Search Documents

By George F. Comstock

It has been reported recently1 that the hardenability and toughness of forging steels may be improved appreciably by the use of complex deoxidizers containing titanium, aluminum, and vanadium. In the light of these results, an investigation of similar deoxidizers containing various hardening and grain-refining elements was proposed. These deoxidizers included three manganese-silicon-aluminum-titanium al- loys with and without boron and calcium, three titanium-aluminum ferroalloys with vanadium, molybdenum and zirconium, respectively, an aluminum-zirconium ferroalloy, and plain ferroboron. Their approximate compositions are given in Table I, and the results obtained so far in the investigation of their effects on steel are presented in the later parts of this paper. Method of Using Deoxidizers The deoxidizers listed in Table I were used in three series of 0.40 per cent carbon steels, which were melted in a small basic-lined induction furnace and cast in the form of I7-lb. ingots about 2.75 in. square. The basic charge for each heat was 10 lb. Armco iron and 6 lb. of clean low-carbon sheet-steel scrap. When melted, this charge was deoxidized with 15 grams of ferrosilicon, then about 500 grams of pig iron was added to give the desired carbon content. After the pig iron was dissolved, nickel shot and high-carbon ferrochromium were added if required, and additions of 80 per cent high-carbon ferromanganese and 50 per cent ferrosilicon were made to provide the proper amounts of those elements. Aluminum was added to every heat 2 min. after the silicon, for grain-size control and preliminary deoxidation. One minute later the special deoxidizer was added to the clean bare surface of the molten steel, and 2 min. afterward the steel was poured into the ingot mold, using a hot-top and pipe-preventing compound after pouring to decrease the depth of the shrinkage cavity. The special deoxidizers were used mostly in the form of about ½ to¾-in. lumps,

Jan 1, 1942

By H. J. LONGMORE

MORE improvements have probably been made in the diamond-drill field in the past decade than were accomplished in the entire prior period since diamond drilling was discovered in 1864 by a French engineer named Leschot. True, machines were improved, and new accessories were adopted from time to time up to 1929, but no radical changes were made. The basic principle instituted during the early years of the industry were still in use in 1929. At that time, practically the only diamonds used were "black diamonds" or "carbon." These came from Brazil through a rather round- about route, and by the time they got to the drilling trade their cost was excessive-$175 per carat and even higher at the peak. Perhaps because of this. or perhaps because of increased competition, the industry sought ways to reduce expenses even in the face of constantly increased labor costs due to higher wages and reduction in hours worked. The solution seemed to rest in two things-use of a cheaper cutting medium to replace the carbon diamond and the development of new machines to drill faster and thus produce more hole per shift of operation.

Jan 1, 1940

[University of Alaska College, Alaska Mining Society TOM CHRISTENSON, President MARY ANN KISER, Secretary H. G. WILCOX, Faculty Sponsor WM. FACKLER, Counselor University of California Berkeley, California Mining Association JOHN W. COOK, President LORETTA A. RICHTER, Secretary W. S. WEEKS, Faculty Sponsor]

Jan 1, 1946

By E. W. Pehrson, J. W. Furness

THE method of presentation in the accompanying charts is based upon the well-known formula: production plus imports minus exports equals apparent consumption. Thus for each area for which data are shown there are two columns of equal height, one representing production plus imports, the other apparent consumption plus exports. Usually no account has been taken of increases or decreases in stocks. All columns and flow lines are drawn to a linear scale which is supplemented by figures showing the actual quantities represented. Owing to the small area occupied by the European countries on the world map it was impossible to show the details of the trade of the individual countries. All of Europe (except Russia), therefore, is shown as a single unit. However, statistics of production and consumption indicated elsewhere on the charts show the degree of self-sufficiency enjoyed by the major European powers.

Jan 1, 1936

By Nicholas J. Grant, David S. Bloom

An investigation of the Cr-Ni system has revealed that there exists a eutectoid transformation at about 1180°C. 'The transformation in chromium implied by the eutectoid reaction was detected at about 100°C below its melting point. Associated with the eutectoid transformation, a metastable transition structure has been found. THE Cr-Ni phase diagram is shown in current literature to be a simple eutectic system in which the high chromium end of the diagram is incompletely determined, Fig. 1. The last investigation of this system was undertaken in 1936.' Because of limitations in metal purity, especially of chromium, and temperature measuring techniques, the diagram should be accepted with some reservations, as was indicated by Hansen.V or examwle.-, the wresence of the compound Cr²Ni, with a tetragonal structure," remains unverified and incompatible with the simple eutectic diagram. The work reported here began because of discrepancies that appeared in experimental studies in the binary alloys Cr-Ni and in the ternary system Cr-Mo-Ni. The existing binary Cr-Ni diagram failed to provide a complete explanation of observations made by microscope, X-ray, and thermal analyses. Experimental Methods The materials employed were of a higher degree of purity than were available to previous investigators. The Mond nickel used had a nominal purity of 99.87 pct, while the electrolytic chromium analyzed 99.03 pct with the major impurity being 0.5 to 0.7 pct oxygen with some iron. All alloys were melted in a vacuum furnace to remove dissolved gases, sufficient carbon was added to the charge to combine with all the oxygen as CO. Final analysis determined the carbon as about 0.02 pct and the oxygen around 0.06 pct or slightly more. Stabilized zirconia crucibles and thermocouple protection tips were used throughout and contributed no detectable impurity to the melts, as long as the oxygen in the chromium was kept low. Heats averaged about 250 g, which permitted about 2 in. of the thermocouple protection tips to be submerged. The temperature drop across the thermocouple protection tip was determined to be about 3°C, a minor source of error. The alloys, after removal from the vacuum furnace, were placed in a furnace operating in a purified argon atmosphere for the thermal analyses. An- nealed wolfram-molybdenum thermocouples provided the primary means of temperature determination, though a Pt-Pt-10 pct Rh couple was employed for a check when possible. A sensitive, high-speed Leeds and Northrup Speedomax recording potentiometer indicated the thermocouple output. X-ray diffraction patterns were obtained primarily on a Norelco unit, which permitted diffraction angles only up to 45" (28 = 90"). In general, either rotated or stationary solid samples were used, because it later evolved that many samples were susceptible to transformation resulting from the cold working involved in pulverizing. It was also noted that rotating a sample sometimes suppressed weak lines, so, usually, samples were not rotated. Later a high temperature camera was adapted for use at temperatures up to 1400 °C on the Norelco unit utilizing both solid and powdered samples. Unless otherwise noted, all X-ray patterns were made with a chromium tube and a vanadium filter. Further-

