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By R. Schuhmann, J. Th. Overbeek, P. L. De Bruyn

THE technique of concentrating valuable minerals from lean ores by flotation depends upon the creation of a finite contact angle at the three-phase contact, mineral-water-air. If the mineral is completely wetted by the water phase, contact angle zero, there is no tendency for air bubbles to attach themselves to the mineral. However, when the contact angle is finite, the surface free energy of the system, water-air bubble-mineral particle, can be diminished by contact between the bubble and the particle, and if not too heavy the mineral will be levitated in the froth. With a few exceptions, all clean minerals are completely wetted by pure water. Thus the art of flotation consists in adding substances to the water to make a finite contact angle with the mineral to be floated, but to leave the other minerals with a zero contact angle. The contact angle concept and experimental measurements of contact angles have played important roles in flotation research for several decades.'-" Nevertheless, there remain unanswered some basic questions as to the scientific significance of the contact angle and the nature of the processes by which flotation reagents affect contact angles. The contact angle is a complex quantity because the properties of three different phases, or rather of three different interfaces, control its magnitude. Considering the interfaces close to the region of ternary contact to be plane, the relation among the contact angle and the three binary interfacial tensions is easily derived. The condition for equilibrium among the three surface tensions, Fig. 1, or the requirement of minimum total surface free energy leads to Young's equation, Eq. I: ysa — ysl = yLA cos 0 [1] According to this equation, the contact angle has one well-defined value. Actually it is found in many experiments that the value of the contact angle depends on whether the air is replacing liquid over the solid (receding angle) or the liquid is replacing air (advancing angle). The receding angle is always the smaller of the two.4 Two explanations have been offered for this experimental fact. According to some investigators,5-8 roughness of the surface causes apparent contact angles that are different for the receding and the advancing cases although the actual local contact angle may be completely determined by Eq. 1. The other explanation involves the hypothesis that the solid-air interface after the liquid has just receded is different from the same interface when no liquid has previously covered it.1,4 Adsorption of constituents of the air or liquid might play a role here. In this discussion the difference between advancing and receding contact angle will be neglected and plane surfaces where Eq. 1 describes the situation will be considered. But there is still a fundamental obstacle to the application of Young's equation. The surface tension of the liquid (rla) can easily be determined, but the two surface tensions of the solid (rsa and ySL) cannot be measured directly. Eq. 1, however, is not without value. By contact angle measurements it is possible to establish how ysl — ysl varies with the addition of solutes to the liquid phase. Also, Eq. 1 affords a convenient starting point for calculating net forces and energy changes involved in the process of bubble-particle attachment.1,2 . If for the moment surface tension of the liquid (yLa) is considered a constant, an increase in ysa — ysL, will tend to decrease the contact angle. A decrease in ySA — ysl, corresponds to an increase of the contact angle. In cases where ySA — ySL > yLa the contact angle is zero; it will only reach finite values when ysa — ysa has been decreased below YLA. Thus on the basis of Young's equation and contact angle measurements alone, it can be learned how flotation reagents affect the difference Ysa — ysl, but no conclusions can be drawn as to the effects of reagents on the individual surface tensions ysa, and ysL, not even as to signs or directions of the surface tension changes resulting from reagent additions. A quantitative relationship between the surface tension or interfacial tension and the adsorption occurring at a surface or an interface is given by the Gibbs equation, which for constant temperature and pressure reads dy = — 2 T, du, [2] where dy is the infinitesimal change in surface tension accompanying a change in chemical potential

Jan 1, 1955

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.' By means of electron diffraction,' 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).' 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.' 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

The pressure of the gas in equilibrium with sphalerite has been determined in the temperature range of 680' 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,' however, concludes from his spectroscopic determinations that the vapor is largely ZnS molecules. In view of the fact that the thermodynamically calculated' dissociation pressure is higher than that experimentally measured by Veselovski, it seemed in order to repeat Veselovski'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' 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'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'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' 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 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.' 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' 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' for phosphorus in iron, and Ainslie et 1.' 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' 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

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 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 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 ' sintering.' 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'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' 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'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'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 Willard C. Lacy

Rather than attempting to present a summary of the many and highly varied papers that have been presented at this symposium on sampling and grade control, I will attempt to extract the general philosophy of analysis and approach, and attempt to identify the trend of future developments. First, the term "sampling" is used with its broadest connotations. A sample consists of a representative portion of a larger mass, and must represent the mass not only in the grade of contained metals or minerals, but also in all other respects in terms of mineralogy and mineral quality (1, 5), deleterious materials, recoverability of economic components, physical behavior, geophysical response (I), and even archaeological and environmental aspects (7, 11). The sample must be taken from a locality and in such a manner and quantity that it is representative of the larger rock mass. This calls for complete and accurate geological control and an understanding of the nature and distribution of the contained chemical and physical elements and a record of the effectiveness of the different sampling methods. Second, value of a given mass of ore material is based upon its profitability - the difference between recoverable value and costs to achieve recovery, beneficiation and sale. There is a strong movement in mining geology control toward more complete analysis in determining cutoff grades and in grade control, as illustrated by the kriging of metallurgical recovery factors as well as grade at the Mercur Mine (8). To achieve a "profit- ability factor" as a guide for economic mining practice requires further integration of: 1) the value of contained metal or mineral, 2) percentage recovery of values, 3) dilution of ore with waste rock, 4) addition to, or loss of value as a consequence of by-product materials or deleterious components, 5) cost of producing a saleable product plus mini- mum profit to justify the effort (cutoff), and 6) cost of land restoration (7, 11). All these parameters vary with the rock type, rock structure, mineralogy, depth, geometry, mining and metallurgical methods, but they must be sampled and analyzed if sampling and grade control are to reflect profitability. A wide variety of deposits has been presented at this symposium; each deposit with its own problems and special solutions. Deposits containing high unit-value components, e.g. precious metals and diamonds, present special problems in the obtaining of accurate samples and generally require statistical analysis control methods or may disregard or modify occasional high or occasional low values, based upon experience (12 ) Grade control may be accurate for the long term but may vary for the short term. Bulk sampling is always essential. Deposits containing metals or minerals with low unit value are very sensitive to transport costs, and they are often very sensitive to small amounts of deleterious components or differences in physical or chemical behavior. Problems of sampling and grade control change with the genetic type of deposit, with the stage of deposit development and with the size of the information base. Precious metal epithermal deposits (2, 6, 8), because of rapid vertical zonation and erratic lateral distribution of values, have always been difficult to evaluate and maintain grade control and ore reserves. On the other hand, evaluation and grade control are relatively easy in bulk-low- grade deposits (4, 13). However, these deposits generally have a low margin of profit and are sensitive to mining and beneficiaton costs, price fluctuations and political costs. Industrial mineral deposits (5) often must be evaluated on the basis of their behavior, rather than by chemical analysis. Environmental impact generally increases with the scale of the operation, but certain elements or minerals have especially high impact effects (7, 11). In the exploration phase there is no production control of sampling procedures and careful geological observations are particularly essential. The greatest number of problems is related to the oxidized outcrop where the chemical environment of the ore body has changed and the contained values may have been enriched, depleted or values left unchanged (2, 6). Present evidence suggests that gold values may be very mobile under certain conditions (2, 6) and stable under others. Everything must be sampled in detail. Principal values and by-product or deleterious elements may vary dependent upon their position within the soil profile. Such factors as geomorphic position, erosion rate, vegetation, climate, etc., may affect the interpretation (1, 3). During the development phase it is equally easy to overtest, to have "paralysis by analysis," as to undertest (3, 6). Bulk samplng and testing are

