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By I. Nilsson, G. Sundkvist, A. Danielsson

A spectrochemical method of slag analysis is descibed which utilizes fusion of the sample with a flux, then cooling and crushing to put all samples into a common form. The powder, is then fed onto adhesive tape and passed through a spark by means of a special device to provide very reproducible discharge conditions. Combined with a direct-reading optical emission instrument, this system provides speed and accuracy of antilysis heretofore unobtainable for. slag samples by any other method. THERE are many suggested methods for spectro-chemical analysis of slags. The most important are fusion-pellet methods and solution methods. In our company the fusion-pellet method has been used earlier, and is described by Lounamaa. W. H. Tingle and C. K. Matocha have made refinements of the pellet method.' A Committee of the British Iron and Steel Research Association has investigated different methods of spectrochemical analysis of slags. Its purpose was to formulate a method which would give the complete analysis in 30 min with an accuracy of ±3 pct. The committee recommends a solution technique, which was thought to correspond best to their requirements. The method to be described is a fusion-tape method. The sample is isoformed by fusion and grinding. The powder is analyzed with the tape technique. The tape machine and some of its applications are described earlier.4'5 The powder to be analyzed is continuously fed onto a moving adhesive tape, which passes through the spark gap. The sparks break through the tape and vaporize and excite the material to be analyzed. The most outstanding feature of the tape machine is its very high reproducibility. The concept of isoformation has been introduced earlier.= It means a pretreatment of the samples to make them uniform from the spectralanalytical point of view, i.e., cancelling systematic differences between the samples. The purpose of the isoformation is to reduce the influence of particle size, as well as chemical and mineral composition. Before describing the isoformation by fusion plus grinding, which is the purpose of this paper, a few words will be said about two other types of isoformations, i.e., grinding with buffer and ion exchange plus grinding as they have been considered during the application of the tape technique for slag analysis. Grinding with buffer might be a sufficient isoformation in some special cases, where the mineral composition is defined. The time of grinding must, however, be rather long, at least 5 to 10 min, even when using such an efficient mill as the Swing Mill from Messrs. Siebtechnik, Mühlheim, Germany. This type of isoformation is not recommended due to its limitations and due to the fact that the gain in time compared with the fusion technique is almost negligible. The principle of ion exchange plus grinding and some of its applications have been described.5 The following elements of interest in slag analysis can be determined in this way: Na, K, Cu, Mg, Ca, Ba, Zn, Pb, Al, Cr, Mn, Fe, Co, Ni. The very important element Si, however, cannot be analyzed with this method. Although the ion exchange isoformation for slag analysis cannot be looked upon as a routine method, it can be used for

Jan 1, 1962

By B. E. Eakin, A. L. Lee, M. H. Gonzalez

Expeximental viscosity and density data of lour natural gases are presented for temperatures from WO to 340F and pressures from 100 to 8.000 psia. A correlation is also discussed and results reported. INTRODUCTION This investigation is one of several efforts by the authors to provide viscosity data for pure hydrocarbons and mixtures. Results of some pure hydrocarbons and their mix tures have been presented." Several predictive methods and correlations have also been reported.'," This paper presents the experimental viscosity and density data of four natural gases and confirms a correlation in a previous study. APPARATUS The viscometer was described previously. 2,6 A bank of stainless steel pycnometers was included to determine density in conjunction with the viscosity measurements (Fig. I). DATA AND MATERIALS The natural gases were furnished by the Atlantic Richfield Co. (Samples 1 and 21, the Continental Oil Co. (Sample 3) and the Pan American Petroleum Corp. (Sample 4). Table 1 shows the composition of these gases obtained by mass spectrometer analysis. The pure component viscosity data used in the correlation have been published."".'"', '" Experimental and calculated viscosity data for the natural gases are presented in Tables 2 through 5. CORRELATION Starling and Ellington"' have reported several semi-empirical expressions based to a certain extent on the theory of viscosity advanced by Born and Green.' The final expression presented by Starling and Ellingto'n is: i'. (micropoise) -- /t,, exp rX(7') p ..... (I) Eq. 1 was modified by Lee et 01." to represent mixture and pure component data simultaneously. This equation has the form: F = K exp [XpY] :.........(2) where K=- (7.77 i 0.0063M)Ti 122.4 + 12.9M +T .................(3) ' ' ' X= 2.57 + ---1914.5+ 0.0095M .... (4) Y = 1.11 -r 0.04 X ........(5) Over the pressure and temperature range studied in this investigation, this equation represents the data on methane, ethane, propane, n-butane and four methane-n-butane mixtures with a standard deviation* of ? 1.89 per cent. Density data by Sage and Lacey" were used to fit the equations. Eq. 2 gives reliable values of the viscosity of light hydrocarbons without prior knowledge of experimental viscosity values, but accurate density data must be available. Eq. 2 was used to predict viscosity values for the natural gases studied in this paper. The particular set of parameters contained in Eqs. 3 through 5 gave values that reproduced experimental data within 25 per cent. Needless to say, in using Eq. 2 better density values will give better viscosity results. However, those engineers who do not have accurate density values at hand or the time, or such facilities as a library, computer, etc. readily available, may feel uneasy in using Eq. 2. Therefore, the authors sought the easiest but by no means best density prediction method, which was reported by Kaf 30 years ago. A set of density values, which were calculated based on Kay's method, and a generalized compressibility factor chart were used to fit Eq. 2. The results are:

