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By R. E. Shinkosk

This paper outlines the development of a program for increasing the life of woolen bags used for filtering Dwight-Lloyd gases by treating the bags and gases with hydrated lime. Methods and apparatus are described for determining alkalinity of dusts, acidity and breaking strength of bag cloth. Procedure and results, based on several years of operation, are presented. DURING 1939, additional facilities were constructed in the Dwight-Lloyd Blast Furnace and Baghouse departments at the Selby, California, Plant of the American Smelting and Refining Co. In order to handle adequately the increased volume of gases from the resultant increase in production, it was necessary to increase gradually the amount of water used for cooling gases ahead of the sinter machine baghouse. As a result of this increased water cooling, the average bag life dropped from 27 months in 1939 to 14 months in 1941. (Table I). This drop in life meant an increased. bag cost, as well as lower recovery of dust and some curtailment of operation. During 1941, it was found new bags showed as high as 0.3 pct acidity* after two weeks of opera- tion and as much as 2.0 pct acidity after some months of operation. This high acidity was present in spite of the fact that free oxide or relative alkalinity of the unburned dust ran from 5 to 6 pct. In view of these circumstances, a twofold program was started in Nov. 1941.t Part one of this program consisted of vigorously dipping all new bags in a weak lime solution, containing 50 lb of hydrated lime per 50 gal of water. Part two consisted of feeding fine, dry, hydrated lime into the gas stream intake of the sinter baghouse fan. Apparatus for feeding this lime is shown in fig. 1. All baghouse chambers are shaken in rotation about once each hour. On alternate hours, the baghouse operator places 50 lb of hydrated lime (one sack) into the lime feeder, starts feeder and immediately starts the bag shaking machinery. The rate at which lime is fed is set to coincide with the approximate time necessary to shake all sinter bag-house chambers, or about 15 min. It is felt this method of lime addition is most effective for getting lime into the woolen bag fabric. The amount of lime so fed averages about 600 lb per day. The amount of lime fed per day is varied to keep a minimum relative alkalinity of 9 pct in the unburned sinter dust. A daily dust sample is taken for alkalinity by allowing dust to accumulate in a sample pipe over a 24-hr period. This sample pipe, placed in any chamber cellar, is 2 in. in diam, 4 ft long, is sealed on the inner end, and capped on the outer end. It has a 1/2 in. slot cut for 18 in. along the tip end. This slot faces upward and allows the pipe to fill gradually with dust as bags are shaken. Breaking strength of bags has, in most cases, been the deciding factor in bag replacement. Bags that normally test 100 psi breaking strength when new are replaced when they test under 35 lb. The method for determining breaking strength is shown in the description accompanying fig. 2. Since the start of the liming program in 1941, bag life has increased from 14 months to an average of over 23 months, with a consequent material decrease in bag cost per year. Acidity, as per cent sulphuric acid, may be determined by means of a Beckman pH meter as follows: From a piece of bag cloth. which has been thoroughly cleaned of dust, a 5 g sample is weighed on a balance. Cut the sample into fine pieces and place in a 400 cc beaker. Add 100 cc (measured) of distilled water and stir vigorously. Filter on suction funnel, holding cloth pulp in beaker with a stirring rod. Wash cloth sample and filter wash water four additional times, each time with 20 cc distilled water, the last time squeezing cloth pulp over funnel. Discard pulp and rinse funnel and filter paper. Pour wash solution jnto measuring graduate and make up to exactly 300 cc with distilled water. Place into clean 600 cc beaker and measure the pH on meter. The per cent acid in bag cloth is read from the following table:—

Jan 1, 1951

By James K. Hallenburg

INTRODUCTION The delayed fission neutron, or DFN technique for uranium ore body analysis uses the first down-hole method for detecting uranium in place quantitatively. This technique detects the presence of and measures the amount of uranium in the formation. DFN TECHNIQUE DESCRIPTION The DFN technique depends upon inducing a fission reaction in the formation uranium with neutrons, resulting in an anomalous and quantitative return of neutrons from the uranium. Since there are no free, natural neutrons in formation, a good, low noise assessment may be made. There are several methods available for determining uranium quantity in situ. The method used by Century uses an electrical source of neutrons. This is a linear accelerator which bombards a tritium target with high velocity deuterium ions. The resulting reaction emits high energy neutrons which diffuse into the surrounding formation. They lose most of their energy until they come to thermal equilibrium with the formation. Upon encountering a fissile material, such as uranium, these thermal neutrons will react with the material. These reactions produce additional neutrons, the number of which is a function of the number of original neutrons and the amount of fissile material exposed. The particular source used, the linear accelerator, has several distinct advantages over other types of sources: 1. It can be turned off. Thus, it does not constitute a radioactive hazard when it is not in use. 2. It can be gated on in short bursts (6 to 8 microseconds). This results in measurements free of a high background of primary neutrons. 3. The output can be controlled. Thus, the neutron output can be made the same in a number of tools, easily and automatically. There are several interesting reactions which take place during the lifetime of the neutrons around the source. During the slowing down or moderating process the neutron can react with several elements. One of these is oxygen 17. This results in a background level of neutrons in any of the measurements which must be accounted for in any interpretation technique. These elements are usually uninteresting economically. The high energy neutrons will also react with uranium 238. However, the proportions of uranium 235 and 238 are nearly constant. Therefore, this reaction aids detection of uranium mineral and need not be seperated out. Upon reaching thermal energy the neutrons will react with any fissile material, uranium 235, uranium 234, and thorium 232. At present, we do not have good techniques for seperating out the reaction products of uranium 234 and thorium 232. However, uranium 234 is a small (.0055%) percentage of the uranium mineral and thorium 232 is usually not present in sedimentary deposits. When the uranium 235 reacts with thermal neutrons it breaks into two or more fragments and some neutrons. This occurs within a few microseconds after the primary neutrons have moderated and is the prompt reaction. One system uses this; the PFN or prompt fission neutron technique. We don't use this method because the neutron population is low and, therefore, the signal is small and difficult to work with, accurately. Within a few microseconds to several seconds the fission fragments also decay with the emmission of additional neutrons. Now, with a long time period available and a large neutron population we gate off the generator and measure the delayed fission neutrons after a waiting period. These neutrons can be a measure of the amount of uranium present around the probe. Thermal neutrons are detected with the DFN technique instead of capture gamma rays to avoid some of the returns from other elements than uranium. LOGGING TECHNIQUE The exact logging technique will depend, to some extent, upon the purpose of the measurement. However, the general technique is to first run the standard logs. These will include: 1. The gamma ray log for initial evaluation of the mineral body and for determining the position of the borehole within the mineral body, 2. The resistance or resistivity log for determining the formation quality, lithology, and porosity. 3. The S. P. curve for estimating the redox state and shale content, and measuring formation water salinity, 4. The hole deviation for locating the position, depth, and thickness of the mineral (and other formations), and 5. The neutron porosity curve. The neutron porosity curve is most important to the interpretation of the DFN readings. The neutrons from this tool are affected in the same way by bore hole and formation fluids as the DFN neutrons are. Therefore, we can use this curve to determine effect of the oxygen 17 in the water. Of course, this curve can be used to determine formation porosity. It can also be used to calculate formation density.

Jan 1, 1979

By John C. Gifford, Charles L. Thomas

A unique and reproducible method is presented for the determination of water vapor in tough pitch wire bar copper. The procedure involves reduction of the water vapor with molten aluminum to form hydrogen, which is subsequently measured by mass spectroscopy. Average water vapor pressures within the porosities of the wire bar samples are calculated. Correlation is to exist between the specific gravities of the samples and their measured water vapor contents. The method should find application as a very sensitive means of detecting hydrogen embrittlement in copper. The nature and quantity of gases evolved and retained during the horizontal casting of tough pitch wire bar copper have long been of interest to the metallurgist. Considerable work has been done at this laboratory on the determination of these gases. The work has involved not only qualitative but also quantitative analysis, so as to provide a basis for a total accounting of the porosity which is associated with the cast product. From a knowledge of the gas-forming elements within the copper, and the practice of melting and protecting it with a reducing flame followed by contact with a charcoal cover in the casting ladle, the gases which one might expect to find in the pores of the cast product are sulfur dioxide, carbon monoxide, carbon dioxide, hydrogen, and water vapor. Hydrogen sulfide, nitrogen, and hydrocarbons would be other possibilities; however vacuum fusion-mass spectroscopy techniques employed at this laboratory have shown that no hydrogen sulfide and only traces of nitrogen and methane are present. It is highly improbable according to phillipsl that any sulfur dioxide could be evolved in wire bar copper with 10 ppm or less sulfur under normal freezing conditions. Mackay and smith2 have noted that porosity due to sulfur dioxide only becomes noticeable at concentrations above 20 ppm S. Investigation of carbon monoxide and carbon dioxide by a variation in the method of Bever and Floe showed that these two gases could only account, at 760 mm and 1064°C (Cu-Cua eutectic temperature), for a maximum of about 25 pct of the total porosity in a wire bar having a specific gravity of 8.40 g per cu cm. phillips' has noted that no normal furnace atmosphere is ever sufficiently rich in hydrogen to cause porosity in copper from hydrogen alone. In addition, using a hot vacuum extraction technique for hydrogen,4 values have never been observed in excess of 10 ppb in tough pitch wire bar. On the basis of the preceding considerations of gases in tough pitch wire bar, only water vapor is left to account for the major portion of the porosity. Direct determinations of water vapor are virtually impossible at low concentrations by any presently known technique, due to adsorption and desorption within the walls of the apparatus used.5 The present investigation deals with a method for the determination of water vapor by an indirect procedure, using molten aluminum as a reducing agent to form hydrogen according to the reaction: 2A1 + 3H2O — A12O3 + 3H2 The evolved hydrogen can then be measured quantitatively by mass spectroscopy. EXPERIMENTAL A 10-g piece of 99.9+ pct A1 was charged into a porous alumina crucible (Laboratory Equipment Co., No. 528-30). Fig. 1 shows the crucible in place at the bottom of an 8-in.-long quartz thimble. A funnel tube with two l1/8-in.-OD sidearms extending at a 90-deg angle from each other was attached to the top of the thimble. One of the sidearms was joined to the inlet system of the mass spectrometer (Consolidated Electrodynamics Corp. Model 21-620A) via a mercury diffusion pump situated between two dry-ice traps. The copper samples were placed in the other sidearm, followed by a glass-enclosed magnetic stirring bar for pushing the samples into the crucible. All ground joints were sealed with vacuum-grade wax. The entire system was evacuated and the aluminum was heated with a T-2.5 Lepel High Frequency Induction Furnace for 21/2 hr at a temperature visually estimated to be 900°C. The temperature was then lowered and the hydrogen was monitored on the mass spectrometer until it was given off at a constant rate of about 4 to 5 1 per hr. This rate corresponded to a slope of 2 to 3 divisions per min on the X3 attenuation of a 10-mv recorder at a hydrogen sensitivity of approximately 100 divisions per 1. A micromanometer (Consolidated Electrodynamics Corp. Model 23-105)

