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The total crude production from the Dutch East Indies, and the British Protectorates of Sarawak and Brunei for 1932 is given in the following table: Production, which had shown a rapid increase from 1921 to 1929, has since been reduced. Exploitation drilling in the majority of the larger fields was suspended and existing production kept within predetermined limits by beaning. Most of the smaller fields were shut down entirely. The depth of producing levels varies widely in the different fields. Shallow fields with drilling depths to 500 meters (1650 ft.) are still numerous as well as those with depths between 500 and 1000 m. (3300 ft.). A few of the newer fields reach greater drilling depths up to 1600 m. (5250 ft.). The deepest exploration well drilled to date (Djambi field, South Sumatra) reached a, depth of 2332 m. (7650 ft.) before abandonment in March, 1930. The specific gravities of the various crude oils produced range between 0.65 (86.2° A.P.I.) and 0.95 (17.5° A.P.I.). The lighter crudes have a high gasoline content and fuel oils are derived from the heavier crudes, which usually are rich in paraffin. Producing companies are: De Bataafsche Petroleum Maatschappij (Royal Dutch-Shell combine in Dutch territory); The Anglo-Saxon Petroleum Co. Ltd. (Royal Dutch-Shell combine in British territory); De Nederlandsch-Indische Aardolie Maatschappij (De Bataafsche Petroleum Maatschappij-Dutch East Indian Government combine); De Nederlandsche Koloniale Petroleum Maatschappij (Standard Oil Co. of

Jan 1, 1933

By Byron W. Cheever

In the Journal of the Iron and Steel Institute for 1881 (vol. ii., page 379), will be found an article upon this subject. The analyses there reported mere of samples taken from an ingot made especially rich in impurities and then cast in a sand-mould in order to give the elements a better chance to redistribute themselves, on account of the steel solidifying more slowly than when cast in an iron mould. Wishing to ascertain how much, if any, segregation takes place when the steel is cast in the ordinary way, I placed the subject in the hands of Mr. W. J. Olcott, a graduate of the Michigan University, who has carefully collected and analyzed samples taken from different parts of ingots and billets. In this undertaking he found himself somewhat embarrassed on account of being distant from steel-works, and therefore unable to supervise the selection of samples. He found, however, a competent and willing aid in the person of E. F. Wood, M.E., chemist of the Pittsburgh Bessemer Steel Company, Limited, who has spared no pains in collecting the desired samples and furnishing drawings and descriptions of the ingots from which the samples were taken. The following are the results obtained: Ingot No. 1. This ingot was about 5 feet long and 13 X 131/4 inches at the point of fracture, which was 2 feet from the bottom. The samples

Jan 1, 1885

Another excellent Swedish underground crushing plant is that at Crusher Station 51 of LKAB in Kiruna, Sweden. Bottom-dump rail cars pass over a dump scroll to discharge magnetite ore into a 288 cu yd (220 m3) pocket equipped with a Hewitt Robins E-13 vibrating grizzly feeder. The grizzly section is 59 in. (1500 mm) long and has bar spacings 3 in. (76 mm) wide. Any oversize boulders which will not pass the chute opening to the feeder are blasted. The primary breaker is a Morgardshammar overhead eccentric 55 in. by 71 in. (1400 by 1800 mm) jaw crusher powered by a 200 hp (149 kw) motor. An open side setting of 10. 8 to 11. 8 in (275 to 300 mm) accounts for the high capacity of this crushing unit (2760 stph or 2500 mtph). This figure compares favorably with the capacity of many gyratory crushers but is typical of the Swedish jaw crushing installations, which are often followed by three more stages of crushing or by autogenous grinding. Crushed ore from the jaw crusher is combined with the grizzly under- size in the crushed ore pocket and is fed by a second vibrating grizzly feeder to the secondary cone crusher. Both crushers are serviced by 50/5 ton overhead bridge cranes. During maintenance operations air -operated gates shut off the flow of ore from the dump and primary crushed ore pockets.

Jan 1, 1978

By L. D. Hall, S. J. Rothman

DIFFUSION coefficients of RaD (PD) and RaE (Bi) in a Bi-0.255 atomic pct Pb alloy have been measured by the methods outlined in a previous paper.' The following changes were made in the experimental procedures: the bath was made up of about 400 to 500 g Bi; and no correction was made for the volume change of bismuth on melting. The data are presented in Table I and Fig. 1. A least-squares analysis of the data yielded D = (0.120 t 0.09) exp ( -9680 + 1070/RT) D- (0.0465 0.001) exp (-8100 + 466/RT). Statistical analysis of the data indicated that the activation energies are significantly different at the 50 pct level, but not at the 75 pct level, while the D values are significantly different at the 97.5 pct level, but not at the 99 pct level. The differences are thus not as clear-cut as they were for lead. However, bismuth diffuses faster than lead in high bismuth alloys as well as in lead. The high activation energies, more than twice the value calculated from viscosities, are in good agreement with the results of Grace and Derge.' No data were obtained at higher temperatures because of gas bubble formation in the capillary. Acknowledgments This work was supported by the Office of Ordinance Research under contract DA-04-200-ORD-178. The authors thank 0. P. Vieria and M. Volpe for help in the experimental work.

Jan 1, 1957

By N. A. D. Parlee, I. D. Shah

THIS communication reports the results of some work done on the effect of gold, palladium, and platinum on the liquid-gas-solid equilibrium in the Ag-O system,1,2 see Fig. 1. Recently several investigators3, 4, 5 have reported the effects of gold, palladium, platinum, and copper on the solubility of oxygen in liquid silver. This communication represents only a portion of the work done here on the solubility of oxygen in liquid silver and liquid Ag-Au, Ag-Pd, and Ag-Pt alloyss—the greater part having already been reprted. The apparatus, materials, and procedures used were similar to those of earlier work.' l The main precautions observed in obtaining these freezing points at constant temperature and PO2 = 1 atm were as follows: 1) The points were approached very slowly from higher temperatures. 2) In the very near vicinity of the L-G-S equilibrium point precautions were taken to see that the melt was in equilibrium with oxygen at one atm. Equilibrium was considered to have been attained when no more absorption of oxygen was observed in the last 15 min at a fixed temperature. 3) In the case of these alloys, the freezing point of the melt at 1 atm of oxygen and the equilibrium amount of oxygen absorbed by the liquid "infinitesimally" above this point were determined by lowering the temperature in very, very small steps, equilibrating at

Jan 1, 1969

By T. R. Cass, M. R. Achter

A technique for measuring the creep and rupture strength of nickel specimens bonded by sintered oxide layers has been developed for the investigation of the role of grain-boundary oxide in the oxidation strengthening of nickel during creep at 817 °C. Prom the inverse dependence of the rate of bonding on the thickness of the starting oxide, it is suggested that the rate of diffusion of Ni++ through NiO is the controlling process at the start of sintering. Cornparisons of the creep and rupture strength and of the microstructure of sintered couples with conventionul creep specimens suggest rupture in nickel is prevented by load-bearikzg, ikltergranular layers of oxide. A number of recent investigations have shown that at high temperature and low stress the creep-rupture strength of nickel and nickel-base alloys may be much greater in air than in an inert gas or vacuum.'-7 For example, at 817°C and 2500 psi, nickep failed in vacuum in 48 hr but not in air after 4000 hr, and a Ni-Cr-A1 alloy5 at 1038°C and 3000 psi failed in vacuum in 4 hr but not in air after 731 hr. Fig. 1 is a typical creep curve in air for materials experiencing this pronounced environmental strengthening. A common feature of these curves is the entry into a third stage of accelerating creep, concurrent with the for-

Jan 1, 1962

ON FEBRUARY 13, 1953 the Miami Copper Company signed a contract with the United States Government, through the Defense Materials Procurement Agency, whereby it under- took, with a base price of 27.359 per lb (subject to escalation based on wage scales and wholesale-price indices), to exploit a new ore body in its mine in Gila County, Arizona. Recoverable ore was estimated at 23,000,000 tons, averaging about 0.677 per cent copper. Moreover, the proportion of oxidized copper, difficult to recover by flotation concentrating, was so high that an average net recovery of only ten pounds per ton (0.5 per cent) was expected. The company proposed to mine the ore by the underground method known as undercut block-caving. Unfortunately the ore body, as has been true of all Miami's ore bodies, lay so deeply buried that no possibility existed of adopting opencut mining as has been done in recent years at Miami's neighbors, Ray and Inspiration. Probably the most distinctive characteristic of Miami's operations has been the extent to which it has developed and improved methods and equipment for underground mining. True, necessity was a potent stimulant, because its available ores since 1926 have been exceptionally low-grade; so, if the enterprise were to survive, it was necessary to keep costs at the minimum. This fact does not detract, however, in any way from the remarkable achievements of Miami engineers in conducting efficient and low-cost underground mining.

