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By M. Cohen, R. T. Howard

The purpose of this paper is to direct attention to the lower part of the austenite transformation diagram, or TTT curves, where considerable uncertainty still exists as to the blending of the bainite and martensite reactions. The dissimilarities between these two low temperature transforrnations have been summarized recently by Troiano and Greninger.' Of particular importance to the present work is the fact that martensite formation, unlike bainite formation, is not suppressed by the fastest quenching rates attainable. However, most of the investigations in this direction have been concerned primarily with the start of the martensite transformation (Mg) during cooling, rather than with the quantitative aspects of the subsequent reaction on either further cooling or isothermal holding. Several years ago, Cohen2 suggested a form of transformation diagram which reconciled, at least schematically, the isothermal characteristics of the bainite transformation with the insuppressible nature of the martensite transformation. In essence, the diagram consisted of a family of horizontal lines (representing the progress of the austenite-martensite reaction during quenching) which were interrupted in point of time by the bainite family of C-curves extending down from higher temperatures. This picture received later support by metal-lographic3 and dilatometric4,5 studies, but the techniques employed were not sufficiently critical to substantiate some of the details. Elmendorf6 and Grange and Stewardi have shown that the Greninger-Troiano metallographic technique8 may be used not only for M, determinations, but for studying the course of further austenite decomposition. In fact, the latter investigators have observed the amount of martensite formed as a function of temperature below M, in fourteen commercial steels. In the present paper, the Greninger-Troiano technique has been combined with the lineal method of Howard and Cohen9 to provide a unique, quantitative treatment of the austenite-martensite reaction. Furthermore, isothermal transformations have been similarly measured on holding at temperatures above and below M., thus permitting a fairly detailed presentation of the lower part of the conventional austenite transformation diagram. Experimental Details Materials The steels, listed in Table I, were selected for study because they exhibit a satisfactory range of M. temperatures and have not received as much attention as the lower carbon grades. They were prepared as 30 lb induction furnace heats in the laboratories of the Vanadium-Alloys

Jan 1, 1949

By K. A. Gschneidner, O. D. McMasters

DgIel-ential thermal, nretallographic, and X-ray paramzetric methods were used to establish the Yb-Pb phase dingram. The terminal solid solubilities in the system are less than 0.2 at. pct. Lead additions lower the 792°C transformation tewzperature of pure ytterbiun~ at 787°C by means of an inverted peritectic veaction. Ytterbium additions raise the 327oC melting point of lead to 329°C by means of a perilectic reaction. Three intermetallic compounds Yb2Pb, YbPb, and YbPb, melt congruerttly at 1246", 1116o, and 740°C. respectively. YbsPb3 forms by a peritectic reacliorz at 1250°C. YbzPb crystallizes in the PbCh (C23) type structure and the orthorhombic latliceo constants are a, = 7.478. bo = 5.225, aud C0 = 9.549A. YbsPbs crystallizes in the hfnsSi3 (D8a) type structure. The hexa~nal lattice constants for this phase are a, = 9.325 and co = 6.92d. YbPbl crystallizes in the cubic AuCU3 (Ll2) type structure with a: = 4.8628 + 0.0008A. YbPb is diwrorphic with the structure change occurring al 507°C upon healing. The crystal structures of the two forttzs were not determined. Eulectic reactions occur at 5.1 at. pct Pb and 769oC, 41.5 at. pct Pb and 1086oC, and G8.0 at. pct Ph and 717°C: RELIABLE information concerning the alloying behavior of most of the rare-earth metals is scarce. A systematic study of the rare earth-lead alloy systems now in progress is an effort to supply such information. Determination of the phase diagram of the Yb-Pb alloys is the first step of this study. Variations in the valence of normally divalent ytterbium may be studied by comparing the characteristics of the Yb-Pb alloys to those of the trivalent rare earth-lead and of the divalent alkaline-earth-lead systems. Of particular interest are the melting temperatures of the intermetal-lic compounds since these are known to be high relative to the melting points of the constituent metals. Crys-tallographic and magnetic susceptibility data for the rare earth-lead alloys are to be included in the comparative study. Results from this and subsequent investigations should provide a better understanding of alloy formation. EXPERIMENTAL PROCEDURE Materials. The major impurities in the ytterbium and lead used in this investigation are listed in Table I. The lead was obtained from Cominco Products, Inc., and was specified to be 99.99 pct pure. The ytterbium was prepared in this laboratory by the lanthanum reduction of the oxide followed by a distillation of the ytterbium. Alloy Preparation. The alloys were prepared by

