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By Louis Raymond, John E. Dorn

The object of this investigation was to analyze the recovery that arises when the stress on a specimen undertaking creep is reduced. For this purpose annealed specimens of high-purity aluminum were precrept under a stress of 1000 bsi to a strain of 0.08 following which the stress was reduced for various periods of time to 10, 250, 500, or 700 psi. When the original stress was reapplied the subsequent creep curve lay above that for the unre-covered state and below that for the original annealed state. Analyses on the kinetics of this recovery as a function of the temperature gave a stress-sensitive activation energy that decreased as the reduced stress was increased from a value of 64,000 cal per mole at 10 psi to 37,000 cal per mole at 750 psi. Recovery was also detected and measured during creep under the reduced stress. Following a short initial period, the creep rate under the reduced stress increased monotonically until it reached the secondary-creep rate for the reduced stress. The temperature dependence of this phenomenon was also shown to be correlatable in terms of the previously deduced activation energy for recovery. The activation energies for creep of most pure metals at high temperatures have been shown to agree well with those for self-diffusion.'j2 Since the true secondary stage of creep is usually due to the steady-state balance between the rate of strain hardening and the rate of recovery, it is generally thought that the activation energy for recovery of the creep-induced substructure equals that for creep itself. A shoft time ago, however, Ludemann, Shepard, and Dorn~ found that the activation energy for recovery of the creep-induced substructure in high-purity aluminum under zero stress was almost twice that for self-diffusion, namely about 65,000 cal per mole; obviously recovery under reduced stresses differs in some significant way from the recovery that accompanies the secondary stage of creep. The major purpose of this investigation is to study the effect of stress on the re- covery of the creep-induced substructure in order to provide a better understanding of the recovery mechanism itself. EXPERIMENTAL TECHNIQUE High purity aluminum, containing 0.004 pct Cu, 0.002 pct Fe, and 0.001 pct Si, used in this investigation, was in the form of 0.100-in.-thick sheet which has been cold-rolled to the H-18 temper. Creep specimens were milled from the sheet with their tensile axes in the rolling direction. All specimens were then heated at 686°K for 1 hr followed by air cooling in order to produce an annealed structure which exhibited a uniform equiaxed grain size of about 4 grains per mm. Tests were run in creep machines fitted with Andrade-Chalmers type of lever arms so contoured as to maintain the stress constant to within 0.05 pct of the reported values. Constant temperatures to *O.l°K were obtained by complete immersion of each specimen in a temperature-controlled and agitated bath of molten KN02-KNOs mixture. Where changes in temperature were involved, the change was effected in less than 2 min by manually replacing one bath by another controlled at the second temperature. Displacements over the gage section were sensed by linear differential transformers, the output of which was autographically recorded. The calculated strain measurements were sensitive to 5x EXPERIMENTAL PROCEDURE The following analyses are based on extensions of the previously announced effect of the temperature on the creep strain,2 namely for a = constant, where e = the total true tensile creep strain for a given applied true tensile stress, t = the duration of the test, R = the gas constant, T = the absolute temperature, Q, = the activation energy per mole for creep which is independent of the stress, / = a function of 8, = and of the stress, and a = the stress. The validity of this correlation for high-purity aluminum is demonstrated in Fig. 1 for temperatures in the near vicinity of 600°K; the activation energy for creep, Q,, which is approximately that for self-diffusion, is insensitive to the applied stress

Jan 1, 1964

By Paul J. Fopiano

SOME relationships between the flow characteristics of iron single crystals of 99.9 pct purity and the behavior of imperfections have been investigated. X-ray rocking-curve measurements and etch-pit counts were made as a function of plastic strain, and compared to the stress-strain curve obtained on a modified Polyani tensile machine. Crystals grown from rolled strips of vacuum-melted iron by the strain-anneal method1 had a high preference for a (110) longitudinal direction and a (211) face normal. The tensile specimens were prepared from 2 by i by 0.040 in. single crystals having a gage area of 3 by \ in. Rocking-curve measurements were carried out with a highly perfect germanium monochromating crystal in which the dazz spacing was matched to that of the dZl1 in the ir0n.l Well-collimated CuKal radiation was used throughout. These procedures practically eliminated errors due to geometrical and wavelength resolution. Inasmuch as the rocking-curve half breadth may vary markedly from point to point in the specimen being irradiated, the crystals were strained in place by mounting a hydraulic loading device on the double-crystal spectrometer. The rocking curves were taken after each increment of strain in the unloaded condition, since no observable difference was found in the rocking-curves between the loaded and unloaded states. The rocking-curve half breadths of the as-grown specimens were in the range 90 to 120 sec of arc when the beam irradiated an area of about -£ by -& in. on the specimen. DeMarco and weiss3 have shown that, for a well-colli- mated X-ray beam, irradiating about 10"! sq in. of the very same material, half breadths within 10 pct of the Darwin natural half breadth were observed. Since the rocking-curve specimens were stressed by the load-unload technique, the strain achieved at any given stress depended on the time of holding because of low-temperature creep. Fig. 1 shows the rocking-curve half breadth (also area/peak height) as a function of plastic strain for a relatively short holding time (2 to 5 min) at each stress level. For strains less than 0.1 pct the rocking-curve breadth is essentially constant; it is only for larger strains that there occurs a significant increase in this breadth. Where the holding times at each stress level were longer (by well over an order of magnitude) there occurs a significant increase in the rocking-curve breadth only after plastic strains of the order of 0.6 pct had been introduced into the specimen. This observation is related to the time dependence of creep phenomena and emphasizes the difficulty in comparing data obtained by two such different straining methods. Etch-pit results were obtained using a 2 pct nital etch on specimens strained in the range of 0 to 1 pct. Prior to etching, all specimens were annealed for 3 hr at 150°C, the carbon content being sufficient to decorate the dislocations for strains of at least 1 p~t.~ The data points were all taken from parts of the same single crystal which had been strained with short holding times at stress in increments of strain of the order of tenths of 1 pct. The (211) plane is particularly difficult to etch-pit in vacuum-melted iron and therefore it is felt that these values are as much as an order of magnitude low. Fig. 2 shows the etch-pit density as a function of plastic strain. The smooth curve passing through the data points is not meant to infer a quantitative correlation with the rocking-curve data. What is of interest, however, is the change in etch-pit density in the region of 0.2 pct plastic strain. The first three increments in strain (points 2,3, and 4) did not produce a measurable change in the etch-pit density while subsequent increments did produce a measurable change. While the absolute values of these results do not appear to be cor-