Jan 1, 1957

By Norman Pilling

The deleterious effect on the mechanical properties of copper, resulting from heating in contact with reducing gases, is well known, but the mechanism of the action does not appear to have been definitely established. The purpose of the present investigation was not to determine the exact variations in physical quality of commercial copper, with its various impurities, as dependent upon the composition of the gas, temperature of exposure, etc., but rather to study the nature of the action and the conditions under which it occurs. The physical result of the so-called "gassing." of copper has been generally described as the development of "brittleness." There appears to be a striking decrease in tensile strength, associated with a considerable loss of ductility, and a small copper strip or wire affected in this way can he snapped in two when bent. Larger articles may have an outer layer alone affected, which can be detached from the core by hammering or bending. Perhaps the earliest reference to this subject is that of Heyn1 in 1900, and among the later observers may be mentioned Archbutt,2 Bengough and Hill,3 Johnson,4 Mathewson and Thalheimer,5 and Ruder .6 Bengough and Hill, while working with arsenical copper, were the first to suggest the possibility of the formation of gases within the copper itself, while Ruder endeavored to explain the action by the reduction of the

Jan 10, 1918

By Eugene S. Meieran

The effects of methods of crystal growing, wafer sawing, polishing, routine handling, diffusion, and epitaxial growth on the defects in GaAs are reviewed and studied using reflection and transmission X-ray topographic techniques. In general, it was found that boat-grown crystals exhibited fewer defects than Czochralski crystals, although all crystals showed large numbers of precipitates visible when examined in the electron microscope. Mechanical surface treatments such as sawing and mechanical polishing introduce damage to a depth of about 5 µ, most of which can be removed by suitable chemical or chem-mechanical polishing. In addition, defects can be introduced through routine handling of wafers, for example with metallic tweezers. These defects can be quite severe, and have been observed 20 µ below the wafer surface. Defects can also be introduced through diffusion and epitaxial growth. These defects, which include precipitates, growth pyramids, stacking faults, dislocations, and so forth, can be detrimental to device fabrication. It is shown that wafers or films which appear defect-free optically can contain defects visible in the X-ray topographs. WHILE the use of GaAs in the semiconductor industry has increased very rapidly in the last few years, due mainly to the recent development of many important GaAs devices,1,2 the major limit to the production of commercial quantities of many GaAs devices remains a severe lack of suitable materials technology. This lack is apparent in two critical areas. First, production quantities of high-quality GaAs crystals, reproducibly doped and precipitate-free, simply are not available commercially, although some reasonable quality material is available on a limited first-come, first-serve basis. Second, in comparison to silicon technology, little is known about the effects of processing variables on the defects either present in as-grown GaAs or introduced through processing and handling of wafers. These areas are now receiving some attention from semiconductor device manufacturers, who are studying defects in GaAs in order to better understand how either to prevent their occurrence or to cope with their existence. Most investigations of the defects in GaAs have been made by optical microscopy3-5 or transmission electron microscopy techniques.'-' Recently, however, the imaging techniques of X-ray topography, electron mi-croprobe analysis, and scanning electron microscopy are being applied to the study of GaAs.9-14 In the case of X-ray topography, a one-to-one image is obtained that must be photographically enlarged. In compensa- tion, the defects within entire wafers may be imaged by simple scanning (Lang technique15) if the wafer is reasonably perfect, or by using the scan oscillation technique developed by Schwuttke16 if the wafer is warped or distorted. The purpose of this paper is to both review and extend the general application of X-ray topographic techniques to GaAs. Emphasis will be placed on the effects of growth and process variables on the quality and perfection of both bulk and epitaxial GaAs. Reference to optical or electron microscopy results will be made when useful. Since the effects on defects of a wide variety of processing variables such as crystal growing, sawing, polishing, diffusion, and epitaxial growth will be somewhat superficially reviewed, a fairly extensive bibliography of the most important recent results in these areas is included. However, for completeness, important defects will be illustrated here, although such defects have been previously shown by others. While this paper is concerned with defects rather than with the physics of X-ray scattering, the mechanisms of contrast formation in the topographs will of necessity be briefly mentioned. EXPERIMENTAL GaAs crystals, both boat-grown18 and Czochralski-grown,'8 containing a variety of dopants of various concentrations, were purchased from outside vendors. Wafers were sliced from the crystals using a Hamco ID saw and were mechanically polished using 1 µ diamond paste. Chem-mechanical polishing was done in bromine-methanol as described by Sullivan and Kolb.18 Chemical polishing was done using a modified sulfuric-peroxide solution, 11 parts H2SO4, 1 part 30 pct H2O2, 1 part DI water.5 Zinc diffusion was carried out in a closed tube, using a 10 pct Zn-In source at 825°C for 1 hr. Oxide masking techniques were used to select the area to be diffused. Epitaxial wafers were either purchased or prepared here. All epitaxial runs prepared here were carried out using a Ga-GaAs-AsC13 source in a closed tube at a substrate temperature of 750°C. Wafers were chem-mechanically polished and gas-etched prior to deposition. The X-ray topographs were taken on a Krystallos Lang camera, operating in the transmission scanning geometry (Lang technique15) or in the reflection scanning geometry (modified Berg-Barrett technique20,21). MoKa, radiation was used for all transmission topographs using a Jarrell-Ash 100-µ spot focus. CuKal radiation was used for all reflection topographs using a General Electric CA-7 1-mm spot focus X- ray tube. Topographs were printed from an intermediate contrast inversion film, so the contrast shown in all figures here is the same as that of the original 50-µ-thick emulsion L4 Iiford nuclear plate used to record the topograph.

