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Jan 1, 1967

Jan 1, 1967

By Richard H. Bube

Photoelectronic radiation detectors may be conveniently classified as homogeneous intrinsic, homogeneom extrinsic, or junction type. Highly photosensitive homogeneous intrinsic photodetectors may be prepared from a number of different II-VI, III-V, or IV materials. Such materials require the presence of specific imperfections that can act as sensitizing centers to provide long majority carrier lifetime. Homogeneous extrinsic photodetectors are of interest primarity for infrared detection, and consist principully of germaniuim and silicon with suitable inpurities. A vaviety of junction photodetectors exist, with silicon us the most common material for them all. The following extrema in performance are found: 1) highest photoconductivity gain in homogeneous intrinsic photodetectors; 2) smallest response time (highest frequency response) in p-i-n junctions; 3) largest gain-bandwidth product in avalanche diodes. ALTHOUGH the total number of materials exhibiting photoconductivity effects is very large, only a relatively few of these have appropriate properties for a practical photoelectronic radiation detector. In fact if one surveys the commercial detectors currently available, one finds that the field is dominated by some ten or fewer materials. These are summarized in Table 1 in terms of the other important variable in detector fabrication, the structure of the material, i.e., whether the material is used as a single crystal or in polycrystalline form, and whether the material is used as a homogeneous detector or whether the detector depends upon the existence of a junction or barrier in the material. The principal wavelength range for each type of material is also shown, together with an indication of the utilization of intrinsic or extrinsic excitation. For the purposes of the present comparison of various detectors, it will be convenient to discuss two main topics: homogeneous photodetectors and junction photodetectors. The performance of a photoelectronic radiation detector is measured in terms of two parameters: the photoconductivity gain, and the response time of the detector. A convenient figure-of-merit is given by the ratio of gain to response time, often called the gain-bandwidth product. The photoconductivity gain is a device parameter since it varies in many cases with the applied voltage and the detector geometry, and should be distinguished from the actual photosensitivity of the material involved. This photosensitivity can be conveniently given as the product of free carrier lifetime and mobility in the material. The photoconductivity gain is defined as the number of charge carriers which pass between the electrodes per second for each photon absorbed per second, where G is the photoconductivity gain, is the pho-tocurrent in amperes, e is the electronic charge in coulombs, and F is the total number of electron-hole pairs created in the photo conductor per second by the absorption of light. The gain may also be expressed as the ratio of the lifetime of a free carrier to the transit time for that carrier, i.e., the time required for the carrier to move between the electrodes. For a material in which one-carrier conductivity dominates, where T is the lifetime of a free carrier, ttr is the transit time for this carrier, p is the carrier mobil-ity, V is the applied voltage, and L is the electrode spacing. From Eq. [2] it follows that the photoconductivity gain is proportional to the photosensitivity of the material (tP), the proportionality constant being The photosensitivity of a material, and hence the photoconductivity gain of a device utilizing the material, depends on the lifetime of the free carriers as a critical parameter. In a homogeneous material this lifetime is determined by the nature of imperfections in the material, and in an inhomogeneous material the lifetime is determined by the specific junction structure. The speed of response, the other basic photodetec-tor parameter, is determined by factors quite similar to those important for photoconductivity gain. Imperfections are of primary importance in homogeneous materials, and the structure of the junction is a determining factor in junction devices. For infrared detectors it has become common to define another quantity designed to indicate directly how effective the detector is in distinguishing between a small photoconductivity signal and random noise due to the detector and its environment. The detectivity D* is given by

Jan 1, 1968

By A. H. Larson, L. G. Twidwell

DURING the roasting of zinc sulfide concentrates and by-product recovery from flue dusts, the zinc sul-fate phase may occur in various stages of the processes. In addition to zinc sulfide, other sulfide phases are usually present such as CdS, PbS, CuzS, and FeS. Consequently, the effect of these sulfides on the thermal decomposition of ZnS04 was investigated. Merck high-purity, anhydrous zinc sulfate was used as the starting material. The sulfate was heated under an argon atmosphere, at a rate of 10°C per min. With the first slight indication of decomposition as determined by gas analysis, the sample was removed from the furnace and allowed to cool. X-ray patterns showed no other phase present. It is estimated that less than 0.05 pct of the ZnS04 decomposed. The metal sulfides were also heated under an argon atmosphere to drive off any adsorbed water. The materials were then sized (-200 to +270 mesh), mixed in predetermined mole ratios, and pressed into pellets. The experimental measurement of the reaction rate consisted of purifying the argon carrier gas, continuous flushing of the heated reacting mixture with argon, and collection and analysis of the SOn product gas. The reaction system was heated to 600°C and al-