Jan 1, 1968

By G. B. Craig, H. E. Rosinger, W. J. Bratina

NUMEROUS studies have been reported on how the mechanical properties of a bcc metal are affected by interstitial solute atoms in the temperature range where appreciable diffusion occurs in finite time. Only recently, though, has work on an "interstitial-free" iron been reported, this by Leslie and sober,' on an Fe-0.15 pct Ti alloy. The data was limited to a maximum temperature of 296°K. This work presents a systematic study of the flow stress and the activation volume of a similar material as a function of plastic deformation over the temperature range 253" to 473°K. The material studied was Ferrovac E iron containing 0.01 wt pct C and 0.0002 wt pct N annealed at 870°C for 1 hr, and the same iron with an addition of 0.15 wt pct Ti annealed at 870°C for 80 hr; cooling rates being 75°C per hr. In the latter alloy depletion of interstitials in the a-iron matrix associated with the second phase formation was expected. The temperature dependence of the flow stress was measured at 6 = 3.33 x 10-5 sec-1 from 253" to 473°K as a function of the deformation in 0.01 intervals from 0.03 to 0.12 strain. Results for two strains, 0.03 and 0.09 are shown in Fig. 1. The flow stress increased and reached a maximum in the Ferrovac E iron with increasing temperature, but showed a consistent gradual decrease in the Ferrovac E-0.15 pct Ti alloy. A yield point and Luders region, observed in Ferrovac E, was absent in the Ferrovac E-0.15 pct Ti alloy. Also, the Ferrovac E exhibited a serrated stress-strain behavior in the temperature range 330" to 400°K; dynamic strain-aging characteristics were absent for the titanium alloy. The activation volume was determined over the same temperature range by the method of stress relaxation at constant strain. Results are shown in Fig. 2. At the temperatures of serrated stress-strain behavior in Ferrovac E the activation volume was found to increase toward infinity and, as the temperature rose, to decrease, to about 100 to 200 v/b3. By contrast, the activation volume for the Ferrovac E-0.15 pct Ti alloy tended to increase continuously from the initial value of approximately 300 at 300°K to more than 800 at 473°K. Above 400°K, then, the corresponding activation volume values of Ferrovac E were less than that of the Ferrovac E-0.15 pct Ti alloy. The activation volumes, Fig. 2, appear to depend on dislocation-interstitial interaction in Ferrovac E and on dislocation-dislocation interaction in Ferrovac E-0.15 pct Ti alloy. The gradual increase in the latter metal is associated with thermally activated motion of dislocation segments.' The decrease of the activation volume in Ferrovac E above 400°K is probably related to the diminishing effect of interstitials on dislocations. Here, the interstitials can be moved along with a dislocation without appreciable drag. As indicated in Fig. 1, removal of interstitials

Jan 1, 1970

By O. L. Culberson, J. J. McKetta

INTRODUCTION The equilibrium constants for methane and for water, and for ethane and water have been calculated from experimental data for the two binary systems.2,3,11,12 These constants are for the two-phase systems for the temperature range of 100" to 340°F and the pressure range of 200 to 10,000 psia. The constants for ethane are greater than those for methane. Equilibrium constants for a natural gas from a natural gas-water system6 are about the same as those for methane. Equilibrium constants for water in all three systems are very nearly the same over the greater part of the range of temperatures and pressures studied. In ideal solutions, the vapor-liquid equilibrium constants for a component should be the same in all such solutions at a given temperature and pressure.5 Few solutions are ideal, and consequently the use of generalized constants often leads to error in situations where they do not actually apply. In decidedly non-ideal systems, it is desirable to have equilibrium constant data for the specific system. Vaporization equilibrium constants may be calculated either by thermodynamic methods or from experimental vapor-liquid equilibrium data. In the thermodynamic calculation. the equilibrium constant K may be shown to be" K = fP,/fl'..........(1) where fPi is the fugacity of the pure component i at the given temperature and at the vapor pressure of the pure component, and fp is the fugacity of pure i in the gaseous state at the temperature and total pressure of the system. At best the thermodynamic calculation results in a K applicable to ideal solutions. Equilibrium constants calculated from Equation (1) have been found to have only limited agreement with experimental values, even though the solutions for which they were calculated approached ideality."'" If the more volatile component is at a temperature above its critical, the accuracy of the equilibrium constant for the component is made more doubtful by the use of the vapor pressure term under such conditions. In order to make the thermodynamic calculation on a component above its critical temperature, one must resort to one of the empirical methods of extrapolation of the vapor pressure to conditions above its critical point.13 Application of the thermodynamic equilibrium constant then to the methane-water or the ethane-water systems would be risky since the systems are non-ideal, and the methane and ethane are above their critical in most petroleum operations. ETHANE-WATER SYSTEM In the temperature range of from 100° to 340°F and for pressures up to 10,000 psia, the vapor-liquid equilibrium data 'References given at end of paper. Manuscript received in the office of the Petroleum Branch March 16, 1951.