Jan 1, 1969

By B. W. Christ, P. M. Giles

Mossbauer effect measurewents have been made on I-mil-thick foils of commercial 1 wt pct C steel and Fe-2 wt pct C alloy. The experimental method required about 3 to 5 vol pct of a phase in the nzultiphase steel sample for detection. Room-temperature Md'ssbauer patterns obtained on austenitized atid quenched samples exhibit fifteen, and possibly twenty-one, lines. A sharp parama&tetic singlet and a quadrupole doublet, poorly resolued from the singlet, are attributed to austenite. Remaining lines are due to tnartensite. Accurate evaluation of austenite line paranzeters is not feasible if significant amounts of other phases such as carbides or martensite occur simultaneously with austenite. This is demonstrated by comparison of hyperfine interactions determined for austenite in multiphase high-carbon samples with those reported for Fe-C austenite in a nearly 100 pct austenitic sanple. Lines from carbides are incompletely resolced from austenite lines, as demonstrated by comparison of austenite line positions with carbide line positions calculated frow published values of hyperfine interactions. One martensite line overlaps an austenite line in the pattern for commercial 1 wt pct C steel. Results of this study suggest that the usefulness of e M6ssbauer pectroscopy for quantitatizle analysis of austenite in bulk samples of quenched and tempered high-carbon steels is restricted by poor resolution. Use of Mossbauer spectroscopy for phase identification and for evaluation of atomic and electronic structures appears quite feasible, however, The Mijssbauer effect has been widely discssed,'-and e Mossbauer effect measurements have been reported on materials of metallurgical interest.7"20 In particular, it has been proposed that e Mossbauer patterns of commercial steels and laboratory-made Fe-C alloys, in the quenched condition, are composed of lines originating in two phases, Fe-C austenite and Fe-C martenite.-' Evidence accumulated in this study demonstrates that three absorption lines found in the central region of the e Mossbauer pattern obtained on quenched steels are attributable to retained austenite. This interpretation is supported by parallel decreases in the intensity of these three lines caused by subambient cooling of commercial 1 wt pct C steel samples after water quenching to room temperature. A second result of this study is to clarify effects of line resolution and sensitivity in the Mossbauer patterns of multiphase steels on the accu- rate determination of austenite line parameters. Experimental line widths (full width at half height) are generally 1.5 to 3 times larger than the natural line width of 0.19 mm per sec. At least two lines, and sometimes more, from a single phase such as cementite (Fe3C), other carbides, martensite, and austenite fall in the energy band, i0.85 mm per sec. hhis band width is employed simply for convenient reference. It represents approximately the energy interval between the + 112 to 112 transitions in ferrite and is expressed as the velocity needed to Doppler shift a 14.4kev 7 ray to the aforementioned ferrite energy levels. This energy band is referred to as "the central region of the Mossbauer atttern:: in this paper. Hence, due to the large number of lines from different phases in a multiphase steel falling in a relatively narrow energy band, absorption lines from the different phases may overlap. Analysis of available data, presented below, indicates that this occurs to a significant extent for phases which commonly occur in quenched and quenched and tempered high-carbon steels. One consequence of limited resolution in the Mdssbauer patterns from multiphase steels is difficulty in accurate determination of such line parameters as position, width, and intensity. In fact, it appears that quantitative analysis for retained austenite in quenched and tempered high-carbon steels is not practical with the present experimental method. Line resolution is influenced to some extent by sensitivity. We point out below that atom or volume fractions of less than about 3 to 5 pct are not detected by the present experimental method. Thus, the presence of a multiplicity of phases does not always lead to impaired resolution. Finally, we report in this paper room-temperature MGssbauer parameters determined for austenite in a freshly quenched, commercial 1 wt pct C steel and in a freshly quenched laboratory heat of an Fe-2 wt pct C alloy. These parameters are compared with others reported in the literature. Three types of hyperfine interactions are detectable in a Mossbauer effect measurement: isomer shift, quadrupole interaction, and magnetic dipole interaction. These interactions are evidenced by one, two, and six line patterns, respectively.'-4 More than one type of interaction has been reported in certain metallurgical phases thus far studied by this method. Isomer shift is the experimentally measured displacement of line position from some arbitrarily defined reference position. In the case of a multiline pattern, isomer shift is given by the displacement of the centroid (center of gravity) of that pattern from the reference position. All isomer shifts measured at finite temperatures contain a second-order Doppler effect characteristic of that temperature. The isomer shift is related to the total s electron density at the nucleus, becoming more negative with increasing s electron density. The first-order quadrupole effect arises from the interaction between the nuclear quadrupole moment and any axially symmetric electric field gradient in

Jan 1, 1969

By A. S. Yue, A. E. Vidoz, F. W. Crossman

Tensile specimens were prepared from a single grain of an epitaxially grown Al-CuAl2 eutectic ingot. The eutectic lanzellae were oriented parallel and perpendicular to the tensile axis of the specimens. Since the composite was of the eutectic composition, the aluminum-rich matrix could dissolve up lo 5. 7 wt pct Cu in solid solution and, therefore, was amenable to strengthening by precipitation hardening. The tensile properties of the eutectic single crystals were determined at room temperature as functions of interlamel-lar spacing, platelet orientation, and thermornechanical trealment. The obserced variations in composite stress and modulus with respect to the level of&apos; composite strain are discussed in terms of premature fracture of CuAlz platelets, a distribution function for the strength of the lamellae, and unequal strains due to localized fracture of&apos; platelets. The discontinuous fiber composite model of Kelly and Tyson is modgied to account for a changing distribution of fiber lengths during composite loading. The tensile properties at elevated temperatures were determined for the direc-tionally solidified eutectic oriented with platelets parallel to the tensile axis. The observed properties are attributed to the onset of plasticity of the CuAL2 phase above 150°C. DURING the investigation of whisker- and fiber-reinforced metallic matrix composites in recent years, two major problem areas have developed: 1) The fabrication of the composite involves tedious handling techniques in order to obtain a unidirection-ally aligned and uniformly spaced set of whiskers in the metal matrix. 2) Due to weak interfacial bond strengths or because of the formation of additional embrittling phases at the metal-fiber interface during long-time exposure or fabrication at elevated temperatures, many composite systems have exhibited considerably lower strengths than those predicted by a law of mixtures analysis.&apos; These problems have been bypassed by the technique of growing whiskers and plates of high-strength materials in a ductile metal matrix by controlled unidirectional eutectic solidification.2 The tensile properties of directionally solidified A1-CuA12 eutectic are presented here. This alloy consists of a ductile aluminum matrix, containing up to 5.7 wt pct Cu in solid solution, which is amenable to precipitation hardening by heat treatment and a reinforcing high modulus CuAlz intermetallic phase. The two phases are present in the form of alternating platelets or lamellae. The microstructural stability of this unidirectionally solidified alloy at elevated temperatures has been studied extensively.3.4 and preliminary tensile and bend tests have been reported. 5-7In the present investigation the tensile properties of the A1-CuA1, eutectic have been studied as a function of several ther-momechanical variables: solidification rate. heat treatment. rolling at elevated temperatures. and lamellar orientation. It was felt that the uniformity of structure and excellent interfacial bonding would give tensile properties concomitant with the metal matrix composite theory of strengthening proposed by Kelly and coson. 8-9 The tensile properties that were obtained point to a wide distribution of strengths for the CuAlZ platelets, which leads to large deviations from the predicted mechanical behavior for this composite. EXPERIMENTAL PROCEDURE Epitaxial Growth of Eutectic Alloy. The A1-CuA1, eutectic alloy was prepared by an epitaxial growth process. Sections of a master alloy ingot (total impurity content <0.008 pct) were placed in an alundum boat, melted. and directionally solidified to obtain a multigrained plate 12 by 2 by 4 in. This plate was tapered at one end to mate with a seed crystal 1; in. long and 4 by $ in. square. Then the seed-plate combination was placed in an alundum boat which sat in a quartz tube passing through the center of a horizontal resistance wound tube furnace. A dried argon atmosphere was maintained. The temperature gradient in the furnace was such that the liquid-solid interface of the eutectic alloy was located near the end of the furnace and could be observed through the quartz tube. Single-crystal plates were formed by melting the material back to the midpoint of the seed crystal of the desired platelet orientation and then epitaxially growing the plate from the seed by withdrawing the alundum boat from the furnace at a constant rate. This technique was used to produce aluminum and CuA12 lamellae parallel and perpendicular to the transverse direction on the plate. Metallographic examination showed that both phases were continuous across the original liquid-solid interface. It was also possible to grow a plate from two seeds placed side by side: and, although the lamellae of one seed were oriented at 90 deg to those of the second seed, the interface between the two grains remained parallel to the growth direction along the entire ingot length. Maintenance of a straight intergranular boundary during the solidification process was possible as long as both seeds were oriented with their original growth direction parallel to the solidification direction of the plate. Eutectic plates were directionally solidified at rates of 0.2, 1.0. and 4.7 cm per hr and sectioned transversely to the solidification direction to determine the apparent inter lamellar spacing of the lamellae. Metallographic examination was also employed

