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This investigation was supported by a grant from the American Iron and Steel Institute and a fellowship given by the Armco Steel Corp. We are indebted to J. F. Butler of the Jones & Laughlin Steel Corp. for a portion of the alloys used in this investigation and to R. L. LeRoy of the Armco Steel Corp. for chemical analyses.

Jan 1, 1962

By M. N. Pathasarathi, P. A. Beck, T. J. Koppenaal

EARLIER work1 indicated that in rolled and annealed copper the volume fraction of the cube-texture component may increase on continued isothermal annealing. Merlini found2 that in rolled copper the well-known increase in the amount of cube-oriented material with increasing annealing temperature3 is concurrent with grain growth, which takes place after recrystallization is complete. It was found that both in copper2 and in aluminum4 the texture resulting from primary recrystallization, even if this is followed by some grain growth, may in general comprise, in addition to the cube-texture component, four symmetrical texture components of the general type (123)[412], i.e., components not very far from the ideal orientations of the rolling texture. The volume fraction of the cube-texture component increases on further annealing at the expense of these four components.' The present work was undertaken to study quantitatively the increase in the volume fraction of the cube-texture component in the course of normal grain growth, under varying conditions. Since the grain-gr3wth behavior of A1-Mn alloys has been already investigated in some detaiL5 an A1 + 0.8 pct Mn alloy was chosen for this work, in which the extent of normal grain growth can be effectively controlled by the inhibiting effect of a dispersed second phase. EXPERIMENTAL PROCEDURE High-purity aluminum of the following composition was used: A1 99.997 pct, Cu 0.0004 pct, Fe 0.0005 pct, and Si 0.001 pct. A high-purity A1-Mn master alloy was prepared by dehydrating manganese chloride at 500°C and reacting it with molten high-purity aluminum at 750°C in a graphite crucible.5 The Mn metal reduced from the salt by the molten aluminum goes into solution in the latter. The master alloy? containing 13 pct Mn, was solidified under conditions suitable for minimizing segregation. An ingot of the final alloy, to be used in the experiments, was per pared by melting high-purity aluminum and the necessary amount of master alloy in a graphite crucible, and by pouring the molten alloy into a graphite mold preheated to 700°C through a box containing graphite baffles designed to intercept most of the aluminum oxide films floating in the melt. The melt was then solidified by placing the mold on a water-cooled copper plate and by keeping the top of the mold hot, in order to minimize piping. The 6 1/2 in. long, 11/4-in.diam ingot consisted largely of a single crystal. Chemical analysis showed that there was little segregation from the bottom (0.83 pct Mn) to the top (0.81 pct Mn). Spectrographic analysis of the ingot showed that the impurity content remained essentially the same as that of the aluminum used. The ingot was processed so as to obtain a uniformly small grain size. without banding. The ini-tial steps of rolling in the form of a flat bar and of annealing, down through a thickness of 0.4 in., are shown in Table I. At this point the bar was cut into four sections, each approximately 3 in. long, which were processed separately as shown in Table If. The table also gives the penultimate grain sizes obtained in each section by annealing prior to the final rolling of 90 pct RA. During the final rolling the strips were reversed end to end between passes. Annealing treatments were given in a salt pot furnace, with the temperature controlled to 1°C. All annealing of section "6s" was carried out above the solvus point to avoid the formation of second-phase precipitate. The penultimate annealing was sufficiently short to give a relatively fine grain size in spite of the high temperature. Section "6L" was processed in a similar manner, except that the penultimate annealing was much longer, so as to allow considerable grain growth. Section "4" was given the two final

Jan 1, 1961

By J. B. Cohen, O. F. Kimball

Aging of specimens of Au-20 at. pct Ni and Au-40 at. pct Ni has been followed with electron microscopy, resistance, magnetic measurements, and X-ray diffraction. The periodic structures observed by electron microscopy in Au-Ni alloys by Fisher and Embury are shown to be due to nickel-rich prepre-cipitates. Assuming this alloy system undergoes spinodal decomposition, our observations and available data on thermodynamic properties and elastic constants indicate that the gradient energy coefficient in this system is of the order of 10-5 ergs cm-1 . Shifts of X-ray peaks in this system have been found to be due primarily to residual stresses, not faulting. A large effect of applied stress during aging was detected by X-rays and electron microscopy for the Au-20 at. pct Ni alloy. The presence of a preprecipitation stage during aging of gold-rich Au-Ni alloys for short times below 300°C has been indicated by several previous investigations. Changes in electrical resistivity,1"3 elastic modulus,4 superparamagnetisrn,5 positions of X-ray diffraction peaks,6 and hardness7 have been studied. Electron diffraction and microscopic studies by Fukano8 and Fisher and Embury9 respectively showed that the precipitation process resulted in the formation of periodic composition fluctuations—perhaps spinodal decomposition. MOSS10 detected a satellite with X-rays in quenched specimens, but the wavelength was much smallerthan that found by Fukano. Recently Woodilla and Averbac11 detected a modulation similar to that found by Moss, but with electron diffraction. They suggested that the microstructures observed by Fisher and Embury may be simply due to a Moir'e effect from a back deposit of gold during electropolishing. Satellites in electron diffraction can be enhanced by small bends or tilts of a foil. Even if this were not the case, asymmetric satellites, as are found in Au-Ni are difficult to analyze. They are principally influenced by the large differences in atomic size in this system, and obtaining information on the associated composition fluctuations depends on assumptions as to whether or not there are isolated clusters or composition "waves", and in the latter case whether or not the entire specimen has decomposed. Assuming that decomposition was complete in alloys aged at low temperatures, Woodilla and Averbach suggested that the satellites indicated a composition which varied about the average value by only± 5 pct, much less than that predicted from the phase limits, Fig. 1; they attributed this small variation to a cessation of the reaction caused by the loss of mobile vacancies. In Ref. 12 it was shown that vacancies are trapped by the preprecipitate in this system. Furthermore, after reversion, aging proceeds at a much slower rate than immediately after a quench from the solution temperature,5'11,13 indicating the importance of vacancies in the kinetics. In some senses Au-Ni is a favorable system for studying preprecipitation and reversion mechanisms because of the simplicity of the equilibrium diagram14 and the availability of good thermodynamic15 and diffusion data.16,l7 The calculations necessary to relate the decomposition process to classical nucleation and growth or to spinodal decomposition, as well as those necessary for understanding reversion mechanisms, can, therefore, be made. As an example, the position of the chemical spinodal, calculated from the thermo-dynamic data of Sellars and Maak15 with the method described by Rundman and Hilliard,18 is included in Fig. 1. It is compared to the chemical spinodal calculated from the equilibrium miscibility gap by Cook and Hilliard's method.l9 Included is the coherent spinodal suggested by Golding and MOSS," using measured elastic constants for their calculations and the same ther-modynamic data. The spinodal estimated by Woodilla and Averbach,11 by determining the temperature limits for the observation of the satellites discussed above, is also given. Actually this last curve is open to question on the following grounds: As a phase boundary is approached the spacing of particles or the wavelength