Jan 1, 1967

By G. C. Kuczynski

Two particles in mutual contact form a system which is not in thermo-dynamical equilibrium, because its total surface free energy is not a minimum. If such a system is left for a certain period of time, the bonding of the two particles will take place in order to decrease the total surface area, even though the temperature is lower than the melting point. This phenomenon of bonding of two or more particles with the application of heat only and at temperatures below melting point of any component of the system will be called sintering, although the powder metallurgists use this term in a broader sense, including the presence of molten phase and pressure. It is the objective of this paper to study this process and the mechanisms involved in it. This problem is of utmost importance to powder metallurgy and powder ceramics, and its technological aspects have been studied for a good many years. However, the powder metallurgical operations are too complex and include too many superimposing mechanisms, and too many variables for a direct study. It was therefore advisable for the purposes of this study to reduce the variables to a minimum. In this work the radius of the interface formed during bonding in a simple system composed of a spherical particle and a large block of the same metal was studied as a function of time and temperature. It is believed that the mechanism involved in this simple process is fundamental to any sintering operations. Previous Work J. Frenkell was the first to make a serious attempt to develop a theory of sintering. He assumed that the process consists of a slow deformation of crystalline particles under the influence of surface tension which reduces to a viscous flow where the coefficient of viscosity n is related to the self-diffusion coefficient D by the following equation 1 _ D D6kT [1] where 6 is interatomic distance, k the Boltzman constant, and T the absolute temperature. This type of viscous flow of a crystalline substance is essentially different from the ordinary plastic flow. The latter is a specific propcrty of crystals and cannot take place in amorphous bodies. According to Frenkel this viscous type of flow is due to the diffusion of the holes or vacancies arising in the lattice. He was able to derive an equation relating the growth of the interface between two spherical crystalline particles or between a particle and a semi-infinite crystal (Fig 1) to time t at constant temperature. This relationship can be written as follows: x²— 3/2 a/nt [2] where x is the radius of the interface assumed to be circular, a the original radius of the sphere and a surface tension of the material. The other assump- x tions are that x is less than 0.3 and that during the period 1 the original radius, a, of the metallic particle did not change appreciably. A. J. Shaler and J. Wulff2 followed closely the ideas of Frenkel in their theory of sintering of a mass of metallic powder. Neither Frenkel nor Shaler has validated his theoretical speculations with conclusive experimental data. Two measurements reported by

Jan 1, 1950

By S. R. Mitchell, M. Weissman, J. H. Brown

It has been demonstrated that for fine grinding operations, the energy input (E) and the product size modulus (k) are related by an equation of the form E = Ak1-n (1) where "A " and "n" are constants.(1) Also, it was suggested that "A" is determined by the comminuting device and that "n" depends primarily on the rock. It was the purpose of this investigation to study these constants by determining the form of an energy-size reduction relationship over an extreme range of energy inputs, and at the same time, to evaluate the minimum particle size produced during grinding. The comminution system consisted of a stainless steel laboratory ball mill in which quartz was ground at 60% solids for times ranging from 1 minute to 700 hours. The mill which measured 10" diameter by 12" long was loaded with about 60 pounds of balls ranging in diameter from 3/8" to 1-1/8" and was driven at 55 r.p.m. Quartz was selected for the tests because it is a brittle material for which size distribution data were expected to show a constant distribution coefficient regardless of product size.(1) This permitted specification of product size by a size modulus alone. The quantities measured were the time of grind (measure of energy input) and the specific surfaces of the products. Since no accurate method is available yet for determining particle size directly in the very fine size ranges specific surface data were used for this purpose. The relationship between these two quantities may be expressed as where Cis a constant which includes a shape factor, p is the specific gravity of the solid,_S is the specific surface and D is an equivalent particle diameter. For a material in which the distribution coefficient does not change with size the term (l/S) may be substituted for k in equation L if the proportionality constant "A" is changed correspondingly. This substitution was made in this work but in addition, screen analyses were made on several products to verify the equivalence of the two methods of size specification. EXPERIMENTAL DETAIL QUARTZ Quartz from two sources was tested. One series of tests was made on a clear, high purity crystalline quartz from Brazil and Arkansas, and another series was run on massive quartz fromNewHampshire. Grinding charges of the high purity material consisted of 2000 gram samples of -3-8 mesh particles, while the charges of the New Hampshire quartz were 500 gram samples of a -35+48 mesh screen fraction. The difference in the sizes of these two series of charges was not significant in view of the extreme size reduction accomplished, and the New Hampshire samples were smaller to permit ginding to finer sizes in shorter times than required for 2000 gram lots. SAMPLING AND ANALYSIS Separate charges were ground for all tests in which the grinding times were less than one hour and for all tests on New Hampshire quartz since these products were difficult to sample accurately without affecting the entire charge. During the long time grinding tests the mill was

Jan 1, 1961

By P. Chiotti, R. J. Hecht

Vapov pressure data of solid and liquid alloys along with phase diagram data WERE employed in calculating the thermodynamic properties for the Ca-Zn system. The relative excess partial molal free energy for zinc in the liquid alloys was fitted to the relation wnvre i is we tvmpvrutiirc in ucgrees neivtu ana in is the tr~ole fvaction. Relations for the standard free energy of fortnation for the eight cornpounds were developed. ThE phase diagram for the Ca-Zn system as determined by Messing et a1.' is reproduced in Fig. 1. Earlier work on the phase diagram has been summarized by Hansen and Anderko2 and the very limited thermodynamic information for this system has been reviewed by Hultgren et a1.3 The heat of fusion for the compound CaZn2 has recently been determined to be 9260 cal per mole.4 Solubility data for calcium in zinc for the temperature range of 650" to 725°C were fitted by the relation and reported by Johnson and Dilloon5 In the present investigation zinc vapor pressure data were obtained by both the dewpoint and Knudsen methods in order to calculate the thermodynamic properties for liquid and solid alloys. The dewpoint method as employed in this work requires that only one of the components be significantly volatile, in this case zinc, and that the vapor pressure of pure zinc be known. For calculation purposes the vapor pressure data for pure zinc3 were fitted to the relations; for temperatures below 692.7K, the melting point of zinc. The corresponding free-energy relations for vaporization and sublimation of pure zinc are respectively. Some of the calculations also required knowledge of the free energy of fusion of zinc and calcium and the free energy for the 721°K allotropic transformation of calcium. The free energy of fusion of zinc was taken to be the difference between Eqs. [4] and [5], or The free energy of fusion and transition for calcium were based on the enthalpy values of 2040 and 222 cal per g-atom, respectively, as reported by Chiotti et al. 4 and the heat capacity data given by Hultgren et a1.3 and are represented by the relations, In the latter expression the difference in heat capacity between the a and 0 forms of calcium was neglected. MATERIALS AND EXPERIMENTAL PROCEDURES The metals used in preparing the alloys were Bunker Hill slab zinc, 99.99 pct pure, and Ames Laboratory calcium. The average impurities in the calcium in parts per million by weight, ppm, were N-17, H—77, Ba—216, Mg-15, Si—7, C—157, and Fe—9. The zinc was cleaned in dilute nitric acid and the calcium was used as received. The calcium was stored and handled in an argon-filled glove box and when removed from the dry box it was kept in an argon-filled desiccator. Tantalum was found to be a suitable alloy container. The alloy components were contained in a 1-in.-diam cylindri-