Jan 1, 1951

By J. J. McKetta, O. L. Culberson

INTRODUCTION The equilibrium constants for methane and for water, and for ethane and water have been calculated from experimental data for the two binary systems.2,3,11,12 These constants are for the two-phase systems for the temperature range of 100" to 340°F and the pressure range of 200 to 10,000 psia. The constants for ethane are greater than those for methane. Equilibrium constants for a natural gas from a natural gas-water system6 are about the same as those for methane. Equilibrium constants for water in all three systems are very nearly the same over the greater part of the range of temperatures and pressures studied. In ideal solutions, the vapor-liquid equilibrium constants for a component should be the same in all such solutions at a given temperature and pressure.5 Few solutions are ideal, and consequently the use of generalized constants often leads to error in situations where they do not actually apply. In decidedly non-ideal systems, it is desirable to have equilibrium constant data for the specific system. Vaporization equilibrium constants may be calculated either by thermodynamic methods or from experimental vapor-liquid equilibrium data. In the thermodynamic calculation. the equilibrium constant K may be shown to be" K = fP,/fl'..........(1) where fPi is the fugacity of the pure component i at the given temperature and at the vapor pressure of the pure component, and fp is the fugacity of pure i in the gaseous state at the temperature and total pressure of the system. At best the thermodynamic calculation results in a K applicable to ideal solutions. Equilibrium constants calculated from Equation (1) have been found to have only limited agreement with experimental values, even though the solutions for which they were calculated approached ideality."'" If the more volatile component is at a temperature above its critical, the accuracy of the equilibrium constant for the component is made more doubtful by the use of the vapor pressure term under such conditions. In order to make the thermodynamic calculation on a component above its critical temperature, one must resort to one of the empirical methods of extrapolation of the vapor pressure to conditions above its critical point.13 Application of the thermodynamic equilibrium constant then to the methane-water or the ethane-water systems would be risky since the systems are non-ideal, and the methane and ethane are above their critical in most petroleum operations. ETHANE-WATER SYSTEM In the temperature range of from 100° to 340°F and for pressures up to 10,000 psia, the vapor-liquid equilibrium data 'References given at end of paper. Manuscript received in the office of the Petroleum Branch March 16, 1951.

Jan 1, 1951

By R. E. Cech

Particles -3 to 30 µ diam of a 29.5 pet Ni, 70.5 pet Fe alloy after being melted and solidified while falling through a hydrogen atmosphere were found to contain a distortion-free body-centered-cubic phase. In most particles, the amount of body-centered-cubic phase approaches or equals 100 pet. It was shown by a series of experiments that this phase forms directly from the liquid. It is pointed out that a mechanism previously suggested in recent papers on the theory of the liquid-solid transformation can account for the solidification of body-centered-cubic phase in a temperature range where face-centered-cubic is the stable phase. IN the course of an operation to produce powders for a study of the martensitic transformation in Fe-Ni alloys, an unexpected phenomenon was observed. Powders of a 29.5 pet Ni alloy, which has a subzero M, temperature, were found to contain considerable amounts of a body-centered-cubic phase at room temperature. The experiments to be described were designed to elucidate the mode of formation of this unexpected phase. Experiments and Discussion An alloy containing 29.5 wt pet Ni, balance iron, was vacuum melted from high purity electrolytic iron and carbonyl nickel shot and solidified in the crucible. The metal was converted to oxide, ground to 100 pet <37µ, pressed to a cake, and homogenized 20 hr at 1350°C. To insure homogeneity the grinding, pressing, and homogenizing heat treatment was repeated a second time. The oxide was then ground to 100 pct <37µ and hydrogen reduced at 500°C. The reduced powder was given a spheroidizing treatment to produce spherical cast particles. The treatment consisted of dispersing the particles in hydrogen and allowing them to drop through a vertical hydrogen furnace set to a temperature higher than the liquidus of the alloy. This operation, commonly known as shotting, has been described more fully elsewhere.&apos; The powder produced by this method was found to contain approximately equal parts of body-centered-cubic and face-centered-cubic phases. A magnetic separation technique, previously described,&apos; was applied, and it was found that particles were either completely face-centered-cubic or nearly completely body-centered-cubic. The microstruc-