Jan 1, 1952

By D. H. Polonis, Gary A. Dreyer

This investigation involved a study of the reactions occurring during aging of supersaturated a phase in a CIL-Si alloy. The aging processes at temperatures below 552°C were studied by means of metallography, microhardness, and electrical resistivity measurements. During the initial stage of aging in the temperature range 275o to 525oC an anisotropic phase, which is visible in the form of striations in the micro-structure formed prior to the appearance of the -phase. This transition phase has been designated k' due to its structural similarity to the k phase. The rate of K' formation during aging increased with aging temperature. At 450o and 525°C the reaction appeared to reach completion before significant amounts of precipitation were observed. The overall precipitation reaction in this system involved four stages: formation of formation of at a grain boundaries, growth of particles in 7-egions, and coagulation of into large particles. The relative role of each of these processes is dependent on the temperature of aging. The formation of phase from supersaturated a during aging increases the electrical resistivity and slightly increases the hardness of a 4.9 pct silicon alloy. The formation of during aging decreases the electrical resistivity and increases the hardness until coalescence of particles begins during the final stages of aging. The coalescence of results in a decrease of both electvical resistivity and hardness. THE available literature does not record a detailed analysis of the aging processes in Cu-Si alloys. Such an analysis has been difficult due to both the limited data available and misleading microstructural observations which have been reported. Previous work suggests that a rather complicated series of processes may be involved in the aging of supersaturated a. The present investigation is concerned with an analysis of the sequence of these aging processes in a 4.9 pct Si alloy. The currently accepted version of the Cu-Si equilibrium diagram, Fig. 1, was published by C. S. Smith in 1940.' This diagram shows a eutectoid at 5.2 pct Si with the high temperature phase, having its composition range entirely to the right of this value. By quenching alloys in the composition range 4.6 to 5.2 pct Si from the region, Smith was able to retain the phase. However, this phase exhibited a striated microstructure which was anisotropic to polarized light. These striations were suggested by Smith to be duplex twins. Using an oscillating crystal X-ray method Barrett2 studied alloys containing 4.0 to 5.4 pct Si quenched from the and regions. The striations observed in quenched specimens were attributed by Barrett to quenching stresses. He found that striations could also be formed by deformation after quenching and identified them as stacking faults in alloys containing 4.0 to 5.0 pct Si. Barrett related the striations in alloys containing 5.0 to 5.4 pct Si both to stacking faults and to a metastable phase which he called The striated structure reported by smith' was also obtained by Hoffmann, et al. 3 By deeply etching their specimens the anisotropic phase on the surface was removed, leaving a structure which was microscopically similar to pure copper. By means of X-ray analysis before and after deformation at room temperature it was concluded that the anisotropic phase was metastable This conclusion has been questioned by Barrett who attributes the discrepancy to superposition of k and lines on the powder patterns of Hoffmann, et alm3 The decomposition of supersaturated according to smith' and Anderson,4 is very sluggish. However, Nony and Dreyer,5 and Hoffmann, et al.' found that deformation increases the rate of decomposition. X-ray studies by Hoffmann, et al on

Jan 1, 1962

By G. G. Brown

Enthalpy-entropy diagrams are presented for natural gases of 0.6, 0.7, 0.8, 0.9, and 1.0 gravity over the pressure range of 5 to 10,000 Ib. per sq. in. and temperature range of 32º to 700°F. The charts indicate directly the work requirement and temperature rise for adiabatic compression or temperature change for free emansion of natural eases. Computation and Uses of Charts The Mollier diagram (Fig. 4), in which the enthalpy (heat content) is plotted against the entropy with lines of constant temperature, pressure,, and in some cases volume, has been found most convenient when dealing with the compression, expansion, and flow of fluids. In dealing with the flow of fluids, the sum of the increase in the enthalpy, AH, plus the increase in kinetic energy, AM —, plus the increase in potential energy, ?MZ, representing the total increase in energy of the fluid in flow, is equal to the sum of the heat q, and work added, —w, to the fluid while flowing between the entrance and exit of the flow system. AH + AM— + ?MZ = q-w [1] In cases where there is no significant change in potential energy or in kinetic energy (velocity), it follows that the increase in enthalpy is equal to the total energy supplied to the fluid. Under such conditions' the changes in the property of the fluid as it flows through a throttling valve, choke, or any other similar arrangement, may be read directly from the enthalpy-entropy diagram by following a horizolltal line between the known pressures' When compressing Or expanding a gas by means of a Or or engine in which no heat is added to or subtracted from the gas, but only work done, the changes in the properties of the gas may be determined along a vertical line of constant entropy between the entering and exit pressures. The power required for the compression of the gas may be readily determined by converting the increase in enthalpy into the desired units. The enthalpy-entropy diagram for natural gases is to the gas engineer what the steam diagram is to the steam-power engineer. For this reason it would be extremely convenient if a reasonably satisfactory enthalpy-entropy diagram could be prepared as a function of the gravity of the gas. A careful study of the known properties of natural gas indicate that this is possible. The effect of temperature upon the enthalpy at constant pressure is expressed as the "heat capacity" or "specific heat" of the gas. The best available data for natural and petroleum refinery gases indicates the relationship shown in Fig. I. From this it is clear that the specific heat of natural gases is a function only of the gas gravity and the temperature at atmospheric pressure. The effect of pressure on the enthalpy of natural gases is dependent upon the pressure-volume-temperature relationships,

Jan 1, 1945

By H. H. Podgurski, J. C. M. Li, Y. T. Chou, F. N. Davis, R. A. Oriani

When an Fe-2 pct A1 alloy is nitrogemted at 500ºC with a gus tnixture (NH3-H2) in which the nitrogen activity has been kept Lou] enough to avoid the formation of iron nitride, a two-phase alloy is generuled which consists of AlN particles and a ferrite phase cotaining a heavy network of dislocations. The amount of nitrogen contained in such an alloy, when equilibrated with the nitrogenating atmosphere, far exceeds both that needed to satisfy the normal solution requirements of a Fe and that needed to convert all of the aluninuwi to AlN. This excess nitrogen is accounted for as being trapped on dislocations, adsorbed at the ferrite -AlN interface, and as an en-Iuznced lattice solubility in strained ferrite. This excessive uptake of nitrogen had previously been attribuled by other investigators to the formation of a nonstoichiornetric aluminum nitride. Isotope exchange experiments revealed various amomts of exchargeable N14 present in the originally nitrided samples that could not be removed by reduction with HS at 500ºC. This exchangeable nitrogen has been identified us that bound to the AlN -ferrite interface. Estimates of inter facial areas in alloys containing -3 pct by weight of A1N are as high as 10 sq m per g of alloy. ThE first1 of this series of papers described the experiments forming the basis for the elucidation of the mechanism of the formation of aluminum nitride particles within an Fe-A1 alloy. It was found that not only are dislocations necessary for the nucleation of the AlN particles but also the nitriding reaction in turn produces a dense network of dislocations in the ferrite matrix. It was also observed that nitrogen in excess of that needed for the formation of stoichiometric AIN is taken up by the alloy without the formation of an iron nitride. The present paper is an analysis of the excess nitrogen sorbed by the nitrided Fe-A1 alloy which considers the structure generated by the formation of the aluminum nitride particles. Because of the high density of dislocations, this system has proved to be quite useful in studying nitrogen-dislocation interactions. Wriedt and Darken2 have already reported such studies in a cold-worked ferritic steel. EXPERIMENTAL Nitrogen Sorption. The specimens of this alloy were cold worked (50 pct reduction) to 0.011 in. thickness and chemically cleaned in a 2:1 concentrated phosphoric acid-50 pct hydrogen peroxide solution before nitrogenation. The flowing nitrogenating atmosphere (11 pct NH3-89 pct H2) was established before the ni- trogenating temperature was reached. The same gravimetric and gas-flow equipment described in an earlier paper1 was also used in this investigation. Changes in nitrogen concentration were followed gravimetrically when establishing the isotherms. In the isotope exchange experiments, concentration changes were followed volumetrically. In some instances chemical analyses (Kjeldahl method) were used to check for material balances in both the gravimetric and volumetric procedures. Our objective was to obtain reversible nitrogen sorption isotherms for alloys equilibrated with NH3-H2 gas mixtures over a large range of nitrogen activity* and temperature. The *Defined as equal to PN H /P3/2 ,where P corresponds to partial pressure in atmospheres. Actually the nitrogen activity, aN, in the alloy equals K PNH3 /p3/2H2, where K is the equilibrium constant for the reaction NH3 = N + 3/2H2. upper limit for nitrogen activity was below that which would produce iron nitride; for reasons which will become apparent later, most of the sorption studies were made in the temperature range between 400" and 500'C. The alloy sample studied most extensively in this investigation was given a series of successive reductions (100 pct H2) and nitrogenating treatments at 500°C to attain a stabilized structure. Presumably some dislocations were lost during these treatments. Throughout most of the sorption studies, temperature was held at ± 1°C and the gas-phase composition was held to k0.2 pet NH3. Based upon the results of numerous diffusion experiments with this alloy* the 'Results to be published in a third paper of this series. times chosen for equilibration were considered more than adequate, Nitrogen Isotope Exchange. The first exchange experiments were carried out by circulating a measured quantity of H2 and NH3* (with a predetermined isotopic *The NH3 was synthesized over a synthetic ammonia iron catalyst using H2, and N2, enriched with N15. AS the NH3 was formed it was removed continuously from the gas phase by circulating the mixture through a refrigerated (78ºK) trap. When most of the nitrogen had been converted to NH3, it was distilled from the trap into a glass storage vessel. composition, N15/N14) over a nitrided specimen at 450°C in a closed glass system. Attempts to reach the exchange limit isothermally, i.e., an isotopic ratio (N15/N14) in the gas phase identical with the extractable nitrogen in the nitrided alloy, were futile. The exchange rate between the gas phase and the alloy was too slow, involving many days. To circumvent the slow exchange with the gas phase, measured amounts of N15 and N14 were introduced into the alloy by nitrogenating with an NH3-H2, mixture at 500°C containing an N15/N14 ratio of 7.76; the charged specimens were then sealed by allowing an oxide film to form over them in air, and in this final condition the specimens were subjected to 500°C in vacuum for various times during which isotope exchange was allowed to proceed within the specimen. It was estab-