Jan 1, 1985

By F. A. Crossley

The titanium-rich end of the Ti-A1 system has been investigated up to 35 at. pct A1 (23 wt pet). One conzpound Ti3Al was found to occur between primary a and TiAl. It is ordered hcp with DO19 structure, it has virtually no solid-solubility range, and it has a closed maximum at about 875°C. OIL either side of the compound are a +Ti3Al two-phase fields. The limiting a1uminum solubility in primary a at the titanium-rich end is indicated to be 7.5 at. pct A1 (4.4 wt pet) at 550°C and about 6.8 at. pct Al fl wt pct) at 500°C. Quenching alloys from above the a + Ti3Al two-phase field produces the following structures with respect to alloy composition: Up to 13 at. pct A1 (7.8 wt pet), a solid solution; from 15 to 18 at. pct A1 (9 to 11 wt pct), shear transformation product or martensite; from 19 to approximately 30 at. pct (11 to 19 wt pet), submicro-scopic coherent Ti3Al in an a malvix. The twin hcp phase fields reported in the literature are the result of nonequilibrium corzdztions. Ti-A1 alloys, once partitioned by dwelling- in the a + ß phase field during either hot working or heat treatment, are extremely difjicult to homogenize at temperatures below 1000°C. Such partitioned alloys exhibit the characteristics or symptoms of two-phase materials, and may be said to suffer the "twin-phase syndrome". THE earliest investigations of the Ti-A1 system by Ogden et al.1 and Bumps et al.2 reported wide solubility of the primary solid solutions. Aluminum was reported soluble in the low-temperature allomorph to the extent of 37 at. pct (25 wt pct), and the first intermediate phase was reportedly TiA1. Somewhat later Kornilov et al.3 reported a similar diagram with phase boundaries displaced towards lower aluminum contents and higher temperatures. Beginning about this time (1956) reports in the literature made it very clear that one or more intermediate phases occurred at lower aluminum contents than TiAl.4-17 These reports included five major investigations of the titanium-rich end of the Ti-A1 diagram.4,12,14,16,17 Three of these diagrams show two two-phase fields below 37 at. pct Al, while two of them show a single two-phase field. The existence of the phase Ti3A1 is firmly established and is included in each of the diagrams, except one—that of Sato and Huang.12 The new phases are reportedly hcp and differ from primary a only slightly when disordered, and when ordered the "a" parameter is approximately one,4,12,15 two, 6-10,13,14 or four14 times that for primary a. Beyond this, however, the diagrams are remarkable for their lack of agreement. Two tacit assumptions are usually made in phase-diagram determinations of metal systems. These are: 1) equilibrium anneals bring the alloy to equilibrium or to indistinguishable closeness to it, and 2) equilibrium conditions established at elevated temperatures are either "frozen" by rapid quenching for evaluation at room temperature, or quench-transformation products are recognized as such. In the current investigation evidence was obtained that over substantial composition ranges neither of these two conditions was met in any of the more recent major investigations. I) MATERIALS, METHODS, AND TECHNIQUES The alloys of this investigation were prepared by nonc on sum able electrode arc melting. Materials used in the preparation of the alloys are summarized in Table I. The investigative tools employed were: optical and electron microscopy, differential thermal analysis (DTA), disatometry, X-ray diffraction, electron diffraction, and resistometry. Alloys for microscopic and X-ray investigations were prepared as 15-g melts. Alloys containing from 7 through 11 at. pct A1 were hot-rolled out of a furnace at 900°C, from 12 through 15 at. pct out of a furnace at 1000°C, and from 16 through 18 at. pct out of a furnace at 1125°C. Alloys containing more than 18 at. pct A1 could not be hot-rolled. The ingots were covered with Markal coating prior to hot rolling to minimize atmospheric contamination. After hot rolling, alloys containing up to 15 at. pct A1 were ground and pickled to remove 7 mils from each surface; alloys containing 16 and 18 at. pct A1 were skinned to a

