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By M. C. Flemings, R. Mehrabian, G. E. Nereo

Analytic expressions were developed and applied in two preuious papers to predict effects of solidification variables on macrosegvegation. In this paper, experiments are reported to test the analyses quantitatively and qualitatively. Good quantitative agreement is obtained between predicted and measured compositions for the following types of segregation (all experiments were on laboratory-size ingots of Al-4.5 pct Cu alloy): (a) inuerse segregation, (b) composition distribution in ingots solidified with unidirectional heat and fluid flow, (c) simulated centerline segregation, and (d) segregation resulting fro, change in solidification cross section. In addition, qualitative agreement of prediction of theory with experiment is obtained for (a) segregation resulting from sudden change in heat extraction during solidification, (b) centerline segregation in a bidirectionally solidified ingot, and (c) under-riser positive segregation. IN two previous papers,''2 it was shown that a variety of apparently different types of macrosegregation result from the same basic mechanism. This mechanism is the flow of solute -rich liquid to feed solidification and thermal contractions. Analytic expressions were derived and examples given of their application to predict macrosegregation. Calculations were for Al-4.5 pct Cu alloy. In this paper, we describe a series of experiments designed to test results of the foregoing calculations. All experiments were on laboratory-size ingots (weighing 10 lb or less), of A1-4.5 pct Cu alloy, nominal composition. APPARATUS AND PROCEDURE A typical mold and chilling arrangement used for unidirectional solidification is shown in Fig. 1. Before casting, this assembly was placed in an electrically heated air furnace, and heated to 680" to 700°C. When the plaster reached the desired temperature (usually 1 to 2 hr after the furnace temperature), the mold was filled through a funnel extending from the outside of the furnace into the cavity. The assembly, with liquid, was held for several minutes to allow convection to subside. The coolant, either air or water (depending on the desired solidification rate), was turned on and the ingot solidified. With the four sides and top insulated, heat removal for solidification was through the bottom surface, resulting in a fully columnar structure. Design of the casting cavity was varied, as discussed below, to produce desired macrosegregation effects. However, for castings solidified with unidirectional heat flow, the top was always open to the furnace atmosphere and temperature. The bottom sur- face rested on a chill through which the heat was extracted, and all other surfaces were insulated with plaster. In one case to be discussed, the mold design was changed to obtain bidirectional solidification. In this case, two vertical chills were used (on two faces of the mold) and the bottom was insulated. In a last case discussed, that dealing with under-riser segregation, experiments were conducted in bottom-chilled sand molds at ambient temperature. Except for this one case, all experiments were conducted in the electrically heated air furnace as described above. In many ingots, thermocouples were used to obtain continuous records of solidification. In all cases, these were chromel-alumel couples with output continuously monitored on a twelve-channel recorder. Where thermocouples were employed, these were embedded in the refractory mold walls during mold-making so their heads extended to the vertical centerline. Six to nine couples were employed, placed at intervals from the chill. METAL-MEL TING PROCEDURE Aim analysis was A1-4.5 pct Cu; alloy was prepared from high-purity aluminum (99.9 pct Al) and A1-Cu master alloy (50 pctOFHC Copper, 50 pct high-purity Al). The A1-4.5 pct Cu alloy was cast in 5- to 7-lb ingots for subsequent remelting for the macrosegregation experiments. Melting for ingot casting was as follows. After preheating the melting crucible for an hour or more, the charge was melted and superheated to a temperature of 690" to 720°C. After stirring, the melt was degassed with chlorine for 10 min. After stirring again for 1 min, a gas sample, under reduced pressure, was taken. When the sample was gas-free, the crucible

Jan 1, 1969

By O. N. Carlson, H. A. Wilhelm, J. M. Dickinson

On the basis of microscopic studies, melting-point observations, and X-ray analyses, a phase diagram is proposed for the Cb-V system. A complete series of solid solutions is formed with a minimum in the solidus at 1810°C near 35 wt pct Cb. No compounds or intermediate phases were found in the system above 650°C. THERE is an ever increasing need for better structural metals and alloys for use in nuclear reactors. In addition to the normal properties of engineering structural materials, such as high temperature strength, resistance to corrosion, and fabric-abil~ty, the nuclear properties of the material must be considered. In a nuclear reactor it is important to conserve neutrons, so a material which removes these neutrons from the reaction excessively is considered to have unfavorable nuclear properties. In nuclear-reactor design the engineer must have nuclear as well as other data available on alloys in order to make a wise selection of materials. Due to the fact that many of the common structural materials have undesirable nuclear properties, it is vital that new alloys of metals having more favorable nuclear properties be investigated. Columbium and vanadium are both high melting metals, both exhibit resistance to chemical attack, and no great difficulty is encountered in fabricating them into desired shapes. With proper treatment both metals can be cold rolled extensively without failure. In addition they have desirable nuclear properties for certain types of reactors. Therefore, the alloys of columbium and vanadium should be of interest in the atomic energy program. Since an alloy-development program is enhanced by a knowledge of the phase equilibria of the components, this investigation was undertaken to establish the phase diagram for the Cb-V system. According to the Hume-Rothery rules for alloying,' the chemical similarity, crystal structure, and atomic-size factor are favorable for a complete series of solid solutions for this system. Both elements are in the same family of group V of the periodic table and thus are quite similar chemically. The crystal structures of columbium and vanadium are compatible for extensive solid solubility, since both have body-centered-cubic structures. The atomic diameters of columbium and vanadium are 2.85 and 2.62Å, respectively. This difference of slightly more than 8 pct is well within the 15 pct maximum difference allowed for extensive solid solubility. Experimental Procedures Source of Materials: Columbium powder and sheet trimmings were obtained from the Fansteel Metallurgical Corp. According to the manufacturer the metal contains less than 1 pct impurity. An analysis of the metal showed approximately 1800 ppm C in the powder while the sheet trimmings contained less than 500 ppm C. Spectrographic analysis showed minor amounts of Ca, Cr, Fe, Mn, Si, Ti, V, and Zr in both forms of the columbium. No commercial source of vanadium having the ductility and purity desired was available to the authors at the beginning of this investigation. As a result, all of the vanadium used in this study was prepared by the bomb reduction of vanadium pen-toxide with calcium employing the method reported by Long.' Yields of massive vanadium normally were about 80 pct. Chemical analysis of the vanadium prepared in this manner showed the presence of 200 to 500 ppm N and 800 to 1000 ppm C. Minor amounts of Ca, Fe, Mn, Si, Zr, Cr, and Cb were detected by spectrographic analysis. This vanadium metal was ductile and was cold rolled into 5 mil sheet. Annealing was not necessary during this rolling and the metal retained its cold-rolling characteristic after are-melting. Preparation of Alloys: The Cb-V alloys were prepared by melting pieces of vanadium sheet togethel-with columbium in the form of sheet or pellets of powder. The melting was carried out under argon in conventional arc-melting equipment employing a tungsten electrode and a water-cooled copper crucible. Each alloy was remelted three or four times, inverting the alloy after each melting in order to assure complete mixing. Alloys normally were obtained as round flat disks, weighed approximately 70 grams, and had roughly the shape of a disk 1 1/2 in. in diameter and 1/4 in. thick. Half of each alloy