Jan 1, 1970

By A. J. Ardell, O. D. Sherby

The elevated-temperature steady-state creep behavior of polycrystalline a Zr was studied in vacuo under constant tensile stress. The experiments were conducted from 660° to 845°C over the stress range 870 to 5400 psi. The steady-state creep rate, is, of a Zr was found to be related to the applied stress, a, and temperature, T, by the equation where C is a constant. Two interpretations of the data are considered, each requiring different values of the parameters ß, Q, and . It is shown that this formula can be derived using Gilman's phenomeno logical equation relating dislocation velocity to the shear stress for the case where dislocation mobility is limited by the nucleation of kink motion. From the fundamental dependence of strain rate on dislocation velocity and the density of mobile dislocations, the factor exp(ßo) can be related to the stress dependence of the dislocation density, and Q to the energy of kink formation. The two interpretations considered are as follows: Case I—the parameters are adjusted so that the equation above describes the data over the entire range of stress and temperature, implying that creep is controlled exclusively by dislocation glide; Case IIthe parameters are adjusted to allow for the possibility that creep is diffusion-controlled at high stresses but glide-controlled at low stresses. For both cases the values of the parameters are within the spread permitted by the data. Factors favorable to the observation of the unusual stress -dependence of i, for a Zr are discussed and it is shown that the anomalously rapid rates of self-diffusion and the low modulus on the principal glide system are primarily responsible for the anomalous creep behavior. Metal-lograflhic observations are presented and discussed with reference to the overall creep behavior. PREVIOUS investigations of the creep behavior of a Zr have been conducted from room temperature to about 500C The results of these investigations are primarily of engineering significance because of the variety of conditions under which they were carried out, and the low testing temperatures employed. Little information is available concerning the fundamental nature of the elevated-temperature creep behavior of a Zr. This investigation of the creep behavior of a Zr above half the melting temperature (the melting temperature, T,, of a Zr is estimated to be about 1860"K was initiated because there are two significant anomalies of the physical properties of a Zr: a) the rate of self-diffusion in polycrystalline a Zr is anomalously rapid: b) both Young's modulus for polycrystalline a Zr and the shear modylus as-sociated with the prismatic slip system { 1010)( 1120) (the principal slip system in a Zr 7) have anomalously large temperature variations The importance of these quantities to the understanding of creep behavior has been demonstrated, and it was hoped that this investigation would provide a critical test of the predictions based upon the analysis of creep data for other pure metals. However, as the investigation progressed it became increasingly clear that the unusual creep behavior of a Zr could not be analyzed in a conventional manner. 1) EXPERIMENTAL EQUIPMENT AND PROCEDURES Two grades of zirconium, designated A and B, were used in this investigation. The analyses, as supplied by the vendors, are presented in Table I. Material A, with a nominal purity of 99.84 pct, was supplied by the Wah Chang Corp. in the form of cold-rolled sheet from which specimens were machined without further fabrication. Material B, with a nominal purity of 99.95 pct, was purchased from the Nuclear Materials and Equipment Corp. in the form of iodide crystal bar.