Jan 1, 1957

By Robert A. Rapp, Jan E. Verfurth

From the measurement of internal-oxidation band widths in Ag-In and Cu-A1 alloys, products for the solubility No and the diffusivity Do of oxygen in silver and apparent NODo products in copper have been obtained over a range of temperature. The NODO products for silver are in agreement with the product of the individually determined No and Do values of Eichenauer and Müller. Although interstitial diffusion of oxygen in silver is expected, the temperature-independent factor Do of the diffusion coefficient does not agree well with the theory of Wert and Zener. For Cu-A1 alloys internal oxidation does not follow the expression derived for ideal kinetics, since the apparent NODo product calculated from internal-oxidation measurements is a function of the alloy content in the specimen. Thus, unequivocal values for No or DO or even the product NODO in copper are not available from the literature. RELIABLE values for the saturation solubility NO and diffusivity DO of oxygen in fcc metals are lacking and difficult to determine, partially because internal-friction measurements are not applicable. Permeability experiments and Sieverts7-type adsorption -desorption methods3 have met with difficulty; however, Eichenauer and Müller4 have recently reported on the solubility and diffusivity of oxygen in silver as determined by the latter technique . Rhines, Johnson, and Anderson,5 and independently Meijering and Druyvesteyn,&apos; introduced the possibility of deriving values for the product NODO in a metal from measurements of the widths of internal-oxidation zones (subscales) formed during the internal oxidation of alloys. In addition to these authors, several other investigators7-9 have made such measurements in various silver and copper alloys at various temperatures, mostly using me- chanically polished specimen surfaces. The purpose of this investigation was to determine over a range of temperature the calculated products NODO in silver and copper from systematic measurements of the widths of internally oxidized zones in Ag-In and Cu-A1 alloys, using specimens with electro-etched surfaces. Consider the internal oxidation of a dilute solid-solution alloy in which the base metal is not oxidized. If diffusion of oxygen is the rate-controlling step, then the depth 5 of the internal-oxidation zone is given by where y is a dimensionless parameter and t the oxidation time. Based on the solution of the diffusion equation, wagnerl0 has shown that under certain conditions several mathematical simplifications may be used in the evaluation of y, so that Eq. [I] may be expressed for the case that y « 1, and y(Do/DB) 1, where N(so)the is the equilibrium solubility (mole fraction) of oxygen in the solvent metal for the given oxidizing atmosphere, is the mole fraction of the solute metal in the bulk alloy, 4 is the diffusivity of the solute in the base metal, and v is the ratio of oxygen atoms to solute atoms in the precipitated oxide BO,. Rearranging Eq. [2] Maak&apos; has derived an expression analogous to Eq. [3] for the case that an external-oxide scale, in addition to an internal-oxidation zone, is formed. For this combined external and internal oxidation for the case that 5/2 where 5 is the distance of the internal-oxidation front from the original external s.urface, and X is the distance of the external oxide/metal interface from the original external interface. The function F is defined by the expression

Jan 1, 1964

I. THE ENERGY CRISIS STILL EXISTS A. ALTHOUGH IMMEDIATE SHORTAGE IS NOW ALLEVIATED BY: 1. Increased imports through lifting of Arabian embargo; a. Imports were 6.5 million bpd during the week ended January 17, or 1.3 million bpd higher than last year. b. 1.7 million bpd higher than March &apos;74. 2. Decreased demand resulting from higher prices and conservation. a. January 1975 demand was 17.9 million bpd, 0.1 million bpd lower than January 1973 and … b. 2.0 million bpd, or 10% under originally projected demands for January 1975. c. Gasoline (excluding tax) went up 59% from 27d/gal in August 1973 to 43c/gal in September 1974. d. Gasoline demand for January 1975 was 6.3 million bpd, 0.1 million bpd higher than 1973 but 0.7 million bpd ( 10% ) under projected demand. 3. Increased crude and major product inventories. a. January 17, 1975 crude and major product inventory was 794 million bbl, 99 million bbl higher than the same date in 1973. b. 105 million bbl increase since April 12, compared to only 38 million bbl increase during the same period in 1973.

Jan 5, 1975

By Per Kofstad, Hallstein Kjöllesdal

The oxidation behavior of niobium (columbium) has been studied in the temperature range 500° to 1200°C and at oxygen pressures of 760,100, 10, 1, and 0.1 mm of Hg. The work comprises kinetic studies of the oxidation as well as structural investigations of the oxidized specimens by means of X-ray diffraction, electron diffraclion, electron microscopy, and metal1ographic techniques. The oxidation reaction has a highly irregular temperature dependence and is unusually sensitive to changes in the oxygen pressure. The complex oxidation behavior is suggested to be due to formation of different Nb2O5 modifications in different temperature regions. During oxidation oxygen is also dissolved in the metal. Formation of oxide whiskers take place on the oxide surface at temperatures above approximately 800°C. VARIOUS aspects of the oxidation of niobium have been studied.l-12 The investigations show that niobium has a complex oxidation behavior, and the oxidation