Jan 1, 1967

By W. S. Gillam

A need for new supplies of fresh water exists today and in many specific areas that need is urgent. One solution lies in saline water conversion, a problem complicated by cost factors. The principles involved in saline water conversion, the status of development, and the estimated costs (present and future) of several processes are presented. Among the methods discussed are distillation, electrodialysis, and freezing. In general, the costs presented are based on a standardized procedure for estimating conversion costs, permitting a valid comparison among the various processes. The need for new supplies of fresh water and the potential benefits to be derived from an abundant supply of converted water are recognized by practically everyone concerned with water problems. The water supply problem exists today; it is urgent in many specific areas in this country and also in the world, and it will become more acute in the future. One answer to the growing problem of adequate water supplies is the development of new sources. Very significant quantities of brackish underground and surface waters exist in certain areas and an inexhaustible supply of ocean water is available. Thus in many areas water resources can be extended through saline water conversion. Congress recognized the need for new sources of fresh water in 1952 and passed the Saline Water Act, Public Law 448, amended it in 1955, and in September 1958, enacted Public Law 85-883, calling for the construction of at least five demonstration plants. The program is administered by the Dept. of Interior through the Office of Saline Water, and its primary objective is to reduce the cost of converted water produced, whether it be by development of new processes or improvement of known processes. This is a most difficult problem and one that will require several years of prodigious effort. It is difficult—not because of any intricate or new chemistry, engineering, or physics involved—but because of the difficulty in converting water at low cost. Whatever the sources of the saline water, the salts which are held tenaciously in solution must be removed before the water becomes suitable for industrial or domestic uses. Saline water is a relatively simple system of salts dissolved in water. It has certain chemical and physical properties that determine the various methods by which the salts may be separated from the water. The system, although not complex, in most instances, has had countless years in which to reach equilibrium and is, therefore, comparatively stable. Because of its stability, separation of saline solutions requires relatively large quantities of energy. The unique properties of water depend on the fact that its molecules are chemically active. The chemical and physical properties of water are associated with the type of bonding involved in the water molecule. Chemical changes such as hydrolysis, or rusting of iron, involve the breaking of chemical bonds between the hydrogen and oxygen atoms. Physical changes, such as evaporation in a boiler, the melting of ice, or the viscous resistance to flow in a pipe, involve breaking of the hydrogen bonds. (The hydrogen nucleus is so small that it can attract two negative atoms.) Thus water molecules not only combine with molecules of other compounds but even with one another; e.g., each molecule may be bounded to four other molecules. Water molecules cling to the ions of dissolved salt to form water-encumbered hydrated ions and they cling to one another to form entangling networks through which hydrated ions can be propelled only by tearing the networks apart. That is one reason why considerable energy still needs to be expended in our simplest procedures for purifying water. If water molecules did not have this habit of clinging so tenaciously to other molecules, and to one another, it would be easy to push salt ions past the water molecule and get a separation. But the water would not then dissolve salt, so the problem would not exist.' Water when heated evaporates very slowly, relative to other liquids having simple molecules. Vaporization involves the separation of molecules from the liquid, and this means overcoming the attraction between molecules which is due to the hydrogen bonding. The heat of vaporization for water is high; consequently, the boiling point of water is also high. Water boils at 100" C; hydrogen sulfide (H2S) at -60" C; oxygen (02) at -183" C; nitrogen (N2) at -196" C; and methane (CH4) at-161°C, even though the latter has about the same molecular weight as water. Because of these peculiar properties of water, it exists as a liquid on earth instead of a gas such as hydrogen sulfide or nitrogen and oxygen.

Jan 1, 1961

By R. E. Ogilvie, A. R. Kaufmann, S. H. Gelles

The solid-solubility limits of iron in beryllium were determined between 850o and 1200oC by analysis of differential type multiphase diffusion couples, using an X-ray absorption technique. The maximum value of the solubility limit was found to be 0.92 ± 0.02 at. pct (5.46 wt pet) at the eutectic temperature 1225°C. The solubilities of nickel and beryllium were determined between 900°and 1200°C by the same technique and the maximum solubility was found to be 4.93 + 0.01 at. pct (25.2 wt pet) at the eutectoid temperature, 1065°C. A previously unreported high-temperature phase which decomposes eutectoidally at 1065 °C was found to exist in the beryllium-nickel system at a composition of approximately 8 at. pct Ni (36 wt pet) by diffision-couple analysis. The presence of this phase was confirmed by thermal analysis and metallo-graphic analysis of the structure resulting from the eutectoid decomposition. G. V. Raynor1 has treated the solid solubilities of some of the elements in beryllium on the basis of the "Hume-Rothery" rules2 which have been modified to include ionic size and ionic distortion effects. It was predicted that the solubility of iron and nickel in beryllium should be slightly less than that of copper. The lowering of the solubility, according to Raynor, is due to a more unfavorable relative valency effect and an ionic size effect. Kaufmann and corzine3 have compiled data on the solubilities of elements in beryllium and have discussed them in the light of the Raynor paper. These authors feel that, because the elements having the greatest solubility in beryllium systematically fall in the Group VIII and IB Columns of the periodic table, the electronic structure greatly influences the maximum solid solubility of elements in beryllium. The solubility of iron in beryllium was determined by Teitel and cohen4 as part of the study of the beryllium-iron phase diagram. The determination was carried out by X-ray and thermal analysis and according to the phase diagram presented, the maximum solubility of iron in beryllium is 0.41 at. pct (2.5 wt pct) at 1225oC. However, it is estimated that the uncertainty in the position of the a-beryllium primary solid-solution boundary is about 0.5 at. pct (3wtpct). Losana and Goria3 in studying the beryllium-nickel phase diagram, determined the solid solubility of nickel in beryllium by thermal analysis. They found the maximum solubility to be between 1.65 and 2.65 at. pct (10 to 15 wt pct) at 1240°C. This value decreased rapidly with decreasing temperature. In determining approximate ranges of solubilities for different elements in beryllium, Kaufmann, et al,8 reported a value of between 1.3 and 1.7 at. pct (7.9 to 10.1 wt pct) for the solubility of nickel in beryllium. The value was obtained by metallographic examination of quenched alloys and lattice-parameter measurements. However, the authors also noted a single-phase structure for a 1.7 at. pct Ni alloy (10 wt pct) on cooling from the liquid. This would indicate a higher solubility range than was reported. ~isch,' in his X-ray studies of beryllium-copper, beryllium-nickel, and beryllium-iron intermetallic compounds, reports the disappearance of a second phase (Ni,Be2) in the beryllium primary solid solution at approximately 4 at. pct (20 wt pct). THEORY The analysis of concentration gradients in diffusion couples has proven to be a useful tool in determining phase equilibria.8-14 In this particular study the diffusion couples were chosen to straddle the expected composition range of the phase boundary, then heat treated at a given temperature and the concentration gradient evaluated. The composition of the phase boundary for a given temperature appears at a point of discontinuity of the composition gradient. Examples of typical phase diagrams and the concentration gradients which should be found in such systems are shown in Fig. 1. In the present work, gradients of the form of Fig. l(c) were obtained in diffusion couples made of pure beryllium and two-phase alloys of beryllium with either iron or nickel. The composition at the point where the gradient becomes discontinuous, Cs, corresponds to the solubility limit of either iron or nickel in beryllium. The analysis of the concentration gradients was carried out by an X-ra absorption method developed and applied by Ogilvie and later used by Moll13 and Hilliard.l4 It depends on the fact that the absorption of X-rays by matter is determined by the concentration and type of the various atomic species present. The relationship for the intensity, I, of a monochro-

