 °C. Erfling3 reported an inflection in the thermal expansivity curve at 36°C. These temperature-dependence curves are reversible with no hysteresis being detected. No discontinuity or inflection has been observed in the heat capacity&apos; or the paramagnetic susceptibility.&apos; Likewise, no one has noted a change in crystal symmetry. In the present investigation Young&apos;s modulus, internal friction, coefficient of expansion, electrical resistivity, illustrated in Fig. 1, lattice constant, Fig. 2, thermal electromotive force, Fig. 3, and paramagnetic susceptibility, Fig. 4, were measured over an extended temperature range. The samples were prepared in two ways: (1) By cold pressing a sintered electrolytic powder compact,&apos; and (2) by electroforming from an aqueous solution. The electroformed samples were prepared by R. A. Ehrhardt and G. Bittrich by plating on copper or nickel tubes from an aqueous solution according to the method of Brenner, Burkhead, and Jennings.&apos; The pressed powder samples (method 1) were finally annealed at 1400°C in purified helium: the electroformed samples, packed in powdered chromium, were vacuum-annealed at 1000°C. From the composition of the original powder&apos; the purity of the pressed powder samples is estimated to be 99.8 pct Cr. Spectrochemical analysis furnished by E. K. Jaycox, revealed only slight traces of impurities in the electroformed sample (less than 0.001 pct), neither iron nor nickel being detected. Chromium deposited by this method is reported to contain approximately 0.05 pct 0.- Methods and Data Young&apos;s Modulus: For measurement of Young&apos;s modulus, an annealed, pressed powder sample, 0.114 x0.237x1.845 in., and an electroformed sample, 0.10 in. od, 0.01 in. id, and 1.86 in. long, were prepared. The resonant frequencies of the rod samples in forced longitudinal vibration were measured at a series of temperatures from —192" to 200°C by a method previously de~cribed.6,7 Because the samples were nonferromagnetic, iron- silicon or molybdenum permalloy tips (0.013 in. thick) were soldered to the ends. The softening temperature of the solder limited the temperature of measurement. Young&apos;s modulus, E,, of chromium at temperature, T, may be calculated from the resonant frequency of the composite rod, f,; the thickness of the tips, t; the length of the chromium sample at 25°C, l,.; the density of chromium at 25°C, pc (7.20 g per cc);5 and the thermal expansivity, Al/l25. Modulus differences for two temperatures, ET and E, are accurate to +0.002x10" dynes per cm2. ET = 4PAlc+2tyf Young&apos;s modulus at 25°C, Fig. la, is 28.2~10" dynes per cm&apos; (40.8xlO" si) in wrought chromium (upper curve). The modulus of the electroformed sample is apparently lower due to cracks. From the modulus measurements two transitions were observed: one with a critical temperature at 37°C, the other at —152°C. Internal Friction: The internal friction, 1/Q, was determined from the width, Af, of the strain amplitude-frequency curve at 0.707 times the strain amplitude at resonance;&apos; the internal friction, 1/Q, then equals Af/f. Fig. lb shows a sharp peak in internal friction at 38°C. The internal friction of the electroformed sample had a similar maximum. Thermal Expansion: The expansivity measurements, shown in Fig. 2, covering the temperature range —195" to +400°C were made by D. MacNair in an interferometric dilatometer8 sing a sample consisting of three pyramids 0.25 in. high prepared from wrought electrolytic chromium. Below —120°C the experimental points deviated up to ±2x10 from the drawn expansivity curve in Fig. 2 because of decreased precision of the quartz wedge thermometer at low temperatures. Near 38°C the thermal expansivity curve, Fig. 2, goes through an inflection corresponding to a minimum in the coefficient of expansion, Fig. ld, and a relative volume decrease on heating. Electrical Resistivity: The variation of electrical resistivity with temperature measured by the po-tentiometric method is also shown in Fig. l. The resistivity of the wrought chromium (upper curve) at 20°C is 13.6 microhm-cm; of electroformed chromium (lower curve) 12.8 microhm-cm. The lower value reflects higher purity and agrees closely with a published value." A minimum at 40°C occurs in the resistivity curves of both samples. No conclusive evidence for a transition near — 150°C was observed, but the points, Fig. 1, appear to deviate from a smooth curve between —120" and —160°C.

Jan 1, 1952