Jan 1, 1967

By Fred H. Perfect

The thermite process, now approximately 80 years old, has heretofore been practiced in ceramic vessels. This paper outlines the successful production of several million pounds of alloy produced in zuater-cooled copper vessels. THE high temperatures necessary to produce fluid high-alumina slags have probably discouraged earlier workers from attempting to utilize water-cooled vessels. The following is a report of the successful application of this method, together with a report of the physical and chemical problems which had to be solved to make the process practical. About 1955, when Reading Alloys was more concerned with the production of ferrotungsten than with most other types of alloys, it was found to be exceedingly difficult to predict the reaction rate that one would experience from one tungsten ore to another . This difficulty was increased for two reasons: the physical state of the ore as received was seldom uniform, and the size of the mass to be reduced is a very important factor in developing the proper mix for the final reaction rate and temperature. For example, the reduction of ferric oxide with aluminum would theoretically produce temperatures of about 2600°C, and, in spite of the losses due to radiation, the energy is quite sufficient to produce a very fluid slag and iron effecting a very clean separation. On the other hand, the reduction of ferrous oxide with aluminum, while it is not as energetic as the first reaction, nevertheless should give temperatures above 2000"C, which should cause separation of alumina and iron, but, in practice, this is not the case and good fluid slag is not obtained. The heat of reaction, calculated from the thermo-dynamic heats of formation of the constituents of the reaction, offers a convenient guide to the feasibility of a thermite reaction. Thus, the values listed in Table I can be used as an approximate guide for estimating the results which may be expected in the reduction of various oxides and combinations of oxides using aluminum as the reducing agent. As a generalization, those reactions releasing approximately 600 cal per g of total mass represent the lower limit at which to expect separation of slag and metal. On the other hand, those reactions which release 1100 to 1200 cal per g are too violent to be used successfully without the addition of considerable inert material, such as lime, fluorspar, and so forth. Reaction rate is also a very important factor. This, of course, is governed very largely by the sizing and intermingling of the oxidant and reductant and by the starting temperature of the reactants. This is why those figures shown in Table I can only be a general guide. Further, it must be understood, that by adding powerful oxidizing agents, such as nitrates, chlorates or metallic oxides, such as molybdenum trioxide or chromium trioxide, certain problems respecting low temperatures are not necessarily solved. This is true in the addition of oxidants, such as sodium chlorate. By adding too high a concentration of sodium chlorate as a booster for a particular reduction, explosions are generated. Generally, they are of an intermittent type so as to give a Roman Candle effect to the reduction itself; with each surge resulting from a minor explosion more and more material is thrown