Jan 1, 1969

By C. G. Dunn, M. Sharp

IN twinned crystals of the face-centered cubic metals the lattice of one twin is a mirror image of the other in a common twin boundary. When several twins appear within large grain in a sheet specimen, the twin one boundaries form a set of lines at the surface of the specimen which coincide with (111) planes of the large grain. Furthermore, for twins of the same orientation, these lines are parallel. Generally, the presence of identically oriented regions with straight parallel boundaries coinciding with a (111) plane of the surrounding crystal is strong evidence for identifying the island regions as twins of the parent crystal. However, Fig. 1, which shows the macrostructure of a large grain of copper with island regions that satisfy these conditions. is not an illustration of (111) twins. Since the reverse side of the specimen has much the same appearance, it was thought at first that these regions, which appear dark in the macrograph, actually were twins. According to X-ray data, however, these regions are second-order twins of the large crystal. With regard to their formation, these second-order twins formed by secondary recrystallization in a cube texture matrix. Growth occurred in the direction of the arrow (see Fig. 1) as the specimen moved slowly into a gradient temperature furnace as described previously.' Nucleation of the second-order twins occurred, therefore, on the ends facing opposite the arrow. If the origin of the second-order twins were due to repeated twinning, some first-order twin structure should be visible on these ends. This proved to be the case, as very small twins were readily found with the aid of a microscope, and probably could have been seen, in some instances, under ideal lighting conditions without aid of a microscope. Fig. 2 shows a cross-section view taken perpendicular to both the surface and the (111) trace of the parent crystal (visible as a straight boundary in Fig. 1) at the beginning point of growth of a second-order twin and where one first-order twin was relatively thick. In the micrograph, A is the large parent grain; B is the first-order twin of A; and C, which is a first-order twin of B; is a second-order twin of A. Between A and B and between B and C the major straight portions are traces of common (111) twin boundaries. The straight portion of boundary between A and C, however, is not a common crystallographic plane to the two lattices; it is a (111) plane of A and a (115) plane of C. Without considering the mechanism of twinning itself, the origin of the second-order twins may be accounted for in terms of repeated twinning and special growth characteristics. After each nucleation, a selective growth process can be thought of as favoring growth of the first-order twin in local spots only and favoring growth of the second-order twin to an extent comparable with that of the parent grain over relatively large areas in a way similar to that described for twinning in aluminum.' It has already been pointed out that the boundary between the large grain (A) and the second-order twin (C), which is responsible for the straight boundary portions in Fig. 1, involves a (111) plane of A and a (115) plane of C. The same combination of planes is not only possible in first-order twins, but actually appears quite frequently.3 Their prevalence in first-order twins and their presence here in second-order twins, together with the necessary occurrence of a large number of common lattice sites at the boundary, is an indication that this combination produces an "energy cusp"' boundary. (Energy cusp boundaries have been described by Shockley and Read.") The configuration of atoms near a {Ill), (115) boundary in first-order twins is of course different from the configuration near the same type of boundary in second-order twins. References 1 M. Sharp and C. G. Dunn: Secondary Recrystallization Texture in Copper. Journal of Metals (January 1952) Trans. AIME, p. 42. 2W. G. Burgers and W. May: Stimulated Crystals and Twinning in Recrystallized Aluminum. Recueil des travaux chimiques des Pays-Bas (1945) 64, p. 5. aD. Whitwham, M. Mouflard, and P. Lacombe: Discussion of W. C. Ellis and R. G. Treuting, "Atomic Relationships in the Cubic Twinned State." Trans. AIME (1951) 191, p. 1070; Journal of Metals (October 1951). 4 W. Shockley and W. T. Read: Dislocation Models of Crystal Grain Boundaries. Physical Review (1950) 78, p. 275.

Jan 1, 1953

By I. R. Kramer, L. J. Demer

T. H. Alden and R. L. Fleischer (General Electric Research Laboratory)— The authors' results indicate clearly and, we believe, significantly that during tensile deformation the surface layers of an aluminum crystal are hardened more severely than the interior of the crystal. A probable explanation of this effect, as the authors indicate, is that dislocations in the primary slip system may be obstructed at the surface or, it should be added, near the surface. The intent of this discussion is to show that the oxide film on aluminum is not likely to be responsible for this effect, but that the results can be understood if it is assumed the secondary slip is more active in the surface layers than in the interior. Prior study has shown that the principal mechanical effect of an oxide film on a single crystal is to raise the yield stress while leaving the rate of strain hardening during the initial deformation relatively unaffected.33 Since the yield stress is unchanged during polishing in the present case, we conclude that continual removal of the oxide film exerts a small effect on the plastic hardening.* It appears that the hardening interactions are occurring not only at the immediate surface, but to an appreciable depth below it, although with decreasing severity. For example, Kramer and Demer found that with removal of 0.004 in. from a specimen, the easy glide region was extended somewhat; but the yield stress did not decrease. The initial yield stress was recovered only after 0.041 in. was removed. Since a very brief polish would permit dislocations trapped behind a surface film to run out,34 extra dislocations must, instead, be trapped to a considerable depth below the surface. The same conclusion is drawn from the observation of decreasing hardening slope with increasing surface removal rates. If the hardening interactions were only at the immediate surface, a full softening effect would be observed at some small removal rate. The view is taken here that strain hardening is principally caused by small amounts of secondary slip.35 The secondary dislocations will interact in various ways with the primaries, interfering with their motion and causing them to accumulate in the crystal. Prior studies of easy glide have shown Diehl's model of hardening to be qualitatively consistent with the effects of impurities,36 of temperature,36 and of crystal size.37 On this basis the enhanced hardening of the surface layers in aluminum arises from increased secondary slip at and to some depth below the surface. Selective removal of this hardened layer is expected to decrease the measurable "bulk" hardening, the effect increasing with the removal rate and decreasing with the applied strain rate. We suggest that the stress on secondary systems is raised by the bending moments arising from interactions with the grips during the deformation. This stress from the grips has been shown to be a maximum37 near the surface, and hence, increased secondary slip should result. Prior investigations of grip effect:; indicate that as the grip stresses are raised by changing the crystal shape, the easy glide slope increases while the extent of easy glide decreases.38-40 It has been shown also that bending moments superimposed during tensile testing may either decrease easy glide, when supporting the moments caused by gripping, or increase it, when cancelling the gripping moments.38 This interpretation of the authors' results, emphasizing the special importance of secondary slip near the surface, is also consistent with the earlier results of Rosi.41 Copper crystals alloyed with silver in the surface layer show greatly increased easy glide compared with pure copper. In addition, the easy glide slope is reduced. The effect of bulk alloying in extending easy glide has been well established and has been interpreted as indicating the relative difficulty of secondary slip in alloy crystals. Since non-basal glide is difficult in zinc crystals, the effects of surface removal during deformation may be less important. Experiments to test this idea are in progress. I. R. Kramer and L. J. Demer (authors' reply)—The authors wish to thank Dr. Alden and Dr. Fleischer for their discussion. Our interpretation of the data in the paper is that dislocation motion is obstructed by "debris" which starts to form at the surface and extends towards the interior of the crystal with further plastic deformation. The fact that we did not find a reversion from Stage II to Stage I by surface removal shows that in Stage II the "debris" fills the entire cross-section of the specimen. Drs. Alden and Fleischer take the view that bending stresses due to the grips are responsible for the