Jan 1, 1967

By L. D. Jaffe

From the limited experimental data in the literature, preliminary values were derived for the thermal diffusivity of titanium alloys and for the quenching severity of various mediums used in heat treating them. The thermal diffusivity, under quenching conditions, was found to be approximately 0.006 sq in. per sec. The quenching severity H for brine is approximately 8, for water 4, for oil 0.9, for air 0.08, and for argon 0.02 in. -' The quenching severity of the Jominy-Boegehold end quench water jet against titanium alloys appeared effectively infinite. Using these results, graphs were prepared showing ideal round sizes and Jominy positions having cooled conditions equivalent to those at the center and surface of titanium alloy rounds, plates, and sheets quenched in various media. THE importance of proper quenching of steel parts has long been recognized, and much quantitative work has been done on the subject. Recently, it has become evident that for titanium alloys, too, the rate of cooling from high temperature is important in obtaining high strength and hardness. Maximum hardness and strength may not be attained in the as-quenched condition but rather after a subsequent tempering or aging treatment. Nevertheless, if cooling from temperatures near the ß transus is too slow, for the hardenability of the alloy used, a soft a-ß structure will be produced which will not harden on subsequent aging.'-" The quenching of parts of practical interest can usually be approximated by simple shapes, such as cylinders (rounds) and plates (or sheets) of various thickness and combinations of these. A standard that has been used extensively for comparison purposes is the center of an ideal round: a cylinder quenched in a medium that instantly cools its surface to room temperature." For steel, considerable practical use has been made of the Jominy-Boegehold end quench hardenability specimen,".' and its value for titanium alloys is becoming evident. It therefore appeared worthwhile to attempt comparison of cooling conditions in titanium alloy rounds and plates of various sizes quenched in various media with cooling conditions in ideal round and Jominy quenches. These comparisons are based upon measurements reported in the literature. Quenching Severity of Jominy Water Jet and of Brine As one starting point, there were available the data of Phillips and Tobin8 for positions on a Jominy specimen having the same hardness as the centers of round bars of various size quenched in 5 pct NaOH brine, shown in Table I. The Jominy speci- men is a round bar, normally 1 in. diam and 4 in. long, quenched by a water jet at one end and air cooled on the other surfaces.' After quenching, hardness is measured vs distance from the quenched end on longitudinal flats ground 0.015 to 0.10 in. deep. The data of Table I were based on a subsized specimen of 3/4 in. diam. Since the positions tested are close to the jet quenched end, cooling there may be taken as due solely to the jet, and air cooling can be neglected. These positions are then equivalent to various locations near one surface of an 8 in. plate jet quenched on both sides. Assuming that the thermal diffusivity is constant and that Newton's law of cooling holds, temperature vs time curves at corresponding positions in Jominy specimen and round bar then may be calculated for various assumed quenching severities of the water jet and of the brine. Cooling curves are generally taken to be metallurgically equivalent when they show the same time to reach a temperature halfway between the initial (heat treating) temperature and the final (quenching medium) temperature.5 On this basis, using Russell's tables,9 upplemented by heat flow tables derived in similar fashion by the author for positions near the surface of severely quenched plates, it was found that the experimentally observed equivalence, Table I, was best matched if the quenching severity H of the Jominy water jet was taken as infinite (H = co). This means that the jet quenched surface should be considered as instantly cooled to the temperature of the water

Jan 1, 1956

By Carl Wagner

AT elevated temperatures. most metals react with oxygen, sulphur, or halogen rather rapidly, although a coherent layer of the reaction product is formed and separates the two reactants from each other. The reaction proceeds, since one of the reactants or both dissolve and diffuse across the reaction product. In many cases the rate is diffusion-controlled as is indicated by the validity of the parabolic rate law of Tammann' and Pilling and Bed-worth. wott and Gurney' have suggested that diffusion processes also play an important part in the case of reduction reactions. This is confirmed by observations reported below. In general, however, the rate of reduction reactions is determined not exclusively by diffusion processes. Reduction of Silver Sulphide Above 179°C silver reacts with liquid sulphur very rapidly. Since the volume of silver sulphide is much greater than that of the equivalent amount of silver, a coherent layer of silver sulphide is formed in accordance with a rule established by Pilling and Bedworth.' According to Wagner','' silver ions and electrons migrate from the Ag-Ag,S interface across the silver sulphide layer to the Ag,S-S interface where the following chemical reaction takes place: 2 Ag +2 e + S Ag,S If silver sulphide is reduced by means of hydrogen, the equivalent volume of the reaction product, silver, is less than that of the initial substance. Accordingly, no coherent layer of silver is formed. Instead, fine filaments of silver grow from the silver sulphide surface into the ambient gas. This phenomenon has been investigated by many authors, especially by Kohlschiitter," who has also reviewed previous literature. Kohlschiitter has interpreted the observed phenomena by migration of "silver" dissolved in silver sulphide. According to Wagner,' silver ions and electrons are the migrating particles. The sequence of consecutive steps involved in the reduction of silver sulphide is shown in Fig. 1. Every- where at the silver sulphide surface, hydrogen reacts with sulphur so that excess silver ions and electrons are formed, which migrate at random until they are captured by a nucleus of metallic silver. If such a nucleus lies on the surface and material is furnished only from the base, the original nucleus is pushed outward and thus a filament grows. Removal of sulphur and formation of metallic silver may also take place at different points when silver sulphide is reduced by copper. The reaction is usually formulated as: AgL,S + 2 Cu = Cu.S + 2 Ag but the actual reaction products are silver and a solid solution of the two sulphides. Primarily, excess silver and electrons dissolved in silver sulphide are formed. After a certain degree of supersaturation has been reached, nuclei of metallic silver may form. If, however, a silver foil is in contact with the silver-sulphide phase, as is shown in Fig. 2, silver ions and electrons will migrate to the Ag,S-Ag interface where they are captured. To confirm this scheme, a composite sample, with the geometry shown in Fig. 2, was made by pressing silver sulphide tablets 0.6 cm in diam and 0.4 cm high together with a copper cylinder and a silver foil in crucibles of iron, which does not react appreciably with silver sulphide at 400°C. After heating for 3 to 7 hr at 400°C. the specimens were cut normal to the phase boundaries. It was found that the thickness of the silver foil had increased from 0.01 cm as initial value to a maximum value of 0.06 cm. In the interior of the silver sulphide tablets only traces of metallic silver were observed. These observations show clearly that silver ions and electrons may migrate over distances of several millimeters before they form metallic silver. Reduction of Cuprous Sulphide by Iron The behavior of cuprous sulphide is very similar to that of silver sulphide, since the lattices of the high temperature modifications are identical, but in Cu.S there is a variable deficit rather than an excess of metal. To investigate the displacement of copper by iron, cuprous sulphide powder covered by a copper foil was pressed in iron crucibles and heated up to 800 °C for 3 to 7 hr. The thickness of the copper foil increased from an initial value of 0.005 cm to a nlaximum value of 0.029 cm. In the interior of the