Jan 1, 1968

Jan 1, 1968

By S. Kobayashi, A. E. Armenákas, C. Mylonas

Earlier work1-4 has shown that commercial mild steels under static loading at the lowest natural operating temperatures fracture in a brittle manner only when damaged by a suitable history of straining. Notched and then compressed plates have fractured in subsequent tension at loads as low as 10 pct of the limit load and precompressed smooth bars at a subsequent extensional strain as low as 0.01. The comparison of the average net fracture stress with the flow limit stress was shown to be an excellent criterion of brittle or ductile behavior of mild steel structures when only loads and general stress levels are known. The present investigation shows that the deformation at fracture is a far more sensitive measure of brittleness than the fracture stress. Both criteria, deformation and stress, are used to determine the amount of uniform precompression of two mild steels, ABS-B and Project E-steel, resulting in brittle fracture under the strong constraint of a subsequently machined severe circumferential groove. The fracture stress equaled or exceeded the theoretical flow limit Ól = 2.680., ('based on the elevated yield stress at each prestrain) up to prestrains of about 0.20 and then fell off to about 1 at prestrains of 0.60. Yet a prestrain of only 0.05 reduced the elongation at the shoulders by a factor of about 4. The total plastic elongation of a region surrounding a sharp notch in prestrained steel deternines whether or not fracture will be initiated in large structures, hence is a direct and realistic measure of the remaining ductility and provides an excellent test of the material's resistance to embrittlement. ALMOST all engineering structures of mild steel which failed in service in a so-called "brittle" manner exhibited appreciable local plastic deformation (more than 1 pct) and a mixture of "cleavage" and "shear" fracture surfaces. Furthermore, the peak local stress in an advancing crack is always high. Such simple criteria as pure cleavage, absence of plasticity, or fracture stress, so useful in other fields, are insufficient to characterize brittleness in engineering structures at ordnary temperatures. The "brittle)' behavior of engineering structures is caused by the limited ductility at cracks or notch regions, as shown by recent work at Brown University summarized and extended in Refs. 1 to 4. It was also shown that the local ductility of mild steel may be catastrophically reduced by a suitable strain and temperature history similar to those which may occur in real structures. Localized yielding begins at the notch roots at low loads but is contained by surrounding elastic regions. The plastic strains are hence small. They increase slowly with the load up to the flow limit or limit load for an ideally plastic material. Unrestricted plastic flow then occurs. At such strains the real material locally st rain-hardens and fractures. With work-hardening materials no flow limit exists, and the transition from low to high plastic strains is gradual but increases more rapidly at loads close to the flow limit of an equivalent perfectly plastic material. When the available ductility is sufficient for reaching the limit load, the behavior is normal or "ductile". With insufficient ductility, fracture will occur at a low load and will be defined as "brittle". Accordingly the sufficiency or not of the ductility at a notch is reflected in the magnitude of the fracture load or average net stress as compared with the limit load or flow limit. The application of this criterion gave surprising results. Service failures under semistatic loading occurred mostly below the limit load, hence were "brittle". On the contrary all static tests with undamaged commercial mild steels reached the limit load in spite of the deepest notches and temperatures below Charpy impact transition (say about -30°F or higher). Notched bars showed no brittleness even under fully dynamic axial loading up to 1.5 x107 psi per sec at -23°F and about 2.5 x l06 psi per sec at -ll0°F, but broke in a brittle manner at -200°F and l03 psi per sec.' To propagate a fracture at low load, experimenters had to trigger it by superimposing a strong dynamic impact at a notch cooled by liquid nitrogen as in the Robertson5 or Esso tests.' Once started the cracks would run through warmer regions of smaller stress. Obviously the ductility of the apparently undamaged commercial steels was "sufficient" in the laboratory tests but not in the service fracture, where it must have been reduced by some embrittling procedure of fabrication or service. These conclusions were confirmed by the achievement of low static stress (brittle) fracture initiation in unwelded steel after a local reduction of ductility by simple prestraining. Symmetrically notched plates of undamaged mild steel cooled below the Charpy V-notch transition range were tested in central static tension: Undamaged plates withstood loads of limit intensity, but when subjected to prior in-plane com-pressive prestraining perpendicular to the notch axis and to accelerated aging, the plates fractured completely or partially at very low static loads, as low as one-tenth of the flow limit.2-4 However, damage could not be easily related to the strongly variable prestrain around a sharp notch. To obtain easier strain measurements, it was decided to use axially precom-pressed bars and bent bars. Final testing was done by

Jan 1, 1970

By K. A. Gschneidner, O. D. McMasters

DgIel-ential thermal, nretallographic, and X-ray paramzetric methods were used to establish the Yb-Pb phase dingram. The terminal solid solubilities in the system are less than 0.2 at. pct. Lead additions lower the 792°C transformation tewzperature of pure ytterbiun~ at 787°C by means of an inverted peritectic veaction. Ytterbium additions raise the 327oC melting point of lead to 329°C by means of a perilectic reaction. Three intermetallic compounds Yb2Pb, YbPb, and YbPb, melt congruerttly at 1246", 1116o, and 740°C. respectively. YbsPb3 forms by a peritectic reacliorz at 1250°C. YbzPb crystallizes in the PbCh (C23) type structure and the orthorhombic latliceo constants are a, = 7.478. bo = 5.225, aud C0 = 9.549A. YbsPbs crystallizes in the hfnsSi3 (D8a) type structure. The hexa~nal lattice constants for this phase are a, = 9.325 and co = 6.92d. YbPbl crystallizes in the cubic AuCU3 (Ll2) type structure with a: = 4.8628 + 0.0008A. YbPb is diwrorphic with the structure change occurring al 507°C upon healing. The crystal structures of the two forttzs were not determined. Eulectic reactions occur at 5.1 at. pct Pb and 769oC, 41.5 at. pct Pb and 1086oC, and G8.0 at. pct Ph and 717°C: RELIABLE information concerning the alloying behavior of most of the rare-earth metals is scarce. A systematic study of the rare earth-lead alloy systems now in progress is an effort to supply such information. Determination of the phase diagram of the Yb-Pb alloys is the first step of this study. Variations in the valence of normally divalent ytterbium may be studied by comparing the characteristics of the Yb-Pb alloys to those of the trivalent rare earth-lead and of the divalent alkaline-earth-lead systems. Of particular interest are the melting temperatures of the intermetal-lic compounds since these are known to be high relative to the melting points of the constituent metals. Crys-tallographic and magnetic susceptibility data for the rare earth-lead alloys are to be included in the comparative study. Results from this and subsequent investigations should provide a better understanding of alloy formation. EXPERIMENTAL PROCEDURE Materials. The major impurities in the ytterbium and lead used in this investigation are listed in Table I. The lead was obtained from Cominco Products, Inc., and was specified to be 99.99 pct pure. The ytterbium was prepared in this laboratory by the lanthanum reduction of the oxide followed by a distillation of the ytterbium. Alloy Preparation. The alloys were prepared by