Jan 1, 1962

By Thomas R. Mager

An in.t!estigation has been made of the equilibrium conditions at 1200°C in the reaction between hydrogen sulfide gas and sulfur dissolved in Fe-Si alloys From this the equilibrium constant, activity coefficient, and activity of sulfur in solution were calculated. A number of studies of the equilibrium of sulfur with iron and iron alloys have given closely agreeing results from which the activity and free energy of the dissolved sulfur may be found. Sherman, El-vander, and chipman1 discussed the significant researches of dilute solutions of sulfur in liquid iron prior to 1950, and the results of this study indicated that the relationship between the ratio of PH2S/PH2 in the environment and the percentage of sulfur in solution is not a linear one. Morris and williams2 studied the equilibrium conditions in the reaction between hydrogen sulfide gas and sulfur dissolved in liquid iron and Fe-Si alloys, and reported that silicon dissolved in iron has a pronounced effect on the equilibrium conditions. They found that the activity of sulfur in iron is increased by the addition of silicon. At a silicon content of 4 pet the activity coefficient of sulfur was about twice that for sulfur dissolved in pure iron. Sherman and chipman3 investigated the chemical behavior of sulfur in liquid iron at 1600°C through the study of the equilibrium: H2 + S = H2S; K = PH2S/PH2 . 1/as [1] From the known equilibrium constant of the reaction between H2, H2S, and S and the experimental data, the activity of sulfur in the melt was determined. They found that the activity coefficient of sulfur defined as fs = as/%s is increased by silicon and decreased by manganese. Morris4 and Turkdogan5 also reported that manganese decreases the activity coefficient of sulfur in liquid iron and iron-base alloys. A recent technique of sulfur analysis developed by Kriege and wolfe6 of the Westinghouse Research Laboratories permits an accurate sulfur analysis of 0.5 * 0.2 ppm in the range of 0.1 to 3 ppm, whereas in the range of 3 to 50 ppm the accuracy is ±1 ppm. This technique of sulfur analysis was utilized in this experiment. Previous unpublished data reported that sulfur analysis by the combustion technique was not accurate below 20 ppm. EXPERIMENTAL PROCEDURE Five 5-lb ingots of high-purity Fe-Si were prepared. Three of these ingots were prepared without the addition of manganese but with a variation of silicon contents from 2 to 4 pet. The remaining two ingots contained 3 pet Si with the addition of manganese. Ingots were made at each of three silicon levels: 2, 3, and 4 pet. No alloys were made with less than 2 pet Si since below approximately 1.8 pet Si the binary alloy exhibits a to ? transformation. The two additional ingots of 3 pet Si-Fe were made at each of two manganese levels: 0.20 and 0.50 pet. To minimize the effects, if any, of impurities on the activity of sulfur on Si-Fe, the best metals available were used for melting. All ingots were vacuum-melted in magnesium oxide crucibles. After obtaining samples for chemical analyses, the ingots were processed. This consisted of hot rolling and subsequently cold rolling the alloys. Each ingot was hot-rolled at 1000°C, reheating between every pass to minimize grain growth. All heating was done in a protective argon atmosphere. The slabs were hot-rolled to strips 50 mils thick. After hot rolling, all the material was pickled to remove the scale formed on the surface of the strip during hot rolling. The material was then cold-rolled to 12-mil strips. Single strips of the material used in this experiment were hydrogen-annealed at 1200°C for 16 hr in an alumina tube. Chemical analyses of strips M-1, M-3, M-4, M-7, and M-8 are given in Table I. Sulfur, silicon, and manganese analyses were made from the millings from the cold-rolled 12-mil strips. The oxygen analyses were made from slugs of the as-cast material. The hydrogen sulfide used in these experiments was supplied from cylinders containing a mixture of argon and 1 pet hydrogen sulfide. The parts per million of hydrogen sulfide were determined from the analysis of the exit gas of the annealing furnace during each anneal. The flow rate of hydrogen was approximately 1 liter per min in all anneals. The

Jan 1, 1964

By J. D. Verhoeven

In the past 15 years a considerable amount of experimental and theoretical work has been done concerning the onset of convection in liquids as a result of interm1 density gradients. This work, which has been doue in many different fields, is reviewed here and extended slightly to give a rrlore quantitative understanding to the probletrz of conzection in liquid metal dlffusion experinletzts. In liquid metal systems the capillary reservoir technique is currently used, almost exclusively, to measure diffusion coefficients. In this technique it is necessary that the liquid be stagnant in order to avoid mixing by means of convection currents. Convective mixing may result from: 1) convection produced as a result of the initial immersion of the capillary; 2) convection produced in the region of the capillary mouth as the result of the stirring frequency used to avoid solute buildup in the reservoir near the capillary mouth; 3) convection produced during solidification as a result of the volume change; and 4) convection produced as a result of local density differences within the liquid in the capillary. The first three types of convection have been discussed elsewhere1-a and are only mentioned for completeness here. This work is concerned only with the fourth type of convection. Local density differences will arise within the liquid as a result of either a temperature gradient or a concentration gradient. It is usually, but not always, recognized by those employing the capillary reservoir technique that the top of the capillary should be kept slightly hotter than the bottom and that the light element should be made to migrate downward in order to avoid convection. In the past 15 years a considerable amount of work, both theoretical and experimental, has been done in a number of different fields which bear on this problem. This work is reviewed here and extended slightly in an effort to give a more quantitative understanding of the convective motion produced in vertical capillaries by local density differences. The Stokes-Navier equations for an incompressible fluid of constant viscosity in a gravitational field may be written as: %L + (v?)v = - ?£ + Wv - g£ [1] where F is the velocity, t the time, P the pressure, p the density, v the kinematic viscosity, g the gravitational acceleration, and k a unit vector in the vertical direction. A successful diffusion experiment requires the liquid to be motionless, and under this condition Eq. [I] becomes: where a is the thermal expansion coefficient [a =-(l/po)(dp/d)], a' is a solute expansion coefficient [a' = -(l/po)(dp/d)], and the solute is taken as that component which makes a' a positive number. Combining with Eq. [3] the following restriction is obtained: Since there is no fixed relation between VT and VC in a binary diffusion experiment, Eq. [5] shows that the condition of fluid motionlessness requires both the temperature gradient and the concentration gradient to be vertically directed. Given this condition of a density gradient in the vertical direction only, it is obvious that, as this vertical density gradient increases from negative to positive values, the motionless liquid will eventually become unstable and convective movement will begin. The classical treatment of this type of instability problem was given by aleih' in 1916 for the case of a thin fluid film of infinite horizontal extent; and a very comprehensive text has recently been written on the subject by handrasekhar.' It is found that convective motion does not begin until a dimensionless number involving the density gradient exceeds a certain critical value. This dimensionless number is generally referred to as the Rayleigh number, R, and it is equal to the product of the Prandtl and Grashof numbers. For the sake of clarity a distinction will be made between two types of free convection produced by internal density gradients. In the first case a density gradient is present in the vertical direction only, and, since the convection begins only after a critical gradient is attained, this case will be called threshold convection. In the second case a horizontal density gradient is present and in this case a finite convection velocity is produced by a finite density gradient so that it will be termed thresholdless convection. Some experimentalists have performed diffusion experiments using capillaries which were placed in a horizontal or inclined position in order to avoid convection. These positions do put the small capillary dimension in the vertical direction and, consequently, they would be less prone to threshold convection than the vertical position. However, if the diffusion process produced a density variation, as it usually does, it would not be theoretically possible to avoid thresh-

Jan 1, 1969

By E. Teghtsoonian, E. M. Schulson

The tensile behavior of bcc ordered P' AuZn single crystals (CsCl structure) has been investigated under varying conditions of temperature, composition, and orientation. Between -0.2 and 0.4 T, multi-stage hardening occurs fm stoichiometric and nonstoichio-metric crystals oriented near the middle of the primary stereographic triangle. At higher and lower temperatures, parabolic type hardening occurs, followed by work - softening at the higher temperatwes. Deviations from stoichiometry give rise to increased flow stresses. Multi-stage hardening was observed for most orientations, except along the [loll-[lll] boundary and near the [001] corner of the stereo -graphic triangle, where parabolic type hardening occurs. Along two slip systems, (hk0)[001] and (, operate simultaneously while in the [001] comer, slip occurs mainly on the system. Electron microscopy of deformed crystals revealed bundles of edge dislocations forming walls approximately Perpendicular to the glide plane. In general the plasticity of 4' AuZn closely resembles the plasticity of bcc crystals. In recent years, considerable interest has arisen concerning the mechanical properties of the CsCl type intermetallic compounds Ag Mg,'- Fe co,' and Ni Al.'-' The compound P'AuZn is structurally similar. It has a low and congruent melting point of 725"~,'" remains ordered up to the melting point,16 and pos-esses a range of solid solubility from 47.5 to 52.0 at. pct Au at room temperature.15 The present paper reports the results of an investigation on the general tensile behavior of material in single crystal form. Some dislocation configurations characteristic of the deformed state are also reported. The results of a detailed study of the slip geometry in AuZn are presented in a separate paper.17 PROCEDURE Alloy preparation, crystal growing techniques, and the procedure followed in selecting specimens of minimum composition variation are reported elsewhere.17 Dumb-bell shaped tensile specimens were prepared by carefully machining single crystals in a jewellers' lathe to a gage length of 0.80 in. and diam of 0.090 in. Back-reflection Laue X-ray patterns and room temperature tensile tests revealed that machining damage could be eliminated by electrochemically polishing 0.005 in. from the machined surface followed by annealing at 300°C for 1 hr. Specimens were polished in fresh 5 pct KCN solution (40°C, 12 v). Experiments were performed by gripping specimens in a self-aligning pin-chuck and threaded collet system, then straining in a floor model Instron tensile machine. All tests were performed in duplicate. Experimental variables included temperature, composition, and orientation. Unless otherwise stated the strain rate was 2.5 x 10"3 per sec. Liquid testing environments included nitrogen (WOK), nitrogen cooled petroleum ether (133" to 293"K), and silicone oil (293" to 488°K). Resolved shear stress-shear strain curves were electronically computed from autographically recorded load-elongation curves. Stress and strain were resolved on the macroscopic noncrystallographic (hkO) [001] system operative under the specific test conditions of temperature, strain rate, and orientation reported earlier.17 RESULTS The temperature dependence of the work-hardening curves is shown in Fig. 1 for gold-rich crystals of 51.0 at. pct Au oriented near the center of the stereo-graphic triangle. Over the range of intermediate temperatures from -200" to 400°K, they are very similar to those classically observed for fcc metals (reviewed by Nabarro et al.).'' The beginning of deformation is characterized by a region of decreasing hardening rate, stage 0, which is followed by a region of low linear hardening, stage I, and then a region of higher linear hardening, stage 11. At the higher temperatures, stage 111 is observed, a region of decreasing hardening rate. Over the intermediate temperature range, the extent of stage 0 and of the slow transition between stages I and I1 decreases with increasing temperature. Total ductility is large, often greater than 300 pct shear. As the temperature is either increased or decreased, the extent of stage I is decreased, giving rise to parabolic type flow and reduced ductility. Similar temperature effects have been reported for bcc ~r~stals.~~-~~ Below -14O°K, hardening is terminated in brittle fracture while above -400°K. initial hardening is followed first by work-softening and then by chisel-edge type ductile fracture. Stoichiometric (50.0 at. pct Au) and Zn-rich (51.0 at. pct Zn) crystals were also tested from 77" to -500°K. The effect of composition on the flow behavior is illus-