Jan 1, 1970

By R. P. Laforce, ZJ. R. Low, A. M. Turkalo, D. F. Stein

The interaction of the crlloying eletnenls, nickel and chromium, with the impurity elements, antimony, pIzosphorus, tin, and arsenic, to producse reversible temper brittleness in a series of high-purity steels containing 0.40 wt pct C has been investigated. The alloyed steels contained approximately 3.5 pcl Ni, 1.7 pct Cr, and 0.05 to 0.08 pct of the particular irnpurity to be investigated. Susceptibility to teirlper embrittlement was measured by comparing the notched-bar transition temperature of each steel after quenching from the final temper and after very slow cooling (step cooling;) following the final temper. A plain carbon steel without alloying elements, bu/ ud/h 0.08 pel Sh, does not embrittle when step-cooled through the emzbrittling range of temperatures. The same embrittling treatment, applied to a steel with about the same antinzony content but with nickel and chvonziunz added, causes a 700°C increase in transition temperature. If chromium or nickel is the only alloying element, the increase in transition temperature is only 50%, again with antimony present. A carbon-free iron containing nickel, chromium, and antimony shou~s a 200°C shift in transition temperature for the same thermal treatment. Specific alloy-impurily interactions are also observed for the other impurity elements, phosphorus, tin, and arsenic. Additional investigations involving electron microscopy, trzicrohard-ness tests of vain boundaries, minor additions of zirconiutn and the rare earth and noble metals, nzainly with negative results, are also described. HE particular type of embrittlement investigated is that which is encountered in alloy steels tempered in the temperature range from about 350" to 525&apos;C or slowly cooled through this range of temperatures when tempered above this range. This type of embrittlement is sometimes called reversible temper brittleness to distinguish it from the embrittlement indicated by a minimum in the room-temperature V -notch Charpy energy vs tempering-temperature curve encountered in the range 28 0" to 350°C. Temper brittle-ness seriously restricts the use of many alloy steels since it precludes tempering or use in the embrittling range of temperatures and may significantly raise the ductile-brittle transition temperature of heavy-section forgings and castings tempered above the embrittling range, since such sections cannot be sufficiently rapidly cooled after tempering to avoid embrittlement. The very voluminous literature of temper brittle-ness up to about 1960 has been reviewed by woodfine&apos; and LOW.&apos; Of particular significance to the present investigation was the demonstration by Balajiva, Cook, and worn3 that high-purity Ni-Cr steel does not exhibit temper brittleness and the subsequent detailed and systematic study by Steven and Balajiva~ of the effect of impurity additions on the susceptibility to embrittlement of Ni-Cr steels. Steven and Balajiva showed that, of the impurities which may be found in commercial steels, Sb, As, P, Sn, Mn, and Si could all produce temper brittleness in a high-purity Ni-Cr steel. The principal purpose of the present investigation was to study the effects of particular alloy-impurity combinations on susceptibility to temper embrittlement. The steels used were high-purity 0.30 to 0.40 wt pct C steels containing 3.5 wt pct Ni and 1.7 wt pct Cr, separately or in combination. The susceptibility of these steels was then determined when approximately 500 ppm by weight of antimony, arsenic, phosphorus, or tin were added as an impurity. The melting, casting, and forging practices used in the preparation of the materials investigated are described in Appendix A. Table A-I in this appendix shows the analysis of all steels to be discussed. The steels were produced as 20- or 2-lb heats. The smaller heats were used after it had been demonstrated (see Appendix B) that a small, round, notched test specimen could be used to measure the shift in the ductile-brittle transition temperature caused by temper brittleness with about the same result as that obtained by Charpy testing. HEAT TREATMENT Unless otherwise noted, all steels were tested for embrittlement in the tempered martensitic condition. A typical heat treatment for a 0.40 C, 3.5 Ni, 1.7 Cr steel was: 1 hr at 870"C, in argon, quench into oil at 100"C, quench into liquid nitrogen, temper 1 hr at 625"C, and water-quench. The warm oil quench was used where quench-cracking was encountered; otherwise the initial quench was into room-temperature oil or water. For other compositions austenitizing temperatures were 50°C above Acs with the remainder of the thermal cycle the same. Steels in this condition, with no further heat treatment, are designated as non-embrittled. The above quenching and tempering cycle for the 0.40 pct C steels resulted in as-quenched hardnesses of 48 to 53 RC and as-tempered hardnesses of 24 to 31 Rc except in the case of the plain nickel or plain carbon steels. In these, the as-tempered hardness was as low as 80 to 90 Rg. No attempt was made to adjust the tempering temperature to obtain the same hardness in ali steels since it was felt that a uniform thermal cycle was more important than exactly equivalent hardness values. Pro- the standard quench and temper described above, the standard embrittling treatment was "step-cooling". For this the thermal cycle was: 593"C, 1 hr; furnace-cool to 538"C, hold 15 hr; cool to 524"C, hold 24 hr; cool to 496"C, hold 48 hr; cool to 468&apos;C, hold 72

Jan 1, 1969

By Y. C. Liu, G. A. Alers

measurements of the anisotropy in Young&apos;s modulus produced in copper by rolling 95 pct reduction in thickness below room temperature have been carried out in order to study the dependence of the texture on rolling temperature. The results clearly show the transition from a copper-type texture to a brass-type texture as the temperature of rolling is lowered. The intermediate textures observed can be described very well as a simple mixture of the two terminal textures. These results cormbined with other texture measurements make possible afresh review of the experimental facts velating to rolling textures in fee metals and, as a consequence, a critical examination of the current theories is presented. PREVIOUS experiments have shown that the transition from the copper- to the brass-type rolling texture is clearly displayed and can be quantitatively analyzed by measurements of the anisotropy of Young&apos;s modulus.&apos; Application of this method to the Cu-Zn alloy system showed that the description of the texture transition as a gradual rotation of the grains from the orientation characteristic of the copper texture to the {110}(112) texture of brass2 was inconsistent with the data. Instead, the data suggested that the texture within the transition region could be described as a simple mixture of the two terminal textures.5 Unfortunately, it was difficult to establish this point conclusively because of the inadequacy of corrections for the composition dependence of the single-crystal elastic constants. Since a rigorous establishment of the nature of this texture transition is essential to our understanding of the formation of rolling textures in fee metals, it is clearly important to undertake an investigation in which the composition dependence of the elastic constants would not enter. A suitable composition-independent texture transition is provided by the well-established variation in the rolling texture of copper with rolling temperature. This temperature-dependent texture transformation has been studied by smallman&apos; in several fee alloys and by Müller5 and others"&apos; in copper. They observed that the texture characteristic of copper rolled at room temperature changed to a brass-type texture when the rolling temperature was lowered to 77°K. Although it is not possible to decide unequivocally from the published pole figures whether or not the 77°K rolling texture of copper is entirely of the brass type,&apos; this complication does not affect the main purpose of the present investigation. In addition to establishing the nature of the texture within the transition region, the modulus data should also provide a determination of the temperature at which the transition occurs as well as the temperature range over which the transition extends. This information when combined with the modulus data on Cu-Zn alloys would then provide a considerable body of new information on textures in fee metals. Since these modulus results and the data obtained from pole-figure studies must be internally consistent, it is appropriate to compile a brief summary of the experimental observations based on all available methods rather than on the pole-figure data alone as has been done in the past. The primary purpose of such a summary would be to yield a more precise definition of the experimental facts on the rolling textures of fee metals, and thus greatly facilitate our evaluation of various proposed theories in this field. The final section of this paper is devoted to this compilation of consistent, experimental facts and their application to the various theories. EXPERIMENTAL PROCEDURE Two 18-lb ingots of cathode copper of 99.99 pct purity were induction-melted under a nitrogen atmosphere in a graphite crucible and chill-cast into a steel mold. The ingots were repeatedly cold-rolled and annealed (I hr at 500°C) into slabs about 1 1/8 in. thick. Blocks 3 1/4 in. wide, 2 1/4 in. long, and 1.000 in. thick were machined from each slab. The rolling schedule used was the same as in the previous investigation1 and the final thickness of the sheet was 0.050 in. with a rolling reduction of thickness of 95 pct instead of 97.5 pct as in the previous work.&apos; The compositions and temperatures of the cold baths used for the low-temperature rolling were as shown in Table I. After each pass the rolled strip was immediately immersed in the cold bath for about 1 min or until the bubbling of the bath had subsided. The modulus data were taken within 2 hr after the rolled strip was warmed to room temperature for the first time, so that effects due to recrystallization were minimized. The modulus specimens were in the shape of flat bars, 3 in. long, 4 in. wide, and 0.050 in. thick, cut with their long dimensions oriented at 15-deg intervals between the rolling direction and the transverse direction. The values of Young&apos;s modulus were deduced from measurements of the frequency at which these long narrow bars were set into longitudinal, resonant vibration as previously described.9 To excite the mechanical vibrations in the specimen, an electromagnetic drive similar to that employed by Thompson and lass" was used. The maximum in the amplitude of