Jan 1, 1968

Jan 1, 1967

By P. O. Metz, E. L. Wiehe, J. S. Foster

BASED on their analyses of growth kinetics, ibbs' and Harris et al.' have recently proposed that the rate of subcritical graphitization of steel is controlled by the rate of diffusion of carbon in ferrite. illert has concluded that this is the rate controlling growth mechanism when the growing graphite particles are relatively large. Harris et al.' also propose that the matrix surrounding a growing graphite particle is plastically deformed to make room for the growing particle. If this plastic deformation does occur, then preferred recrystallization may be observed around the larger particles. If diffusion of carbon in ferrite is rate controlling, then a graphite particle growing in a matrix of, say, tempered martensite or pearlite, should be surrounded by a carbide depleted zone. Thus, the simple metallographic experiments reported here are sufficient to determine whether or not diffusion of carbon in ferrite is the dominant mechanism for the growth of graphite in steel. The present investigation considers graphitization in five alloys. The compositions are given in Table I. Graphitization was carried out at three temperatures: 900°K (627"C), 950°K (677"C), and 990°K ('717°C). Samples were reacted in both the martensitic and the pearlitic initial conditions. In the hypoeutectoid steels, Sets 7 and 8, .and in the hypereutectoid Fe-C binary, Set 6, graphitization was very limited after 1000 hr at temperature. In these samples graphite particles observed were less than 10 p in diam. In the hypereutectoid commercial steels, Sets 1 and 3, large graphite particles (100 p) did not appear until after 250 hr at temperature. The limited number of graphite particles observed and the range of sizes obtained prohibited measurement of growth rates in this investigation. However, in all samples showing graphitization, neithe a carbide depleted zone nor recrystallization was observed. illert's model predicts that diffusion of carbon in ferrite should become dominant for graphite particles larger than about 300 p in diam. The largest graphite particle observed in this investigation was approximately 450 p in diam. No evidence of carbide depletion was noted. Fig. 1 shows a typical graphite particle. This nodule is from an initially pearlitic sample from Set 3 after

Jan 1, 1970

By D. G. Westlake

Levine has concluded that prismatic slip in titanium is controlled by two distinct thermally activated processes in the temperature ranges O° to 220°K and 220" to 300°K. We feel that his evidence is not convincing. He has drawn a curve through the plot of activation volume u* vs temperature T with a break at 22WK, Fig. 2. We believe that considering the possible experimental error apd the fact that there are only four points between 4° and 196°K, one could be justified to represent these points by the smooth curve we show as a solid line in Fig. 12. Levine believes that (8rf/aT)+ is discontinuous at 220°K, but he hints that he recognizes the possible doubt. We doubt seriously that the results shown in Fig. 1 indicate a discontinuity at 220°K. Certainly if he were justified in drawing a straight line through only four points on Fig. 2, we are even more justified in drawing straight lines through the seven points between 77" and 297°K on Figs. 1 and 3, because a straight line touches all seven of his circles. If his measurements of t were so precise that his decision to draw two curves in Fig. 3 is valid, then the experimental technique is deserving of further mention. We have drawn a smooth curve through the results in Fig. 3 and used this and the curve in Fig. 12 to obtain the values of shown in Fig. 13. Finding continuous is a contradiction of Fig. 4, which deserves comment itself. Levine obtained the solid circles in Fig. 4 using Fig. 2 and a smooth curve through the points shown in Fig. 3. We have replotted these solid circles in Fig. 14. We submit that there is no basis for his two straight lines (dashed in Fig. 14), and therefore have superimposed the curve that we believe is indicated. Certainly, there is reasonable doubt that a discontinuity in dH/dT exists at 220°K, even when one accepts Fig, 2. Furthermore, if one draws the straight lines indicated by least-squares treatments of the results presented as open circles in Fig. 4, their intersection is at about 196°K, not at 220°K as shown. Thus the alleged agreement between Figs. 3 and 4 with respect to "break" temperatures is not apparent. In Fig. 9, Levine has drawn a straight line very much different from the one indicated by a least-squares treatment. Obviously, he has done this because he recognized that H = 0 when T = 0. He is consistent in drawing the low-temperature curves in