Jan 1, 1970

By G. Sachs, A. W. Dana, C. C. Chow

A great number of the investigations on the plastic flow of metals have been concerned with the establishment of a "universal" stress-strain relation. In such a relation some stress function when plotted against a strain function should yield identical curves for the various stress states. In the first investigation of this type, Ludwik and Scheu1 plotted the maximum shearing stress as a function of the maximum principal strain. Later Ros and Eichinger2 introduced two universal stress-strain relations, the one relating the maximum shearing stress to the maximum shearing strain, and the other relating a stress invariant, suggested by von Mises and Haigh, to the corresponding strain invariant. (In more recent investigations the stress and strain invariants are frequently supplemented with some factor to render their meaning more lucid.) A further suggestion which has not attracted appreciable attention is that by Baranski³ who used stress and strain deviators. The most common means of experimentation to determine the relation between stress and strain consists in subjecting thin walled tubes to combined internal pressure and axial tension.4a,4b,4c This method allows the study of plastic flow under stresses which are variable in two directions. However, the plastic flow which can be obtained in this manner is comparatively small, being limited by either tension failure or instability. For copper,'. only the relation between maximum shearing stress and maximum shearing strain yielded good agreement. On the other hand, tests on a stee14b and on an aluminum alloy4c. resulted in systematic deviations if any of the discussed universal stress-strain relations were used. It would seem, therefore, that the agreement mentioned above for copper is only incidental and explained by its high rate of strain hardening compared to that of other metals. Much larger strains than experienced in the tube tests can be obtained by subjecting a thin membrane of a ductile metal, which is restrained at its periphery, to a uniform hydraulic pressure. The thin sheet forms a deep bulge before it fails. The stresses and strains in such a bulge increase with increasing distance from the edge of the clamping "die," the maximum stresses and strains occurring at the pole (crown) of the bulge. While the stress and strain states are determined by the contour of the bulge, the absolute magnitude of the stresses and strains depends upon the hydraulic pressure. The bulge contour is in turn correlated with the geometry of the die opening. The deformation and fracture characteristics of circular bulges, that is, bulges formed with circular clamping dies, have been the subject of numerous experimental and analytical investi-gations.5,6,7 It has been shown that plastically deformed circular bulges develop large and comparatively uniform strains before failure by instability"6b,6c,6d and closely assume a spherical shape.6d Also the distribution of strains across the contour of the bulge is dependent on the metal being investigated and is correlated with, but cannot be predicted from, the metal's stress-strain characteristics. On the other hand, oblong or elliptical bulges, that is, bulges formed with elliptical clamping dies, are not as susceptible to analytical analysis and have not been investigated to the extent that circular bulges have. The few available data6c,7c indicate that stress states are obtained at the poles of the bulges, varying between plane strain and balanced biaxial tension, depending upon the geometry of the die opening. In this paper, the strain state and curvatures exhibited by three bulge shapes, a circular and two elliptical bulges, Fig 1, are analyzed experimentally using methods described in previous publications.6a,6c An attempt is made to derive the stress-strain relations for these bulges, which represent strain states in which the ratio of the two positive principal strains varied between 1.0 and 0.35. In addition, tension tests yielded data for a value of —0.5 for this strain ratio. Such an analysis should indicate the applicability of the various laws correlating stress with strain to the stress and strain states occurring in bulged shapes. Definitions and Nomenclature The definitions of the major stress and strain quantities used in this paper are as follows: s1, s2, s3 = principal normal stresses Sl > s2 > S3 t = shear stress e = conventional (unit) strain e = In (1 + e) El, E2, E3 = principal natural strains 7 = shear strain The maximum shear stress: , _ S1 — S3 lmax = 2 Frequently, the flow stress, s1 — s3 = 2lmax rather than the maximum shear stress is used.

Jan 1, 1950

By Rex O. McLellan

The formation of a monovacancy in a metal is simulated in an elastic model by the displacement of the surface of a small spherical cavity in a large elastic continuum. The application of linear elasticity to this distortion results in a well- known formula for the energy and an expression for the concomitant entropy change due both to the shear strain in the continuum and also to the dilation of the solid resulting from the boundary conditions at the surface of the solid. Elastic data (the sliear modulus and its temperature coelficient) are used to calculate the entropy and energy of formation for many metals. Despite the simplicity of the assumptions involved, the agreement between the calculated entropies and energies and experimental values is remarkably good. In recent years there has been a large increase in measurements of the absolute concentration of mono-vacancies in metals as a function of temperature. Hence new data for both the energy and the noncon-figurational entropy of formation of monovacancies has become available. Recent measurements&apos; of the anomalous (non-Arrhenius) self-diffusion in many bcc metals has also focused interest on the prediction of the thermodynamic parameters of mono- and multi-vacancies in those metals for which no data are available. Damask and Dienes&apos; have discussed the various theoretical calculations of the energy of formation EL, of a monovacancy. These include simple models involving the breaking of atomic bonds on moving atoms from the interior of a crystal to the surface, models combining elastic calculations with surface-energy terms and detailed quantum mechanical calculations. The simler models give the correct order of magnitude of &, but tend to overestimate it by a factor of about two. The quantum mechanical calculations4"7 have been carried out for the noble and alkali metals with generally reasonably good agreement with the available Ef data. The calculation of entropy of formation Sfv14 lnvolves a fundamental calculation of the perturbation of the phonon spectrum caused by the creation of a vacancy. Huntington, Shirn. and wajda8 have given an approximate evaluation of sJV by considering an Einstein model for the localized vibrations in the immediate neighborhood of the defect and then using elastic theory to calculate the entropy associated with the shear stress field in the distorted crystal (as originally proposed by Zenerg). They also included a term due to the dilation of the crystal. They obtained a value of 1.47k for copper, in good agreement with the experimental value (1.50k). However, Nardelli and Tetta- manzi1° have recently shown that neglecting the coupling between atoms (Einstein Model) may lead to a serious error so the agreement may be somewhat fortuitous. In this work simple linear elastic theory is used to calculate the entropy and energy of formation of mono-vacancies. Despite the simplicity of some of the assumptions involved, the agreement with the available experimental data is remarkable. However. the reasonable degree of success in the application of linear elastic calculations to the excess entropy of a solute atom in a dilute solid solution1&apos; indicates that the application of elastic theory to vacancies. where the interaction of different atomic species is not involved, may not be inappropriate. THE ELASTIC MODEL The metal is assumed to be a spherical elastic continuum. A small spherical cavity of volume V = 4i;v:&apos;/3 is cut from the center. removed. and dissolved rever-sibly in the bulk of the material. TO a good approximation no net atomic bonds are broken and the material does not undergo a volume change although the externally measured volume of the body would increase by V. The radius of the sphere of metal is much larger than r Next a negative pressure is applied to the cavity causing its surface to be displaced inward by an amount simulating the relaxation of the lattice around a monovacancy. In this model the energy and entropy accompanying the distortion are taken as 4, and <. As a first approximation the equation of state for the solid is taken as: r = ro(i + *~D LiJ where K is the bulk modulus. P the hydrostatic pressure. Vo the volume of the material at 0°K and zero pressure. and d+/dT = 30. where 0 is the linear thermal expansion coefficient. The variation of entropy with hydrostatic pressure is given by the Maxwell equation: These equations give the entropy change resulting from increasing the hydrostatic pressure from 0 to P as: and since • we have: This is the entropy arising from the dilation resulting