Jan 1, 1960

By A. W. Adamson

A model for transport processes in liquid mixtures is discussed which supposes that the elementary act involves a position exchange between two species and that the exchange is so confined by the solvent cage as to occur nearly isosterically. The rate-determining step, thus, is likened to a bi-molecular reaction and is so treated, using absolute rate theory. The cage model has been applied to diffusion, thermal diffusion, sedimentation and viscosity, but only the first and last of these phenomena are emphasized in the present paper. The model leads to semi-empirical relationships between the absolute value for a digusion coefficient and the activation energy for diffusion, between mutual and self-diffusion coefficients and for the variation of the viscosity of a binary mixture with composition. These are discussed in relation to experimental data for various systems, including hydrocarbon mixtures. It is shown that the proposed viscosity equation and seven other commonly used ones all may be regarded as special cases of a single general relationship; a brief critical analysis is made of the basis of selection of one or the other for data fitting or interpolation. INTRODUCTION AND GENERAL THEORY The present paper covers a brief discussion of a cage model for transport processes in liquid mixtures and how this model may be useful in treating the diffusional behavior and the viscosity of such systems. Since diffusion requires the more detailed treatment, it will be taken up first, and the model then applied to viscosity. There are two types of diffusion coefficients that may be measured experimentally, apart from thermal diffusion quantities. The first is the mutual or binary diffusion coefficient, D which may be defined in terms of Fick's first law. This states that the permeation, or flux P, is proportional to the concentration gradient. In the usual experiment, P is measured relative to a frame of reference fixed with respect to the medium (e.g., the diaphragm in a diffusion cell); as a consequence, the same value of D is obtained regardless of whether P and C refer to Component 1 or to Component 2; i.e., there is only one independent mutual diffusion coefficient for a binary system. In addition to D there will be various self-diffusion coefficients. defined in terms of the gradient in labelled species i and its permeation in an otherwise uniform medium. The thermodynamic approach to mutual diffusion supposes that the actual driving force is the gradient of the chemical potential, i.e., that In the case of a dilute solution of solute, Eqs. 1 and 3 lead to the Einstein equation, If the solution is ideal and the friction coefficient is taken to be then the familiar Stokes- Einstein equation results. Mutual and self-diffusion coefficients can not be related on general thermodynamic grounds; it is necessary to invoke some additional assumptions, i.e., a model; several such have been proposed. Hartley and Crank' supposed the existence of separate, intrinsic diffusion coefficients (Dl and D2) for each component, essentially corresponding to the two self-diffusion coefficients. The two flows can not be independent, however, but must be coupled through the usual restriction that there be no net volume flow. For an ideal solution. one then obtains' Glasstone, et al' treated diffusion in terms of absolute rate theory, but their approach otherwise resembled the previously mentioned one in that each species was considered to move with respect to the general medium in a manner determined by its individual jump distance and specific rate constant. For other than dilute solutions, a coupling of flows leading to an equation such as Eq. 6 would again be present. However, as required by Eq. 6, one does expect that the self-diffusion coefficient for the solute and the mutual-diffusion coefficient for the system become identical at infinite dilution. Lamm4 recognized that there should be three distinctive interactions in a two-component system-1-1, 1-2 and 2-2 — and, therefore, proposed three rather than two fundamental friction coefficients. Mutual diffusion resulted from 1-2 interactions only, and self-diffusion resulted from 1-2 plus either 1-1 or 2-2 interactions. Again, a collective coupling between all motions was imposed to meet the condition of no net volume flow. Laity' has shown how to convert the Onsager equations to a form very similar to Lamm's. Cage Model For Diffusion Work in this laboratory on diffusion in aqueous sucrose solutions made it apparent that three, rather than two, interactions were indeed needed," but considera-

By J. W. Clark, C. T. Sims

Wrought chromium-base alloys containing yttrium, cubic monocarbides of the Ti(Zr)C type, and similay alloys containing manganese and rhenium have been melted and fabricated. Strength has been studied by hot hardness and elevated-temperature tensile and rupture measurements, low-temperature ductility by tensile testing, and surface stability by oxidation testing. In additiod, studies have been conducted of the carbide stability, and of aging behavior. The carbide dispersion generates effective elevated-temperature strength, which is further enhanced hv strain-induced precipitation. The dispersion exhibits classical dissolution and aging response. The ductile-to-brittle transition temperature of these alloys is above room temperature. The alloys reported show fairly good oxidation resistance, but nitrogen contamination can cause fortnation of a hard Cr2N layer under the oxide scale. Manganese does not appear to be a promising alloying element in chromium. In the years 1945 to 1950, the metal chromium was considered as a possible base for alloy systems due to its considerably higher melting point than superalloys, its low density, its high thermal conductivity, and its apparent capacity for strengthening. However, this interest in chromium was short-lived. It was found difficult to melt and cast, to be exceptionally sensitive to the effect of minor imperfections, to have a lack of ductility at both room and elevated temperatures, and to be subject to a deleterious effect of alloying elements upon the ductile-to-brittle transition temperature.' Since then, chromium, as a practical alloy base, has remained virtually unstudied. Further, purposeful ignoring of chromium has been promoted by statements that its bcc structure would not allow it to be strengthened to useful values, when compared to the "austenitic" alloys.2 Recently, a new look has been taken at chromium-base alloy systems. Study of the literature will show that chromium, providing some of its disadvantages could be eliminated or minimized, actually has a rather attractive potential as an alloy-system base. Analysis of rather scattered data suggests that chromium is quite capable of being strengthened to high levels. Also, significant strengthening of its two sister elements in Group VI-A, molybdenum and tungsten, has been demonstrated in a number of commercial and exploratory alloys. Chromium should be similar. Since chromium does not readily form a volatile oxide like tungsten or molybdenum, it offers a much higher probability of giving birth to alloy systems with useful oxidation resistance. Concerns about possible high elemental vapor pressure have been mitigated by recent data.3 In addition, the physical properties exhibited by chromium are attractive for application as a high-temperature structural material. For instance, its thermal conductivity varies from 49 to 36 Btu-ft/hr-sq ft-°F over its range of usefulness (which is two to four times higher than most superalloys), its density is about 7.2 g per cc (20 pct less than most nickel-base alloys), its coefficient of thermal expansion varies from 4 to 8 x 10-6 per OF, and it has a relatively high modulus of elasticity, approximately 42 x 10' psi.4 Alloying studies on a chromium base in the past have usually encompassed rather sweeping solid-solution alloy additions for strengthening. This is not consistent with contemporary alloying practice in Group VI-A. For instance, molybdenum, also in Group VI-A, is primarily alloyed for strength improvement by use of heat-treatable carbide dispersions.5 Chromium and molybdenum are similar in their chemical activity and other properties. Thus, strengthening of chromium by carbide dispersions was studied. Chromium-base alloys are plagued with room-temperature brittleness, although high-purity unal-loyed chromium can be made ductile.4,8 Use of yttrium as a scavenger has done much to improve ductility and resistance to nitrogen embrittlement in chromium systems,7 so it was utilized in this program. It has also recently been found8 that small rhenium additions (1 to 5 pct) create improvement in the ductility of Type 218 tungsten wire. This is apparently related to the remarkable effect of rhenium additions near its terminal solid solubility in all Group VI-A metals.9'10 Investigation to establish if dilute concentrations of rhenium would also be effective in chromium appeared to be logical for this program. Since rhenium is too expensive to be practical in alloys for application as structural components, ductility improvements through solid-solution alloying were also sought by substitution of manganese for rhenium; manganese, like rhenium, exists in Group VII of the periodic system. The optimum amount of carbide dispersion for chromium-base alloys was obtained by analogy with molybdenum. Strengthening in molybdenum is achieved by use of Ti-Zr carbide dispersions. A