Jan 1, 1970

By Donald Warren, J. F. Libsch

HIS investigation originated as a result of a pre-vious experimental study' of the magnetic properties of Fe-Co alloys fabricated by the powder metallurgy technique. Densities of powder compacts prepared for the magnetics investigation varied from 7.45 to 7.70 g per cu cm or from 93 to 95 pct of the experimental value of 8.08 g per cu cm for a fused alloy of the same composition.' While this range of density is considered sufficiently high for most applications, the highest possible density is to be desired for maximum magnetic properties. By applying a technique similar to the one described above to a pure electrolytic iron powder, Rostoker³ was able to achieve a density of 7.895 g per cu cm, which is the highest density ever reported for sintered iron. While Rostoker's work involved the sintering of an elemental powder rather than a mixture, it was believed that higher densities should also have been obtained for alloys using the above technique because of the recoining operation and the high sintering temperature. Consequently, it was decided to investigate the various factors affecting the density of this alloy with the idea that such a study might lead to higher densities and, as a result, powder alloys having magnetic properties identical with those of the fused alloys. It was believed that the principal reason that near-theoretical densities for the powdered alloy were not obtained was the interference of gases with the normal sintering mechanism. When present during the sintering operation, gases can exert several harmful effects: they can remain on the particle surface and interfere with surface diffusion and plastic flow; they can be released and, under certain conditions, expand the void spaces through gas pressure; or they can remain trapped in the pores and exert a hydrostatic pressure that retards elimination of the pores. Jones,4 Rhines,5 Goetzel," and others have given the effect of gases in the sintering of powder compacts an extensive treatment. Among the more important sources of gases in the sintering process are dissolved gases, adsorbed gases, air entrapped during pressing, and gaseous products of chemical reactions. During sintering adsorbed gases are partly released at a relatively low temperature, while those gases entrapped during pressing cannot escape until their pressure is increased sufficiently through increasing temperature to expand the interpartjcle openings. The remaining adsorbed gases, gaseous reduction products, and dissolved gases produce a similar effect at the higher temperatures. If, in the sintering process, gas evolution occurs after the interpore channels have been sealed, an exaggerated expansion of the void spaces results. This is particularly true if the temperature is high enough for extensive plastic flow. In his fabrication of powder bars from tantalum, Balke7 had to consider the effect of adsorbed hydrogen and provide for its escape during sintering by limiting the compacting pressure to a maximum of 50 tons per sq in. The effect of gases entrapped during pressing was first noted by Trzebiatowski8 when he found that gold and silver powders decrease in density with increasing sintering temperature if pressed at 200 tsi, while they exhibit the usual increase when pressed at 40 tsi. Recent investigators9-11 have also noted that entrapped gases have an effect on the expansion of copper compacts during sintering. Proper provision for the escape of gaseous products of reduction must be made in order to avoid deleterious effects. Myers" states that in the sintering of electrolytic tantalum powder, the temperature was gradually raised to 2600°F with a pause at 2000°F to permit reduction of the oxides. Experimental Details For the present study, 50 pct Co-50 pct Fe compacts in the form of circular disks 1½ in. in diam and 0.15 in. thick were fabricated by the pressing and sintering of a mixture of the elemental powders. It was decided to follow the sintering process by means of liquid permeability measurements, because it was thought that such measurements might serve as a measure of relative pore sizes, as well as a possible indication of the point at which most of the interpore channels become sealed. However, since the permeability as measured by the flow of a liquid, such as ethylene glycol, does not give an absolute indication of the point where the pores have become isolated, a method for determining the percentage of pores connected to the surface was set up. As an additional cross check on the permeability measurements, metallographic methods were used to study the relative pore size. Finally, the property of ultimate interest, the density, was measured. Raw Materials: The powders used consisted of an annealed, 99.9 pct pure, —150 mesh grade of electrolytic iron powder, and a 98 pct pure, —200 mesh grade of reduced and comminuted cobalt powder. The cobalt powder was not further processed either by hydrogen reduction or annealing. The screen analyses for the iron and cobalt powders are given in Table I, while the chemical analyses for each type of powder are listed in Table 11. Table 111 gives the hydrogen loss measurements for the powders according to the M.P.A. Standard Method and for a higher temperature as well. Preparation of Compacts: Equal amounts of the elemental powders were mixed by rotation for 1 hr and then pressed into compacts approximately 0.15

Jan 1, 1952

By L. F. Mondolfo, W. T. Collins

A study of the relationship between undercooling for nucleation and structure in Sn-Cu alloys with 0.1 to 5 pct Cu has shown that in hypereutectic allojls the halo of tin that surrounds the primary crystals of Cu3Sn5 is larger, the larger the undercooling for nucleation o,f the tin. This increase of halo size results in a decrease of coupled eutectic, and, in alloys far from the eulectic composition, may produce its complete disappeavance, with the formation of a divorced eutectic structure. This was confirnred by the excrrnination of other alloys in which divorced eutectic slructuves are formed, and leads to the conclusion that they ave only an extrenle case of halo forrtzalion , which results when the two phases freeze one at a time and solidification of the first is completed Defove the second starts. It was also found that under proper conditions of nucleation all types of eutectic structures can be formed in the sartte system , and therefore divorced eutectics, like normal and anomalous, are not characteristic of the syslett~, but are mainly controlled by nucleatiorz. Dizlovced eutectics are formed when the phase that tutcleates the eulectic vequires a large undevcooling for ils nucleation and when the cotnpositiorz of the alloy is far from the eutectic., on the side of the primary phase that does not nucleate the other phase. It is recommended that the tevm "divorced" be used in preference to degenerate because it is more desct-iptice of the morphology and mode of forinalion of the structures. ThE variety of structures found in eutectic alloys has been extensively investigated and classified. The most accepted classification is the one by ~cheil,' in which three different types of eutectic were distinguished: 1) normal, 2) anomalous, 3) degenerate (divorced). ATornlal eutectics are typified by the simultaneous growth of the two phases ("coupling") by which the two phases appear as interpenetrating crystals. The presence of a crystallization front, in which the two phases grow side by side, creates the eutectic grains, with the boundaries where the fronts meet. The presence of eutectic grains is the .distinguishing feature of a normal eutectic, according to Scheil. Straumanis and Brakss2 examined the Cd-Zn system and showed that there was a crystallographic relationship between the phases. Later, others4 also investigated additional systems and found definite crystallographic relationships in the coupled eutectics. The anornalous eutectic shows much less coupling than the normal; the two phases are intimately mixed but 'grow without a uniform crystallization front—a consistent crystallographic relationship— and the eutectic grain is conspicuously absent. As in the normal eutectics faster rates of growth result in a finer structure, but there is not the typical uniform spacing of normal eutectics. The degenerate eutectic shows no coupling; in fact the two phases attempt to minimize their area of contact and to form separate crystals. It has been suggested5" that slow cooling may favor this type of structure. Scheil believes that normal eutectics are formed when the two solid phases are present in more or less equal proportions, whereas both anomalous and degenerate eutectics form when one of the phases is present only in small amounts. spengler7 extended much farther this qualitative relationship between the eutectic type and the ratio of the two phases, and added a relationship to the melting point of the constituents. On this basis he proposed two equations for determining into which of Scheil's classifications an alloy belongs. The first equation is: where TI is the melting temperature of the lower-melting component, Tp of the higher-melting component, and Te the eutectic temperature. The second equations is: where is the volume percent of the lower-melting phase and $2 of the higher-melting phase at the eutectic composition. If 0 and/or 4 are in the range 0.1 to 1, a normal eutectic is formed; if in the range 0.01 to 0.1, anomalous; if less than 0.01, degenerate. Although the examples given by Spengler show a good agreement with the formulas, chadwick found that the Zn-Sn eutectic is normal to all growth rates, even though the volume ratio is 12/1, and Davies9 reports that the A1-AlgCo2 eutectic is normal, with a volume ratio of more than 30/1. Many more discrepancies of this type can also be found. Neither Scheil nor most of the other investigators have considered nucleation as a factor in the formation of divorced eutectics. Daviesg states that divorced eutectics form when neither phase acts as