Jan 1, 1970

By Robert Baschwitz, John Chipman

The distribution of silicon between liquid silver and Fe-Si-C alloys has been studied at 1420oand 1530°C. The data are consistent with earlier studies. New data of Hager on the liquidus lines of the system Ag-Si and the distribution data are used to obtain the activity coefficient of silicon in both liquid phases. Data on the heat of mixing in iron permit accurate extension to 1600°C. Equilibrium data involving SiO2 and silicon in liquid iron together with revised data on the free energy of SiO2 are used to fix the activity of silicon in the infinitely dilute solution. The binary system exhibits strong negative deviation from ideality. At infinite dilution ? Si at 1600" is 1.25 x 10&apos;3, and at concentrations up to NSi = 0.4 the slope d InySi/dNSi has a constant value of r; = 13. It is found that logysi in the ternary solutzon is approximately but not exactly the same function of Nsi + NC as of NSi in the binary. The results are consistent with currently available data on the free energy of Sic and its solubility in molten iron. LIQUID solutions of the system Fe-Si-C have acquired considerable importance as the laboratory prototypes of blast furnace hot metal. Equilibrium studies involving such solutions and slags approximating those of the blast furnace have yielded useful information concerning the thermodynamic properties of blast furnace slags. In studies of this kind great importance attaches to a knowledge of the thermodynamic activity of silicon in the solution as a function of temperature and composition. An attempt was made by Chipman, Fulton, Gokcen, and askey&apos; to evaluate all of the pertinent data on this system and to deduce the desired relation between activity, composition, and temperature. These authors published data on the solubility of graphite and Sic in molten Fe-C-Si solutions and on the distribution of silicon between liquid iron and liquid silver. They showed further how the activity of silicon in very dilute solutions in liquid iron could be calculated from equilibrium data involving the molten alloy and solid SiO,. These calculations rested on the published thermodynamic properties of SiO, in- cluding its heat of formation which at that time was recorded as -209.8 kcal. This value has been under suspicion for some time and has recently been replaced by the concordant results from two independent laboratories2,3 which place the heat of formation of a-quartz at -217.6 kcal. This revision necessitates a re-evaluation not only of the activity of SiO2 in slag but also of silicon in molten iron. It is the purpose of this paper, therefore, to recalculate the activity of silicon, and in furtherance of this objective to present new data on its distribution between liquid Fe-Si-C alloys and liquid silver. HEAT OF SOLUTION OF SILICON IN IRON In order to determine the effect of temperature upon the activity coefficient it is necessary to know the heat of solution of silicon in iron as a function of composition. This is found in the data of Korber and Oelsen4 shown in Fig. 1. The curve corresponds to the following equation, which is of a form suggested by Wagner:5 Here AH is the heat absorbed in kilocalories in forming one gram atom of molten alloy from its molten elements and the N&apos;s are atom fractions. The relative partial molal enthalpies of the components, each referred to its pureliquid state and defined as zFe = aFe - PFe and zsi = HSi — -psi, are shown graphically. At low concentrations zSi = -28.5 kcal, in agreement with Kijrber and Oelsen&apos;s computation. This is in good agreement with the value of -29.3 kcal obtained by Chipman and Grant6 using an entirely different method. ACTIVITY AT INFINITE DILUTION From the known free energy of SiO, it is possible to obtain the activity of silicon in dilute solution in liquid iron from equilibrium studies. The heat of formation of a-quartz is —217.6 kcal and the heat capacity and entropy data are given by Kelley and ~ing.&apos; The free energy of formation of ß-cristo-balite at temperatures above the melting point of silicon is expressed by the following equation: Si(Z) + O2(g) = SiO2 (crist); AF" =-226,500 + 47.50T [I] The value of the deoxidation product for silicon [%Si] x [%O]2 at 1600°C according to Gokcen and chipmans is 2.8 x 10"5, in agreement with results of Hilty and Crafts.9 More recent works of Matoba,

Jan 1, 1963

By P. Duwez, C. B. Jordan

A. J. SHALER*—I should like to congratulate the authors for having carried out such a precise set of experiments. It has been found useful, in sintering experimental compacts in vacuo, to make certain that the residual gas is not one which reacts with the metal. Since traces of oxygen can be kept away only with great difficulty, the technique is often adopted of using a "getter " of powder in the vicinity of the compacts, and, in addition, of permitting a small hydrogen leak to flow into the vacuum chamber. Did the authors use similar devices? This paper brings up a question concerning the definition of the word &apos; sintering.&apos; The authors restrict its use to the adhesion between particles. Kuczynski, in a paper presented at this meeting, applies the word to the growth of areas of contact between particles. I have used it to mean both these phenomena and also the dimensional changes which continue to take place after the first two have run their course. May I suggest that we should come to an agreement on the use of these words ? Fig 1 and 2 show an interesting feature: extrapolation of the curves to zero time does not give a densification parameter of zero. The higher the temperature, the higher is the intercept on that axis. These observations agree with the concept of a practically instantaneous densification taking place while the compact is being brought to heat. Such a change may be brought about by plastic deformation and primary creep. The stress pattern causing this first rapid flow is, to my mind, due to the force of attraction between the surfaces of opposite particles in the regions immediately flanking their common areas of contact. The stress is not temperature-sensitive, but at room temperature plastic deformation only proceeds until the metal in the area of contact can support it elastically. As the metal is heated, the elastic limit falls, and further plastic flow occurs. At the higher temperatures, this is followed by primary creep, and finally by the steady-state rate-reaction which the authors are seeking. If they were to recalculate their densification-parameter values, using, not the initial density of the cold compact, but the density after the compacts have been brought to temperature, the systematic deviations from linearity in Fig 3 and 4 might be eliminated. Such initial densities might be obtained by extrapolating the curves of Fig 1 and 2 to zero time. I am naturally pleased to see that such a very well done series of experiments leads to a heat of activation (for the densification process in hydrogen) that is much higher than that for self-diffusion, in confirmation of the less elaborate results reported by Wulff and myself (Ind. and Eng. Chem., (1948) 40, 838). J. T. KEMP*—I would like to comment on Dr. Shaler&apos;s remarks. There are apparently different interpretations of the word "sintering." It seems to me that an accurate definition of our word is essential in all metallurgy. May I point out, in this connection, that in practical metallurgy the word "sintering" has been applied to a bonding process in the preparation of ores and flue dust for fur-nacing. It would be unfortunate if in the area of powdered metallurgy we should establish a definition that is essentially different in meaning. F. N. RHINES*—I think that I can answer the question by saying that I see no essential difference between the use of the term "sintering" in extractive metallurgy and in powder metallurgy; physically the same things are going on. I admit sintering is used for different end purposes in the two cases. When we resort to the sintering of lead ore mixture we are doing so to obtain a chemically reactive, loose texture of some rigidity. This is only a difference in use. After all, in powder metallurgy we sometimes deliberately produce a very porous material which has just a little strength, just as in the case of sinter cake. P. DUWEZ (authors&apos; reply)—We agree that it would be helpful to have well-established definitions of such terms as "sintering." Since the question has now been raised, the time might be appropriate for its consideration by some suitable committee of one or more of the metallurgical societies. In answer to Dr. Shaler&apos;s first question, no getter nor hydrogen leak was used in our vacuum experiments, except insofar as the guard disks (used to reduce friction between specimens and trays) may have acted as getters. Dr. Shaler&apos;s statement that extrapolation of the curves of Fig 1 and 2 does not lead to zero densification at zero time apparently overlooks the logarithmic

Jan 1, 1950

By C. N. Garman

Homestake-New Mexico Partners operate a 750-tpd carbonate leach uranium concentrate mill near Grants, N.M. The highly mineralized water available as process water leaves much to be desired. The 628 ppm as CaCO 3 makes the use of raw water very troublesome in pipes and on filter cloths. However, the residual sodium carbonate in the final filter cake going to tails makes an ideal softening agent. To take advantage of this fact, all makeup water used in the mill is first used as tailing slurry dilution water and comes to the mill from the tailings pond. The 5-acre tailings pond serves as a thickener and 100 to 150 gpm of nearly clear solution is decanted to a pump to be returned to the mill. Since this tailings water has small quantities of uranium in the solution an ion exchange scavenger unit was installed to remove as much uranium as possible. The ion exchange raffinate is then used as final filter wash ahead of the tailings slurrying step. In spite of the large settling area this return water is not clean enough for ion exchange feed. The solids present are very fine and composed of approximately 15 pct (by weight) burnable carbonaceous material common to the sandstone uranium ores in the area, 40 pct SiOz plus 45 pct CaC03. Laboratory work showed that this material responds very well to flotation. Before deciding to use flotation, various clarifying systems such as pressure leaf filters, sand filters, and continuous vacuum pre-coat filters, were considered. Each of these could have solved the problem but with much more operating labor, more reagents and greater installation costs than the flotation step. About 100 to 150 gpm of fouled water is fed to two 66-in. Fagergren cells, in series. Reagents used at the beginning were Arquad 2HT75 and Arquad C50, at the rate of about 1% lb per 8-hr shift, or about 0.0053 lb each per ton of ore. This did not completely remove the solids but does an acceptable job. Approximately 75 pct of the slimes are a size that can be caught on a 41-Whatman Paper are removed. Removal of these slimes also allows much better settling of the coarse nonfloatable material. Advantage is also taken of this fact in a small settling tank ahead of precipitation. Removal of this amount of the slimes makes the ion t:xchange feasible. PREGNANT SOLUTION CIRCUIT The carbonate? leach-caustic precipitation method of uranium concertration does not provide for any process purification step ahead of precipitation. Therefore, any fine solids getting into the pregnant solution through the filter cloth show up in the final concentrate. This, of course, lowers the grade, and, at times, the slimy nature of these very fine solids rendered final filtration of the concentrate difficult if not impossible At Homestake-New Mexico Partners a 75-ft thickener was available for gravity clarification of 100 to 120 gpm of this pregnant solution. However this did not sufficiently remove the slimes. Laboratory investigation of the whole range of flocculants that were suggested by literature, salesmen, and friends failed to turn up anything of consequence. A continuous vacuum pre-coat filter would do the job and was investigated. The capital cost and the operating labor and materials made this a last chance choice. Following work done in the metallurgical laboratory on the tailings return water, it was found that some changes in the reagent strengths and combinations made a very definite decrease in the solids in the pregnant solution. Concentrate grade improved about 5 pct anti the final product after drying had an appreciably greater bulk density. Compared to a cost of about 2.2e per ton for pre-coat filter opelation for cleaning just one circuit, flotation costs less than 1.0 per ton of ore for cleaning two circuits. While a pre-coat filter would do a more thorough job, the flotation does all that is required for either circuit. Gravity causes the froth produced to run back into the leach circuit. This does not appear to result in a build-up of objectionable slime. No extra manpower is required; the operators in the separate areas can observe the operation of the cells and mix the small quantities of reagents as needed. Normally the 66-in. Fagergren cell requires 15 hp per cell, but this very dilute slurry needs only 10 hp for both cells. Originally, a combination of the two Arquads mentioned previously served as frothers and promoters. As further testing