Jan 1, 1968

By C. N. J. Wagner, E. N. Aqua

The bee alloys of the terminal solid solution of rhenium in niobium were investigated by X-ray diffraclion methods. The analysis of the broadening of the powder pattern peaks from the niobium-rich alloys, cold-worked by filing, showed evidence for faulting on (211) planes. The alloys with larger rhenium concentrations were brittle, and fractured when filed. The analysis of the shapes of the diffraction profiles from powders produced by hammering with steel plates) of these brittle alloys indicated that the broadening was due predominantly to platelets bounded by clibe planes. This result was consistent with the Laue patterns from single-cleazlage facets which showed evidence for clearage fracture on {100} planes. The influence of faulting on the mechanical properties of face-centered cubic (fcc) metals and alloys, i.e., how faulting affects the various dislocation configurations in fcc metals, is well-known.' The important question is whether or not faulting occurs in body-centered cubic (bcc) metals and alloys, and, furthermore, if faulting does occur, to what extent it affects the mechanical properties. Theoretical calculations indicate that the stacking-fault energy is rather high in most bcc metals.2 The results of sev-eral studies using transmission electron microscopy have indicated that faults are present in well-annealed and lightly deformed bcc metals and alloys, such as tungsten,3 vanadium,4 and Mo-Re alloys.5,6 These faults, as have been observed, may be stabilized by the presence of interstitial (C,N,O) atom segregation to the extended dislocations. In addition, there is considerable X-ray diffraction evidence for faulting in bee metals,7-8 using the same techniques that were successful for the study of faulting in fee metals and alloys.7 This X-ray diffraction method is currently the only technique applicable for the investigation of faulting in metals with high stacking fault energies, i.e., when severe amounts of deformation are required to produce sufficient amounts of faulting to be detectable by X-rays. The thinned films of heavily deformed metals will become opaque in the electron microscope due to the high dislocation densities present. In a previous investigation of faulting8 in bee metals of group Vb and VIb, the authors observed that niobium was ductile when filed at room temperature, and showed the highest density of faults on the (112) family of planes. A niobium-base alloy system was therefore chosen to study the effect of alloying on the occurrence of faults in a bcc alloy. There is a large range of terminal solid solubility of rhenium in the alloys of the bee refractory metals. All of the Group VIb-rhenium alloys, i.e., Cr-Re, Mo-Re, and W-Re, exhibit solid solution hardening without loss of ductility.9 In fact, the ductility reaches a maximum for alloys with compositions near to the limits of solubility. This enhanced ductility has been attributed to the lowering of the stacking-fault energy in the alloy; i.e., the increased amounts of twinning provide an additional mode of deformation, resulting in the observed high ductility.' Because of this demonstrated effect of rhenium in bee alloys, it was decided to study a series of alloys of Nb-Re. Using the same X-ray diffraction methods as described in the previous investigation,' one is able to identify the possible causes of broadening of the powder pattern peaks, e.q., particle size, microstrains, stacking faults, and/or twin faults. The separation of these various contributions to the measured peak shape is accomplished by the Warren-Averbach analysis.7 I) EXPERIMENTAL PROCEDURE A series of seven bee solid solution alloys containing from 2.5 to 21.5 at. pct Re (remainder niobium) were prepared at the General Electric Research Laboratory. Buttons, weighing 25 g, were alloyed from electron-beam-melted niobium chips and random rhenium sheet by arc melting in a water-cooled copper crucible, under an argon atmosphere. The buttons were inverted and completely remelted to promote homogenization. The nominal alloy compositions are listed in Table I. Samples of the cold-worked alloys were obtained by filing to produce a powder for the X-ray study. Alloys too brittle to be filed were hammered with steel plates. The resultant filings or brittle impactings were screened through 150-mesh classifiers and compacted to the required shape for the diffractometer holder. The X-ray diffraction peaks were automatically recorded by fixed time counting at uniform 28 intervals, or by continuous registration with a ratemeter. In all cases, the (110)-(220) and (200)-(400) pairs of peaks and the (211), (310), (222), and (321) reflections were measured on a GE-XRD5 diffractometer using CuKa radiation with nickel filter or MoKct radiation and a LiF monochromator in the diffracted beam. The first part of the data reduction, completed with the aid of a series of programmed computations,10 was the analytical resolution of the Kol component of the diffraction profile. Subsequent computations were then made with reference only to the resolved Kol peak shape. These included the measurement of the