Jan 1, 1965

By Nicholas J. Grant, John S. Benjamin

A series of composite alloys containing a high volume of NiAl, Ni3Ah or CoAl, bonded with 0 to 40 vol pct of a ductile metal phase, were prepared by powder blending and hot extrusion. The binder metals were of four types: pure nickel or cobalt, near saturated solid solutions of aluminum in nickel and cobalt, type 316 stainless steel, and niobium. Sound extrusions were obtained in almost all instances. Studied or measured were the following: interaction between the alunzinides and the binders, room-temperature modulus of rupture values, 1500° and 1800°F stress rupture properties, hardness, structure, and oxidation resistance. Stable structures can be produced for 1800°F exposure, with interesting high-temperature strength and good high-temperature ductility. Oxidation resistance was excellent. A large number of experimental investigations have been made of the role of structure on the properties of cermets and composite materials. Gurland,1 Kreimer et al.,2 and Gurland and Bardzil3 have indicated the preferred particle size in carbide base cermets to be about 1 µ, with a hard phase content of 60 to 80 vol pct. The optimum ductile binder thickness was noted to be 0.3 to 0.6 µ.1 Complete separation of the hard phase particles by the binder is important in reducing the severity of brittle fracture.' The purpose of the present study was to produce structures comparable to the conventional cermets, using a series of relatively close-packed intermetal-lic compounds rather than carbides as the refractory hard phase, and to study the effects of binder content and composition on both high- and low-temperature properties. The selected intermetallic compounds were particularly of interest because of the potential they offered in yielding room-temperature ductility. The highly symmetrical structures are known to possess high-temperature ductility and room-temperature toughness. Based on a ductile binder, the alloys were prepared by the powder-metallurgy route to avoid melting and subsequent alloying of the matrix, and were extruded at relatively low temperatures. It was expected that the composite alloy would retain useful ductility. In contrast, infiltration and high-temperature sintering led to alloying of the matrix and to decreased ductility. The systems Ni-A1 and Co-A1 were selected for this study. In the Ni-A1 system the compounds NiA1, having an ordered bcc B2 structure, and Ni3Al(?1), having an ordered fcc L12 structure, were chosen. In the system Co-A1 the intermetallic compound CoAl with an ordered bcc B2 structure was used. ALLOY PREPARATION The intermetallic compounds, see Table I, were prepared by using master alloys of Ni-A1 and CO-A1, with additions of either cobalt or nickel to achieve the desired compositions. The master alloy in crushed, homogenized form, was melted with pure nickel or cobalt in an inert atmosphere, cold copper crucible, nonconsumable tungsten arc furnace. The resultant intermetallic compounds were homogenized at 2192°C in argon, crushed, and dry ball-milled in a stainless mill to -100 and -325 mesh for the Ni-A1 compounds and to -325 mesh for the CoAl compound. Finer fractions were separated for some of the composite alloys. Several ductile binders were utilized. These included Inco B nickel, 5µ ; pure cobalt, 5 µ, from Sher-ritt Gordon Mines, Ltd.; fine (-325 mesh) niobium hydride powder; fine (15 µ) type 316 stainless-steel powder; and near-saturated Ni-A1 and Co-A1 solid-solution alloys, also in fine powder form. The niobium hydride was decomposed above about 700°C in processing of the compacts in vacuum to produce niobium powder. The Ni-7.1 pct A1 and the Co-5.3 pct A1 solid-solution alloys were prepared from pure nickel or cobalt and pure aluminum by nonconsumable tungsten arc melting under an inert atmosphere. The ingots were homogenized, lathe-turned to fine chips, and dry ball-milled in air to -325 mesh powder. These solid-solution alloys are designated NiSS and CoSS; see Table I. Subsequently the hard and ductile phases were dry ball-milled as a blend. Experiments clearly established the need to coat the hard particles with the ductile binder to optimize subsequent hot compaction by extrusion. Ordinary dry mixing usually resulted in nonhomogeneous alloys which were quite brittle. Conventional cermets are consolidated by liquid phase sinteiing or infiltration, which resulis in undesirable and uncontrolled alloying of the binder phase. For this study, a loose (unsintered) powder-extrusion process was emploved, minimizing reactions between binder and hard particle, thereby permitting much greater control of composition and structure. The constituent powders were first mixed in the desired

Jan 1, 1967

By M. Kaufman, A. E. Palty

The phase structure and aging characteristics of two nickel-base alloys, Inconel 718 and 702, were investigated. Wrought and cast Inconel 718 showed ArisCb as the major hardening phase, as well as ture was positively identified. Inconel 702 had Ni3Al as the major hardening phase, as well as Cr7C3-of the phase amounts with temperature was determined. 1 HE value of phase structure and aging rate information in the utilization of alloys has been well established. As part of a continuing program, the results obtained on two nickel-base alloys, hconel 718 and Inconel 702, are presented here. Inconel 718 is a relatively new alloy considered for application in both sheet and cast form. Its novelty lies in the use of high columbium tantalum levels in a Ni-Cr-Fe base. No previous work on this type of alloy has been reported in the literature. Inconel 702 is being used now as an oxidation-resistant sheet alloy. Occasional fabrication problems spurred work on its basic structural behavior. The composition of 702 is essentially nichrome plus aluminum and titanium for strengthening and added oxidation resistance. This puts it in the basic class of other nickel-base alloys which have already been studied. The two alloys will be treated separately. INCONEL 718 Phase Study Methods. The phases present after the various conditions investigated were determined by X-ray diffraction analysis of two different elec-trolytically extracted residues and microscopic examination. The extraction methods, using 10 pct HC1 in alcohol and 10 pct H3PO, in water and the X-ray technique have been described previously.' The effect of temperature on the phase behavior was studied on sheet material using the following heat treatments: 2250°F, 2 hr, ice brine quench 1300 100 hr, water quench 2250°F, 2 hr, ice brine quench 1400°, 100 hr, water quench 2250°F, 2 hr, ice brine quench 1500°F, 100 hr, water quench 2250o 2 hr, ice brine quench 1700°F, 48 hr, water quench 2250°F, 2 hr, ice brine quench 1850°, 24 hr, water quench 2250°, 2 hr, ice brine quench 2000°F, 6 hr, water quench The high solution temperature was an attempt to dissolve as many phases as possible without causing melting. Aging exposures were made longer than normal heat treating times in order to approach equilibrium conditions more closely. Cast material was examined only in the as-cast condition and with the normal heat treatment (1700°F, 1 hr, air cool 1325oF, 16 hr,air cool). Materials. The full phase study was performed on sheet specimens approximately 1/2 in. by 2 in. by 1/16 in. The cast pieces were cut from the gating system of a vacuum-cast test piece. An all weld-metal specimen was made by running many weld beads on the sheet using filler made from the sheared edges of the sheet. All external contamination was removed by belt-sanding followed by an electro-polishing treatment. The typical composition of wrought Inco 718 is as follows: The alloy is air-melted. Cast hco 718 may have higher molybdenum. Due to the expectation of finding a Ni-Cb phase (Ni, Cb) for which no X-ray diffraction pattern was available in the ASTM Index, a special 5 Ib heat to the stoichiometric composition of this intermetallic