Jan 1, 1962

By F. E. Werner, B. L. Averbach, Morris Cohen

THAT bainite formed near the M, temperature bears a striking r esemblance to martensite tempered at the same temperature has been shown by the electron microscope.&apos; By means of electron diffraction,&apos; it has been established that carbide and cementite are present in bainite formed at 500°F (260°C); these carbides are also found in martensite tempered at 500°F (260°C).&apos; The investigation reported here is concerned with an X-ray study of the matrix phases in lower bainite and tempered martensite. These phases have turned out to be dissimilar in structure; the matrix of bainite is body-centered-cubic while that of tempered martensite is body-centered-tetragonal. A vacuum-melted Fe-C alloy containing 1.43 pct C was studied. Specimens of 16 in. diam were sealed in evacuated silica tubing and austenitized at 2300°F (1260°C) for 24 hr. One specimen was quenched into a salt bath at 410°+7 °F (210°+4°C), held for 16 hr, and cooled to room temperature. The structure consisted of about 90 to 95 pct bainite, the re: mainder being martensite and retained austenite. A second specimen was quenched from the austen-itizing temperature into iced brine and then into liquid nitrogen. It consisted of about 90 pct martensite and 10 pct retained austenite. The latter specimen was tempered for 10 hr at 410°+2°F (210°+1°C). The specimens were then fractured along prior austenite grain boundaries (grain size about 2 mm diam) by light tapping with a hammer. Single aus-tenite grains, mostly transformed, were etched to about 0.5 mm diam and mounted in a Unicam single crystal goniometer, which allowed both rotation and oscillation of the sample. Lattice parameters were measured by the technique of Kurdjumov and Lyssak. This method takes advantage of the fact that martensite and lower bainite are related to austenite by the Kurdjumov-sachs orientation relationships Thus, the (002) and the (200) (020) reflections can be recorded separately, permitting the c and a parameters to be determined without interference from overlapping reflections. According to these findings, the matrix phase in bainite is body-centered-cubic and, within experimental error, has the same lattice parameter as ferrite (2.866A). On the other hand, martensite, tempered as above, retains some tetragonality, with a c/a ratio of 1.005t0.002. Most workers in the past have assumed that bainite is generated from austenite as a supersaturated phase, but the nature of this product has not been established. The question arises as to whether bainite initially has a tetragonal structure and then tempers to cubic, or if it forms directly as a cubic structure. If it forms with a tetragonal lattice, it might well be expected to temper to the cubic phase at about the same rate as tetragonal martensite. The martensitic specimen used here was given approximately the same tempering exposure, 10 hr at 410°F, as suffered by the greater part of the bainite during the isothermal transformation. About 50 pct bainite was formed in 6 hr at 410°F. On tempering at this temperature, martensite reduces its tetragonality within a few minutes to a value corresponding to 0.30 pct C.&apos; Further decomposition proceeds slowly, and after 10 hr the c/a ratio is still appreciable, i.e., 1.005. Thus, even if the bainite were to form as a tetragonal phase with a tetragonality corresponding to only 0.30 pct C, which might be assumed to coexist with e carbide, it would not be expected to become cubic in this time. It seems very likely, therefore, that bainite forms irom austenite as a body-centered-cubic phase and does not pass through a tetragonal transition. The carbon content of the cubic phase has not been determined, but it could easily be as high as 0.1 pct, within the experimental uncertainty of the lattice-parameter measurements. It has been postulated that retained austenite decomposes on tempering into the same product as martensite tempered at the same temperature. There is now considerable doubt on this point. The isothermal transformation product of both primary and retained austenite at the temperature in question here is bainite," and the present findings show that bainite and tempered martensite do not have the same matrix. Acknowledgments The authors would like to acknowledge the financial support of the Instrumentation Laboratory, Massachusetts Institute of Technology, and the United States Air Force.

Jan 1, 1957

By Y. Nakada, U. F. Kocks, B. Chalrners

THE orientation dependence of work hardening has previously been studied over the entire range, i.e., including special orientations of high symmetry, in aluminum1-3 and silver.* The differences between various orientations are substantial, and the trend is the same in all fcc crystals. However, there are quantitative differences in behavior between aluminum and silver at room temperature, particularly in the (100) orientation. While many experiments on gold have been reported,5&apos;9 none included the special orientations. We therefore undertook tension tests on gold crystals of the axial orientations (100),(110), (Ill), (211), and, as a representative of single slip, one whose Schmid factor was 0.5 (hereafter referred to as m = 0.5). Single crystals of dimensions 1/4 by .1/4 by 6 in. were grown from the melt1&apos; at a rate of 4 in. per hr, using gold of 99.99 pct purity obtained from Handy and Harman. A growth rate of 1/8 in. per hr, or a purity of 99.999 pct,ll produced no difference in results. The crystals were annealed at 1000°C in air for 24 hr and furnace-cooled down to room temperature, after which they were electro-polished in a solution of potassium cyanide (40 g), potassium ferrocyanide (10 g), soda (20 g), and enough distilled water to make 1 liter of solution, at a current density of 0.02 amp per sq mm. After 2 or 3 hr, a very smooth surface was obtained by this method. Nine m = 0.5, two (loo), one (110), three (Ill), and one (211; crystals were tested at room temperature in a floor-model Instron machine at a tensile strain rate of 0.5 pct per min. The accuracy of the stress measurement was k10 g per sq mm up to 500 g per sq mm, k2 pct for higher stresses. The strain measurement was accurate to t2 pct. The scatter of stress at a given strain among the crystals of the same orientation was small, *5 pct being the largest. The representative stress-strain curves for various orientations are shown in Fig. 1. Table I summarizes the work-hardening parameters as used by seeger.12 Results of other investigators are also included in this table for comparison. There are no previous data for the corner orientations. Values for m = 0.5 crystals agree fairly well with those of Berner. Tm is defined as the stress at which the stress-strain curve begins to deviate from linearity of Stage 11. However, in practice this is a very difficult value to estimate because each investigator has a different idea of linearity. Therefore, the comparison with the values of other investigators may not be valid. In the present experiment, (100) crystals had the highest 111, followed by (111) crystals. The work-hardening rate in Stage I1 was highest for the (111) crystal followed by the (100) crystals. Our value for 0x1 of m = 0.5 orientation agrees very well with those of other investigators. 1) Single-Slip orientation (m = 0.5). These crystals were oriented so that the primary-slip vector was contained in one side face. The dimension perpendicular to this side should then not change if single slip indeed takes place. Within the accuracy of measurement, this dimension did not change during the deformation. Since the accuracy is k0.2 pct, the amount of secondary slip is less than 0.7 pct of the slip on the primary-slip system at 30 pct tensile strain. This is in conformity with the results obtained by Kocks" for aluminum crystals. The tensile axis moved toward (211) during the deformation. Slip bands, Fig. 2(a), were very fine and closely spaced. Some deformation bands were observed. There were no clear-cut cross-slip traces such as the ones observed on aluminum m = 0.5 crystals. 2) (111) Crystals. The orientation of the tensile axis was stable during the deformation. From this observation, one can deduce that at least three slip systems must have operated, and probably all six because the remaining three all have cross-slip relation to one of the first three.&apos; It was very difficult to observe the slip markings. Consequently, we could not confirm by this method that six systems were operative during the deformation. At high strains (above 50 pct shear strain), this orientation had the highest stress-strain curve. At 80