Jan 1, 1968

By R. A. Graham, R. W. Rohde

The effect of hydrostatic pressure upon the austenite start temperature of a commercial Fe-28.4 at. pct Ni-0.5 at. pct C alloy has been determined. For pressures to 20 kbar, the austenite start temperature decreased from its atmospheric pressure value of 380°C at the rate of about 4°C per kbar. These data are analyzed by two different thermodynamic approaches; first, considering the transformation as an isothermal process, and second, considering the transformation as an isentropic process. It was found that both these approaches fit the experimental data equally well. The effect of hydrostatic pressure upon the austenite start temperature is best described by considering the mechanical work done during the transformation as that work obtained by multiplying the applied pressure with the gross volume change of the transformation. It is widely recognized1 that strain has an important effect on the initiation of martensitic transformations.* For example, the martensite start tempera- *In this paper, use of the term martensitic transformation implies the reversal of martensite to austenite as wen as the formation of martensite from austenite. ture, M,, may be increased by plastic deformation. Similarly, plastic deformation is observed to lower the austenite start temperature, A,. The effect of uniaxial stress on the M, of iron-nickel alloys has been studied by Kulin, Cohen, and Averbach.2 They found that the martensite start temperature was significantly changed by stresses well within the elastic region. Moreover, the effect of tensile and compres-sive stresses differed. These effects were explained in terms of the interaction of the applied stress with both the dilational and shear components of the transformation strain. The magnitudes of the influence of uniaxial tension, compression and hydrostatic pressure on Ms were measured in 30 pct Ni 70 pct Fe by Pate1 and Cohen.3 Their thermodynamic calculations and similar calculations by Fisher and Turnbull4 predicted the experimental results when the transformation was assumed to occur isothermally at some fixed driving force. This driving force was assumed to be supplied by a combination of the chemical free energy difference between the austenitic and martensitic phases and the work performed during transformation by the applied stress. More recently, Russell and winchel15 reported the effect of rapidly applied shear stress on the reversal of martensite to austenite in iron-nickel-carbon alloys. They performed a thermodynamic analysis of this transformation based upon the assumption that the re- versal occurred adiabatically. They concluded that the applied shear stress did not significantly interact with the transformation strain and thus did not assist in inducing the reversal. Rather they concluded that the reversal was effected by localized strain heating which resulted from the gross local shear deformation of the experiment. In either the adiabatic or isothermal analysis it is necessary to compute the work performed by the interaction of the applied stress and the transformation strains. In the case of hydrostatic pressure this interaction has been treated by two different methods. In either case the applied pressure is assumed to remain constant during the transformation. In one treatment the applied pressure is assumed to interact directly with the dilatational strain associated with the formation of an individual martensite plate.3&apos;4 This local strain has been measured at atmospheric pressure in iron-nickel alloys by Machlin and Cohen.6 In the above treatment this local strain is assumed invariant with temperature and pressure changes. In the other treatment the applied pressure is assumed to interact with the gross volume change of the transformation.7,8 The usefulness of this latter treatment has been demonstrated by Kaufman, Leyenaar, and Harvey7 who calculated the effects of pressure upon the martensite and austenite start temperatures of Fe-10 at. pct Ni and Fe-25 at. pct Ni alloys. Excellent agreement was obtained between their calculations and their experimental data on an Fe-9.5 at. pct Ni alloy. However, this treatment suffers from the fact that the data required to calculate the volume change of the transformation (i.e., the initial specific volumes, the thermal expansion and compressibility data for both the austenitic and martensitic phases) is, in general, not available for any material except pure iron. Thus the calculations of Kaufman et al.7 were necessarily performed by assuming that the volume change of the martensitic transformation in the iron-nickel alloys was that same volume change occurring during the a-? transformation in pure iron. While this approximation may suffice for very dilute alloys it is likely to be inaccurate in high nickel alloys. We have performed measurements of the effect of hydrostatic pressure to 20 kbar on the A, temperature of an Fe-28.4 at. pct Ni-0.5 at. pct C alloy. The composition is similar to the alloy used by Pate1 and Cohen3 to determine the effect of pressure upon the M, temperature. The present measurements permit calculation of the interaction between the applied pressure and the transformation strain. Additionally, measurements have been made which allow precise determination of the gross volume change of the transformation. The data allow direct comparison between the alternate hypotheses of the interaction between the applied pressure and a dilatational transformation strain characterized by either the formation

Jan 1, 1970

By Arcy A., J. G. Davis, D&apos Shock

Results of a pilot operation in the Carlsbad Basin are discussed. After hydrafracing between wells, a block of potash was removed by solution techniques. The distance between frac wells was about 200 ft, the thickness of potash mineralization, 5 ft. By proper manipulation, a feed of concentrate brine was obtained. The ex-periment showed that the thin-bedded potash could be removed by the solution techniques. The details of well construction, method of operation, and removal rates are discussed. Continental Oil Co.&apos;s laboratory research on the fundamentals of potash solution mining has been expanded by means of a series of field tests, and subjects such as well completion and hydraulic fracturing were added to the investigation. Both single-well and multi-well systems were studied in the field work. Discussion Background: The current paper discusses one field test in which potash was solution mined by a two-well system from a thin sylvinite zone. The potential economic value of solution mining evolves from (1) the use of drilled holes and solution techniques instead of excavated shafts and caverns and (2) the ability to mine both land and marine deposits through any type of overburden geology and below conventional mining depths. Recent interest has been focused on potash&apos; and other soluble minerals, such as trona. Solution extraction minerals, such as copper and uranium, are also worthy of important consideration. In addition, many of the techniques are directly applicable to the construction of horizontal underground storage carverns in salt. There are two general approaches to potash solution mining. The first is to mine on a single-well basis, in which the same well bore is used for both injection and production. This method is slow, and the areal extent may be quite limited in other than very thick ore zones. The second, and the preferred approach, is to mine on a multi-well basis in which the solvent is circulated between wells. This technique, if applied in a manner which allows the ore zone to be mined from the bottom upward, results in nearly all the solution taking place from the cavern roof. Salt removal rates, therefore, are very much higher than from a single-well system.l Wells can be interconnected into a multi-well pattern by several means. One is to join single-well caverns in the lower part of an ore zone. Another is to use the hydraulic fracturing techniques developed in the oil fields.&apos; We preferred the fracture approach because of its potential for creating the greatest area of salt exposure. Test Site Description: The field tests were conducted in New Mexico&apos;s Carlsbad Basin, where the potash deposits are flat and uniform over reasonable distances. Here, 12 potash zones are present in the massive Salado Salt section. The specific target was the Third Ore Zone which is about 4 ft thick at our location and about 1150 ft deep. The test pattern was designed in the shape of an equilateral triangle with a fourth well located in the center, 200 ft from each of the vertex wells. This configuration allowed the ore zone to be hydraulically fractured from the center well with good assurance that the fracture would intersect the bore of at least one outside well. Several multi-well test patterns would be available if the fracture connected all wells. Well Completion: Surface casing was set in the top of the Salado Salt at 600 ft to shut off water flows from the surface sands, and the salt section was drilled and diamond-cored to a point below the Third Ore Zone. A drilling fluid made of diesel oil with a small amount of emulsified water was used to drill and core the salt. This fluid was highly successful in preventing enlargement of the drilled hole and in promoting good core recovery. The three outside wells were completed by setting 51/2-in. casing at the base of a streak of anhydrite about 20 ft above the ore zone. Pipe was set high so that the intersection point of the fracture could be detected even if the fracture migrated above the ore zone as it progressed outward from the center well. The center well itself was completed by cementing 51/2-in. casing through the Third Ore Zone. Cement bond logs run on the center well have shown excellent bonding. Fracturing Practice: A mechanical tool was used to cut a notch through the casing and into the salt at a point about 1 ft below the ore zone in the center well. The purpose of this notch was to fix the point of fracture entry into the salt. The fracturing was done with water at injection rtaes as high as 30 bbl per min. The salt parted at 1450 psi; and it required only 5 min for the fracture to reach the well which was 200 ft to the south. It took about 5 min more to reach the other two wells. Caliper surveys were run to locate the point of fracture entry in the three outside wells. The fracture appears to have drifted downward slightly, entering the outside wells at the top of a streak of carnallite 8 or 9 ft below the ore zone. A cross section of the wells selected for the multi-well test is shown in Fig. 1. The figure includes KC1 values based on core analysis and the trace of the fracture plane between the wells.