Jan 1, 1962

By H. W. Meyer, H. N. Lander, F. D. Delve

A method of calculating the changes in blast-furnace performance brought about by burden and/or blast modifications is presented. Essentially the method consists of three simultaneous equutions derived from materials and heat balances. These equations can be used not only to evaluate quantitatively the effect of changes in process operating variables on furnace performance, but also to provide a useful means of evaluating changes in process variables which cannot be measured directly. It has been customary for a number of years to use simple heat and materials balances as a basis for assessing blast-furnace practice. A good example of the method used to set up these balances is that proposed by Joseph and Neustatter.1 This approach to process assessment has limited utility, however, in that it cannot be used to predict the furnace coke rate or production under new operating conditions. Using an approach based on multiple correlation of blast-furnace variables, R V. Flint2 has developed an equation which may be used to predict the change in coke rate that will result from some changes in operating conditions with a reasonable degree of accuracy. Although this equation has useful applications in production planning, it cannot be used to study the relationships between the operating variables and the fundamental thermochemi-cal characteristics of the process. In attempting to analyze the blast-furnace process quantitatively, the idea of dividing the furnace into zones3 may at first appear attractive. In our present state of knowledge, however, it is not possible to define with any accuracy the physical limits of such zones in relationship to their temperatures or to the reactions which may occur in them. Although its application is restricted, the zonal approach to blast-furnace analysis is useful in some instances. For example, the change in the calculated flame temperature in the "combustion zone" caused by injecting steam constitutes information which is helpful in understanding why the addition of steam to the blast is best accompanied by an increase in blast temperature. The zonal approach cannot, at the present time, be used to establish the relationships between process variables and process performance if the whole process rather than part of it is to be considered. One of the earliest approaches to the problem of relating blast-furnace operating variables to pro- duction and coke rate was that developed by Marshall.4 Essentially Marshall's work showed that it was possible to estimate the performance of a furnace by solving three simultaneous equations which consisted of rudimentary carbon and heat balances plus a further equation relating the production, wind rate, and the carbon burned at the tuyeres. Although these equations did not include all of the chemical and thermal variables of the process, their derivation and application seems to be the earliest attempt which achieved any success in relating prior furnace operating data to the calculation of furnace performance under different blast conditions. Work carried out in Germany has been directed mainly towards prediction of coke rates using material and thermal balances rather than statistical methods. wesemann5 used prior furnace operating data as part of the basis for predicting the change in coke rate accompanying a change in burden composition. This author employed a method of successive approximations to estimate the secondary changes in slag volume and stone rate brought about by the change in coke rate. The most recent analysis, which seems to have been developed concurrently with the thermochemical model presented in this paper, has been described by Georgen.6 This author has succeeded in improving on Wesemann's approach by expressing the total changes in the slag volume and stone rate in terms of the change in coke rate itself. This is accomplished in a manner similar to that used in the thermochemical model described in this paper. Although Georgen makes use of a calculated furnace heat loss, he does not relate the heat loss per unit of hot metal to the production rate as is done in the present work. Georgen's approach may be used to calculate the changes in materials requirements accompanying changes in furnace operation; it cannot be used to assess the resulting changes in production. The fact that blast-furnace behavior can be interpreted by consideration of the heat requirements of the process was demonstrated by Dancy, Sadler, and Lander.7 In the analysis of blast-furnace operation with oxygen and steam injection these authors showed that it was possible to account for the changes in production and coke rate

Jan 1, 1962

By O. C. Baptist

The Umiat oil field is in Naval Petroleum Reserve No. 4 between the Brooks Range and Arctic Ocean in far-northern Alaska. The Umiat anticline has been tested by 11 wells, six of which produced oil ; however, [lie productive capacity and recoverable reserves of the field are subject to considerable speculation because of unusual reservoir conditions and because several wells appear to have been .seriously damaged during drilling and completion. Oil is produced at depths of 275 to 1,100 ft; the depth to the bottom of the permanently frozen zone varies from about 800 to 1,100 ft, .so that most of the oil reserves are in the permafrost Reservoir pressures are estimated to range from 50 to 350 psi, increasing with depth, and the small amount of gas dissolved in the oil is the major source of energy for production. Laboratory tests were made on cores under simulated permafrost conditions to estimate oil recoverable by solution-gas expansion from low saturation pressures. The cores were also tested for clay content and susceptibility to productivity impaiment by swelling clays and increased water. content if exposed to fresh water. The results indicate that oil can be produced fronz reservoir rocks in the permafrost and that substantial amounts of oil can be produced from depletion-drive reservoirs by a pre.s.r~lrr drop of as little as 100 psi below the saturation pressure. Freezing of formation water reduces oil productivity much more than that due to increased oil viscosity: Failure of we1ls drilled with rtuter-base mud to produce is attributed to freezing of water in the urea immediately surrounding the wellbore. Swelling clays apparently contributed very little to the plugging of the wells. INTRODUCTION Naval Petroleum Reserve No. 4 lies between the Brooks Range and the Arctic Ocean in northern Alaska. The Umiat oil field is located in the southeastern part of the Reserve and is about 180 miles southeast of Point Barrow (the only permanent settlement in the Reserve and the primary supply point for drilling of the wells at Umiat). Eleven wells were drilled for the U. S. Department of the Navy, Office of Naval Petroleum and Oil Shale Reserves, between 1944 and 1953 to test the oil and gas possibilities of the Umiat anticline. Six of these wells produced oil in varying quantities and the best one pumped about 400 B/D.' Estimates of recoverable oil range from 30 to 100 million bbl. The main oil-producing zones are two marine sandstone beds in the Grandstand formation of Cretaceous age: these are referred to as the upper and lower sands. Good oil shows were found throughout the sand settions in the first three wells drilled on the structure, but the highest rate of oil production obtained on any 01 the many tests was about 24 BOPD. These first wells were drilled with conventional rotary methods using water-base mud; later wells were drilled either with cablc tools using brine or rotary tools using oil or oil-base mud. These experiments were successful as is shown by comparing the oil production from Well No. 2 with that from No. 5. These two wells are only 200 ft apart and are located at about the same elevation on the structure. Well No. 2. drilled with a rotary rig using water-base mud, was abandoned as a dry hole after all formation tests were negative. Well No. 5. drilled with cable tools and reamed with a rotary using oil, pumped 400 BOPD which was the maximum capacity of the pump and less than the capacity of the well. These field results indicated that the producing sands were extremely "water sensitive" and it was assumed that the cause of this sensitivity was the presence of swelling clays in the sands. Because of the very unusual reservoir conditions and the difficulties encountered in completing oil wells in the permafrost. the Navy asked the U. S. Bureau of Mines to make laboratory studies under simulated permafrost conditions to assist them in estimating the production potential of the field and the recoverable reserves. These tests were designed to determine the cause of the plugging of wells in the permafrost and to test oil recovery from frozen sand by solution-gas expansion with the oil gas-saturated at very low pressures. EXPERIMENTAL METHODS AND PROCEDURES Samples Analyzed Core samples were analyzed that represent the lower sand in Umiat Well No. 7, the upper sand in No. 3. and both the upper and lower sands in No. 9. These sands should be productive in all of the wells because of their location on the structure. Core samples from