Jan 1, 1955

By E. J. Ripling

A NY mechanism proposed to explain hydrogen embrittlement in titanium and its alloys must, of course, be consistent with the experimental data that characterize this embrittlement. Unfortunately, however, the mechanical behavior of hydrogen-bearing titanium, at least a-ß titanium, has not been unequivocally defined. Lenning, Craighead, and Jaffee have clearly shown that hydrogen cmbrittles a-titanium by elevating its transition temperature, probably as the result of the formation of titanium hydrides. Therefore, in these alloys, hydrogen acts like at least one other interstitial contaminant, namely, nitrogen.&apos; On the other hand, ß-titanium has been shown by these same investigators to tolerate very large amounts of hydrogen without suffering severe mechanical damage.2 Mixing these two phases, however, to form the most important class of commercial alloys, the a-ß alloys, again results in severe hydrogen embrittlement, although the mechanism by which the embrittlement is produced is not of the same type as that which causes brittleness in a alloys. Ductility damage due to hydrogen increases as the strain rate is reduced in a-ß alloys, while embrittlement in a alloys increases as the strain rate is increased, since the latter is a transition temperature behavior. Steel, like the a-ß alloys, becomes more hydrogen sensitive at slow strain rates, suggesting that the mechanism producing the embrittlement in these two metals is similar. Brown and Baldwin&apos; described the hydrogen-produced ductility depression in steel as a function of testing temperature and strain rate by defining the slope of the two surfaces that produced the depression in a three dimensional chart, Fig. 1. One of these surfaces was given by the equations (de/de)4 > 0 (de/dT) < 0 [1] while the other was defined by the pair of equations Surfaces of the type given by Eq. 1 are suggested in two ways. One is an embrittlement mechanism wherein the diffusion rate of hydrogen is competitive with the rate at which the material is being deformed, as suggested by the planar pressure theory of Zapffe and his co-workers.&apos; The other is the diffusion controlled extension of Orowan&apos;s theory on delayed fracture in glass by Petch and Stables.&apos;&apos; Surfaces of the type given by Eq. 2 are also compatible with a mechanism of pressure build up in voids, according to de Kazinczy,&apos; since the solubility of hydrogen in the metal increases with testing temperature so that as the temperature is raised, the pressure in the voids is reduced. Kotfila and Erbin recently presented some data on the dependence of ductility on testing temperature and strain rate for the a-ß 3 pet Mn complex alloy at four different hydrogen levels.H Although data were presented for only three testing temperatures at three different strain rates, their results indicated that surfaces of the types defined in Eqs. 1 and 2 are produced in the alloy when the hydrogen level is sufficiently high—200 and 300 ppm—Fig. 2. Jaffee and his co-workers presented data on a number of different a-ß alloys which indicated the existence of surfaces of the type described by Eq. 2, but the ductility recovery at low temperatures as given by Eq. 1 was not found." In an attempt to aid in crystallizing this description of the ductility dependence of hydrogen-bearing a-ß alloys, tests were conducted by the author on a-ß titanium 140A with three different hydrogen contents. The tensile properties of two as-received rods and one vacuum annealed rod* were obtained over a range of testing temperatures and strain rates as shown in Figs. 3 and 4. Hydrogen analyses were made by the Battelle Memorial Institute on four pieces of the rod whose properties are shown as solid circles in Fig. 4. The hydrogen content of these pieces, taken at widely spaced intervals within the rod, were 280, 270, 289, and 270 ppm, indicating that the hydrogen content within a single as-received rod was quite uniform. One of the broken test pieces whose properties are shown in Fig. 3 as solid circles was also analyzed, and found to have a hydrogen content of 310 ppm. The analyses obtained on two of the vacuum annealed specimens were 92 and 170 ppm.

Jan 1, 1957

By E. G. Woods, A. G. Comer

A mathematical model is presented and solved for determinution of the saturation distribution and pressure in a radial gas-storage reservoir. The model consists basically of two parts: (I) the growing gas-bubble core, and (2) the surrounding aquifer. Since the total pressure at the injection well is a function of the two-phase flow in the gas bubble and the unsteady single-phase flow in the aquifer, the resistance to flow in both zones was taken into consideration. The assumptions involved for both the radial equivalent of the Buckley-Leverett two-phase-flow equation and the injection-pressure equation are as follows: (I) the geometry is radial, (2) the gas bubble is free to expand or contract, (3) compression or expansion of the gas within the bubble may occur at the beginning of a time step, (4) the fluids are immiscible, (5) water is incompressible within the gas-storage region whereas it is compressible outside of this region, (6) a stable gas-water interface exists and (7) gas injection occurs at a constant rate or a series of constant rates. The mathematical model was solved numerically using an IBM 650 computer. A comparison is presented between the predicted results of the model, the results assuming steady-state flow and the actual initial injection-pressure history of an operating reservoir. Using the initial field pressure for a basis, the average deviation between the predicted pressure and the actual field pressure was less than 4.3 per cent. INTRODUCTION Underground storage of natural gas in abandoned oil fields, in abandoned coal mines, in caverns and in aquifers has had varying degrees of success. The lack of these first three facilities in the vicinity of most major gas marketing areas is leading more and more to storage in virgin aquifers. For the most part, this storage in aquifers has led to the study upon which this paper is based. Several excellent papers have been published discussing the problems involved in the underground storage of natural gas; however, to these writers&apos; knowledge, none of them considered the two-phase flow of fluids in the gas bubble and the fact that the gas bubble will grow or shrink in size depending upon the injection and withdrawal history of the reservoir. Not only will both of these factors influence the required injection pressure, but also they will have a definite bearing upon the amount of water production occurring upon withdrawal from the storage area. To solve this problem on an IBM 650 computer, several simplifying assumptions were made. Only the drainage portion of the relative-permeability curve was considered, thereby neglecting any hysteresis effects introduced during an imbibition cycle. The quantity of gas withdrawn during the period of the field study was only a small percentage of the total gas in place; therefore, this assumption is justified. For the problem studied in this case, the combined effects of capillarity and gravity were assumed to be negligible. This assumption becomes less valid as the formation thickness or the average pore diameter increases markedly. The justification of these assumptions should be studied in each case. THEORY This mathematical model for the prediction of saturation distribution and pressure for a radial gas-storage aquifer is based upon the equations for radial two-phase fluid flow&apos;.&apos; and for radial, unsteady-state, single-phase fluid flow.&apos; The two-phase flow is considered to take place in a "core" whose radius is equal to the maximum radius that the gas zone will attain, as indicated in Fig. 1. The injection-withdrawal history is approximated by a series of constant flow rates. Within the "core", gas is assumed to behave as a semi-compressible fluid; that is, the gas is assumed to have a constant density based on the average gas-zone pressure for the flow period. Between each constant-rate time increment, the gas density is allowed to change. This involves an iterative solution of the pressure equation. The liquid phase in this zone is considered to be incompressible, whereas it behaves as a compressible liquid outside of this region. The equations of incompressible flow in the "core" may be written for the gas phase

By D. P. Kedzie, R. A. Dodd

The fatigue strength, fatigue hardening, and effect of fatigue deformation on subsequent age hardening of supersaturated Al(Cu) solid solutions have been determined as functions of alloy composition and temperature. The fatigue strength/tensile strength ratios, determined at 150°, 25°, and -195°C, decreased with increase in alloy content for all temperatures, but the F.S./U.T.S. ratios at -195°C decreased much more rapidly than did the ratios for 150" and 25°C. This suggested that strain aging and/or age hardening occurred during tests at higher temperatures. Additionally, the F.S./U.T.S. ratios at 150°C exceeded those at 25°C for all compositions, indicating greater strengthening during fatigue at 150°C. The effect of fatigue or tensile deformation at 150°, 25°, and -195°C on subsequent age hardening showed that the deformatzon increased the rate of precipitation and indicated that mechanically produced vacancies were probably formed during deformation. Fatigue hardening was studied at 150°, 25°, and -195°C, and the effect of room-temperature rests after 10 and 100 cycles was examined. The results confirmed that strain aging occurred at the higher temperatures . DURING the last 15 years various mechanisms of fatigue crack nucleation and growth based on dislocation and vacancy interactions, operating singly or collectively, have been proposed. The probable consensus of present opinion is that the fatigue process in pure metals essentially involves dislocation interactions, and that vacancies formed by such interactions play a minor or inconsequential role. However, there is some evidence that age-hardened alloys tend to overage during fatigue, probably by local vacancy-enhanced diffusion, and strain aging also might be important in selected cases. Furthermore, it has been established that the behavior of quenched-in vacancies in Al(Cu) and other solid solutions is composition-sensitive. Therefore, it seemed worthwhile to investigate various aspects of the fatigue behavior of supersaturated Al(Cu) alloys and to examine the results in terms of vacancy-enhanced effects. EXPERIMENTAL The alloys used in this investigation were prepared from 99.994 wt pet A1 and OFHC copper, the latter containing 0.04 pet 0 as the principal impurity. Six alloys were made, containing, by actual analysis, 0.58, 0.96, 1.96, 2.85, 4.45, and 5.51 wt pet Cu. They were induction-melted in air in graphite crucibles, cast as 7-in. by 7/8-in.-diam rods in graphite molds, and hot-rolled to 5/8 in. diam. The experimental work was a three-part program involving the determination of a) 10&apos; cycle fatigue stresses as a function of alloy composition and temperature; b) the effect of fatigue deformation on subsequent aging of the supersaturated alloys; and c) fatigue hardening as a function of alloy composition and temperature. For determining the 105 cycle fatigue stresses, a portion of the 5/8-in. stock was machined into Krouse rotating cantilever beam fatigue specimens, 2 in. in length by 1/4 in. minimum diameter. These were tested at +150°, 25°, and -195°C (liquid nitrogen) on a Krouse high-speed machine, with special weights to provide lower -than-normal loading ranges, this being necessitated by the small load requirements of those alloys of lower copper content. For the high-temperature tests a small resistance heater was designed to clear the collets and fit in between the chucks, while for the low-temperature tests a hollow nylon cylinder was used, having closed ends drilled to pass a fatigue specimen, and positioned similarly to the heater. A plexiglass container completely enclosed the fatigue machine; rubber gloves fixed to ports in the walls enabled the machine to be operated from outside the box. A tray of magnesium perchlorate dried the air in the container and prevented both atmospheric corrosion fatigue at room and elevated temperatures and troublesome ice build-up at low temperatures. A total of ten to fifteen specimens was used for each combination of alloy composition and temperature. The result of each test was plotted on a standard S-N diagram, and the next stress was selected on the basis of the trends indicated by previous specimens. In this manner a small portion of the S-N diagram was constructed, and the 105 cycle fatigue stress obtained. Small Krouse fatigue specimens were also used to study the effect of cyclic prestrain on subsequent