Jan 1, 1967

By R. W. Bartlett

DISILICIDE coatings, MeSi2, on refractory metals are usually grown by a solid-state diffusion reaction similar to the parabolic oxidation of metals. Two or more silicide compounds occur in each of the binary silicon-refractory metal systems. When an adequate source of silicon is available during the coating process, the lower silicide compounds are usually not formed.1 However, during service application of the coated refractory metal, continued heating causes further diffusion and growth of the lower silicides. A study of the growth kinetics of the lower silicides in the molybdenum and tungsten binary systems with silicon was previously reported.2 The present study is concerned with the growth kinetics of Ta5Si3 and Cb5Si3, formed as conjugate layers between TaSi2/Ta and CbSi2/Cb diffusion couples, respectively. Diffusion couples consisting of disilicide-coated tantalum and columbium sheet coupons were subjected to diffusion heat treatments, cross-sectioned, and scanned by an electron-microprobe analyzer to determine the relative X-ray intensity profiles. These profiles correspond closely to concentration profiles for the constituents, and were used to identify phases and locate phase boundaries. The analysis of a single specimen after a diffusion-anneal provides only one value for the Me5Si3 phase thickness, corresponding to the time and temperature of the diffusion treatment. Consequently, each isothermal curve showing the rate of accumulation or depletion of silicide phases required several samples with identical diffusion temperatures but different diffusion times. The disilicide coatings were prepared by pyrolizing commercially pure silane in a hydrogen carrier gas. The hydrogen was purified by passage through a palladium membrane. Commercially pure tantalum and columbium sheet metal coupons, 0.005 by $ by 2 in., were used. All of the coated specimens for each system were prepared together in a single group in order to minimize coating variations among individual specimens. During pyrolysis, the brightness temperature was maintained at 1200°C. With suitable flow rates at this temperature, only the disilicides of tantalum and columbium were formed in detectable amounts. Coated specimens were subjected to diffusion an-neals in a tube furnace with automatic proportional control within ±10°C. A H2-Ar mixture, deoxidized and dried over anhydrous magnesium perchlorate, was used to prevent oxidation during the diffusion anneal. Diffusion couples were cross-sectioned, polished and examined with an A.R.L. electron-microprobe X-ray analyzer. Both metal constituents in each specimen were simultaneously scanned using a 1-p beam spot size and either 1- or 2-µ steps. Analyses were performed using the Ka radiation of columbium and silicon, and the La, radiation of tantalum. For all specimens, an acceleration voltage of 30 kv was used, and integration was made against the specimen current. LiF and A DP monochrometers were used to isolate characteristic radiations. Three or more traverses were made on each specimen. The intensity scales were adjusted using the constituent metals and standard samples of the silicide compounds which had been verified by X-ray diffraction. The relative intensity profiles were used directly, in conjunction with the standard samples, to identify compounds and phase boundaries. The stoichiometries differ considerably between the silicide compounds in both systems, and the homogeneity range for these silicide compounds is narrow. These characteristics result in marked differences in relative intensities of the fluorescing elements—differences which permit ready identification of phases and phase boundaries. Projections of both electron back-scattered images and characteristic

Jan 1, 1967

By R. D. Pehlke, J. V. Gluck

AS a portion of a general study of liquid-metal thermodynamics, high-temperature galvanic cells were used to study the effect of added dilute solutes on the activity of zinc in dilute solution with molten bismuth in the range 450° to 650°C. 1,2 The cells were of the type The experimental procedures have been described previously.1 In the course of studies conducted in the bismuth-rich corner of the systems Bi-Zn-Au and Bi-Zn-Cu, marked discontinuities were observed in the electromotive force-temperature relations for a number of the experimental compositions. Above the discontinuity temperatures, the relations, which were linear, exhibited slopes similar to the binary alloy of the same mole fraction zinc, as shown in Figs. I and 2. Below the discontinuity temperatures, the slopes were markedly different, but the relations remained linear. In general, as the amount of the third eletnent was increased, the slope of the lower portion of the relations tended to become less positive. The slopes of electromotive force-temperature relations reflect the entropy of mixing of zinc and, therefore, the occurrence of a discontinuity marks the transition between two different equilibrium situations involving the solution of zinc. A number of other investigators3-5 have used such observations to determine liquidus boundaries in binary alloys. Since the upper portions of the ternary electromotive force-temperature relations of this study had the same slopes as the binary alloys, the discontinuities are interpreted as marking the transition between the single-phase liquid region and the underlying two- and three-phase regions. Reference to the Zn-Bi equilibrium diagram6 confirms that the binary alloys were single-phase liquids at the experimental temperatures. Fig. I shows the electromotive force-temperature relations for the alloys containing Bi-Zn-Au and Fig. 2 shows similar relations for the Bi-Zn-Cu alloys. In the lower portion of each figure, the temperature of the discontinuity is plotted vs the mole fraction of the added solute at the various levels of zinc concentra- tion. Cross-plots from these curves at constant temperature permitted the construction of the partial liquidus surface isotherms shown in Figs. 3 and 4.