By A. E. Blakeslee

Morphological and electrical properties of GaAs epitaxial layers are influenced not only by changes in the nominal substrate orientation but also by small amounts of misorientation from the exact crystal planes. Deviations up to 5 deg from {11IA}, {11IB}, and (100) planes were investigated. Growth rates increase progressively with angle, approximately I u per hr per deg. Size and density of growth pyramids fall off with increasing angle, but other effects that are deleterious to the surface may occur which are heightened by increased misorientation. Carrier concentration decreases and electron mobility consequently increases as the angular offset increases, except in the case of strong compensation, where the mobility trend is reversed. It has long been known that changes in the crystallo-graphic orientation of the substrate may cause pronounced effects on the morphological properties of vapor grown semiconductor films. Reports of orienta-tion-dependent growth rates and surface characteristics are as old as the literature on epitaxy itself. shawl has recently published a comprehensive study of the dependence of growth rate on substrate temperature and orientation in epitaxial GaAs. It is also well-known that misorienting the substrate surface a few degrees away from the nominal low-index crystal-lographic plane often produces a much smoother epitaxial surface. This was reported by Tung2 for silicon, Reisman and Berkenblit3 for germanium, and by Kontrimas and Blakeslee4 for GaAs, and use is commonly made of this fact in the semiconductor industry to help guarantee smooth vapor deposits. The effects of substrate orientation on the carrier concentration and mobility of vapor grown GaAs were first documented by williams5 in 1964 and have been observed by several other authors since then,6,7 but no one has yet reported a careful study of how small changes influence these properties. We have made such a study and have found that sizable differences in growth rate, morphology, carrier concentration, and mobility can indeed be observed for epitaxial films grown on substrates that are oriented by progressive small increments away from the exact crystal plane. EXPERIMENTAL Early in the investigation an arsine synthesis system of conventional design8 was employed to produce growths on {111A}-oriented GaAs substrate crystals. In that early work, pronounced effects on carrier concentration and electron mobility were observed as a function of slight misorientation from this low index plane. That observation led to the more careful study that is reported here. An AsC13 system, differing in major aspect from those commonly in use today9 only in that the reactor is vertical rather than horizontal, was used for the detailed study. The gallium source was at 900°C and the substrates were at 750°C. The flow rate of pal-ladium-diffused H2 through the AsCl3 bubbler was 200 cu cm per min, and the flow rate of bypass H2 was also 200 cu cm per min. The substrates consisted of chro-mium-doped semiinsulating GaAs to facilitate elec-trical evaluation of the overgrowth by means of Hall and conductivity measurements on conventional eight-legged Hall bridges. They were misoriented by 0 to 5 deg from the {111A}, {111B}, and (100) planes, toward the (100) from the {111A} and {111B} and randomly toward the <111A> or <111B> from the {loo). The crystals were oriented for sawing by the Laue back-re-flection technique, which is good only to about ±1/2 deg; but after polishing or sometimes after epitaxial growth the wafers were checked by a diffractometer technique which is accurate to about * 0.1 deg. After lapping, the wafers were polished with NaOCl after the technique of Reisman and Rohr,10 and just before use they were cleaned in NaOC1, thoroughly rinsed with de-ionized water, and blown dry with nitrogen. Each run employed four wafers, each misoriented by differing amounts from one of the three major faces, and at least two runs were made for each orientation. The runs were continued long enough to provide at least a 15-µ or thicker layer. SURFACE MORPHOLOGY The appearance of all the films that were grown in a given run always changed from wafer to wafer as a function of increasing misorientation, but not always in the same regular fashion. At least three different trends were observed. These are more easily seen than described, and reference to the series of photo-

Jan 1, 1970

By J. J. Gilman

The data presented indicate that the critical shear stress and strain-hardening Thedatapresentedrate of a zinc monocrystal depend on the orientation of its slip direction with respect to its external boundaries. The tendency of a crystal to form deformation bands also depends on its shape. THE plastic behavior of pairs of zinc monocrystals in which both members of the respective pairs had the same orientation with respect to the longitudinal axis, but each had different orientations with respect to their rectangular external shapes, were compared in this investigation. The purpose of the investigation was to see what influence the shape or surface of a zinc crystal has on its mechanical properties. In a previous investigation of triangular zinc monocrystals,1 anomalous axial twisting was observed which seemed to be related to the triangular shape of the crystals. Wolff,&apos; in 400°C tensile tests of rectangular rock-salt crystals bounded by cubic cleavage planes, found that, of the four equivalent slip systems, the two with the "shorter" slip directions yielded and produced slip lines at lower stresses than the other two. This observation and the work of Dommerich³ as formulated by Smekal4 as a "new slip condition" for rock-salt: "among two or more slip systems permitted by the shear stress law, with reference to the formation of visible slip lines by large individual glides, that slip system is preferred which has the shortest effective slip direction." More recently, Wu and Smoluchowski5 reported essentially the same effect for ribbon-like (20x2x0.2 mm) aluminum crystals at room temperature. Experimental Chemically pure zinc (99.999 pct Zn), purchased from the New Jersey Zinc Co., was the raw material. Glass envelopes, containing graphite molds and zinc, were evacuated while hot enough to outgas the graphite but not melt the zinc. At a vacuum of about 0.2 micron the envelopes were sealed off and then lowered through a furnace at 1 in. per hr so as to melt and resolidify the zinc and produce mono-crystals. One-half of one of the molds is shown in Fig. la. Each mold consisted of four pieces from a cylindrical graphite rod that was split longitudinally and transversely at its midpoints. Rectangular milled grooves 0.050 in. deep and % in. wide formed the mold cavity when the split halves were assembled with twisted wires. Fig. lb shows the specimen shape obtained when the top and bottom mold-halves were rotated 90" with respect to each other. Good fits prevented leakage and excess zinc was necessary to provide enough liquid head to fill the mold completely. In removing soft crystals from the molds it was impossible to avoid small amounts of bending. However, manipulations were carried out whenever possible with the crystals protected by grooved brass blocks. All specimens were annealed prior to testing. From the top and bottom sections of each crystal, X-ray specimens and tensile specimens 7 to 8 cm long were sawed. The tensile specimens were annealed inside evacuated tubes for 1 hr at 375°C. Next the crystals were cleaned and polished by 2-min dips in a solution of 22 pct chromic acid, 74 pct water, 2.5 pct sulphuric acid, and 1.5 pct glacial acetic acid.&apos; Cleaning was followed by a 10-sec dip in a 10 pct caustic solution, then washed in water and alcohol, and dried. This treatment results in a bright surface covered by an invisible oxide film. The testing grips were a slotted type with set screws and were supported in a V-block during the mounting operations in order to avoid bending the crystals. A schematic diagram of the recording tensile-testing machine is shown in Fig. 2. The machine has been described elsewhere.&apos; The head speed was 0.3 mm per sec for all tests. The crystal orientations were determined by the Greninger X-ray back-reflection method with an estimated accuracy of 1. Description of Crystal Geometry A schematic picture of a rectangular zinc mono-crystal is shown in Fig. 3. ABD designates the front edge of a basal plane (0001) of the crystal, the only active slip plane for zinc at room temperature. Of the three possible (2110) slip directions, the active one is indicated by an arrow. Cartesian coordinates are taken parallel to the specimen edges. The normal, n, to the basal plane (n is parallel to the hexagonal axis) has the direction cosines a, ß and ?. X0 = 90 — y is the angle between the longitudinal axis and