Jan 1, 1953

By G. Wakefield, A. H. Daane, D. H. Dennison, F. H. Spedding

Preparation of pure scandium metal was accomplished by calcium reduction of the fluoride by two methods; a low-temperatzdre alloy process and direct reduction with subsequent distillation of the product. The following properties were determined: melting point: 1181OK; boiling point (calculated): 3000°K; lattice constants at 298°K (hexagonal lattice): a = 3.308 * 0.001 A, c = 5.267 * 0.003A; calculated density at 298°K, g per cm3: 2,990 + 0.007; electrical resistivity, ohm-cm: 299°K, 66.6 ± 0.2 x 10 -6; 373oK, 77.4 * 0.2 x 10 -6; thermal coefficient at 299°K, ohm-cm per deg: 5.4 X x; heat of sublimation at 298°K, kcal pel- mole: 80.79. The vapor pressure was determined as a function of temperature between 1505o and 1748°K, with the data fitted to a straight line shielding the equation: Log Pmm = -1.718 X 104/ToK + 8.298. SCANDIUM, element number 21, was first discovered by Nilson in 1879 and was recognized as the Ekaboron as predicted by Mendeleff. As it is in group III of the periodic table, the general properties are a little like aluminum and also resemble quite closely the properties of yttrium and the rare-earth metals, in both the metallic and ionic form. Although the earth's crust contains approximately 5 ppm of scandium (the element is as abundant as arsenic and twice as abundant as boron) it generally occurs so widely distributed that it has earned the reputation of being very rare. The one exception to this is the mineral thortveitite, which has been found in Madagascar (20 pct Sc2O3) and in Norway (35 pct Sc2O3). Scandium also occurs in small but distinct amounts in uranium and rare-earth ores; the recent larger scale processing of these materials has made some scandium available from these sources. As with other naturally-occurring monoisotopic elements (except Be), scandium contains an odd number of protons and an even number of neutrons. Scandium metal was first prepared by Fischer and coworkers1 in 1937 by electrolysis of scandium chloride in a molten eutectic mixture of lithium and potassium chlorides, using molten zinc as a cathode and collector of the scandium metal produced. The zinc was removed from the Zn-2 pct Sc alloy by vacuum distillation, leaving a product reported to be 94 to 98 pct Sc, with the main impurities being iron and silicon. They reported a melting point of 1400° C for this material. Scandium has also been prepared by the reduction of scandium chloride with potassium metal in a glass apparatus by Bommer and Hohmann in 1941,' resulting in a mixture of metal and potassium chloride; these workers did not isolate the metal proper, but the X-ray diffraction of the slag-metal mixture showed it to be hexagonal with a = 3.30A, c = 5.45A. petru3, 4 and coworkers have recently reported the preparation of the metal in a compact form by the reduction of either ScF3 or ScC13 with calcium metal and subsequent distillation of the product. This process probably yielded a metal of high purity, but they list no chemical analysis nor do they list any of the properties of their product. Previous related work in this Laboratory has been concerned with the production of yttrium and the rare-earth metals and the determination of their physical properties. Because of its similarity to these metals, scandium is being included in this study. PREPARATION OF SCANDIUM METAL The preparation of yttrium and the rare-earth metals may be accomplished by reduction of their fluorides with calcium metal in tantalum crucibles.5 This process leads to the introduction of tantalum (up to 0.5 pct) as an impurity in the higher melting rare earths, but since the tantalum occurs as dendrites, uncombined with the rare-earth metals, its presence is not objectionable in some cases. The preparation of scandium metal in this manner, however, was found to yield a product containing 2 to 5 pct Ta. To obtain a purer product, the following two methods were developed for the reduction of scandium fluoride with calcium metal: i) a low-temperature process utilizing zinc to form a low

Jan 1, 1961

By Gearge C. Wallick

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

Jan 1, 1953

By P. G. Winchell, M. Cohen, J. Woodilla

The kinetics of stabilization have been studied with respect to the isothermal component of the martensitic reaction in ivon-nickel alloys. Although the carbun (or nit-vogen) content may be very low in these alloys, it plays an important vole in the stabilization phenomenon. In fact, if these interstitial elements are removed by moist-hydrogen treatment, no evidence of stabilization is found. In the pvesence of 0.007 wt pct C, the activation energy for the stabilization process is comparable to that -for the diffusion of carbon (or nitrogen) in ferrite) rather than in austenitc. This suggests that the interstitial diffiision controlling the stabilization occurs within the martensitic embryos rather than in the matrix of the parent phase. The kinetics furthev indicate that the interstitial atoms tend to diffuse -from the embryos toward the suvroundings, thereby immobilizing the interface. This appears to be the origin of the stabilizing effect -found in these alloys. PHILIBERT' has shown recently that interstitial elements play an important role in the stabilization of martensitic transformations in iron-base alloys. However, determination of the temperature dependence for this stabilization process has produced values (- 11,000 cal per mol) that cannot be readily identified with the activation energy for diffusion of carbon (or nitrogen) in either ferrite or austenite. In the present work, the isothermal mode of the martensitic reaction was employed to derive the activation energy for stabilization, and this was found to agree fairly well with that for diffusion of carbon (or nitrogen) in ferrite. An iron-nickel alloy with 30.8 wt pct Ni and 0.007 wt pct C was used in these experiments, the specimens being in the form of rods 0.074 in. in diam and 2 1/2 in. long. The specimens were sealed in evacuated vycor tubes and austenitized at 1100°C for 1/2 hr prior to quenching in water. A second series of specimens was decarburized (and denitrogenized) prior to the above austenitizing treatment by annealing in moist hydrogen for 182 hr at 1215°C. In order to achieve a "standard" condition in the austenitic state, all the specimens were transformed to martensite by cooling to - 195°C before the final austenitizing. Both the 0.007 pct C and the decarburized series then had an Mi temperature of about - 30°C. The isothermal martensitic transformation was carried out at a reference temperature of -88"C, and was traced by means of electrical resistance measurements. Each resistance value was normalized by dividing by the resistance of the austenitic specimen at O°C, thus giving a resist- ance ratio R. The rate of decrease of the resistance ratio (-dR/dt) provided a measure of the transformation rate. It was found that a plot of R vs log t at the reference temperature resulted in a slightly curved line which could be closely approximated by two straight lines intersecting at a time of 10 min, as shown in Fig. 1. If G = dR/dlogt is the linear slope of either straight segment, the rate of transformation at any given time is: The ratio of G for 1 less than 10 min to that for t greater than 10 min was 1.19 for the 0.007 pct C series and 1.24 for the decarburized series. These ratios made it possible to compute the transformation rate at any time greater than 10 min, knowing the transformation kinetics prior to 10 min. The stabilization treatment was introduced by up-quenching from the reference temperature, usually after 10 min of isothermal transformation: to one of four stabilizing temperatures between 0" and 22"C. After holding for various times, the specimens were down-quenched to the reference temperature for additional isothermal transformation. In this way, it was possible to compare the transformation rate (- dR/dt), immediately after stabilization with the transformation rate (- dR/dt)U in the unstabilized condition. The degree of stabilization was defined as: Ordinarily, Gs and Gu would be measured after the same time of prim isothermal transformation.