Jan 1, 1968

By F. Weinberg

The distribution of solute during the progressive solidification of dilute Sn-Ag alloys was determined in both solid and liquid as a jbnction of growth rate, rod diameter, temperature gradient, and solute concentration, using radioactive tracer techniques. The results indicated that partial mixing occurs in the liquid during freezing for most of the conditions investigated, approaching complete mixing for larger rod diameters (>2mm) and slow growth rates. Measured solute distributions were fitted to distribution functions based on diffusion controlled solute transport in the liquid state, by treating the normal solute diffusion constant and distribution coefficient as variable solute transport and segregation parameters. Values of these parameters were determined for the growth conditions investigated. Observations were made to determine whether mixing in the liquid was primarily due to volume changes in freezing or to differences in density between solute and solvent. The former effect was investigated using a selected Sn-Bi alloy as a solvent with zero volume change on fusion, and the density effect was studied using vertical solidifications of Sn-In and Sn-Ag alloys, systems having, respectively, very small and large differences in density between solute and solvent. The result indicate that mixing in the liquid cannot be primarily attributed to either volume changes or density differences. THE solute distribution in progressively solidified rods of dilute alloys, in which there is complete mixing in the liquid, is given by the expression1 where C, is the solute concentration in the solid at the point where the fraction g of the material has solidified, C, is the initial solute concentration and k is the equilibrium distribution coefficient. If it is assumed that there is no significant con-vective mixing in the liquid, then the solute distribution in the initial portion of a semi-infinite rod has been calculated to be2 where t is the time from the start of freezing and x1 is the distance in the liquid ahead of the solid-liquid interface at time t. Expressions [1] and [2] describe the limiting cases of solute distribution in the liquid, i.e., complete mixing, or no convective mixing; in practice it might be expected that the solute distribution would be somewhere between these two cases. pfann3 has suggested that Eq. [I] is applicable for partial mixing, as well as complete mixing, if an effective distribution coefficient k1 is used instead satisfactorily describes solute distributions in rods of magnesium, n-g, and b-sb6 alloys that were progressively solidified from the melt without mechanical mixing. The values of kl were observed to depend on the growth rate4-6 and the temperature gradient.6 The purpose of the present investigation was to measure solute distributions in both the solidified rods and the melt as a function of the solidification variables, and to determine to what extent the observed solute distributions could be described by the above theoretical expressions, keeping in mind that expressions [2] and [3] have been derived solely for the case of diffusion controlled solute transport in the liquid and are not theoretically applicable to

Jan 1, 1963

By R. A. Swalin, B. W. Roberts

ONSIDERABLE controversy has centered about the existence of an order-disorder transformation in alloys in the composition vicinity of 75 atomic pct Ni and 25 atomic pct Cr. All the evidence to date for such a transformation is of a rather indirect nature, since direct X-ray confirmation is difficult to obtain as a result of the small difference in X-ray scattering factors between nickel and chromium. As a result, superlattice lines would be very faint if present. Several investigators1-5 have observed anomalies in the electrical resistivity vs temperature curves of alloys in the vicinity of 25 atomic pct Cr. The observed change is characterized by a rather broad resistivity peak at about 540°C. A discontinuous drop in resistivity is normally associated with ordering. However, Taylor and Hin-ton4 have rationalized that the discrepancy could conceivably be caused by changes in the Brillouin zones upon ordering. The same investigators measured the specific heat as a function of temperature and have also observed anomalies in the vicinity of 540°C. Nordheim and Grant5 have proposed, because of the absence of superlattice lines and the absence of a discontinuity in the resistivity vs temperature curve, that long-range order is improbable and that the results could be better explained by short-range ordering. To date, however, the question has not been resolved satisfactorily. Neutron diffraction techniques provide a method for partially resolving the question since nickel and chromium have considerably different atomic scattering factors in the case of neutrons and, therefore, superlattice peaks, if ordering occurs, should be readily discernable. The results of a neutron diffraction investigation* coupled with a dilatometric in- vestigation are reported here. An alloy containing 28.0 atomic pct Cr was prepared from carbonyl nickel and electrolytic chromium in the form of a 1-in. diam x 10-in. ingot. Melting was done in an alumina crucible and casting in a graphite mold. The ingot was homogenized at 1250°C for seven days, after which a 5-in. long piece was cut into shavings in order to obtain a suitable neutron diffraction sample. Hardness samples and samples suitable for dilatometric analysis (% in. diam x 2 in.) were prepared from the remainlng piece. The neutron diffraction sample, along with a hardness sample, was annealed at 501°C for seven days, at which time the temperature was lowered to 479°C and held at this temperature for twenty-four more days. The hardness was observed to increase from 133 to 147 Vhn during the course of the annealing treatment, in qualitative agreement with Nordheim and Grant." Neutron diffraction traces taken at room temperature and at 42 °K are shown in Fig. 1. The insert shows the trace taken at room temperature with short counting times and thus only fair statistical accuracy. No evidence for superstructure reflections of Cu3Au-type ordering is observed. The statistical accuracy of the 4.2°K trace was improved by using longer counting periods. Several weak intensities are observed at small angles, including an aluminum reflection from the cryostat assembly and a very