Jan 1, 1968

By H. Papazian, Robert A. Lad

SINCE it can be shown both theoretically'' : and experimentally that disordered Cu3Au has a greater internal energy than the ordered alloy, it might be expected that its chemical reactivity would be greater. Experiments were initiated' to test this hypothesis. This note will describe some recent results obtained in the study of the rates of corrosion of ordered and disordered polycrystalline Cu3Au by aqueous FeCl., The work of Robertson' has indicated that the attack by such solutions is intergran-ular and highly orientation dependent. It was of interest to see whether the order-disorder effect was great enough to overcome the variations from specimen to specimen which would arise from this type of attack. The polycrystalline alloy specimens (0.020 in. wire, 1¾ in. long) were preannealed in ½ atm H, at 700°C for 2 hr. The subsequent heat treatments to produce order and disorder were performed in M atm He. Disordered specimens were obtained by heating at 450°C for 2 hr and quenching in water. Ordered specimens were produced by heating for 2 hr at 450°, furnace cooling to 380°, and soaking at 380°C for 80 to 90 hr. Room temperature was reached by cooling to 150" at a rate never exceeding 7" per hr, and water quenching from 150°C. The specimens so produced exhibited electrical resistivities characteristic of the two states; of the order of 11.27x10 ohm-cm for the disordered, and 4.28x10 ohm-cm for the ordered specimens. The corrosion rates were studied at 30°C in ferric chloride solutions of several concentrations (?,¼, 1.5, and 1 molar). The specimens were attached to a stirrer during the tests. Tests were interrupted at intervals, and the specimen weight change was determined after washing and drying. Analysis of the solutions showed that the attack was limited to the copper. The higher corrosion rate for the disordered alloy, as illustrated in Fig. l, is typical of the results obtained. The effect of variation in ferric chloride concentrations is indicated in Table I for typical specimens. The variations in the ratio of reactivity for disordered and ordered specimens at the different concentrations are within experimental error, and the best value is estimated to be near 1.4. The experimental variation was due in part to the mode of attack, as mentioned, and in part to the fact that films were formed on the specimens under some conditions. Electron diffraction patterns of the films were very complex and indicated the presence of metallic iron, as did chemical analysis. In an attempt to eliminate the difficulties arising from the mechanism of attack, single crystal specimens were used. A greater reaction rate was again measured for the disordered specimens, but the reproduci-bility was no better than that for the polycrystalline metal. This would indicate that film formation was primarily responsible. It was concluded that the reaction with aqueous ferric chloride served to demonstrate that the reactivity of disordered Cu3Au is greater than that of the ordered alloy. However, the difficulties with film formation did not permit determinations of a precision great enough to warrant extension of the study to partially ordered alloy systems.

Jan 1, 1957

By P. G. Shewmon, F. N. Rhines

THE determination of the self-diffusion coefficient of magnesium has been made possible recently by discovery1-1 of a radioactive isotope, Mg28 having a half-life of 21.3 hr,1 and subject to manufacture in useful quantity. In the present research this material was condensed from the vapor phase upon a surface of high purity magnesium. The progress of diffusion of the tracer atoms into polycrystalline magnesium was followed by machining layers and measuring the change in the intensity of radiation as a function of the distance of each layer from the surface. The self-diffusion coefficient was found to be 2.1 X 10-8 sq cm per sec at 627°C, 3.6 X 10-9 sq cm per sec at 551°C, and 4.4 X 10-10sq cm per sec at 468°C; the activation energy is about 32,000 cal per mol. Experimental Procedure Since there was no other published measurement of a diffusion velocity in any magnesium-base material, is was necessary to employ a number of new experimental techniques. The short half-life of Mg28 made it necessary to complete the entire experimental procedure within three or four days. This meant that the work had to be done where a cyclotron was readily accessible and that all operations, prior to the diffusion heat treatment, had to be so designed as to minimize their time requirements. Unusual problems were imposed also by the chemical reactivity of magnesium, its high vapor pressure, and the fact that no satisfactory method for elec-trodepositing magnesium on magnesium is presently available. Finally, the machining and handling of the easily air-borne radioactive-magnesium chips involved certain health hazards, resulting in the need for further experimental restrictions. Preparation of Mg28 The Mg28 was produced in the Carnegie Institute of Technology syncrocyclotron by the neutron spallation of chlorine.5 his involved bombarding a 2 gram crystal of high purity NaCl with a beam of 350 mev protons for a period of 2 hr, after which the crystal was dissolved in warm water and the Mg28 was concentrated and purified by chemical means (see Appendix). About 50 microcuries of Mg28 thus were obtained in the form of magnesium oxinate (8 hydroxyquin-olatc?), which was ignited in air to produce MgO. This in turn was reduced to magnesium metal vapor, by the method of Russell, Taylor, and Cooper," in the vacuum apparatus shown schematically in Fig. 1. Here the essential part is a tantalum ribbon, slightly dished to receive the MgO. The ribbon, pre- viously outgassed at high temperature, is heated to about 1700°C by passing an electric current through it, whereupon tantalum oxide is formed, magnesium vapor is released almost instantaneously, and condensed partly upon the diffusion sample. Diffusion-Sample Preparation: Hot-extruded magnesium rod, 21/32 in. round was used in making the diffusion specimens. The magnesium analyzed as follows: 0.004 pct Al, 0.027 pct Fe, 0.040 pct Mn, 0.0004 pct Cu, 0.0002 pct Ni, and less than 0.01 pct Ca, 0.0004 pct Pb, 0.0011 pct Si, 0.001 pct Sn, and 0.001 pct Zn. A brief study of the crystal texture of this material revealed a sharp fiber texture with the (001) plane roughly parallel to the extrusion axis. Cylindrical samples 1/2 in. long by 5/8 in. were machined from this rod, the end faces dressed on 3/0 emery, and lightly etched with 20 pct HC1 in water. These samples then were annealed for at least twice the intended time of diffusion, at the intended diffusion temperature, in order to stabilize the grain structure at about 1 mm average diameter. The annealing treatments were conducted in argon in the same apparatus and in the same manner as the subsequent diffusion treatments, which will be described presently. Thus, a strain-free plane surface was produced, but there remained a layer of MgO which had largely to be removed before the layer of Mg28 was deposited. Most of this layer was taken off by two light passes over 3/0 emery paper. The balance of the oxide and a thin layer of metal were then removed by etching 5 to 10 min in 4 pct nital (4 pct HNO3 and 96 pct ethyl alcohol) made with absolute alcohol. There followed immediately three quick rinses in: 1-49 1/2 pct methanol, 49 1/2 pct acetone, and 1 pct formic acid, 2-50 pct methanol and 50 pct acetone, and 3-pure benzene. This procedure is essentially that of Sturkey.7 The resulting surface, which was of almost elec-tropolished brightness, remained plane and was free of cold work. It could be kept clean by storing under benzene, or in a desiccator; short exposure