Jan 1, 1962

By S. S. Nosseir

INTRODUCTION General Quebec Cartier Mining Co. (QCM) , Canadian subsidiary of U . S. Steel Corporation, operates two iron ore concentrators at the North Shore of Quebec, Canada. In its own operation at Mount-Wright, over 45 million tons per year of crude ore are ground, fully autogenous in six 32&apos; x 11-1/2 Dominion Engineering mills. At Fire Lake/Lac Jeannine operation, property of Sidbec Normines Incorporated (SNI), over 11 million tons per year of crude ore are now ground, fully autogenous, in twelve 18&apos; x 5&apos; Hardinge mills. The Lac Jeannine concentrator was originally constructed by Q.C.M. in 1960 and. was since operated, in beneficiating the Lac Jeannine ore deposit until fully depleted in 1977, at an annual capacity of 18 million tons of crude per year. In 1976, the new company, Sidbec Normines, was incorporated by Sidbec Corporation, British Steel Corporation and Quebec Cartier Mining Co., in proportional holding of approximately 51 to 41 to 8 percent respectively. Since the acquisition of the Fire Lake ore deposit and Jeannine concentrator by the new company, QCM has continued managing the mine and concentrator operaions for S idbec Normines Inc . Flow-Sheet Design Jeannine concentrator comprises twelve identical mill lines grinding the ROM from 8" top size to 14 mesh nominal size range. Driven by 600 H.P. motor at 13.4 RPM corresponding to 73% critical speed, each mill is operated in a closed circuit with two 5&apos; x 12&apos; vibrating screens dressed with 2.5 x 25 mm slot opening cloth. The ground product or screen undersize at an average grade of 33.4% Fe is beneficiated by gravity separation to produce 66.5% Fe concentrate grade at 42.7% weight recovery. At the rougher stage, each mill line con-tains 128 Humphrey 13- 1/2" pitch spirals which produce a final tailing, middling and concentrate products. The rougher concentrate - so produced - is classified by 64 Humphrey pocket sizers into an underflow, which is a final concentrate product and an overflow which is further upgraded by 64 cleaner spirals. The cleaner concentrate together with the pocket sizer &apos;s underflow are dewatered on a 13&apos;[6] horizontal filter. The rougher middling together with the cleaner tailing are recycled in a closed circuit with the spirals. The rougher tailing after dewatering by hydrocyclones is disposed of by a

Jan 1, 1982

By P. M. Robinson, M. B. Beaer

The heat of formation at 0°C of the intermetallic compound AgMg as a function of composition has been determined by tin-solution calorimetry. In this technique, the heat of formation is determined as the difference in the heat effects, at infinite dilution, on dissolution of the compound and a mixture of its components added alternately to liquid tin from a reference temperature, such as 0°C. The experimental technique and method of calculation have been described elsewhere.&apos; The compound AgMg is a 0 electron compound with a CsCl structure.2 In nonstoichiometric compositions of the compound, atoms of the excess component are substituted on the AgMg lattice.3,4 The homogeneity range is approximately 43 to 52 at. pct Mg at room temperature and 38 to 65 at. pct Mg at 500Oc.2 Single-phase alloys outside the homogeneity range at room temperature can be prepared by quenching. The specimens used in this investigation contained from 39.0 to 54.8 at. pct Mg. They were made available, together with analyses, by the Research Laboratory of the General Electric Co. Metallographic examination of the quenched specimens containing 39.0 and 54.8 at. pct Mg gave no indication of the presence of a precipitated phase. The measured heats of formation are listed in Table I. In a calorimetric run, four to six additions each of a specimen and the corresponding mixture of the components were made. One or two calorimetric runs were carried out with each composition depending on the amount of samples available. The limits stated in the table represent the experimental precision and not the absolute accuracy of the measurements. Thermodynamic investigations of the compound AgMg have been made by Trzebiatowski and Terpilowski 5 and Kachi.6,7 The heats of formation at 500°C derived from the combined electromotive-force values of these investigators8 are more exothermic than the heats of formation at 0°C reported here, Fig. 1. The differences cannot be explained by the difference in temperature. This is shown by the published heat capacity of the stoichiometric compound, which obeys the Kopp-Neumann rule from 0" to 500oC.9 The heats of formation derived from equilibrium data are generally considered to be less accurate than those determined by calorimetric methods because small errors in the change of the equilibrium values with temperature have a large effect on the calculated value of the heat of formation. The heats of formation derived from the combined published electromotive-force values reach a maximum at 56.0 at. pct Mg.8 Kachi reports a maximum of -4.7 kcal per g-atom at the stoichiometric composition.7 The present results also indicate a maximum at or near the stoichio metric composition, Fig. 1.