Jan 1, 1964

By R. D. Russell, R. M. Farquhar

SOTOPIC tracers have become an important aid in following the progress of chemical processes in the laboratory. It has recently been found possible to utilize a system of naturally existing iso-topic tracers to obtain information about the geological history of lead ores. Common lead, such as is found in lead deposits, is a mixture of four stable isotopes having atomic weights 204, 206, 207, and 208. Of these, the last three are identical with the lead isotopes produced as stable end products of the radioactive decays of uranium and thorium: the first, lead-204, is not known to be produced on the surface of the earth by any process. Since uranium and thorium Occur in the surface regions of the earth in amounts comparable with lead, and since the half-lives of uranium and thorium isotopes are of the same order as the age of the earth, they produce the radiogenic lead isotopes in amounts comparable to the amount of nonradiogenic lead present. Every significant exposure of a sample of lead to uranium and thorium will therefore lead to the permanent alteration of the lead isotope ratios in that sample. It is this unique property of lead that serves as a means of tracing the history of a lead sample in terms of its contacts with the radioactive elements. If lead from a lead mineral has been analyzed with a mass spectrometer, the measured isotope ratios are determined entirely by the isotope ratios of primeval lead, which are identical for all minerals, and by the particular history of the sample. It follows that for samples from any particular geological area, observed differences in the isotopic composition are enough to distinguish different geological histories. An illustration of the qualitative application of this statement is given in Table I by analyses of some galenas from the western Cordillera. Samples from deposits in Pre-Cambrian sediments have very different lead isotope ratios from those of the ores in the Paleozoic sediments. Although the two types are associated closely geographically, it is apparent that they have had quite different histories and have probably been emplaced at quite different times, as the ideas outlined in the following section suggest. Even when applied quantitatively, a lead isotope analysis can never indicate a unique history of any lead sample. However, it greatly restricts the choices available and combined with other geological and geochemical data can lead to a much better understanding of the genesis of lead ores. General Character of Lead Isotope Variations: Early isotopic analyses of common leads by Nierl showed that geologically younger leads were generally richer in isotopes of masses 206, 207, and 208, with respect to that of mass 204. This regularity of measured lead isotope ratios can be easily observed by plotting each of the ratios Pb207/Pbm and Pb205/ Pb204 against the ratio Pb206/Pb204. In both graphs the points lie scattered closely about a well defined mean curve. It was immediately supposed that this regularity resulted from the growth of all leads from a common primeval lead present at some time, To, early in the earth&apos;s history. Lead in the outer part of the earth would become continually enriched in the radiogenic isotopes as a result of the uranium and thorium intimately associated with it. The subsequent extraction of some of this lead and formation of a lead mineral free of the radioactive parents provide samples of lead existing in the earth at the time of mineralization, T. Younger leads in general will be richer in the radiogenic isotopes because they have been associated with uranium and thorium for a longer time. Then lead ratios would be given by the formulae: (Pb206/Pb204)YTPb^/Pb204) +f "(u201) kdt T (Pb207Pb204)r = (Pb208/Pb204)T0 + J (U295/Pb204) k&apos; dt T (Pb208/Pb204) t = (Pb208Pb204 + J (Th204) A." dt [1] where A, A&apos;, and A" are the decay constants of U238, U285 and Th232. Summaries of measurements of radioactivity of surface rocks, such as given by Faul,2 are of limited

Jan 1, 1958

By J. B. Clemmer, P. E. Churchward, C. Rampacek

A cyclic method for processing manganese ores using sodium sulphate as the basic reagent is described. Sodium sulphate is electrolyzed in a diaphragm cell to give an anolyte-containing agentisdescribed.Sodiumsulphuric acid and a catholyte-containing iselectrolyzedincaustic soda, a part of which is carbonated in a cyclic manner to form soda ash. The reduced ore is leached with the anolyte to dissolve the manganese and the impurities are precipitated by addition of catholyte. Manganese in the pregnant todissolvethemanganeseandtheimpuritiessolution is precipitated as synthetic rhodochrosite with a carbonated catholyte, and sodium sulphate is regenerated. Calcining the manganese carbonate gives a high grade product and the carbon dioxide is returned to the carbonation step. The results of laboratory tests using cells equipped with synthetic-fiber diaphragms and permselective membranes, which permit transfer of anions or cations during electrolysis, are described. A METHOD which has been considered for pro-cessing manganese ores under emergency conditions to recover a product suitable for smelting to ferromanganese involves reductive roasting of the ore for leaching with dilute sulphuric acid followed by precipitation of the manganese from the purified pregnant solution with sodium carbonate. Although the process is operable and involves no novel or untried steps, it is noncyclic in regard to reagents. Research seemed warranted to determine if a cyclic process could be developed by electrolyzing the sodium sulphate formed during precipitation of the manganese with soda ash to regenerate sulphuric acid and sodium hydroxide. Reduced ore would be leached with acid anolyte to dissolve the manganese. The iron, alumina, and other dissolved impurities would be precipitated with the caustic catholyte, and the manganese in the purified pregnant solution would be precipitated as artificial rhodochrosite with carbonated catholyte to regenerate simultaneously sodium sulphate for electrolysis. Waste boiler gas and CO, from calcining the manganese precipitate would be recycled for carbonation of the caustic catholyte. There is a surprising lack of data in the technical literature on diaphragm electrolysis of sodium sulphate. One of the more comprehensive papers, by Atwell and Fuwa, describes the problem involved in electrolyzing byproduct sodium sulphate from the viscose process with particular regard to the design and operation of asbestos diaphragm cells. Frisk&apos; also investigated electrolysis of sodium sulphate from viscose-spinning baths to recover caustic soda and sulphuric acid. The objective in these investigations was to produce, by electrolysis and evaporation, a concentrated caustic-soda solution and a sulphuric-acid solution low in sodium sulphate. Because of these requirements, the process was not too attractive despite a favorable yield of caustic and acid with a reasonable power consumption. Because strong acid and base are not required for processing manganese ores and nonelectrolyzed sodium sulphate is not objectionable in either reagent, the present problem was somewhat simpler than that of these investigators. The purpose of this report is to summarize briefly the results of laboratory electrolysis of commercial sodium sulphate in diaphragm cells of different types and the application of the acid and caustic for processing manganese ores in a cyclic manner. Electrolysis in Permeable Diaphragm Cells The principal problem in electrolyzing sodium sulphate in a permeable diaphragm cell is to prevent mixing of the acid and base formed in the anode and cathode compartments. It is not enough to separate the anolyte and catholyte by a single diaphragm. A three-compartment cell in which the sodium-sulphate solution is fed to the center compartment and flows through the permeable diaphragms to the anode and cathode compartments is more efficient electrically. Diaphragms should be porous enough to allow flow of solution to outer compartments but tight enough to prevent gross mixing of the anolyte and catholyte. In addition to being resistant to acid and alkali, the diaphragms should offer minimum resistance to the flow of current. It is also important that the electrolyte contain a high concentration of sodium sulphate to furnish an abundance of sodium and sulphate ions and thereby decrease the current carried by hydrogen and hydroxyl ions. A practical limitation of sodium-sulphate concentration is imposed by the decreased solubility of sodium sulphate with decreasing temperature. At 20°C, a concentration of 200 g per liter sodium sulphate can be maintained. Little is gained in conductivity by using a concentration greater than 200 g per liter. The small-scale continuous electrolysis tests were made in a 12x12 in. open-top bakelite cell equipped