Jan 1, 1958

By H. C. Gatos, J. T. A. Pollock, V. Sadagopan

PREVIOUS investigations on the effect of neutron irradiation on the superconducting properties of bulk transition metals indicated that little or no change occurred in poly crystalline disordered materials, particularly if cold-worked prior to irradiation. Only the superconducting properties of ordered intermetal-lic A15 compounds were found to be sensitive to irradiation, and the observed reduction in Tc and increased magnetic hysteresis or critical current densities were attributed to the introduction of flux pinning vacancy clusters and disorder in the lattice. Since the former would also be expected to be present after irradiation in disordered alloys, the general conclusion was that the disordering effect of the irradiation was the primary reason for the changes observed. The present note reports some unusual effects of neutron irradiation on cold worked Nb(Cb)-Ti-V alloys. Specimens with nominal composition as given in Table I were prepared by repeated arc melting, taking the usual precautions, and homogenized at 1000°C for 24 hr prior to a reduction of 15:l by cold rolling. Samples of size 1 by 1 by 25 mm were sliced from the rolled strips. The sarmples chosen for irradiating were sealed in a cadmium tube (wall thickness 1 mm) prior to exposure in a fast neutron flux of 3.7 x 1019 neutrons per sq cm in the swimming pool reactor of Union Carbide Corporation&apos;s Sterling Forest Reactor. During irradiation the temperature of the samples was maintained below 70°C. The irradiated samples were stored at room temperature for approximately 1 year allowing the radiation levels to reach a safe handling value. Tc measurements were made employing a self-inductance coil technique. Jc data at 4.2"Kwere obtained using a four lead resistivity method. The latter measurements were made at the Francis Bitter National Magnet Laboratory by setting the perpendicularly applied field and increasing the direct current through the sample until a potential of about 10-6 v was detected. Tc values for the nonirradiated samples are listed in Table I together with the change in Tc after irradiation. A decrease in Tc of 0.25 or 0.3"K was measured for three of the four samples, the exception being the Nb-30 at. pct V-20 at. pct Ti alloy which showed a decrease of 0.5K. Jc vs applied field data are plotted in Fig. 1 for the Nb-20 at. pct V-30 at. pct Ti alloy. These curves are typical of the four alloys and surprisingly indicate that a degradation occurred in all the samples after irradiation. Hc2 values, obtained by extrapolation to zero current, are 8 to 15 pct lower than those measured for the nonirradiated specimens. The reasons for the degradation of Tc and jc ob-served in the alloys on exposure to neutron irradiation are not immediately apparent. The temperature at which the samples were maintained during bombardment (<70°C) and stored (20°C) make it highly unlikely that any significant change in the dislocation density of the materials through thermally activated migration and annihilation would occur. The explanation must lie in the type of damage produced by the bombardment. McEvoy et a1. reported no change in superconduct-

Jan 1, 1970

By Rex E. Cheek, Donald E. Menzie

The fluid mapper, a model relatively new to the petroleum industry, was used to study the effect of various mobility ratios on the areal sweepout efficiency for two typical spacing patterns. The experiments indicated the areal sweepout for the five-spot pattern varied from 51.8 per cent at a mobility ratio of 0.093 to 89.1 per cent for a ratio of 24.4. In the direct line-drive pattern, the areal sweepout varied from 58.7 per cent at a mobility ratio of 0.089 to 92.0 per cent at a ratio of 24.8. The fluid mapper may prove useful to determine quickly the effective areal sweepout for any well spacing. INTRODUCTION The effect of the mobility ratio on the areal sweepout efficiency during primary depletion of a water drive reservoir and during secondary recovery operations has recently become more widely recognized The mobility ratio, M, is defined as: _ kjp (displaced fluid) k/u (displacing fluid) where k = effective permeability u = viscosity When the displacing fluid is miscible with the displaced fluid, the permeabilities are equal and the mobility ratio may be expressed as a ratio of viscosities: u (displacing fluid) u (displaced fluid) The mathematical analysis for calculating the areal sweepout efficiency with mobility ratio other than unity is quite tedious.5,11 An experimental method by using a potentiometric model2 in a step-wise manner and in using an X-ray shadowgraph model3,12 with continuous fluid flow has been published. Several types of models have been used to study oil and gas reservoir performance. These models may be classified as physical, electrolytic, potentiometric, X-ray shadowgraph, and fluid mapper. The X-ray shadowgraph and fluid mapper are models relatively new to petroleum engineering and have not been used extensively for specific applications. The theory of the fluid mapper indicates it may be adapted to a study of areal sweepout efficiencies for various mobility ratios as well as other complex problems in petroleum engineering. The use of the model and results obtained when using displacing and displaced fluids of equal densities but different viscosities are discussed. FLUID MAPPER DEVELOPMENT H. S. Hele-ShawV eveloped a fluid mapper model based on the hydrodynamic theory. This theory is that in the steady flow of a viscous fluid between horizontal parallel plates spaced a very small distance apart, the effect of viscous shear is of such magnitude in the vertical plane that the motion in the horizontal plane is practically undisbrted streamline flow. Techniques developed by Hele-Shaw were not expanded and no additional developments in fluid mapper models were forthcoming until A. D. Moore, University of Michigan, published a series of papers starting in 1949, which describe the construction of a flexible and inexpensive model&apos;h nd the fluid mapper technique for the solution of certain problems.&apos; The essential components of the fluid mapper as developed by Moore are::.ln 1. A base slab. made of dental stone (calcium sulphate and water). which is shaped to represent the prototype. 2. Plate glass slightly larger than the slab and spaced