Jan 1, 1964

By J. S. Kirkaldy, R. C. Dorward

The solubilities of copper and gold in silicon and germanium have been determined by equilibration of the solid semiconductors with coexisting liquidus phases. The experimental results were combined with estimated activity coefficients of the liquidus alloys to determine the following partial molar enthalpies and excess entropies of solution for the systems: Solubility measurements were also carried out below the eutectic temperature of the copper-silicon system. The partial molar enthalpy and excess entropy of copper in silicon with respect to Cu3Si were determined to be 40.2 ±0.5 kcal per mole and 9.7 i 0.5 cal per mole — OK, respectively. It has been dentonstrated that nu-cleation and growth of intermediate compounds are retarded in copper-silicon diffusion couples. This points out the dangers of determining equilibrium solubilities by dijjusion methods. THE thermodynamic requirement for equilibrium between contiguous phases at constant pressure and temperature is that the chemical potential of each component be constant throughout the system. Equating expressions for the chemical potential of solute, component 2 say, in two phases, alpha and beta, gives

Jan 1, 1969

Executive R. C. ALLEN, Chairman ERLE V. DAVELER, Vice-chairman WILDER JUDSON R. M. ROOSEVELT G. B. WATERHOUSE Finance H. G. MOULTON, Chairman HENRY KRUMB PAUL D. MERICA H. T. HAMILTON, Consultant Admissions W. B. HEROY, Chairman GEORGE D. BARRON, Vice-chairman H. T. HAMILTON PHILIP KRAFT W. A. SINSHEIMER

Jan 1, 1937

By D. O. Noel, E. B. Gebert, J. D. Shaw

IT. is, of course, expected that manufacture of the various metal powders should involve numerous methods adapted to the specific characteristics of the metals themselves. Several methods for powdering a particular metal may be available, but the physical characteristics of the powder as produced by the several methods may be quite different. In the final analysis, the application of the metal powder dictates the physical properties that the metal powder must possess, and this, to a very large extent, will determine which of several possible methods of. manufacture is adopted. Sometimes powder of the same metal is produced by more than one method, each powder having specific characteristics .that make it particularly suitable for its application. For purposes of classification, the various methods of powdering metals can be grouped as follows: 1. Machining. 2. Milling (by ball mills, stamps, attrition mills, etc.). 3. Shotting (pouring molten metal into water, air, etc.). 4. Granulation (by stirring molten metal while solidifying). 5. Atomizing (disintegrating by steam, compressed air, etc.). 6. Condensation of metal vapor. 7. Reduction of oxide powders. 8. Chemical precipitation. 9. Electrolytic deposition. 10. Sintering (for producing alloy powders). 11. Formation of an alloy followed by dissolving or otherwise removing one of the alloying constituents. MACHINING Turnings, cuttings, filings, etc., are, of course, produced as by-products in machine work of various kinds, and find certain applications as metal powders. For example, iron powder so produced is used extensively as a reducing or precipitating agent and as a constituent in certain hard-finish cements.

Jan 1, 1938

By J. D. Shaw, E. B. Gebert, D. O. Noel

It is, of course, expected that manufacture of the various metal powders should involve numerous methods adapted to the specific characteristics of the metals themselves. Several methods for powdering a particular metal may be available, but the physical characteristics of the powder as produced by the several methods may be quite different. In the final analysis, the application of the metal powder dictates the physical properties that the metal powder must possess, and this, to a very large extent, will determine which of several possible methods of manufacture is adopted. Sometimes powder of the same metal is produced by more than one method, each powder having specific characteristics that make it particularly suitable for its application. For purposes of classification, the various methods of powdering metals can be grouped as follows: 1. Machining. 2. Milling (by ball mills, stamps, attrition mills, etc.). 3. Shotting (pouring molten metal into water, air, etc.). 4. Granulation (by stirring molten metal while solidifying). 5. Atomizing (disintegrating by steam, compressed air, etc.). 6. Condensation of metal vapor. 7. Reduction of oxide powders. 8. Chemical precipitation. 9. Electrolytic deposition. 10. Sintering (for producing alloy powders). 11. Formation of an alloy followed by dissolving or otherwise removing one of the alloying constituents. Machining Turnings, cuttings, filings, etc., are, of course, produced as by-products in machine work of various kinds, and find certain applications as metal powders. For example, iron powder so produced is used extensively as a reducing or precipitating agent and as a constituent in certain hard-finish cements.