Jan 1, 1967

By C. G. Dunn

James, Jones, and Leak (JJL) conclude that growth-rate and orientation-distribution data obtained in a steep thermal gradient should be used with caution to account for isothermally produced recrystallization textures. I agree. However, it is not clear whether they conclude a) that the orientation dependence of rate of growth of secondaries in (commercial) silicon iron depends on both temperature and a temperature gradient or b) that, in applying rate data obtained in a steep thermal gradient, the isothermal annealing temperature must be close to the temperature for which the rate data apply. Also, it is not clear what tem-

Jan 1, 1967

Jan 1, 1968

By V. O. Struebing, K. A. Johnson, F. H. Ellinger, C. C. Land

The system Pu-Ce-(Pu,Ce)Co2 has been investigated by thermal, micrographic, and X-ray diffraction methods. The principal characteristics of this system were found to be: 1) a "eutectic" valley extending across the phase diagram from the 6 Pu/Pu3 Co eutectic at 408°C to the 7 Ce/Ce24Co11 eutectic at 42Z°C, the maximum temperature in the valley being 443°C; 2) three ternary inter mediate phases, each having an appreciable mutual solid solubility for plutonium and cerium; 3) eleven four-phase equilibria comprising one eutectic, four peritectics, and six Class II reactions; 4) a pseudobinary section between the ternary phase A and the y Ce solid solution, and another such section between the 6 Pu solid solution and the 77 phase (Pu3 CO) containing cerium in solid solution. ALLOYS of plutonium that are completely molten at temperatures in the vicinity of 500°C are of interest for use in the liquid state as fast reactor fuels. Since exploratory work1,2 on Pu-Ce-Co alloys had indicated that a "eutectic" valley extended across the ternary system between the Pu-co[6 Pu/n (Pu3co)] and the Ce-CO[? Ce/? (ce24co11)] binary eutectics, thus allowing alloys of any plutonium concentration to be selected within the temperature limitation, the present investigation was begun for the purpose of accurately defining the "eutectic" valley. However, when the existence of three ternary intermediate phases and the consequent complexity of the liquid-solid equilibria were discovered, the investigation was expanded to cover the complete system, except for the corner between the isomorphous series PuCo2-CeCo2 and pure cobalt. 1) EXPERIMENTAL The purities of the alloy components are given in Table I. Alloy buttons, each weighing about 5 g, were prepared by arc melting the accurately weighed components (±0.1 mg) under a purified He-Ar atmosphere. Weight loss as a result of melting was usually less than 1 part in 1000; therefore, the buttons were not chemically analyzed. One half of each button was re-melted by induction heating in vacuo in a magnesia crucible and furnace-cooled, so as to produce a somewhat coarser microstructure than that of the rapidly cooled arc-melted button. About 20 g of each alloy to be examined by thermal analysis were remelted in vacuo in a tantalum crucible of the inverted thermocouple well type. The thermal, X-ray diffraction, heat treating, and micrographic methods used in this investigation were essentially the same as have been described in earlier reports.3,4 Two etchants for these alloys were found to be satisfactory: A) 8 (parts by volume) orthophosphoric acid, 5 ethanol, and 5 glycol; B) 7 tetraphosphoric acid, 36 water, and 57 2-ethoxyethanol. Etchant A was used at 5 to 6 v for 15 to 20 sec, and etchant B was used at 10 v for 10 sec. 2) THE LIMITING BINARY SYSTEMS 2.1) Plutonium-Cerium. The phase relationships of this system have been investigated by Ellinger, Land, and Cramer,5 and by Selle and Etter.6 The over-all agreement of their phase diagrams is good, but there are some discrepancies in the positions of some phase

Jan 1, 1967

By R. G. Aspden, D. A. Colling

THE {120}(001) texture has been reported in near 50 pct Ni-Fe alloys <0.004 in. thick1,2 and in copper 0.014 in. thick.3 This texture is one of two types of secondary recrystallization textures observed to replace a cube texture matrix and is desirable in Fe-Ni alloys because of its unique magnetic characteristics.1,2 However, the conditions for and mechanism of the {120}(001) formation are not well understood.3 During an investigation of the influence of alloying elements on the magnetic characteristics of 49 pct Ni-Fe alloys with a sheet thickness of 0.014 in., we have observed the growth of {120}(001) secondaries when strong deoxidizers such as aluminum and calcium were present. It is believed that this observation in unique and provides additional insight with regard to the mechanism for {120}(001) texture formation. Three alloys were prepared with electrolytic iron and nickel and contained a nominal 49 wt pct Ni. Alloy A was a straight binary Fe-Ni alloy, whereas intentional additions of 0.013 wt pct Ca to alloy B and of 0.59 wt pct A1 to alloy C were made. Ingots were forged at 1100 C, hot-rolled to 0.150 in. gage, pickled, and cold-rolled to 0.014 in. Standard U-shaped magnetic test specimens were punched, deburred, sprayed with MgO, and annealed in ~50°C dew point hydrogen. Textures were identified with the aid of etch pits, the