Jan 1, 1956

By B. C. Allen

Surface free energy and surface self-diffusion of solid molybdenum were studied in the temperature range 1600" to 2400°C using pressure-sintered bi-crystals. Comparison of groove angles formed in various surfaces perpendicular to the grain boundary indicate a maximum of 1 pct variation in surface free energy with crystallographic mientation. This anisotropy tends to decrease with increasing temperature. The surface diffusion of the bicrystals is equivalent to that of sheet with a mild (100) Preferred orientation. Anomalously low values found for bi-crystals with surface orientations of (OOl), (012), and (011) are rationalized in terms of anisotropy in surface free energy. THE effect of crystallographic orientation on surface free energy1,&apos; and surface self-diffusion3,4 has been primarily studied in fcc metals. The object of this work was to study the effect of orientation on surface diffusion and surface free energy of bcc molybdenum using pressure-sintered bicrystals. EXPERIMENTAL WORK Materials and Crystal Preparation. Arc-melted molybdenum rod was obtained commercially and electron beam zone refined at 50 cm per hr at 10- 5 torr to form single crystals about 8 cm long and 0.65 cm diam. Three crystals were prepared with axial orientations about 1 deg from [001.], [011], and. [111]. To reduce the carbon content, the crystals were annealed 2 hr in 1.4 atm flowing wet hydrogen at 2050°C. Then the oxygen content was reduced by annealing for 2 hr in -30°C dewpoint hydrogen at 2020°C. The resulting impurity analysis is given in Table I. Bicrystal Preparation. The single crystal rods were cut into transverse slices with a thin silicon carbide abrasive wheel to produce specimens about 0.6 cm long. They were mounted in epoxy and surrounded by stainless steel washers. Cutting in half was done longitudinally at various angles to known crystallographic planes containing the cylinder axis according to Fig. 1. To reduce surface deformation resulting from the cutoff wheel and thus reduce parasitic grain boundary formation on subsequent annealing, about 0.003 cm was manually ground off each cut surface with 600 grit paper. Care was taken to keep the surface flat. After removal from the mounts, one half was generally ro-tated 180 deg with respect to the other to give a po- tential symmetrical tilt grain boundary between the two halves. In the other cases when low misorienta-tion angles were desired, the crystals were not rotated. On the basis of symmetry, sufficient bicrystals were prepared to cover the entire range of misorientations for symmetrical tilt boundaries. The misorientations, +, ranged from 0 to 45, 0 to 90, and 0 to 60 deg for [001], [011], and [111] bicrystals, respectively. One [Ill] twist bicrystal was prepared from 2 single crystal discs rotated 17 deg relative to each other. Each specimen consisted of two pieces which were placed in a cylindrical tantalum can. Sharp edges were rounded and the fit was made as snug as possible to reduce subsequent deformation during bonding. The assembly and crystals generally were vicuum outgassed at 900" or 1700°C and then electron beam welded in the can at l x 10-4 torr. After being leak checked, the samples were placed in an autoclave and hydrostatically gas-pressure bonded5 in four batches under helium at 10,000 to 18,000 psi at 1650°C for 3 hr. Satisfactory bonds were obtained in many cases, and most of the crystals bonded after two exposures. The results did not appear to be affected by the various pressures used, preannealing conditions, crystal orientation, or time-pressure-temperature route taken to the final bonding condition. After bonding, the tantalum cans were selectively removed in cold concentrated HF. Measurements indicated overall deformation was under 1 pct. The bicrystals were metallographically ground and polished flat and perpendicular to the axis. Examination showed the boundaries were straight and almost free of parasitic grains caused by extraneous local deformation. Annealing. In preparation for thermal grooving, the bicrystals were cleaned and annealed by outgassing at 10-5 torr at 1900°C and heating at 2300°C under 1 atm 99.996 Ar for 0.5 hr. The crystals were held in a closed 4-deck box made of molybdenum sheet, and were heated in a Ta-1OW resistance furnace. The ar-

Jan 1, 1970

By W. J. Moore, E. M. Swiggard, H. Shenker

Beryllium is the most soluble of the double-acceptor impurities in germanium; the solubility is at least 1 x 10 19 atoms cm3. Photoconductive, optical, and electrical measurements were made on a set of Ge:Be samples grown without other intentionally added imp~ivities. The excited-states spectrum of beryllium in germanium was determined both in zero field and in magnetic fields up to 30 kc. It was found that the absorption 0-f lines in the vicinity of the G line increased markedly with increasing magnetic field. The ioniza-tion energy of neutral beryllium in germanium was determined to be 24.3 * 0.1 mez: using the excited-states spectrum. An infrared detector using a Ge:Be element was found to have a normalized detectivity D* (50 p, 1000 cps, 1 cps) of 9 X ldO cm-~ps"~-w-&apos;. The Ge:Be sambles exhibited low-field breakdown and inibzirity band conduction effects at low temperatures. WHILE the properties of most dopants in germanium have been thoroughly investigated, the properties of beryllium in germanium have had relatively little attention. Much of the work that has been done&apos; has been concerned with the properties of singly ionized beryllium in germanium. The purpose of the present work was to determine the photoconductive and electrical properties of beryllium-doped germanium (Ge:~e) with a small concentration of other electrically active impurities. Ge:Be is of special interest since beryllium is the most soluble of the known double-acceptor impurities in germanium. MATERIALS PREPARATION A set of &:Be crystals with beryllium concentrations between 5 x 1013 and 1x 10&apos;&apos; cmm3 and without other intentionally added impurities was grown in a hydrogen atmosphere by the horizontal zone-refining techniques described by Swiggard and Shenker.2 The segregation coefficient of beryllium in germanium was found to be about 0.07 and its maximum solubility was found to be greater than 1 x 10&apos;&apos; atoms cm-3 which is in agreement with previous results.&apos; Most of the crystals were found to have a small net concentration of shallow acceptors; for example, a crystal having a beryllium concentration of 1.1 X 1016 cm-3 was found to have a net concentration of shallow acceptors of less than 1.5 X 1013 cm-3. Ingots designed to have a beryllium concentration of 1 x 1017 cm-3 or less exhibited a loss of electrically active beryllium. That is, the beryllium concentration was lower than expected and the decrease of the beryllium concentration along the length of the ingot was much larger than expected. It is hypothesized that this effect is due to the oxidation of the beryllium by the residual oxygen in the system. This effect is similar to but not as severe as the effect observed with boron in germani~m.3, 4 OPTICAL ABSORPTION SPECTRUM The excited-states spectrum of neutral beryllium atoms in germanium was studied by measuring the infrared optical absorption spectrum. A typical spectrum is shown in Fig. 1; the labeling of the absorption lines follows that of Fisher and Fan.* The spacings between the observed lines and the D line are nearly equal to the corresponding spacings for the other ac -ceptors in germanium.5-7 By the use of the theoretical result67&apos; that the neutral impurity ionization energy is 2.53 mev larger than the energy of the D line, the ionization energy of neutral beryllium in germanium was determined to be 24.3 * 0.1 mev at 4.2"K. This method yields a more precise value of the impurity ionization energy than the usual Hall effect method.&apos; The absorption spectrum of Ge:Be was also measured as a function of magnetic field using a magnetic field parallel to the incident radiation and using unpolarized radiation. A typical spectrum taken at 30 kG is shown in Fig. 2. An unexpected observation is that the absorption of the lines in the vicinity of the G line (there is a strong possibility that some of these lines may be associated with lines other than the G line) increases with magnetic field. RESISTIVITY AND HALL DATA Measurements of Hall voltage vs temperature were made on several of the samples. The thermal activa-

Jan 1, 1968

By L. D. Mullins, W. R. Bartlett, R. H. Barham, W. L. Wahl

This paper describes an electrical model and its applicution to the analysis of four reservoirs in Saudi Arabia. The model has 2,501 mesh points and represents 35,000 sq miles of the Arab-D member. Details of modeling such as mesh size, control problems and standards of performance in matching reservoir history are discussed. The particular performance match achieved for the Arab-D rnember is presented. Details such as permeability barriers, aquifer depletion and iwerference between oil fields are given. The performance match realized in the Abqaiq pool is presented in detail. INTRODUCTION The resistor-capacitor network and associated control equipment described in this paper comprise an electrical analog of a reservoir system. Similar equipment has been used to study the transient response of reservoirs for many years. The unique feature of the model and application to be described is the extremely large size of the model and reservoir system, and the detail observed in simulating the reservoir with the model. The Arabian American Oil Co. first became interested in analog computers for simulation of oil reservoirs in 1949. Since that time, several models have been developed, each more elaborate and refined so that the reservoir system might be more closely simulated. The current model is the latest in a series designed, built and operated by the Field Research Laboratory of Socony Mobil Oil Co. in collaboration with Aramco. It has been and continues to be used to study the regional performance of the Arab-D member limestone reservoir. The Arab-D member is one of the Middle East&apos;s most prolific producing horizons.. THE MODEL The theory of simulating a reservoir system with an electrical system has been presented in the literature.1-5 Therefore, this paper will not discuss the theoretical aspect of the problem except to point out the correspondence between the fluid system and electrical system, as shown in Table 1. In general, the complete model is made up of input devices, output devices, central control and a resistance-capacitance (RC) network. At times, the RC network alone is referred to as the "model". However, it should be evident from the text which meaning is attached to the word "model". A discussion of the equipment follows. THE RESISTANCE-CAPACITANCE NETWORK The RC network consists of 2,501 capacitance decades interconnected through 4,900 resistance decades. The components are arranged to form a rectangular network of 2,501 mesh points in a 41- X 61-mesh array. Imposing the mesh grid system on the continuous reservoir system divides the reservoir into discrete areal segments. These discrete segments may be of various sizes. More precisely, the mesh size need not be uniform throughout the model. The RC network is fabricated in two sections which are connected at the top. An inside view of the "tunnel" formed by the two sections is shown in Fig. 1. The height and width of the tunnel are shown in the figure. Numerals appear along the bottom and along the back opening of the tunnel. These numbers denote the x and y co-ordinate positions of mesh points. Fig. 2 presents a rear view of one-half the model. The length dimensions of the model, as well as a rear view of the capacitor decades, are shown in this figure. The control dials used in adjusting the resistance and capacitance values on the model can be seen in the enlarged portion of the model shown in Fig. 3. The electrical capacity at any mesh point can range from 0 to 1.0 microfarads set to the nearest tenth of a microfarad. The electric resistance connecting any two mesh points can range from 0 to 9,990,000 ohms set to the nearest 1,000 ohms. External capacitors may be added to any or all mesh points if the need arises. The values of electrical resistance and capacitance are adjusted manually by manipulating the two types of decade units. INPUT EQUIPMENT A considerable quantity of equipment is used to control the input to the RC network. These input devices are