Jan 1, 1964

By C. M. Loeb, A. M. Gaudin, J. H. Brown

THE object of most size reduction operations in the mineral industry is to liberate the grains of valuable minerals in the ore from those of the gangue. This is usually accomplished by crushing and grinding the entire mass of ore until there is only a small probability that any single particle contains more than one mineral. During this size reduction only limited control exists over size or composition of the particles exposed to the breaking action, and there is no control over the paths followed by cracks generated during the operation. This lack of control usually results in overgrinding and in production of large quantities of very fine material. The first detriment, overgrinding, is costly in itself, but when combined with the second factor it is doubly so. Not only is the fracture of a free particle unnecessary—the fracture of these particles may also make subsequent separation operations difficult, inefficient, and wasteful. It has been pointed out' that if the object of size reduction is to liberate the valuable mineral component of the ore then, ideally, fracture should follow intergranular paths to the exclusion of trans-granular ones. This would result in liberation of the valuable minerals with as little size reduction as possible. This ideal comminution operation is referred to as intergranular comminution, and it was the object of the investigation to determine the extent to which it could be developed by heat treatments. There are many indications in the literature that heating rocks prior to crushing may be favorable. Reports by Holman,2 Yates3 and Myers' are pertinent. These investigators showed that heating certain rocks prior to crushing them did, in fact, improve their crushing characteristics in that fewer fines were produced, although the fact that intergranular comminution was being effected apparently was overlooked. In addition, Sosman noted that if there is appreciable anisotropism in the thermal coefficients of expansion of even a pure mineral, then considerable permanent separation of the grains of the rock can be expected as a result of heating the rock to a high temperature.' By the same token, if there are ap- preciable differences in the thermal expansion coefficients of the various minerals of a multi-component rock, similar results should be obtained by heating this rock. This has been tested, partially, by Brenner," who obtained patents covering the heat treatment of some pegmatitic rocks in order to facilitate comminution of these materials. It has also been demonstrated that this may occur in taconite." Also, the possibility of causing decomposition of one mineral in a rock as a means of promoting intergranular fracture has been considered. Seigle2 and Schiffman et al. have obtained patents on such processes as applied to calcareous iron ores. These reports all indicate that heat treatments prior to crushing may contribute materially to intergranular comminution, but they also indicate that no organized attempt has been made to determine the controlling factors of the method or to determine its applicability in general. The present article is a report on the initial phase of such an investigation. The authors have reviewed the claims of prior investigators and have attempted, also, to establish the factors that might determine the applicability of heat treatments in the mineral industry. In this work 2000-g samples of various rocks were heated in a small laboratory furnace and crushing and sizing operations were carried out in standard laboratory equipment. All samples of each rock were as nearly identical as possible in particle size, grain size, and composition and contained only lumps coarse enough to contain many grains each. Tests on Granite A number of tests were made on a coarse grained Finnish granite obtained in the form of coarse chips from a local monument yard. This rock exhibited little variation from piece to piece in either composition or grain size. The minerals contained were quartz, orthoclase, small amounts of hornblende, and minute quantities of mica. Grain size ranged from about 1 mm to about 3 mm. Temperature of the Heat Treatment: In some cases the granite was heated to a particular temperature and crushed, hot, immediately upon withdrawal from the furnace—in others the rock was allowed to cool before crushing, but without quenching to room temperature after heating. In most tests on granite the heating period was about 2 hr with the furnace at the highest temperature for about 1 hr. Cases in which these periods were varied greatly will be presented separately.