Jan 1, 1956

By D. B. Jugle, J. B. Darby

The equilibrium solubility limits of the transition elements Ti, V, Cr, Fe, and Co in liquid tin were determined in the temperature range from 827" to 1211°K. Since the equilibrium concentration of Ti, V, or Cr in liquid tin is moderately dilute in the temperature range investigated, the solubility data were employed to estimate the partial molar enthalpy and the partial molar excess entropy of solution for the three solutes in liquid tin. The reference state s for titanium or vanadium were selected relative to the solute in an intermetallic compound. The thermodynamic quantities for chromium were calculated relative to pure chromium. RECENT experiments1 have shown that the rate of solution of several transition elements in liquid aluminum can be qualitatively predicted from a knowledge of the temperature dependence of the solvent-rich liquidus. In order to extend the analysis to binary tin systems,&apos; it was necessary to know the equilibrium solubility of several of the first-long-period transition elements in liquid tin. Since solubility data for the tin-rich portion of binary systems are meager, the present investigation was undertaken to determine the equilibrium solubility limit of Ti, V, Cr, Fe, and Co in liquid tin in the temperature range from 850" to 1200°K. EXPERIMENTAL PROCEDURE A useful technique3 to determine the equilibrium solubility of a solute in a liquid-metal solvent consists essentially of placing a pure solute in contact with a pure solvent and allowing sufficient time for equilibrium to be attained at the temperature of interest. Subsequently, the solvent is removed and the equilibrium solute concentration in the solvent is determined by an appropriate analytical procedure. The technique can be applied very successfully to relatively simple systems such as the copper-germanium system.3 However, in many binary systems of transition elements with tin, several intermediate phases are known to exist and precipitation of the phases from solution, in solubility experiments, can lead to an erroneously large solubility. The difficulty was avoided in the present work by equilibration of the most solvent-rich intermediate phase with the pure solvent over a range of temperatures. The compounds employed in this investigation, and the temperature range over which each compound is in equilibrium with liquid tin are listed in Table I. No stable intermediate phases are reported in the chromium-tin system4 and, therefore, pure chromium was utilized as the solute. The stoichiometric compounds were prepared by arc-melting on a water-cooled copper hearth in a helium-argon atmosphere. The elemental tin employed in.this investigation was 99.998 pct pure, and typical lot analyses of the transition elements were reported elsewhere.&apos; A sample of an appropriate solute, in the form of a button and in sufficient quantity to exceed the equilibrium solubility, and the pure tin solvent were placed in a high-purity recrystallized-alumina crucible. The crucible and contents were placed in a horizontal furnace under a high vacuum (10-6 torr) and held at temperature for periods varying between 24 and 72 hr. A water-cooled quench tank, partially filled with diffusion-pump oil, was attached close to one end of the

Jan 1, 1970

By M. G. Benz, J. F. Elliott

A new type of solution calorimeter has been constructed to measure heats of mixing, enthalpy increments, and heats of fusion, formation, and reaction at temperatures above 1000°C. With it, measurements hove been made of the partial molar heats of mixing joy tin in the Cu-Sn system from 0 to 13 at. pet Sn at 1127°C and fir nickel in the Cu-Ni system from 0 to 9 at, pet Ni at 1200°C. The integral molar heat of mixing for the Cu-Sn system is strongly negative with a calculated minimum value corresponding to the composition Cu3Sn. The integrnl molar heat of mixing for the Czi-Ni system is slightly positive at these temperatures. These integral molar values have been measured with a probable error less than 30 cal per g-atom at an atom fraction of nickel or tin equal to 0.1. The estimated probable error for the Cu-Sn system is approximately 450 cal for a measured heat of solution of approximately 9000 cal. For the Cu-Ni system it is approximately 700 cal for a measured heat of&apos; solution of approximately 60 cal. A new type of heat of solution calorimeter was designed and constructed to provide a more accurate experimental method for determining heat effects at high temperatures, (T > 1000°C). With it, partial molar heats of mixing, enthalpy increments, heats of fusion, and heats of formation and reaction for metal systems can be measured directly. Previously it has been necessary to rely upon the use of the dropping calorimeter or the van&apos;t Hoff equation applied to equilibrium data for this information. For the most part equilibrium data are scarce and difficult to obtain in many cases, and it has been necessary to rely upon the "method of mixtures" for the desired enthalpy measurements. An earlier study1 provides a prototype for this work. The "method of mixtures" as outlined by K. K. kelley,2 "consists of dropping the substance under investigation from a furnace, with controlled and measured temperature, into a calorimeter operating at or near room temperature.&apos;&apos; In this manner the change in the total heat content (enthalpy) of the substance between the temperature of the furnace and that of the calorimeter is determined. In order to determine high-temperature heats of mixing, enthalpy increments, and heats of fusion, formation, and reaction, it is necessary to take the difference between two such measurements, one before and one after the occurrence of the enthalpy change being studied. The desired enthalpy change is then the small difference between two large numbers. Small errors in the large numbers are included as large errors in this small difference. It is this difficulty that must be avoided for accurate measurement of the desired functions. The procedure with the new solution calorimeter consists of dropping the substance under investigation from a furnace, with controlled and measured temperature, into a solution calorimeter operating in the temperature range of 1000o to 1200°C. Measurements in the range of 1200" to 1600°C are planned for the future. By selection of the composition and