By C. E. Lundin, D. T. Klodt

The binary alloy systems, Y-Ti, Y-Zr, and Y-Hf, have been investigated throughout their entire composition regions. There is no compound formation in any of the systems, and each system is characterized by a single eutectic reaction. The eutectic compositions and temperatures are as follows: A eutectoid reaction pct Y and 870°C occurs in the Y-Ti system, whereas a peritectoid reaction,: pct Y and 880°C occurs in the Y-Zr system. Peri-tectic-type reactions at temperatures above the eutectic levels are postulated for the yttrium and hafnium transfovmations. The development of the technology of yttrium has been given considerable attention during the past few years, and studies of binary phase equilibria have, of course, taken a prominent position in this development. In many respects yttrium, in the third group of metals of the periodic table, is similar to the adjacent group of metals, titanium, zirconium, and hafnium, and the knowledge of the phase relationships of yttrium with these metals is basic to their technology. MATERIALS AND EXPERIMENTAL PROCEDURES Materials. The metals for this investigation were supplied by the General Electric Co., Aircraft Nuclear Propulsion Department. The yttrium was in the form of an arc-melted ingot, and the other metals were in the form of high-purity, iodide-Process crystal bar. Table I lists the purities of these materials. Alloy Preparation. Melting was done by conventional techniques in a nonconsumable electrode arc furnace in an atmosphere of purified argon. Melting conditions for each binary system were the same. Each alloy button was inverted and remelted several times to assure homogeneity. Accurate weights of the charges and resultant alloy buttons were obtained to indicate deviations from intended compositions. No chemical analyses were obtained since melting weight losses were consistently in the range of 0.1 to 0.2 pct of the total weight. 10- or 20-g buttons for each 5.0 wt pct composition increment were melted to survey the three individual alloy systems. Additional alloys differing in composition by 1.0 or 0.1 wt pct increments were also melted to study selected regions of the systems. Metallograpllic Techniques. Standard metallo-graphic techniques were followed for mounting and rough grinding. Preliminary polishing was accomplished using 6-u diamond paste as an abrasive on a Metcloth Lap. Final polishing was done on a Microcloth-covered wheel using 1-u diamond abrasive paste. Purified kerosene was used as a lubricant for both polishing stages. • sothermal- Annealing. Alloys were sectioned for as-cast structlure examinations and then homogenized in preparation for isothermal-annealing treatments. Homo{:enization was accomplished by cold pressing the alloy buttons followed by 72-hr anneals at 1100c. The alloys were encapsulated in Vycor or quartz for the homogenization treatments or for isothermal anneals. Resistance-wound or resistance-element tube furnaces were used for the annealing treatments. The homogenized alloy buttons were cold rolled until cracking occurred or until a -in. specimen thickness was obtained. Small -in. square) specimens for the isothermal anneals were then sawed from the alloys. Each specimen was wrapped in tantalum foil before being sealed in the capsule. Temperatures during the anneals were controlled The time at temperature necessary to equilibrate the structures during the anneals was determined for each alloy system by holding triplicate specimens of alloys at a constant temperature for three different periotls. The specimens were quenched and examined microscopically to determine the number and amounts of phases present in the micro-structure as a function of time. Melting Studies. Eutectic temperatures of the three alloy systems were established from the results of incipient-melting studies conducted on as-cast alloys. Specimens to be melted were suspended on a tungsten wire inside a graphite cylinder placed in a glass vacuum chamber. An optical pyrometer was used to follow the temperature of the specimen as it was inductively heated in a high vacuum. The temperatures were corrected for emissivity losses by standardizing the pyrometer with known-melting-point metals. Accuracy of the temperature measurements is estimated to be + 10°C. The melting point of the yttrium was determined to be 1550°C by this technique. The invariant-temperature levels were also checked by an anneal-quench technique. This technique consists of annealing a series of

Jan 1, 1962

By S. A. Levy, J. D. Wood, J. F. Libsch

A study of the preparation and characteristics of a sevies of Fe-C-N alloys has been conducted. X-ray, microhardness, and metallographic data from a series of single-phase alloys produced by controlled nitriding of carburized foils are reported. The existence of a, y&apos;, and E carbonitride has been confirmed and it has been found that y&apos; is the hardest of these three phases. These results have been used to interpret the structures observed in nitrided steels. ThIS paper represents the second phase of a three-part study concerning combined nitriding and induction hardening. The first part&apos; reports the engineering properties (depth hardness, fatigue, corrosion, and tempering resistance) obtained when seven different steels were subjected to the combined surface hardening treatment. The present paper concerns the nature of a series of single-phase alloys prepared by controlled nitriding of carburized foils. The objective was to simulate the various layers present in the as-nitrided steels. A third paper will report data concerning induction hardening of the foils and will attempt to set forth a mechanism for the high hardness obtained with the combined treatment. Studies of nitriding generally fall into three categories. The first were performed with bars using conventional nitriding procedures, thus yielding a nitrogen gradient and a series of phase layers near the surface.2-6 Such studies were complicated by the presence of the nitrogen gradient, for it was difficult to correlate microhardness and X-ray data with chemical composition. There were always complications due to penetration of the X-ray beam which yielded data influenced by underlying layers. Similar difficulty was experienced in taking hardness readings on a tapered specimen or in isolating various layers by machining. A second group of investigations was primarily concerned with establishment of regions in the Fe-N7,9,10 or the Fe-C-N phase diagram.8 These were performed with powder specimens which were nitrided in am-monia-hydrogen mixtures to produce a series of single-phase specimens. Unfortunately, the powder specimens were not amenable to either microstruc-ture or microhardness study. The third was essentially a compromise of the first two. Bose and Hawkes," using 0.005-in.-thick iron foils, were able to obtain microhardness, microstruc-ture, composition, and X-ray data on Fe-N alloys of eutectoid composition by equilibration in ammonia-hydrogen mixtures at 1290°F. To obtain a better understanding of the nitriding process and the characteristics of the phases in ni-trided plain carbon steels, the present investigation was undertaken. The investigation was patterned after that of Bose and Hawkes with the exception that nitriding was performed at 930°F (below the eutectoid temperature), as in the case of conventional nitriding. In addition to gaining some understanding of the source of high hardness of nitrided and induction-hardened steels, found in a previous study,&apos; this investigation was concerned with the possibility of controlling the phases present in a nitrided case through the use of ammonia-hydrogen mixtures. Knowledge of the properties of each phase and the conditions necessary for its production could make it possible to avoid undesirable surface layers in nitriding. EXPERIMENTAL PROCEDURE Specimen Preparation. A specimen thickness of 0.005 in. was selected to insure that a uniform nitrogen concentration could be obtained in a reasonable length of time by equilibration in an ammonia-hydro-gen atmosphere. The starting material was a vacuum-cast ingot of pure iron, the major impurities of which were: C, 0.003 pct; 0, 0.014 pct; S, 0.004 pct; N, 0.003 pct.* &apos;Compositions are quoted as weight percent except as noted. The ingot was heated to 1500°F in a nitrogen atmosphere and forged to a slab 0.25 in. thick. The material was then reduced to 0.005 in. by alternately cold rolling and annealing in a nitrogen atmosphere at 1600°F. This material was divided into five groups of specimens which, after carburizing, yielded the following carbon contents: 0.003, 0.28, 0.36, 0.46, and 0.74 pct C. The carburizing was performed by the Leeds and Northrup Research Laboratories, using a controlled "carburizing potential" which yielded a uniform carbon concentration throughout the cross section.