Jan 1, 1971

By C. A. Stickels, C. M. Yen

Secondary recrystallization was investigated in vacuum-melted electrolytic iron to which 70 pm N was vacuum-meltedadded. The secondary texture is "near {554}<225>" for material cold-rolled 75 to 90 pct, the sharpness of the texture increasing with increased rolling reduction and with decreased annealing temperature. At reductions of 95 and 97.5 pct the secondary texture is &apos;"near {322)(296)". Both secondary orientations also exist as major components of the primary re-crystallization texture. Development of a strong "near {554) (225)" secondary texture appears to depend on the evolution of the Primary texture to a transition texture depleted in orientations near the secondary orientation before the onset of secondary growth. A variety of qualitative experinzents have been used to show that nitrogen is important in limiling primary grain growth. The presence of nitrogen does not seem essential for the establishment of a transition texture, but a loss of nitrogen during annealing may facilitate growth of grains in the secondary orientation. Secondary grains we shown to form initially at the specimen surface. This is not thought to indicate that surface energies are important in the growth process. It is proposed that the quasi-two-dimensional character of surface grains permits discontinuous growth parallel to the surface before secondary growth of interior grains is possible. An earlier study of recrystallization textures in 90 pct cold-rolled electrolytic iron showed that secondary recrystallization occurred after annealing for several days at 700C1 This type of secondary recrystallization, which had not been reported previously, results in the formation of a strong texture, best described by the indices "near {554}(225)". The purpose of the present work was to investigate the effect of various processing variables on secondary recrystallization in this material and determine the mechanism of secondary grain growth. LITERATURE REVIEW An understanding of the mechanism of a secondary recrystallization process depends on knowing: 1) how grains in the secondary orientation come to be in the primary recrystallization texture; 2) why normal grain growth does not occur; and 3) what factors determine the strength of the secondary texture. For secondary growth of grains of a particular orientation, a certain minimum fraction of the grains must be in that orientation after primary recrystallization. This requirement is apparently satisfied "naturally" in certain systems, i.e., when the primary texture obtained by rolling and recrystallizing material initially randomly oriented contains a sufficient fraction of primaries in the secondary orientation. However, in other cases, e.g., {110}<001> and {100}<001> secondary growth in silicon iron,2 it is necessary to enhance the fraction of primary grains in the secondary orientation by rolling and recrystallizing textured material. In the present case, the "near {554}<225>" orientation is contained within the spread of orientations found in the primary recrystallization texture of iron or bbc iron-base alloys. In systems where the main driving force for secondary growth is the reduction in total grain boundary energy, secondary growth is observed only when normal grain growth is minimized. Four ways in which normal grain growth can be limited are: 1) Limitation by a strong primary texture. When a very strong primary texture consisting of a single component or twin-related components develop, most primary grains are separated from one another by relatively immobile small-angle grain boundaries. The classic instance of this is secondary growth into the primary cube texture in some fcc metals. 2) Limitation by precipitates. Precipitates present in the proper volume fraction with a suitable dispersion will limit primary grain growth. The role of MnS inclusions in impeding normal grain growth in Si-Fe is well-documented.5 3) Limitation by sheet thickness. Normal grain growth slows drastically when the mean grain diameter is of the same order as the sheet thickness. This effect has been used to obtain secondary recrystallization in thin sheets of high-purity silicon iron.&apos; 4) Limitation by solute impurities. It is well-established that certain impurity elements in solution can have a large effect on grain boundary mobility.&apos; However, there does not seem to be any secondary recrystallization process in which primary grain size stabilization has been shown to be due to the drag exerted on grain boundaries by dissolved impurities. In certain systems, e.g., secondary recrystallization in silver,&apos; the means by which normal grain growth is limited has not been identified, and solute-impurity limitation might be suspected. In order to understand the factors which determine secondary texture strength in three-dimensional growth, it is necessary to examine in more detail the current picture of general secondary recrystallization processes. Following Cahn,9 it is assumed that the primary grains have a range of sizes and that secondary growth of one of the large grains in this distribution is possible when it exceeds a critical size with respect to its neighboring grains. The critical size depends on the ratio ?S/?p, where ?s is some sort of average grain boundary energy between the potential secondary and the primary grains and ?p is some sort of average grain boundary energy between primary grains. For a constant primary grain size, the critical size for secondary growth increases as ?$/?p increases. May and Turnbull5 have incorporated the