Jan 1, 1964

Time as a Factor in the Making and Treating of Steel. By John Johnston. (Howe Lecture) (T.P. 1478) Influence of Chromium and Molybdenum on Structure, Hardness and Decarburization of 0.35 Per Cent Carbon Steel. By R. F. Miller and R. F. Campbell. (T.P. 1345) Mechanical Properties of Iron-manganese Alloys. By F, M. Walters, Jr., I. R. Kramer and B. M. Loring. (T.P. 1369) The Instability of Low-expansion Iron-nickel-cobalt Alloys. By I. R. Kramer and F. M. Walters, Jr. (T.P. 1370) The S-curve of a Chromium-nickel Steel. By B. M. Loring. (T.P. 1383) A Magnetic Determination of the A3 Transformation Point in Iron. By B. A. Rogers and K. 0. Stamm. (T.P. 1388) Rapid Tension Tests Using Two-load Method. By A. V. de Forest, C. W. MacGregor and A. R. Anderson. (T.P. 1393) Effects of Scrap in the Blast-furnace Burden. By C. L. T. Edwards. (T.P. 1270) Temperature Gradients through Composite Carbon Columns and their Application to Blastfurnace Linings. By F. J. Vosburgh and M. R. HatFIeld. (T.P. 1363) Correlations of Some Coke Properties with Blast-furnace Operation. By H. W. Johnson. (T.P. 1402) The Technical Cohesive Strength and Yield Strength of Metals. By D. J. McAdam, Jr. (T.P. 1414) Effects of Eight Complex Deoxidizers on Some 0.40 Per Cent Carbon Forging Steels. By George F. Comstock. (T.P. 1417) Significance of the Bessemer End Point. By H. T. Bowman. (T.P. 1428) Structural Diagrams of Nickel Irons and Steels. By J. T. EasE and N. B. Pilling. (T.P. 1432) Carbides in Low-chromium Steel. By Walter Crafts and C. M. OFFenhauer. (T.P. 1436) Hardenability Calculated from Chemical Composition. By M. A. Grossmann. (T.P. 1437) Recrystallization of Silicon Ferrite in Terms of Rate of Nucleation and Rate of Growth. By J. K. Stanley and R. F. Mehl. (T.P. 1438) The Solubility of Carbon as Graphite in Gamma Iron. By R. W. Gurry. (T.P. 1440) An Evaluation of the Factors Affecting Iron Oxide in the Open-hearth Bath. By J. E. Gould and H. J. Hand. (T.P. 1442) Precision in Creep Testing. By J. A. Fellows, Earnshaw Cook and Howard S. Avery. (T.P. 1443) Lattice Relationships in the Decomposition of Austenite to Pearlite, Bainite and Martensite. By G. V. Smith and R. F. Mehl. (T.P. 1459) Rate of Nucleation and Rate of Growth of Pearlite. By F. C. Hull, R. A. Colton and R. F. Mehl. (T.P. 1460) Weight Change as a Criterion of Extent of Decarburization or Carburization. By R. W. Gurry. (T.P. 1470) Observations in the Making and Use of Sulphite-treated Steels. By E. L, Ramsey and L. G. Graper. (T.P. 1476) Diffusion in Metal Accompanied by Phase Change. By L. S. Darken. (T.P. 1479) Engineering Properties of Heat-resistant Alloys. By Howard S. Avery, Earnshaw Cook and J. A. Fellows. (T.P. 1480) Distribution of Manganese and of Sulphur between Slag and Metal in the Open-hearth Furnace. By L. S. Darken and B. M. Larsen. (T.P. 1481)

Jan 1, 1942

By D. F. Kaufman, E. D. Levine, L. R. Aronin

ThE primary slip plane in hcp metal crystals can usually be inferred from the c/a ratio. Basal slip is the primary system at room temperature for zinc, cadmium, magnesium, cobalt, and rhenium, all of which have a c/a ratio greater than 1.60. Zirconium and titanium, on the other hand, each with a c/a ratio of 1.59, deform primarily by (1010) [ll20] prismatic slip. The increasing tendency toward prismatic slip in these materials is generally thought to be associated with the increase in atomic density on (10i0) as c/a is reduced. Beryllium, with a c/a ratio of 1.57, would, according to the above scheme, be expected to exhibit prismatic slip. However, in beryllium of commercial purity, the primary deformation mode at room temperature is basal slip.&apos; The ratio of critical shear stress for prismatic to that for basal slip at room temperature is approximately 5:l. The critical shear stress for prismatic slip decreases quite rapidly with increasing temperature, however, and at 500°C the ratio is approximately unity. The apparently anomalous behavior of beryllium has led to speculation that prismatic slip in this material may be restricted by an anisotropic effect of impurities,&apos; and that, if beryllium could be purified to a sufficient extent, prismatic slip might predominate at room temperature. In such a case, the conditions under which basal cleavage can occur would be suppressed, and enhanced low-temperature ductility might result. With the advent of the floating-zone method for purification of beryllium,3 it has become possible to determine experimentally the effects of purification on basal and prismatic slip. Tensile tests were performed at room temperature on single crystals oriented favorably either for basal or prismatic slip. The crystals had been purified to various extents by zone-refining, employing procedures that have been described elsewhere.4 Relative purities of the crystals were estimated by measuring the ratio of electrical resistance at 4.2"K to that at 298°K. Typical values ranged from 0.40 for commercial-purity material to 7.7 X 10"4 for crystals prepared from vacuum-distilled beryllium and subjected to six zone-refining passes. The purity dependence of the critical shear stresses for basal and prismatic slip is shown in Fig. 1. In the purity range studied, the critical shear stress for basal slip is decreased from 1.4 to 0.14 kg per sq mm. The critical shear stress for prismatic slip is decreased from 6.7 to 5.3 kg per sq mm. The decreases in critical shear stress for each deformation mode, therefore, are approximately equivalent. No purification anisotropy is observed. Because of the absolute magnitudes of the critical stresses, the ratio of the critical shear stress for prismatic slip to that for basal slip actually increases from 5:l to 38:l at the highest purity level studied. It is evident, therefore, that basal slip is the primary deformation mode in high-purity beryllium. This fact must be accounted for in any rationalization of the deformation behavior of hcp metals. This work was performed at Nuclear Metals, Inc., under the sponsorship of the U.S. Atomic Energy Commission.

Jan 1, 1964

By M. S. Seltzer

In a previous paper1 it was shown that activation energies for steady-state creep in lead sulfide single crystals varied with the concentration of electronic defects. For n-type lead-excess crystals, values for the creep activation energy, Qc, decreased from 2.9 to 2.2 ev as the carrier concentration increased from 1017 to 1017 electrons per cu cm, while Q, increased from 1.6 to 2.3 ev for p-type sulfur-excess crystals as the concentration of holes increased from 6 X 1017 to 1018 per cu cm. It was also shown in Ref. 1 that the Q, values could best be correlated with the sum of the self-diffusion activation energies in PbS, i.e., for a given concentration of electrons or holes Q, = QPb + Qs, where Qpb and QS are the activation energies for diffusion of pb210 and s35 in PbS. This result suggests that creep rates are controlled by at least two atomic defects, one for lead migration and one controlling sulfur transport. It would be most desirable to interpret these creep energies in terms of an atomic defect disorder model for lead sulfide. Studies of creep rates as a function of sulfur pressure2 indicate that the defects controlling creep rates are singly ionized lead and sulfur vacancies, VPb and Vs. This hypothesis can be more firmly established by determination of Q, under conditions of constant atomic defect concentration rather than constant electronic defect concentration as has been previously done. Therefore, several measurements have been made of Q, on crystals whose con- centration of lead vacancies or sulfur vacancies has been fixed by equilibration under appropriate sulfur pressure and temperature. It was indicated in Ref. 2 that high-temperature isotherms are available which show how the concentrations of various atomic point defects in PbS vary with composition over the entire homogeneity range for the compound. Two such isotherms are shown in Fig. 1 for 900" and 1000°K. Utilizing these isotherms the concentration of a given defect, say Vbb, may be held constant over a particular temperature range by adjusting the sulfur pressure as the temperature is changed. Thus the concentration of one defect suspected of controlling creep rates may be held constant while other defect concentrations vary. The apparent activation energy for formation of these defects can be obtained from the isotherms giving their concentration as a function of sulfur pressure. Activation energies for creep obtained under conditions of constant defect concentration can then be compared with defect formation energies and certain possible atomic defects can be eliminated as creep-controlling species. We have performed creep experiments as a function of temperature at a stress of 244 g per sq mm under conditions which fixed the lead vacancy concentration, V&apos;, at 1016, 3.2 x 1016, and 3.2 x 1017 cm-3, and one were the sulfur vacancy concentration, Vs, was held at lo17 cm-3. The results are presented in Fig. 2 as creep rate, k, vs reciprocal of absolute temperature, 1/T. Specimens 34, 35, and 36 were deformed under constant lead vacancy concentration while specimen