Jan 1, 1970

By E. R. Helderman, C. Y. Ang, C. C. Nealey

Beryllium-base alloys have been successfully p7.-epared by the liquid-phase sintering technique. Depending orz the composition and amount of the intended liquid please, microstructures either single -phase or duplex in feature with randomly oriented grains have been obtained. Qualitative and semi-qualitative determination of distribution of alloying elements md microsegregations in some experi)uzental alloys haue been made by electron-micro-probe analysis in conjunction with microliardness testing and standard metallography. Tensile tests reoealed that some conzpositions possess attractive elastic properties with Young&apos;s mot1uli greater than 40 X 10&apos; psi. In powder metallurgy, liquid-phase sintering is a process or phenomenon that has proven to be of practical value. For example, the heavy metals such as tungsten-copper-nickel, high-strength heavy gyro alloys,&apos; heavy-duty electrical contacts,&apos; and tungsten carbide tool materials are all products of liquid-phase sintering. Mechanisms involved in liquid-phase sintering, however, are not completely understood. Questions regarding the exact roles played by rearrangement of particles, liquid/vapor surface energy, solution and precipitation, and so forth, have not been completely answered. Evidence has been cited3 that at least volume shrinkage in the densification process is diffusion-controlled. It is possible that the predominant mechanisms for the complete densification and grain growth in a liquid-phase sintered system depend primarily on the alloy systems involved, in addition to processing conditions. Despite the lack of sound understanding of the mechanisms of the process, liquid-phase sintering has been and is being used to advantage either to synthesize microstructures for their special properties or to prepare alloys which are difficult to form by fusion process. Liquid-phase sintering of beryllium alloys was first attempted by Jones and Williams.4 They first tried infiltrating beryllium with magnesium, and then succeeded in preparing the alloy by sintering Be-Mg powder compact in molten magnesium bath. This investigation resulted in the identification of some Mg-Be intermediate phases. Crossley et a1.5 also used liquid-phase sintering technique in an attempt to produce ductile beryllium alloys for structural applications. The major liquid-phase components investigated by Crossley et al. were aluminum and silver with minor additives of germanium, calcium, lanthanum, cerium, and yttrium. The lack of wetting and the bleeding out of liquid phase were the difficulties encountered during experimentation. According to Hodge,6 the two compositions investigated by Crossley, which showed some promise based on compression tests, involved large amounts (over 35 wt pct) of silver, thus making them too heavy to be of practical interest to the aerospace industry and military users. The present investigation was prompted by the search for light-weight (density less than aluminum) beryllium alloys exhibiting small anisotropy in physical properties for precision inertial navigation instrument applications. In addition to isotropy, good structural properties such as high elastic modulus or desirable combination of mechanical and physical properties are also objectives of this investigation. The technique of liquid-phase sintering was chosen for its versatility in producing either duplex or homogeneous microstructures. This report is concerned with the use of copper and aluminum, with or without silicon addition, as the intended liquid phase, and the resultant micro-structures and some physical properties of the beryllium alloys. Qualitative and semiquantitative electron-microprobe analyses of some of the alloys are presented to illustrate the usefulness of this microanalytical technique for the identification of microconstituents and their distribution. EXPERIMENTAL PROCEDURES The beryllium powder used was -200 mesh Brush Beryllium Co. QMV NP-50 grade. Typical chemical analysis of the powder is shown in Table I. Commercial high-purity copper, aluminum, and silicon powders all screened to —100 mesh were used as additions. Mixed powder compositions were ball-milled in ceramic jars for 1/2 hr to ensure thorough blending. Milled powder was loaded either in 3/4-in.-diam button die or in Metal Powder Association flat tensile specimen die and compacted under top and bottom pressure. No lubricant or organic binder was used. Sintering was carried out in a quartz tube under a vacuum of 50 to 100 µ pressure. Surface hardness and some microhardness readings were taken on sintered specimens. Sintered density was

Jan 1, 1965

By D. F. Stein

In spite of the apparent usefulness of autoradiography in demonstrating segregation, it has had very limited success in demonstrating grain boundary segregation. Because of this limited success, a model system amenable to mathematical analysis was devised to determine which variables in the experiment are most important. As a result of these calculations, it was concluded that autoradiograPhy is a rather insensitive technique to measure grain boundary segregation. The range (energies) of the emitted particles (ß and a) must be low, and the concentration of the radioactive speczes at the grain boundary must be (in general) two or three orders of magnitude greater than the concentration within the grain. Because of these very restricted conditions , the limited success of the technzque is not surprising. In spite of the apparent usefulness of autoradiography in demonstrating segregation, it has had very limited success in demonstrating grain boundary segregation. Sulfur segregation in iron1 and possibly polonium segregation in Pb-5 pct B1 2 alloys are the only autoradi-ography experiments that have demonstrated segregation to grain boundaries in metals without the formation of a second phase. Segregation to a boundary without the formation of a second phase is often called Gibbs&apos; absorption and is discussed by McLean in Chapter 5 of Ref. 10. There are several394 experiments showing grain boundary diffusion of a radioisotope, but this type experiment is not representative of equilibrium between the concentration of an element at a grain boundary and that in bulk, so it will not be discussed in this paper. In an attempt to determine if temper embrittle ment of low-alloy steels was associated with segregation of antimony to grain boundaries, a program to use auto-radiography (using Sb-125) was initiated. Even though other measurements strongly suggest that segregation is occurring during embrittlement, no evidence of grain boundary segregation was observed in the auto-radiography experiments. An attempt was also made to detect segregation of carbon (using C-14) to grain boundaries in iron during slow cooling. There is again strong indirect evidence7 that segregation occurs during such a treatment, but the autoradiography experiment gave no evidence of such segregation. Because of the failure of these experiments and the general lack of success by our metallographic unit in measuring grain boundary segregation using autoradi-ography, a model system amenable to mathematical analysis was devised to determine which variables in the experiment are most important. MODEL SYSTEM The model system is illustrated in Figs. 1 and 2. It is a semi-infinite bicrystal with a single grain boundary perpendicular to the top face. The width of the grain boundary, W, is defined as the region to which segregation has occurred and it is assumed that this region is of constant composition. The assumption of an infinite dimension is for mathematical reasons but the results of such an analysis (as will be demonstrated later) are valid for even very small specimens. It also is assumed that the measurement is not limited by means of detecting the radiation (photographic emulsions, counters, and so forth) but that the means of detecting radiation is linearly sensitive to the radiation and has infinite resolving power. The distance y is the perpendicular distance from the center of the grain boundary to the point at which the background radiation is measured and x is the integration variable. CALCULATION FOR BETA PARTICLES Two values of the intensity of radiation will be calculated, the intensity at the center of the grain boundary and the intensity far from the grain boundary. The radiation at the center of the grain boundary can be calculated in the following way. Assume a grain boundary of width W having a uniform concentration equal to Cgb + Cb where CGb is the excess concentration of the element under consideration at the grain boundary and Cb is the bulk concentration. The intensity of radiation, IgB, at the center of the grain boundary, is a consequence of the amount of radioactive material, decay rate, absorption, and geometrical factors which can be represented mathematically by the following expression:

Jan 1, 1968

By R. Schuhmann

Available thermodynamic data applicable to copper smelting systems are collected and tabulated, and the important gaps are pointed out. A few examples are given of estimations which can be made from the available data. An experimental research program is proposed to supply the thermodynamic data that appear most essential to better quantitative understanding of the chemistry of copper smelting. The proposed program is designed also to shed specific light on the practical problems of slag losses and magnetite behavior. OPPER smelting, from flotation concentrates to V-/ blister copper, is conspicuous among the large scale chemical processes which are conducted with only relatively incomplete knowledge of the physical chemistry involved. For example, no common metallurgy text explains adequately why copper enters the matte phase while iron enters only to the extent that sulphur is left over after satisfying the copper. Explaining this important phenomenon as a manifestation of the greater affinity of sulphur for copper than for iron is unsound because the affinities of copper and iron for sulphur are about the same at smelting temperatures. As will be shown, other affinities are really decisive in the relatively clear-cut separation of copper in the matte. In contrast, thermodynamic studies have contributed much to our understanding of copper refining, zinc oxide reduction, magnesium production, steelmaking, etc. For each of these processes, the important reactions are clearly recognized, and fair to good quantitative values of the free-energy changes and equilibrium constants are available. Such data have proved to be of constantly increasing practical value in process development and improvement. Work in other fields has furnished much thermodynamic data applicable to copper smelting systems. In fact, several promising starts have been made in applying such data to specific smelting problems. Kelleyl made an appraisal of the possibilities of recovering elementary sulphur from low grade matte. His calculations and compilations of thermodynamic data have represented the starting point for much of the work in this field, including the present survey. Huang and Hayward2 nd Aksoy3 used thermodynamic methods in the study of copper losses in reverberatory slags. Peretti4 used thermodynamic data in explaining the chemistry of converting. The recent publications of Darken and Gurry"." and of Darken,&apos; dealing with the iron-oxygen and iron-silicon-oxygen systems, respectively, present equilibrium data that are applicable to copper smelting systems. Also additional data were reported recently on the affinity of sulphur for copper, manganese, and iron8 and on the sulphur pressures of iron-sulphur melts." The survey presented in this paper was made as the basis for planning an experimental program on the thermodynamics of copper smelting. Few researches on the chemistry of copper smelting have been reported in recent years, so that a reappraisal and coordination of old and new data are essential if further work is to take the directions of maximum value. The experimental program is now in progress, and the plan of attack is outlined at the end of this paper. Progressive Oxidation and Desulphurization of Copper-bearing Liquid Phases: The chemical activities of sulphur and oxygen are two of the most important thermodynamic yardsticks to be applied to copper smelting processes. Virtually the entire smelting and refining sequence involves a series of systems characterized by decreasing sulphur activity and increasing oxygen activity. In this section, therefore, an attempt is made to define and explain these activities in terms of equilibrium partial pressures of SO2, s2, and O,. Also, estimates of these quantities are presented, the estimates being based largely on calculations presented in a later section of this paper. The sequence of steps from raw flotation concentrate to fully oxidized copper ready for poling involves progressive and controlled oxidation. Iron is oxidized and enters the slag. Sulphur is oxidized and leaves in the gas. Table I summarizes several important features of this oxidation and desulphurization sequence, starting at the beginning of the matte blow in the converter. The top line gives in order the principal smelting and refining stages up to fully oxidized copper, plus the additional step, not used commercially, of oxidizing all the way to Cu,O. In the second line are shown the principal copper-bearing liquid phases which characterize the process. Through most of the sequence the copper