Jan 1, 1969

By C. S. Sivaramakrishnan, Manjit Singh, Rajendra Kumar

Regarding liquid nzetals structurally as a suspm-sion of clusters , having derivated solid-state coordination, in truly liquid atoms, the recently developed Kuvlar-Samarin technique of centrifuging- in liquid state enabled the determination of the cluster sizes inAl-Cu mid Pb-Sn systems. It is shown that the colutne fraction of the clusters does not exceed 9 pct and the energy of their formation in Al-Cu is about 5.5 kcal per g-atom and in Pb-Sn eutectic alloys about 25 kcal per y-atoni. STRUCTURAL investigations of liquid state have principally followed the following three courses: i) studies with X-ray, electron, or neutron diffraction; these investigations have shown that there is a certain amount of regularity in the structure of liquid metals which can be defined by a coordination number and that the structure is a derived function of that in the solid state; ii) thermodynamical investigations which are based on the concept of ideal behavior; these describe the liquid state in terms of free-energy values and other thermodynamic functions; although these investigations are of help in the study of the general effects of alloying, they do not provide any structural insight into the precise atomic distribution in liquid state; iii) measurements of surface tension and viscosity; although it is natural to expect that the viscosity is related to the structure in liquid state, these investigations have so far only provided information which can be used by the foundry technologists and has been little utilized in formulating models of the structure of liquid state. As it happens, investigators in the three groups have worked almost independently of each other and there is practically no structural correlation between the results of one group with those of another. The purpose of the present paper is to indicate that the experimentally measured parameters of these three groups of research are closely related to the structure in the liquid state. STRUCTURE OF LIQUID METALS Although atomic distribution in solid and gaseous states is rigorously known, that of the liquid state is only appreciated on the fringes. There is no universal model of atomic distribution in liquid state, but two diverse models are at present hotly contested. The first, largely expounded by ~ildebrand,&apos; regards the liquids as condensed gas since many of their properties and much of their behavior can be adequately described by regarding them as fluids. The second mode12j3 considers that some form of near-solid as- sociation of large number of atoms exists in the liquid state. On the other hand, ~ernal&apos; was able to predict rather precisely the radial distribution functions in liquids on the basis of statistical geometric approach which considered that liquids are "homogeneous, coherent, and essentially irregular assemblage of molecules containing no crystalline regions or holes large enough to admit another molecule". He introduced the concept of pseudonuclei in the otherwise random structure as aggregates of closely packed tetrahedra which gradually merge into irregularity and continually replace each other. To what extent the pseudonuclei can be regarded as regions of near-solid association is indefinite but Bernal suggested that the concept of pseudonuclei can be compromised with the latter model if the near-solid associations are regarded as extremely dense and not necessarily crystalline. The difficulty in projecting the structure of liquid metals arises because they exhibit duplicity of character as some of their properties are closer to those of crystalline solids and others to fluids. There is an increasing tendency to discuss the structure of liquid metals in terms of the second concept according to which the structure of liquid metals may be conceived as consisting of i) clusters of atoms where the aggregation is a close derivative of that in the crystalline state, ii) individual atoms which behave like true liquids in respect to degrees of freedom and iii) excess number of vacancies. It is noteworthy that the introduction of only 5 pct vacancies is sufficient to transform crystalline matter into the liquid state. At any instant of time thermodynamic equilibrium exists between i, ii, and iii, but the relative proportion of the clusters and random atoms is not known. That this is so can be appreciated by the fact that, when liquid metal is rapidly cooled, liquid state vacancies may condense in the form of dislocation loops and vacancies in excess of their equilibrium number in solid state. These dislocation loops have been observed in thin foils of aluminum prepared from rapid cooled aluminum. As temperature increases above melting point the number and volume fraction of clusters decrease but those of vacancies and random atoms increase. Clusters are transient in nature. In pure metals the cluster is an aggregation of the metal itself. In alloys, however, the nature of the cluster largely depends on the interaction between solvent and solute atoms. If the interaction between unlike atoms is greater than between like atoms: the clusters are then aggregates of unlike atoms. Examples of this kind of system are A1-Cu, Mg-Pb, and so forth, i.e., systems which exhibit negative departures from the Raoult&apos;s law. In systems where the interaction between unlike atoms is smaller than be-

Jan 1, 1968

By J. E. Dutrizac, R. J. C. MacDonald, T. R. lngraham

When sintered disks of synthetic chalcopyrite (CuFeS2) were leached in acidified aqueous solutions of ferric sulfate, the following reaction stoichiometry was obtained: CuFeS2 + 2Fe2(SO4)3 = CuSO4 + 5FeSO4 + 2S Over the temperature range from 50º to 94ºC, the reaction displayed parabolic kinetics. The parabolic rate constant for the dissolution of copper is given by the equation: log.k(mg2/cm4-hr)= 11.850 - 3780/T The activation energy for the dissolution process is 17 ± 3 kcal per mole. The parabolic kinetics have been attributed to the progressive thickening of a sulfur film on the surface of the chalcopyrite. When the leaching solutions contain less than 0.01 molar Fe+3 , the Fe concentration influences the rate of leaching, probably through a mechanism involving the diffusion of ferric sulfate through the sulfur layer. At higher Fe+3 concentrations, the rate control in the leaching. reaction has been attributed to the diffusion of ferrous sulfate through the sulfur. The rate of reaction is insensitive to changes in acid concentration and in disk rotation speed. ThE reaction of acidic ferric sulfate solutions with various sulfide minerals is of practical interest for both bacterial and heap leaching. This leaching medium is generally used with low-grade ores that cannot be treated profitably by conventional means. In both bacterial leaching1-3 and heap leaching, the active agent for sulfide dissolution is ferric sulfate. Although the reactions of ferric sulfate with chalcocite, covellite, and bornite have been investigated,4*7 the kinetics of leaching chalcopyrite with ferric sulfate have not been thoroughly studied. This paper reports a study of that reaction. EXPERIMENTAL Reagent-grade sulfur was purified by the method of Bacon and FanelliB and then it was vacuum-distilled to remove any soluble magnesium salts that had been introduced during the purification procedure.9 From stoichiometric quantities of the purified sulfur and hydrogen-reduced electrolytic copper sheet (99.90 pct Cu), CuS was synthesized at 450°C in a vacuum-sealed, pyrex vessel. About 24 hr was required for the completion of the reaction. A similar procedure involving hydrogen-reduced iron wire (99.90 pct Fe) was used to synthesize FeS1.002. A 2-furnace arrangement was required. The iron was heated to 800°C while the sulfur was maintained at about 400°C. Although the reaction to consume the sulfur was rapid, the material required additional heating (1 week) in a sealed silica ampoule at 800°C before it was homogenized. X-ray powder diffraction analysis confirmed that the copper sulfide was covellite and that the iron sulfide was troilite. The composition of the iron mineral was confirmed by wet chemical analysis. The two sulfides were ground to minus 100 mesh, weighed in equimolar amounts, mixed thoroughly, and pressed into pellets at 80,000 psi. The pellets were vacuum-sealed in pyrex ampoules and then sintered for 3 days at 550°C after an initial heating at 450°C for a few hours. The pellets were then cooled, polished with 3/0 emery paper, rinsed in acetone, and stored. The material had the characteristic brassy color of chalcopyrite and was shown by X-ray diffraction to be CuFeS2. Microscopic examination of the polished surfaces revealed small inclusions of pyrite (approximately 0.5 vol pct) as the only impurity. The presence of small amounts of a second iron compound will not alter the amount of dissolved copper but might increase the amount of ferrous ion slightly. It was calculated that dissolution of all of the pyrite and 100 mg of Cu (a typical value) would change the expected ferrous concentration by only 4 pct. Microscopic examination of a pellet after leaching revealed that the pyrite was not preferentially solubilized; only those pyrite grains at the surface were attacked. Hence, the pyrite is unlikely to alter the rate of copper dissolution. The chalcopyrite disks were about 1.7 mm thick and 27 mm in diam. They were about 80 pct of theoretical density, and for this reason their true reaction area was somewhat larger than the 5.8 sq cm area presented by the polished face. The disks were cemented to lucite cylinders in such a way that only the polished face was exposed. The disks were then leached by methods previously described.6,7 RESULTS AND DISCUSSION Stoichiometry and Kinetics. The initial experiments were directed to the problem of resolving the stoichiometry of the leaching reaction. Disks of CuFeS2 were leached at 80°C for various periods of time in acidified ferric sulfate solutions that were protected from oxidation by a cover of flowing nitrogen. When the disks had been partly leached, they were removed, their soluble salts were washed out, and then they were treated with CS2 in a Soxhlet extraction apparatus. The ratio of elemental sulfur to dissolved copper thus obtained was approximately 2 to 1. After the extraction of elemental sulfur from the pellet, the residue consisted of unreacted chalcopyrite only. For runs in which an appreciable amount of copper was dissolved, the ratio of ferrous ion to cupric ion in the solution was