Jan 1, 1938

By D. O. Noel, E. B. Gebert, J. D. Shaw

It is, of course, expected that manufacture of the various metal powders should involve numerous methods adapted to the specific characteristics of the metals themselves. Several methods for powdering a particular metal may be available, but the physical characteristics of the powder as produced by the several methods may be quite different. In the final analysis, the application of the metal powder dictates the physical properties that the metal powder must possess, and this, to a very large extent, will determine which of several possible methods of manufacture is adopted. Sometimes powder of the same metal is produced by more than one method, each powder having specific characteristics that make it particularly suitable for its application. For purposes of classification, the various methods of powdering metals can be grouped as follows: 1. Machining. 2. Milling (by ball mills, stamps, attrition mills, etc.). 3. Shotting (pouring molten metal into water, air, etc.). 4. Granulation (by stirring molten metal while solidifying). 5. Atomizing (disintegrating by steam, compressed air, etc.). 6. Condensation of metal vapor. 7. Reduction of oxide powders. 8. Chemical precipitation. 9. Electrolytic deposition. 10. Sintering (for producing alloy powders). 11. Formation of an alloy followed by dissolving or otherwise removing one of the alloying constituents. Machining Turnings, cuttings, filings, etc., are, of course, produced as by-products in machine work of various kinds, and find certain applications as metal powders. For example, iron powder so produced is used extensively as a reducing or precipitating agent and as a constituent in certain hard-finish cements.

Jan 1, 1938

By M. J. Duggin

Evidence is produced below to show that the carbide "Mn,C," reported by Kw,and ersson&apos; is really a mixture of two carbides. One is an Fe-Mn carbide which is probably isomorphous with hexagonal Mn15C4, reported by Bouchand and Pruchart.&apos; The othev is a tetragonul Fe-Mn carbide with cell dzmensions close to those reported previously by the author3 for an Fe0.1Mno. carbide. IN a recent investigation several Fe-Mn carbide specimens, each weighing about 5 g, were prepared in a manner previously reported.3 Debye-Scherrer X-ray photographs were taken of at least two samples from different parts of each specimen using a 114.6-mm-diam camera. The exposure time was 2 hr using filtered chromium radiation, a tube voltage of 35 kv and a tube current of 10 ma. In the case of one specimen, which contained nominal weight percentages of iron, manganese, and carbon of 9.40, 84. 6,, and 6.0, respectively, duplicate chemical analyses indicated f1.2 and 5.5 wt pct C to be present in the two samples taken of the specimen: no graphitic carbon was detected. One powder-diffraction pattern of this specimen indicated the presence of the tetragons: carbide reported earlier,3 (a = 6.772A, c = 9.427A), plus cementite (O), plus an unknown structure. The second powder-

Jan 1, 1970

By Sidney Diamond, Charles Wert

The anelastic behavior of Ni-C alloys of nominal 0.5 wt pct C has been determined over a wide temperature range. The most prominent damping effect, that due to the presence of C-C pairs, was measured both above and below the Curie temperature. The effect on diffusion of the C-C pairs by magnetic spin ordering is deduced to be very small. The diffusion coefficient for C-C pairs in nickel is given by the expression ThE solid-state diffusion of both interstitial and sub-stitutional solutes in ferromagnetic solvents has been the subject of a number of investigations. Measurements of the self-diffusion of iron,"&apos; the diffusion of nickel1,4 gold,"1 and cobalt4 into iron, the diffusion of vanadium in an Fe-V alloy,&apos; and the diffusion of cobalt and nickel in Co-Ni alloys9 show an unusual behavior. In both the ferromagnetic and paramagnetic regions, far removed from the Curie temperature, the diffusion coefficient, D, obeys an Arrhenius relation. However, in the ferromagnetic region D is markedly lower than would be expected from an extrapolation of D from the paramagnetic region; this results in an increase in the activation energy in the ferromagnetic region. The cause of this effect for self- and substitutional diffusion is uncertain. Since substitutional diffusion is presumed to proceed by the vacancy mechanism, it has been variously suggested that the increase in activation energy below the Curie temperature is due to one of the following causes: 1) an increase in the energy of formation of a vacancy because of local demagnetization around the vacancy,2,3 2) a decrease in vacancy mobility in the ferromagnetic state,&apos; or 3) an effect on the entropy of motion of a vacancy in the ferromagnetic state.5 Quite another explanation for the observed effect has been proposed by Borg.7 The diffusion coefficient is a function of the elastic constants of a metal. Since the temperature dependence of the elastic constants often changes at the Curie temperature, the temperature dependence of D might be expected to undergo a universal change in the vicinity of the Curie temperature. The magnitude of the change in D would, of course, depend on the magnitude of the change in the elastic constants. One might attempt to differentiate between vacancy effects and elastic effects by making measurements on interstitial alloys, in which vacancy effects play no role. Some measurements of this type have been made; this paper reports additional observations. Measurements of the diffusion of carbon in a Fe by a tracer technique8 and of the diffusion of carbon in a Co and in Co-Fe alloys by a decarburization method9 reveal a small effect, if any, on D of the magnetic change at the Curie temperature. If the effect of spin ordering is only to alter the parameters controlling formation or mobility of a vacancy, interstitial diffusion should indeed remain unchanged through the magnetic transition, as these measurements imply.