Jan 1, 1954

By S. J. McCarroll

Fig. 1—The belt shown at right carries filter cake to mixing station over calciner. Crude ore conveyors appear in right background. THE Three Kids mine, some six miles east of Henderson, Nev., is in a typical southwest desert area, with high dry summer heat and cool to cold winter seasons. The manganese deposit was located during World War I.&apos; During this period 15,000 to 20,000 tons of ore assaying up to 41 pct manganese were shipped. Interest in the deposit was not revived until the middle thirties, when experiments on the ore were initiated. Test work indicated possible recovery of only 70 pct by flotation, but in 1941 additional work was done at the Boulder City pilot plant of the U. S. Bureau of Mines and also by M. A. Hanna Co. As a result, the Manganese Ore Co. was formed and a plant utilizing the SO2 process was constructed. Numerous operation difficulties ensued, and the plant. was closed when the manganese situation in the country eased. In 1949 Hewitt S. West initiated negotiations to acquire the plant. In 1951 Manganese, Inc., was formed and contract entered into with the General Services Administration to supply 27 million units of metallurgical grade manganese in the form of nodules to the national stockpile. A second contract was made to upgrade 285,000 tons of stockpile ore. Test work was undertaken by the Southwestern Engineering Co. and likewise by the Boulder City pilot plant at the U. S. Bureau of Mines. Results obtained indicated the commercial feasibility of the flotation process. Construction of the plant, which is shown in Figs. 1 and 2, was started in June 1951, and operations on a break-in basis began in September 1952. Apart from the usual starting difficulties two major disasters caused serious setbacks, one a kiln failure in February 1953, and the other a fire that destroyed the flotation building in June of the same year. The nodulizing section of the plant resumed operation in November, and the flotation section in January 1954. The ore minerals are chiefly wad,* with minor amounts of psilomelane, and occur in sedimentary beds of volcanic tuff. The ore is overlain with beds of gypsum which outcrop or may be covered with surface gravel. Intermediate beds of red and white tuff occur frequently with lenses of red and green jasper and stringers of gypsum and calcite. Small amounts of iron are present; lead content averages about 1.0 pct and minute amounts of copper and zinc are found. Barite, celestite, and bentonite are present. Since these are made up of minute asicular crystals, moisture content is very high, averaging about 18 pct. Ore reserves have been estimated at 3 million tons averaging 18 pct Mn2 and up to 5 million tons after grade is dropped to 10 pct Mn. A good part of the orebody was stripped of overburden by the previous operating company . Approximately 50 pct of the ore, representing more than 60 pct of the manganese, can be mined by open-cut methods. A system for underground min- . ing has not yet been decided on. Open-cut mining with benches of 20 ft has proved satisfactory. Although the ore is soft and appears dry and dusty it has a certain resilience, probably due to the porosity and moisture which makes drilling and fragmentation difficult. Wagon drills have been abandoned in favor of the Joy 225-A rotary drill which will put down a 43/4 -in. hole at the rate of 2 ft per min. Holes are spaced in a pattern with 8 to 9-ft centers. Forty percent powder has been used, but better breaking to 2-ft size is obtained with low velocity bag powder of 30 pct strength. Loading is done with one 21/2-yd shovel, and cleanup follows with one D-7 bulldozer. The ore is hauled with Euclid trucks about 1000 ft from the pit to a blending pile, where the daily mine production is spread in layers by bulldozing until approximately one month&apos;s mill feed is accumulated. A new pile is then started and mill feed is drawn from the first pile by one 13/4-yd shovel and Euclid trucks, with a haul of approximately 500 ft. Mining is performed by an independent contractor with engineering and supervision by the company staff. Early test work indicated that the manganese could be floated with soap, a wetting agent, and fuel oil to give a recovery of better than 75 pct with a grade of 43 pct Mn. The concentrate when nodulized with coke would upgrade to 46 pct Mn or over, and the lead volatilized to 0.6 pct residual.

Jan 1, 1955

By K. P. Mukherjee, K. P. Gupta, J. Bandyopadhyaya

Even though a lot of work has been done in the past to establish phase equilibrium in the Cu-Si system a re cent investigation casts some doubt about the existence and crystal structure of some of the phases that form in the composition range of 15 to 25 at. pct Si in Cu. The present investigation was carried out using high temperature X-ray diffraction technique along with other standard techniques to study the phases in this composition range. The high temperature 6 phase appears to be tetragonal with parameters a,, = 8.815A, c, = 7.903A, and co/ao = 0.896. The reported bcc E phase exists at room temperature and at least up to 780°C and appears to undergo a transformation near 600°C. The phase appears to be cubic but not of the bcc type. The ? phase appears to undergo a transformation, as has been indicated by earlier investigators, and the low temperature form of .? phase is tetragonal with parameters a, = 7.267A, co = 7.8924, and co/ao = 1.086. THE Cu-Si binary system has been investigated by several investigators1" and several intermediate phases,?,e,?&apos; at lower temperatures and ?,ß,0,e, and ? at higher temperatures, were observed between terminal solid solutions of copper and silicon. Even though the existence of the e phase and the transformation in the ? phase were reported in many early works, in a recent study of this system Nowotny and Bittner6 doubted the existence of the e phase and phase at 550°C. Among the high temperature phases, the 6 phase was reported to have a complex cubic structure with parameter a, = 8.805A.7 Nowotny and Bittner, however, suggested that the structure of the 6 phase might be of CsCl type. In order to check these contradictory reports the present study was taken up to investigate the Cu-Si binary system in the composition range of 17 to 25 at. pct Si. EXPERIMENTAL PROCEDURE Weighed amounts of copper (99.99 pct) and silicon (99.9 pct) were induction melted in recrystallized alumina crucibles under argon gas atmosphere. The alloys containing 17, 18, 20, 21, 21.2, 22, and 24 at. pct Si were annealed in evacuated and sealed quartz capsules at 700°C for 3 days and subsequently water quenched. Other than this annealing, the 21.2 at. pct Si and 24 at. pct Si alloys were annealed at 550°C for 10 days, the 17 at. pct Si alloy was annealed at 750°C for 3 days, and the 22 and 24 at. pct Si alloys were annealed at 780°C for 2 days. All annealing temperatures were controlled to within *l°C. Alloys after quenching were subjected to metallographic and X-ray diffraction investigation. A solution containing 5 g FeC13 + 10 cc HCl + 120 cc H2O diluted with six times its volume with water was used as etching reagent. A 114.6 mm diam Debye Scherrer camera was used for obtaining diffraction patterns. The 17, 21.2, and 24 at. pct Si alloys were subjected to high temperature diffractometry using a Tempress Research High temperature attachment and a GEXRD VI diffractometer. For the 6 phase (17 at. pct Si alloy) powder specimen from a 750°C annealed alloy was reheated to 750°C in the high temperature attachment for 1½ hr before taking a diffraction trace. A 550°C annealed and slowly cooled phase (24 at. pct Si) alloy was first reheated to 550°C. a diffraction trace was made after annealing it for 2 hr, and subsequently it was heated to 716OC and kept at this temperature for 2 hr before taking a diffraction trace. For the e phase (21.2 at. pct Si alloy) a 550°C annealed and slowly cooled specimen was heated first to 425°C and annealed at this temperature for 2 hr before taking a diffraction trace. Subsequently, the specimen temperature was raised to 495", 540°, 603", 635", 682", 720°, and 748°C and homogenized at each temperature for 1 hr before taking diffraction traces. The powder specimen temperature was controlled to within +2oC at each temperature and argon gas, purified by passing it at slow rate through a fused CaC12 column, hot (800°C) copper and titanium chips and finally through a P2O5 column, was used to prevent oxidation of the powder. For all X-ray work copper-radiations at 25 kv, 15 ma (for Debye Scherrer technique), and 40 kv, 20 ma (for diffractometer tech-nique) were used. RESULTS AND DISCUSSION At 700°C the alloys containing 17 to 21 at. pct Si showed two phases while the 21.2 at. pct Si alloy was found to be single phase. The X-ray diffraction patterns of the two-phase alloys were consistent with the phase (ßP-Mn type structure) and the phase (21.2 at. pct Si) patterns. The diffraction patterns of the 17 at. pct Si alloy quenched from 750" and 700°C were identical. According to the accepted Cu-Si phase dia-gram4,5,10 the 17 at. pct Si alloy at 750°C should be in the (k + 6) two-phase region and very close to the -phase boundary. The identical patterns possibly resulted from the decomposition of the 6 phase on