Jan 1, 1960

By J. L. Bray

ACCELERATED programs, as discussed in this paper, refer to the year-around operation of a college or university with three sixteen-week or four twelve-week terms per year, with pauses between sufficient only to allow the administrative group to perform their functions. In most engineering schools, under such conditions, the conventional four-year course can be condensed to two years and eight months. These accelerated programs of study, which now exist in most of our technical institutions, were forced upon us under the influence of the outburst of wartime patriotism, perhaps one might go so far as to call it hysteria. Many of us felt and hoped that these programs were temporary expedients, but the cold hard threat of permanency now rears its head; consequently we should examine the merits and shortcomings of the new order, and evaluate its effect on our educational system, since it may persist in some form in the postwar years, because of the changing political and financial situation. Many engineering students are being educated in state supported institutions

Jan 1, 1944

By Frederick Bacorn

THE paper of Reno H.. Sales on Ore Deposits at Butte, Mont.,' is a careful and painstaking work, an important contribution to the literature of the subject. As is almost inevitable in a work of such magnitude, a few misstatements occur; and some of these I shall endeavor to point out. Mr. Sales has also fallen into error, as I think, in the application of his theory of ore deposition; and while the modification which I venture to suggest is a minor one geologically, it is of great economic importance. I refer to Mr. Sales's idea of an outward migration of mineralizing solutions from a supposed central copper zone. In order that I may make myself plain, I must first epitomize the theory as a whole, which I understand to be this: All of the Butte ore deposits (save those along the Continental fault, not herein considered) are derived from solutions of magmatic origin, which, at the beginning of their ascent from a great depth, were, acid in reaction, hot, and were carriers of copper, zinc, manganese and other metals. As they departed from their point of origin they lost heat, and, because of reactions with the wall rock, became more and more alkaline. The metal deposited from these solutions in their hot condition was principally copper; and, as they cooled, they deposited less and less copper and more and more manganese and zinc, until the copper practically disappeared. With this theory in general I have no fault to find, nor have I, it must be stated, the qualifications necessary to enable me to pass upon it critically; but it seems to me to be reasonable and probable, and is one which I am quite willing to accept.

Jan 8, 1914

By Klaus Schwerdtfeger

The equilibria in the Cr-N system have been investigated in the temperature range 1100° to 1310°C by reacting chromium powder with Nz-Hz gas mixtures. The solubility of nitrogen in chromium in equilibrium with chromium subsitvide ("Cr,N") is given by Chromium subnitride is nonstoichiotnetric; its nitrogen content is always less than that corresponding to the formula CrzN. The lattice paranzeters of quenched samples have been measured; c, and a. parameters are found to increase with increasing nitrogen content. The growth rate of the subnitride layer on chromium plates was measured by a thermogravimetric technique, using a silica spring balance. The self-diffu-sivity obtained from the theoretical analysis of the parabolic rate constant is found to decrease with increasing nitrogen content, i.e., with decreasing vacancy concentration in the nitrogen sublattice. The intrinsic nitrogen diffuivity is derived from another series of rate measurements using "CrzN" plates; the intrinsic diffusivity, DN = 3.2 X 10-a cmZ sec-' at 1200 C, is found to be essentially independent of- the subnitride composition. The concentration gradient was measured in a chromium subnitride layer by the X-ray method and found to be consistent with the derived diffusivity value. TWO chromium nitrides are known to exist:' the nonstoichiometric subnitride "CrzN" and the nitride CrN. In the present work the kinetics of the formation of chromium subnitride from chromium and nitrogen have been investigated at 1100" and 1200°C. In additional experiments the relevant equilibria have been measured. The data are used to evaluate the diffusivity of nitrogen in chromium subnitride. Since chromium nitrides are often found in chromium-containing steels, the results are expected to be helpful in the interpretation of the chemical reactions between chromium steels and nitrogen. Equilibria in the Cr-N system have been determined by several investigators.2"3 The rate of nitriding of chromium was measured by Arkharov et a1 .' in ammonia in the temperature range 800" to 1200°C. The parabolic rate law was observed. Due to the undefined nitrogen activity of the ammonia atmosphere, it is dif- ficult to interpret these rate data theoretically. An additional difficulty arises from the fact that the two-layer scale consisting of CrN and "Cr2N" was formed at the temperatures below 1030°C. The rate of nitriding of technical chromium (95 pct) was measured by Zaks in nitrogen at -1 atm in the temperature range 800" to 1300°C. EXPERIMENTAL METHODS The chromium samples were reacted with Nz-HZ gas mixtures in a vertical tube furnace, wound with Pt-10 pct Rh resistance wire. The gas-tight reaction tube was of high-purity recrystallized alumina. In nitrogen solubility measurements the nitrogen content of chromium was determined by the Kjeldahl method, on samples quenched in the cold part of the furnace. The nonstoichiometry of the subnitride and the nitriding rates were measured thermogravimetrically using a sensitive (+0.1 mg) silica spring balance. For the equilibrium measurements samples of 1 g of chromium powder contained in high-purity alumina crucibles were used. In order to remove most of the oxygen and nitrogen impurities from the chromium, the samples were initially annealed in purified hydrogen until a constant weight was obtained. The chromium plates (approximate dimensions 2 by 1 by 0.08 cm) used for the rate measurements were machined from ingots obtained by arc-melting of iodine-processed chromium. According to manufacturers' specifications the purity of the chromium powder was 99.9 pct Cr and that of the iodine-processed chromium 99.99 pct Cr. Our own spectroscopic analysis of the chromium powder yielded 0.02 pct Fe, 0.05 pct Mn, 0.05 pct Si, and 0.02 pct Ti as major impurities with all the other detectable elements below 0.005 pct. The nitrogen partial pressure of the gas phase was controlled by mixing prepurified hydrogen and nitrogen with constant pressure head capillary flowmeters. Oxygen and water vapor were removed from the mixed gas by passing it through columns of platinized asbestos (450°C) and anhydrone. The gas flowed upward in the furnace with flow rates of 300 to 500 cu cm per min (25"~). Gas tightness of the furnace system was ensured by pressure checks at regular intervals. The furnace temperature was controlled electronically in the usual manner. The reported temperatures were measured with a Pt/Pt-10 pct Rh thermocouple and are estimated to be accurate within ±$C The X-ray measurements were made with a Debye-Scherrer camera and a diffractometer using chromium radiation {\Ka = 2.29092A). EQUILIBRIUM MEASUREMENTS The experimental results of the equilibrium measurements are contained in Tables I to In. Fig. 1 shows the solubility of nitrogen in solid chromium in the temperature range 1100" to 1310°C. In Figs.