Jan 1, 1964

By J. Chipman, M. N. Dastur

The importance of dissolved oxygen as a principal reagent in the refining of liquid steel and the necessity for its removal in the finishing of many grades have stimulated numerous studies of its chemical behavior in the steel bath. From the thermodynaniic viewpoint the essential data are those which determine the free energy of oxygen in solution as a function of temperature and composition of the molten metal. A number of experimental studies have been reported in recent years from which the free energy of oxygen in iron-oxygen melts can be obtained with a fair degree of accuracy for temperatures not too far from the melting point. Certain discrepancies remain, however, which imply considerable uncertainty at higher temperatures; also several sources of error were recognized in the earlier studies. It has been the object of the experimental work reported in this paper to reexamine these sources of uncertainty and to redetermine the equilibrium condition in the reaction of hydrogen with oxygen dissolved in liquid iron. The reaction and its equilibrium constant are: H2 (g) + Q = H2O (g); K1 _ PH2O / [1] Ph2 X % O Here the underlined symbol Q designates oxygen dissolved in liquid iron. The activity of this dissolved oxygen is known to be directly proportional to its concentrationl,2 and is taken as equal to its weight percent. The closely related reaction of dissolved oxygen with carbon monoxide has also been investigated:3,4,5 co (g) +O = CO?(g); K _ Pco2___ [2] K2= pco X % O [2] The two reactions are related through the wat,er-gas equilibriuni: H2 (g) + CO2 (g) = CO (g) + H2O (g); K2 = PCO X PH2O [3] PH2 X PCO2 and with the aid of the accurately known equilibrium constant of this reaction, it has been shown5 that the experimental data on reactions [1] and 121 are in fairly good, though not exact, agreement. Experimental Method Great care was taken to avoid the principal sources of error of previous studies, namely, gaseous thermal diffusion and temperature measurement. The apparatus was designed to provide controlled preheating of the inlet gases and to permit the addition of an inert gas (argon) in controlled amounts, two measures found to be essential for elimination of thermal diffusion. A known mixture of water vapor and hydrogen was obtained by saturating purified hydrogen with water vapor at controlled temperature. This mixture, with the addition of purified argon, was passed over the surface of a small melt (approximately 70 g) of electrolytic iron in a closed induction furnace. After sufficient time at constant temperature for attainment of equilibrium the melt was cooled and analyzed for oxygen. GAS SYSTEM A schematic diagram of the apparatus is shown in Fig 1. Commercial hydrogen is led through the safety trap T and the flowmeter F. The catalytic chamber C, held at 450°C, was used to convert any oxygen into water-vapor. A by-pass B with stopcocks was provided so that the hydrogen could be introduced directly from the tank to the furnace when desired. From the catalytic chamber the gas passed through a water bath W, kept at the desired temperature by an auxiliary heating unit, so that the gas was burdened with approximately the proper amount of water vapor before it was introdvced into the saturator S. All connections beyond the catalytic chamber were of all-glass construction. Those connections beyond the water bath were heated to above 80°C to prevent the condensation of water vapor. After the saturator, purified argon was led into the steam-hydrogen line at J, and finally the ternary mixture was introduced into the furnace. THE SATURATOR The saturator unit comprised three glass chambers, as shown in Fig 1, the first two chambers packed with glass beads and partially filed with water and the third empty. Each tower had a glass tube with a stopper attached for the purpose of adjusting the amount of water in it. The unit was immersed in a large oil bath, which was automatically controlled with the help of a thermostat relay to constant temperature, ± 0.05ºC, using thermometers which had been calibrated against a standard platinum resistance thermometer. The performance of the saturator over the range of experimental conditions was checked by weighing the water absorbed from a measured volume of hydrogen; the observed ratio was always within 0.5 pct of theoretical.

Jan 1, 1950

The pressure of the gas in equilibrium with sphalerite has been determined in the temperature range of 680&apos; to 825°C, using the Knudsen orifice method. A comparison of these experimental pressures with those calculated from thermal data and from other equilibrium measurements shows that the vapor above sphalerite is predominantly dissociated ZnS. Equations have been given for correctly calculating dissociation pressures using the Knudsen orifice method. It has been shown that the experimentally determined pressure is the same, whether the zinc sulphide is sphalerite or not, or a mixture of wurtzite and sphalerite. CONFLICTING points of view appear in the literature on the constitution of the vapor in equilibrium with solid zinc sulphide in the vicinity of 800°C. By comparing the dissociation pressure calculated from thermodynamic data and the vapor-pressure determination of ZnS by Veselovski,1 Lumsden2 has concluded that the vapor consists largely of dissociated ZnS. Sen Gupta,&apos; however, concludes from his spectroscopic determinations that the vapor is largely ZnS molecules. In view of the fact that the thermodynamically calculated&apos; dissociation pressure is higher than that experimentally measured by Veselovski, it seemed in order to repeat Veselovski&apos;s measurements. Experimental Procedure The method used for the determination of the pressures in this papel- is the Knudsen effusion cell. The apparatus and procedure were described in a previous paper- from this laboratory on the determination of the vapor pressure of silver. The only difference is that the Knudsen cell in this work is made from platinum and there is no external cover around the cell. The cell is an ordinary platinum crucible of 2.2 cm top diameter with a capsule cover. It was thought that platinum might stand up at these temperatures to the solid and gaseous ZnS, since it was found that the weight of the platinum cell itself did not change appreciably on heating ZnS in it at the working temperatures. To insure that reaction of the zinc sulphide with the cell was not giving&apos; a false value, a stabilized zirconia cell was employed for check runs. Fig. 1 shows the comparison, which is satisfactory. Veselovski previously had measured the vapor pressure of ZnS using a silica Knudsen effusion cell. On repeating his experiment in this laboratory, it was found that ZnS at-tacked the silica cell, giving it a marked frosty appearance. This led to the belief that Veselovski&apos;s result:; may be in error. Also, he was operating at pressures above the range ordinarily considered safe for the Knudsen method. The effusion rate was measured by weighing the cell before and after each run. The weight loss during heating to temperature and cooling down was measured and subtracted from the weight loss during the actual run. The zinc sulphide used in this investigation was from two sources: Fisher cp grade, and a sample of pure sphalerite supplied by Mr. E. A. Anderson of the New Jersey Zinc Co. Before and after the series of runs with Fisher ZnS, X-ray analysis showed that both wurtzite and sphalerite were present. However, the ratio of sphalerite to wurtzite increased. All measurements were made below the transition temperature which has been reported" to be 1020°C. The data obtained in this investigation are tabulated in Table I. The pressure was calculated by the usual Knudsen formula" on the assumption that ZnS molecules were effusing. From these data, using pure sphalerite in the platinum Knudsen cell, the vapor pressure of ZnS, in mm of Hg, as a function of temperature is given by the solid line in Fig. 1. The best straight line, as determined by the method of least squares, is given by 14405 logpzns =-14405/T +11.032. A comparison of these results with Veselovski&apos;s shows that his results are about 50 pct lower. Discussion The vapor in equilibrium with solid zinc sulphide in the temperature range of this study will consist of Zn, S2, and ZnS mol, since other species of zinc and sulphur&apos; are relatively unstable. The question to be settled is whether or not ZnS is largely dissociated. The derivation8 which follows gives the method of calculating the pressure of zinc and sulphur over solid ZnS, assuming complete dissociation, from Knudsen cell data. The free energy of the reaction 2 ZnS(solid) ? 2 Zn(gas) + S2(gas) is given by ?F?° = -RT In K = —RT In p12p2 where p1 is the zinc pressure and p is the sulphur pressure. If dissociation occurs in a closed system,

Jan 1, 1955

By W. M. McKewan

Magnetite pellets were reduced in flowing hydrogen at pressures up to 40 atm over a temperature range of 350° to 500°C. The rate of weight loss of oxygen per unit area of the reaction surface was found to be constant with time at each temperature and pressure. The reaction rate was found to be directly proportional to hydrogen pressure up to 1 atm and to approach a maximum rate at high pressures. The results can be explained by considering the reaction surface to be sparsely occupied by adsorbed hydrogen at low pressures and saturated at high pressures. PREVIOUS investigation1,2 have shown that the reduction of iron oxides in hydrogen is controlled at the reaction interface. Under fixed conditions of temperature, hydrogen pressure, and gas composition, the reduction rate is constant with time, per unit surface area of residual oxide, and is directly proportional to the hydrogen pressure up to one atmosphere. The reduction rate of a sphere of iron oxide can be described3 by the following equation which takes into account the changing reaction surface area: where ro and do are the initial radius and density of the sphere; t is time; R is the fractional reduction; and R, is the reduction rate constant with units mass per area per time. The quantityis actually the fractional thickness of the reduced layer in terms of fractional reduction R. It was found in a previous investigation2 of the reduction of magnetite pellets in H2-H,O-N, mixtures, that the reaction rate was directly proportional to the hydrogen partial pressure up to 1 atm at a constant ratio of water vapor to hydrogen. Water vapor poisoned the oxide surface by an oxidizing reaction and markedly slowed the reduction. The enthalpy of activation was found to be + 13,600 cal per mole. It was also found that the magnetite reduced to meta-stable wüstite before proceeding to iron metal. The following equation was derived from absolute reaction-rate theory4,8 to expfain the experimental data: where Ro is the reduction rate in mg cm-2 min-&apos;; KO contains the conversion units; Ph2 and PH2O are the hydrogen and water vapor partial pressures in atmospheres; Ke is the equilibrium constant for the Fe,O,/FeO equilibrium; Kp is the equilibrium constant for the poisoning reaction of water vapor; L is the total number of active sites; k and h are Boltzmann&apos;s and Planck&apos;s constants; and AF is the free energy of activation. Tenenbaum zind Joseph5 studied the reduction of iron ore by hydrogen at pressures over 1 atm. They showed that increasing the hydrogen pressure materially increased the rate of reduction. This is in accordance with the work of Diepschlag,6 who found that the rate of reduction of iron ores by either carbon monoxide or hydrogen was much greater at higher pressures. He used pressures as high as 7 atm. In order to further understand the mechanism of the reduction of iron oxide by hydrogen it was decided to study the effect of increasing the hydrogen pressure on rebduction rates of magnetite pellets. EXPERIMENTAL PROCEDURE The dense magnetite pellets used in these experiments were made in the following manner. Reagent-grade ferric oxide was moistened with water and hand-rolled into spherical pellets. The pellets were heated slowly to 550°C in an atmosphere of 10 pct H2-90 pct CO, and held for 1 hr. They were then heated slowly to 1370°C in an atmosphere of 2 pct H2-98 pct CO, then cooled slowly in the same atmosphere. The sintered pellets were crystalline magnetite with an apparent density of about 4.9 gm per cm3. They were about 0.9 cm in diam. The porosity of the pellets, which was discontinuous in nature, was akrout 6 pct. The pellets were suspended from a quartz spring balance in a vertical tube furnace. The equipment is shown in Fig. 1. Essentially the furnace consists of a 12-in. OD stainless steel outer shell and a 3-in. ID inconel inner shell. The kanthal wound 22 in. long, 1 1/2, in. ID alumina reaction tube is inside the inconel inner shell. Prepurified hydrogen sweeps the reaction tube to remove the water vapor formed during the reaction. The hydrogen is static in the rest of the furnace. The sample is placed at the bottom of the furnace in a nickel wire mesh basket suspended by nickel wire from the quartz spring. The furnace is then sealed, evacuated, and refilled with argon several times to remove all traces of oxygen. It is then evacuated, filled with