Jan 1, 1968

By Michel Elle, John Chipman

The equilibrium of gaseous H20-H2 mixtures with liquid iron-columbium alloys in the range 0.2 to 2.4 pct Cb shows that the activity coefficient of oxygen is strongly diminished by Cb. The oxide in equilibvium with the melt is identified as CbO2. The equi-librium constant at 1600 deg C is [%Cb][fo. % 0]2 = 2.9 x 10-4 where log fo= —0.14 [%Cb]. The deoxidizing power is comparable to that of vanadium, COLUMBIUM (niobium) is finding increasing use as an alloying element in steel. It forms several stable oxides and is removed from the steel bath by oxidizing slags. The specific oxide formed in steel-making is not known although Sazanov snd Shvarts-man,&apos; in their study of its partition between slag and metal, have assumed this to be Cb2O5 It is the purpose of this study to determine its deoxidation equilibrium at 1600°c, the nature of the oxide formed and the interaction between columbium and oxygen in the melt. The investigation was based on the reaction H2 + O=H2O; Kl=ph2o/(pH2[%O]) [1] in which 0 signifies oxygen dissolved in the liquid Fe-Cb alloy. The method and equipment of recent studies2,3 were utilized. In brief, the molten alloy, contained in a small crucible, is held in a controlled mixture of steam, hydrogen, and argon until equilibrium is reached. It is then cooled quickly in a stream of helium and analyzed for columbium and oxygen, the latter by vacuum fusion.4 The solubility of the oxide in liquid iron is determined by visual observation of scum formation on the surface of the melt. Crucibles of MgO proved unsatisfactory because of partial reduction at the metal interface. Pure A12O3 crucibles were heavily blackened by the melt and could not be used at high concentrations of columbium. Silica crucibles were more satisfactory but these introduced some silicon into the melt thus necessitating an uncertain but small correction for its effect on the activities of the other solutes. In silica it was possible to observe the scum formation when the melt became saturated and there was no apparent reaction between scum and crucible. The procedure for saturated melts was as follows. The alloy charge composed of electrolytic iron and columbium grain was melted and held at constant temperature. The gas ratio H2O:H2 was progressively increased by raising the saturator temperature until a point was reached at which a bright film appeared on the surface and covered it completely. This film could be dissolved by raising the temperature of the melt and it reappeared when the temperature was lowered. Under a partial film covering, utilizing bare areas for temperature measurement, the melt was brought to equilibrium with the atmosphere by holding at temperature for about 1 hr as suggested by Dastur.5 INTERACTION IN THE MELT Several experiments were made without columbium to establish a base in silica crucibles. The results are shown in Table I. Except for one run, they agreed well with Floridis1 comparable results in other crucibles. The value of log K1 is taken from Table I as 0.58. This is slightly larger than Floridis&apos; result, probably because of less elaborate safeguards against thermal diffusion. Results obtained in the Fe-Cb-0 series in which no scum formed are shown in Table 11. The equilibrium ratio Kl is defined as ph2o/(ph2[%O]). It is related to the equilibrium constant K1 as follows:

Jan 1, 1962

By H. O. Pickard

The Hoover and Mason Phosphate Co. made extensive tests with a rig manufactured by the Paris Manufacturing CO. as shown in Fig I. It is powered by a 20-hp Wisconsin air-cooled gasoline engine, through a single reduction gear and flexible coupling. The machine operates a 4-in. auger with a detachable cutting head, as shown in Fig I. The auger length is 3 ft 4 in. and the cutting head averages 7 in. The unit has a drilling capacity of 100 ft and can drill an average prospect hole in less than 15 min. The weight of the motor is used for the thrust feed for driving the auger and head down, aided by the spiral shape of the auger. The rate of downward motion—i.e., the feed—is regulated by a handwheel. The string of tools is withdrawn from the hole by a gear train controlled by a hand-operated clutch. In fairly dry strata, not of the consistency of mud, the action of the auger

Jan 1, 1948