Jan 1, 1968

By R. H. Richman, F. Borik

A 0.3 pct C-12 pct Cr-6 pct Ni steel was rolled to 93 pct reduclion in area as austenite at 510°C, and then partially transformed as desired to ~rlartensite by qnenching to - 196°C. Pole figures for the austenitic matrix and for the martensitic product were separately determined by an X-ray transmission method. The deforitration texture of' the warm-worked austenite is characlerized by (110)(225) components, and is thus closely similar to those produced in a brasses. The pole jigure of the martensite in partially transformed material agrees well with that which can be constructed by transfortnation of the {110)(225) orientations according to either the Kuvdjuniov- Sacks or the Nishi-yatuu relatiotship. Howeuer, an important result of this construction is that me-third of the predicted orientations are missing. A graphical analysis can then be used to show that in deformed austenite certain crystallographic variants of martensite (related to the most probable austenite slip systems) are suppressed, resulting in this preferred transformation. The evidence for preferred transformation is corroborated by the measured elastic anisotropy of warm-rolled and fully transformed H-11 steel. EXTENSIVE plastic deformation of a polycrystal-line aggregate in a manner that causes flow predominantly in one direction results in a preferred orientation of the constituent crystallites. The particular orientations that are produced depend upon the crystal structure and composition of the material, as well as upon the temperature, mode, and degree of deformation; in any case, the preferred crystallo-graphic orientations, or textures, are reflected in directionality of mechanical properties. Although such anisotropy may be exploited in certain specialized applications, it is more commonly diminished or eliminated by heat treatment lest it interfere undesirably in subsequent forming operations or in structural design. In the recently developed thermomechanical treatments that significantly enhance the strength of some steels,1,2 considerable deformation of the metastable austenite prior to the martensite transformation is essential to the strengthening process. If the austenite is textured by the deformation, and if the transformation to martensite proceeds according to one of the relationships established for transformation in annealed austenite, then it must be expected that the martensite will also possess a preferred orientation even though the multiplicity of martensite orientations possible in a given austen- ite crystal will tend to restore some degree of randomness. The existence of a residual anisotropy, both mechanical 3-6 and crystallographic,' has been substantiated. In the latter crystallographic investigation, preferred orientations were determined for the martensitic structure of an SAE 4340 steel rolled 72 pct as austenite at 833°C and then quenched. However, the choice of a composition that transformed almost completely to martensite during the quench to room temperature did not permit direct measurement of the prior austenitic texture. In fact, when the "ideal orientations'' associated with well-known fcc rolling textures were converted, alone or in combination, to martensite according to the Kur-djumov-Sachs (K-s)' or Nishiyama8 relations, the agreement obtained with the observed martensite texture was only fair at best. Recently a pertinent aspect of the austenite to martensite transformation was reported by Bokros and parker,10 who found that certain habit-plane variants of martensite were suppressed by tensile deformation of Fe-31.7 Ni single crystals prior to the necessary subzero cooling. It might be anticipated that the consequences of such preferred transformation are sustained during the formation of martensite in warm-worked austenite that has a well-developed deformation texture. The present investigation was undertaken first to establish more firmly the relation between preferred orientations in plastically deformed austenite and in the resulting martensite, and second to examine the textures for evidence of deformation-induced preferred transformation. EXPERIMENTAL PROCEDURES An alloy containing 0.3 pct C, 12 pct Cr, 6 pct Ni, and the balance iron, was selected because the mar-tensite-start temperature (M,) of about -100°C allowed convenient experimental manipulation of either austenite or martensite at room temperature. Furthermore, this composition can be readily deformed as metastable austenite at moderately elevated temperatures without intervention of appreciable isothermal or athermal decomposition products. The alloy was austenitized at 1150°C, aircooled to 510°C, rolled unidirectionally at this temperature to 93 pct reduction of cross-sectional area, and finally oil-quenched to room temperature. Partial transformation to martensite was accomplished by quenching to -196°C as needed. The rolled stock was reduced in thickness from 0.067 to 0.010 in. by etching in a solution of 5 pct HC1, 45 pct HNO3, and 50 pct water, and further thinned by careful mechanical polishing to maintain the two sides of the sheet parallel within 0.0003 in. After mechanical polishing to 0.005 in., electropolishing in 1:9 perchloric-acetic acid solution produced a final thickness of 0.002 in. The preferred orientations were determined from

Jan 1, 1968

By J. Weijdema, D. N. Dietz

In the conventional underground combustion process (dry combustion) much heat is left behind in the swept formation and goes to rva.rte. Econonmy can be improved by heat recuperation through water injection. This is most advantageous if done at the earliest opportunity before much heat is dissiputed to cap and base rock. Water injected simultaneously with the air will flash to superheated steam, which passes through the combustion front together with the nitrogen from the air. A condensation front traveling up to three times as fast as the combustion front drives out the oil. In this type of wet combustion, the water evaporates before it reaches the combustion zone. The evaporation front travels more slowly than the combustion zone. If so much water is injected that the evaporation front overrun the combustion front, combustion in that spot will be quenched and some unburned fuel will be left behind. Air reacts with the oil farther down-stream where steam temperatures occur; at steam temperature, the air reacts rapidly with the oil. Velocity of the combustion front is increased thereby and is governed essentially by the water-injection rate. In the extreme case of high water-injection rate, a short heat wave of constant length is driven through the formation by water injection. Once this wave has been established, no more heat need be generated than that required to make up the heat losses from the short heat wave; a relatively low rate of air injection will suffice. The feasibility of partially quenched combustion has been confirmed in tube experiments. A heat wave at steam temperature is observed. Chemical analyses of flue gas indicate preferential burning of hydrogen while a carbonaceous residue is left in the formation. Introduction A disadvantage of so-called dry in situ combustion is that air-compression costs are rather high. An air consumption of about 400 std cu m/cu m (400 scf/cu ft) of formation swept is an accepted figure. This high consumption is mostly wasted since much heat is left behind in the depleted oil sand. Methods were investigated for recuperating as much as possible of the heat left behind. This paper deals only with basic principles and is confined mainly to one-dimen- sional flow without lateral heat losses; experiments were conducted in relatively narrow, well insulated tubes. If some water is injected with the air, it will turn to superheated steam in an evaporation front, which should travel behind the combustion front. The steam having passed the combustion front causes a steam drive by a condensation front that can travel faster than the combustion front. The latter needs to travel only part of the distance covered by the oil-displacing condensation front, and thus consumes less air. The water-air ratio would seem limited to that at which cold water overruns the combustion. This limitation was deliberately exceeded considerably in theory and experiments. It was found that combustion is then indeed quenched, but only locally. Farther downstream, the oxygen finds residual oil at steam temperature, which is suficiently high to ensure rapid oxidation. Thus, the combustion front uses only part of the available fuel because it is chased through the formation faster than its normal velocity. No heat is left behind. This new process is called "partially quenched combustion". At the upper limit of the water-air ratio, a small heat slug is moved through the formation by the flow of water and steam. Only a small flow of air is needed since it has only to generate sufficient heat to make up for the lateral heat losses of the short heat slug. Theory Although many factors complicate underground combustion, the processes will be presented in their simplest form. For this reason, one-dimensional flow without lateral heat losses is assumed. Heat conduction in the direction of flow also is disregarded. Under these conditions, dry combustion causes very high temperatures. The heat-carrying capacity of the gas stream is small. Heat generated by oxidation of a residual oil saturation is retained in the sand. The available fuel determines the air requirement and the temperature obtained. Accepting the often-mentioned air consumption of 400 std cu m/cu m (400 scf/cu ft) formation, we calculate a temperature of the swept sand of 1,200C (2,192F) (Fig, I). If water is injected at a modest rate with the air, it will flash to superheated steam upon contact with the heated sand. One cu m (35.31 cu ft) of hot formation will evaporate about 0.5 cu m (17.66 cu ft) of water, and thereafter will accommodate (at an estimated 0.80 saturation and an assumed 0.40 porosity) another 0.3 cu m (10.59 cu ft) of water in cold condition. As long as less than 0.5 + 0.3 = 0.8 cu m (28.25 cu ft) of water is injected for every 400 std cu m (14,125 scf) of air (water-