Jan 1, 1970

By K. Lundholm, M. Andersson, L. Nevatalo, Å. Holmström, L. Hedström

"Several copper deposits have the arsenic bearing minerals enargite (Cu3AsS4) and tennantite (Cu12As4S12). Since this mineral contain both arsenic and copper, it is not possible to remove the arsenic by selective flotation. Partial roasting is a method to clean the concentrate from arsenic and thereby produce a suitable raw material for the copper smelter. Also mercury is commonly found in copper bearing minerals. In the partial roasting process both arsenic and mercury are volatilized, while part of the sulphur is used as energy during the roasting. This gives a calcine suitable as feed to a smelter, much like an ordinary copper concentrate. A typical analysis of calcine would be: S content 18–22%, As content below 0.3%, Cu content – not affected, grain size – not affected. However, if the concentrate contains mercury, the mercury will be volatilized during the roasting process. This will lead to very high mercury contents in the subsequent process of sulphuric acid, unless the SO2-rich process gas is pre-treated by the Boliden/Norzink process that recovers mercury as calomel, Hg2Cl2. The volatilized arsenic in the gas phase can be recovered either by dry or wet methods. However, one problem is that arsenic also contaminates the fine grained calcine dust that is captured downstream in the gas cleaning system. The calcine dust contains valuable metals such as copper, gold and silver that needs to be separated from the arsenic. The ultimate aim is to stabilize the recovered arsenic as a compound with minimum impact on the environment and recover the valuable metals. By using wet recovery methods it is possible to selectively precipitate copper, gold and silver as sulphides while arsenic is kept dissolved in the liquid. The precipitated metal sulphides can be mixed with the calcine and sent to the copper smelter. The dissolved arsenic is sent to effluent treatment and is further treated in Outotec ferric arsenate precipitation process, where arsenic is precipitated as the stable compound basic ferric arsenate, formula FeAsO4·xFe(OH)3. All concentrates containing arsenic are unique and the process design of a roaster plant must be based on relevant process data. The process data is collected during tests in lab scale and pilot scale, where the influence of critical process parameters on the quality of calcine is studied. The ability to run tests both in lab scale and pilot scale is crucial to the development of the correct roasting process. Outotec is well equipped for this."

Jan 1, 2014

By K. Barr, F. Grenier, J. Hettinger

"To reach the Hemlo Gold Inc. objective of becoming the lowest cost, safest Canadian gold producer while causing no environmental damage, management embarked from its inception on a strategy to build an organization which would foster employee involvement wherever possible. This paper describes the organizational structure and how various innovative concepts have been implemented.BACKGROUNDThe Golden Giant Mine in the Hemlo area of North Western Ontario is owned by Hemlo Gold Mines Inc. 50.1% is controlled by Noranda. The mine commenced operations in March 1985 at a modest production rate. After it reaches full production by the end of 1988, with completion of underground development work, the mine is expected to operate at 3000 t/d and produce in excess of 300,000 ounces of gold per annum. This makes it one of the three largest producing gold mines in North America.Two other companies, Lac Minerals and Teck-Corona, operate mines on either side of the Golden Giant Mine. Noranda Geco Division is located 55 kilometers to the north at Manitouwadge. The majority of Hemlo Gold's 270 employees live in Manitouwadge where a mix of accommodation was constructed and made available for purchase by employees. Bus transportation is provided to and from the workplace with modern coaches. The housing and transportation policies were implemented prior to start-up and designed to help attract and maintain good employees.As shown in Figure 1, the operation is divided into five operating and support departments with each reflecting a high degree of departmental control and responsibility. For example, the mill department is responsible for all aspects directly relating to the milling complex including safety, operation, mechanical maintenance, assaying, instrumentation, and environmental control. Similarly, the mine department is responsible for all phases of operation, mine engineering, geology, and maintenance of underground mobile equipment. The plant department is responsible for maintenance of hoisting plants and conveyances, underground stationary equipment, all site electrical services, surface buildings, housing and yard maintenance. Efforts are made at every level of the organization, from departmental organization to individual job assignments, to ensure responsibilities are clear and that they encompass complete functions."

Jan 1, 1988