Jan 1, 1959

By C. J. Guntner, R. P. Reed

A commercial 18Cr-8Ni iron alloy (AISI 304L) was examined in tension at 300°, 76°, 20°, and 4°K. Continuous stress-strain recordings were made, X-ray analyses at periodic stress (strain) intervals were obtained, and the magnetic measurements were taken. From this data the percentage of martensitic products [bcc(a) and hcp (E)] was computed as a function of stress (strain). It was found thatup to 15 pct E phase forms at low temperntures. The amount of E formed increases to a maximum at about 5 pct strain, then decreases. This decrease indicates the additional transformation of E to a'. The total amount of E and a' was suppressed at constant stress (strain) at 4°K as compared to 76°K. It is proposed that the suppression of E and a' is associated with the decreased mobility of extended dislocations at very low temperatures. The yield strength decreased as the temperature was depressed below room temperature and then increased rapidly near 4°K. SOME ferrous alloys which are austenitic (fcc ?) at room temperature appear to be unique in that two martensitic products (hcp e and bcc a') may form on cooling to lower temperatures or on application of mechanical stress. The most common room-temperature austenitic ferrous alloys are 18Cr, 8Ni stainless steels. Most aspects of the spontaneous transformations have been previously described for these steels.' Several previous papers have described special aspects of the stress-induced transformations at low temperatures for the stainless steels, such as the existence of the hcp phase (c) after straining at 76oK,2-7 the morphology after straining using electron microscopy,7 and the decrease in E at higher strains at 76oK.4 However, for a complete representation, one must know the stress-strain characteristics and the dependence of both martensitic products on applied stress and temperature. It is the intention of this paper to provide that documentation. To accomplish this, continuous stress vs strain recordings were made at four temperatures: 300°, 76", 20°, and 4°K for annealed AISI 304L (a commercial 18Cr-8Ni alloy). At periodic stress intervals at each temperature the integrated X-ray line intensity of a selected peak for each phase (y, E, and a') was measured. In addition, photomicrographs of the strained surfaces were taken and magnetic measurements were made. The magnetic readings can be directly converted into percent a'.',e With these measurements the percentage of each phase may be plotted as a function of stress (or strain) and test temperature. It was found that up to 15 pct E phase forms upon stressing the AISI 304L alloy at low temperatures. The E percentage increases abruptly after the alloy yields, but then decreases gradually at higher stresses. The rapid increase in e at 76°K is associated with an "easy-glide" portion of the stress-strain curve. The total amount of a' + .G is suppressed below 76°K at a constant stress or strain. The yield strength decreases down to 76°K but increases rapidly below 20°K. EXPERIMENTAL PROCEDURE Tensile test specimens were cut parallel to the rolling direction from 0.1-in.-thick sheet. Continuous stress vs strain recordings were obtained at each test temperature (300°, 76o, 20°, and 4°K) using equipment and methods described elsewhere.' The specimens which were used in the X-ray analysis were stressed to successive increments of strain at each temperature, analyzed at room temperature, then restressed at the test temperature. This procedure was repeated until approximately ten X-ray analyses had been performed with approximately 1.0 pct strain increments. The specimens had a reduced section 1 in. long, 1.2 in. wide, and 0.1 in. thick. They were electro polished prior to testing and after each strain increment. Table I lists the chemical composition, grain size, and hardness for the alloy which was used. This is the same alloy for which extensive mechanical-property tests3 and morphological studies of the spontaneous transformations' have previously been made. For the low-temperature tests (76o and 4°K) below the Ms temperature the specimens were initially cooled to the test temperature, held for 1/2 hr, then warmed and X-rayed at room temperature. The results are listed in Table 11. From earlier work8 it was known that additional transformation on the second cycle would be considerably less (-0.1 pct

Jan 1, 1964

By M. Shelef, A. W. Fletcher

The effect of alkali sulfates in promoting the sul-fation of nickel and copper in a bulk sulfide flota -tion concentrate by fluidized bed roasting has been studied in the laboratory, and it was shown that the various alkali sulfates promote sulfation to approximately the same extent. The sulfation of a mixture of synthetically prepared iron and nickel oxide and of nickel ferrite has also been studied. Nickel sulfation was promoted by high ratios of Fe:Ni and by the presence of sodium sulfate. THE work described in this paper was a continuation of earlier studies into the role of alkali sulfates in promoting the sulfation roasting of nickel sulfides1,2 in an endeavor to determine how the system was affected by the presence of compounds of iron and copper. The earlier work1 showed that, in the sulfation of NiO at 680°C, the reaction was limited by the formation of an impermeable film of nickel sulfate on the oxide surface. The relative effect of the various alkali sulfates in promoting nickel sulfation varied in the order: Li > Na >Cs > Rb > K A study of alkali sulfate/ nickel sulfate interactions at high temperatures showed that the promoting action was due to the fact that the nickel sulfate product layer sintered and agglomerated only when the more active additives were present. This resulted in the formation of discontinuities in the nickel sulfate layer so that diffusion of the sulfating gases to the NiO surface was no longer impeded and the reaction could proceed to completion. A similar explanation was used for the observation that sodium and lithium sulfates promote the oxidation of NiS to NiO at temperatures below 750°C since small amounts of nickel sulfate were formed during oxidation.2 It was of interest to study the effect of alkali sulfates on the sulfate roasting of a sulfide flotation concentrate which is typical of material treated commercially. In order to control temperature it is essential to roast sulfides in a fluidized bed and this technique was therefore used, although the batchwise operation of a small-scale laboratory reactor does not reproduce all conditions which prevail in full-scale continuous plant. The results obtained are therefore only comparative, and cannot be used for predicting the optimum conditions for metal extraction. The sulfation of synthetically prepared mixed oxides of nickel + copper and nickel + iron and of nickel ferrite was also studied to evaluate the relative effects of alkali sulfates with more complex systems. SULFATION ROASTING OF A SULFIDE FLOTATION CONCENTRATE The bulk sulfide flotation concentrate used in this work contained 7.92 pct Ni, 1.74 pct Cu, 35.66 pct Fe, and 31.28 pct S. The sulfide minerals present in order of abundance were pyrrhotite FeS, pyrite FeS2, pentlandite (FeNi)S, and chalcopyrite CuFeS2. Two samples described as coarse and fine were used. The coarse sample, which was a flotation concentrate (58 pct plus 300 mesh), was ground to 100 pct minus 350 mesh to produce the fine sample. Before roasting, the sample of sulfide concentrate was agglomerated by wetting witli a solution of the alkali sulfate (or water), thoroughly mixing, and drying at 110°C. This gave a cake which was gently crushed and screened, the -18 +100 mesh fraction being used for fluidized bed roasting. A similar-size fraction had been used by the authors in pilot plant work with a 4-in.-diam fluidized bed reactor.' In this work it was found that the molar ratio of additive to the total iron + nickel + copper content of the sulfide sample should be adjusted to a value of approximately 0.06, as this was the optimum amount necessary for nickel sulfation. Experimental. The fluidized bed reactor consisted of a quartz tube approximately 60 cm long and 30 mm in diameter resting in a vertical tube furnace. The sulfide bed (30 g) was supported on a bed of -4 +12 mesh quartz particles 3 cm high, which rested on a sintered quartz disc welded to the tube. The temperature of the furnace was controlled with a variable transformer to give a final bed temperature of 680°C. The bed was fluidized with air or mixtures of air + 10 pct v/v SO2, at a total apparent gas velocity of 60 to 65 cm per sec at 680°C. The SO2 was introduced into the fluidizing air stream only when the oxidation of the sulfides was completed. At the end of the roasting period the calcine was leached with boiling water and the