Jan 1, 1951

By D. A. Munro, G. P. Contractor

Fourteen 50-lb laboratory melts were investigated to determine the effect of uranium on the tenpering characteristics of loo-carbon (0.06 to 0.1 pct C) steels. It was found that uranium additions, particularly in the range 0.30 to 0.45 pct, enhanced the hardness and both ultimate and yield strength of the experivzental steels in the quenched and tempered condition. The structural and morphological chazges indicated that uranium retarded tempering of the tnartensite, thereby hindering the normal formation of polygonal ferrite formed in the late stages of tempering. The effect of this was to make possible the re-tension of the acicilar ferritic structure in the uranium-bearing&apos; steels. The iraniuin-bearing steels also showed IVidnzanstatten-type growth of ferrite plates and had large prior austenite grains containing assenzblies of fine ferrite grains, mainly acicular in geometry. The fine-grained ferrite structure and the presence of more numerous and apparently smaller precipitates in the uranium-bearing steels are thought to he principally responsible for the itnproved tensile strength and hardness of the experinzental uranium-bearing steels. At ternperirzg temperatures above 455% (850&apos;F) the ferrite in the higher-uraniun steels nzaintained acicularity and, hence, its strength and resistance to tempering. Uranium did not produce a secondary hardening peak. However, it retarded softening during the third stage of tempering because of its effect of inhibiting the grouth of cementite particles and of retaining the acicularity of ferrite plates. The resistance to coalescence accounted for the slow grocth of the ferrite grains in the uranium-modified steels and, hence, fov the persistence of the acicular ferrite structure. IT had been found previously1 that uranium additions up to about 0.45 pct had no significant effect on the tensile properties of low-carbon steel (0.06 to 0.10 pct C) in the as-rolled and normalized conditions, Fig. 1. On the other hand, it was observed that uranium in excess of about 0.30 pct had an embrittling effect as revealed by Charpy V-notch impact results. It was also noted that, as the uranium content increased, the morphology of pearlite changed from lamellar to feathery and the ferrite grains showed an etching effect resembling striated or dashed markings, suggestive of precipitation. The sharp drop in the impact properties shown in Fig. 2 warranted an assumption that the uranium content of about 0.30 to 0.45 pct may produce some secondary hardening reaction on tempering, analogous to that associated with a Cr-Mo-V steel, which shows very poor CVN toughness at the secondary hardness peak in the tempering curve.1&apos; With this background and the reported findings of Hasegawa and noda that low-carbon uranium-treated steel showed signs of secondary hardening, the present investigation was undertaken to determine the effect of uranium additions on the mechanical properties of 0.10 pct C steels. No attempts were made to investigate in detail the mechanisms of hardening, although some suggestions based on the experiments are made. MATERIALS AND PROCEDURES A series of 50-lb induction-furnace melts was made using AISI 1008 rimming steel billets as the melting stock. The melting, forging, and rolling techniques proven satisfactory in previous projects&apos;-3 were employed as a guide for this investigation. The steel was deoxidized with aluminum (2 lb per ton) prior to the addition of high-purity uranium. The analysis of each melt is given in Table I. Properties were evaluated as a function of heat treatment and are presented in terms of hardness and tensile strength vs tempering temperatures. The variation of hardness with the tempering temperature was studied on the quenched and tempered specimens, some of which measured 0.50 by 0.25 in. diam and the others 0.40-in. cubes. Before quenching, the specimens were vacuum-sealed in glass tubes and normalized at 900°C (1650°F) for 20 min. Following this treatment, the sealed specimens were hardened by austenitizing at 955°C (1750°F) for 20 min and water quenching, and then tempered for 1 hr in the range 150 to 730°C

Jan 1, 1967

By Pradeep K. Rohatgi, Clyde M. Adams

Structures obtained on freezing of several hypo-and hypereutectic Al-Cu alloys over a range of solidification rates have been examined. Dendrite spacing, L, increases linearly with solute concentration and with the square root of the inverse freezing rate. The relationship for hypoeutectic alloys is: where rate of change of fraction solid with time, is freezing rate, C is solute concentration, (pct Cu)=1. Mass transport in inter dendritic liquid during solidification is analyzed; the experimental observations suggest maximum concentration differences and constitutional supercooling in the inter dendritic liquid increase with an increase in the solute concentration. The dendrite morphology changes with freezing rate and alloy composition. The dendrites of the a phase are parallel, uniformly spaced plates with slow freezing and rods with rapid freezing. Nonor-thogonal side branching has been observed in phases with cubic and tetragonal structures. Side branches in a dendrites are orthogonal with slow freezing and at 60 deg with rapid freezing. Formation of second-phase envelopes around the Primary phase is also discussed. DENDRITIC structure is characteristic of many types of phase transformation. The most extensively studied so far has been solidification of liquid solutions. chalmersl and coworkers have interpreted the formation of dendrites in terms of the breakdown of a planar interface. Most of the work done concerns itself with the development of an instability at the interface. Little theoretical work has been done quantitatively to relate the parameters of dendritic structure to mass transport in the liquid phase. A few empirical relations based on the experimental2&apos;3 observations exist in the literature. Several workers2 including Brown and Adams1 have studied dendrite spacing in A1-Cu system as a function of solidification variables. In most cases, dendrite spacing has been found to increase linearly with the square root of some parameter proportional to the freezing time. The effect of solute concentration is not clear; some workers report the dendrite spacing increases with solute concentration4 whereas others report vice versa.&apos;&apos;&apos; ~ohatgi&apos; has observed an increase in the spacing between ice dendrites with an increase in solute concentration in water. Tiller has also suggested that dendrite spacing should increase with solute concentration. In the present work dendrite spacing and morphology have been examined as a function of solute concentration and freezing rate. The freezing rate is defined as the fraction of liquid solidified per unit time, dfs/dß?, where f, is the fraction solid and 8 the time. The fastest freezing rate studied was 4550 times the slowest freezing rate. THEORETICAL CONSIDERATIONS It is of interest to analyze the concentration distribution in the liquid phase between growing dendrites during solidification, Fig. 1. Since this distribution is a direct consequence of the rejection of solute by the growing solid, a diffusional process, the concentration gradients increase with the freezing rate. However, when solidification rate is the only variable in a series of experiments, the interdendritic liquid regions become smaller (i.e., the dendrites become more closely spaced) with an increase in freezing rate. The main purpose of the analytical treatment of interdendritic liquid diffusion will be to reveal a tendency for dendrite spacing to decrease with increasing solidification rate in just such a way that the maximum concentration differences developed in the liquid phase are remarkably independent of freezing rate. Two rather different analyses are set forth, one pertaining to the one-dimensional diffusion which obtains in the interdendritic liquid between parallel plate-shaped dendrites, and the other to the cylindrically symmetrical diffusion around rod-shaped dendrites during early stages of solidification. The results of the two analyses are quantitatively similar, correlating dendrite spacing, maximum concentration difference, and freezing rate. First consider the simpler one-dimensional case. Two parallel plate-shaped dendrites are separated by a distance, L, between centers, Fig. 1. Solidification takes place by the thickening of these plates, with solute being rejected into the liquid. It is assumed there is no diffusion in the solid. This thickening process is slow enough and the dendrite spacing small enough that the concentration differences which develop, although interesting and important, are very small (an important assumption which is verified ex-