Jan 1, 1955

By B. F. Decker, D. Harker

The recrystallization reaction in OFHC and spectroscopically pure copper has been followed by X ray diffraction determinations of the amount of material with the cold-worked and recrystallized textures in specimens which had been given various heat treatments. The heats of activation for recrystallization are found to be: 29.9 kcal per rnol for OFHC copper and 22.4 kcal per rnol for spectroscopic copper. A reliable and speedy method for measuring the amount of recrystallized material in a piece of rolled metal, together with a new scheme for low temperature heat treatment, have made possible a determination of the activation energy for recrystallization in rolled copper. This method for studying recrystallization rates is different from others reported in the literature.' Oxygen free, high conductivity copper—OFHC copper—of purity 99.98 pct gave an activation energy for recrystallization of 29.9 kcal per mol with recrystallization data taken in the temperature range 208-245°C. Copper of purity 99.999 pct from the American Smelting & Refining Co. gave an activation energy for recrystallization of 22.4 kcal per mol with recrystallization data taken in the temperature range 43-135°C. The great change in activation energy for recrystallization as the small amount of impurity in the metal is decreased suggests that the motion of grain boundaries is conditioned by the impurities which must be concentrated in them. Thus the OFHC copper recrystallizes with an activation energy of the order of magnitude to be expected for the diffusion of the kind of impurities likely to be present. The high purity copper, on the other hand, recrystallizes with a much lower activation energy, in accord with the notion that the grain boundaries need not wait for the impurities to diffuse with them. Copper when strongly cold-rolled has a well-defined crystallite texture which can be described by saying that the [111] direction of the face- centered cubic crystals is aligned approximately in the cross-rolling direction and that faces of the form {110} lie approximately in the rolling plane. When such rolled copper is completely annealed the crystallite texture is quite different: the [loo] direction is in the cross-rolling direction and faces of the form (100) are in the rolling plane. At intermediate stages of annealing, the copper contains both textures. If a Geiger counter X ray diffraction spectrometer is set up so as to measure the intensity of the (200) Bragg reflection from the rolled surface of a piece of copper strip, this intensity increases from a very small value in the cold-worked state to a maximum as the specimen is annealed. The same would be true for (200) reflections from planes normal to the rolling or cross-rolling directions. In view of this fact, the amount of material in the copper which has recrystallized can be measured by comparing the intensities of one of these reflections just mentioned with the intensity from a similar specimen after complete annealing. In much the same way the intensity of (111) reflections from planes normal to the cross-rolling direction could be used to measure the amount of unrecrystallized material in the specimen. In the work to be described here the amounts of recrystallized and unrecrystallized material in partly annealed rolled copper specimens were determined in this way. Experimental Procedure Samples of oxygen free high conductivity (OFHC) copper, cold-rolled 99.7 pct to 0.002 in. thickness, were subjected to heat treatments for various chosen times at four different temperatures, and high purity copper samples, cold-rolled 98.0 pct to 0.0015 in. thickness, at six different temperatures. Each sample, after treatment, was placed in an orientation sample holder for use with a Geiger counter X ray spectrometer and a reading was taken of the intensity of the (200) reflection from planes perpendicular to

Jan 1, 1951

By Frank S. Gardner, Morris Cohen, Dara P. Anita

There is a conspicuous lack of information in the literature on the precise role played by residual quantities of austenite in heat-treated steels. While retained austenite may be expected to have significant effects on such important properties as the impact resistance of tools and carburized parts and the dimensional stability of dies and gauges, attempts to study these problems systematically have suffered from the absence of a reliable technique for ascertaining absolute amounts of austenite. With such a technique, not only could the relation between residual austenite and service behavior be investigated, but the chemical and heat-treatment factors that control the quantity of the austenite could also be established. Thus, the retained austenite content could then be regulated in accordance with the results desired. The presence of retained austenite may be inferred by observing the changes in properties (hardness, specific volume, electrical resistance, etc.) due to its decomposition on subatmospheric cooling1 or on tempering,2 but these indirect measurements are not readily translated into percentages of austenite. Quantitative determinations have been achieved in special cases by means of dilatometric3 and magnetic4,5 measurements; but the former procedure requires that observations be conducted during the hardening and involves difficult calibrations that limit its applicability, while the magnetic procedure is very sensitive to extraneous influences and depends upon assumptions that can only be considered as approximate in an absolute sense. Furthermore, both the dilatometric and magnetic effects are volume properties, and cannot be conveniently utilized to explore variations in retained austenite from point to point in a suecimen. Some microstructures lend themselves to accurate evaluation of retained austenite contents by metallo-graphic methods,3 but such instances are not too common, and at best, the procedure is tedious. The characteristic X-ray diffraction lines of austenite permit positive identification of its presence; and in the present paper the intensities of these lines, or the corresponding densities* on photographic film, are employed to measure the amount of the retained austenite. In order to cancel the effect of uncontrollable variations in the X - ray exposure and film - development conditions, a standard reference line is Superimposed on each film by maintaining a foil of aluminum in the path of the incident X-ray beam during the entire exposure. With different samples of steel and the same foil, the ratio of the density of some one austenite line to the density of