Jan 1, 1970

By Forrest V. Williams

The dislocatiorl and stacking-fault structuve of epitaxial GaAs1-,PX lms been examined by chemical etching. The layers were groun in the (100) direction and etch Pils were developed on (111} planes which nad been lapped and polished on the epiLaxia1 layevs. Tile effecL of the jollolcing cariables on the quality of the epilaxial layers has been examined: doping leuel, grouth rate, and composition. High stacking-faullL densilies weve found in the GuAsi_xpx layers. These are not observed in heavily dolled epitaxial layers tzar in layers with low phosphorus compositions. The dislocatiorz density in GuAsi-x px was highest at the sub-stvate- epilaxia1 layer interface. Composilion changes introduced dislocations in the epitaxial layers. ManY semiconductor p-n junction lasers of Group TIT-Group V compounds and their alloys have been reported in the past several years. Laser action at visible wavelengths in GaAsl-x,Px was first reported by Holonyak and Bevacqua. GaAs, a direct transition semiconductor which lases, and Gap, an indirect transition semiconductor which does not lase, form a continuous series of solid solutions.2 Laser junctions can be fabricated in GaAsl-xPx crystals with phosphorus compositions up to about 40 mol pct. In addition to the production of coherent radiation in these crystals, the efficient recombination radiation of p-n junctions in this material has equally important potential in the development of low-power semiconductor lamps. To achieve a high conversion efficiency of electrical to optical energy in p-n junctions in this material, the relation of physical properties of the crystal to luminescence efficiency must be better understood. Although the electrical, optical, and device properties of GaAsl -xPx junction lasers are understandably of considerable interest, the work to date indicates that the more serious problems are the chemical and metallurgical difficulties encountered in the growth of this material.3 In addition to the problems of chemical purity, crystal imperfections, such as dislocations and stacking faults, can be expected to affect both the efficiency of the radiative recombination process and the perfection of the p-n junction.3 The last requirement, i.c., that of the perfection of the p-n junction, is a particularly troublesome one in the fabrication of laser diodes. To obtain good laser diodes, the p-n junction must be flat, which permits the radiation to be reflected from the resonant cavity boundaries. Junction planarity is extremely sensitive to the crystal perfection of the semiconductor material. Also, it is known that at high dislocation densities (-105 per sq cm) it has not been possible to build laser junctions in GaAsl-xPx . Few studies have been reported on the crystal defect structure of GaAsl-,P,. The first serious study seems to be that of Wolfe, Nuese, and Holonyak,3 who examined the dislocation structure of monocrystalline bulk (nonepitaxial) material grown by halogen vapor transport. In this paper are reported some observations on the dislocation and stacking-fault structure of GaAsl_,P, crystals grown by a vapor transport process on substrates of GaAs. EXPERIMENTAL Crystal Growth. The GaAsl-xPx crystals were grown in an open-tube flow system, using two sets of reagents. GaAs, Pr(red), and HC1 were employed in one method. The transport reaction is =950JC GaAs+HC1 = GaCl +1/4As4 +1/2H2 and the deposition reaction is 2GaAs1-xPx +GaCl3 Composition control is obtained by the flow rate of the HC1 and the vapor pressure of the P4, which is maintained in a separately controlled furnace. The second method has been described by Ruehr-wein4 and utilizes gallium, AsH3, PH3, and HC1. The same transport and deposition reactions as above are involved. Composition control is obtained solely by the flow rates of the three gases involved. All of the crystals were grown on chemically polished GaAs substrates oriented on the (100) plane. The thicknesses of the epitaxial layers were typically 100 to 300p. Revealing of Dislocations. Dislocations were re-vealed on both the( 111 ) and { l l l }b faces by chemical etching. The specimen to be examined was mounted at 54.7 deg, lapped on glass with 3-p alumina, polished on cloth with 3-p diamond paste, and, to remove work damage, chemically polished at room temperature for

Jan 1, 1968

By S. L. Craig, C. Orr, H. G. Blocker

WITHIN recent years gas adsorption methods have been developed for measuring the surface area of finely divided materials and have become extremely valuable in research on the corrosion and the catalytic activity of metals. Rather elaborate apparatus is required, and a single determination is so time-consuming that these methods have not been utilized to the fullest extent; the methods are un-suited for most routine control work such as that encountered in powder metallurgical operations and in processes employing metal catalysts. These difficulties are largely eliminated, and surface area is reduced to a routine determination if the liquid-phase adsorption of a surface-active agent such as a fatty acid can be used. When the affinity of the fatty acid carboxyl group for the solid surface is greater than its affinity for the solvent, a unimolec-ular layer of orientated fatty acid molecules will be formed at the solid-liquid interface in a manner similar to that of a compressed fatty acid film on a water surface. The measurement of surface area is then reduced to a measurement of fatty acid adsorption. This propitious circumstance, first investigated by Harkins and Gans,¹ has been employed with somewhat inconclusive results by a number of investigators in evaluating the surface properties of metals, metal catalysts, and metal oxides. The specific surface area values for nickel and platinum catalysts, determined from the adsorption of a number of fatty acids from various solvents, were found by Smith and Fuzek² to agree with values calculated by the gas adsorption technique of Brunauer, Emmett, and Teller," he so-called BET technique. And recently Orr and Bankston4 have also reported good agreement between nitrogen gas and stearic acid adsorption results in the measurement of the surface areas of clay materials. On the other hand, Ries, Johnson, and Melik5 found only order-of-magnitude agreement between these two methods in studying supported, cobalt catalysts having specific surface areas as great as 420 sq m per g; the reason is partially attributable to the very porous nature of the materials. Greenhill,6 investigating the adsorption of long-chain, polar compounds in organic solvents on a number of metal powders, concluded that a uni-molecular layer of stearic acid was formed on exposure of the solid to the acid solution and that the presence of an oxide or another film did not alter this result. Furthermore, the adsorption process appeared to be the same whether or not the sample was degassed prior to exposure to the solution. Greenhill estimated the surface area of one of the powders he investigated from microscopic diameter measurements, and obtained a rough check with surface area evaluation. Russell and Cochran7 found moderate agreement for alumina surface area results by fatty acid and gas adsorption methods. In addition, they also found that the prolonged heating and evacuating pretreatments previously used by investigators were unnecessary. The present work, however, considerably extends these previous investigations, shows that fatty acid adsorption can be used to determine the surface area of a variety of metals and metal compounds, offers further confirmation of the correctness of gas adsorption methods, and presents a simplified technique for the determination of the metal surface area which is suitable for routine work. Experimental Technique Basically, the fatty acid adsorption method is quite simple. It consists of exposing a sample of the material of which the surface area is desired to a fatty acid solution of known concentration. By analysis of an aliquot of the solution, the concentration after adsorption has occurred may be determined. The difference between the initial quantity of acid in solution and the final quantity is that quantity of acid adsorbed by the sample. The specific surface area of the adsorbent material may be calculated from the quantity adsorbed and the weight of the sample. In agreement with the findings of others as outlined above, it was found entirely unnecessary to degas or pretreat the nonporous materials employed other than by drying them thoroughly. However, precaution was necessary so that the dried sample entered the fatty acid solution with little exposure to moisture. The effect of moisture on the interaction of stearic acid with finely divided materials has been thoroughly investigated by Hirst and Lancaster." They found the presence of water merely reduced the amount of acid adsorbed by powders such as TiO2, SiO2, Tic, and Sic. With reactive materials such as Cu, Cu2O, CuO, Zn, and ZnO, however, water was found to initiate chemical reaction. Only with ZnO was reaction observed when the solid and the solu-

Jan 1, 1953