Jan 1, 1968

By H. Rumfelt

The author demonstrates an approach for analyzing overcasting requirements for a stripping project. This approach to the problem employs indicated trends in the relationship of the weight of the machine (only electropowered draglines are considered) to its ability to do stripping work. Strip mining in the domestic coal industry is contributing a greater proportion of overall production than at any time in the past. U.S. Bureau of Mines statistics show that the percentage of total production mined in 1958 by stripping methods increased to more than 30 pct, up from 22 pct in 1951. The tendency to favor strip and open pit mining exists in other mining activities, but such trends are not as clearly outlined as in the coal industry. Significantly, the depth of overburden in coal properties considered suitable for stripping is also increasing. A depth of 50 to 60 ft of overburden was once considered to be about the maximum amount that could be handled economically. Today, a depth of 80 to 100 ft is frequently taken into consideration for simple overcasting operations (i.e., excavating overburden and dumping it in the spoil position where it remains indefinitely). Accompanying these changes, there have been continuing studies searching for procedures to appraise deeper stripping problems. Although the only known practicable way to accurately evaluate a proposed stripping venture is to make studies (including cost estimates) step by step, it is nevertheless desirable to have a quick method for making preliminary evaluations of potential strip mining operations. A recently devised digital computer method of providing a preliminary simple overcasting analysis for a stripping prospect relates the indicated trends in the relationship of the weight of the machine to its ability to do stripping work. In addition, this approach employs situations where the geometry of each cut-and-spoil section is assumed to take certain defined relationships for varying overburden depths. Slopes are necessarily considered to be stable, which means such practical factors as the mechanics of soils are neglected for the sake of convenience. Definition of Terms: Photographs A and B show a shovel-type and a dragline-type of simple overcast operation, plus terms used to denote certain parts of a pit. Pits that follow a straight line when projected on a horizontal plane are referred to as straightaway. It has been assumed that all section drawings accompanying this article pertain to straightaway cuts which, in turn, means that areas can be compared by relative volume. The volume of overburden is measured in virgin cubic yards. Material displaced from its virgin state to a spoil area normally occupies a larger volume than when in situ; the difference is termed swell, and it is expressed as a percentage of the original volume. For example, if the original cut volume is denoted V cu yd, and the spoil volume is 1.2 V cu yd, the swell is a positive 20 pct. MAXIMUM USEFULNESS FACTOR (MUF) CONCEPT Men experienced in overburden removal give emphasis to the weight of the machine in relation to its value as an excavator. From the writer&apos;s understanding, value has been only loosely defined and discussed generally and vaguely. The concept is considered important, however, since it is based upon the observations of many men. It provides a clue for the approach taken to solve deep stripping problems. In the procedure described below, value has been arbitrarily established as the product of the nominal dipper size of the shovel (or dragline) times a functional dumping reach. MUFs—Shovel: A shovel&apos;s capability to handle thick overburden is usually limited by its ability to dispose of the spoil. Thus the dumping reach, as well as its dumping height, is significant. The angle of repose of spoil material must also be considered. Although such slope angles vary among different mines, jobs, and materials, a slope of 1.25 to 1, frequently found in both planning and practice, is used by this author in the illustration of this analytical method. The relationship of the shovel&apos;s geometry to a stripping operation is shown in Fig. 1. To gain maximum advantage in constructing spoil piles, the shovel is placed so that its tracks adjacent to the spoil are as close as possible to the rib of the coal deposit. A vertical plane passing through the line

Jan 1, 1961

By F. Bonaccorso, G. Venturello, C. Antonione

A study of the effect of small amounts of interstitial impurities on recovery and re crystallization in high-purity iron (99.995 pct) has been undertaken. This paper gives results on the effect of carbon, introduced in small-dosed amounts @om 0.0005 to 0.0086 wt pct) by heating iron specimens of high purity m a static atmosphere of CO + ,. The materials prepared in this way, cold-worked 80 pct and subjected to a series of isochronal and isothermal annealings, were submitted to examinations by X-rays, micrographs, and hardness tests. It was observed that effect of carbon is remarkable in the sense of blocking recovery of mechanical properties up to the temperature at which re crystallization begins. On the contrary, carbon has a negligible effect on primary re crystallization temperature, when compared with the known effect of substitutional impurities. This is in agreement with the high mobility of interstitials. In effect, only a slight decrease, -20 pct, of the grain-boundary motion rate was noted, due to the interaction between the grain boundary and the carbon atoms. On the other hand, in the samples in which the carbon content is above the solubility limit at the temperature at which re crystallization occurs, a slight increase of nucleation frequency is noted due to the presence of precipitated carbides. ThE effect of purity on recovery and recrystalli-zation phenomena has been known for a long time; however, recently, new attention has been given to the problem since new methods for obtaining metals with extremely low impurity contents have become available. Most of this research work essentially concerns the effects of impurities which cause precipitates or give rise to substitutional solid solutions. The works of Bolling and winegard1 and Aust and utter&apos; on lead, and Vandermeer and orddon&apos; on aluminum, 01sen4 on nickel, and Abrahamson and Blakeney on iron should in particular be referred to. On the effect of this type of impurities a quantitative theory has been formulated by Detert and Lucke6 and has lately been discussed critically by Cahn7 and Gordon and vandermeer.&apos; Interstitial solid solutions in iron have not as yet been studied. As the study of the effect of interstitials is of great interest both from a theoretical and practical standpoint, it was deemed useful to examine the effects of carbon, nitrogen, and boron on recovery and primary recrystallization of iron. There already is some work by Chaudron et al.,1-" on the effect of interstitials on iron; their work, however, mainly concerns secondary recrystallization. The present paper refers in particular to the effect of carbon. EXPERIMENTAL PART Preparation of Materials. The pure iron used for this research was obtained from FeCIS recrystal-lized and purified by extraction with isopropylic ether. From the ferric chloride purified in this manner, hydroxide was precipitated with a solution of very pure ammonia, and, by calcination in pure sintered alumina crucibles, oxide was obtained. Reduction of the oxide to iron sponge was performed in sintered alumina tubes with very pure hydrogen at 650°C; at the end of the operation temperature was increased to 900°C. Specimens for the experiments were obtained by sintering the sponge at 1480°C in pure Hz after a pre sinter ing treatment at 900° C. It is important to note that the treatment at 1480°C in Ha produces a further purification from more volatile elements such as zinc, cadmium, arsenic, lead, and tin. Details on the preparation and characteristics of this type of very pure iron are given in a previous work.&apos;&apos; Only the complete analysis performed by neutron-activation methodsz3 is given here, Table I. Some of the specimens prepared in this way were carburized in the 0 region with very low amounts of carbon by treating them at 700°C in a static atmosphere of Ha containing a definite amount of CO. The set-up used is described in Fig. 1. A gas-tight quartz tube containing the specimen to be carburized and an internal friction control specimen, after being evacuated, was filled with Hz

Jan 1, 1963

By Johnson Crawford, Alan H. Hoagland

Zinc mining at Jefferson City began in 1854 with small scale production of oxidized ore from open pits. Significant production began in 1913 with the development of the Mascot Mine by the American Zinc Company. In 1965, four flotation mills were treating crude ore at the rate of about 12,000 tons per day and the Mascot-Jefferson City district currently is the largest zinc producing area in the United States. The producing mines, Mascot, Young, North Friends Station, New Market Zinc, Grasselli, Davis-Bible, Coy, and the Jefferson City, are 15 to 28 miles northeast of Knoxville in the Valley and Ridge physiographic province, an area of gentle relief lying between the more rugged Great Smoky Mountains and the Cumberland escarpment. The ore is strata bound, occurring essentially within a stratigraphic range of 200 feet in the Lower Kingsport and Upper Longview formations of Lower Ordovician age. In the lower strata of the mineralized section, the ore generally is found in coarsely crystalline elastic dolomite breccia containing silica and chert debris within a sequence of aphanitic limestones. The upper strata of the mineralized section are fine-grained primary dolomites that, in the ore zones, have been fragmented in mosaic patterns by solution collapse with the interfragmental space filled by white gangue dolomite and sphalerite. The mineralogy is unusually simple, sphalerite being the only important sulfide. Pyrite is very sparse. Dolomite gangue is abundant, silica was deposited in quite small amounts, calcite is minor, and fluorite and barite are very rare. The principal ore controlling structures are rubble breccia zones that were formed during the post-Lower Ordovician-pre-Middle Ordovician erosion interval by supergene solution as a phase of the development of a regional karst system of underground drainage that extended to' depths of at least 800 feet. The major Appalachian orogenic structures are post-ore and have been superimposed on the ore bodies. There is substantial evidence to support the current thinking of the writers that the ore deposits were formed at a depth not exc1~eding 800 feet.

Jan 1, 1968