Jan 1, 1969

By A. C. Fieldner

COKE and by-products have prime importance in the war program. The past year was marked by the construction of new and the rehabilitation of old by-product and beehive ovens and by the increase of production from existing plants to practically 100 per cent capacity. W. T. Brown's paper, "Plan to Improve Blast-Furnace Coke," presented at the Fuels Conference in St. Louis, Sept. 30, showed how the blending of low- and high-volatile coal improved the physical properties of coke and that improvement of the quality and uniformity of coke favored an increase of blast-furnace capacity- a vital need in the war program. Others in the industry have likewise increased production of coke or recovery of by- products by various innovations or changes in operation. Recognition of the value of pooling these developments for the benefit of the entire industry led to organization of a survey of coke- oven plants and blast furnaces by the Bureau of Mines in co-operation with the Solid Fuels Co-ordinator with a view to improving metallurgical coke production and recovery of by-products. This survey is being conducted with the help of an advisory committee consisting of representatives of the Bureau of Mines, the Solid Fuels Co-ordinator, the War Production Board, and the blast-furnace, coke-oven, and coal-mining industries.

Jan 1, 1943

By CARLE R. HAYWARDC

BEFORE discussing this subject it is necessary to define somewhat the meaning of the tern metallurgical.. When I was a student at M. I. T. ore-dressing was not thought of as metallurgy in any sense of the word, but it was considered as more closely allied to mining, enabling the miner to market an otherwise unmarketable product, or at least to market it to better advantage. Progress in ore-dressing has gradually allied it more closely with smelting, for at present the raw material received by the smelter is predominantly concentrates and smelting practice has been profoundly affected by the products of the concentrating mill. It is not surprising therefore that in many schools ore-dressing is frankly considered as metallurgy and although there may be some dissenting opinions it will be considered as such in this discussion. Twenty-five years ago at M. I. T. metallography was taught in a few optional lectures. This subject has developed rapidly and been expanded into the new and increasingly important field of physical metallurgy. To many metallurgists the -term metallurgy means only physical metallurgy. They know little or nothing about smelting and leaching and even less about ore-dressing. Having then broadened the scope of the term metal¬lurgy to cover ore-dressing, smelting, leaching and phys-

Jan 1, 1930

By CHARLES M. A. STINE

IN fighting a war the all-absorbing intent is to win. There is little time to analyze the rush of events or to appraise their consequences beyond the war's end. The united objective is, rightly, success for our arms. Yet under the pressures of a great war there may be compressed scientific, economic and social developments that might have taken many decades to achieve under less urgent conditions. Their effects on our lives and all civilization may be more wide-reaching and lasting than any military conquest. They constitute one of the must imperative incentives to victory. No American, least of all any scientist worthy the name, conceivably could endorse war as a justifiable means to progress. The destruction of life and property wrought by the present war far exceeds the havocs of a century's earthquakes. Most of mankind is burning itself on an altar of paganism. Nonetheless, one fact is inescapable Despite the recurrent malady of war history's overall course is forward. Mankind has the habit of rising phoenix-like from its own ashes to attain greater heights. Progress is immortal.

Jan 1, 1942

By K. M. Rolls, F. W. Reuter, J. Wulff

The superconducting behavior and microstructural characteristics of a nominal Nb-40 wt pct Ti-0.239 wt pct O alloy were studied as a function of ther mo -mechanical processing treatment. Critical current density us applied transverse magnetic field was obtained for 0.010-in.-diam wires at 4.2°Kin steady fields 14 to 110 kG. Both optical metallogvaphy and transmission electron microscopy were used to delineate the micros tructures of the same wires. It wan found that a 1-hr 500°C precipitation heat treatment after cold drawing to final size led to the highest critical current density. Heat treatment at 600°C also led to a high critical current density, but the precipitate differs in kind and form from that at 500°C. The resistire critical field was also found to be sensitive to precipitation heat treatment since the effective composition of the superconducting phase changes. This is discussed in terms of the oxygen in interstitial solid solution. Two types of high-field superconducting wire are at present used in the construction of high-field superconducting solenoids. These types are solid-solution alloy wire such as Nb-Zr and Nb-Ti and composites of the brittle inter metallic compound Nb3Sn. The latter generally have a high super cur rent-carry ing capacity which is difficult to vary if properly made. The supercur rent- carry ing capacity of the former can be varied drastically and often predictably by suitable thermomechanical processing treatments. In general, the critical current density Jc of the solid-solution type of alloy is increased by cold work and by additions of interstitial elements along with aging heat treatments. The imperfections which result are be-iieved to be responsible for the observed increase in Jc. In 1962 Kneip and coworkers1 found that the critical faurrent density of Nb-Zr alloys could be increased by proper heat treatment preceded and followed by cold work. Betterton and coworkers2 using a Nb-25 at. pct Zr alloy found that small additions of oxygen or carbon enhanced the effect of this heat treatment. They suggested that the interstitials present aided precipitation in the alloy, leading to a filamentary structure with superior properties. If the precipitation heat treatment was omitted, interstitial additions had a negligible effect on Jc. wong3 showed that higher heat-treatment temperatures lowered Jc. Walker and co-workers,4 who studied microstructure (by transmission electron microscopy) as well as superconductivity, found that the Jc anisotropy introduced by cold rolling was itself affected by heat treatment. They were unable to clarify the relation between microstructure and critical current density, although evidence of precipitation was indicated. More recent investigation of Nb-Zr alloys,5,6 besides showing that structural defects and fiber ing due to cold work and precipitation serve to raise Jc, also elucidate important optically observable microstructural changes which occur upon precipitation. In these reports, coarsening of the microstructural features was found to decrease Jc. Vetrano and Boom,7 who studied Ti-20.7 at. pct Nb, found that Jc was increased to a maximum by a 415°C, 3-hr heat treatment following quenching from 800°C and cold working. Heat treatments can also affect the resistive critical field Hr. Final-size heat treatments of Nb-Zr wire can lower Hr drastically if gross phase decomposition occurs5'* or moderately if the effects of cold work are eliminated without changing significantly the composition of the phase of interest.3,5,6,8 The percentage of oxygen which can be added to Nb-Zr alloys to enhance Jc is limited by the difficulty of subsequent cold drawing. Since Nb-Ti and Ta-Ti alloys in contrast can tolerate appreciably higher percentages of oxygen, it was decided to investigate the superconducting behavior of various alloys in these systems. The present paper describes the results of adding oxygen to a nominal 40 wt pct Nb alloy as a function of thermomechanical treatment. I) EXPERIMENTAL PROCEDURE A small alloy ingot was prepared from high-purity niobium, iodide, crystal-bar titanium, and Nb2O5 powder by arc melting on a water-cooled copper hearth in a gettered argon atmosphere. The ingot was turned and remelted fourteen times to insure homogeneity. After final melting and rapid cooling, it was machined round to 0.415 in. diam, jacketed in stainless steel, and cold-swaged to 0.117 in. diam. The jacket was removed and swaging continued to 0.051 in. diam followed by wire drawing in carbide dies to 0.010 in. diam. Although it was intended that about 1500 ppm O (by weight) be added, inert gas fusion analysis indicated a 2390 ppm 0 content, apparently due to additional oxygen pickup in the arc furnace. Even so, the alloy was sufficiently ductile to be cold-worked to greater than 99.9 pct reduction