Jan 1, 1964

By J. Philibert, A. G. Guy

Diffusion in the presence of deformation was studied by the method of vacuum dezincification of copper-rich and silver-rich solid solutions containing 7 to 30 pct Zn. The specimens were designed to permit the study of diffusion in separate portions of a given specimen characterized by strain rates ranging from essentially zero to approximately 10 sec-. No effect of deformation on diffusion was observed. BEGINNING with the work of Buffington and Cohen: interest in the question of the effect of stress or strain on diffusion has largely been concentrated on the enhancement of diffusion in specimens subjected to Continuous plastic deformation. The present research is a contribution to this limited area. However, as a preliminary to focusing attention on this special topic, it will be desirable to make a broad survey of the larger question, especially since there has been considerable foreign work in areas outside those of current interest in the United States. Since most of the topics referred to in the following section are both complex and imperfectly understood at present, it has been expedient in most instances to offer only a guide to the general nature of the work rather than a critical evaluation. PREVIOUS WORK The effect of elastic stress on diffusion has received considerable attention, especially with regard to the thermodynamic driving force for diffusion. The thermodynamic treatments have been based on the work of Gibb, Voigt, Planck, and Leontovich.' Konobeevskii and Selisski6 made a first attempt at treating the problem in 1933, and Gorskii7 a few years later gave a solution applicable to single crystals as well as to polycrystalline specimens. In 1943 Konobeevski8 published treatments that have been the basis of much Russian work up to the present. For example, Aleksandrov and Lyubov used his work in explaining the velocity of lateral growth of pearlite. Early work in the United States was that of Mooradian and Norton, which showed that lattice distortion tends to be relieved before it can significantly affect the diffusion process. Druyvesteyn and Berghoutl1 observed a slight effect of elastic strain on self-diffusion in copper, while de Kazinczy12 found that both elastic and plastic deformation increased the rate of diffusion of hydrogen in steel. On the other hand, Grimes58 observed no effect of either elastic or plastic straining on the diffusion of hydrogen in nickel. High-frequency alternating stresses have been reported by various investigator s13-l5 to increase the rate of diffusion. A special form of elastic stressing is the imposition of hydrostatic pressure, a condition that is amenable to Conventional thermodvnamic analysis. Most of the experimental results in this area are consistent in showing a slight decrease in diffusion rates at high pressures.16-l8 Although Geguzinl reported a pronounced effect of relatively small pressures, Barnes and Mazey20 failed to Corroborate this finding, while Guy and Spinelli21 advanced an explanation of the phenomenon observed by Geguzin. It has been recognized that the thermodynamic treatment of diffusion phenomena in an arbitrarily stressed body is complicated by the fact that the desired state of quasi-equilibrium of the shear stresses cannot be maintained during a general diffusion process. However, attempts have been made by Meix-ner22-24 and Fasto to treat certain restricted cases, such as relaxation. FastovZ7 has also incorporated the general stress tensor into the thermodynamics of irreversible processes. The lattice strain that accompanies the formation of a solid solution has been the subject of much study,28-s0 and indirectly it has entered into many recent theories of diffusion. However, some Russian investigators31'32 have taken other views of this matter and have predicted large effects on diffusion rates because of concentration stresses.o In completing this brief resume of previous work involving elastic strains and before proceeding to a consideration of the effect of continuous plastic deformation, it should be pointed out that deformation of various additional types may also influence diffusion. The effect of cold-working on subsequent diffusion has been studied directly by AndreevaS and by Schumann and Erdmann-Jesnitzer, while indirect evidence has been obtained by Miller and Guarnieri and by Vitman.38 Thermal stresses may also influence diffusion, contributions to this subject having been made by Fastovs7 and by Aleksandrov and Lyubv. The work of Johnson and Martin,o Dienes and Damask,3Band DamaskS considered the question of radiation-enhanced diffusion. In considering previous work on the subject of plastic deformation and diffusion, attention will be directed to those studies concerned primarily with diffusion rather than with its relation to Creep, e.g., the work of Dorn, or to the acceleration of diffusion -controlled reactions. Observations of the effect of