Jan 1, 1969

By J. H. Frye, J. L. Scott, J. W. Woods

THERE are at least two effects of temperature on the stress required to produce plastic flow. 1) Metals plastically strained for the same amount at different temperatures possess different stress-strain curves and different internal structures (e.g., nature and density of imperfections). 2) There is both experimental and theoretical evidence that the stress necessary to produce plastic flow may vary with temperature even though the internal structure remains constant.1-5 Much of the data in the literature is difficult to interpret because these two effects are simultaneously present. In the work reported here, they are considered separately. Standard 0.505 in. tensile specimens were prepared from high-purity Mond nickel and OFHC copper. Nickel specimens were annealed in hydrogen at 1150°C for 30 min and furnace cooled. This produced an average grain diameter of 1.4 mm. Copper specimens were annealed in vacuum at 525°C for 6 hr and furnace cooled. This produced an average grain diameter of 0.11 mm. Nickel tensile specimens were prestrained in tension at 500°K, whereas copper tensile specimens were prestrained in tension at room temperature. All specimens were cooled to 74°K and heated to room temperature. After these treatments all of the tensile specimens of a given set possessed identical structures. All specimens were returned to the tensile machine and tested at various temperatures and constant cross-head speed. Many of the tensile curves showed well defined yields but, in order to minimize personal judgment, the yield was always chosen as the stress corresponding to 0.2 pct plastic deformation. Results are summarized in the graph of yield vs temperature of test, Fig. 1. All of the specimens corresponding to any one line have had identical treatments and, therefore, have identical internal structures. Thus, each line gives the relation between yield stress and temperature for one constant internal structure. The lowest line, which is for annealed nickel, shows no measurable variation of yield with temperature. This is similar to Blewitt&apos;s6 results on copper single crystals, but a variation in yield with temperature was found by Andrade and Henderson&apos; in single crystals of nickel, and by Car-recker and Hibbard8 in annealed polycrystalline copper. The three other lines, which are for prestrained materials, show that there is a linear relation between yield stress and temperature when the internal structure remains constant. Such a relation might be expected from a rate equation such as that of Kauzmannl on the assump- tion that the activation energy is either independent of or linear with temperature. These results would also be expected if the shear moduli of copper and nickel are linear functions of temperature for, as Cottrell and Stokes have pointed out," in all dislocation theories of strength the flow stress appears proportional to the shear modulus of the material. In copper and nickel, since the modulus of elasticity decreases roughly linearly with increase in temperature," it is probable that the modulus of shear also decreases approximately linearly with increasing temperature. A detailed discussion of the effect of these and other factors on the temperature dependence of flow stress in face-centered-cubic metals has recently been given by Seeger.4

Jan 1, 1958

Time as a Factor in the Making and Treating of Steel. By John Johnston. (Howe Lecture) (T.P. 1478) Influence of Chromium and Molybdenum on Structure, Hardness and Decarburization of 0.35 Per Cent Carbon Steel. By R. F. Miller and R. F. Campbell. (T.P. 1345) Mechanical Properties of Iron-manganese Alloys. By F, M. Walters, Jr., I. R. Kramer and B. M. Loring. (T.P. 1369) The Instability of Low-expansion Iron-nickel-cobalt Alloys. By I. R. Kramer and F. M. Walters, Jr. (T.P. 1370) The S-curve of a Chromium-nickel Steel. By B. M. Loring. (T.P. 1383) A Magnetic Determination of the A3 Transformation Point in Iron. By B. A. Rogers and K. 0. Stamm. (T.P. 1388) Rapid Tension Tests Using Two-load Method. By A. V. de Forest, C. W. MacGregor and A. R. Anderson. (T.P. 1393) Effects of Scrap in the Blast-furnace Burden. By C. L. T. Edwards. (T.P. 1270) Temperature Gradients through Composite Carbon Columns and their Application to Blastfurnace Linings. By F. J. Vosburgh and M. R. HatFIeld. (T.P. 1363) Correlations of Some Coke Properties with Blast-furnace Operation. By H. W. Johnson. (T.P. 1402) The Technical Cohesive Strength and Yield Strength of Metals. By D. J. McAdam, Jr. (T.P. 1414) Effects of Eight Complex Deoxidizers on Some 0.40 Per Cent Carbon Forging Steels. By George F. Comstock. (T.P. 1417) Significance of the Bessemer End Point. By H. T. Bowman. (T.P. 1428) Structural Diagrams of Nickel Irons and Steels. By J. T. EasE and N. B. Pilling. (T.P. 1432) Carbides in Low-chromium Steel. By Walter Crafts and C. M. OFFenhauer. (T.P. 1436) Hardenability Calculated from Chemical Composition. By M. A. Grossmann. (T.P. 1437) Recrystallization of Silicon Ferrite in Terms of Rate of Nucleation and Rate of Growth. By J. K. Stanley and R. F. Mehl. (T.P. 1438) The Solubility of Carbon as Graphite in Gamma Iron. By R. W. Gurry. (T.P. 1440) An Evaluation of the Factors Affecting Iron Oxide in the Open-hearth Bath. By J. E. Gould and H. J. Hand. (T.P. 1442) Precision in Creep Testing. By J. A. Fellows, Earnshaw Cook and Howard S. Avery. (T.P. 1443) Lattice Relationships in the Decomposition of Austenite to Pearlite, Bainite and Martensite. By G. V. Smith and R. F. Mehl. (T.P. 1459) Rate of Nucleation and Rate of Growth of Pearlite. By F. C. Hull, R. A. Colton and R. F. Mehl. (T.P. 1460) Weight Change as a Criterion of Extent of Decarburization or Carburization. By R. W. Gurry. (T.P. 1470) Observations in the Making and Use of Sulphite-treated Steels. By E. L, Ramsey and L. G. Graper. (T.P. 1476) Diffusion in Metal Accompanied by Phase Change. By L. S. Darken. (T.P. 1479) Engineering Properties of Heat-resistant Alloys. By Howard S. Avery, Earnshaw Cook and J. A. Fellows. (T.P. 1480) Distribution of Manganese and of Sulphur between Slag and Metal in the Open-hearth Furnace. By L. S. Darken and B. M. Larsen. (T.P. 1481)