Jan 1, 1967

By L. Zwell, H. A. Wriedt

X-ray Lattice parameter and bulk density measurements were made at room temperature on solid solutions of nitrogen in purified a, iron, containing from 0 to 0.053 wt pct N. When nitrogen is dissolved in a, iron, the fractional increase in unit cell volume (X-ray diffraction) and the fractional increase in volume per gram of iron (bulk density measurement) are the same within experimental error. The two sets of results when averaged show that the unit cell edge of a, iron increases linearly by 3.2 X 10-2 A per wt pct N dissolved and that the partial gram-atomic volume of nitrogen at infinite dilution in a, iron is the same as it is in nitrogen-martensite, that is, 5.90 cm3. RECENT reviews17' of literature dealing with the lattice parameter of ferritic iron-nitrogen alloys indicate that the dilating effect of dissolved nitrogen on the bcc unit cell of a iron is not known with confidence. Five investigations3-7 either revealed no effect of nitrogen or were inconclusive. Among other investigations, where quantitative data were obtained, the results of Osawa and Iwaizumi8 indicate too large an effect, and those of Eisenhut and Kaupp9 show the lattice parameter of ferrite varying with the nitrogen content at concentrations far beyond the known solubility limit. It is not clear in Burns'10 data what value should be used for the lattice parameter of nitrogen-free ferrite in calculating the dilation. Burdese11 shares with Burns the shortcoming that the iron-nitrogen alloys were solution-treated and quenched with undissolved nitride present. The consequence is that the un- certainties in the solubility limit1 are added to those in the lattice parameter measurements when the variation of lattice parameter with nitrogen content is computed. For a system of such widespread interest, we felt that a new investigation was merited. Both X-ray diffraction and bulk density measurements were used to determine the dilation of the , a iron lattice by dissolved nitrogen. EXPERIMENTAL PROCEDURE Material. The bulk density and X-ray diffraction experiments were both made with specimens of a purified iron (denoted as "B iron")- from Battelle Memorial Institute. The impurity contents were as follows: As, C, Cu and Pb—50 ppm; Ag, Al, Co, Ga, 0, Ti-11 to 50 ppm; B, Be, Ca, Cd, Cr, Mg, Mn, Mo, N, Ni, S, Sb, Si, Sn, V, W-10 ppm or less. The vacuum-cast ingot was hot-forged and annealed. Pieces of the forged stock were machined and cold-rolled to give strip stock 0.08 and 0.002 in. thick for density and for X-ray measurements, respectively. Preparation of Specimens. Starting with a piece of the cold-rolled iron stock, the general sequence of operations, after an initial recrystallizing anneal, was to measure the density or lattice parameter, nitrogenize, remeasure the property, age or denitrogenize, and again measure the property. Duplicate analyses for nitrogen were made by a modified Kjeldahl method on portions of the actual density specimens and on pieces nitrogenized to equilibrium alongside the X-ray specimens. Annealing of the density specimens was performed by sealing the pieces with the as-rolled surfaces removed by filing into evacuated "Vycor" capsules and treating for 1 hr at 750°C followed by furnace cooling. (TO check the effect, if any, of annealing procedure, a specimen also was

Jan 1, 1962

By C. S. Barrett

The torsional after-effect in polycrystalline cadmium is interrupted by an abnormal twisting when the film is removed by etching. This is accounted for by the pile-up of dislocations beneath anodic or other oxide films during the original twisting, and the relief of the residual stresses from these by elastic and plastic processes during etching. WHEN an iron or zinc wire coated with oxide is plastically twisted and then released, it untwists according to a law of equal increments in equal intervals of log time, but if the oxide is suddenly removed by etching this normal after-effect is interrupted by a transient, and thereafter the untwisting is slower.' Etching may even reverse the strain temporarily. Related effects caused by hydrated oxide films on cadmium in creep tests have been observed by Phillips and Thompson:' definite transient strains caused by the removal of the film from single crystals of cadmium were observed; however, negligible effects were obtained with polycrystalline wires. Pronounced effects of both transient and steady nature were observed by Phillips and Thompson when anodic films 10-4 cm thick were etched off, and weak effects, transient only, occurred with hydrated oxide films 3x10 cm thick such as were formed by standing in water for 10 min. It was estimated that 10.' cm films, formed in water in 1 or 2 min, could be detected in the apparatus of Phillips and Thompson. Since abnormal after-effects were observed by the torsion method equally well with polycrystalline and monocrystalline iron and zinc wires,' it was anticipated that the torsion method would also disclose them in polycrystalline cadmium, in spite of the fact that analogous effects had not been observed in tensile creep tests on polycrystalline zinc by Pickus and Parker3 r on polycrystalline cadmium in tensile studies by Andrade and Randall." Because of its high sensitivity, particularly its sensitivity to surface conditions, the torsion method was thought to merit further study as a method of investigating thin films of various types and the action of various reagents on these films. Materials and Methods Cadmium wires were cold-drawn to a diameter of 0.039 in. from a stock that was determined spectro-scopically to contain 0.05 pct Bi, 0.005 pct Cu, 0.008 pct Pb, and 0.007 pct Si. The wires were cut to about 17/8 in. lengths and were annealed 1 hr in air at 290°C, after which the grain size was such that there were 5 to 7 grains across a diameter. The wires were then mounted as indicated in the insert in Fig. 1, one end being fastened with laboratory wax into a small diameter brass tube and the other end being waxed to a thin glass fiber which supported, on its upper end, a small galvanometer mirror. The wires were waxed with care to avoid straining or annealing, for it was found that the after-effect rate was sensitive to prior cold work in a wire. The assembly was then clamped above a beaker into which the reagent could be poured. The lower end of the wire was held in the hand while the wire was twisted 180"; it was then released immediately. Since precision timing of these manipulations was not attempted, the different wires cannot be compared as to exact rates of untwisting. The length subjected to plastic strain was 15/8 in. The surface coatings studied were those produced by the annealing in air, by anodic treatment of