Jan 1, 1962

By W. A. Oates, J. A. Lambert, P. T. GaIIagher

Monte Carlo methods have been used to compute the arrangements of interstitial atoms dissolved in tetrahedral sites in bcc lattices. It is assumed that the presence of an interstitial atom "blocks " a certain number of neighboring sites and prevents their occupancy. Sites "blocked" by more than one filled site are allowed for. The computed values of. the mean occupation number (defined as the ratio of the total number of sites blocked to the number of solute atoms are used to calculate the configurational entropies of the solutions. These entropies are compared with those resulting from previous theoretical studies of this problem and also with available experin~ental data for the p Zr-H, Nb-H, V-H, and Ta-H systems. Evidence is also given that the "blocking" explanation of low limiting compositions in these systems, rather than this being due to initial limitations on the number of sites available, is probably correct. THE ideal partial configurational entropy of mixing of an interstitial solute in a metal is given by: where p is the number of interstitial sites per metal atom and Xi is the atomic fraction of the interstitial. For the bcc lattice. which we shall be concerned with in this paper, the interstitial positions are shown in Fig. 1. It can be seen that for the tetrahedral sites, p=6. whereas for the octahedral sites, p = 3. Different emphasis has been placed on the relative importance of energy and entropy effects in determining deviations from ideality in interstitial solid solutions. In some cases the same system, e.g., Fe-C, has been described by the contradictory regular and athermal solution models indicating that the enthalpy and entropy functions, derived from equilibrium data, are frequently not accu.rate enough to differentiate between these treatments. However, for certain metal-hydrogen solutions the equilibrium data is available over sufficiently wide ranges of temperature and composition to permit a reasonably accurate determination of the compositional variation of the heats and entropies. Hoch&apos; has attempted to interpret the results of interstitial solid solutions in terms of a regular solution model. In the case of the Ta-H system where 13 = 6, this model entails fitting the experimental relative partial entropies of solution, asH, to the equation: where ASgs is the relative partial excess entropy of solution of hydrogen. Hoch found that the results of Mallett and Koeh1 could be fitted to this equation with an approximately constant value of AF up to XH = 0.25. However, it is apparent from the solubility isotherms in this system which become asymptotic to the composition TaH that, since (Xh /6 - ~Xh ) becomes infinite only at TaH6, it is necessary that AS<&apos; tends to infinity at TaH. In other words, the low saturation composition of TaH, instead of the anticipated TaH,, eliminates the possibility of applying regular solution theory to such systems. Rather large negative excess configurational entropies must exist at higher hydrogen concentrations in order to explain the lower saturation values. To account for these low limiting compositions and excess entropies two distinctly different approaches have been followed. Rees and many others1-l2 have assumed that not all interstitial sites are crystallographically equivalent with respect to the interstitial addition; that is, in Eq. [I] p is less than the value anticipated from geometrical considerations. To describe, say, a bcc metal-hydrogen system with a limiting composition of MH by this approach one would consider that p = 1 in the first instance instead of p = 6.&apos;j3 In some cases, nonintegral values of B have been taken in order to improve the fit with the experimental data over limited ranges of composition. The other approach which has been used to explain the low saturation compositions is to assume that, although all sites are available for occupancy, strong repulsive interactions exist between the neighboring interstitial atoms, and hence occupancy of any site excludes or blocks a certain number of neighboring sites from being occupied. Earliest treatments of this concept considered the exclusion of an integral number, of nearest-neighbor sites from being occupied at all concentrations. In this case, the partial configurational entropy is given by: These early treatments failed to allow for the overlap of the blocked sites which will arise at all but the very lowest concentrations. More recently attempts have been made to calculate the effect of this decrease in the number of blocked sites on the configurational entropy. Using the quasichemical treatment of interstitial solid solutions as given by Lacher and assuming that an infinite repulsive interaction energy existed between the solute atoms. atom obtained an approximate configurational entropy applicable to the blocking with overlap case:

Jan 1, 1970

By S. T. Wlodek

The surface and subscale oxidation reactions were followed by means of continuous weight-gain and metallographic techniques over the range 1600" to 2200°F (871° to 1204 °C) for up to 400 hr. Full identification of all scale and subscale reaction products was obtained by electron and X-ray diffraction. At or below 1800°F (982°C) a linear rate of reaction (QL = 46.0 kcal per mole) governed the oxidation process, extending for up to 100 hr at 1600°F (871 "C). During linear oxidation the surface scale consisted of an amorphous SiO2 film overgrown with Cr 2O 3 and NiCr204. This initial linear process was followed, and above 1800°F completely replaced, by two successive parabolic rate laws (Qp = 60 and 57 kcal per mole). This parabolic reaction involved the formation of a complex scale consisting of Cr2 O3 and smaller amounts of NiCr2O4. Parabolic oxidation appeared to coincide with the disruplion of the silica film present during linear oxidation and was followed by subscale (internal) oxidation of crystobalite and NiCr2O4. The balance between the subscale and surface oxidation reactions controls the oxidation of this commercial alloy. The amorphous silica film appears to result in the linear rate and diffusion through Cr2O3 is the more likely rate-limiting step during parabolic oxidation. THE oxidation of a multicomponent composition is a complex phenomenon not presently amenable to a rigorous classical interpretation. Nevertheless, even a qualitative understanding of the scaling and subscale reactions that occur in a commercial composition can illuminate the reactions that limit its high-temperature stability in an oxidizing environment. This study of the oxidation of Hastelloy Alloy X presents the first of a series of studies with the above approach in mind. Hastelloy X exhibits one of the best combinations of strength and oxidation resistance available in a wrought, solution-strengthened, nickel-base alloy. Although during long time exposure some precipitation of M6C and M23C8 carbides as well as a complex Laves phase occurs, the amounts are probably small enough to have no appreciable effect on the chemistry of the matrix. Radavich has identified the oxidation products on Hastelloy X oxidized for 5 min to 10 hr at 1115°F as NiO and the NiCr2O4 spinel. Oxidation for 5 to 15 min at 1500°F produced a scale of spinel, NiO, and a rhombohedra1 phase, probably Cr2Os. Sannier et 2. have reported continuous weight-gain data for Hastelloy X at 1650" and 2010°F and internal-oxidation measurements after 150 hr at 2010°F. In addition, much of the data on binary Ni-Cr alloys recently reviewed by Kubaschewski and okins&apos; and Ignatov and Shamgunova4 as well as studies of binary Ni-Mo alloys5 are also pertinent to the oxidation of this composition. EXPERIMENTAL Continuous weight-gain measurements and metallographic measurements of subscale reactions were the main experimental techniques used in this study. X-ray and electron diffraction backed up by a limited amount of electron-microprobe analysis served to characterize the nature of the scale- and subscale-reaction products. Two heats of commercial sheet of the composition given in Table I and identified as A and B were used in the bulk of this study. Internal-oxidation measurements were made on a third heat of material in the form of a 0.5-in.-diam bar. In order to assure homogeneity, all heats were reannealed 4 hr at 2175°F prior to sample preparation. weight-Gain Measurement. All specimens (1.5 by 0.4 by 0.03 in.) were abraded through 600 paper, electropolished, and lightly etched in an alcohol-10 pct HCl solution. An electrolyte of 150 cu cm H,O, 500 cu cm HsPO4 (85 pct conc), and 3 g CrO3 at a current density of 0.9 amp per sq cm or a solution of 10 pct HaW4 in alcohol used at 4 v and 0.3 amp per sq cm was used for electropolishing. The resultant surface exhibited a finish of 3 ± 1 p rms. Continuous weight-gain tests were made at 1600°, 1700°, 1800°, 1900°, 2000", and 2200°F on auer&apos; type balances capable of recording a total weight change of 110 mg with an accuracy of k0.1 mg. All tests were made in air dried to a dew point of -70°F and metered into the 2-in.-diam reaction

Jan 1, 1964