Jan 1, 1942

By John W. Cahn, John E. Hilliard

Single crystals of copper and silicon-iron were cold rolled in orientations chosen to produce individually the major components of the poly crystalline deformation texture. The orientation dependence of the recrystallization kinetics was studied relative to the primary recrystallization textures of these specimens. It is not possible to rationalize the poly crystalline re crystallized textures on this basis. The results suggest an important role of the interaction of adjacent crystals during the deformation of aggregates. BRICK and Williamson1 first noted evidence that the recrystallization temperature was different for straight-rolled and cross-rolled copper given comparable reductions. Michalak and Hibbard2 investigated quantitatively the effect of rolling procedure on the kinetics of recrystallization of cold-rolled copper and found that material cold-rolled the same amount by straight-pass, cross-pass, and compression-pass techniques developed different deformation and recrystallization textures, different 1-hr recrystallization temperatures, and different recrystallization kinetics. However, the temperature dependence of the rate of recrystallization was essentially the same for all techniques. Similar results were obtained on iron. Dunn and Koh3 in studying the recrystallization textures of cold-rolled silicon-iron single crystals reported large differences in the annealing temperatures and times required to recrystallize specimens of different orientations cold-rolled the same amount. Liu4 and Liu and Hibbard5 utilized the concept of different recrystallization temperatures for different orientations of cold-rolled crystals to rationalize the difference in recrystallization textures found in various face-centered-cubic metals. Their explanation predicted a significant difference in the recrystallization kinetics of deformed (110) [112] crystals as compared to (358) [352] crystals, and, in fact, Liu4 reported evidence that this difference existed at least qualitatively. Verbraak6 who rolled and recrystallized a series of copper single crystals found that the double (112) [111] rolled orientation is related to recrystallized cube texture formation. The effect of rolling intensity on the relative amount of cube component suggests an important role of recrystallization kinetics. Walter and Hibbard7 while trying to relate the behavior of rolled and recrystallized single crystals of silicon-iron to the behavior of polycrystal-line silicon-iron again noted that differences in recrystallization kinetics were necessary to rationalize their results. The purpose of the present investigation was to obtain quantitative kinetic data on the recrystallization of single crystals of both copper and silicon-iron rolled to initial isolated individual orientations corresponding to those found in the polycrystalline deformation textures. These data might indicate the importance of the deformation orientation, i.e. the range of available nuclei, the deformed matrix orientation into which these nuclei grow, and the orientation dependence of their growth rate during primary recrystallization on this texture. I. COPPER Experimental Procedure—A single crystal of 99.98 pet purity copper, grown by the Bridgman technique, was purchased in the form of a bar 1 1/2

Jan 1, 1962

Pulverized Fuel Combustion The apparatus shown in the illustration is a complete equipment f r supplying coal to boiler and metallurgical furnaces, cement and lime kilns, dryers, and to all furnaces requiring a temperature and quality of fame under prompt and easy control. It is claimed that the Aero Pulverizer makes practicable the highest efficiency obtainable from burning coal. It makes coal burn like a gas, a flame, the physical and chemical character of which is regulable- a flame that may be elongated or shortened, thus placing the zone of highest temperature where needed- a flame that may be made oxidizing, reducing, or neutral, as occasion may require. The coal is burned as pulverized, and there is no storage of the powder with its attendant hazard. Artificial drying before pulverizing is not necessary if the coal supply be sheltered from rain and snow. Where the Aero is used, it is wholly a furnace question whether a dryer should be installed; it is not at all a pulverizing or storage question. Labor is reduced to a minimum. Slack coal at low cost yields its last B.T.U.

Jan 11, 1915

By Richard Rice

TURBO blowers for blast-furnace blowing have now been in use for some years, and a review of the experience gained and the present state of progress may be interesting. FIG. 1.-STEAM TURBINE DIRECT CONNECTED TO AIR COMPRESSOR. The first blast-furnace blowing engine of this type to be installed in America was built by the General Electric Co. and is installed at the oxford Furnace of the Empire Iron & Steel Co. This machine was put into operation March 8, 191.0. A duplicate unit is situated at the furnace of the Northern Iron Co. at Port Henry, N. Y., which unit was put into opera ion June 30, 1910 (Fig. 1).

Jan 5, 1914

By P. B. Leavenworth

Development in the Upper Texas Gulf Coast during 1945 resulted in the discovery of 23 new fields; one Miocene, eight Frio, three Cockfield-Yegua and eleven along the Wilcox trend. The Wilcox trend appears to he the most active at this time. Producing depths of the new Wilcox discoveries range from 2050 to 10,050 ft. Drilling along the Frio trend has been, for the greater part, projected to the deeper downdip sands. Eighteen of the new discoveries are classified as oil or condensate and five are gas discoveries.

Jan 1, 1946