Jan 1, 1943

By Morris Cohen, Dara P. Anita, Frank S. Gardner

There is a conspicuous lack of information in the literature on the precise role played by residual quantities of austenite in heat-treated steels. While retained austenite may be expected to have significant effects on such important properties as the impact resistance of tools and carburized parts and the dimensional stability of dies and gauges, attempts to study these problems systematically have suffered from the absence of a reliable technique for ascertaining absolute amounts of austenite. With such a technique, not only could the relation between residual austenite and service behavior be investigated, but the chemical and heat-treatment factors that control the quantity of the austenite could also be established. Thus, the retained austenite content could then be regulated in accordance with the results desired. The presence of retained austenite may be inferred by observing the changes in properties (hardness, specific volume, electrical resistance, etc.) due to its decomposition on subatmospheric cooling1 or on tempering,2 but these indirect measurements are not readily translated into percentages of austenite. Quantitative determinations have been achieved in special cases by means of dilatometric3 and magnetic4,5 measurements; but the former procedure requires that observations be conducted during the hardening and involves difficult calibrations that limit its applicability, while the magnetic procedure is very sensitive to extraneous influences and depends upon assumptions that can only be considered as approximate in an absolute sense. Furthermore, both the dilatometric and magnetic effects are volume properties, and cannot be conveniently utilized to explore variations in retained austenite from point to point in a suecimen. Some microstructures lend themselves to accurate evaluation of retained austenite contents by metallo-graphic methods,3 but such instances are not too common, and at best, the procedure is tedious. The characteristic X-ray diffraction lines of austenite permit positive identification of its presence; and in the present paper the intensities of these lines, or the corresponding densities* on photographic film, are employed to measure the amount of the retained austenite. In order to cancel the effect of uncontrollable variations in the X - ray exposure and film - development conditions, a standard reference line is Superimposed on each film by maintaining a foil of aluminum in the path of the incident X-ray beam during the entire exposure. With different samples of steel and the same foil, the ratio of the density of some one austenite line to the density of

Jan 1, 1943

By R. L. Frink

I FEAR that my treatment of this subject may not, in all instances, meet the approval of those who read my opinion as to the utility and efficiency of pyrometers in the making of glass, or bottle-glass. It may be superfluous for me to add that this opening statement is based on over 15 years' experience in an endeavor to successfully apply pyrometers, or heat-measuring instruments, to glass-melting furnaces, particularly the type known as tank-furnaces, and that such endeavors have proved more or less futile. It is my desire, therefore, to herein set forth the problems encountered, hoping thereby to stimulate further effort in the successful application of pyrometers. It is my desire, also, to invite criticism of the methods used and suggestions from those who may have a wider experience, or who may have better methods. Fin. 1.-GLASS TANK FURNACE. In general, I believe that the inefficient application of pyrometry to the melting of glass in tank furnaces is not due as much to the instruments as to the conditions under which they must be used and the character of the medium in which they must reside. Most of those present no doubt, are familiar, at least in a general way, with the methods and apparatus used in the melting and making of bottle-glass, or glass that is formed into containers of various shapes and designs. Most of this glass is made in what is known as a tank furnace, these furnaces vary widely in design and construction. In general, they consist of a rectangular tank-like form, built up of fireclay blocks, and range from 2 by 4 ft. (0.6 by 1.2 m.) to 24 by 140 ft. (7.3 by 42.6 m.) with a depth of glass from 2 to 5 ft.; or a capacity of from approximately. 2 to 700 tons. Fig. 1 shows the general