Jan 1, 1962

By Robert E. Green, Wolfgang Sachse

Simullaneous ultrasonic attenuation measurements of both quasishear waves propagating in single cryslals of aluminum indicate that, in the undeformed annealed state, the dislocation density is generally not uniform on all slip systems. Change oof attenuation measurements made during plastic defortnation of crystals , which possessed specific orientations ideal for studying the orientation dependence of dislocation damping, indicate that, for low strain levels, dislocation motion occurs on additional slip systems besides the primary one, even for crystals oriented for plastic deformation by single slip. THE sensitivity of internal friction measurements permits such measurements to be used successfully in studying the deformation characteristics of metal crystals. On the basis of experimental observations, T. A. Read1 was the first to associate internal friction losses with various dislocation mechanisms. Since that time further work2-&apos; has been performed and a dislocation damping theory has been formulated by Granato and Lucke.6 In the amplitude independent region, this theory predicts the attenuation a to be dependent on an orientation factor O, a dislocation density A, and an average loop length L. if is a constant, independent of crystallographic orientation. For a given crystallographic orientation, changes in dislocation density and loop length give rise to the observed attenuation changes accompanying plastic deformation. The Granato-Liicke theory suggests the investigation of the orientation dependence of attenuation measurements in hopes of obtaining information to separate dislocation motion losses from other losses.7 An experimental study of the orientation dependence of attenuation in undeformed annealed single crystals should yield an insight into the uniformity of dislocation distribution throughout the entire specimen. A similar study on crystals plastically deformed in a prescribed fashion should give information about the alterations in the dislocation distribution on the slip systems activated during plastic deformation. The possible modes of elastic waves which can be propagated in aluminum,8 copper,9 zinc,10 and other hexagonal metals" have been calculated. Associated with each mode of wave propagation are dislocation damping orientation factors, which are based on the resolution of the stress field of that particular elastic wave onto the various operative slip systems in the material. These orientation factors have also been calculated as a function of crystallographic orientation in the papers cited above. Einspruch12 obtained agreement between predicted and observed attenuation values of longitudinal and shear waves in (100) and (110) directions of two undeformed aluminum crystal cubes. He ascribed the slight deviations between predicted and observed values to a nonuniform dislocation distribution, or to other loss mechanisms. In shear deformation of zinc crystals, Alers2 found that the attenuation of shear waves having their particle displacements in the slip plane was very sensitive to the deformation, while the longitudinal wave attenuation was affected only when the wave propagation direction was not normal to the slip plane. Using aluminum single crystals oriented for single slip, Hikata3 et al. found that during tensile deformation the change of attenuation of the shear wave (actually quasishear) having particle displacements nearly perpendicular to the primary slip direction exhibited the easy-glide phenomena, while longitudinal waves did not. Similar results were reported by Swanson and Green5 during compressive deformation of aluminum crystals. These results are in qualitative agreement with the calculated orientation factors for specimens of this orientation. In well-annealed, undeformed aluminum crystals, the damping is expected to be due to dislocations vibrating on all twelve slip systems. The orientation factors associated with this initial damping will be designated by O2 and O3, where a, represents the average orientation factor for the slow shear (or quasishear) wave and O3 represents the average orientation factor for the fast shear (or quasishear) wave. The calculation of these values for aluminum crystals by Hinton and Green8 shows that they vary very little as a function of crystallographic orientation—at most, by a factor of 2.47. If the dislocation density and loop length are uniform, then in the initial undeformed state, Here the subscript zero refers to the initial value of the attenuation. Also for aluminum, the calculations8 show that the orientation factors for primary slip only, associated with each shear wave, exhibit a sharp minimum for particular crystallographic orientations. A composite plot of the two shear wave orientation factors for primary slip only is shown in Fig. 1. Since these orientation factors are associated with dislocation motion occurring on the primary slip system only, the proper condition to check these factors might be attained by slightly deforming a single crystal oriented for primary slip. For dislocation motion on the primary slip system only,

Jan 1, 1969

By F. Weinberg

The relative concentration of 1" at grain boundaries in controlled orientation bicrystals has been examined by autoradiographic techniques, and by activity measurements of grain boundary surfaces exposed by preferential ,melting. The autoradio-graphs indicate that thallium is concentrated at grain boundaries in as-grown bicrystals, but not in zcell-annealed bicrystals. They also indicate that the solute concentration and the distribution on as-grown bicrystal surfaces are markedly different than that of the bulk material. The boundary surface measurements are in agreement with the autovadiographic evidence. On the basis of these measurements, as-grown bicrystals containing approximately 100 ppm of Tl, solidified at rates between 5 and 30 cm per hr and with tilt boundaries greater than 10 deg, exhibited grain boundary segregation equivalent to roughly 10 atomic planes of pure solute. Higher solute concentrations (equivalent to 140 atomic planes of pure solute) were obtained in bicrystals solidified slowly (0.6 cm per hr); slightly higher values were obtained in specimens containing a large angle nantilt boundary. Annealing for various times over a range of temperatures eliminated grain boundary segregation within the experimental uncertainty of the results (equivalent to 1 atomic Plane of pure thallium at the boundary). The results for the as-grown bicrystals can be qualitatively accounted for by assuming the presence of a groove on the solid-1iq;id interface, at the grain boundary. SOLUTE segregation at grain boundaries may be considered in two parts, namely, nonequilibrium segregation associated with the solidification process, and equilibrium segregation in fully annealed materials.&apos; There is much indirect evidence for nonequilibrium segregation, based on preferential etching at grain boundaries and the mechanical properties of as-cast alloys. In addition, some direct observations have been reported in which radioactive tracers were used as solute additions and segregation detected at the grain boundaries by autoradiographic techniques. However, there is little detailed quantitative data on solute concentrations related to grain boundaries, particularly for different freezing conditions and grain boundary configurations. Equilibrium segregation at grain boundaries has been considered both theoretically and experimentally. cean&apos; has made an estimate of the maximum equilibrium solute concentration that might be expected at a grain boundary, based on the lattice distortions in the boundary region. He arrived at a concentration which was equivalent t a monatomic layer of pure solute. A similar value, based on thermodynamic arguments, was calculated by Cahn and Hilliard for the segregation of phosphorus in iron. Experimentally, much higher values of solute concentration at grain boundaries have been reported recently by both Inman and iler&apos; for phosphorus in iron, and Ainslie et 1.&apos; for sulfur in iron. They observed concentrations equivalent to as much as 20 to 100 atomic layers of pure solute at the grain boundaries. However, in both cases it was shown that the observed segregation was not due solely to equilibrium segregation at the grain boundary. In the former case, precipitation effectss due to trace impurities in the material were believed to account for the large amount of solute present at the grain boundary. In the latter case it was shown that a high density of dislocations in the boundary region could provide a large number of additional sites for solute atoms, other than at the grain boundary. Thomas and chalmera have reported on the equilibrium segregation of po210 in grain boundaries of Pb-5 pct Bi alloys. Using autoradiographic techniques, they observed a concentration of polonium along the boundary trace on the surface of annealed bicrystal specimens grown from the melt. The concentration only appeared after annealing, and varied with boundary angle, increasing as the boundary angle increased. Their conclusions have been questioned by Ward," who pointed out that the segregation they observed along the boundary trace was much too wide to be compatible with the usual concepts of the thickness of a grain boundary of several lattice spacings. Also, Maroun et al.,l1 with specimens similar to those of Thomas and Chalmers, found that segregation could only be detected on the specimen surface, suggesting that Thomas and Chalmers&apos; results were associated with an oxidation effect of polonium, and not equilibrium segregation. Thomas and Chalmers replied12 that they did observed segregation at the grain boundary in the bulk material and suggested further experiments were necessary to resolve the difference. The purpose of the present investigation was to examine both nonequilibrium and equilibrium grain boundary segregation in melt grown bicrystal specimens as a function of boundary angle, growth rate, and solute concentration, and to de-

Jan 1, 1963