Jan 1, 1954

By T. R. Ingraham

The thermal decomposition of cupric sulfate and of cupric oxysulfate has been examined by deter-mining the equilibrium gas pressure generated over each pure compound. The equilibrium data have been used to calculate the thermodynamic properties of both compounds, and the results have been combined with previously published data to establish a predominance-volume diagram for the Cu-S-O system over the normal range of roasting temperature and gas composition used in the treatment of copper minerals. Copper oxysulfate recrystal-lizes at 1100°K with an apparently irreversible en-dothermic heat requirement of 2.9 * 0.3 kcal per mole. WHEN mixed sulfide minerals containing copper are roasted, either of two procedures may be followed, depending on the nature of the other metals. In the first procedure, the roasting is controlled to produce a maximum amount of water-soluble copper sulfate which may be removed selectively from the water-insoluble compounds of the other metals. In the second, control is directed toward the production of water-soluble sulfates of the other metals and an insoluble salt of copper. Effective control of either of these processes would be enhanced by knowledge of the thermal stability of cupric sulfate and of cupric oxysulfate. In this paper, established data for the thermal stability of cuprous and cupric sulfide and of cuprous and cupric oxide will be combined with new data for the therma1 stability of cupric sulfate and cupric oxysulfate, to designate the volumes of stability of the significant compounds in the Cu-S-O system in the temperature range 750° to 950°K. EXPERIMENTAL Materials, Apparatus, and Procedure. The copper source material used in all experiments was Fisher Certified Reagent Grade "anhydrous" cupric sulfate, for which the following analysis was supplied by the manufacturer: chloride, 0.002 pct; alkalies and earths, 0.2 pct; insoluble, 0.002 pct; iron, 0.020 pct. Cupric oxysulfate was prepared from the normal sulfate by two procedures. In the first, equimolar quantities of cupric oxide (prepared by decomposition of the sulfate) and cupric sulfate were ground together and heated under vacuum in a sealed silica vessel at 800°C. This procedure was subject to failure because of explosion of the silica vessels due to the pressure of the final traces of water released from the sulfate on the initiation of its decomposition. The second procedure involved roasting the cupric sulfate, with frequent rabbling, in a muffle furnace at 725°C for a period of 48 hr. This produced an excellent product, which was confirmed as copper oxysulfate by X-ray and chemical analyses. The experimental technique used was that of measuring the total pressure exerted by the sulfur trioxide, sulfur dioxide, and oxygen produced by the following reactions: 2 CuS04 = cuo • CuSO4 + SO3 [1] cuo . CuSO4 = 2 cuo + SO3 [2] SO3= SO2 + 1/2 O2 [3] The apparatus used in the experiments has been described previously.1-3 Its main feature of a U-tube mercury manometer in which the mercury is protected from corrosion by enclosure in a flexible Pyrex bellows. Before beginning a run, the interior of the apparatus, including the bellows-manometer and all connecting tubing, was thoroughly washed with hydrochloric acid and distilled water to remove any

Jan 1, 1965

By Howard Martens

HARDNESS behavior of commercial titanium alloys following various heat treating processes has been studied for some time. However, the hardness of such alloys following a definite measured cooling rate from the single phase region has not been reported. Therefore, a series of experiments was conducted on the three alloys: 150A, 130A, and 130B. The manufacturers gave the following compositions for these alloys: Ti 150A: 0.045 pct C, 0.069 pctN, 1.32 pct Fe, and 2.68 pct Cr; RC 130A: 0.13 pct C, and 7.9 pct Mn: RC 130B: 0.11 pct C, 3.5 pct Mn, and 3.2 pct Al. A method similar to that described by Greninger was used, and the cooling rate varied from 6" to 9200°F per sec. The cooling curves were recorded on a Speedomax or a recording oscillograph, and the hardness was measured on a Tukon hardness tester using a Knoop indentor with a 500 g load. The curves of Fig. 1 show the variation of the Knoop hardness with the cooling rate for these three alloys. The alloys 150A and 130B behave in a similar manner, as shown by the curves. In neither of these alloys was it; possible to retain the high temperature ß phases, even by the most rapid cooling rate. The slow cooling rates produced a typical ß structure in these two alloys. As the cooling was increased, these phalses became more finely dispersed and the hardness increased. This increase in hardness with cooling rate continued for the 150A alloy until a rate of approximately 1000°F per sec was reached. At this cooling rate, the structure showed the result of the dif-fusionless transformation of the ß solid solution to the supersaturated a solid solution which is referred as a'. At higher cooling rates, the structure showed no change and the hardness showed no great change, as indicated by the dotted portion of the curve in Fig. 1. The increase in hardness with cooling rate for the 130B alloy continued until a rate of approximately 300°F per sec was reached. At this rate the structure was produced and the maximum hardness was reached. Higher cooling rates caused no marked changes in hardness or structure. The alloy 130A behaved in a different manner, as shown by the curve of Fig. 1. The structure produced by the slow cooling rates was a typical a+ß structure, which became finely dispersed as the cooling rate increased. The hardness increased with cooling rate and reached a maximum at approximately 120°F per sec. At this maximum hardness, the structure showed the result of the decomposition of the ß phase into the ß + phase structure. As the cooling rate was increased, smaller amounts of the at phase were formed and the alloy became softer. For cooling rates of 3000°F per sec and above, the high temperature ß phase was retained and the hardness of the alloy did not change markedly. Acknowledgment This paper presents the results of one phase of research carried out at the Jet Propulsion Laboratory, California Institute of Technology, under Contract No. DA-04-495-0rd 18, sponsored by the Dept. of the Army, Ordnance Corps.

Jan 1, 1957