Jan 9, 1919

By J. L. Walter

Extra-low carbon iron sheet, when deformed and annealed, undergoes exaggerated or abnormal grain growth in the critically deformed regions of the sheet. This exaggerated pmth occurs, for low strains (3 to 6 pct), only in sheet which has a fine dispersion of precipitates in the subsurface region of the sheet and fewer and coarser precipitates in the sheet interior. These particles have been identified as manganese sulfides. Wing the anneal, grains near the surface are gvowth-inhibited by the fine particles but the grains in the interior are free to grow normally. With the additional driving force provided by the strain energy, the interior grains first grow into the small subsurface pains. Eventually, these growing grains grow completely through the sheet. Calculations of limiting grain sizes at various values of strain indicate that a volume fraction of precipitates in excess of 10-' would be required to eliminate exaggerated growth in material strained to 10 pct. OPEN-COIL annealing has made decarburization of sheet steel both efficient and economically practical. Use of such material for porcelain-enameling stock is one of many possible applications of low-carbon iron since carbon in steel is deleterious to porcelain-enameling properties.' However, the extra-low carbon iron presently available presents another problem, that of exaggerated grain growth in regions where the sheet has been deformed as by bending or stretching. In exaggerated grain growth a few grains start to absorb their neighbors and these may become very much larger than the average grains of the sheet. Other names for exaggerated grain growth are coarsening, critical grain growth, secondary recrystallization, abnormal grain growth, or discontinuous grain growth. While this grain growth does not affect the enameling qualities of the low-carbon iron sheet, their presence results in a marked reduction of tensile-yield strength in the region of the sheet containing the large grains. The loss of tensile yield strength may render the material unsuitable for many applications . This report describes the results of a study undertaken to determine the cause of the exaggerated grab growth in extra-1ow carbon iron and, if possi- ble, to prescribe practical procedures for its prevention. GENERAL THEORY The driving force for exaggerated grain growth is the grain boundary free energy. This driving force is proportional to (l/rl + l/r2) where rl and r2 are the mutually perpendicular radii of curvature of the boundary between the growing grain and the grain being consumed. Thus, the greater the difference in size between the growing grain and the matrix grains, the higher the driving force for grain growth. If, however, the matrix grains are free to grow simultaneously, the driving force for exaggerated growth will be diminished. Stability toward growth of the matrix grains may be caused by a) a strong single orientation texture (texture inhibition),' b) a dispersed second phase,3"5 c) the thickness effect,' or d) intergranular segregation.7"9 As exaggerated growth occurs when growth of the matrix grains has been slowed by the stabilizing processes mentioned above, there must be an additional factor acting to promote growth of a few of the grains to the point where they are enough larger than the matrix grains that boundary energy driving forces are sufficient to cause continued growth. For instance, at the annealing temperature, growth-inhibiting inclusions may slowly dissolve and coalesce. Eventually, a grain boundary becomes unlocked and migration occurs at the expense of neighboring grains. Or, an additional driving force may be supplied to some of the grains if the material is strained. Then, since some grains will be strained less than others, the difference in strain energy between adjacent grains may be sufficient to overcome boundary locking and allow growth of the grains with lower strain energies. In the present study, therefore, such factors as the presence or absence of dispersed phases, the nature of the exaggerated growth, and the effect of strain have been considered. EXPERIMENTAL PROCEDURE I) Material and Processing. Three types of low-carbon iron were used in this study; types A and B were commercial grades, each from a different supplier.* The precise details of the processing of

Jan 1, 1963

By J. R. Sturgul, J. C. Wynn

Computers and computer applications in geophysics are fairly recent innovations. The area of data handling is one that has found many immediate applications. Initially, the use of computers involved in almost all cases, a batch-input arrangement requiring a trip to the local computer center, a wait and another trip. Even though the computer worked at incredible speeds, the time-delays, in case an error in the input data developed, often could be several days. Use of computers in mining exploration by medium-sized or small-sized companies and contractors has been almost non-existent beyond simple data reduction work. The principle J. R. Sturgul, SME member, and J. C. Wynn, both affiliated with the Dept. of Geosciences, University of Arizona, Tucson, presented this paper at the Eleventh International Symposium on Computer Applications in the Minerals Industry, reasons appear to be the time factor and the inconvenience of the frequent trips involved. With the advent of time sharing systems, whereby a company need only have access to a terminal, the main objections can be overcome. An Apparent Solution An apparent solution to the time and distance inconveniences which would permit more extensive use of computers in geophysical exploration, is to buy or lease/purchase a medium-sized computer and install it in or near one's office. This is a formidable proposition in terms of cost and programming time on both respects, but several mining companies of larger size have already gone this route. In most cases, the computers have been expensive devices with a limited memory or core size. For routine data handling this memory size is not restrictive; however, for many geophysical applications such as three-dimensional modeling this can be critical.

Jan 10, 1973

By Horace Pops, L. L. Isaacs

SOLID-STATE transformations at low temperatures are frequently studied by optical metallography. Greninger and ooradian' and Hull and Gar-wood2 have examined the martensitic-type transformations in 0-phase Cu-Zn alloys during continuous cooling below room temperature. The lowest range of temperatures in the apparatus described by these investigators is limited to approximately 110°K. Present studies of reversible first-order phase transitions near liquid-helium temperatures have necessitated the design of a new cold-stage cryo-stat for optical metallography. The features of interest of this instrument are as follows: a) temperatures down to 4.Z°K are obtainable; b) the specimen is always under vacuum, thereby avoiding the condensation of moisture; c) specimens can be rapidly cooled or heated at varying rates, thus permitting studies of transformation hysteresis. Fig. 1 shows the details of the cryostat. All major parts are made of type 304 stainless steel. Both helium and nitrogen cans are Heliarc-welded to the helium filling tube, permitting the use of a common shielding can. Hence, the over-all weight and the area of the evacuated surface are mini- mized. The copper sample holder is screwed into the bottom of the helium can. To assure good thermal contact between the sample holder and the refrigerant, an indium ribbon is placed on the threads