Jan 1, 1970

By Donald E. Meyer

High current-low voltage EB-gun evaporation in an oil-free ultra-high vacuum system was found to be necessary, though not sufficient, for stability (300°C, 106 v per on) of aluminium gate MOSFET&apos;s and MOS capacitors not stabilized by a phosphorous glaze. five characteristics of the equipment used: 1) Vacuum purification of the aluminum charge, 2) Ionization of the evaporant by the electron beam, 3) X-ray formation, 4) Residual gases during evaporation, and 5) Metal film structure were studied as Possibly significant in MOS fabrication. EVAPORATION of contact metals common to the semiconductor industry historically has been accomplished with oil diffusion pump systems and various resistance heated evaporant sources as dictated by the type of metal evaporated. To meet a need for greater reliability of semiconductor devices, other metallization methods were developed. A good example would be application of the moly-gold contact system to integrated circuits with deposition by RF or triode sputtering.&apos; More recently, fabrication of stable metal-oxide-silicon devices and circuits has put new demands on metallization. The purity of the thin metal films composing MOS structures is critical, particularly at the metal-oxide interface, and ultra-high vacuum metallization using sputter-ion pumping and electron beam gun (EB-gun) evaporation are well suited for the task. At this laboratory aluminum has been the most common contact-gate metal for both MOS capacitors and MOSFET&apos;s. In the earliest work with MOS capacitors, aluminum was evaporated from wetted tungsten filaments using both diffusion pump and ion pump vacuum systems. In spite of clean oxide techniques these capacitors were unstable under bias-tempera-ture stressing. Only after a switch to EB evaporation of aluminum were stable capacitors produced. Using the same techniques it was possible to make MOSFET&apos;s with equivalent stability. Stability data for a discrete MOSFET is shown in Fig. 1. This is a "clean" oxide gate (no phosphorus stabilization or no etch back of a thicker gate) having a thickness of lOOO? thermally grown on the (111) plane. Gate length after diffusion was 0.24 mils, and the devices were hermetically sealed. Stressing conditions were 300°C and 106 v per cm applied alternately as a positive and negative field for 10 min, 50 min, and 4 hr for a total stress time of 10 hr. An initial shift in turn-on voltage of 0.1 v was detected for 10 min of positive bias. All evidence at this laboratory indicated that while EB-gun evaporation of ultra-high purity aluminum was not sufficient for 300°C stability, it did seem to be necessary. There may well then be something inherent in the EB-gun deposition used which enhanced stability, and probably no single factor existed but rather a series of factors. It is the purpose of this paper to report on some of the investigations carried out to learn more about EB-gun evaporation in ultra-high vacuum systems. EXPERIMENTAL DESCRIPTION The EB-gun was self accelerated, had a maximum power rating of 10 kw, and used a water-cooled copper crucible able to hold a 20-g aluminum charge. The electron beam was bent 180 deg and focused by an electromagnet which also provided movement of the beam across the crucible. Normal power conditions in this work were 9 kv and 300 to 600 mamp. The gun can be described as high-cur rent/low-voltage and was quite different in its mechanism of operation from EB-guns with much higher acceleration potentials. An oil-free vacuum system capable of 5 x 10- l0torr, a quartz crystal rate and thickness monitor and a quadruple mass spectrometer completed the evaporation system, Fig. 2. A typical evaporation cycle consisted of a 3 to 4 hr pumpdown to the upper l0-9 range and evaporation at l0? per sec with the pressure in the bell jar not rising above 1 x 10"7 torr. Thickness control was 5 pct or less and could be automatically monitored and controlled. Five phenomena associated with the EB evaporation and considered as possible contributors to Ma performance included a purification effect, ionization of evaporating aluminum, X-rays, constitution of vacuum ambient during evaporation, and film structure dependence upon evaporation rate. These phenomena are now discussed. Vacuum Purification. The design of the EB-gun permitted purification of the aluminum charge by vacuum outgassing. Particular features included an efficiently water-cooled copper hearth with a capacity of over 20 g of aluminum and the capability for sweeping the beam across the charge. Such capacity meant that aluminum had to be added only after about every fifth evaporation. A new charge was not required each evaporation as is necessary with filament evaporation. An oxide "scum" which appeared on the charge could be completely cleared from the top hemisphere of the charge by sweeping with the beam prior to opening the shutter. An indication of the purifying effect was obtained by a series of analytical measurements on incoming aluminum, after melting but with little vacuum out-gassing, after 30 min outgassing, and the evaporated film itself. Either a solids (spark source) mass spectrometer or an emission spectrometer were used for analyzing the aluminum charge. Analysis of the evapo-

Jan 1, 1970

By Ursula E. Wolff

The precipitation of carbides from an alloy containing 33 pct Ni, 21 pct Cr, balance iron, was investigated electron microscopically by means of extraction replicas and thinned metal foils. Annealing temperatures ranged from 565°to 870°C and up to several thousand hours. M23C6 precipitated in pain boundaries, incoherent and coherent twin boundaries in that sequence. The orientation relationship between carbides and austenite matrix was determined and correlated with the morphology of the carbides and with the type of boundary in which precipitation occurred. In large-angle grain boundaries, as well as in coherent twin boundaries, the carbides had the same orientation as one of the adjacent pains. These carbides formed sheets of individual flakes with shapes related to the orientation of the boundary. In incoherent twin boundaries carbides precipitated in ribbons composed of pavallel rods. An unidentified subcarbide was found to precede precipitation of M23C6 in these boundaries. The M 23 C6 rods had a kind of fiber texture with (110) parallel to the long dimension of the rods and ribbon, and with orientations of both of the adjacent twin-related austenite crystals Predominant in the texture of the carbide. A hard sphere crystal model has been used to discuss orientation and morphology of the carbides in terms of free volume and vacancies available in the boundaries. A number of papers have dealt with the morphology of chromium carbide (M23 C6) precipitated in austenitic stainless steels.1"7 In all these investigations, the carbides were examined in the electron microscope by means of extraction replicas. With this technique, the carbides retain the spatial distribution they had in the bulk sample. However, since the matrix is dissolved in the process, the particles can turn in an unpredictable way; and the orientation relationship between matrix and carbides cannot be established. In this paper the results of studies on extraction replicas and on thinned metal foils are reported. These studies were undertaken to determine the matrix-to-car bide orientation relationship, and to correlate the orientation of the carbides with their morphology. PROCEDURE The material used was an austenitic alloy with 33 pct Ni, 21 pct Cr, balance iron, containing approximately 0.05 pct C. Coupons of 1.25-mm sheet were first solution-annealed at 1050°C for 15 min and air-cooled. Then, to precipitate the carbides, samples were isothermally annealed in the range from 565" to 870°C for times up to several thousand hours. All further specimen-preparation procedures were carried out after the final anneal. Carbon extraction replicas from polished and etched surfaces were made with 10 pct bromine in methyl alcohol.&apos; Thin foils were prepared from punched-out 3-mm-diam disksg which fit into the electron-microscope holder. The disks were prethinned by grinding to approximately 0.5 mm thickness, and then electro-polished in a polytetrafluoroethylene holder1&apos; with a solution containing 5 pct perchloric acid in acetic acid to which 10 g per 1 Cro3 and 5 g per 1 nickel chloride were added (etchant modified from that of Briers et al."). This solution dissolves neither the carbides nor the austenite around the carbides preferentially. By using extraction replicas, electron micrographs and selected-area electron-diffraction patterns were taken from the same carbide arrays. By using thin foils, electron micrographs were made from a grain boundary area containing carbides. Electron-diffraction patterns were then taken from the same area and from each of the adjacent grains separately. In this manner, the orientation of each grain could be determined without interference by the carbide pattern. A peculiarity of extraction replicas should be pointed out. After the matrix is etched away, the carbide arrays float freely in the etching and washing solutions, and are held in place only at the anchoring points in the carbon replica. When the replica is picked up with a screen the carbide arrays tend to flip to one side. Thus, while the surface features are preserved, the original arrangement of the carbides may severely and unpredictably be disturbed whenever the specimen contains large amounts of interconnected carbides. Nevertheless, it is possible to correlate the different morphologies of the carbides with the type of boundary in which they have precipitated. RESULTS 1) Extraction Replicas. Fig. 1 shows that the grain boundaries usually are curved, multicornered surfaces of random orientation. The coherent twin boundaries (which are (111) planes) cut a grain into parallel slices. Incoherent twin boundaries occur at the ends and on the steps of twins and are often narrow, parallel-sided strips which are much longer than they are wide. Different morphologies can clearly be distinguished for the M23Ce carbides precipitated in each of these types of boundaries, and agree well with those observed by kinzel.2 The kinetics of this precipitation has been investigated." The first carbides precipitate in junctions of three grain boundaries and fan out from there into the adjoining boundary surfaces, Fig. 2(a). These carbides are oriented randomly, Fig. 2(b), and become coarser and thicker as annealing time increases. The large-angle grain boundaries are next to fill

Jan 1, 1967