Jan 1, 1964

By William E. Ford, Edward Salisbury Dana

1. Normal Class (6) Zircon Type 2. Hemimorphic Class (7) Iodusuccinimide Type 3. Pyramidal Class (8) Scheelite Type 4. Pyramidal- Hemimorphic Class (9) Wulfenite Type 5. Sphenoidal Class (10) Chalcopyrite Type 6. Trapezohedral Class (11) Nickel Sulphate Type 7. Tetartohedral Class (12) Mathematical Relations to the Tetragonal System THE TETRAGONAL SYSTEM includes all the forms which are referred to three axes at right angles to each other of which the two horizontal axes are equal to each other in length and interchangeable and the third, the vertical axis, is either shorter or longer. The horizontal axes are designated by the letter a; the vertical axis by c (see Fig. 166). The length of the vertical axis expresses properly the axial ratio of a : c, a being uniformly taken as equal to unity. The axes are orientated and their opposite ends designated by plus and minus signs exactly as in the case of the Isometric System. Seven classes are embraced in this system. Of these the normal class is common and important among minerals; two others have several representatives, and another a single one only. It may be noted that in four of the classes the vertical axis is an axis of tetragonal symmetry; in the remaining three it is an axis of binary symmetry only.

Jan 1, 1922

By L. S. Castleman, R. Resnick

The self-diffusion coefficient of columbium has been measured over the temperature range 1535o to 2120°C, using the radioactive isotope Nb95 as the tracer. The data for the temperature dependence of the self-diffusion coefficient are best fitted by the equation D = (12.4 ± 0.8) exp [(-105,000 ± 3000)/RT] cm2/sec. THE increasing need for high-temperature structural materials has stimulated considerable interest in the properties of columbium, since it is a strong, ductile metal with a high melting point—in the vicin- ity of 2450oc.l One fundamental property which has not heretofore been measured is the self-diffusion coefficient. The self-diffusion coefficient plays an important role in theories of oxidatioq2 creep,' sintering4 and other phenomena. Accordingly, the investigation reported in this paper was undertaken to determine the value of the self-diffusion coefficient and its temperature dependence by the use of Mb95.

Jan 1, 1961

By R. M. Bird

FEW persons outside of the foundry trade have any conception of the great variety of sands now regularly specified and furnished, nor of the differences in foundry practice frequently resulting from apparently minor changes in sands. Conditions in each locality where the sands are mined and prepared necessitate special equipment designed to perform the work efficiently. The description that follows of the foundry sand producing operations is typical of the areas in New York and New Jersey, and particularly refers to the operations of the George F. Pettinos Co.

Jan 1, 1926

By A. W. Schlechten, W. J. Kroll

As this paper is written, the only method for the commercial production of lithium metal is by the fusion electrolysis of LiCl-KCl mixtures, as first proposed by Guntz.2 The details of the industrial process have not been made public but Osborg8 stated that an efficiency of more than go pct is obtained with a lithium metal recovery above 95 pct and that the metal is 99.5 pct pure. Pletenev9 gave a power consumption of 66 kw-hr per kilogram of lithium metal produced with a salt consumption of 9 kg of LiCl and 0.4 kg of KCI. Pletenev's figures indicate a somewhat lower recovery than that given by Osborg. In a recent report on German lithium plants, Motock7 gives figures that are also less favorable than those reported by Osborg. At the German plants the lithium recovery in electrolysis was only 83.4 pct, with a power consumption of too kw-hr per kilogram of lithium. The metal averaged only 97 pct pure, with sodium and potassium as the chief impurities. The presence of SO2, SO3, Si02, Ba, Ca, and appreciable amounts of Na and Fe2O3 were especially disturbing in the electrolysis, and it was necessary to run the cell for some time to eliminate most of these impurities and condition the bath. It may be noted here that most of these impurities do not interfere with a thermal reduction process. These facts show that the present method of producing lithium is fairly efficient and would suggest that other methods could not compete. However, fusion electrolysis has certain drawbacks, such as the necessity of supplying costly, low-voltage direct current and of collecting and processing the anodic chlorine if it is not to be wasted. An even greater disadvantage is that lithium chloride must be used for an electrolyte. It is one of the most expensive salts of lithium, since it cannot be produced in anhydrous form by precipitation from aqueous solution, as is done with LiF or Li2CO3. The reduction of lithium by a vacuum process similar to the ferrosilicon method for the production of magnesium would have certain advantages. Small units, such as Pidgeon retorts, would give flexibility of production to meet demand, and a small plant would require moderate capital outlay. There is a possibility of using idle Government-owned magnesium plants for such purposes. The raw materials needed for vacuum thermal reduction would not have to be extremely pure, especially with respect to the iron and calcium content. This paper describes a series of experiments to determine the effectiveness of various agents for the reduction of lithium compounds. The investigation is part of a program of vacuum metallurgy con-

Jan 1, 1947

By L. R. Gremillion, G. H. McClellan

Chemical and mineralogical data on 25 apatitic phosphate ores of diverse origins from commerical deposits around the world were widely variable. From 14 to 18 chemical constituents of each sample were determined, and each sample contained four to ten mineral phases. Altogether there were 34 different mineral phases, many of which occurred in only one sample. Samples contained varieties of francolites, fluorapatites, and mixed hydroxyfluorapatites. The apatite compositions calculated from chemical data showed that certain types of substitutions were characteristic of different apatite varieties, but every variety had some degree of substitution of (CO3)2- for (PO4)3-, and the P205 contents of the apatites ranged from 35.1 %for a North Carolina apatite to 41.9% for a Missouri apatite. In a concentrate with a grade of 32% P2O5, the North Carolina sample should contain about 8.7% gangue and the Missouri sample about 23.7% gangue. To optimize the digestion process of a phosphate ore, it is necessary to consider the dissolution behaviors of different mineral phases during acidulation and their possible effects on chemical properties of a fertilizer intermediate. In the production of wet-process acids or superphosphates, acidulation of certain minerals may cause foaming and formation of sludge, filtration, corrosion, and handling problems.

Jan 1, 1982

By E. W. Dewing

The solubility of A14C3 in cryolite melts has been determined by holding the melt with aluminum in a graphite crucible for 4 hr at temperature, quenching the crucible in water, and analyzing the frozen salt by CH4 evolution. The solubility is given by log(wt % Al4C3) = 1.298 -1.l74 (wt ratio NaF/AlF3) independent of temperature in the range 985" to 1080°C. Additions ofA1203, .NaCl, CaF2, or LiF up to 10 wt pct do not affect the solubility, provided that LiF and CaF2 are replaced by equivalent NaFin calculating the ratio. The variation with ratio can be accounted for quantitatively on the hypothesis that the Mechanism of solution is Al4C3 + 2AIF3 = 3Al2CF2 since the solubility varies as (aAIF3) 2/3; Mechanisms involving C", C2- , AlC- , A12C2+, and Al2CF+are re-jected. At ratios above 1.5 truces of c;- are also present. It was noticed some years ago in this laboratory that carbon specimens made cathodic in a cryolite bath decrease in diameter over a period of time.' It was also found that cryolite which had been melted with aluminum in a graphite crucible was sometimes yellow, and that A14C3 could be seen intimately embedded in the cryolite crystals. CH, could be recovered by acid attack. There was thus reason to believe that A14C3 had a genuine solubility in molten cryolite, and the object of this work was to measure it. EXPERIMENTAL AND RESULTS A graphite crucible, 2 in. ID and 2.5 in. deep with a close-fitting lid, contained about 150 g of cryolite (plus additives) and about 20 g of 99.99 pct pure Al. It was placed in a vertical Inconel furnace tube, covered with refractory brick to provide thermal insulation, and the top of the tube was roughly sealed with disks of asbestos paper. An alumina tube led argon down to the level of the top of the crucible to expel air, and an alumina sheath with a Pt-10 pct Rh ther-mocouple inside was positioned between the crucible and the wall of the Inconel tube. The tube was heated in an electrical resistance furnace for the desired length of time; the refractory bricks were then removed, and the crucible was lifted out with tongs and quenched in water. After such a crucible was broken open it could be seen that its inner surface was yellow owing to the presence of A1,C3, and sometimes needle-like crystals could be observed (at X75 magnification) growing out into the nielt. It was concluded that Al,C, had been present as a separate phase when the crucible was at tenlperature and that the melt had been saturated: in

Jan 1, 1970

In the death of Major A. B. desaulles at South Bethlehem, Pa., on Dec. 24, 1917, the Institute lost a valued and esteemed member, one of the last few of those who, in May, 1871, at Wilkes-Barre, attended the first meeting of the Institute, in response to a call that had been issued by R. P. Rothwell, Eckley B. Coxe, and Martin Coryell. Major desaulles was born Jan. 8, 1840, at New Orleans, and was the son of Louis and Armidc Longer desaulles, each of French descent. As a boy he was privately tutored in the South, and then schooled in New England to enter the Rensselaer Polytechnic Institute, where he was graduated with honors in the class of '59. Following his graduation he studied in France and Germany for two years, returning to New Orleans in 1861 to enter the Confederate Army, in which he served with distinction as an engineer throughout the Civil War, rising to the rank of Major and the command of his Corps. After the close of the war he pursued his studies abroad for another year, returning to New York in April, 1866, when he became engineer for the New York & Schuylkill Coal Company, and remained at its plant near Wilkes-Barre until October, 1871, when it was sold to the Philadelphia & Reading Coal and Iron Company. He then returned to New York where for several years he was engaged in the practice of his profession. In 1869, he married Catharine M. Heckscher, daughter of Charles Heckscher of New York City. Mrs. desaulles survives him. In 1876, Major desaulles became superintendent of the Dunbar Furnace Company in Fayette County, Pa., with which company he remained until 1883. As a feature of the Pittsburgh meeting in May, 1879, Major and Mrs. desaulles entertained the members of the Institute at dinner at Dunbar on the occasion of the Institute's visit to Dun-bar furnace. In 1883, he took up the formation of the Oliphant Furnace Company in that county, of which he became and remained the head for five years. He was then called to the superintendency of the New Jersey Zinc Company at South Bethlehem, Pa., in which position he remained until his retirement,from active service in 1911, maintaining, however, hi.; keen interest in the metallurgy and smelting of zinc, in collaboration with his son Charles, a member of the Institute and a Director of the American Smelting and Refining Company. From 1888, until his death, he resided in South Bethlehem, a prominent and highly respected citizen of that progressive and stirring community. Major desaulles was a member of the Protestant Espicopal Church, and at the time of his death was a vestryman of the Church of the Nativity at South Bethlehem. He was president of the Men's Club of the Church of the Nativity. His kindly nature led him to take an active interest in the promotion of healthful sports among young men. He was ever active in all things pertaining to the general welfare and uplift of his fellowmen. In his death the profession has lost a distinguished engineer and mctallurgist and our country a patriotic able citi-zen and man of affairs—one greatly beloved by those whom he honored with his friendship, and respected and looked up to by all who were privileged to know him.

Jan 1, 1920